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 A Light Metal Systems Part 3 Selected Systems from Al-Fe-V to Al-Ni-Zr Editors G. Effenberg and S. Ilyenko
Authors Materials Science and International Team, MSIT®
ISSN 1615-2018 (Physical Chemistry) ISBN-10: 3-540-25013-1 Springer Berlin Heidelberg New York ISBN-13: 9783540-25013-5 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/11A3: 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 2005 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
SPIN: 10915998
63/3020 - 5 4 3 2 1 0 – Printed on acid-free paper
Editors:
Günter Effenberg Svitlana Ilyenko
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 authored by MSIT® in which data and knowledge are contributed by many individuals and accumulated over almost twenty years. Authors for the evaluations in this volume are: Zoya Alekseeva, Moscow, Russia
K.C. Hari Kumar, Chennai, India
Sergiy Balanetskyy, Jülich, Germany
Vasyl Kublii, Kyiv, Ukraine
Christian Bätzner, Stuttgart, Germany
Viktor Kuznetsov, Moscow, Russia
Georg Beuers, Hanau, Germany
Hans Leo Lukas, Stuttgart, Germany
Natalia Bochvar, Moscow, Russia
Evgeniya Lysova, Moscow, Russia
Oksana Bodak, L’viv, Ukraine
Pierre Perrot, Lille, France
Marina Bulanova, Kyiv, Ukraine Gabriele Cacciamani, Genova, Italy
Alexander Pisch, Grenoble, France Dmitriy Petrov†, Moscow, Russia
Nirupan Chakraborti, Kanpur, India
Qingsheng Ran, Stuttgart, Germany
Tatyana Dobatkina, Moscow, Russia
Peter Rogl, Wien, Austria
Oleksandr Dovbenko, Kyiv, Ukraine
Lazar L. Rokhlin, Moscow, Russia
Olga Fabrichnaya, Stuttgart, Germany
Rainer Schmid-Fetzer, Clausthal-Zellerfeld, Germany
Gautam Ghosh, Evanston, USA
Elena Semenova, Kyiv, Ukraine
Joachim Gröbner, Clausthal-Zellerfeld, Germany Jean-Claude Tedenac, Montpellier, France Benjamin Grushko, Jülich, Germany
Ludmila Tretyachenko, Kyiv, Ukraine
Frederick H. Hayes, Manchester, UK
Vasyl Tomashik, Kyiv, Ukraine
Volodymyr Ivanchenko, Kyiv, Ukraine
Volodymyr Turkevich, Kyiv, Ukraine
Konstyantyn Kornienko, Kyiv, Ukraine
Tamara Velikanova, Kyiv, Ukraine Andy Watson, Leeds, UK
Institutions The content of this volume is produced by Materials Science International Services GmbH and its international team of materials scientists, MSIT®. Contributions to this volume have ben made from the following institutions: The Baikov Institute of Metallurgy, Academy of Sciences, Moscow, Russia
Materials Science International Services GmbH, Stuttgart, Germany
Degussa AG, Hanau, Germany
Max-Planck-Institut für Metallforschung, Institut für Werkstoffwissenschaft, Pulvermetallurgisches Laboratorium, Stuttgart, Germany
ENSEEG, Laboratoire de Thermodynamique et Physico-Chimie Metallurgiques, Domaine Universitaire Saint Martin d’Heres, Cedex, France 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 Indian Institute of Technology Madras, Department of Metallurgical Engineering, Chennai, India Indian Institute of Technology Department of Metallurgical Engineering, Kanpur, India Institute for Semiconductor Physics, National Academy of Sciences, Kyiv, Ukraine Institute for Superhard Materials, National Academy of Sciences, Kyiv, Ukraine G.V. Kurdyumov Institute for Metal Physics, National Academy of Sciences, Kyiv, Ukraine Laboratorie de Physico-chimie de la Materiere Universite de Montpellier II, Montpellier, France
Moscow State University, Chemical Faculty, Moscow, Russia National University of L’viv, Kathedra of Inorganic Chemistry, L’viv, Ukraine Northwestern University, Department of Materials Science and Engineering, Evanston, USA Technische Universität Clausthal, Metallurgisches Zentrum, Clausthal-Zellerfeld, Germany Universita di Genova, Dipartimento di Chimica, Genova, Italy Universite de Lille I, Laboratoire de Métallurgie Physique, Villeneuve d’ASCQ, Cedex, France Universität Wien, Institut für Physikalische Chemie, Wien, Austria University of Leeds, Department of Materials, School of Process, Environmental and Materials Engineering, Leeds, UK
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. The first four volumes contain evaluation reports on selected ternary systems of importance to industrial light alloy development and systems which gained in the recent years otherwise scientific interest in the area of light metal systems. In a ternary materials system, however, one may find alloys for various applications, not only light alloys, 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 which works together since almost 20 years, now. 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 calculations for individual equilibria, driving forces or complete phase diagram sections. 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 a materials system best. The conclusions on the 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 mass-spectrometry, thermo-gravimetry, measurement of electro-motive forces, Xray 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 and the interpretation of the results with regard to the material’s chemistry has to be verified. 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 ternary AB-X systems. Therefore combining systematically binary and ternary evaluations can lead to a new level of 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 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 evaluations and on many different specific questions, adding the competence of a team 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. Special features of the compendium and the standard format are explained in the Introduction to the volumes. 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 an as useful tool 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 I want to thank all those who contributed their comments and insight during the evaluation process. In particular we thank the reviewers. Their names are as follows: Pierre Perrot, Hans Leo Lukas, Hari Kumar, Peter Rogl, Gabriele Cacciamani, Riccardo Ferro, Lazar Rokhlin. We all gratefully acknowledge the dedicated work of the editorial team: Dr. Oleksandra Berezhnytska, Dr. Larisa Plashnitsa, Mrs. Irina Korolkova, Ms. Natalya Bronska.
Günter Effenberg and Svitlana Ilyenko
Stuttgart, June 2004
Contents IV/11A3 Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data Subvolume A Part 3
Light Metal Systems
Selected Systems from Al-Fe-V to Al-Ni-Zr
Introduction Data Covered . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . XI General . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . XI Structure of a System Report . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . XI Literature Data . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . XI Binary Systems . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . XI Solid Phases . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . XII Pseudobinary 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 Aluminium – Iron – Vanadium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1 Aluminium – Iron – Yttrium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11 Aluminium – Iron – Zinc . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21 Aluminium – Iron – Zirconium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 42 Aluminium – Germanium – Lithium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 52 Aluminium – Hydrogen – Lithium. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 58 Aluminium – Hydrogen – Magnesium. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 64 Aluminium – Hydrogen – Titanium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 71 Aluminium – Hafnium – Nickel. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80 Aluminium – Lithium – Magnesium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93 Aluminium – Lithium – Silicon . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 109 Aluminium – Lithium – Zinc . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 124 Aluminium – Lithium – Zirconium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 134 Aluminium – Magnesium – Manganese. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 142 Aluminium – Magnesium – Nickel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 150 Aluminium – Magnesium – Scandium. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 157 Aluminium – Magnesium – Silicon . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 165 Aluminium – Magnesium – Tin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 178
Aluminium – Magnesium – Titanium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 187 Aluminium – Magnesium – Zinc . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 191 Aluminium – Magnesium – Zirconium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 210 Aluminium – Manganese – Palladium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 215 Aluminium – Manganese – Titanium. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 253 Aluminium – Molybdenum – Nickel . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 266 Aluminium – Molybdenum – Titanium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 287 Aluminium – Nitrogen – Silicon . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 317 Aluminium – Nitrogen – Titanium. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 322 Aluminium – Niobium – Titanium. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 334 Aluminium – Niobium – Zirconium. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 380 Aluminium – Nickel – Ruthenium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 389 Aluminium – Nickel – Silicon . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 400 Aluminium – Nickel – Tantalum . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 425 Aluminium – Nickel – Tungsten . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 440 Aluminium – Nickel – Zirconium . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 451
CD-ROM providing interactive access to the system reports of this volume
Introduction
XI
Introduction 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 selfsufficient. 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: Literature Data 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.
Landolt-Börnstein New Series IV/11A2
MSIT®
XII
Introduction
Heading Literature Data Binary Systems Solid Phases Pseudobinary Systems Invariant Equilibria Text
Liquidus, Solidus, Solvus Surfaces Isothermal Sections Temperature-Composition Sections Thermodynamics Notes on Materials Properties and Applications Miscellaneous
References Miscellaneous Notes on Materials Properties and Applications Thermodynamics Temperature-Composition Sections Tables and diagrams
Isothermal Sections Liquidus, Solidus, Solvus Surfaces Invariant Equilibria Pseudobinary Systems Solid Phases Binary Systems
Fig. 1: Structure of a system report
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., J, J' – 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.
MSIT®
Landolt-Börnstein New Series IV/11A2
Introduction
XIII
Pseudobinary Systems Pseudobinary (quasibinary) sections describe equilibria and can be read in the same way as binary diagrams. The notation used in pseudobinary 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) For transition reactions the letter U (Übergangsreaktion) is used in order to reserve the 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). 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-pseudobinary 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. 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.
Landolt-Börnstein New Series IV/11A2
MSIT®
MSIT®
Ag-Tl
144 e9 (Tl)(h) Tl3Bi+(Tl)(r)
192 e8 l Tl3Bi+Tl2Bi3
202 e7 l (Bi)+Tl2Bi3
303 e1 l (Tl)(h)+Tl3Bi
Tl-Bi
294 e2 (max) L (Ag) + Tl3Bi
Ag-Tl-Bi
144 (Tl)(h) Tl3Bi + (Tl)(r),(Ag)
equation of eutectoid reaction at 144°C
(Ag)+(Tl)(r)+Tl3Bi
E2
D1
(Ag)+Tl3Bi+Tl2Bi3
188 L (Ag)+Tl3Bi+Tl2Bi3
(Ag)+(Bi)+Tl2Bi3
197 L (Ag)+(Bi)+Tl2Bi3
207 e6 (max) L (Ag) + Tl2Bi3
(Ag) + (Tl)(h) + Tl3Bi
E1
ternary maximum
289 L + Tl3Bi (Ag) + (Tl)(h) U1 289 e4 (min) L (Ag) + (Tl)(h)
first binary eutectic reaction (highest temperature)
Figure 2: Typical reaction scheme
234 d1 (Tl)(h) (Tl)(r),(Ag)
291 e3 l (Ag)+(Tl)(h)
second binary eutectic reaction
261 e5 l (Ag) + (Bi)
Bi-Ag
second ternary eutectic reaction
monovariant equilibrium stable down to low temperatures
reaction temperature of 261°C
XIV Introduction
Landolt-Börnstein New Series IV/11A2
Introduction
XV
C
Data / Grid: at.% Axes: at.%
δ
p1
700
20
80
500°C isotherm, temperature is usualy in °C primary γ -crystallization
γ
40
400°C
300
estimated 400°C isotherm
e2
U
e1
40
300
300
400
α
0 40
80
β (h)
E
50 0
60
liquidus groove to decreasing temperatures
60
0 40
binary invariant reaction ternary invariant reaction
50 0
0 70
20
limit of known region
20
A
40
60
80
B
Fig. 3: Hypothetical liquidus surface showing notation employed
C
Data / Grid: mass% Axes: mass%
phase field notation estimated phase boundary
20
γ
80
γ +β (h)
40
phase boundary
60
three phase field (partially estimated) experimental points (occasionally reported)
L+γ 60
40
tie line
L+γ +β (h)
β (h)
L
80
L+β (h)
L+α
20
limit of known region
α
Al
20
40
60
80
B
Fig. 4: Hypothetical isothermal section showing notation employed Landolt-Börnstein New Series IV/11A2
MSIT®
XVI
Introduction
750
phase field notation
Temperature, °C
L 500
L+β (h)
L+α
concentration of abscissa element
32.5%
250
β (h)
L+α +β (h)
temperature, °C β (h) - high temperature modification β (r) - room temperature modification β (r) alloy composition in at.%
188
α α +β (h) 0
A B C
80.00 0.00 20.00
60
40
Al, at.%
20
A B C
0.00 80.00 20.00
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 crossreferences. 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.
MSIT®
Landolt-Börnstein New Series IV/11A2
Introduction
XVII
General References [C.A.] [Curr.Cont.] [E] [G] [H] [L-B]
[Mas] [Mas2] [P] [S] [V-C] [V-C2]
Landolt-Börnstein New Series IV/11A2
Chemical Abstarts - 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)
MSIT®
Al–Fe–V
1
Aluminium – Iron – Vanadium Gautam Ghosh Literature Data [1960Gup] studied the effect of additions of Al on the stability of the )-phase (FeV). They prepared a number of alloys, using electrolytic grade elements, in an induction furnace under He atmosphere. The alloys were homogenized at 1175°C for 72 h. Metallographic observations and X-ray diffraction were performed to identify the phases. [1987Sok] and [1988Sok] reported the phase equilibria in the Al-rich ternary alloys containing up to about 50 at.% Fe. The alloys were prepared using the metals of following purity: 99.95 mass% Al, 99.95 mass% Fe and electrolytic V. A number of alloys were prepared by arc melting under Ar followed by homogenization at 500°C in evacuated silica capsules. The V-rich alloys (0 to 75 at.% Al) were heat treated for 1800 h at 1000°C followed by 600 h at 500°C, whereas Al-rich alloys (75 to 100 at.% Al) were annealed at 500°C for 1430 h [1988Sok]. The phase analysis was performed by means of microstructural, thermal analysis, microhardness and X-ray diffraction techniques. Apart from conventional casting, a number of ternary alloys were also subjected to rapid solidification by melt-spinning which were subsequently annealed at 250 and 450°C for 50 h. An additional rapidly quenched alloy was investigated by Mössbauer spectroscopy [1989Sok]. These results were assessed by [1992Gho] and [1992Rag]. Recent experimental results are primarily related to phase separations [1989Zha, 1989Koz, 1993Miy, 1994Koz] and bcc-based ordering in Fe-rich alloys [1983Bus, 1985Okp, 1995Ant, 1997Nis1, 1997Nis2, 2001Nis1, 2001Nis2]. An update summarizing some of these results has been reported by [2002Rag]. Binary Systems The Al-Fe and Al-V binary phase diagrams are accepted from [2003Pis] and [2003Sch], respectively. The Al-Fe phase diagram has undergone slight modification due to recently established congruent melting behavior of the Fe4Al13 phase [1986Len]. The Fe-V phase diagram is accepted from [1982Kub], which has also been adopted in [Mas]. Solid Phases The maximum equilibrium solid solubilities of V and Fe in (Al) are about 0.3 at.% at 660.4°C [1989Mur] and 0.03 at.% at 652°C [1982Kub], respectively. However, by rapid solidification, the corresponding solid solubilities can be enhanced up to about 1.25 at.% V and 4.4 at.% Fe [1976Mon] and in the ternary regime, the solid solubility can be up to 0.5 at.% V and 2 at.% Fe [1987Sok]. The lattice parameter of supersaturated (Al) containing about 4.4 at.% Fe is about 401.2 pm [1976Mon]. Also, the lattice parameter of (Al) decreases linearly to 404.2 pm at 1.2 at.% V [1976Mon]. The substitution of Fe by V in Fe3Al increases both the D03 (Fe3Al) B2 (FeAl) and the B2 (FeAl) A2 (Fe) transition temperatures [1969Bul]. Recently, the effect of V on the D03 B2 ordering of Fe 3Al has been determined by several investigators [1997Nis1, 1997Nis2, 2001Nis1, 2001Nis2]. These results are summarized in Fig. 1. The D03 B2 temperatures reported by [1969Bul] differ significantly from those of Nishino and co-workers, as a result the data of [1969Bul] are not considered in Fig. 1. Along the Fe3Al-V3Al section, solid solutions (Fe1-xV x)3Al have been prepared [2003Kaw1]. The D03 lattice of Fe3Al (x = 0) has three sublattices labeled Al (4 sites), FeI (4 sites) and FeII (8 sites). V has a strong tendency to occupy the FeI sublattice as shown by X-ray absorption fine-structure [1997Nis1], and this leads to the formation of Heusler phase at the ideal composition of VFe2Al [1976Vla, 1983Bus, 1985Okp, 1997Nis1, 1997Nis2, 2001Nis1, 2001Nis2]. While the addition of V in Fe3Al increases D03 B2 ordering temperature, the Curie temperature of D03 decreases monotonically [2001Kan]. This is shown in Fig. 2. Another consequence of substitution Fe by V is the decrease of lattice parameter of Fe3Al down to a
Landolt-Börnstein New Series IV/11A3
MSIT ®
2
Al–Fe–V
minimum at the ideal Heusler composition of VFe2Al beyond which it increases [2001Nis1, 2001Nis2]. This behavior shown in Fig. 3. [1969Bul] also reported the D03 B2 and B2 A2 ordering temperatures along Fe3Al-VFe3 section, both showing increasing trend as V is substituted for Al as shown in Fig. 4. However, in view of the above mentioned discrepancy, further measurements are needed to verify the results of [1969Bul]. As expected, V also increases D03 B2 ordering temperature of other Al-Fe alloys in the vicinity of Fe3Al. For example, [1995Ant] prepared three alloys VFe73Al26, V2Fe72Al26 and V 4Fe70Al26 and measured the ordering temperature using DTA. The D03 B2 temperature transition of these alloys are 585, 624 and 695°C, respectively. [1997And] determined site occupancy of V in V5Fe50Al45 (1) by ALCHEMI (Atom Location by CHanneling Enhanced MIcroanalysis) in TEM. [1997And] observed that about 80% of the “Al-site” is occupied by V, and the residual “Fe-site” is attributed to the kinetics of site-equilibrium mechanism. The Fe4Al13 phase can dissolve about 5 at.% V at 500°C [1987Sok] and about 2 at.% V at room temperature [1981Yin]. At 500°C, the VAl3, V4Al23, V 7Al45 and V2Al21 phases can dissolve up to about 6.5, 2.0, 1.7 and 4.5 at.% Fe, respectively [1987Sok]. The V solubilities in Fe2Al5, FeAl2 and FeAl were reported to be about 3, 1.7 and 10 at.% V, respectively [1988Sok]. However, [2000Sah] uses, in the Al-rich corner at 475°C, a diagram in which V4Al23 dissolves up to 4 at % Fe and Fe4Al13 dissolves up to 8 at.% V. In contrast to the results of [1987Sok], Skinner et al. [1988Ski] reported that melt-spinning of Al-rich alloys containing up to 16 at.% Fe and 10 at.% V gives rise to a quasicrystalline icosahedral phase. Also, [1988Ski] suggested that the lattice parameter of such an icosahedral phase is dependent on the Fe:V ratio in the alloy. Rapidly quenched alloys of the compositions 94Al-6Fe (at.%) and 95.3Al-4Fe-0.7V (at.%), which consisted of (Al) + slight amounts of FeAl6 were investigated by Mössbauer spectroscopy. Two kinds of coordination of Fe atoms in the Al lattice, a symmetric and an asymmetric one, were observed in the V containing alloy. In contrast to this result the Al-Fe alloy had shown only one kind of coordination [1989Sok]. The details of the crystal structures and lattice parameters of the solid phases are listed in Table 1. Isothermal Sections [1960Gup] reported the phase boundaries involving ) and (Fe) phases in the form of a partial isotherm at 1175°C. Al is a strong ) phase destabilizer; about 0.5 at.% Al at 1175°C is reported to be sufficient to suppress the ) phase completely. [1994Koz] prepared ribbons of Fe rich alloys by melt-spinning. The samples were annealed at 500C for 240ks, and were examined in a transmission electron microscope. Figure 5 shows the partial Al-Fe-V isothermal section at 500°C from [1987Sok] and [1988Sok]. It should be mentioned that the Al-V binary phases VAl6, VAl7 and VAl11 as designated by [1987Sok, 1988Sok], correspond to V4Al23, V 7Al45 and V 2Al21 in the presently accepted Al-V phase diagram. [2000Sah] presents, in the Al rich corner at 475°C a diagram in which the solubility of Fe in V2Al21 is very low so that Al may be in equilibrium with VAl10 and V7Al45 phases, which contradicts the observations of [1987Sok, 1988Sok]. Figure 1 also includes the results of TEM analyses on Fe-rich samples annealed at 500°C [1989Zha, 1994Koz]. Three types of phase separation sequences from the single phase regions of the , 1 and 2 phases into the +1 phase region have been distinguished [1989Zha]. [1987Sok] also reported the phases obtained in the as-melt-spun condition as well as after annealing at 250 and 450°C for 50 h. Their results are summarized in Table 2. It was noted that, except for the ternary alloy containing more than 16.5 at.% Fe and 3.6 at.% V which was annealed at 450°C for 50 h, equilibrium was not reached in the rest of the alloys after the annealing treatments used by the authors. For example, after annealing the binary Al-V and Al-Fe melt-spun alloys at both 250 and 450°C, the authors obtained (Al+VAl3+V2Al21) and (Al+Fe4Al13+FeAl6) phases, respectively. In the latter case, FeAl6 represents a metastable phase.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Fe–V
3
Thermodynamics [2003Kaw2] measured various thermophysical (dilatability, compressibility) and thermochemical properties of VFe2Al, and proposes, for the heat capacity, the following expression: Cp/J#mol-1#K-1 = 229 - 0.328 T + 2.50 # 10-3 T 2 - 5.63 #106 T - 2. [1994Koz] constructed the free energies of A2, B2 and D03 phases by a statistical approach employing Bragg-Williams-Gorsky approximation. They considered both atomic and magnetic interaction energies up to second nearest neighbor. Based on the model description of free energies, they calculated isothermal section at 500°C which is good agreement with the experimentally observed microstructures of Fe-rich alloys. Notes on Materials Properties and Applications Magnetic and electrical properties of V1-xFe2+xAl alloys have been studied extensively [1985Okp, 1997Nis1, 1997Nis2, 1998Kat, 1998Weh, 2000Kat, 2000Zar, 2001Fen, 2001Han, 2001Kan, 2001Lue, 2001Mak, 2001Nis1, 2001Nis2, 2001Sum, 2003Kaw1]. An important finding is that VFe2Al is nonmetallic with respect to transport properties while it is metallic with respect to its thermodynamic properties. For example, [1997Nis2] observed an anomalous negative temperature dependence of electrical resistivity such that it behaves almost like a semiconductor. This is despite the fact that it has a large density of states at the Fermi level as revealed by the photoemission valence-band spectra. VFe2Al is non-magnetic semimetal with a sharp pseudogap at the Fermi level [2000Kat]. It has been reported that a strong hybridization of Feand V-3d states causes a broadening of the d-states and their shift to the higher binding energy. As a result long-range magnetic order disappears and a narrow energy gap near the Fermi level is formed [2000Zar]. The unusual electron transport is mainly attributed to the effect of strong spin fluctuations, in addition to the existence of very low carrier concentrations [2000Kat]. [1962Min] studied the effect of V addition on the properties of Fe3Al. Addition of V increases hardness, electrical resistivity and also improves the high temperature mechanical properties. [2001Nis1] reported the mechanical properties of the (VxFe1-x)3Al alloys. In the composition range 0 x 0.38, the room temperature yield stress exhibits a double-well behavior starting from 550 MPa for Fe3Al with a first minimum at 150MPa for x = 0.02, a maximum at 300 MPa for x = 0.15 and a second minimum at 150 MPa for x = 0.333 corresponding to the composition VFe2Al. Furthermore, [2001Nis1] observed a correlation between the yield stress peak at higher temperature and the loss of D03 order. [2000Ino] reported a significant increase in strength of rapidly solidified Al-Fe-V alloys containing nano-quasicrystalline phase. Miscellaneous From a preliminary investigation of the section Fe4Al13-V 2Al21 , a eutectic reaction was claimed to exist at ~610°C with an invariant composition at ~83 at.% Al [1988Sok]. References [1960Gup]
[1962Min]
[1969Bul]
[1976Mon]
Landolt-Börnstein New Series IV/11A3
Gupta, K.P., Rajan, N.S., Beck, P.A., “Effect of Si and Al on the Stability of Certain ) Phases”, Trans. Met. Soc. AIME, 218, 617-624 (1960) (Equi. Diagram, Experimental, #, *, 18) 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), Dop. Akad. Nauk Ukrain. RSR, 144, 1324-1327 (1962) (Experimental, 1) Bulycheva, Z.N., Svezhova, S.I., Kondrat’ev, V.K., “Change in the Ordering Temperature of Fe3Al on Adding a Third Element” (in Russian), Ukrain. Fiz. Zhur., 14, 1706-1708 (1969) (Crys. Structure, Experimental, 5) Mondolfo, L.F., “Aluminum-Vanadium System”, in “Aluminium Alloys: Structure and Properties”, Butterworths, London, 392-394 (1976) (Review, 46)
MSIT ®
4 [1976Vla]
[1981Yin]
[1982Kub] [1983Bus]
[1985Okp]
[1986Len] [1987Sok]
[1988Ski]
[1988Sok]
[1989Koz]
[1989Mur] [1989Sok]
[1989Zha]
[1992Gho]
[1992Rag]
[1993Miy]
[1994Koz]
MSIT®
Al–Fe–V Vlasova, E.N., Prokoshin, A.F., “Formation of L12 Substructure and Stratification in Solid Fe-Cr Solutions Doped with Al and V” (in Russian), Dokl. Akad. Nauk SSSR, 231, 599-602 (1976) (Crys. Structure, Experimental, 2) Ying-Hong, Z., Jing-Qi, L., Jiang-Xuang, Z., Cheng, C.S., “A Room-Temperature Section of the Phase Diagram of TiAl 3-VAl3-MAl3 of the System Alloys of Al-Ti-V-M (M = Ni, Fe)”, Acta Phys. Sin. (Chin. J. Phys.), 30, 972-975 (1981) (Crys. Structure, Experimental, Equi. Diagram, 4) Kubaschewski, O., “Fe-V”, in “Iron-Binary Phase Diagrams”, Springer Verlag, Berlin, 160-164 (1982) (Equi. Diagram, #, 15) Buschow, K.H.J., van Engen, P.G., Jongebreur, R., “Magneto-Optical Properties of Metallis Ferromagnetic Materials”, J. Magn. Magn. Mater., 38, 1-22 (1983) (Magn. Prop., Optical Prop., 23) Okpalugo, D.E., Both, J.G., Faunce, C.A., “Onset of Ferromagnetism in 3d-Substituted Fe-Al Alloys. I: Ti, V and Cr Substitutions”, J. Phys. F, Met. Phys., 15, 681-692 (1985) (Crys. Structure, Experimental, 21) Lendvai, A., “Phase Diagram of Al-Fe Sytem up to 45 mass% Iron”, J. Mater. Sci. Lett., 5, 1219-1220 (1986) (Equi. Diagram, Experimental, #, *, 7) Sokolovskaya, E.M., Badalova, L.M., Kazakova, E.F., “Phase Composition of Rapidly Quenched Alloys of the System Al-Fe-V”, Izv. Akad. Nauk SSSR, Met., (5), 212-215 (1987) (Equi. Diagram, Experimental, #, *, 7) Skinner, D.J., Ramanan, V.R.V., Zedalis, M.S., Kim, W.J., “Stability of Quasicrystalline Phases in AlFeV Alloys”, Mater. Sci. Eng., 99, 407-411 (1988) (Crys. Structure, Experimental, 8) Sokolovskaya, E.M., Badalova, L.M., Kazakova, E.F., Stroeva, N.V., “Interactions of Intermetallic Compounds in the Ternary System Aluminum-Iron-Vanadium” (in Russian), Vestn. Mosk. Univ., Ser. 2: Khim., 29(3), 303-306 (1988) (Experimental, 5) Kozakai, T., Zhao, P.Z., Miyazaki, T., “Phase Separations in Fe-Rich Fe-Base Ternary Ordering Alloy Systems”, Met. Abstr. Light Metals and Alloys, 23, 32-33 (1989/1990) (Crys. Structure, Equi. Diagram, Experimental, 0) Murray, J.L., “Al-V (Aluminum-Vanadium)”, Bull. Alloy Phase Diagrams, 10(4), 351-357 (1989) (Crys. Structure, Equi. Diagram, Review, 34) Sokolovskaya, E.M., Badalova, L.M., Kazakova, E.F., Reiman, S.I., Ryaskyi, G.K., Sorokin, A.A., Philipova, A.A., Chaldieva, G.M., “Investigation of Chemical Composition Microcrystalline of an Al Alloys with Transition Metals” (in Russian), Vestn. Mosk. Univ., Ser. 2: Khim., 30(2), 162-165 (1989) (Crys. Structure, Experimental, 6) Zhao, P.Z., Kozakai, T., Miyazaki, T., “Phase Separation into A2+D03 Two Phases in Iron-Aluminium-Vanadium Ternary Ordering Alloys” (in Japanese), Nippon Kinzoku Gakkai Shi, 53(3), 266-272 (1989) (Crys. Structure, Equi. Diagram, Experimental, #, *, 23) Ghosh, G., “Aluminium-Iron-Vanadium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.19022.1.20, (1992) (Crys. Structure, Equi. Diagram, Assessment, 14) Raghavan, V., “The Al-Fe-V (Aliminium-Iron-Vanadium) System”, in Phase Diagram of Ternary Iron Alloys, Part 6A, Ind. Inst. Metals, Calcutta, 204-207 (1992) (Review, Equi. Diagram, 7) Miyazaki, T., “Phase Diagrams of Iron-Base Ternary Ordering Alloy Systems”, Comput. Aided Innovation New Mater. 2, Proc. Int. Conf. Exhib. Comput. Appl. Mater. Mol. Sci. Eng., 2nd 1992 (Pub. 1993) (Pt.1), 707-712., 2ND1992 (1993) (Equi.Diagram) Kozakai, T., Miyazaki, T., “Experimental And Theoretical Investigations on Phase Diagrams of Fe Base Ternary Ordering Alloys”, ISIJ Int., 34(5), 373-383 (1994) (Calculation, Equi. Diagram, Experimental, Magn. Prop., #, *, 18)
Landolt-Börnstein New Series IV/11A3
Al–Fe–V [1995Ant]
[1997And]
[1997Nis1]
[1997Nis2]
[1998Kat]
[1998Weh] [2000Ino]
[2000Kat]
[2000Sah]
[2000Zar]
[2001Fen]
[2001Han] [2001Kan]
[2001Lue]
[2001Mak]
[2001Nis1]
[2001Nis2]
Landolt-Börnstein New Series IV/11A3
5
Anthony, L., Fultz, B., “Effects of Early Transition Metal Solutes in the D03-B2 Critical Temperature of Fe3Al”, Acta Metall. Mater., 43, 3885-3891 (1995) (Crys. Structure, Experimental, 35) Anderson, I.M., “Alchemi Study of Site Distributions of 3d-Transition Metals in B2-Ordered Iron Aluminides”, Acta Mater., 45(9), 3897-3909 (1997) (Calculation, Crys. Structure, Experimental, Theory, 26) Nishino, Y., Kumada, C., Asano, S., “Phase Stability of Fe3Al with Addition of 3d Transition Elements”, Scr. Mater., 36, 461-466 (1997) (Crys. Structure, Equi. Diagram, Experimental, 26) Nishino, Y., Kato, M., Asano, S., Soda, K., Hayasaki, M., Mizutani, U., “Semiconductor-Like Bahavior of Electrical Resisitivity in Heusler-Type Fe2VAl Compound”, Phys. Rev. Lett., 79(10), 1909-1912 (1997) (Crys. Structure, Experimental, 18) Kato, M., Nishino, Y., Asano, S. Ohara, S., “Electrical Resistance Anomaly and Hall Effect in (Fe1-xVx)3Al Alloys” (in Japanese), J. Japan. Inst. Met., 62(7), 669-674 (1998) (Crys. Structure, Experimental, 23) Weht, R., Pickett, W.E., “Excitonic Correlations in the Intermetallic Fe2VAl”, Phys. Rev. B, 58(11), 6855-6861 (1998) (Calculation, Crys. Structure, Mechan. Prop., 21) Inoue, A., Kimura, H.M., Zhang, T., “High-Strength Aluminium- and Zirconium-Based Alloys Containing Nanoquasicrystalline Particles”, Mater. Sci. Eng. A, 294-296, 727-735 (2000) (Crys. Structure, Experimental, Mechan. Prop., 28) Kato, M., Nishino, Y., Mizutani, Y., Asano, S., “Electronic, Magnetic and Transport Properties of (Fe1-xVx)3Al Alloys”, J. Phys.: Condens. Matter, 12, 1769-1779 (2000) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop., Phys. Prop., 33) Sahoo, K.L., Sivaramakrishnan, C.S., Chakrabarti, A.K., “Solidification Characteristics of the Al-8.3Fe-0.8V-0.9Si Alloy”, Metall. Mater. Trans. A, 31A(6), 1599-1610 (2000) (Experimental, #, 21) Zarek, W., Talik, E., Heimann, J., Kulpa, M., Winiarski, A., Neumann, M., “Electronic Structure, Magnetic and Electrical Properties of Fe3-xVxAl Compounds”, J. Alloys Compd., 297, 53-58 (2000) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop., 15) Feng, Y., Rhee, J.Y., Wiener, T.A., Lynch, D.W., Hubbard, B.E., Sievers, A.J., Schlagel, D.L., Lograsso, T.A., Miller, L.L., “Physical Properties of Heusler-Like Fe2VAl”, Phys. Rev. B, 63(16), 165109-1-165109-12 (2001) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop., Phys. Prop., 30) Hanada, Y., Suzuki, R.O., Ono, K., “Seebeck Coefficient of (Fe,V)3Al Alloys”, J. Alloys Compd., 329, 63-68 (2001) (Electr. Prop., Experimental, 18) Kanomata, T., Sasaki, T., Hoshi, T., Narita, T., Harada, T., Nishihara, H., Yoshida, T., Note, R., Koyama, K., Nojiri, H., Kaneko, T., Motokava, M., “Magnetic and Electrical Properties of Fe 2+xV1-xAl”, J. Alloys Compd., 317-318, 390-394 (2001) (Crys. Structure, Electr. Prop., Experimental, 19) Lue, C.S., Ross, J.H., Rathnayaka, Jr., K.D.D., Naugle, D.G., Wu, S.Y., Li, W.-H., “Supermagnetism and Magnetic Defects in Fe2VAl and Fe2VGa”, J. Phys.: Condens. Matter, 13, 1585-1593 (2001) (Crys. Structure, Experimental, Magn. Prop., 25) Maksimov, I., Baabe, D., Klauss, H.H., Litterst, F.J., Feyerherm, R., Toebbens, D.M., Matsushita, A., Suellow, S., “Structure and Magnetic Order in Fe2+xV1-xAl”, J. Phys.: Condens. Matter, 13, 5487-5501 (2001) (Crys. Structure, Experimental, Magn. Prop., 25) Nishino, Y., “Electronic Structure and Transport Properties of Pseudogap System Fe2VAl”, Mater. Trans., JIM, 42(6), 902-910 (2001) (Crys. Structure, Electr. Prop., Equi. Diagram, Experimental, 58) Nishino, Y., Makino, Y., “Effect of Vanadium Substitution on Strength Properties of Fe3Al-Based Alloys”, Mater. Sci. Eng. A, 319-321, 368-371 (2001) (Equi. Diagram, Experimental, Mechan. Prop., #, *, 29)
MSIT ®
Al–Fe–V
6 [2001Sum]
[2002Rag] [2003Kaw1]
[2003Kaw2]
[2003Pis]
[2003Sch]
Sumi, H., Kato, M., Nishino, Y., Asano, S., Mizutani, U., “Electrical Resistivity Anomaly and Magnetic Properties in Heusler-Type Fe2VAl Alloy” (in Japanese), J. Jpn. Inst. Met., 65(9), 771-774 (2001) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop., Thermodyn., 16) Raghavan, V., “Al-Fe-V (Aliminum-Iron-Vanadium) System”, J. Phase Equilib., 23, 439-440 (2002) (Equi. Diagram, Review, 7) Kawaharada, Y., Kurosaki, K., Yamanaka, S., “High Temperature Thermoelectric Properties of (Fe1-xVx) 3Al Heusler Type Compounds”, J. Alloys Compd., 349(1-2), 37-40 (2003) (Electr. Prop., Experimental, Mechan. Prop., Phys. Prop., 27) Kawarahada, Y., Kurosaki, K., Zamanaka, S., “Thermophysical Properties of Fe2VAl”, J. Alloys Compd., 352, 48-51 (2003) (Thermodyn., Phys. Prop., Mechan. Prop., Experimental, 22) Pisch, A., “Al-Fe (Aluminum-Iron)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2003) (Equi. Diagram, Review, 58) Schuster, J.C., “Al-V (Aluminium-Vanadium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2003) (Equi. Diagram, Review, 31)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) 660.452 , (Fe) 912 (V) 1910 V5Al8 1408
VAl3 1270
V4Al23 736
MSIT®
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cI2 Im3m W cI2 Im3m W cI52 I43m Cu5Zn8 tI8 I4/mmm TiAl3 hP54 P63/mmc V4Al23
Lattice Parameters Comments/References [pm] a = 404.96
pure Al at 25°C [Mas2]
a = 286.65
pure Fe at 25°C [Mas2]
a = 302.40
pure V at 25°C [Mas2]
a = 923.0 a = 921.8
[2003Sch], Al-rich [2003Sch], V-rich solid solubility ranges from 60.0 to 66.0 at.% Al [2003Sch], Al-rich limit
a = 378.14 c = 832.2 a = 378.07 c = 830.9 a = 769.28 c = 1704.0 a = 769.9 c = 1705.3
[2003Sch], V-rich limit solubility ranges from 74 to 75 at.% Al [1989Mur] [2003Sch]
Landolt-Börnstein New Series IV/11A3
Al–Fe–V Phase/ Temperature Range [°C] V7Al45 730
V2Al21 690 V3Al(r) 650 1, Fe3Al 547 2, FeAl 1310 J, Fe2Al3 1102 - 1232 FeAl2 1156
Pearson Symbol/ Space Group/ Prototype mC104 C2/m V7Al45
cF184 Fd3m V2Al21 cP8 Pm3n Cr3Si cF16 Fm3m BiF3 cP2 Pm3m CsCl cI16? aP18 P1 FeAl2
Fe2Al5 1169
oC24 Cmcm
Fe4Al13 1160
mC102 C2/m Fe4Al13
), VFe 1252
tP30 P42/mnm )CrFe cF16 Fm3m BiF3
VFe2Al
Landolt-Börnstein New Series IV/11A3
7
Lattice Parameters Comments/References [pm] a = 2540 b = 759 c = 1100 = 127 a = 2563.0 b = 763.7 c = 1108.8 = 128.83 a = 1449.2 a = 1452.1
[1989Mur]
a = 482.9
[2003Sch]
[2003Sch]
[1989Mur, 2003Sch]
a = 578.86-579.30 [2003Pis], solid solubility ranges from ~24 to ~37 at.% Al a = 289.76-290.78 [2003Pis], at room temperature solid solubility ranges from 39.7 to 54.5 at.% Al a = 598.0 [2003Pis], solid solubility ranges from 54.5 to 62.5 at.% Al [2003Pis], at 66.9 at.% Al a = 487.8 solid solubility ranges b = 646.1 c = 880.0 from 65.5 to 67.0 at.% Al = 91.75° = 73.27° = 96.89° a = 765.59 [2003Pis], at 71.5 at.% Al b = 641.54 solid solubility ranges c = 421.84 from 71.0 to 72.5 at.% Al a = 1552.7-1548.7 [2003Pis], 74.16 to 76.7 at.% Al solid solubility ranges b = 803.5-808.4 c = 1244.9-1248.8 from 74.5 to 75.5 at.% Al = 107.7-107.99° Sometimes called FeAl3 in the literature [2003Pis], at 76.0 at.% Al a = 1549.2 b = 807.8 c = 1247.1 = 107.69° a = 895.6 [V-C2], solid solubility c = 462.7 ranges from 33.5 to 64.0 at.% V a = 576.1 a = 576.16 a = 576.0 a = 576.19
[1983Bus] Heussler Alloy [1998Kat] [2001Lue] [2001Nis1]
MSIT ®
Al–Fe–V
8
Table 2: Phases Present in the as-melt-spun Condition and after Annealing Treatments Composition (at.%) Al Fe 98.0 96.0 94.0 91.0 87.0 98.0 2.0 95.0 5.0 92.0 e 8.0 14.0 86.0 2.0 97.5 6.5 92.5 8.5 90.0 10.0 88.0
As-melt-spun V 2.0 4.0 6.0 9.0 13.0 0.5 1.0 1.5 2.0
(Al) (Al)+VAl3 (Al)+VAl3 (Al)+VAl3 (Al)+VAl3 (Al)+FeAl6 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al) (Al)+FeAl6 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13
After annealing for 50 h, at [°C] 250 450 (Al)+VAl3+V 2Al21 (Al)+VAl3+V2Al21 (Al)+VAl3+V 2Al21 (Al)+VAl3+V2Al21 (Al)+VAl3+V 2Al21 (Al)+VAl3+V2Al21 (Al)+VAl3+V 2Al21 (Al)+VAl3+V2Al21 (Al)+VAl3+V 2Al21 (Al)+VAl3+V2Al21 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeA l6+Fe4Al13 (Al) (Al)+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+Fe4Al13 (Al)+FeAl6+V2Al21 (Al)+Fe4Al13+V 2Al21
1100
Fig. 1: Al-Fe-V. Variation of D03 B 2 ordering temperature along V3Al-Fe 3Al section
VFe2Al 1000
B2(α 1)
Temperature, °C
900
800
DO3(α 2) 700
600
500
V Fe Al
MSIT®
30.00 45.00 25.00
50
60
Fe, at.%
70
V Fe Al
0.00 75.00 25.00
Landolt-Börnstein New Series IV/11A3
Al–Fe–V
9
500
Fig. 2: Al-Fe-V. Variation of Curie temperature of D03 phase along V3Al-Fe3Al section
400
Temperature, °C
300
200
100
0
-100
-200
-300
V Fe Al
60
25.00 50.00 25.00
70
V Fe Al
Fe, at.%
0.00 75.00 25.00
579.5
Fig. 3: Al-Fe-V. Variation of lattice parameter of D03 phase along Fe3Al-V 3Al section
579.0
Lattice parameter, pm
578.5
578.0
577.5
577.0
Fe2VAl
576.5
576.0
Al Fe V
Landolt-Börnstein New Series IV/11A3
45
25.00 75.00 0.00
50
55
60
Fe, at.%
65
70
75
Al 25.00 Fe 45.00 V 30.00
MSIT ®
Al–Fe–V
10
900
Fig. 4: Al-Fe-V. Variation of order-disorder reaction temperature as a function of V content along the VFe3-Fe3Al section
(α Fe)
800
Temperature, °C
CsCl-type (α 2 ) 700
600
BiF3-type (α 1 ) 500
400
V Fe Al
20
25.00 75.00 0.00
10
V Fe Al
V, at.%
Al Fig. 5: Al-Fe-V. Partial isothermal section at 500°C
VAl3
Data / Grid: at.%
(Al)
VAl10 V7Al45 V4Al23
0.00 75.00 25.00
20
Axes: at.%
80
Fe4Al13 Fe2Al5 FeAl2
V5Al8 40
60
α2
60
40
V3Al
α1
80
20
α +α 1
V
MSIT®
α
(V)+σ
(V) 20
σ 40
60
80
Fe
Landolt-Börnstein New Series IV/11A3
Al–Fe–Y
11
Aluminium – Iron – Yttrium Gabriele Cacciamani Literature Data The Al-Fe-Y phase equilibria have been systematically investigated by [1972Zar] in the 0-33 at.% Y composition range. Structural and magnetic properties of the Al-Fe-Y phases have been studied by several authors: investigations mainly concerned the solid solutions at the Y2(Fe,Al) 17 ratio [1976McN, 1996Kuc, 1998Che, 1998Kam, 2001Vor] and the Y(Fe1-xAlx)12 ternary phase [1966Zar, 1974Viv, 1976Bus, 1978Bus, 1980Fel, 1995Sch, 2000Sch, 2001Wae2]. Binary and ternary phases at the Y(Fe,Al)2 atomic ratio have been mainly investigated by [1972Ryk, 1973Zar, 1975Bus, 1975Dwi, 1976Gro, 1977Mur, 1986Sec, 1988Cun, 2001Wae2]. The YFe2Al10 phase has been studied by [1998Thi, 2001Wae2]. Samples have been generally prepared by arc melting the pure elements (usually 99.9 mass% pure) under an inert atmosphere. In a few cases other methods were used: synthesis in Al2O3 at 400 to 800°C [1998Thi] or induction melting of Al-Fe master alloys with appropriate amounts of rare earth [1975Dwi]. Samples were generally annealed at appropriate temperatures and then quenched. This evaluation incorporates and continues the critical evaluation made by [1992Gri] considering new published data. Binary Systems The binary systems Al-Fe and Al-Y are accepted from [2003Pis] and [2003Cor], respectively. The Fe-Y phase equilibria are accepted from the assessment by [1992Zha]. Solid Phases Crystal structure data are reported in Table 1. Al-Fe binary compounds and phases are not reported to dissolve Y. Al-Y and Fe-Y phases may show more or less extended solubility ranges due to substitution between Al and Fe. The binary Laves phases YAl2 and YFe2 (isostructural, MgCu2 type) dissolve more than 20 at.% of the third element. At intermediate compositions, however, a different Laves phase (-1, MgZn2 type) is formed: the solubility ranges have been studied by [1975Dwi] and crystal structures by [1972Ryk, 1972Zar, 1973Zar, 1975Bus, 1976Gro, 1977Mur, 1986Sec, 1988Cun]. The solid solutions at the Y2(Fe,Al) 17 ratio have been studied by different authors [1976McN, 1996Kuc, 1998Che, 1998Kam, 2001Vor]: both Th2Ni17 and Th2Zn17 structures have been reported, but their composition and temperature ranges of stability are still not well assessed. The -2 phase has been studied by several authors either at the YFe4Al8, [1976Bus, 1978Bus], YFe6Al6 [1980Fel, 1988Che] or different compositions [1995Sch, 1998Sch, 2000Sal, 2000Sch, 2001Wae2]. Also in this case the solubility range seems to vary appreciably with temperature. Finally, with the same Y(Fe,Al)12 ratio, a different ternary phase (-3, at the composition HoFe2Al10) was first identified by [1972Zar] and then studied by [1998Thi, 2001Wae2]. Isothermal Sections The partial isothermal section at 500°C is reported in Fig. 1. Determined by [1972Zar], it has been adapted considering the more recent indications concerning the solubility ranges of the solid solutions (homogeneity ranges of YFe2 and YAl2 after [1975Dwi]) and the accepted binary systems. The 800°C isothermal section has been recently investigated by [2001Wae2] in the 50-100 at.% Al region: it resulted to be consistent with the section by [1972Zar].
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Fe–Y
12 Thermodynamics
Thermodynamic properties of the liquid phase have been studied by [1982Erm, 1983Erm1, 1983Erm2]. Notes on Materials Properties and Applications Mössbauer measurements have been carried out on the -1 [1975Dwi, 1977Mur], -2 [2000Wae, 2001Wae1, 2003Kal] and -3 [2001Wae2] phases. Magnetic properties have been studied for -2 at different compositions: YFe4Al8 [1978Bus, 1998Hag, 1998Sch, 2000Sik, 2000Wae, 2001Pai, 2001Wae1], YFe6Al6 [1981Fel], YFe5Al7 [1995Sch], YFe7Al5 [2000Sch] and variable composition [2000Wae, 2003Kal], and for -1 [1975Bus, 1976Gro, 1977Mur, 1986Sec], -3 [1998Thi], and the phases at the Y2(Fe,Al)17 ratio [1986Plu, 1996Kuc, 1998Che, 1998Kam, 1999Kuc, 2001Kny, 2001Vor]. [2001Kny] investigated also the optical properties of the Y2(Fe,Al)17 solid solution. [1988Cun] carried out resistivity measurements on -1 and [1992Joh, 1998All] studied the formation of amorphous and nano-crystalline alloys in the system. References [1958Tay]
[1961Lih]
[1966Zar]
[1972Ryk]
[1972Zar]
[1973Zar]
[1974Viv]
[1975Bus] [1975Dwi]
[1976Bus]
[1976Gro]
MSIT®
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, Magn. Prop., Experimental, 49) Lihl, F., Ebel, H., “X-Ray Examination fo the Constitution of Iron-Rich Alloys of the Iron-Aluminium System” (in German), Arch. Eisenhuettenwesen, 32, 483-487, (1961) (Crys. Structure, Magn. Prop., Experimental, 12) Zarechnyuk, O.S., “Ternary Compounds with a ThMn12 Superstructure in the Systems Yttrium-Transition Metal-Aluminium”, Dop. Akad. Nauk Ukr. RSR, 6, 767-769 (1966) (Crys. Structure, 2) Rykhal, R.M., “Crystal Structures of the Ternary Compounds YFeAl and YCoAl” (in Russian), Vestn. L'vov. Univ., Ser. Khim., 13, 11-14 (1972) (Crys. Structure, Experimental, 4) Zarechnyuk, O.S., Rikhal', R.M., Ryabov, V.R., Vivchar, O.I., “The Y-Fe-Al Ternary System in the Region 0 - 33.3 at.% Y”, Izv. Akad. Nauk SSSR, Met., (1), 208 (1972) (Crys. Structure, Equi. Diagram, Experimental, 12) Zarechnyuk, O.S., Rikhal, R.M., Vivchar, O.I., “Laves Phases in Ternary Systems of the Type Rare-Earth Metal-Transition Metal-Al” (in Russian), Akad. Nauk Ukr. SSR, Metallofizika, 46, 92-94 (1973) (Crys. Structure, Experimental, 22) Vivchar, O.I., Zarechnyuk, O.S., “Compounds of the ThMn12-Type Structure in R-Fe-Al Systems” (in Russian), Tezisy Dokl. - Vses. Konf. Kristallokhim. Intermet. Soedin., Rykhal, R.M. (Ed), Vol. 2, L'vov. Gos. Univ.: Lvov, USSR, 41 (1974) (Crys. Structure, Experimental, 0) Buschow, K.H.J., “Crystal Structure and Magnetic Properties of YFe2xAl2-2x”, J. Less-Common Met., 40, 361-363 (1975) (Crys. Structure, Experimental, 6) Dwight, A.E., Kimball, C.W., Preston, R.S., Taneja, S.P., Weber, L., “Crystallographic and Moessbauer Study of (Sc, Y, Ln)(Fe, Al)2 Intermetallic Compounds”, J. Less-Common Met., 40, 285-291 (1975) (Crys. Structure, Moessbauer, Experimental, 8) Buschow, K.H.J., van Vucht, J.H.N., van den Haagenhof, W.W., “Note on the Crystal Structure of the Ternary Rare Earth 3d Transition Metal Compounds of the Type RT4Al8”, J. Less-Common Met., 50(1), 145-150 (1976) (Experimental, Crys. Structure, 2) Groessinger, R., Steiner, W., Krec, K., “Magnetic Investigations of Pseudobinary RE(Fe,Al)2 Systems (RE = Y, Gd, Dy, Ho)” (in German), J. Magn. Magn. Mater., 2, 196-202 (1976) (Magn. Prop., Experimental, 20)
Landolt-Börnstein New Series IV/11A3
Al–Fe–Y [1976Mcn]
[1977Mur]
[1978Bus]
[1980Fel] [1981Fel]
[1982Erm]
[1983Erm1]
[1983Erm2]
[1985Gan]
[1986Plu]
[1986Sec] [1987Ric] [1988Che]
[1988Cun]
[1989Gsc]
[1992Gri]
Landolt-Börnstein New Series IV/11A3
13
McNelly, D., Oesterreicher, H., “Structural and Low-Temperature Magnetic Studies on Compounds Sm2Fe17 with Al Substitution for Fe”, J. Less-Common Met., 44, 183-193 (1976) (Crys. Structure, Magn. Prop., Experimental, 26) Muraoka, Y., Shiga, M., Nakamura, Y., “Magnetic Properties and Moessbauer Effects of A(Fe1-xBx)2 (A = Y or Zr, B = Al or Ni) Laves Phase Intermetallic Compounds”, Phys. Status Solidi A, 42A, 369-374 (1977) (Crys. Structure, Magn. Prop., Moessbauer, Experimental, 15) Buschow, K.H.J., van der Kran, A.M., “Magnetic Ordering in Ternary Rare Earth Iron Aluminium Compounds (RFe4Al8)”, J. Phys., F: Met. Phys., 8, 921-932 (1978) (Experimental, Magn. Prop., 9) Felner, I., “Crystal Structures of Ternary Rare Earth-3d Transition Metal Compounds of the RT6Al6 Type”, J. Less-Common Met., 72, 241-249 (1980) (Crys. Structure, 10) Felner, I., Seh, M., Rakavy, M., Nowik, I., “Magnetic Order and Hyperfine Interactions in RFe6Al6 (R = Rare Earth)”, Phys. Chem. Solids, 42, 369-377 (1981) (Crys. Structure, Magn. Prop., Experimental, 6) Ermakov, A.F., Esin, Yu.O., Gel'd, P.V., “Partial and Integral Enthalpies of Formation of Liquid Alloys of Iron Monoaluminide with Yttrium, Lanthanum and Cerium” (in Russian), Izv. Akad. Nauk SSSR, Met., (5), 69-60 (1982) (Thermodyn., Experimental, 3) Ermakov, A.F., Esin, Yu.O., Levin, E.S., Petrusevskij, M.S., “Estimation of the Enthalpy of Formation of Liquid Ternary Alloys Fe-Y-Si and Fe-Y-Al from the Data of Characteristic Boundaries of the Binary Systems” (in Russian), Fiz. Svoistva Met. Splavov (Sverdlovsk), (4), 68-71 (1983) (Thermodyn., Experimental, 9) Ermakov, A.F., Esin, Yu.O., Levin, E.S., Petrusevskij, M.S., “Assessment of the Enthalpy of Formation of Iron, Yttrium, Silicon and Iron-Yttrium, Aluminum Liquid Ternary Alloys” (in Russian), Fiz. Svoistva Met. Splavov (Sverdlovsk), (4), 71-74 (1983) (Experimental, Thermodyn., 4) Gan, R.J., Littlewood, N.T., James, W.J., “Magnetic Structures of Y 6(Fe1-xAlx)23 Compounds”, IEEE Trans., Magn., 21(5), 1984-1986 (1985) (Crys. Structure, Magn. Prop., Experimental) Plusa, D., Pfranger, R., Wyslocki, B., Mydlarz, T., “Magnetic Properties of Y 2(Fe1-xAlx)17 Pseudobinary Compounds”, J. Less-Common Met., 120, 1-7 (1986) (Crys. Structure, Experimental, 11) Sechovsky, V., Nozar, P., “Magnetic Phase Diagram of the System Yttrium - Iron Aluminum (Y(FexAl1-x)2)”, Acta Phys. Slovaca, 36(3), 210-211 (1986) (Magn. Prop., 3) Richter, R., Altounian, Z., Strom-Olsen, J.O., “Y5Al3, A New Y-Al Compound”, J. Mater. Sci., 22, 2983-2986 (1987) (Experimental, Thermodyn., Crys. Structure, 7) Chelkowska, G., Chelkowska, A., Winiarska, A., “Magnetic Susceptibility and Structural Investigations of Rare Earth-Aluminium-Iron (REAl 6Fe6) Compounds for RE = Yttrium, Terbium, Dysprosium, Holmium, and Erbium”, J. Less-Common Met., 143, L7-L10 (1988) (Crys. Structure, Magn. Prop., Experimental, 12) Da Cunha, S.F., Souza, G.P., Takeuchi, A.Y., “Electrical Resistivity of YFeAl (Y(Fe1-xAlx)2) in the Spin Glass”, J. Magn. Magn. Mater., 73(3), 355-360 (1988) (Crys. Structure, Electr. Prop., Experimental, 18) Gschneidner Jr, K.A., Calderwood, F.W., “The Al-Y (Aluminium-Yttrium) System”, Bull. Alloy Phase Diagrams, 10, 44-47 (1989) (Calculation, Equi. Diagram, Crys. Structure, Review, #, 33) Grieb, B., “Aluminium-Iron-Yttrium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.17517.1.20, (1992) (Crys. Structure, Equi. Diagram, Assessment, 18)
MSIT ®
14 [1992Joh]
[1992Zha] [1993Kat]
[1994Bur]
[1994Fol]
[1994Gri]
[1995Sch]
[1996Kuc]
[1997Kog]
[1998Ali]
[1998All]
[1998Che]
[1998Hag]
[1998Kam]
[1998Sch]
[1998Thi]
MSIT®
Al–Fe–Y Johnson, E., Johansen, A., Sarholt-Kristensen, L., “On Glass Formation in Rapidly Solidified Aluminium-Based Alloys”, J. Mater. Res., 7(10), 2756-2764 (1992) (Crys. Structure, Experimental, Phys. Prop., 35) Zhang, W., Liu, G., Han, K., “The Fe-Y (Iron-Yttrium) System”, J. Phase Equilib., 13(3), 304-308 (1992) (Equi. Diagram, Thermodyn., Review, #, 29) Kattner, U.R., Burton, B.P., “Al-Fe (Aluminum-Iron)“, in “Phase Diagrams of Binary Iron Alloys”, Okamoto, H. (Ed), ASM International, Materials Park, Ohio 44073-0002, 12-28 (1993) (Equi. Diagram, Review, 99) Burkhardt, U., Grin, J., Ellner, M., Peters, K., “Structure Refinement of the Iron-Aluminium Phase with the Approximate Composition Fe2Al5”, Acta Crystallogr., Sect. B: Struct. Crystallogr. Crys. Chem., B50, 313-316 (1994) (Crys. Structure, Experimental, 9) Foley, J.C., Thoma, D.J., Perepezko, J.H., “Supersaturation of the Al2Y Laves Phase by Rapid Solidification”, Metall. Mater. Trans. A, 25A, 230-233 (1994) (Crys. Structure, Experimental, 8) Grin, J., Burkhardt, U., Ellner, M., Peters, K., “Refinement of the Fe4Al13 Structure and its Relationship to Quasihomological Homotypical Structures”, Z. Kristallogr., 209, 479-487 (1994) (Crys. Structure, Experimental, 39) Schaefer, W., Kockelmann, W., Will, G., Fischer, P., Gal, J., “Neutron Diffraction on YFe5Al7 as Reference of the f-Magnetism of Isostructural Rare Earth - Iron - Aluminium Compounds”, J. Alloys Compd., 225, 440-443 (1995) (Crys. Structure, Experimental, Magn. Prop., 17) Kuchin, A.G., Kourov, N.I., Knyazev, Yu.V., Kleinerman, N.M., Serikov, V.V., Ivanova, G.V., Ermolenko, A.S., “Electronic, Magnetic, and Structuralproperties of the Alloys Y 2(Fe1-xMx)17 where M = Al and Si”, Phys. Status Solidi A, A155, 479-483 (1996) (Crys. Structure, Experimental, 4) Kogachi, M., Haraguchi, T., “Quenched-in Vacancies in B2-Structured Intermetallic Compound FeAl”, Mater. Sci. Eng. A, A230, 124-131 (1997) (Crys. Structure, Experimental, 23) Aliravci, C.A., Pekgueleryuez, M.O., “Calculation of Phase Diagrams for the Metastable Al-Fe Phases Forming in Direct-chill (DC)-Cast Aluminium Alloy Ingots”, Calphad, 22, 147-155 (1998) (Calculation, Equi. Diagram, 20) Allen, D.R., Foley, J.C., Perepezko, J.H., “Nanocrystal Development During Primary Crystallization of Amorphous Alloys”, Acta Mater., 46(2), 431-440 (1998) (Calculation, Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 39) Cheng, Z., Shen, B., Yan, Q., Guo, H., Chen, D., Gou, C., Sun, K., de Boer, F.R., Buschow, K.H.J., “Strcuture, Exchange Interactions, and Magnetic Phase Transition of Er2Fe17-xAlx Intermetallic Compounds”, Phys. Rev. B, 57B(22), 14299-14309 (1998) (Crys. Structure, Experimental, 35) Hagmusa, I.H., Brueck, E., de Boer, F.R., Buschow, K.H.J., “Magnetic Properties of RFe4Al8 Compounds Studied by Specific Heat Measurements”, J. Alloys Compd., 278, 80-82 (1998) (Thermodyn., Magn. Prop., Experimental, 9) Kamimori, T., Koyama, K., Mori, Y., Asano, M., Kinoshita, K., Mochimaru, J., Konishi, K., Tange, H., “Preferential Site Occupation of M Atoms and the Curie Temperature in Y 2Fe17-xMx (M = Al, Si, Ga)”, J. Magn. Magn. Mater., 177/181, 1119-1120 (1998) (Crys. Structure, Experimental, 4) Schobinger-Papamantellos, P., Buschow, K.H.J., Ritter, C., “Magnetic Ordering and Phase Transitions of RFe4Al8 (R = La, Ce, Y, Lu) Compounds by Neutron Diffraction”, J. Magn. Magn. Mater., 186, 21-32 (1998) (Crys. Structure, Experimental, Magn. Prop., 13) Thiede, V.M.T., Ebel, T., Jeitschko, W., “Ternary Aluminides LnT2Al10 (Ln = Y, La-Nd, Sm, Cd-Lu and T = Fe, Ru, Os) with YbFe2Al10 Type Structure and Magnetic Properties of the Iron-Containing Series”, J. Mater. Chem., 8(1), 125-130 (1998) (Crys. Structure, Magn. Prop., Experimental, 31) Landolt-Börnstein New Series IV/11A3
Al–Fe–Y [1999Dub]
[1999Kuc]
[2000Sal]
[2000Sch]
[2000Sik]
[2000Wae]
[2001Ike]
[2001Kny]
[2001Pai]
[2001Vor]
[2001Wae1]
[2001Wae2]
[2003Cor]
[2003Kal]
[2003Pis]
Landolt-Börnstein New Series IV/11A3
15
Dubrovinskaia, N.A., Dubrovinsky, L.S., Karlsson, A., Saxena, S.K., Sundman, B., “Experimental Study of Thermal Expansion and Phase Transformations in Iron-Rich Fe-Al Alloys”, Calphad, 23(1), 69-84 (1999) (Equi. Diagram, Experimental, 15) Kuchin, A.G., Medvedeva, I.V., Gaviko, V.S., Kazantsev, V.A., “Magnetovolume Properties of Y 2Fe17-xMx Alloys (M = Si or Al)”, J. Alloys Compd., 289, 18-23 (1999) (Crys. Structure, Experimental, 16) Salamakha, P., Sologub, O., Waerenborgh, J.C., Goncalves, A.P., Godinho, M., Almeida, M., “Systematical Investigation of the Y-Fe-Al Ternary System. Part 1. Single Crystal Studies of the YFexAl12-x Compound”, J. Alloys Compd., 296, 98-102 (2000) (Crys. Structure, Experimental, 16) Schaefer, W., Barbier, B., Halevy, I., “ThMn12 -Type Magnetic ErFe7Al5 and Non-Magnetic YFe7Al5 Studied by X-ray and Neutron Diffraction”, J. Alloys Compd., 303-304, 270-275 (2000) (Crys. Structure, Experimental, Magn. Prop., 7) Sikora, W., Schobinger-Papamantellos, P., Buschow, K.H.J., “Symmetry Analysis of the Magnetic Ordering in RFe4Al8 (R = La, Ce, Y, Lu and Tb) Compounds (II)”, J. Magn. Magn. Mater., 213, 143-156 (2000) (Calculation, Crys. Structure, Magn. Prop., 8) Waerenborgh, J.C., Salamakha, P., Sologub, O., Goncalves, A.P., Cardoso, C., Serio, S., Godinho, M., Almeida, M., “Influence of Thermal Treatment and Crystal Growth on the Final Composition and Magnetic Properties of the YFexAl12-x (4 x 4.2) Intermetallics”, Chem. Mater., 12, 1743-1749 (2000) (Crys. Structure, Experimental, Magn. Prop., 17) Ikeda, O., Ohnuma, I., Kainuma, R., Ishida, K., “Phase Equilibria and Stability of Ordered BCC Phases in the Fe-Rich Portion of hte Fe-Al System”, Intermetallics, 9, 755-761 (2001) (Equi. Diagram, Thermodyn., Experimental, 18) Knyazev, Yu.V., Kuchin, A.G., Kuz'min, Yu.I., “Optical Conductivity and Magnetic Parameters of the Intermetallic Compounds R 2Fe17-xMx (R = Y, Ce, Lu; M = Al, Si)”, J. Alloys Compd., 327, 34-38 (2001) (Crys. Structure, Experimental, Magn. Prop., Optical Prop., 23) Paixao, J.A., Silva, M.R., Waerenborgh, J.C., Concalves, A.P., Lander, G.H., Brown, P.J., Godinho, M., Burlet, P., “Magnetic Structures of MFe4+ Al8- (M = Lu, Y)”, Phys. Rev. B, 63B(5), 054410-1 - 054410-12 (2001) (Crys. Structure, Experimental, Magn. Prop., 29) Voronin, V.I., Berger, I.F., Kuchin, A.G., Sheptyakov, D.V., Balagurov, A.M., “Real Disordered Crystal Structure and Curie Temperature of Intermetallic Compounds Y 2Fe17-xMx (M = Si or Al)”, J. Alloys Compd., 315, 82-89 (2001) (Crys. Structure, Experimental, Magn. Prop., 17) Waerenborgh, J.C., Salamakha, P., Sologub, O., Goncalves, A.P., Serio, S., Godinho, M., Almeida, M., “Fe Moessbauer Spectroscopy Study of the AFexAl12-x Intermetallics (A = Y, Tm, Lu and U, 4 x 4.3)”, J. Alloys Compd., 318, 44-51 (2001) (Crys. Structure, Experimental, Moessbauer, 21) Waerenborgh, J.C., Salamakha, P., Sologub, O., Serio, S., Godinho, M., Goncalves, A.P., Almeida, M., “Y-Fe-Al Ternary System: Partial Isothermal Section at 1070 K Powder XRay Diffraction and Moessbauer Spectroscopy Study”, J. Alloys Compd., 323-324, 78-82 (2001) (Crys. Structure, Experimental, Moessbauer, 9) Cornish, L., Cacciamani, G., Saltykov, P., “Al-Y (Aluminium-Yttrium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Product ID: 20.14305.1.20, (2003) (Crys. Structure, Equi. Diagram, Assessment, 23) Kalvius, G.M., Wagner, F.E., Noakes, D.R., Schreier, E., Waeppling, R., Zimmermann, U., Schaefer, W., Kockelmann, W., Halevy, I., Gal, J., “Magnetic Behavior of YFexAl12-x”, Physica B, 326B(1-4), 460-464 (2003) (Experimental, Magn. Prop., Moessbauer, 7) Pisch, A., “Al-Fe (Aluminum-Iron)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2003) (Equi. Diagram, Crys. Structure, Assessment, 58) MSIT ®
Al–Fe–Y
16 Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(Al)
hP2 P63/mmc Mg
a = 269.3 c = 439.8
at 25°C, 20.5 GPa [Mas2]
(Al) < 660.452
cF4 Fm3m Cu
a = 404.96
at 25°C [Mas2]
(JFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, 13 GPa [Mas2]
( Fe) 1538 - 1394
cI2 Im3m W
a = 293.15
[Mas2]
( Fe) 1394 - 912
cF4 Fm3m Cu
a = 364.67
at 915°C [V-C2, Mas2, 1993Kat] dissolves up to 1.2 at.% Al
(Fe) < 912
cI2 Im3m W
a = 286.65
a = 286.64 to 289.59 a = 286.60 to 289.99 a = 286.60 to 290.12
pure Fe at 25°C [Mas2] dissolves up to 45.0 at.% Al at 1310°C 0 - 18.8 at.%Al, HT [1958Tay] 0 - 19.0 at.% Al, HT [1961Lih] 0 - 18.7 at.% Al, 25°C [1999Dub]
(Y) 1522 - 1478
cI2 Im3m W
a = 407
[Mas2]
(Y) < 1478
hP2 P63/mmc Mg
a = 364.82 c = 573.18
at 25°C [Mas2]
Fe4Al13 < 1160
mC102 C2/m Fe4Al13
a = 1552.7 to 1548.7 74.16 - 76.70 at.% Al [2003Pis] also called FeAl3 in the literature b = 803.5 to 808.4 c = 1244.9 to 1248.8 = 107.7 to 107.99° a = 1549.2 at 76.0 at.% Al [1994Gri] b = 807.8 c = 1247.1 = 107.69°
Fe2Al5 < 1169
oC24 Cmcm -
a = 765.59 b = 641.54 c = 421.84
MSIT®
at 71.5 at.% Al [1994Bur]
Landolt-Börnstein New Series IV/11A3
Al–Fe–Y
17
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
FeAl2 < 1156
aP18 P1 FeAl2
a = 487.8 b = 646.1 c = 880.0 = 91.75° = 73.27° = 96.89°
at 66.9 at.% Al [1993Kat]
J 1102 - 1232
cI16? -
a = 598.0
at 61 at.% Al [1993Kat]
FeAl < 1310
cP8 Pm3m CsCl
a = 289.48 to 290.5 a = 289.53 to 290.9 a = 289.81 to 291.01 a = 289.76 to 190.78
34.5 - 47.5 at.% Al [1961Lih] 36.2 - 50.0 at.% Al [1958Tay] 39.7 - 50.9 at.% Al [1997Kog] 500°C quenched in water room temperature
a = 579.30 to 578.86 a = 579.30 to 578.92
~24 - ~37 at.% Al [2001Ike] 23.1 - 35.0 at.% Al [1958Tay] 24.7 - 31.7 at.% Al [1961Lih]
Fe3Al < 547
cF16 Fm3m BiF3
Fe2Al9
mP22 P21/c Co2Al9
a = 869 b = 635 c = 632 = 93.4°
metastable 81.8 at.% Al [1993Kat]
FeAl6
oC28 Cmc21 FeAl6
a = 744.0 b = 646.3 c = 877.0 a = 744 b = 649 c = 879
metastable 85.7 at.% Al [1993Kat] [1998Ali]
FeAl4+x
t**
a = 884 c = 2160
(0 < x < 0.4) metastable [1998Ali]
YAl3 980 - 654(?)
hR36 R3m BaPb3
a = 620.4 0.2 c = 2118.4 0.7
[V-C2]
YAl3 < 645(?)
hP8 P63/mmc Ni3Sn
a = 627.6 0.2 c = 458.2 0.1
[V-C2] Metastable phase?
Y(FexAl1-x)2
cF24 Fd3m MgCu2
a = 783.4 to 768.9
0x0.41 [1975Dwi] x = 0 - 0.25, T = 800°C [2001Wae2]
a = 785.5 0.7
x = 0 [1989Gsc]
a = 778 to 786
x = 0 [1994Fol]
a = 388.4 0.2 b = 1152.2 0.4 c = 438.5 0.2
[V-C2]
YAl2 < 1485 YAl < 1130
Landolt-Börnstein New Series IV/11A3
oC8 Cmcm CrB
MSIT ®
Al–Fe–Y
18 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
Y3Al2 < 1100
tP20 P42/mnm Zr3Al2
a = 823.9 0.3 c = 764.8 0.4
[V-C2]
Y5Al3
hP16 P63/mcm Mn5Si3
a = 878.7 c = 643.5
Metastable [1987Ric] from recrystallized rapidly quenched alloys
Y2Al < 985
oP12 Pnma Co2Si
a = 664.2 2 b = 508.4 1 c = 946.9 2
[V-C2]
Y(Fe1-xAlx)2
cF24 Fd3m MgCu2
a = 735.5 to 751.0 a = 736.3
YFe2 < 1125
0x0.33 [1975Dwi] at x = 0 - 0.30 annealed at 1000°C [1977Mur] at x = 0 [V-C2]
YFe3 1350
hR36 R3m PuNi3
a = 513.7 c = 2461
[V-C2]
Y6(Fe1-xAlx)23
cF116 Fm3m Th6Mn23
a = 1208.4 a = 1208.4
at x = 0.09, refined at 250°C [1985Gan] at x = 0 [V-C2]
a = 850.1 to 856.6 c = 831.2 to 833.7 a = 852.13 to 852.61 c = 832.86 to 833.44
0 x 0.24 (Th2Zn17 at x > 0.24) [1998Kam] at x = 0.06 - 0.18, annealed at 950°C [1986Plu] at x = 0 - 0.1, annealed at 1300°C X-ray and neutron diffr. [2001Vor]
Y6Fe23 1300 Y2(Fe1-xAlx)17(HT)
hP38 P63/mmc Th2Ni17
a = 846.3 c = 828.2
Y2Fe17(HT) ? < T < 1400 Y2(Fe1-xAlx)17(RT)
Y2Fe17(RT)
MSIT®
hR19 R3m Th2Zn 17
a = 874.6 to 880.0 c = 1266.6 to 1274.9 a = 874.46 c = 1267.28 a = 860.4 to 872.4 c = 1256.8 to 1264.7 a = 846.0 c = 1241.0
at x = 0.0 [V-C2] 0 x 0.45 at 500°C [1972Zar] at x = 0.45 - 0.56, as cast [1976McN] at Y2Fe9Al8, T = 10 K [1998Che] at x = 0.23 - 0.41, annealed at 950°C [1986Plu] at x = 0 [V-C2]
Landolt-Börnstein New Series IV/11A3
Al–Fe–Y Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
* -1, Y(Fe1-xAlx)2 YFeAl
hP12 P63/mmc MgZn2
* -2, Y(FexAl1-x)12 YFe4Al8
tI26 I4/mmm ThMn12
YFe6Al6
* -3, YFe2Al10
Landolt-Börnstein New Series IV/11A3
oP52 Cmcm YbFe2Al10
Lattice Parameters [pm]
a = 536.5 to 540.2 c = 873.9 to 877.5 a = 541 c = 861 a = 534.1 c = 880.5 a = 541 c = 881 a = 872 c = 504 a = 872.2 c = 503.6 a = 874.0 c = 504.5
19 Comments/References
0.35 x 0.54 [1975Dwi] at x = 0.40 - 0.50 [1975Bus] at x = 0.5 [1973Zar] at x = 0.4, annealed at 1000°C [1977Mur] at x = 0.33 [1972Ryk] 0.257 x 0.58 at x = 0.33, annealed at 600°C [1966Zar] at x = 0.33 [1974Viv] at x = 0.33 [1976Bus]
a = 864.6 c = 499.2 a = 873.2 c = 501.8 a = 871.2 c = 503.6
at x = 0.5 [1980Fel]
a = 871.6 c = 502.4
at x = 0.5, T = 210 K neutron diffraction [1998Sch]
a = 869.83 c = 504.30 a = 868.7 c = 503.2
at x = 0.42 [1995Sch]
a = 882.6 to 871.6 c = 506.3 to 503.2
at x = 0.257 - 0.382 single crystal [2000Sal]
a = 864.67 to 876.04 c = 503.74 to 505.04
at x = 0.33 - 0.46, T = 800°C [2001Wae2]
a = 861.7 to 862.9 c = 503.1 to 504.0
at x = 0.58, T = 20 - 127°C [2000Sch]
a = 896.9 b = 1015.6 c = 901.8
[1998Thi]
a = 896.49 b = 1015.68 c = 901.13
at T = 800°C [2001Wae2]
at x = 0.5 [1988Che] at x = 0.5, annealed at 800°C [1988Che]
at x = 0.42 neutron diffraction [1995Sch]
MSIT ®
Al–Fe–Y
20
Al
Fig. 1: Al-Fe-Y. Isothermal section at 500°C
Data / Grid: at.%
(α Al)
YAl3
Axes: at.%
20
80
Fe4Al13 Fe2Al5
τ3
YAl2
FeAl2
40
60
τ2 60
40
FeAl Fe3Al
τ1 80
20
(α Fe)
Y
MSIT®
20
40
60
YFe3 80 YFe2 Y6Fe23
Y2Fe17
Fe
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn
21
Aluminium – Iron – Zinc Gautam Ghosh Literature Data Constitutional equilibria in the Al-Fe-Zn system is very important for the production of high quality Zn-coatings in steels by a process commonly known as hot-dip galvanizing. As a result, a large number of experimental studies have been carried out to determine the phase equilibria. The earlier results [1922Fue, 1934Fue, 1945May, 1947May, 1953Geb, 1953Ray, 1961Ren] on the phase equilibria were reviewed several times [1943Mon, 1952Han, 1961Phi, 1969Wat, 1976Mon]. [1953Ray] studied the solidification using about 150 ternary alloys, and also reported isothermal sections at 350 and 370°C. [1961Ren] investigated the phase equilibria in alloys containing up to 20 mass% Al and 20 mass% Fe. They reported isothermal sections at 600, 400°C and at room temperature. The most comprehensive study was carried out by [1970Koe] and [1971Koe]. They investigated a large number of alloys containing up to 60 mass% (Fe+Zn). The alloys were prepared using Armco-grade Fe and 99.99 mass% Al and Zn. The ternary alloys were prepared by adding either Fe or Zn to a master alloy of Fe:Al 50:50 or to pure Al. The solidification path and the isothermal sections were determined by means of thermal analysis, X-ray diffraction and microstructural investigations. They presented a reaction scheme, liquidus surface, nine isothermal sections in the temperature range of 250 to 700°C, and four temperature-composition sections. [1973Ure1] investigated the partial isothermal section at 450°C by means of metallography and electron microprobe analysis. They carried out equilibration experiments using solid Al-Fe intermetallic (FeAl, FeAl2, Fe2Al5, or Fe4Al13) and either liquid Zn or Zn-1.71Al (mass%) alloy. Prepared samples in evacuated capsules were held at 450°C for 800 h followed by quenching in iced water. These results were critically assessed by [1992Gho] and [1992Rag]. Recently, there has been a renewed interest in the phase equilibria, particularly the Zn corner around 450°C, due to very stringent quality control requirements of galvanized steel sheets for the automotive industry. As a result, recent studies are focused primarily in experimental determination [1990Che, 1992Per, 1994Tan, 1995Tan2, 1996Tan, 1997Gyu, 1997Uwa1, 1999Tan] and CALPHAD modeling [1991Bel, 1992Per, 1999Cos, 2001Gio, 2002Bai] of phase equilibria of the Zn corner in the temperature range of 450 to 470°C. Due to rapid interfacial reaction between steel and liquid Al-Zn alloys, the importance of metastable equilibria [1991Bel, 1992Per, 2002Bai], diffusion path [1992Per, 1998Ada, 1998Uch1, 1998Uch2, 2002Bai], and the mechanism of phase transformations [1994Lin, 1995Lin1, 1995Lin2, 1995Tan1, 1995Yam2, 1997Mcd, 1997Mor, 1997Ser, 1998Ada, 1998Uch1, 1998Uch2, 1998Yam, 2002Bai] during interfacial reaction have also been elucidated. [1990Che] prepared three ternary alloys using Al, Fe and Zn powders of unspecified purity. The final heat treatment of the alloys was annealing at 450°C for about 10 h. The phase equilibria were determined by XRD and SEM/EDX techniques. [1991Bel] determined the stable and metastable solubility limits of Fe in liquid (Zn) 447 to 480°C. [1992Per] determined the metastable and stable isothermal sections at 450°C based on the interfacial reaction studies between solid Al-Fe and liquid Al-Zn alloys. They used Al-Fe alloys containing 5, 29 and 36 at.% Al, and liquid Al-Zn alloys containing 0.12, 0.22, 0.39 and 11.2 at.% Al. Both short time (less than 30 min) and long time (1000 h) experiments were carried out. The phase compositions were determined by SEM/EDX technique. [1994Tan] reported an isothermal section of Zn corner at 470°C. Tang [1995Tan2, 1996Tan] reported the phase equilibria at 450°C by combining the results of [1990Che] and his experimental data of the Zn-corner. [1997Uwa1] prepared four ternary alloys by dry ball milling. They used elemental powders of following purity: 99.5% Al, 99.9+% Fe and 99.9% Zn. The ball milled powders were annealed at 300, 400 and 570°C for 3 h. They used DSC to study phase transformations, and XRD to identify the phases. Some of the controversial results of [1997Uwa1] have been the subject of extensive discussions [1997Tan, 1997Uwa2, 1998Tan, 1998Uwa]. [2000Tan] determined the Fe solubilities in dilute liquid Al-Zn alloys in the temperature range of 450 to 480°C. He prepared 16 ternary alloys containing up to 0.1 mass% Fe and up to 0.23 mass% Al using 99.5% pure Fe
Landolt-Börnstein New Series IV/11A3
MSIT ®
22
Al–Fe–Zn
and Al, and special high grade Zn. The final equilibrations of encapsulated samples were carried out at 450, 465 and 485°C for 40 h followed by water quenching. The phase equilibria information were extracted from SEM/EDX analysis. [2002Tan] re-investigated the phase equilibria of the Zn corner at 435°C using six ternary alloys. They were annealed at 450°C for 15 days, and composition of phases were determined by SEM-EDS analysis. [2002Bai] reported a calculated isothermal section at 450°C. These recent results have been reviewed by [2003Rag]. Binary Systems The Al-Fe, Al-Zn and Fe-Zn binary phase diagrams are accepted from [2003Pis], [2003Per] and [1982Kub], respectively. There are some differences between the presently accepted binary phase diagrams and those accepted by the previous investigators [1953Ray, 1970Koe, 1971Koe]. For example, [1970Koe] and [1971Koe] accepted an Al-Fe phase diagram in which all the order-disorder transitions involving (Fe), 1 and 2 phases were considered to be first order, whereas in this assessment, (Fe) 2 and 1 2 reactions have been considered to be second order [1982Kub] reflected by the absence of the corresponding two-phase fields. Furthermore, the Al-Fe phase diagram has undergone slight modification due to recently established congruent melting behavior of the Fe4Al13 phase [1986Len]. In the case of the Fe-Zn phase diagram, [1953Ray, 1970Koe] and [1971Koe] considered the phase to be stable between 672 and 620°C and the 1 phase to be stable below 640°C [1953Ray, 1970Koe, 1971Koe, 1973Ure1]. However, according to [1982Kub] the phase (which is the 1 phase as designated by the above authors) is stable below 665°C. It is worth mentioning that [1970Koe] and [1971Koe] convincingly established the phase at temperatures above the 1 phase field near the Zn corner, but later on [1973Ure1] failed to identify the phase above the 1 phase field. Also, according to [1982Kub], the and phases react to form the 1 phase at 550°C. This feature was also absent in the Fe-Zn phase diagram accepted by the previous studies [1953Ray, 1970Koe, 1971Koe, 1973Ure1]. Very recent study of solid-state equilibria of Zn rich alloys [2001Mit], and thermodynamic modeling of phase equilibria [2000Reu, 2001Su] are consistent with the Fe-Zn phase diagram assessed by [1982Kub]. In the Al-Zn phase diagram, the phase designated by [1953Ray, 1970Koe, 1971Koe, 1973Ure1] is identical to (Al) in the phase diagram given by [1983Mur]. All these features are taken into account in this critical assessment of phase equilibria. Solid Phases Available data suggest that the solubility of Zn in (Fe,Al) is a function of time of heat treatment at 450°C, with less Zn after shorter time compared to longer time. For example, [1990Che] gives 2 mass% Zn after 10 h at 450°C, while [1992Per] gives 2.26 mass% Zn after less than 30 min at 450°C and 4.85 mass% Zn after 1000 h at 450°C. The equilibrium solubility of Zn in Fe4Al13 at 450°C are 7 mass% [1973Ure1], 5.5 mass% [1990Che], 7.61 mass% [1992Per], while under metastable equilibrium Fe4Al13 can dissolve up to 13.92 mass% [1992Per] and 15.2 mass% [1997Gyu]. [1953Ray] noted that the X-ray diffraction pattern of Zn containing Fe4Al13 is slightly different from that of pure Fe4Al13 which might be due to the slight structural alteration caused by the non-random occupation of the Zn atoms. [1992Per] reported that the presence of Zn in FeAl2 is hardly detectable. The solubility of Zn in Fe2Al5 () has been determined several times by reacting Fe with liquid Al-Zn bath containing varying amounts of Al [1971Ghu, 1973Har, 1973Ure1, 1973Ure2, 1984Nit, 1990Che, 1991Sai, 1992Per, 1997Gyu]. Available data fall in the range of 11 to 23 mass% Zn, and also show a systematic trend that the Zn-content in Fe2Al5 () is a function of reaction time. Due to rapid interdiffusion, the data after short time reaction show higher solubility of Zn in Fe2Al5 compared to long time experiments. For example, [1992Per] found 22.87 mass% Zn in Fe2Al5 after reaction at 450°C for less than 30 min compared to 18.7 mass% Zn after reaction at 450°C for 1000 h. [1971Ghu] noted a scatter of 14 to 17 mass% Zn in Fe2Al5 after reaction at 600°C for 10s. It is important to note that while short time reaction data is relevant to industrial galvanizing process, long time data is appropriate to construct the equilibrium phase diagram. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn
23
Accordingly, we have accepted the solubility of 18.7 mass% Zn (11 at.%) at 450°C [1992Per] as equilibrium value. X-ray diffraction and density measurement show that Zn atoms reside on the Fe site for up to 6.7 at.% Zn giving the formula Fe4Zn10Al, and beyond this composition Zn atoms also reside on the Al sites giving the formula Fe4Zn9Zn2 [2001Koe]. [1973Ure1] reported a solid solubility of 3.6 mass% Al in the phase (FeZn10) at 450°C, which is in qualitative agreement with that of [1956Hor]. On the other hand, [1990Che] and [1992Per] reported solid solubilities of 2.8, 3.71, and 1.84 mass% Al at 450°C. Since the latter value was obtained after long time (1000 h) heat treatment, it is considered as equilibrium solid solubility while other values correspond to metastable equilibria. The phase (FeZn13 ) dissolves 0.78 mass% Al at 450°C [1992Per], but [1961Ren] gives a much lower value of 0.2 mass%. The solid solubilities of Al in and 1 phase at 450°C are similar to that in phase [1992Per]. On the other hand, Tang’s [1996Tan] isothermal section at 450°C show much higher solubility of Al in these two phases which may correspond to industrial galvanizing conditions. [1992Per] reported two Phases, 1 (denoted as 2 by [1992Per]) and 2 (denoted as 3 by [1992Per]), after equilibration for 1000 h at 450°C. However, [1973Ure1] did not detect any 2 after 800 h equilibration at 450°C. On the hand, [1995Yam2] reported continuous solid solubility ( 1) and [1996Tan] reported continuous solid solubility ( ´) in the isothermal sections at 440 and 450°C, respectively. It is possible that these conditions are realized during galvanizing process, and may not represent equilibrium. Later, [1998Yam] synthesized single phase alloys corresponding to 2 and 3 compositions of [1992Per], and diffusion annealing (conditions are not specified) of mechanically pressed 2 and 3 did not show any evidence of continuous solid solubility. Even though the crystallographic data of 2 is lacking, available results suggest that it may be a ternary phase. The details of the crystal structures and lattice parameters of the solid phases are listed in Table 1. Invariant Equilibria Based on the results of [1970Koe] and [1971Koe], the reaction scheme is summarized in Fig. 1. A number of changes have been made to comply with the binary phase diagrams accepted here. The reaction scheme proposed by [1970Koe] contained fourteen invariant reactions. However, three invariant reactions proposed to occur at 485, 440 and 320°C [1970Koe, 1971Koe] are not considered in Fig. 1 as they are not compatible with the presently accepted binary phase diagrams. The assessed reaction scheme is consistent with all the phase diagram information available until now. [1961Ren] proposed a ternary U type invariant reaction L+FeAl2 +Fe2Al5 at 592°C; however, subsequent detailed investigations by [1970Koe, 1971Koe] and [1973Ure1] failed to detect this reaction. Liquidus Surface Figure 2 shows the liquidus surface from 20 to 70 mass% Al and 0 to 40 mass% Zn and Fig. 3 shows the liquidus surface of the Zn corner, both according to [1970Koe] and [1971Koe]. Results of solidification studies of Zn rich ternary alloys by [1945May, 1947May] and [1962May] and of Al/Zn rich alloys [1953Geb] agree quantitatively with those of [1970Koe] and [1971Koe]. Isothermal Sections Figures 4, 5 and 6 show the isothermal sections at 700, 575 and 500°C, respectively, after [1970Koe] and [1971Koe]. Figure 7 shows the isothermal section of the Zn corner at 500°C [1970Koe, 1971Koe]. Figures 8 and 9 show partial isothermal section at 470 [1994Tan] and 460°C [2000Tan], respectively, depicting the solubility limits of Fe in liquid-Zn with respect to (FeZn13), (FeZn10), and (Fe2Al5) phases. The isothermal section at 450°C has been investigated several times. There is substantial agreement between the earlier results of [1970Koe], [1971Koe] and [1973Ure1]. Recent significant results are due to [1990Che, 1992Per, 1995Tan2, 1996Tan]. Except for [1992Per] and [1996Tan], others did not consider 1 phase in the 450°C isothermal section. Figure 10 shows the isothermal section at 450°C [1992Per]. Figure 11 shows the isothermal section of Zn corner depicting the phase fields involving liquid, , , 1 and 2 [1992Per]. [2002Tan] labelled as 2 phase T. Despite qualitative agreement between the results of Landolt-Börnstein New Series IV/11A3
MSIT ®
24
Al–Fe–Zn
[1992Per], [1996Tan] and [2002Tan] at 450°C, the isothermal section of [1992Per] is preferred because the authors used much longer annealing time. Figure 12 shows the isothermal section at 450°C depicting the saturation limits of Fe with respect to , , 2 and phase in liquid Zn [1996Tan]. Contrary to the suggestion of [1962Cam] that the solubility of Fe should decrease with Al content in liquid Zn, [1973Ure1] proposed that the solubility of Fe in liquid Zn at 450°C is 0.029 mass%, irrespective of the Al content. In fact, [1991Bel] showed that when phase is in equilibrium with liquid Zn, indeed the Fe solubility decreases with increasing Al content in liquid Zn which is seen in Figs. 8, 9 and 12. Thermodynamic calculations also predict a similar behavior [2002Bai]. The isothermal section of the Zn-corner at 400°C [1970Koe, 1971Koe] is shown in Fig. 13. The Fe4Al13-Al-Zn partial isothermal sections at 350, 330, 300 and 250°C are shown in Figs. 14, 15, 16, 17, respectively according to [1970Koe] and [1971Koe]. A number of adjustments have been made in the isothermal sections in order to comply with the binary phase diagrams. [1961Ren] studied the isothermal sections of the Zn corner with up to about 20 mass% (Fe+Al) at 600°C, 450°C and room temperature. At 600°C, [1961Ren] observed three-phase fields L+ +FeAl2 and L+Fe2Al5+FeAl2, and proposed a ternary U type invariant reaction L+FeAl2 +Fe2Al5 at 592°C. However, more detailed investigations by [1970Koe, 1971Koe] and [1973Ure1] failed to observe these features. The partial isothermal section at 450°C given by [1961Ren] agrees qualitatively with that of [1973Ure1], but the exact locations of the phase boundaries differ significantly. Because of these reasons, the results of [1961Ren] are not accepted here. Temperature – Composition Sections Figures 18, 19, 20 and 21 show isopleths at 30, 90, 95 and 98 mass% Zn, respectively [1970Koe, 1971Koe]. In Fig. 18, several changes have been made to comply with the accepted Al-Zn phase diagram. Thermodynamics [1995Yam1] reported the activity coefficient of Al in liquid Al-Zn alloys containing up to 10 mass% Zn, and in liquid Al-Fe-Zn alloys containing up to 1 mass% Al at 480°C. [1995Yam2] determined the chemical potential of Al in liquid Zn, in equilibrium with Fe4Al13, (Fe2Al5), 2, (FeZn10), and (FeZn13) in the temperature range of 432 to 510°C. [1971Ghu] reported that the heat formation of Fe(Al,Zn)3 is much more negative compared to the heat of formation of Fe4Al13 and Fe2Al5 phases; however, the actual values reported by [1971Ghu] are very doubtful. [2000Tan] reported that the solubility product of Fe2Al5 in liquid Zn can be expressed as ln(mass% Al)5(mass% Fe)2 = 28.1 - 33066/T where T is the temperature in Kelvin. Besides, [2000Tan] has also discussed a procedure to calculate the solubility limits of Fe in liquid Zn with respect to saturation of , and phase. [1991Bel] reported solubility products of Fe4Al13, Fe2Al5, FeAl2, FeAl and FeZn 13 in liquid Zn. Using the experimental solubility data, [2001Gio] has derived the Gibbs energy of formation of Fe2Al5Znx (). [2002Feu] measured the standard enthalpy of formation of phase at Fe0.07Zn0.93 using solution calorimetry technique. Several attempts have been made to calculate phase diagrams by CALPHAD method [1991Bel, 1992Per, 1999Cos, 2001Gio, 2002Bai]. Of particular interest is the prediction of solubility of Fe and Al in liquid Zn around 450°C, and also the diffusion path during hot-dip galvanizing process. [1991Bel] calculated metastable solubilities in liquid-Zn with respect to +Fe2Al5, +FeAl, +Fe4Al13 and +FeAl2 saturations at 447 and 477°C, and did not consider the phase. On the other hand [1992Per] calculated the solubility of Fe in liquid-Zn at 450°C considering all binary phases, and found a slightly higher solubility of Fe in liquid-Zn compared to [1991Bel] due to participation of the phase. [1999Cos] calculated the 465°C isothermal section, but only the Zn-corner to understand the limiting factor controlling solubility of Fe in liquid Zn. They did not consider any ternary interaction parameter in the liquid phase and also the ternary solubility of Fe-Zn intermetallics. Nonetheless, the calculated activity coefficients of Al in liquid Zn-0.01 mass% Fe-xAl alloys are in good agreement with the experimental data of Yamaguchi et al. [1995Yam1, 1995Yam2]. Even though their calculated solubility limit of Fe2Al5 is in
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn
25
good agreement with experiment, the calculated three phase equilibrium L+Fe2Al5+Fe4Al13 differs significantly from the experimental data [1995Yam2, 1998Yam]. [2002Bai] calculated the entire isothermal section at 450°C, and it appears that they overestimated the solid solubility of Al in , and 1 phase compared to the experimental data of [1992Per]. Also, they did not consider the 2 phase. Nonetheless, their calculation clearly shows a decrease in solubility of Fe in liquid Zn when it is in equilibrium with the phase (Fe2Al5). Miscellaneous The solubility of Fe in a liquid Zn-4Al (mass%) alloy, in the temperature range of 400 to 675°C, was determined by [1963Fri]. The solubility can be expressed as log(mass% Fe) = 3.6359 - 5149/T log(at.% Fe) = 3.6825 - 5150/T where T is the temperature in K. Additions of Al to a liquid Zn bath inhibit the reaction between solid Fe and liquid Zn during the normal galvanizing process. It is believed that Al causes the formation of an inhibition layer, consisting of Fe2Al5, at the substrate/coating interface [1995Tan1]. However, detailed experiments using TEM/SEM/XRD techniques clearly show that the inhibition layer actually consists of Fe2Al5 and Fe4Al13. The details of the reactions and the formation sequences of the different binary intermetallic phases during the hot dip galvanizing process have been reported by [1965Sou, 1971Ghu, 1973Har, 1973Ure2, 1975Gut, 1984Nit, 1991Sag, 1995Lin1, 1995Lin2, 1995Tan1, 1997Mcd, 1997Ser, 1998Uch1, 1998Uch2]. Addition of Si also suppresses the rapid exothermic reaction between liquid Al-Zn and Fe by forming a solid reaction layer [1989Sel] which acts as a diffusion barrier. A comprehensive review of physical metallurgy of the galvanizing process has been presented by Marder [2000Mar]. [1998Akd] proposed that the value of activity coefficient of Al in (Fe,Al,Zn) alloys has a strong influence on the formation and growth kinetics of interfacial diffusion layer. Besides, [2002Bai] compiled the diffusion data in , , and 1 phases which were then used to model the mobility of components in these phases within CALPHAD formalism. [1977Sho] investigated the effect of pressure on the reaction kinetics between solid Fe and liquid Zn-1.5Al (mass%) at 501°C. An applied pressure was found to cause the intermetallic compounds to become unstable and change the overall reaction rate from linear to non-linear. The stability of phase, compared to other phases, under pressure is markedly affected by the presence of the Al in the melt. References [1922Fue] [1924Fus] [1934Fue] [1943Mon] [1945May] [1947May] [1952Han]
[1953Geb]
Landolt-Börnstein New Series IV/11A3
Fuess, V., “Aluminium-Zinc-Iron” in Metallography of Aluminium and its Alloys (in German), 157-159 (1922) (Equi. Diagram, Review, 2) Fuss, V., “On the Constitution of Ternary Alloys of Aluminium” (in German), Z. Metallkd., 16, 24-25 (1924) (Equi. Diagram, Experimental, 5) Fuess, V., “Aluminium-Zinc-Iron” in “Metallography of Aluminium and its Alloy” (in German), 157-159 (1934) (Equi. Diagram, Review, 1) Mondolfo, L.F., “Aluminium-Iron-Zinc”, in “Metallography of Aluminum Alloys”, John Wiley and Sons, Inc., New York, 98-99 (1943) (Equi. Diagram, Review, 1) Mayer, A., “Investigation of the Ternary Zinc-Aluminium-Iron System” (in Italian), Metallurgia Italiana, 37, 95-98 (1945) (Equi. Diagram, Experimental, 33) Mayer, A., “The Ternary System: Zinc-Aluminium-Iron” (in Italian), Gazz. Chim. Ital., 77, 55-66 (1947) (Equi. Diagram, Experimental) Hanemann, H., Schrader, A., “Aluminium-Zinc-Iron” in “Ternary Alloys of Aluminium” (in German), Atlas Metallographicus, Verlag Stahleisen, Düsseldorf, 3(2), 157-159 (1952) (Review, 1) Gebhardt, E., “Investigation on the Ternary Aluminium-Iron-Zinc” (in German), Z. Metallkd., 44, 206-211 (1953) (Equi. Diagram, Experimental, 18)
MSIT ®
26 [1953Ray]
[1956Hor]
[1961Phi] [1961Ren] [1962Cam]
[1962May] [1963Fri] [1965Sou] [1969Wat] [1970Koe] [1971Ghu]
[1971Koe] [1973Har] [1973Ure1]
[1973Ure2]
[1975Gut]
[1976Mon] [1977Sho] [1982Kub] [1983Mur] [1984Nit]
[1986Len]
MSIT®
Al–Fe–Zn Raynor, G.V., Faulkner, C.R., Noden, J.D., Harding, A.R., “Ternary Alloys Formed by Aluminium, Transitional Metals and Divalent Metals”, Acta Met., 1, 629-648 (1953) (Equi. Diagram, Experimental, *, 32) Horstmann, D., Malissa, H., “Electrolytic Isolation of Intermetallic Fe-Zn Compounds and Determination of the Solubility of Several Metals in These Compounds” (in German), Arch. Eisenhüttenwesen, 27, 423-428 (1956) (Experimental, 4) Phillips, H.W.L., “Al-Fe-Zn” in “Equilibrium Diagrams of Aluminium Alloy Systems”, The Aluminium Development Association, London, 97 (1961) (Equi. Diagram, Review, 1) Rennhack, E.H., “Zinc-Rich Corner of the Zn-Fe-Al System”, Trans. AIME, 221, 775-779 (1961) (Equi. Diagram, Experimental, *, 13) Cameron, D.I., Ormay, M.K., “The Effect of Agitation, Cooling, and Al on the Alloying in Hot-Dipping in Zn”, 6th Int. Conf. on Hot Dip Galvanizing, Interlaken, Zinc Development Association, London, 276-311 (1962) (Experimental) Mayer, A., Morandi, F., “Investigation of Zn-Al-Fe Alloys” (in Italian), Gazz. Chim. Ital., 92, 1005-1020 (1962) (Experimental, 15) Friebel, V.R., Lantz, W.J., Roe, W.P., “Liquid Solubilities of Selected Metals in Zinc-4% Aluminium”, Trans. ASM, 56, 90-100 (1963) (Experimental, 12) Southin, R.T., Wright, D.A., “Fe2Al5 and FeSi in Zinc Alloys”, J. Inst. Metals, 93, 357-358 (1965) (Experimental, 12) Watanabe H., Sato E., “Phase Diagrams of Aluminum-Base Systems” (in Japanese), Keikinzoku, 19(11), 499-535 (1969) (Equi. Diagram, Review, 232) Koester, W., Goedecke, T., “The Fe-Al-Zn Ternary System” (in German), Z. Metallkd., 61, 649-658 (1970) (Equi. Diagram, Experimental, #, *, 13) Ghuman, A.R.P., Goldstein, J.I., “Reaction Mechanisms for the Coatings Formed During Hot Dipping of Fe in 0-10% Al-Zn Baths at 450-700°C”, Metall. Trans., 2, 2903-2914 (1971) (Experimental, 18) Koester, W., Goedecke, T., “The Iron-Aluminium-Zinc Ternary System”, Proc. 9 th Int. Conf. Hot Dip Galvanizing, 128-139 (1971) (Equi. Diagram, Experimental, #, *, 13) Harvey, G.J., Mercer, P.D., “Aluminium-rich Alloy Layers Formed During the Hot Dip Galvanizing of Low Carbon Steel”, Metall. Trans., 4, 619-621 (1973) (Experimental, 8) Urednicek, M., Kirkaldy, J.S., “An Investigation of the Phase Constitution of Iron-Zinc-Aluminium at 450°C”, Z. Metallkd., 64, 419-427 (1973) (Equi. Diagram, Experimental, #, *, 21) Urednicek, M., Kirkaldy, J.S., “Mechanism of Iron Attack Inhibition Arising from Additions of Aluminium to Liquid Zn(Fe) during Galvanizing at 450°C”, Z. Metallkd., 64, 899-910 (1973) (Experimental, 26) Guttman, H., Niessen, P., “Galvanizing Si Steels in Al-containing Baths”, Proc. Seminar Galvanizing Si-containing Steels, Int. Lead Zinc Research Organisation, Inc. New York, USA, 198-218 (1975) (Experimental, 10) Mondolfo, L.F., “Aluminium-Iron-Zinc” in Metallography of ALuminium Alloys, John Wiley and Sons, Inc., New York, 98-99 (1976) (Review, 1) Short, N.R., Mackowiak, J., “The Effect of Pressure on the Reactions between Fe(s)-Zn: 1.5% Al(l) at 501°C”, Corrosion Science, 17, 397-404 (1977) (Experimental, 13) Kubaschewski, O., “Iron-Aluminium” and “Iron-Zinc”, in “Iron-Binary Phase Diagrams”, Springer Verlag, Berlin, 5-9 and 172-175 (1982) (Equi. Diagram, Review, #, 26, 13) Murray, J.L., “The Al-Zn (Aluminum-Zinc)”, Bull. Alloy Phase Diagrams, 4(1) 55-73 (1983) (Equi. Diagram, Review, #, 194) Nitto, H., Yamazaki, T., Morita, N., Yabe, K., Bandooo, S., “Effect of Aluminium in Zinc on Alloying of Zinc Coating of Galvanized Steel” (in Japanese), Tetsu-to-Hagane, 70, 1719-1726 (1984) (Experimental, 20) Lendvai, A., “Phase Diagram of Al-Fe Sytem up to 45 mass% Iron”, J. Mater. Sci. Lett., 5, 1219-1220 (1986) (Equi. Diagram, Experimental, #, *, 7) Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn [1989Sel]
[1990Che]
[1991Bel]
[1991Sag]
[1991Sai]
[1992Gho]
[1992Per]
[1992Rag]
[1994Lin]
[1994Tan] [1995Lin1]
[1995Lin2] [1995Tan1] [1995Tan2] [1995Yam1]
[1995Yam2]
[1996Tan] [1997Gyu]
Landolt-Börnstein New Series IV/11A3
27
Selverian, J.H., Marder, A.R., Notis, M.R., “The Effects of Silicon on the Reaction Between Solid Iron and Liquid 55 wt.% Al-Zn Baths”, Metall. Trans. A, 20A(3), 543-555 (1989) (Experimental, 16) Chen, Z.W., Sharp, R.M., Gregory, J.T., “Fe-Al-Zn Ternary Phase Diagram at 450°C”, Mater. Sci. Technol., 6(12), 1173-1176 (1990) (Assessment, Equi. Diagram, Experimental, #, *, 16) Belisle, S., Leson, V., Gagne, M., “The Solubility of Iron in Continuous Hot-Dip Galvanizing Baths”, J. Phase Equilib., 12(3), 259-265 (1991) (Equi. Diagram, Experimental, Thermodyn., 7) Sagiyama, M., Inagaki, J.-I., Morita, M., “Fe-Zn Alloying Behavior and the Coating Microctructure of Galvannealed Steel Sheets”, NKK Technical Review (Japan), (63), 38-45 (1991) (Abstract, Experimental, 14) Saito, M., Uchida, Y., Kittaka, T., Hirose, Y., Hisamatsu, Y., “Formation Behavior of Alloy Layer in Initial-Stages of Galvanizing” (in Japanese), Tetsu to Hagane, 77(7), 947-954 (1991) (Experimental, 7) Ghosh, G., “Aluminium-Iron-Zinc”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.17658.1.20, (1992) (Crys. Structure, Equi. Diagram, Assessment, 27) Perrot, P., Tissier, J.C., Dauphin, J.Y., “Stable and Metastable Equilibria in the Fe-Zn-Al System at 450°C”, Z. Metallkd., 83(11), 786-790 (1992) (Calculation, Equi. Diagram, Experimental, #, *, 12) Raghavan, V., “The Al-Fe-Zn (Aluminium-Iron-Zinc) System”, in Phase Diagrams of Ternary Iron Alloys, Part 6A, Indian Institute of Metals, Calcutta, 215-223 (1992) (Equi. Diagram, Review, 24) Lin, C.S., Meshii, M., “The Effect of Steel Chemistry on The Formation of Fe-Zn Intermetallic Compounds of Galvanneal-Coated Steel Sheets”, Metall. Mater. Trans. B, 25B(5), 721-730 (1994) (Experimental, Kinetics, 31) Tang, N., “Comment on Fe-Al-Zn (Iron-Aluminium-Zinc)”, J. Phase Equilib., 15(3), 237-238 (1994) (Theory, 10, #, *, 10) Lin, C.S., Meshii, M., Cheng, C.C., “Microstructural Characterization of Galvanneal Coatings by Transmission Electron-Microscopy”, ISIJ Int., 35(5), 494-502 (1995) (Experimental, Kinetics, 43) Lin, C.S., Meshii, M., Cheng, C.C., “Phase Evolution in Galvanneal Coatings on Steel Sheets”, ISIJ International, 35(5), 503-511 (1995) (Experimental, Kinetics, 28) Tang, N., “Modeling Al Enrichment in Galvanized Coatings”, Metall. Mater. Trans. A, 26A(7), 1699-1704 (1995) (Theory, Kinetics, 23) Tang, N., “Refined 450°C Isotherm of Zn-Fe-Al Phase Diagram”, Mater. Sci. Technol., 11(9), 870-873 (1995) (Equi. Diagram, Experimental, *, 23) Yamaguchi, S., Fukatsu, N., Kimura, H., Kawamura, K, Iguchi, Y., O-Hashi, T., “Development of Al Sensor in Zn Bath for Continuous Galvanizing Processes” in Proc. Galvatech’95, ISS-AIME, Warrendale, Pa, 647-655 (1995) (Experimental, Thermodyn., *, 12) Yamaguchi, S., Makino, H., Sakatoku, A., Iguchi, Y., “Phase Stability of Dross Phases in Equilibrium with Liquid Zn Measured by Al Sensor” in Proc. Galvatech’95, ISS-AIME, Warrendale, Pa, 787-794 (1995) (Experimental, Thermodyn., *, 11) Tang, N.-Y., “450°C Isotherm of Zn-Fe-Al Phase Diagram Update”, J. Phase Equilib., 17(5), 396-398 (1996) (Equi. Diagram, Experimental, #, *,13) Gyurov, S., “The Reaction Between Solid Iron and Liquid Zn-Al Baths”, Z. Metallkd., 88(4), 346-352 (1997) (Equi. Diagram, Experimental, Kinetics, 33)
MSIT ®
28 [1997Mcd]
[1997Mor]
[1997Ser]
[1997Tan]
[1997Uwa1]
[1997Uwa2] [1998Ada]
[1998Akd]
[1998Tan]
[1998Uch1]
[1998Uch2]
[1998Uwa] [1998Yam]
[1999Cos]
[1999Tan] [2000Mar] [2000Reu] [2000Tan] [2001Gio]
MSIT®
Al–Fe–Zn McDevitt E., Morimoto Y., Meshii M., “Characterization of the Fe-Al Interfacial Layer in a Commercial Hot-Dip Galvanized Coating”, ISIJ Int., 37(8), 776-782 (1997) (Experimental, 24) Morimoto Y., McDevitt E., Meshii M., “Characterization of the Fe-Al Inhibition Layer Formed in the Initial Stages of Hot-Dip Galvannealing”, ISIJ Int., 37(9), 906-913 (1997) (Experimental, 28) Sere, P.R., Culcasi, J.D., Elsner, C.J, Di Sarli, A.R., “Factors Affecting the Hot-dip Zinc Coatings Structure” (in Spanish), Rev. de Metall., 33(6), 376-381 (1997) (Experimental, Kinetics, 11) Tang, N.-Y., “Discussion of “Kinetics and Phase Transformation Evaluation of Fe-Zn-Al Mechanically Alloyed Phases”, Metall. Mater. Trans. A, 28A(11), 2433-2434 (1997) (Theory, 11) Uwakwen, O.N.C., Liu, Z., “Kinetics and Phase Transformation Evaluation of Fe-Zn-Al Mechanically Alloyed Phases”, Metall. Mater. Trans. A, 28A(3), 517-525 (1997) (Equi. Diagram, Experimental, *, 26) Uwakwen, O.N.C., Liu, Z., “Authors’ Reply”, Metall. Mater. Trans. A, 28A(11), 2434-2435 (1997) (Theory, 7) Adachi Y., Arai M., “Transformation of Fe-Al Phase to Fe-Zn Phase on Pure Iron During Galvanizing Layer”, Mater. Sci. Eng. A, 254(1-2), 305-310 (1998) (Crys. Structure, Experimental, 8) Akdeniz, M.V., Mekhrabon, A.O., “The Effect of Substitutional Impurities on the Evolution of Fe-Al Diffusion Layer”, Acta Mater., 46(4), 1185-1192 (1998) (Calculation, Thermodyn., 55) Tang, N.-Y., “Discussion of “Kinetics and Phase Transformation Evaluation of Fe-Zn-Al Mechanically Alloyed Phases”, Metall. Mater. Trans. A, 29A(10), 2643-2644 (1998) (Equi. Diagram, Theory, 9) Uchida Y., Andoh A., Komatsu A., Yamakawa K., “Changing Process from Center Dot Fe-Zn Phase to Al-Fe Intermetallic Compounds in Molten Zn-5mass%Al Alloy Bath” (in Japanese), Tetsu to Hagane, 84(9), 632-636 (1998) (Experimental, 6) Uchida Y., Andoh A., Komatsu A., Yamakawa K., “Changing Process from Center Dot Fe-Zn Phase to Al-Fe Intermetallic Compounds in Molten Zn-5mass%Al Alloy Bath” (in Japanese), Tetsu to Hagane, 84(9), 637-642 (1998) (Experimental, 4) Uwakwen, O.N.C., Liu, Z., “Authors’ Reply”, Metall. Mater. Trans. A, 29A(10), 2644-2645 (1998) (Equi. Diagram, Theory, 5) Yamaguchi, S., “Thermochemical Stability and Precipitation Behavior of Dross Phases in CGL Bath” in Proc. Galvatech’98, Chiba, Japan, The Iron and Steel Institute of Japan, 84-88 (1998) (Experimental, Thermodyn., *, 8) Costa e Silva, A., Avillez, R.R., Marques, K., “A Preliminary Assessment of the Zn-rich Corner of the Al-Fe-Zn System and Its Implications in Steel Coating”, Z. Metallkd., 90(1), 38-43 (1999) (Calculation, Equi. Diagram, Thermodyn., *, 25) Tang, N.-Y., “Characteristics of Continuous-Galvanizing Baths”, Metall. Mater. Trans. B., 30(1), 144-148 (1999) (Equi. Diagram, *, 26) Marder, A.R., “The Metallurgy of Zinc-Coated Steel”, Prog. Mater. Sci., 45, 191-271 (2000) (Equi. Diagram, Phys. Prop., Review, 188) Reumont, G., Perrot, P., Fiorani, J.M., Hertz, J., “Thermodynamic Assessment of the Fe-Zn System”, J. Phase Equilib., 21(4), 371-378 (2000) (Thermodyn., *, 26) Tang, N.-Y., “Determination of Liquid-Phase Boundaries in Zn-Fe-Mx Systems”, J. Phase Equilib., 21(1), 70-77 (2000) (Equi. Diagram, Experimental, Thermodyn., #, *, 29) Giorgi, M.-L., Guillot, J.-B., Nicolle, R., “Assessment of the Zinc-Aluminium-Iron Phase Diagrams in the Zinc-Rich Corner”, Calphad, 25(3), 461-474 (2001) (Equi. Diagram, Thermodyn., *, 36)
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn [2001Koe]
[2001Mit] [2001Su] [2002Bai] [2002Feu]
[2002Tan] [2003Per]
[2003Pis]
[2003Rag]
29
Koester, M., Schuhmacher, B., Sommer, D., “The Influence of the Zinc Content on the Lattice Constants and Structure of the Intermetallic Compound Fe2Al5”, Steel Res., 72(9), 371-375 (2001) (Crys. Structure, Experimental, 29) Mita, K., Ikeda, T., Maeda, M., “Phase Diagram Study of Fe-Zn Intermetallics”, J. Phase Equilib., 22(2), 122-125 (2001) (Experiment, Equi. Diagram, #, *, 14) Su, X., Tang, N.-Y., Toguri, J.M., “Thermodynamic Evaluation of the Fe-Zn System”, J. Alloys Compd., 325(9), 129-136 (2001) (Thermodyn., *, 49) Bai, K., Wu, P., “Assessment of the Zn-Fe-Al System for Kinetic Study of Galvanizing”, J. Alloys Compd., 347, 156-164 (2002) (Equi. Diagram, Thermodyn., Kinetics, *, 40) Feutelais, Y., Legendre, B., de Avillez, R. R., “Standard Enthalpy of Formation of the -Phase in the Fe-Zn System at 298 K”, J. Alloys Compd., 346, 1-2 (2002) (Experimental, Thermodyn., Kinetics, *, 20) Tang, N.Y., Su, P., “Assessment of the Zn-Fe-Al System for Kinetic Study of Galvanizing”, J. Alloys Comp., 347, 156-164 (2002) (Equi. Diagram, Experimental, #, *, 16) Perrot, P., “Al-Zn (Aluminium-Zinc)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2003) (Equi. Diagram, Assessment, Crys. Structure, 41) Pisch, A., “Al-Fe (Aluminum-Iron)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2003) (Equi. Diagram, Assessment, Crys. Structure, 58) Raghavan, V., ”Al-Fe-Zn (Aluminum-Iron-Zinc)”, J. Phase Equilib., 24, 546-550 (2003) (Equi. Diagram, Review, *, 33)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype
(Al)
cF4 Fm3m Cu
a = 404.88 a = 403.52 a = 403.29 a = 403.14
pure Al at 24°C [V-C] at 63.0 at.% Zn and 360°C [1983Mur] at 64.8 at.% Zn and 360°C [1983Mur] at 70.1 at.% Zn and 360°C [1983Mur]
(Fe)
cI2 Im3m W
a = 286.65
pure Fe at 20°C [V-C]
(Zn)
hP2 P63/mmc Mg
a = 266.46 c = 494.61
pure Zn at 22°C [V-C]
1, Fe3Al 552.5
cF16 Fm3m BiF3
a = 578.86 to 579.3
[2003Pis], solid solubility ranges from 22.5 to 36.5 at.% Al
2, FeAl 1310
cP2 Pm3m CsCl
a = 289.76 to 290.78 [2003Pis], at room temperature solid solubility ranges from 22.0 to 54.5 at.% Al
J, Fe2Al3 1102 - 1232
cI16?
a = 598.0
Landolt-Börnstein New Series IV/11A3
Comments/References
[2003Pis], solid solubility ranges from 54.5 to 62.5 at.% Al
MSIT ®
Al–Fe–Zn
30 Phase/ Temperature Range [°C]
Pearson Symbol/ Lattice Parameters [pm] Space Group/ Prototype
FeAl2 1156
aP18 P1 FeAl2
a = 487.8 b = 646.1 c = 880.0 = 91.75° = 73.27° = 96.89°
[2003Pis], at 66.9 at.% Al solid solubility ranges from 65.5 to 67.0 at.% Al
, Fe2Al5 1169
oC24 Cmcm
a = 765.59 b = 641.54 c = 421.84
[2003Pis], at 71.5 at.% Al solid solubility ranges from 71.0 to 72.5 at.% Al. Equilibrium solubility is up to 11 at.% Zn at 450°C [1992Per]. [2001Koe], at Fe4Al10 Zn
a = 764.14 b = 642.76 c = 421.87 a = 762.23 b = 646.25 c = 423.00 Fe4Al13 1160
mC102 C2/m Fe4Al13
Comments/References
[2001Koe], at Fe4Al9Zn2
a = 1552.7 to 1548.7 b = 803.5 to 808.4 c = 1244.9 to 1248.8 = 107.7 to 107.99° a = 1549.2 b = 807.8 c = 1247.1 = 107.69
[2003Pis], 74.16 to 76.7 at.% Al solid solubility ranges from 74.5 to 75.5 at.% Al [2003Pis], at 76.0 at.% Al sometimes called FeAl3 in the literature
, Fe3Zn 10 782
cI52 I43m Fe3Zn10 ? Cu5Zn8
a = 897.41 a = 901.8
[V-C], solid solubility ranges from 68.0 to 82.5 at.% Zn
1, Fe11 Zn39 550
cF408 F43m Fe11 Zn39
a = 1796.3
[V-C2], solid solubility ranges from 75.5 to 81.0 at.% Zn
, FeZn 10 665
hP555 P63mc FeZn10
a = 1283.0 b = 5770.0
[V-C], solid solubility ranges from 86.5 to 92.0 at.% Zn. Equilibrium solubility is up to 4.3 at.% Al at 450°C [1992Per].
, FeZn13 530
mC28 C2/m CoZn13
a = 1342.4 b = 760.8 c = 506.1 = 127.3°
[V-C], solid solubility ranges from 92.5 to 94.0 at.% Zn. Equilibrium solubility is up to 1.85 at.% Al at 450°C [1992Per].
2, AlFe14Zn 1.5 450 (?)
-
-
[1992Per, 1998Yam]
MSIT®
Landolt-Börnstein New Series IV/11A3
Landolt-Börnstein New Series IV/11A3
379
(αFe)+FeAl2+ε
U4
E3
U6
Fe4Al13+(Al)+(Zn)
L Fe4Al13+(Al)+(Zn)
L+Fe4Al13+(Zn)
L + δ η + (Zn)
ca.425 max L δ + (Zn)
L + (αFe) δ + η
E2
L + η Fe4Al13 + (Zn)
U7
δ+(Zn)+ζ
L δ + (Zn) + ζ
U5
(αFe)+α2+FeAl2
E1
?
274 (Al´)(Al´´)+Fe4Al13+(Zn) E4 (Zn)+Fe4Al13+(Al)
ca.351 Fe4Al13+(Al´)+(Al´´) η+Fe4Al13+(Zn)
409
U2
?
ε (αFe) + α2 + FeAl2
L+η+(Zn)
ca.420
U1 (αFe)+α2+ε
L+δ+η
(αFe)+δ+η
553
A-B-C L+(αFe)+α2
L+(αFe)+η
1038
L + ε (αFe) + η
(αFe)+L+ε
L + α2 (αFe) + ε
Al-Fe-Zn
U3
L+(αFe)+δ
L + Γ (αFe) + δ
η+δ+(Zn)
418
Γ+(αFe)+δ
ca.660
1130 (αFe)+η+ε
ca.1200
ε + η (αFe) + FeAl2
(αFe)+FeAl2+η
1065
Fig. 1: Al-Fe-Zn. Reaction scheme
665 e4 l (Al) + Fe4Al13
1102 e3 ε α2 + FeAl2
1156 p2 ε + η FeAl2
1160 e2 l η + Fe4Al13
1232 p1 l + α2 ε 1165 e1 lε+η
Al-Fe
381 e5 l (Al) + (Zn) 277 e6 (Al´) (Al´´) + (Zn)
Al-Zn
425 p7 l + ζ (Zn)
530 p6 l+δζ
550 p5 Γ + δ Γ1
665 p4 l+Γδ
782 p3 l + (αFe) Γ
Fe-Zn
Al–Fe–Zn 31
MSIT ®
Al–Fe–Zn
32
Fe Zn Al
Fig. 2: Al-Fe-Zn. Partial liquidus surface
20.00 0.00 80.00
Data / Grid: at.% Axes: at.%
e2
Fe4Al13
30
70
e1
η ε p1
40
U2
60
U1
50
α2
50
(α Fe) ? Fe Zn Al
60.00 0.00 40.00
Fig. 3: Al-Fe-Zn. Liquidus surface of the Zn corner
10
20
Fe Zn Al
30
0.00 90.00 10.00
Fe Zn Al
20.00 40.00 40.00
Data / Grid: at.% Axes: at.%
to E3
from e2 3 Al 1 Fe 4
from U2 U7
η (αFe)
(Zn)
U4
from p3
δ
Γ Fe Zn Al
MSIT®
10.00 90.00 0.00
p4
U5
U6 E2 p6 ζ p7
Zn
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn
33
Al
Data / Grid: at.% Axes: at.%
Fig. 4: Al-Fe-Zn. Isothermal section at 700°C 20
Fe4Al13
80
η
FeAl2 (αFe)+FeAl2+η
40
60
60
40
L+(α Fe)+η
α2
L+Fe4Al13+η 80
20
(α Fe) L+Γ+(αFe) 20
Fe
40
L 60
Γ
Al Fig. 5: Al-Fe-Zn. Isothermal section at 575°C
Zn
Data / Grid: at.% Axes: at.%
(Al) L+(Al)+Fe4Al13
20
80
Fe4Al13 FeAl2
80
η
40
60
(αFe)+FeAl2+η
L 60
40
α2
L+(α Fe)+η
(αFe)+α 2
L+Fe4Al13+η
80
20
(α Fe) (αFe)+Γ+δ
Fe
Landolt-Börnstein New Series IV/11A3
20
40
L+δ+(α Fe) 60
Γ
80
δ
Zn
MSIT ®
Al–Fe–Zn
34
Al Fig. 6: Al-Fe-Zn. Isothermal section at 500°C
Data / Grid: at.% Axes: at.%
(Al) 20
Fe4Al13
80
η FeAl2
L+(Al)+Fe4Al13
40
60
(α Fe)+FeAl2+η
L+Fe4Al13+η
60
40
(α Fe)+δ+η
α2
α1
L
(α Fe)+α1
L+δ+Fe+η
80
20
(αFe)
Γ+δ+(αFe) 20
Fe
40
Fe Zn Al
Fig. 7: Al-Fe-Zn. Partial isothermal section at 500°C
δ 60
80
Γ
0.00 70.00 30.00
Γ1
Zn
ζ
Data / Grid: at.% Axes: at.%
L+Fe4Al13+η
10
20
L
20
δ+η+(αFe) 10
L+δ+η (αFe)+Γ+δ
Γ+Γ1+δ Fe Zn Al
MSIT®
30.00 70.00 0.00
80
Γ1
δ 90
ζ
L+ζ+δ
Zn
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn Fe Zn Al
Fig. 8: Al-Fe-Zn. Partial isothermal section at 470°C
35 0.00 99.45 0.55
Data / Grid: at.% Axes: at.%
L+ δ
L+η
+η
L+δ L+ζ+δ
L
L+ζ
Fe Zn Al
0.55 99.45 0.00
Fig. 9: Al-Fe-Zn. Partial isothermal section at 460°C
Zn Fe Zn Al
0.00 99.45 0.55
Data / Grid: at.% Axes: at.%
L+η
L+ η
+δ
L+δ L+δ+ζ L
L+ζ
Fe Zn Al Landolt-Börnstein New Series IV/11A3
0.55 99.45 0.00
Zn
MSIT ®
Al–Fe–Zn
36
Al
Data / Grid: at.%
Fig. 10: Al-Fe-Zn. Isothermal section at 450°C
Axes: at.%
20
Fe4Al13
(Al)+Fe4Al13
80
(Al)
FeAl2
η
40
L+ Fe
60
4A
α2
l1
3 +(
Al )
60
40
(αFe)+δ+η
α1
L+ Fe
4
(α Fe)+δ
80
Al 1
3
20
(αFe)
L
Γ2
(αFe)+Γ
Fe
20
δ 40
60
Fe Zn Al
Fig. 11: Al-Fe-Zn. Partial isothermal section at 450°C
Γ1
Γ
80
0.00 80.00 20.00
ζ
Zn
Data / Grid: at.% Axes: at.%
L+Fe4Al13
η+L
L
η+Γ2+L
10
10
η+Γ2
η+δ
Γ2 Γ2+L
(αFe)+δ
δ+L Γ1 Fe Zn Al
MSIT®
20.00 80.00 0.00
Γ1+δ
δ
ζ 90
L+ζ
Zn
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn
37
Fe Zn Al
Fig. 12: Al-Fe-Zn. Partial isothermal section at 450°C
0.00 99.45 0.55
Data / Grid: at.% Axes: at.%
L+η L+Γ2+η
L+δ
L+Γ2
L+δ+Γ2
L+ζ+δ
L
L+ζ
Fe Zn Al
0.55 99.45 0.00
Zn Fe Zn Al
Fig. 13: Al-Fe-Zn. Partial isothermal section at 400°C
0.00 60.00 40.00
Data / Grid: at.% Axes: at.%
(Al)
10
30
L+Fe4Al13+(Al)
(Z n) +η
20
L+(Al)
+F e
4
Al 1
20 3
L
(αFe)+Γ+δ
Γ +Γ 1+δ Fe Zn Al Landolt-Börnstein New Series IV/11A3
40.00 60.00 0.00
L+Fe4Al13+(Zn)
(α Fe)+η+δ
30
70
Γ1
80
10
(Zn)+η+δ
δ
(Zn) 90
ζ
(Zn)+δ+ζ
Zn
MSIT ®
Al–Fe–Zn
38
Al
Data / Grid: at.%
Fig. 14: Al-Fe-Zn. Partial isothermal section at 350°C
Axes: at.%
Fe4Al1320
80
(Al")
TK
(Al')+(Al")+Fe4Al13 (Al')+(Al")
40
60
(Al') 60
40
(Al')+(Zn)+Fe4Al13 80
20
(Zn) 20
Fe
40
60
80
Al
Zn
Data / Grid: at.%
Fig. 15: Al-Fe-Zn. Partial isothermal section at 330°C
Axes: at.%
(Al")
20
Fe4Al13
80
(Al')+(Al")+Fe4Al13 40
60
(Al') 60
40
(Al')+(Zn)+Fe4Al13 80
20
(Zn)
Fe
MSIT®
20
40
60
80
Zn
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn
39
Al
Data / Grid: at.%
Fig. 16: Al-Fe-Zn. Partial isothermal section at 300°C
Axes: at.%
(Al")
20
Fe4Al13
80
(Al')+(Al")+Fe4Al13 40
60
(Al')
60
40
(Al')+(Zn)+Fe4Al13 80
20
(Zn) 20
Fe
40
60
80
Al Fig. 17: Al-Fe-Zn. Partial isothermal section at 250°C
Zn
Data / Grid: at.% Axes: at.%
(Al)
20
Fe4Al13
80
(Al)+(Zn)+Fe4Al13
40
60
60
40
80
20
(Zn)
Fe
Landolt-Börnstein New Series IV/11A3
20
40
60
80
Zn
MSIT ®
Al–Fe–Zn
40
1000
Fig. 18: Al-Fe-Zn. Section at a constant Zn-content of 30 mass%
900
L
700
600
η+Fe4 Al13+L
Temperature, °C
800
Fe4Al13+L L+(Al) Fe4Al13+(Al)+L
500
Fe4Al13+(Al)
379°C
Fe 4 Al13+(Zn)
300
409°C
η+Fe 4 Al13+(Zn)
400
Fe4Al13+(Al) +(Zn)
Fe4Al13+(Al')+(Al'')
(Al)
274°C
Fe 4Al13+(Al')+(Zn)
(Al')+(Zn)
200
Fe 21.67 60 Zn 18.51 Al 59.82
Fig. 19: Al-Fe-Zn. Vertical section at a constant Zn-content of 90 mass%
70
0.00 Fe Zn 15.03 Al 84.97
80
Al, at.%
1000
L 900
L+(α Fe) 800
700
L+η
660°C L+(α Fe)+δ
600
500
L+δ
L+η +(α Fe)
L+Fe4Al13
553°C
L+η +Fe4Al13
L+η +δ
L+δ +ζ 420
409° L+Fe4Al13+(Zn)
418°C
400
δ +ζ
η+δ +(Zn)
η+(Zn)
δ 300
δ +(Zn) δ +ζ +(Zn)
η+Fe4Al13+(Zn)
Temperature, °C
L+Γ +(α Fe) L+Γ
379°C Fe4Al13+(Al)+(Al'')
(Al)+(Zn)
274°C
(Al)+(Al'')+(Zn) (Al")+(Zn)
Fe4Al13+(Al'')+(Zn)
200
Zn 88.49 Fe 11.51 0.00 Al
MSIT®
10
Al, at.%
Fe4Al3+(Zn)
L+(Al) L+(Al)+(Zn)
20
Zn 78.78 0.00 Fe Al 21.22
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zn
Fig. 20: Al-Fe-Zn. Vertical section at a constant Zn-content of 95 mass%
1000
L 900
L+(α Fe)
800
L+(α Fe)+Γ
L+Γ
Temperature, °C
41
700
L+Γ+δ
660°C
L+η
L+(α Fe)+δ L+(α Fe)+η
600
500
L+ζ L+ζ +δ
L+η +Fe4Al13
553°C
L+δ L+δ +η L+δ +ζ
L+Fe4Al13
L+Fe4Al13+(Zn)
L+(Zn)+δ
420°C
L+η+(Zn)
418°C
409°
400
ζ +(Zn)
δ +ζ +(Zn)
δ +η+(Zn)
300
η+(Zn) η+Fe4Al13+(Zn) Fe4Al13+(Zn)
δ +(Zn) 200
Zn 94.20 5.80 Fe 0.00 Al
Fig. 21: Al-Fe-Zn. Vertical section at a constant Zn-content of 98 mass%
379°C Fe4Al13+(Al)+(Zn) 274°C
L+(Zn) L+(Zn)+(Al) (Al)+(Zn) (Al)+(Al'')+(Zn)
Fe4Al13+(Al'')+(Zn) 10
Al, at.%
(Al")+(Zn)
Zn 88.69 0.00 Fe Al 11.31
800
L
700
600
L+δ
L+η +(α Fe)
553°C
L+η 500
L+ζ L+ζ +(Zn)
L+ζ +δ
L+η +δ
420°C
L+η +(Zn)
418°C
ζ +(Zn)
ζ +δ +(Zn)
η+(Zn)+δ
300
Fe4 Al13+(Zn)
400
409°C
379°C
η+Fe4 Al13+(Zn)
Temperature, °C
L+(α Fe) L+(α Fe)+δ
Fe4 Al13 +(Al) +(Zn) 274°C Fe4 Al13 +(Al'') +(Zn)
200
Zn 97.67 2.33 Fe 0.00 Al
Landolt-Börnstein New Series IV/11A3
δ +(Zn)
η+(Zn)
Al, at.%
L+(Zn) L+(Al)+(Zn) (Al)+(Zn) (Al)+(Al'')+(Zn) (Al")+(Zn)
Zn 95.29 0.00 Fe 4.71 Al
MSIT ®
42
Al–Fe–Zr
Aluminium – Iron – Zirconium Zoya M. Alekseeva, updated by Viktor Kuznetsov Literature Data [1966Mar] investigated alloys along the section ZrAl2-ZrFe2 by X-ray diffraction; the alloys studied were prepared by arc-melting and annealed at 900°C for 20 d. Two ternary Laves type compounds 1 and 2 were found with extended homogeneity regions along the section studied. [1968Gru] investigated, essentially by metallography, alloys in the Zr corner along the sections with Al to Fe ratios 2:1 and 1:2 up to 14 mass% Al and 14 mass% Fe, using alloys that were quenched from 1350, 1200, 1100, 900, 800 and 700°C. Partial isothermal sections at 1200, 900, 800 and 700°C were constructed. However, in the isothermal sections below 1200°C the existence of the ternary compound Zr6FeAl2 reported by [1969Bur] has not been taken into account. Isothermal sections at 900 and 800°C do not contain the binary compound Zr2Fe and the binary compound Zr3Fe is missing in the isothermal sections at 800 and 700°C. [1969Bur] investigated, mainly by X-ray diffraction, 116 alloys which were prepared by arc melting and annealed at 900°C for 2100 h. Two more ternary compounds have been found in addition to 1 and 2 reported earlier: (i) a "line" compound ZrFe7-4 Al5-8 with Al content varying from 37 to 61 at.% and (ii) a stoichiometric compound Zr6FeAl2. An isothermal section at 900°C has been constructed. [1970Kri] established the crystal structure of Zr6FeAl2 compound; the structure was later that refined by [1997Yan]. [1973Ath] investigated (by EMPA, X-ray and electron diffraction) the ternary compound ZrFe3.3Al1.3 occurring in a two-phase alloy (the other phase was Fe3Al) which was prepared by substituting 5 at.% Zr for Fe in the alloy Fe76Al24. The alloy studied was annealed at 950°C for 24 h. [1974Dwi] investigated the ternary equiatomic compound ZrFeAl which was prepared by arc melting and annealed in a Vycor capsule. [1974Kuz] prepared alloys from the elements with a purity of 99.99% and annealed them at 500°C for 50 days. By the measurement of the lattice parameters they determined the existing phases and their solubility ranges on the section ZrAl2-ZrFe2. These results are in agreement with [1966Mar]. [1977Mur] studied the crystal structure, magnetic properties and Fe Moessbauer effect on the Laves phase Zr(Fe1-xAlx)2 in the stoichiometric range x = 0 to 0.4. [1987Bla] investigated the solubility of Al in ZrFe2 by means of X-ray powder diffraction and measurements of microhardness on alloys melted and heat treated for at least 24 h in the range 800 to 1500°C. The samples were either quenched or cooled at 1.7 K·min-1. The substitution of Al for Zr changed the unit cell parameter from 706.8 pm for ZrFe2 to 702.3 pm for (Zr0.87Al0.13 )Fe2. A brief review of the system mainly concerning intermetallics formation may be found in [1990Kum]. [1991Des] found no evidence for the presence of L12 phase in mechanically alloyed sample with composition of Al-12.5Fe-25Zr (at.%). [1991Sok] studied partial section from Al corner with Zr to Fe ratio being 1:3, isopleth of 25 at.% Zr and partial isothermal section at 500°C for Al < 25 at.%. They used Al 99.9% purity, iodide-purified Zr 99.9% and Fe 99.9%. Alloys were prepared in arc furnace with water-cooled Cu bottom in Ti gettered Ar atmosphere with subsequent annealing in evacuated silica tubes at 500°C for 1000 h and water quenching. Samples were studied by DTA, metallography and X-ray analysis. No ternary phases were found in the region studied. [1992Sle] investigated temperature dependence of lattice spacing at 0 to 300 K and magnetic susceptibility at 80 to 600 K of the Laves phase with composition of ZrFe1.2Al0.8.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zr
43
[1993Nov], [1994Isr] and [1997Isr] studied bonding characteristics in Laves phases Zr(AlxFe1–x )2 with various x values. Experimental techniques included Moessbauer spectroscopy, nuclear-resonant-photon-scattering and neutron diffraction; all were used to determine effective Debye temperature which measured bonding strength. Minimal value of that at x = 0.2 was found to coincide with maximal hydrogen absorption power. [1994Kle] measured standard enthalpies of formation calorimetrically for Zr(Fe(1-x)Alx phase at x = 0, 0.0833, 0.2, 0.5, 0.7 and 1 by measuring heat of dissolution in acid mixture (HF+HNO3). [1996Gon] used these data (among much others) to test their generalization of well-known Miedema model to ternary intermetallics with moderate success. [1999Zav] investigated structural changes of Zr6FeAl2 under hydrogen treatment. [1999Mek] performed ab initio calculation of interatomic potentials and influence of Zr additions on the ordering in intermetallics of the Al-Fe system. [2000Biz] studied in great detail kinetics of crystallization of Al-Fe and Al-Fe-Zr rapidly solidified alloys. In particular, a number of kinetic models were tried. Mechanical alloying of sample with Zr3Fe7Al90 composition was studied by [2001Rod] who found a mixture of amorphous and unspecified nanocrystalline phases and studied their crystallization behavior using DSC and X-ray techniques. This evaluation incorporates and continues the critical evaluation made by [1992Ale] considering new published data. Binary Systems For the Al-Fe and Al-Zr binary systems recently updated versions of [2003Pis] and [2003Sch] were accepted, respectively. Fe-Zr system is from [Mas2]. Solid Phases Five ternary compounds have been found in the system. ZrFe2 extends into the ternary to about 10 at.% Al. The existence of an additional ternary phase - with AuCu3 structure was claimed by [1989Sch] at the composition Zr25Fe5.5Al and 1100°C; the temperature and composition range of existence is still unknown, so it could not be included in the phase diagram. Crystallographic data of all the phases are listed in Table 1. Invariant Equilibria The partial vertical section from Al corner with Zr to Fe ratio of 1:3 [1991Sok] crosses a plane of invariant reaction between L, (Al), Fe4Al13 and Al3Zr phases at about 650°C (see below Fig. 4), but neither its nature, nor phase compositions are provided (the temperature value was taken by present author from small-scale figure). Isothermal Sections The partial isothermal section at 1200°C, presented in Fig. 1, is based on the results of [1968Gru]. To bring that into agreement with accepted version of Fe-Zr binary, the boundary (Zr)+L/L was shifted; also some modification of position of L corner of (Zr)+L+Zr5Al3 tie-triangle was necessary. These changes necessitate certain boundaries given in the original work as uncertain. Figure 2 displays the isothermal section at 900°C based on the results of [1969Bur]. In both isothermal sections the - phase of [1989Sch] is not included since [1968Gru, 1969Bur] did not detect this phase. To adapt to the accepted binary systems, changes were made as following: the three-phase field Zr+Zr+Zr3Al was inserted; a liquid single-phase field in the Al corner and the corresponding two- and three-phase fields were added. The ternary compound Zr18Fe59Al23 [1973Ath] was also included with the corresponding three-phase fields. Extension of the 2 phase field is shown according to the stoichiometry reported in [1966Mar, 1969Bur, 1974Kuz]. It should be noted that in the isothermal section reported by [1969Bur] it has been shown up to 60 at.% Al, which however, contradicts the tabulated results of Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Fe–Zr
44
[1966Mar], where a sample with 60 at.% Al contained a second phase, ZrAl2. Some minor shifts of position of tie-lines of equilibria with that phase, which do not contradict to real phase compositions of the alloys studied, had to be done. Figure 3 presents the partial isothermal section at 500°C in the Al rich corner [1991Sok]. Temperature – Composition Sections Figure 4 displays partial vertical section from Al corner with Zr to Fe ratio of 1:3 [1991Sok]. Figure 5 presents the isopleth at 75 at.% Al, taken from the same source. According to the accepted Al-Fe binary, the L/L+Fe4Al13 boundary line (given in [1991Sok] as dashed line) must approach the temperature axis a bit higher than point c, but this may hardly be seen in the scale of original figure. Thermodynamics [1994Kle] measured standard enthalpies of reactions: 2x Al + 2(1-x)Fe + Zr = Zr(AlxFe(1-x))2 using acid -solution calorimetry at 25°C. The results are: for x = 0 H = –718 kJ, for x = 0.0833 H = –749 kJ, for x = 0.2 H = –8310 kJ, for x = 0.5 H = –12513 kJ, for x = 0.7 H = –23221 kJ, and for x = 1 H = –15413 kJ. Theoretical results of [1996Gon] are not in very good agreement with these. Miscellaneous [1988Vig] compared the microstructural stability of Al-8Fe and Al-8Fe-1.5Zr (mass%) alloys. The ribbons used were produced by melt spinning and were about 40 to 60 m thick and 4 to 5 mm wide. Fine ZrAl3 precipitates appear in the Al matrix during ageing at 200 to 400°C along with FeAl6. The substitution of Al for Fe rapidly reduces the Fe magnetic moment of the compound ZrFe2 [1977Mur] and the substitution of Al for Zr reduces microhardness values of the compound from 8329 to 6818 N·mm-2 [1987Bla]. [1991Sik] studied possible techniques of industrial treatment of Fe3Al intermetallic, including that with Zr additions. [1999Mek] performed ab initio calculation of influence of a number of elements (including Zr) on ordering in FeAl compound. It has been shown that Zr atoms substitute preferentially for Fe sublattice sites in FeAl compound. References [1961Now] [1966Mar]
[1968Gru]
[1969Bur] [1970Kri]
[1973Ath]
MSIT®
Nowotny, H., Schob, O., Benesovsky, F., “The Crystal Structure of Zr2Al and Hf2Al” (in German), Monatsh. Chem., 92, 1300-1304 (1961) (Crys. Structure, Experimental, 10) Markiv, V.Ya., Kripyakevich, P.I., “Compounds of the Type R(X', X'') 2 in Systems with R = Ti, Zr, Hf; X' = Fe, Co, Ni, Cu and X'' = Al, Ga and Their Crystal Structures”, Sov. Phys.-Crystallogr. (Engl. Transl.), 11, 733-738 (1967), translated from Kristallografiya, 11, 859-864 (1966) (Crys. Structure, Experimental, 25) Gruzdeva, N.M., Zagorskaya, T.N., Raevskii, I.I., “Structure and Properties of Alloys in the Zirconium Corner of Al-Fe-Zr System” (in Russian), in: Fiziko-Khimiya Splavov Tsirkoniya (Physical Chemistry of Zirconium Alloys), Moscow: Nauka, 5-9 (1968) (Equi. Diagram, Experimental, #, 3) Burnashova, V.V., Markiv, V.Ya., “Study of Al-Fe-Zr System”, Dopov. Akad. Nauk Ukr. RSR, A, (4), 351-353 (1969) (Crys. Structure, Equi. Diagram, Experimental, *, 16) Kripyakevich, P.I., Burnashova, V.V., Markiv, V.Ya., “Crystal Structure of the Compounds Zr 6FeAl2, Zr6CoAl2, and Zr6NiAl2”, Dopov. Akad. Nauk Ukr. RSR A, (9), 828-831 (1970) (Crys. Structure, Experimental) Athanassiadis, G., Dirand, M., Rimlinger, L., “X-Ray Diffraction and Electron Diffraction Study of the Compound of Al1.3Fe3.5Zr” (in French), C. R. Seances Acad. Sci. (Paris), 277, C915-C917 (1973) (Crys. Structure, Experimental, 3) Landolt-Börnstein New Series IV/11A3
Al–Fe–Zr [1974Dwi]
[1974Kuz]
[1977Mur]
[1987Bla]
[1988Vig]
[1989Ale]
[1989Sch] [1990Kum]
[1991Des]
[1991Sik] [1991Sok]
[1992Ale]
[1992Sle]
[1993Nov]
[1994Isr]
[1994Kle]
[1996Gon]
Landolt-Börnstein New Series IV/11A3
45
Dwight, A.E., “Alloying Behavior of Zr, Hf and the Actinides in Several Series of Isostructural Compounds”, J. Less-Common Met., 34, 279-284 (1974) (Crys. Structure, Experimental, 6) Kuz’menko, P.P., Suprunenko, P.A., Markiv, V.Ya., Butsik, T.M., “Magnetic Properties of Laves Phases in the Zr-Fe-Al and Zr-Co-Al Systems” (in Russian), Akad. Nauk Ukr. SSR, Metallofizika, 52, 58-61 (1974) (Crys. Structure, Equi. Diagram, Experimental, 10) Muraoka, Y., Shigas, M., Nakamura, Y., “Magnetic Properties and Mössbauer Effect of A(Fe1-xBx)2 (A =Y or Zr, B = Al or Ni) Laves Phase Intermetallic Compounds”, Phys. Status Solidi, 42A, 369-374 (1977) (Crys. Structure, Experimental, 15) Blarzina, Z., Trojko, R., “On Friauf-Laves Phases in the Zr1-xAlxT2, Zr 1-xSixT2 and Zr 1-xTixT2 (T = Mn, Fe, Co) Systems”, J. Less Common Met., 133, 277-286 (1987) (Crys. Structure, Experimental, 10) Vigier, E., Ortez-Mendez, U., Merles, P., Thaller, G., Fouguet, F., “Microstructural Stability of Rapidly Quenched Al, Fe Alloys: Influence of Zirconium”, Mater. Sci. Eng., 98, 191-195 (1988) (Experimental, 11) Alekseeva, Z.M., Korotkova, N.V., “Phase Diagram of the Fe-Zr System” (in Russian), Izv. Akad. Nauk SSSR, Met., (4), 202-208 (1989) (Crys. Structure, Equi. Diagram, Experimental, #, 21) Schneibel, J.H., Porter, W.D., “High Temperature Order Intermetallic Alloys III”, Mater. Res. Soc. Symp. Proc., Stoloff, N.S. (Ed.), 335-340 (1989) (Crys. Structure) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mat. Rev., 35, 293-327 (1990) (Crys. Structure, Equi. Diagram, Review, 158) Desch, P.B., Schwarz, R.B., Nash, P., “Formation of Metastable L12 Phases in Al3Zr and Al-12.5% X-25% Zr (X = Li, Cr, Fe, Ni, Cu)”, J. Less-Common Met., 168, 69-80 (1991) (Crys. Structure, Experimental, 25) 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) Sokolovskaya, E.M., Kazakova E.F., Grigorovitch E.V., Matveyev I.N., “Phase Equilibria in Alloys of the Al-Fe-Zr System”(in Russian), Vestn. Mosk. Univ., Ser. 2: Khim., 32, 478-481 (1991) (Equi. Diagram, Experimental, *, #, 7) Alekseeva, Z.M., “Aluminium - Iron - Zirconium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16088.1.20, (1992) (Crys. Structure, Equi. Diagram, Assessment, 15) Slebarski, A., Hafez, M., Zarek, W., “Spin Fluctuations in ZrM1.2Al0.8 with Transition Metal M of the 3d Type”, Solid State Commun., 82(1), 59-61 (1992) (Crys. Structure, Experimental, 12) Novik I., Yacob B., March R., “Moessbauer Study of Crystallographic and Magnetic Phase Transitions, Phonon Softening, and Hyperfine Interactions in Zr(AlxFe1–x)2”, Phys. Rev. B, 47, 723-726 (1993) (Phys. Prop., Experimental) Israel A., Yacob I., March R., Shanal O., Wolf A., Fogel M., “Correlation Between Anomalous Hydrogen Absorption and 56Fe-Bonding Strength in the Zr(AlxFe1-x)2 System”, Phys. Rev. B, 50, 3564-3569 (1994) (Phys. Prop., Experimental, 29) Klein, R., Jacob, I., O'Hare, P.A.G., Goldberg, R.N., “Solution-Calorymetric Determination of the Standard Molar Enthalpies of Formation of the Pseudobinary Compounds Zr(AlxFe(1-x))2 at the Temperature 298.15 K”, J. Chem. Thermodyn., 26, 599-608 (1994) (Thermodyn., Experimental, 22) Goncalves, A.P., Almeida, M, “Extended Miedema Model: Predicting the Formation Enthalpies of Intermetallic Phases with More than Two Elements”, Physica B (Amsterdam), 228, 289-294 (1996) (Thermodyn., Theory, 19)
MSIT ®
Al–Fe–Zr
46 [1997Yan]
[1997Isr]
[1999Mek]
[1999Zav]
[2000Biz] [2001Rod] [2003Pis]
[2003Sch]
Yanson, T.I., Manyako, M.B., Bodak, O.I., Cerny, R., Pacheko, J.V., Yvon, K., “Crystal Structure of Zirconium Iron Aluminide, Zr6FeAl2”, Z. Kristallogr. NCS, 212, 504 (1997) (Crys. Structure, Experimental, 5) Israel, A., Jacob, I., Soubeyroux, J.L., Fruchart, D., Pinto, H., Melamud, M., “Neutron Diffraction Study of Atomic Bonding Properties in the Hydrogen-Absorbing Zr(AlxFe1-x)2 System”, J. Alloys Compd., 253-254, 265-267 (1997) (Phys. Prop., Experimental, 12) Mekhrabov, A.O., Akdeniz, M.V., “Effect of Ternary Alloying Elements Addition on Atomic Ordering Characteristics of Fe-Al Intermetallics”, Acta Mater., 47, 2067-2075 (1999) (Thermodyn., Theory, 63) Zavaliy, I.Yu., Pecharsky, V.K., Miller, G.J., Akselrud, L.G., “Hydrogenation of Zr6MeX 2 Intermetallic Compounds (Me=Fe, Co, Ni, X=Al, Ga, Sn): Crystallographic and Theoretical Analysis”, J. Alloys Compd., 283, 106-116 (1999) (Crys. Structure, Experimental, 31) Bizjak, M., Kosec, L., “Phase Transformations of Al-Fe and Al-Fe-Zr Rapidly Solidified Alloys”, Z. Metallkd., 91, 160-164 (2000) (Kinetics, Electr. Prop., Experimental, 12) Rodriguez, C.A.D., Botta F., W.J., “High-Energy Ball Milling of Al-Based Alloys”, Key Eng. Mater., 189-191, 573-578 (2001) (Crys. Structure, Experimental, 10) Pisch, A., “Al-Fe (Aluminum-Iron)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, to be published, (2003) (Equi. Diagram, Assessment, Crys. Structure, 58) Schuster, J.C., “Al-Zr (Aluminium-Zirconium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram, Assessment, 151)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 ( Fe) 1538 - 1394 ( Fe) 1394 - 912 (Fe) < 912 (Zr)(h) 1855 - 863 (Zr)(r) < 863
MSIT®
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cI2 Im3m W cF4 Fm3m Cu cI2 Im3m W cI2 Im3m W hP2 P63/mmc Mg
Lattice Parameters Comments/References [pm] a = 404.88
pure Al [V-C]
a = 293.15
[Mas2]
a = 364.67
at 915°C [V-C2, Mas2]
a = 286.65
pure Fe at 20°C [V-C]
a = 362
[P]
a = 323.2 c = 514.7
[V-C]
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zr Phase/ Temperature Range [°C] Fe4Al13 (FeAl3.2, FeAl3) 1160
Pearson Symbol/ Space Group/ Prototype mC102 C2/m Fe4Al13
Fe2Al5 1169
oC24 Cmcm
FeAl2 < 1156
aP18 P1 FeAl2
2 Fe100-xAlx < 1310
cP2 Pm3m CsCl tI16 I4/mmm ZrAl3 hP12 P63/mmc MgZn2 oF40 Fdd2 Zr2Al3 oC8 Cmcm CrB hP7 P6/mmm Zr4Al3 tP20 P42/mnm Zr4Al3 tI32 I4/mcm W5Si3 hP6 P63/mmc Ni2In
ZrAl3 < 1580 ZrAl2 < 1660 Zr2Al3 < 1590 ZrAl < 1275 25 Zr4Al3 1030 Zr3Al2 < 1480 Zr5Al3(h) < 1400 Zr2Al < 1350
Landolt-Börnstein New Series IV/11A3
47
Lattice Parameters Comments/References [pm] a = 1552.7 - 1548.7 b = 803.5 - 808.4 c = 1244.9 - 1248.8 = 107.7 - 107.99° a = 1549.2 b = 807.8 c = 1247.1 = 107.69 a = 765.59 b = 641.54 c = 421.84 a = 487.8 b = 646.1 c = 880.0 = 91.75° = 73.27° = 96.89° a = 290.9
[2003Pis], 74.16 to 76.7 at.% Al solid solubility ranges from 74.5 to 75.5 at.% Al [2003Pis], at 76.0 at.% Al
[2003Pis], at 71.5 at.% Al solid solubility ranges from 71.0 to 72.5 at.% Al [V-C] 65.5 to 67 at.% Al [Mas]
28.0 x 52.5 at 900°C at x = 50 [V-C]
a = 399.93 0.05 c = 1728.3 0.02
[2003Sch]
a = 528.24 c = 874.82
[2003Sch]
a = 960.1 0.2 b = 1390.6 0.2 c = 557.4 0.2 a = 335.9 0.1 b = 1088.7 0.3 c = 427.4 0.1 a = 543.3 0.2 c = 539.0 0.2
[2003Sch]
a = 763.0 0.1 c = 699.8 0.1
[2003Sch]
a = 1104.4 c = 539.1
[2003Sch]
a = 489.39 0.05 c = 592.83 0.05
[2003Sch]
[2003Sch]
[2003Sch]
MSIT ®
48
Al–Fe–Zr
Pearson Symbol/ Space Group/ Prototype cP4 Pm3m Cu3Au cF24
, (Fe1-xAlx)2-z(Zr1-yAly)1+z Fd3m Cu2Mg
Lattice Parameters Comments/References [pm]
Phase/ Temperature Range [°C] Zr3Al < 1019
a = 437.2 0.3
a = 706.8 a = 707.4 a = 709.4 a = 713.5 a = 712.4 a = 702.3 a = 701.0 0,3 a = 704.0 0.3
* 1, Zr(Fe1-xAlx)2
* 2, Zr(Fe1-xAlx)2
* Zr6FeAl2
* Zr18Fe59Al23 * , Zr(Fe1-xAlx)12
* -, Zr25Fe5.5 Al
MSIT®
hP12 P63/mmc MgZn2
cF24 Fd3m Cu2Mg hP9 P62m K2UF6 tI52 I4/mcm tI26 I4/mmm ThMn12
cP4 Pm3m AuCu3
a = 508.7 c = 827.7 a = 524.3 c = 852.5 a = 743.0 a = 746.1 a = 792.1 0.2 c = 336.03 0.09 a = 837 c = 998 a = 859.5 c = 496.7 a = 849.3 c = 488.9
[2003Sch]
0 x 0.20, 0 y 0.133, -0.17 z 0.03 at x = 0, y = 0, z = 0 [1987Bla] at x = 0, y = 0, z = 0 [1977Mur] at x = 0.1, y = 0, z = 0 [1977Mur] at x = 0.15, y = 0, z = 0 [1966Mar] at x = 0.2, y = 0, z = 0 [1977Mur] at x = 0, y = 0.133, z = 0 [1987Bla] at x = 0, y = 0, z = -0.17 [1989Ale] at x = 0, y = 0, z = 0.03 [1989Ale] 0.375 x 0.75 [1966Mar] at x = 0.375 [1974Kuz] at x = 0.75 [1974Kuz] 0.15 x 0.175 x = 0.175 [1966Mar,1974Kuz] x = 0.15 [1970Kri], [1997Yan]
[1973Ath] 0.416 x 0.667 [1969Bur] at x = 0.416 at x = 0.667 claimed by [1989Sch]
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zr Zr Fe Al
Fig. 1: Al-Fe-Zr. Partial isothermal section at 1200°C
49 60.00 0.00 40.00
Data / Grid: at.% Axes: at.%
Zr5Al3 (β Zr)+Zr2Al+Zr5Al3 Zr2Al 70
30
L+Zr5Al3 (β Zr)+Zr2Al 80
20
(β Zr)+L+Zr5Al3 (β Zr)+Zr5Al3 90
10
(β Zr)+L
L
(β Zr)
10
Zr
20
30
Al
Fig. 2: Al-Fe-Zr. Isothermal section at 900°C
Zr Fe Al
60.00 40.00 0.00
Data / Grid: at.% Axes: at.%
L
L+ZrAl3
20
80
ZrAl3
Fe4Al13 Fe2Al5
ZrAl2
FeAl2 Zr2Al340
60
λ2
ZrAl Zr4Al3
γ
λ1
Zr3Al260
α2 40
Zr2Al Zr3Al
Zr6FeAl2
Zr18Fe59Al23
80
20
(αFe)
(αZr)
Zr Landolt-Börnstein New Series IV/11A3
(β Zr)
20
Zr2Fe 40
60
ZrFe2 ZrFe3 80
Fe MSIT ®
Al–Fe–Zr
50
Al
Data / Grid: at.%
(Al)
Fig. 3: Al-Fe-Zr. Partial isothermal section at 500°C
Axes: at.%
10
90
Al 13 Fe 4 l)+ (A
ZrA l3 +( Al)
(Al)+Fe4Al13+ZrAl3
20
80
Fe4Al13
ZrAl 3 +Fe4Al13 Zr Fe Al
25.00 0.00 75.00
Fig. 4: Al-Fe-Zr. Vertical section from Al corner with Zr/Fe=1:3 (in at.%)
10
ZrAl3
20
Zr Fe Al
0.00 25.00 75.00
1200
1100
L
Temperature, °C
1000
900
L+Fe4Al13
800
L+Fe4Al13+ZrAl3
(Al)+L 700
600
(Al)+L+Fe4Al13 (Al)+Fe4Al13
(Al)+Fe4Al13+ZrAl3
(Al) 500
Fe4Al13+ZrAl3 400
Al Zr, at.%
MSIT®
6.25 Zr Fe 18.75 Al 75.00
Landolt-Börnstein New Series IV/11A3
Al–Fe–Zr
Fig. 5: Al-Fe-Zr. Vertical section at 75 at.% Al
51
1580°C 1500
L L+ZrAl3
Temperature, °C
1250
L+Fe4Al13
c ZrAl3
L+Fe4Al13+ZrAl3 1000
Fe4Al13 750
Fe4Al13+ZrAl3 500
Zr Fe Al
Landolt-Börnstein New Series IV/11A3
0.00 25.00 75.00
10
20
Zr, at.%
Zr Fe Al
25.00 0.00 75.00
MSIT ®
52
Al–Ge–Li
Aluminium – Germanium – Lithium Oksana Bodak Literature Data Studies on the Al-Ge-Li ternary system are confined to the identification and characterization of a few ternary compounds. Literature data up to 1986 was reported by [1989Goe] and discussed in [1995Pav]. The first report of a ternary phase emanated from [1952Boo] who added Li to a hypereutectic Al-Ge alloy giving a ternary alloy with 38.30 mass% Ge, 6.22 mass% Li. Metallographic analysis clearly indicated the presence of an unidentified phase, probably the ternary -5 compound (LiAlGe). This compound was synthesized by [1960Now] who heated stoichiometric mixtures of the elements in an Fe crucible at temperatures between 800 and 950°C, and found that at 800°C the reaction was incomplete. At higher temperatures the compound LiAlGe was identified, Table 1, together with a very small amount of an unidentified phase of lower crystallographic symmetry. By the same way [1976Sch] prepared the compound LiAlGe, heating stoichiometric amounts of the elements in a tantalum crucible under argon for 15 min at 1000°C. The sample was subsequently annealed for 24 h at 600°C, cooled slowly to room temperature and then the crystal structure was characterized by neutron diffraction analysis, Table 1. The chemical analysis of the compound, 6.6Li-25.2Al-68.3Ge (mass%), agreed well with the calculated values 6.52Li-25.33Al-68.15Ge (mass%) for the composition LiAlGe. A second ternary compound was identified as Li2AlGe by [1974Boc] using the same preparation technique as [1978Ble]. [1978Ble] used 99.98 % Li, 99.999 % Al and Ge, to prepare a third ternary compound whose composition was given as Li5.3Al0.7Ge2 with 1 formula unit in the elementary cell. This compound showed superlattice reflections, which were ascribed to the presence of a phase of the same composition containing 3 formula units in the elementary cell with enlargement of the “a” axis by 3. Due to the reactivity of the alloys high temperature X-ray diffraction analysis could not be employed to determine whether Li16Al2Ge6(-1´) with 3 formula units, is a low temperature polymorph of Li5.3Al0.7Ge2 (-1). [1981Kis] examined three compositions on the section Li(Al1-xGex) with x = 0.02, 0.066 and 0.11. Alloys were prepared by melting 99.999 % Al, 99.9 % Li and an Al-Ge master alloy under argon. The ingot was encapsulated in a Pyrex glass ampoule under 0.5 atm Ar for annealing it 7 days at 500°C and then cooling it slowly down to room temperature. Metallographically the alloys showed a eutectic structure dispersed throughout the sample. X-ray diffraction analysis showed the presence of LiAl in the alloy with x = 0.02 and Li-rich ternary fcc-phase with a = 620 pm. It is the ternary compound -3, Table 1, with a = 616.3 pm according to [1974Boc]. Alloys of nominal weight composition Al-2Li-0.2Ge [1986Cas] were solution heat-treated, quenched and aged for various holding-times at 200°C. The microstructure and deformation behavior were compared for two alloys revealing that the solubility of lithium was increased when germanium was in solid solution, however, lithium decreased the solubility of germanium at 200°C resulting in small germanium precipitates which were homogeneously distributed throughout the matrix. These precipitates had a very positive effect on the deformation behavior and ductility of the alloy. [1992Pav, 1993Pav1, 1993Pav2, 1996Dmy] constructed an isothermal section at 200°C. They prepared their alloys in an electric arc furnace under an argon atmosphere (1.1 # 105 Pa) and determined the crystal structures of the compounds. The purity of lithium was 98 mass%, the purity of silicon and aluminum was better than 99.9 mass%. After melting all alloys were homogenized in evacuated quartz ampoules, at 200°C for 500 h and subsequently quenched into ice water. X-ray powder analysis was used. The authors confirmed the composition and the structure of Li5.3Al0.7Ge2 and LiAlGe compounds, determined the crystal structure of new ternary compounds Li2AlGe and LiAl2Ge and concluded that further new compounds Li9Al2Ge3 and Li6Al3Ge with unknown structures do exist.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ge–Li
53
[1994Hos] studied the effect of some ternary additions (among them Ge) in the L12 type metastable LiAl3 phase calculating the heat of formation by Miedema semi-empirical formula. Binary Systems The description of the Al-Li phase diagram is accepted as given by [2003Gro], that of Al-Ge as given by [Mas2]. For the system Ge-Li it is necessary to note the following. Since long there is a contradiction in number of compounds reported in the phase diagram by [Mas2] and results of X-ray investigations on the crystal structure of compounds. The authors [1997San] made an attempt to resolve this contradiction by compiling the available data and constructing a hypothetic phase diagram, which subsequently was published as a confirmed one by [2000Oka]. [1997San] however, missed the work of [1982Gru], in which the phase diagram has been constructed in detail, using DTA and X-ray investigations. The investigations on the Al-Ge-Li isothermal section by [1993Pav1, 1993Pav2, 1996Dmy] confirm the binary diagram given by [1982Gru], which hence is accepted in the present evaluation and shown in Fig. 1. Remaining discrepancies concern the composition of the Li-richest compound (Table 1) may be due to the difficult Li refinement in the compounds during the X-ray investigation. In the present evaluation the composition Li4Ge is accepted, as given by [1982Gru]. Solid Phases Crystallographic data for the solid phases of this system are presented in Table 1. Isothermal Sections The isothermal section at 200°C shown in Fig. 2 is based on [1993Pav1, 1993Pav2, 1996Dmy]. However the homogeneity regions of the Al-Li binary phases are adjusted to match the accepted binary diagrams. Solubilities of a third component in the binary and unary phases were not determined by [1993Pav1, 1993Pav2, 1996Dym], and hence are not reproduced in Fig. 2 for this evaluation. The same applies for the homogeneity ranges of the ternary phase presented by [1996Dmy]. Thermodynamics Thermodynamic calculations of Li vapor pressures over Al-Li and Al-Li-Me, (Me=Ag, Zn, Cd, Ga, In, etc.) are reported by [1986Lee]. References [1952Boo]
[1960Now] [1974Boc]
[1976Sch]
[1978Ble]
[1981Kis]
Landolt-Börnstein New Series IV/11A3
Boom, E.A., “New in the Systems Aluminium-Germanium-Sodium and Aluminium-Germanium-Lithium” (in Russian), Dokl. Akad. Nauk SSSR, 84(4), 697-699 (1952) (Equi. Diagram, Experimental, 4) Nowotny, H., Holub, F., “Investigation of Metallic System with Fluorspar Phases” (in German), Monatsh. Chem., 91, 877-887 (1960) (Crys. Structure, Experimental, 15) Bockelmann, W., Schuster, H.-U., “Crystallographic Aspects of Ternary Phases of Li with Group III A and IVA Elements in Ionic and Non-Ionic Compositions” (in German), Z. Anorg. Allg. Chem., 410, 241-250 (1974) (Crys. Structure, Experimental, 5) Schuster, H.-U., Hinterhauser, H.-W., Schäfer, W., Will, G., “Neutron Diffraction Investigations of the Phases LiAlSi and LiAlGe” (in German), Z. Naturforsch. B, 31, 1540-1541 (1976) (Crys. Structure, Experimental, 3) Blessing, J., “Synthesis and Studies of Ternary Phases of Li with Elements of the 3 and 4 Sub Groups” (in German), Thesis, Univ. Cologne, 167 pp. (1978) (Crys. Structure, Experimental, 87) Kishio, K., Brittain, J.O., “Phase Stability of Doped -LiAl”, Mater. Sci. Eng., 49, P1-P6 (1981) (Crys. Structure, Experimental, 14)
MSIT ®
54 [1981Gru] [1982Gru]
[1986Lee]
[1986Gas]
[1987Eve]
[1989Goe] [1992Pav]
[1993Pav1]
[1993Pav2]
[1994Hos]
[1995Pav]
[1996Dmy]
[1997San] [2000Oka] [2001Gow]
[2003Gro]
MSIT®
Al–Ge–Li Gruttner, A., Nesper, R., Schnering, H.G., “New Phases in the Li-Ge System: Li7Ge12, Li 12Ge7, Li14Ge6”, Acta Crystallogr., 37A, 161 (1981) (Crys. Structure, Experimental, 5) Gruttner, A., “About the Lithium-Germanium System and Formation of Metastable Germanium-Modifications from Li-Germanides” (in German), Diss. Dokt. Naturwiss., Chem. Fak. Univ. Stuttgart, 1-102 (1982) (Equi. Diagram, Crys. Structure, Experimental) Lee, J.J., Sommer, F., “Thermodynamic Properties of Lithium in Liquid Aluminium Alloys” (in Korean), Tachan Kunsok Hakhoechi, 24(10), 1185-1189 (1986) (Thermodyn., Theory, 19) Cassada, W.A., Shiflet, G.J., Starke, Jr, E.A., “The Effect of Germanium on the Precipitation and Deformation Behavior of Al-2Li Alloys”, Acta Metall., 34(3), 367-378 (1986) (Crys. Structure, Equi. Diagram, Experimental, 25) Evers, V.J., Oehlinger, G., Sextl, G., Becker, H.-O., “High Pressure LiGe with Layers of Two- and Four-Bond Germanium Atoms” (in German), Angew. Chem., 99(1), 69-71 (1987) (Crys. Structure, Experimental, 11) Goel, N.C., Cahoon, J. R., “The Al-Li-X Systems (X = Ag, As, P, B, Cd, Ge, Fe, Ga, H, In, N, Pb, S, Sb and Sn)”, Bull. Alloy Phase Diagrams, 10(5), 546-548 (1989) (Review, 25) Pavlyuk, V.V., Dmytriv, G.S., Starodub, P.K., “Crystal Structure of the Compounds of the Li-M-X (M = Mg, Al; X = Si, Ge, Sn) Systems” (in Russian), Cryst. Chem. Inorg. Coord. Compounds, VI Conf. (Abstact), L’viv, 210 (1992) (Crys. Structure, Experimental, 6) Pavlyuk, V.V., Dmytriv, G.S., Bodak, O.I., “Phase Equilibria in the Li-Al-Ge System at 470 K” (in Ukrainian), Dop. Akad. Nauk Ukrainy, (8), 84-86 (1993) (Equi. Diagram, Experimental, #, 6) Pavlyuk, V.V., “Synthesis and Crystal Chemistry of Lithium Intermetallic Compounds” (in Ukrainian), Summary of the Thesis for Doctor Science Degree, L’viv Univ., 1-35 (1993) (Crys. Structure, Experimental, Review, 49) Hosoda, H., Sato, T., Tezuka H., Mishima Y., Kamio A., “Substitution Behaviour of Additional Elements in the L12-Type Al3Li Metastable Phase in Al-Li Alloys”, J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Calculation, 26) Pavlyuk, V., Bodak, O., MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.14593.1.20, (1995) (Crys. Structure, Equi. Diagram, Assessment, 15) Dmytriv, G.S., “Phase Equilibria and Crystal Structure of Compounds in Mg-Li-Si, Ca-Li-{Si, Ge}, Al-Li-{Si, Ge, Sn}, Zn-Li-{Al, Sn}” (in Ukrainian), Summary of the Thesis for the Degree of Candidate of Science, 1-23 (1996) (Crys. Structure, Equi. Diagram, Experimental, 10) Sangster, J., Pelton, A.D., “The Ge-Li (Germanium-Lithium) System”, J. Phase Equilib., 18(3), 289-294 (1997) (Calculation, Crys. Structure, Review, Thermodyn., 31) Desk Handbook: Phase Diagrams for Binary Alloys, Okamoto, H., (Ed.), ASM (2000) (Equi. Diagram, Crys. Structure, Review) Goward, G.R., Taylor, N.J., Souza, D.C.S., Nazar, L.F., “The True Crystal Structure of Li 17M4 (M = Ge, Sn, Pb) - Revised from Li22 M5”, J. Alloys Compd., 329, 82-91 (2001) (Crys. Structure, Experimental, 14) Groebner, J., “Al-Li (Aluminium-Lithium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.13517.1.20, (2003) (Crys. Structure, Equi. Diagram, Assessment, 21)
Landolt-Börnstein New Series IV/11A3
Al–Ge–Li
55
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Li) < 180.6 (Al) < 660.452 (Ge) < 938.3
, Li9Al4 < 347 - 275
`, Li9Al4 < 275 , Li3Al2 < 520 , LiAl < 700 `, LiAl3 < 190 - ~120 Li7Ge12 < 510 LiGe < 540
Li12Ge7 < 510 Li9Ge4 < 740
Li14Ge6 < 770
Landolt-Börnstein New Series IV/11A3
Pearson Symbol/ Space Group/ Prototype cI2 Im3m W cF4 Fm3m Cu cF8 Fd3m C (diamond) mC26 C2/m Li9Al4 ? hR15 R3m Li3Al2 cF16 Fd3m NaTl cP4 Pm3m Cu3Au oP* Pnm21 Li7Ge12 tI32 I41/a MgGa tI24 I41/amd LiGe oP152 Pnma Li12Si7 oC52 Cmcm Na9Sn4 h** hR21 R3m Li14Si6
Lattice Parameters Comments/References [pm] a = 351.0
pure Li at 25°C [V-C2]
a = 404.96
pure Al at 25°C [Mas2] Dissolves up to 15 at.% Li pure Ge at 25°C [Mas2]
a = 1915.51 b = 542.88 c = 449.88 = 107.671° ?
[2003Gro]
a = 450.8 c = 1426
[2003Gro] 60 to 61 at.% Li [Mas2]
a = 637
at 50 at.% Li [2003Gro] 45 to 55 at.% Li [Mas2] 46 to 52 at.% Li at 200°C [1993Pav1] Metastable [2003Gro]
a = 403.8
a = 1154.1 0.3 b = 807.3 0.2 c = 1535.9 0.4 a = 975 2 c = 578 2 a = 981.0 0.3 c = 580.7 0.2 a = 405.29 0.01 c = 2328.2 0.3 a = 876.3 b = 2011.5 c = 1464 a = 449 b = 787 c = 2444 a = 449 c = 2444 a = 449.4 0.1 c = 1843.9 0.4
[Mas2]
[1981Gru, 1982Gru]
[1987Eve] [1982Gru] high pressure phase [1987Eve] [1981Gru, 1982Gru]
[V-C2, 1982Gru]
[1982Gru] [1981Gru, 1982Gru]
MSIT ®
Al–Ge–Li
56 Phase/ Temperature Range [°C] Li13Ge4 < 780 Li15Ge4 < 720 Li4Ge < 640
* -1, Li5.3Al0.7Ge2
Pearson Symbol/ Space Group/ Prototype oP34 Pbam Li13Si4 cI76 I43d Cu15Si4 cF416 F43m Li20Si5 cF432 F23 Li22Pb5 cF419 F43m Li17Ge4 hP8 P63/mmc Na3As
* -1´, Li16 Al2Ge6 hP24 * -2, Li9Al2Ge3 * -3, Li2AlGe
* -4, Li6Al3Ge * -5, LiAlGe
? cF F43m CuHg2Ti ? cF16 F43m LiAlSi
Lattice Parameters Comments/References [pm] a = 924 b = 1321 c = 463 a = 1072 a = 1082.5
[V-C2, 1982Gru]
a = 1892.9 0.1
[1982Gru]
a = 1886
Li22Ge5 [V-C2, Mas2]
a = 1875.6 0.2
Li17Ge4 [2001Gow]
a = 438.0 c = 816.2
[1978Ble] 'm = 2.42 g·cm-3 'x = 2.46 g·cm-3 [1993Pav1]
a = 438.0 c = 816.2 a = 758.6 c = 816.2 ? a = 616.3 a = 597.5 ? a = 598.9
a = 598.9 a = 597.7
* -6, LiAl2Ge
MSIT®
cF16 Fd3m NaTl cF16 Fm3m MnCu2Al
a = 599.8
[V-C2] [1982Gru]
[1978Ble] [1993Pav1] [1974Boc] 'm = 2.848 g·cm-3 [1992Pav, 1993Pav1] [1993Pav] [1976Sch] 'm = 3.27 g·cm-3 'x = 3.29 g·cm-3 [1992Pav, 1993Pav1] [1960Now] [1981Kis]
[1992Pav, 1993Pav1]
Landolt-Börnstein New Series IV/11A3
Al–Ge–Li
Fig. 1: Al-Ge-Li. Phase diagram of the Ge-Li system after [1982Gru]
1000 900
L
800
780
770 740 730 690
720
700
Temperature, °C
57
640 600
530
540
510
500
530
500
510
400 300 200 100
Li14Ge6 Li9Ge4
180
Li4Ge Li15Ge4 Li13Ge4
Li7Ge12
Li12Ge7 LiGe
0
Li
80
60
40
Ge
20
Li, at.%
Ge
Data / Grid: at.%
(Ge)
Fig. 2: Al-Ge-Li. Partial triangulation of the Al-Ge-Li ternary system
Axes: at.%
20
80
Li7Ge12 40
60
LiGe
Li12Ge760 Li9Ge4 Li14Ge6 Li13Ge4 Li15Ge4 Li4Ge
40
τ5 τ3
τ1
τ6
80
20
τ2 τ4
Li
Landolt-Börnstein New Series IV/11A3
20
δ´
40 γ
β
60
80
(Al)
Al
MSIT ®
58
Al–H–Li
Aluminium – Hydrogen – Lithium Oksana Bodak, Pierre Perrot Literature Data Two ternary hydrides have been prepared and characterized, LiAlH4 and Li3AlH6. The hydride LiAlH 4 is available as a commercial product. Crystal structure data for Li3AlH 6, obtained by the reaction of LiAlH4, LiH and Al(C2H5)3 in C6H5CH 3, were given by [1966Chi], Table 1. The crystal structure data for the Li3AlH 6 are given in [1985Bas2]. The thermal stability of LiAlH4 was studied by [1970Bra] using DTA, by [1972Dil] using DTA and thermogravimetric analysis and by [1985Bas1] using DSC. The first critical review of literature data, published until 1990, was made by [1993Fer, 1995Pav], followed by the present evaluation. The influence of mechanochemical processing of polycrystalline LiAlH4 was studied in [1999Zal, 2000Bal]. The enthalpy of formation LiAlHx was calculated using the Miedema’s model [2002Her]. Binary Systems The Al-Li system reported by [2003Gro] and the Al-H system as described by [2003Per] are accepted as terminal descriptions of the ternary Al-H-Li phase diagram. The H-Li is accepted from [Mas2]. Solid Phases All authors completely agree that two hydrides, LiAlH4 and Li3AlH6, are formed in this system. Their crystal structures were reported by [1967Skl, 1970Gor, 1985Bas1, 1985Bas2, 2000Bal] and are given in Table 1. [1967Skl] proposed a unit cell with an “a” parameter only half of what was adopted by the other workers. For the remaining cell parameters there is good agreement between the reported data. Mechanochemical processing of polycrystalline LiAlH4 revealed good stability of this complex aluminohydride during high-energy ball-milling in a helium atmosphere for up to 35 h. The decomposition of lithium aluminohydride into Li3AlH6, Al and H2 is observed during prolonged mechanochemical treatment for up to 110 h and is most likely associated with the catalytic effect of a third material, iron, which is introduced into the hydride as a contaminant during mechanical treatment [2000Bal]. According to [2000Bal] the attempts to solve the crystal structure of Li3AlH6 by X-ray powder diffraction data were unsuccessful because of the strong pseudosymmerty found in this compound. The unit cell volume of the rhombohedral lattice is 1.5 times greater than that of both primitive and base centered monoclinic lattices. Isothermal Sections LiAlH4 has a melting point of 163.7°C and decomposes at 160-180°C [1999Zal] according to the reaction: 3LiAlH4(liquid) Li3AlH 6(solid) + 2Al + 3H2. The standard Gibbs energy of this reaction at 298K was assessed to be -27.7 kJ·mol-1 [2000Bal]. At temperatures above 250°C the hydride Li3AlH6 decomposes: Li3AlH 6 3LiH + Al + 3/2H 2. According to [2000Bal] the temperature of decomposition is in the range 207-260°C which is in good agreement with data of [1999Zal]. The analogous ternary deuteride, LiAlD4, has its melting point at 167.5°C and decomposes at 195°C [1985Bas1]. The phase stability diagram at 500°C calculated by [1988Cro] is based on the assumption that only the LiAl phase occurs in the Al-Li binary system. The threephase regions identified were: Li+LiH+LiAl, H+LiH+LiAl and Al+H+LiAl. It should be pointed out that, if the hydrides are considered to be unstable at 500°C [1985Bas1], Al would react with LiH following the reaction: Al + LiH LiAl + 1/2H2 and the tie line LiAl-H of the Al-Li-H stability diagram would be stable. However, hydrides are stable under large hydrogen pressure and the existence of the LiAl-H tie line contradicts the decomposition of LiAlH4 into Li3AlH6+Al and the subsequent decomposition of Li3AlH6 into LiH+Al [1999Zal]. Figure 1 shows a stability diagram taking into account experimental observation. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–H–Li
59
Each of the triangles numbered 1 to 5 is characterized by a hydrogen pressure depending on the given temperature and decreasing from p1 to p5: p1 = p(AlH 3/Al), p2 = p(LiAlH 4/Li3AlH6 + Al), p3 = p(Li3AlH6/LiH + Al), p4 = p(LiH + Al/AlLi), p5 = p(LiH/Li). In Fig. 1 the dashed lines correspond to tie lines never observed experimentally. Between p4 and p5 one should actually observe the following equilibria: LiH + LiAl/Li3Al2 and LiH + Li 3Al2/Li9Al4. The solubility of hydrogen in equiatomic LiAl alloys was measured at 500°C as a function of hydrogen pressure between 204 and 716 mbar (204#102 and 716#102 Pa) by [1976Tal]. Sieverts’ Law was obeyed, with an average value of Sieverts’ constant of 2.20#104 0.15 mbar1/2/atomic fraction H 2 (Table 2). [1988Any] determined the solubility of H2 in molten Al-Li alloys containing 1, 2 and 3 mass% Li (3.8, 7.4 and 10.7 at.% Li, respectively) from 670°C to 800°C and from 5.3#104 Pa to 10.7#104 Pa. Sieverts’ Law was obeyed for all three alloys; the solubility of H2 increases with increasing Li content (Table 3). [1990Fed] quoted data for the solubility of H2 in the Al-2Li (mass%) alloy. At 700°C the data are in good agreement with [1988Any]. The solubility of H2 in molten Al-Li alloys containing up to 4 mass% Li was measured by [1989Lin] for temperatures of 700, 800, 900 and 1000°C. At 700°C the calculated H2 solubilities are lower than determined by [1988Any, 1990Fed]. Interaction parameters for Al-H-Li melts were calculated for 927°C by [1986Lee]. A more general expression of the first order interaction parameter of Li upon H has been proposed by [2003Ma]: eH(Li) = (d ln H/ d (mass% Li)) = -0.138 - 158.2/T. A negative value of the interaction parameter means that the presence of Li increases the solubility of H in liquid Al; this result is already confirmed experimentally by [1988Any] and theoretically by [1989Lin]. Thermodynamics The molar heat capacity of LiAlH4 [1978Cla, 1979Bon, 1985Bas1], of LiAlD 4 [1985Bas1] and Li3AlH6 [1978Cla, 1979Bon] at 298.15 K are given in Table 4. According to [2002Her] the calculated enthalpy of formation using the Miedema’s model was -69 kJ·mol-1 for LiAlH4 and -86 kJ·mol -1 for Li3AlH 6. Notes on Materials Properties and Applications Besides its well-known application as a reducing agent in organic synthesis, LiAlH4 contains 10.5 mass% H, which is one of the highest values among hydrides. Thus LiAlH4 is of considerable interest as potential ultra-high capacity hydrogen storage solid. Miscellaneous [1982Wak] determined the electrical resistance of LiAlH4 at pressures up to 125 kbar. The resistance decreases with applied pressure up to 75 kbar and remains virtually constant from 75 to 125 kbar. Adsorption and desorption of hydrogen in Al and Al-Li alloys were presented and discussed by [1988Wat]. References [1965Ame] [1966Chi]
[1967Skl] [1970Bra]
Landolt-Börnstein New Series IV/11A3
Amendola, A., Index Inorganic to the Powder Diffraction File 1965, American Society for Testing and Materials, Philadelphia, Pa, n.12473, p.469 (1965) as quoted in [1970Gor] Chini, P., Baradel, A., Vacca, C., “The Reaction of Aluminum with Hydrogen and Natriumfluoride” (in Italian), La Chimica e l’Industria, Special, 48(6), 596-601 (1966) (Crys. Structure, Experimental, 23) Sklar, N., Post, B., “The Crystal Structure of LiAlH 4”, Inorg. Chem., 6, 669-671 (1967) (Crys. Structure, Experimental, 4) Brachet, F.-G., Etienne, J.-J., Mayet, J., Tranchant, J., “Structure and Properties of LiAl Hydrides. III. Differential Thermal Analysis and Isothermal (70 and 130°C) Thermal
MSIT ®
60
[1970Gor]
[1972Dil] [1976Tal] [1978Cla]
[1979Bon]
[1981Gor]
[1982Wak]
[1985Bas1]
[1985Bas2]
[1986Lee]
[1988Any] [1988Cro]
[1988Wat] [1989Lin] [1990Fed] [1993Fer]
[1995Pav]
MSIT®
Al–H–Li Decomposition of LiAlH” (in French), Bull. Soc. Chim. Fr., (11), 3799-3807 (1970) (Experimental, 14) Gorin, P., Marchon, J. C., Tranchant, J., Kovacevic, S., Marsault, J. P., “Structure and Properties of LiAl Hydrides. II. Structure of LiAlH4 in the Crystalline State and in Diethyl Ether Solutions” (in French), Bull. Soc. Chim. Fr., (11), 3790-3799 (1970) (Crys. Structure, Experimental, 27) Dilts, J.A., Ashby, E.C., “A Study of the Thermal Decomposition of Complex Metal Hydrides”, Inorg. Chem., 11(6) 1230-1236 (1972) (Experimental, 27) Talbot, J. B., Smith, F. J., Land, J. F., Barton, P., “Tritium Sorption in Li-Bi and Li-Al Alloys”, J. Less-Common Met., 50, 23-28 (1976) (Experimental, 10) Claudy, P., Bonnetot, B., Letoffe, J.M., Turck, G., “Determination of Thermodynamic Constants of Simple Hydrides of Aluminium. IV. Enthalpy of Formation of LiAlH4 and Li 3AlH 6” (in French), Thermochim. Acta, 27, 213-221 (1978) (Thermodyn., Experimental, 11) Bonnetot, B., Claudy, P., Diot, M., Letoffe, J.M., “LiAlH 4 and Li3AlH 6: Molar Heat Capacity and Thermodynamic Properties from 10 to 300K”, J. Chem. Thermodyn., 11, 1197-1202 (1979) (Thermodyn., Experimental, 8) Gorbunov, V.E., Gavrichev, K.S., Bakum, S.I., “Thermodynamic Properties of LiAlH 4 in the Temperature Range 12-300 K”, Russ. J. Inorg. Chem. (Engl. Transl.), 26, 168-169 (1981) (Thermodyn., Experimental, 8) Wakamori, K., Sawaoka, A., Filipek, S.M., Baranowski, B., “Electrical Resistance of Some Alkaline Earth Metal Hydrides and Alkali Metal Al Hydrides and Borohydrides Under High Pressure”, J. Less-Common Met., 88, 217-220 (1982) (Experimental, 6) Bastide, J.-P., Bonnetot, B., Letoffe, J.-M., Claudy, P., “Comparative Study of LiAlH4 and LiAlD4. I. Preparation, Crystallography and Thermal Behaviour- Evidence for a Metastable Form of LiAlD4”, Mater. Res. Bull., 20, 999-1007 (1985) (Crys. Structure, Experimental, 16) Bastide, J.-P., Bonnetot, B., Letoffe, J.-M., Claudy, P., “Structural Chemistry of Some Complex Hydrides of Alkaline Metals”, Stud. Inorg. Chem., 3, 785-788 (1983) (Crys. Structure, Experimental, 16) Lee, J.J., Sommer, F., “Thermodynamic Properties of Li in Liquid Aluminum Alloys” (in Korean), Taehan Kumsok Hakhoechi, 24(10), 1185-1189 (1986) (Thermodyn., Theory, Experimental, 23) Anyalebechi, P.N., Talbot, D.E., Granger, D.A., “The Solubility of H 2 in Liquid Binary AlLi-Alloys”, Metall. Trans. B, 19, 227-232 (1988) (Thermodyn., Experimental, 24) Crouch-Baker, S., Huggins, R.A., “Phase Behaviour in the Li-Al-O-H System at Intermediate Temperatures”, Solid State Ionics, 28-30, 611-616 (1988) (Equi. Diagram, Thermodyn., Theory, 21) Watson, J.W., “Hydrogen in Aluminum and Aluminum-Lithium Alloys”, Thesis, Northwestern University, 1-366 (1988) (Experimental, 147) Lin, R.Y., Hoch, M., “The Solubility of Hydrogen in Molten Aluminum Alloys”, Metall. Trans. A, 20(9), 1785-1791 (1989) (Equi. Diagram, Thermodyn., Calculation, Theory, 31) Fedosov, A.S., Danilkin, V.A., Makarov, G.S., “The Interaction of Al-Li Alloy Metals with Hydrogen” (in Russian), Tsvetn. Met., (8), 88-90 (1990) (Experimental, 4) Ferro, R., Saccone, A., Delfino, S., “Aluminium-Hydrogen-Lithium”, in “Ternary Alloys: A Comprehensive Compendium of Evaluated Constitutional Data and Phase Diagrams” Petzow, G., Effenberg, G. (Eds.), Vol. 6, VCH, Weinheim, 111-112 (1993) (Crys. Structure, Review, 9) Pavlyuk, V., Bodak, O., “Aluminium-Hydrogen-Lithium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.12744.1.20, (1995) (Crys. Structure, Equi. Diagram, Assessment, 17) Landolt-Börnstein New Series IV/11A3
Al–H–Li [1999Zal]
[2000Bal]
[2002Her] [2003Gro]
[2003Ma] [2003Per]
61
Zaluski, L., Zaluska, A., Ström-Olsen, J.O., “Hydrogenation Properties of Complex Alkali Metal Hydrides Fabricated by Mechano-Chemical Synthesis”, J. Alloys Compd., 290, 71-78 (1999) (Experimental, 22) Balema, V.P., Pecharsky, V.K., Dennis, K.W., “Solid State Transformations in LiAlH4 during High-Energy Ball-Milling”, J. Alloys Compd., 313, 69-74 (2000) (Equi. Diagram, Crys. Structure, Experimental, 22) Herbst, J.F., “On Extending Miedema’s Model to Predict Hydrogen Content in Binary and Ternary Hydrides”, J. Alloys Compd., 337, 99-107 (2002) (Thermodyn., Calculation, 20) Groebner, J., “Al-Li (Aluminium-Lithium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 21) Ma, Z., Janke, D., “Solution Behawior of Hydrogen in Aluminium and ist Alloys Melts”, Metall, 57(9), 552-556 (2003) (Thermodyn., Calculation, Review, 14) Perrot, P., “Al-H (Aluminium-Hydrogen)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 21)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Li) < 180.6
cI2 Im3m W
a = 351.0
pure Li at 25°C [V-C2]
(Al) < 660.45
cF4 Fm3m Cu
a = 404.96
pure Al at 25°C [Mas2] dissolves up to 15 at.% Li
Li9Al4 < 347 - 275
mC26 C2/m Li9Al4
a = 1915.51 b = 542.88 c = 449.88 = 107.671°
[2003Gro]
Li9Al4 ( ´) < 275
?
?
[Mas2]
Li3Al2 ( ) < 520
hR15 R3m Li3Al2
a = 450.8 c = 1426
[2003Gro] 60 to 61 at.% Li [Mas2]
LiAl () < 700
cF16 Fd3m NaTl
a = 637
at 50 at.% Li [2003Gro] 45 to 55 at.% Li [Mas2]
LiAl3 (´) < 190 - ~120
cP4 Pm3m Cu3Au
a = 403.8
metastable [2003Gro]
LiH
cF8 Fm3m NaCl
a = 408.3
[V-C2]
AlH 3 < 110
hR24 R3c
a = 445.6 c = 1183
[2003Per] metastable
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–H–Li
62 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
AlH3 < 80 * LiAlH4
[2003Per] metastable “Aluminum hydrogenoaluminate” Al(AlH4)3 mP48
mP48
mP48
mP24
a = 960 b = 786 c = 790 = 112.5° a = 967.9 b = 781.0 c = 792.5 = 112.53° a = 967.9 b = 788.1 c = 791.2 = 111.88° a = 484.5 b = 782.6 c = 791.7 = 112.5°
[1965Ame] 'x = 0.908 g·cm-3 'm = 0.917 g·cm-3 [1970Gor] 'x = 0.911 g·cm-3 'm = 0.95 g·cm-3 [1985Bas1] 'x = 0.900 g·cm-3 'm = 0.907 g·cm-3 [1967Skl] 'x = 0.904 g·cm-3 'm = 0.92 g·cm-3
* Li3AlH6
m**
a = 571.5 a = 539.1 c = 569.4 = 91.33°
[1966Chi]
* -Li3AlH 6
mP* P21/m LiAlSi2O6
a =790.5 b = 812.5 c = 567.5 = 92.7° a = 566.7 0.1 b = 810.7 0.2 c = 791.7 0.2 = 92.07 0.01° a = 791.7 0.2 b = 810.7 0.2 c = 566.7 0.1 = 92.07 0.01° a = 811.3 0.1 c = 957.0 0.1
[1985Bas2] high-pressure phase, 500°C, 50 kbar, pseudo-cubic [2000Bal] prepared mechano-chemically
a = 1114 b = 1145 c = 1034
[1985Bas2]
mP* P21/c
mC* C2/m
hR* R3m * -Li3AlH6
MSIT®
o* Li3Al2 (LiF4)3
[2000Bal]
[2000Bal]
Landolt-Börnstein New Series IV/11A3
Al–H–Li
63
Table 2: Solubility of H2 in LiAl at 500°C [1976Tal] H2 Pressure, (p) [mbar] 204 307 420 716
H 2 Concentration, (N) [atomic fraction] 6.04#10-4 8.07#10-4 9.94#10-4 12.2#10-4
Sieverts’ Constant, (p/N) 1/2 [mbar]-1/2 atomic fraction H2) 2.36#104 2.17#104 2.06#104 2.19#104
Table 3: Solubility of H 2 in Molten Al-Li Alloys. S is the solubility expressed in cm3 H2 measured at 273 K and 101.325 Pa; S° is the standard value: S° = 1cm3 measured at 273 K and 101.325 Pa; p is the pressure expressed in Pa; p° is the standard pressure: p° = 101.325 Pa log(S/S°) - 1/2 log(p/p°) = -2113/T + 2.568 log(S/S°) - 1/2 log(p/p°) = -2997/T + 3.329 log(S/S°) - 1/2 log(p/p°) = -2889/T + 3.508
1 mass% Li: 2 mass% Li: 3 mass% Li:
Table 4: Molar Heat Capacity of LiAlH4, LiAlD 4 and Li3AlH 6 Phase
Molar Heat Capacity, Cp, [J#K-1#mol-1] at 296.15 K
Reference
LiAlH4
89.2 83.19 83.01 82.60
[1978Cla] [1979Bon] [1981Gor] [1985Bas1]
LiAlD4
92.70
[1985Bas1]
Li3AlH6
131.0 127.75
[1978Cla] [1979Bon]
H Fig. 1: Al-H-Li. Stability diagram. The triangles 1 to 5 are characterized by hydrogen pressure at equilibrium decreasing from p1(AlH3/Al) to p5(LiH/Li) (see text)
Data / Grid: at.% Axes: at.%
20
80
AlH3 LiAlH4
40
60
Li3AlH6
LiH
1
2
60
40
3 4
80
20
5
(Al)
Li Landolt-Börnstein New Series IV/11A3
20
Li9Al4
40 Li Al LiAl 3 2
60
80
Al MSIT ®
64
Al–H–Mg
Aluminium – Hydrogen – Magnesium Lazar Rokhlin, updated by Volodymyr Ivanchenko Literature Data The solubility of hydrogen in Al-Mg alloys was measured for different temperatures and composition ranges using a range of different experimental techniques. [1973Hua] used a modified Sieverts apparatus for determination of solubility of hydrogen in pure magnesium and its alloys including Al-Mg system. It was shown that alloying of magnesium with 10 at.% Al lowered the solubility of hydrogen at 700°C and pH2 = 10 5 Pa from 50 cm3 H2/100 g to 40 cm3 H2/100 g (hydrogen volumes measured at 273 K under 101325 Pa). These values are very close to the values calculated by [1965Bur]. [1974And] studied the solubility of hydrogen in (Al) solid solution with 0.45 and 4.75 at.% Mg at 500°C using saturation and vacuum extraction and showed that alloying with Mg raised the hydrogen solubility from 0.012 cm3 H2/100 g (for pure Al) to 0.04 0.01 (for 0.45 at.% Mg) and to 0.06 cm3 H 2/100 g (for 4.75 at.% Mg). These results are significantly lower than those presented by [1976Wat]. [1974Gab] studied the solubility of hydrogen in phase (Mg 2Al3) in temperature interval from 380 to 560°C using high pressure Sieverts apparatus and high temperature vacuum extraction. Under crystallization the hydrogen solubility in Mg2Al3 dropped from 5.9 cm3 H2/100 g to 1.45 cm3 H2/100 g. [1976Lev] studied the porosity of Al-Mg alloys which is caused by hydrogen. [1977Che] investigated permeability, diffusivity and solubility of hydrogen at temperatures from 650 to 800°C in liquid Al-Mg alloys containing up to 16 mass% Al. [1981Tuc] studied the hydrogen saturation of Al-Mg alloys exposed to water-vapor saturated air at elevated temperature. Reversible hydrogen storage in magnesium alloys was reviewed by [1978Gui]. They reported that phase (Mg17 Al12 , sometimes designated as Mg3Al2) did not hydride at 350°C under hydrogen pressure from 3 to 5 MPa. These results are in contradiction with those presented by [1980Min, 1981Gav], who studied the reactions of hydrogen with Mg2Al3 and Mg17Al12 and reported their main features: hydrogenation of the intermetallic Al-Mg compounds resulting in disproportionation; namely, for Mg17Al12 the reaction may be written as: Mg2Al3+2H 22MgH 2+3Al; while for Mg3Al2, the reaction may be written: Mg17Al12+9H29MgH2+4Mg2Al3 [1983Sem] pointed that Al-Mg alloys dissolved only a very small quantity of hydrogen due to very low rate of process. Differential scanning calorimetry and gas chromatography were used to investigate and quantify the reactions occurring when Al-5Mg (mass%) alloy, previously exposed to water-vapor saturated air, were heated from ambient temperature to 600°C. [1984Lue] measured the equilibrium hydrogen pressure at 142 and 170°C of the three phase fields MgH2+(Mg)+ , MgH2+ +, and MgH2++(Al). The H was introduced into Al-Mg alloys by electrolysis in an organometallic melt, NaAlEt4, containing dissolved Na+H- as electrolyte. [1985Lue1, 1985Lue2] discussed the results thermodynamically. [1978Cla1] prepared a ternary hydride Mg(AlH4)2 by reaction of NaAlH 4 with MgCl2 dissolved in tetrahydrofurane and measured its heat capacity at room temperature, to be 136 J#(mol#K)-1. Since [1985Lue1, 1985Lue2] did not find this ternary hydride, it may possibly be stable only under high hydrogen pressure, an assumption supported by the method of sample preparation. A new theoretical method of describing and investigating metal hydrides has been developed by [1987Lue]. It involves thermodynamics and interprets the hydrogenation reaction by ternary phase diagrams. [1987Lue] showed that intermetallic compounds formed by elements of the boron group with magnesium form two phase regions with MgH2 in the ternary phase diagrams. Thus the hydrogen pressures of the resulting three phase equilibria will be higher than or equal to the value of Mg/MgH2 equilibrium. In these systems, including Al-H-Mg, ternary hydrides are not taken into account. The solubility of hydrogen in molten aluminium alloys containing magnesium has been calculated from the solubility of hydrogen in pure metals and binary metal-metal interaction parameters by [1989Lin].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–H–Mg
65
The structure and hydrogen absorption properties of Al-Mg alloys prepared by high-energy ball milling were studied over the whole composition range in their as milled and Al-leached forms by [2000Bou]. The latter were obtained from the milled materials by leaching out of Al in a 1N NaOH solution. Their results on the interaction of intermetallic phases with hydrogen are in good agreement with those of [1980Min, 1981Gav]. [2002Her] used Miedema’s model to predict the hydrogen content and the enthalpy of formation of hypothetical ternary hydrides in the Al-H-Mg system. Binary Systems The binary systems Al-H [2002Per], Al-Mg [2003Luc] and H-Mg [2001Per] are accepted to present the best boundary systems for the Al-H-Mg ternary system. Solid Phases One ternary phase has been reported, Mg(AlH4)2, which is stable under high hydrogen pressure. The phase AlH3 is known to have two polymorphic modifications which are both metastable [1978Cla1, 1978Cla2, 1979Cla]. Chemically AlH3 is stable at room temperature and decomposes when heated at 110°C [1980Her]. Under high hydrogen pressures (2 GPa at 300°C and 6 GPa at 600°C), it is possible to synthesize AlH3 reversibly [1992Kon]. All solid phases are listed in Table 1. Isothermal Sections Figure 1 shows the isothermal section between 140 and 170°C [1984Lue, 1985Lue1, 1985Lue2]. The section is corrected to the accepted homogeneity ranges of the Al-Mg phases: (Al), (Mg), and . Figure 1 shows that MgH2 is in equilibrium with , , , and (Mg) phases. Mg(AlH4)2 and AlH3 hydrides are stable phases at these temperatures under hydrogen pressure higher than 100 kPa. Figure 2 shows the solubility of hydrogen in liquid Al-Mg alloys at 500, 700 and 800°C. It is taken from [1976Wat, 1989Lin] with small corrections to match the solubility in Al given in the Al-H system by [2002Per]. From the activity coefficients of hydrogen in molten Al-Mg alloys at 827°C [1989Lin], the interaction coefficient of Mg upon H in liquid Al may be assessed: JH(Mg) = (dln H/dxMg ) = -8.12 at 827°C. This negative value means that Mg in liquid Al increases the solubility of H. Al-Mg alloys show liquid-solid two-phase fields at 500°C. There, the hydrogen solubility must be represented by a straight line. The pressure-composition isotherms of the Mg2Al3-H system for temperature interval from 335 to 410°C are presented in Fig. 3 [1980Min]. These isotherms are in fair agreement with the measurements of [2000Bou] at 350°C who observed a plateau towards 0.8 MPa for Mg75Al25, corresponding to the Mg-MgH2 equilibrium and a plateau towards 1 MPa for Mg58Al42, corresponding to the equilibrium /+MgH 2. Figure 3 shows that under 5 MPa H2 the global composition of the hydride is Mg 2Al3H7. The corresponding point lies inside the (Al)-MgH2-MgAl2H 6 triangle in Fig. 1, which confirms the formation of the ternary phase under high hydrogen pressures. Thermodynamics The dependence of the equilibrium pressure on temperature for the disproportionation reaction 1/2Mg2Al3+H2=MgH2+3/2Al was reported by [1980Min, 1981Gav] as log10 (p/Pa) = -3306/T+11.47. Hydrogen activity and Gibbs energy changes for the three-phase reactions in the Al-H-Mg system, as measured electrochemically at 142°C were presented by [1985Lue1, 1985Lue2, 1987Lue], as Mg17Al12 - Mg - MgH2 aH2 = 2.7 #10 -3 G = - 20080 J#(mol H 2)-1 -2 Mg2Al3 - Mg17Al12 - MgH 2 aH2 = 1.1 #10 G = - 15481 J#(mol H 2)-1 -2 Al -Mg2Al3 - MgH 2 aH2 = 2.3 #10 G = - 12970 J#(mol H 2)-1. These values are about 1.5 kJ lower than the accepted values. For instance, the first figure ( G = -20080 J#(mol H2)-1) which corresponds to the Mg/MgH2 equilibrium has to be compared with the value (-18833 J#(mol H 2)-1) accepted by [2001Per] at 142°C. The last value ( G = -12970 J#(mol H2)-1) has to be Landolt-Börnstein New Series IV/11A3
MSIT ®
66
Al–H–Mg
compared with ( G = -11870 J#(mol H2)-1) calculated from the expression of [1980Min, 1981Gav] given above. [2002Her] predicted the enthalpy of formation of some virtual hydrides as Mg17Al12Hx xcalc = 29.64 Hcalc (Xcalc ) = - 65000 J#(mol f.u.)-1 Mg2Al3H x xcalc = 5.83 Hcalc (Xcalc ) = - 47000 J#(mol f.u.)-1 MgAl2Hx xcalc = 3.37 Hcalc (Xcalc ) = - 43000 J#(mol f.u.)-1 The molar heat capacity of the hydride Mg(AlH4)2 has been measured at 25°C by means of a Calvet microcalorimeter as Cp = 136 J#mol-1#K-1. Notes on Materials Properties and Applications The Al-Mg system is of great importance for developing of many of the Al based and Mg based multicomponent light alloys used in avionic and space industry. The Al-Mg alloys are also of potential interest as materials for hydrogen storage. Miscellaneous The alloying of Al with Mg dramatically raises the absorption capacity of Al [1976Lev]. [1981Tuc] showed evidence for the formation of MgH2 on the grain boundaries of Al-Mg alloys when exposed to water-vapor saturated air at 70°C and for about 50 days. These authors suggest that its presence plays a prominent role in the pre-exposure embrittlement and stress-corrosion cracking of Al-Mg alloys. The diffusion of hydrogen in liquid Al-Mg alloys at temperatures from 650 to 800°C is slowly changed by raising Al contents up to 5.5 mass% Al [1977Che]. The faster rise of DH was observed in concentration interval of 5.5 to 12 mass% Al; after that D H raised slowly up to 16 mass% Al. For pure magnesium DH(650°C) = 1.5#10-8 m2#s-1 and activation energy is ED = 31380 1670 J#mol-1. For Al-Mg alloys: DH( 5.5 mass% Al, 650°C) = 1.7#10-8 m2#s-1 and ED(5.5 mass% Al) = 34730 1670 J#mol-1; DH(12 mass% Al, 650°C) = 8#10-8 m2#s-1 and ED(12 mass% Al) = 33470 1670 J#mol-1; DH(16 mass% Al, 650°C) = 1#10 -7 m2#s-1 and ED(16 mass% Al) = 33470 1670 J#mol-1. But at 7.5 mass% the Al activation energy has a maximum at ED(7.5 mass% Al) = 50210 1670 J#mol-1. [2000Bou] showed that the measured hydrogen capacity of the as milled material decreases with Al content, from H/M = 1.74 for pure un-milled Mg, to 1.38 for Mg/Al = 90/10, and then to 1.05 for Mg/Al = 75/25. In each case, there is a further 10-15% decline of the hydrogen absorption capacity after leaching. In the case of Mg/Al = 58/42, which basically contains a nanocrystalline Mg17Al12 intermetallic phase, only, hydriding leads to the formation of MgH2 and Al. This reaction is totally reversible and Mg17 Al12 is recovered upon dehydriding. In each case, there is an increase in the kinetics of hydrogen absorption and desorption following leaching. This change in the sorption kinetics is thought to arise as a consequence of the presence of Al solutes in the hexagonal structure of Mg, rather than to be due to purely geometric effects, such as the increase of the surface area. References [1965Bur] [1973Hua] [1974And]
[1974Gab]
MSIT®
Burylev, B.P., “The Solubility of Hydrogen in Magnesium Alloys” (in Russian), Liteynoe Proizvod., 9(1), 25-26 (1965) (Thermodyn., Theory, 11) Huang, Y.C., Watanabe, T., Komatsu, R., “Hydrogen in Magnesium and its Alloys”, Proc. 4th Internat. Conf. Vacuum Metallurgy, 176-179 (1973, published 1974) (Experimental, 8) Andreev, L.A., Levchuk, B.V., Gel’man, B.G., Danilkin,V.A., Kharin, P.A., Myagkov, E.A, “The Solubility of H in Al-Mg Alloys” (in Russian), Tekhnol. Legk. Splavov, Nauch. Byul. VILSa., (4), 58-62 (1974) (Experimental, 8) Gabidullin, R.M., Shvetsov, I.V., Kolachev, B.A., Archakov, Yu.I., “The Solubility of Hydrogen in Intermetallic Compounds of Aluminium with Magnesium, Copper, Manganese, Titanium and Zirconium” (in Russian), in “Constitution, Properties and Application of Metallides”, Kornilov I.I., Matveeva N.M., (Eds.), Nauka, Moscow, 188-190 (1974) (Experimental, 2) Landolt-Börnstein New Series IV/11A3
Al–H–Mg [1976Lev] [1976Wat]
[1977Che]
[1978Cla1]
[1978Cla2]
[1978Gui]
[1979Cla]
[1980Her]
[1980Min]
[1981Gav]
[1981Tuc]
[1982Mur] [1983Sem]
[1984Lue]
[1985Lue1]
[1985Lue2]
Landolt-Börnstein New Series IV/11A3
67
Levchuk, B.V., Andreev, L.A., “Interaction of Al-Mg Alloys with H” (in Russian), Metalloved. Term. Obrab. Met., (7), 23-27 (1976) (Experimental, 10) Watanabe, T., Tachihara, T., Huang, Y.C., Komatsu, R., “The Effect of Various Alloying Elements on the Solubility of Hydrogen in Magnesium” (in Japanese), J. Jpn. Inst. Light Met., 26(4), 167-174 (1976) (Experimental, 25) Chernega, D.F., Gotvyanskii, Yu.Ya., Prisyazhnyuk, T.N., “Permeability, Diffusion and Solubility of Hydrogen in Magnesium-Aluminum Alloys” (in Russian), Liteinoe Proizvod., (12), 9-10 (1977) (Experimental, 4) Claudy, P., Bonnetot, B., Letoffe, J.M., Turck, G., “Determination of the Thermodynamic Constants of Simple and Complex Al Hydrides. II. Measurements of Molar Heat Capacities at 298 K” (in French), Thermochim. Acta, 27, 199-203 (1978) (Thermodyn., Experimental, 10) Claudy, P., Bonnetot, B., Letoffe, J.M., “Determination of Thermodynamic Constants of Simple and Complex Aluminium Hydrides. III. Enthalpy of Formation of AlH3 and AlH3” (in French), Thermochim. Acta, 27, 205-211 (1978) (Thermodyn., Experimental, 12) Guinet, P., Halotier, D., Perroud, P., “Hydrogen Sorage by Means of Reversible Magnesium Alloys”, Eur. Communities Rep., EUR 1978, EUR 6085. Semin. Hydrogen Energy Vector: Prod., Use, Transp., 373-391 (1978) (Experimental, 20) Claudy, P., Bonnetot, B., Letoffe, J.M., “Preparation, Physicochemical Properties and Enthalpy of Formation of Aluminium Hydride -AlH3” (in French), J. Therm. Anal., 16(1), 151-162 (1979) (Thermodyn., 16) Herley, P.J., Christofferson, O., Todd, J.A., “Microscopic Observations on the Thermal Decomposition of -Aluminum Hydride”, J. Solid State Chem., 35, 391-401 (1980) (Experimental, 15) Mintz, M.H., Gavra, Z., Kimmel,G., “The Reaction of Hydrogen with Magnesium Alloys and Magnesium Intermetallic Compounds”, J. Less-Common Met., 74, 263-270 (1980) (Thermodyn., Experimental, 16) Gavra, Z., Hadari, Z., Mintz, M.H., “Effects of Nickel and Indium Ternary Additions on the Hydrogenations of Mg-Al Intermetallic Compounds”, J. Inorg. Nucl. Chem., 43, 1763-1768 (1981) (Thermodyn., Review, 11) Tuck, C.D.S., “Evidence for the Formation of Magnesium Hydride on the Grain Boundaries of Al-Mg and Al-Zn-Mg Alloys During their Exposure to Water Vapour”, in “Hydrogen Eff. Met. ”, Proc. 3rd Int. Conf., 1980 (Publ. 1981), 503-511 Bernstein I.M., Thompson, A.V., (Eds.), Metall. Soc. AIME, Warrendale, USA, (1981) (Experimental, 23) Murray, J.L., “The Al-Mg (Aluminum-Magnesium) System”, Bull. Alloy Phase Diagrams, 3, 60-74 (1982) (Review, Equi. Diagram, Thermodyn., 112) Semenenko, K.N., Verbettskii, V.N., Kotchukov, A.V., Sytnikov, A.N., “Reaction of Magnesium Containing Intermetallic Compounds and Alloys with Hydrogen” (in Russian), Vestn. Mosk. Uni., Ser. 2: Khim., 24(1), 16-27 (1983) (Thermodyn., Review, 46) Luedecke, C.M., Deublein, G., Huggins, R.A., “Use of Electrochemical Methods to Study and Control Hydrogen Storage in Solid Metal Hydrides”, Adv. Hydrogen Energy, 4, (Hydrogen Energ. Prog. 5, Vol. 3) 1421-1431 (1984) (Equi. Diagram, Thermodyn, Experimental, #) Luedecke, C.M., Deublein, G., Huggins, R.A., “Electrochemical Investigation of Hydrogen Storage in Metal Hydrides”, J. Electrochem. Soc.: Electrochem. Sci. Techn., 132(1), 52-56 (1985) (Thermodyn., Experimental, 29) Luedecke, C.M., Deublein, G., Huggins, R.A., “Investigation of Metal Hydrides with Thermodynamic Calculations and Electrochemical Experiments”, Hydrogen Syst. Pap. Int. Symp Meeting Date, 1, 363-377 (1985) (Equi. Diagram, Thermodyn., Experimental, #, 18)
MSIT ®
Al–H–Mg
68 [1987Lue]
[1989Lin] [1992Kon]
[1992San] [1998Lia]
[2000Bou]
[2001Per]
[2002Her] [2002Per]
[2003Luk]
Luedecki, C.M., Deubleiun,G., Huggins, R.A., “Thermodynamic Characterization of Metal Hydrogen Systems by Assessment of Phase Diagrams and Electrochemical Measurements”, Int. J. Hydrogen Energy, 12(2) 81-88 (1987) (Equi. Diagram, Thermodyn., Review, #, 18) Lin, R.J., Hoch, M., “The Solubility of Hydrogen in Molten Aluminium Alloys”, Metall. Trans. A, 20(9), 1785-1791 (1989) (Theory, Thermodyn., 31) Konovalov, S.K., Bulchev, B.M., “High Pressures in the Chemistry of Beryllium and Aluminium Hydrides”, Russ. J. Inorg. Chem., 37(12), 1361-1365 (1992), translated from Zh. Neorg. Khim., 37, 2640-2646 (1992) (Equi. Diagram, Experimental, 16) San Martin, A., Manchester, F.D., “The Al-H (Aluminum-Hydrogen) System”, J. Phase Equilib., 13(1), 17-21 (1992) (Equi. Diagram, Review, 45) Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G., Seifert, H.J., Lukas, H.L., Aldinger, F., “Experimental Investigation and Thermodynamic Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89(8), 536-540 (1998) (Experimental, Assessment, Calculation, Equi. Diagram, Thermodyn., 33) Bouaricha, S., Dodelet, J.P., Guay, D., Huot, J., Boily, S., Schulz, R., “Hydriding Behavior of Mg-Al and Leached Mg-Al Compounds Prepared by High-Energy Ball-Milling”, J. Alloys Compd., 297, 282-293 (2000) (Equi. Diagram, Crys. Structure, Experimental, 27) Perrot, P., Schmid-Fetzer, R., “Hydrogen-Magnesium”, in “Ternary Alloys: A Comprehensive Compendium of Evaluated Consitutional Data and Phase Diagrams”, Effenberg, G., Aldinger, F., Rogl, P. (Eds.), Vol. 18, MSI, Materials Science International Services GmbH, Stuttgart, 3-4 (2001) (Thermodyn., Assessment, Equi. Diagram, #, 6) Herbst, J.F., “On Extending Miedema’s Model to Predict Hydrogen Content in Binary and Ternary Hydrides”, J. Alloys Compd., 337, 99-107 (2002) (Calculation, Thermodyn., 20) Perrot, P., “Al-H (Aluminium-Hydrogen)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14832.1.20, (2002) (Equi. Diagram, Crys. Structure, Assessment, 21) Lukas, H.-L., Lebrun, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Al)
hP2 P63/mmc Mg
a = 269.3 c = 439.8
at 25°C, 20.5 GPa [Mas2]
(Al) < 660.452
cF4 Fm3m Cu
a = 404.96
at 25°C [Mas2] 100 to 81.4 at.% Al at 450°C [1982Mur]
(Mg) < 650
hP2 P63/mmc Mg
a = 320.94 c = 521.07
at 25°C [Mas2] 0 to 11.5 at.% Al at 437°C [1982Mur]
, Mg17Al12 458
cI58 I43m Mn
a = 1054.38
at 41.4 at.% Al [V-C2] 39.5 to 51.5 at.% Al [1998Lia] 40 to 52 at.% Al [2003Luk]
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–H–Mg
69
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, Mg 2Al3 452
cF1168 Fd3m Mg2Al3
a = 2816 to 2824
60-62 at.% Al [2003Luk] 1168 atoms on 1704 sites per unit cell [2003Luk]
J, Mg23Al30 410 - 250
hR159 R3 Mn44Si9
a = 1282.54 c = 2174.78
54.5-56.5 at.% Al [2003Luk] Structure: 159 atoms refer to hexagonal unit cell [2003Luk]
AlH 3 < 110
hR24 R3c
a = 445.6 c = 1183
[1992San], metastable
AlH 3 < 80
-
-
Metastable “Aluminum hydrogenoaluminate” Al(AlH4)3 [1978Cla2]
MgH2
tP6 P42/mnm TiO2
a = 451.68 c = 302.05
[P]
* Mg(AlH4)2
-
-
[1978Cla1], stable above 5 MPa H2 at 410°C [1980Min]
H
Data / Grid: at.%
Fig. 1: Al-H-Mg. Isothermal section at temperatures between 140 and 170°C
Axes: at.%
20
80
MgAl2H8
AlH3
MgH2 40
60
60
40
80
20
β +γ +MgH2
(Mg)+γ +MgH2
Mg
Landolt-Börnstein New Series IV/11A3
(Mg)
20
40
γ
(αAl)+β +MgH2 60
β
80
(α Al)
Al
MSIT ®
Al–H–Mg
70
12
Fig. 2: Al-H-Mg. Hydrogen solubility in liquid Al-Mg alloys under 1 bar at 500, 700 and 800°C
10
0° C
6
0°C
70
80
(H, at.%)@10-3
8
4
C
500°
2
0 0
20
Al
80
60
40
100
Mg
Mg, at.%
10
Fig. 3: Al-H-Mg. Pressure-composition isotherms of the Mg2Al-H system
410°C
375°C
PH2 (MPa)
350°C 335°C 1
0 0
0.5
1.0
1.5
2.0
H/Mg
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–H–Ti
71
Aluminium – Hydrogen – Titanium Viktor Kuznetsov Literature Data The major works on this system has been concentrating on H in Ti-rich phases to investigate H embrittlement and related phenomena and using Al-Ti alloys for hydrogen storage. [1981Ive] mentioned Ti3Al as one of most promising candidate systems for hydrogen storage. Unfortunately the equilibrium usually was obtained only between H2 gas and metal surface, if at all, but not within the metal sublattice, which corresponds to paraequilibrium conditions. True phase equilibria were achieved and investigated very rarely and the information on them remains very limited. [1958Ber] studied by metallography the H embrittlement of Ti and alloys with 2.5, 5 and 7 mass% Al prepared between 675 and 940°C and concluded that Al increases the H solubility. Later [1971Pat] re-investigated this using alloys of iodide-purified Ti with 1, 3 and 10 at.% Al. Using resistometrical methods and direct observation of hydride formation by electron microscopy, he showed the increase of H solubility to be due to self-stresses around the hydride particles; plastic flow of the matrix causes a strong hysteresis. This hysteresis state is rather stable and the apparent equilibrium is not disturbed for several weeks from 20 to approximately 150°C. This was confirmed by [1976Che] who showed that such supersaturated solutions of H in Ti-4Al alloys do decompose, giving TiH2 after annealing for 40 d under stress conditions. [1974Sch1, 1974Sch2] investigated in great detail the solubility of H in Ti and its alloys with 5, 7 and 10 at.% Al from 800 to 900°C and described the coexisting phase configurations in three partial isothermal sections. The main impurities were up to 0.03 mass% Fe, 0.04% C and 0.4% O. Analogous work was conducted by [1981Buk] from 500 to 800°C, but the results at 800°C agree rather poorly. [1981Buk] also displayed the position of three-phase triangles -2-. The claim of the three-phase state of the products of hydridation of alloys with Al content from 7.5 up to 18.4 mass% based on metallography are corroborated to some degree by the observations which [1989Ili] made on the formation of the 2 phase in hydridated alloys with more than 7 mass% Al. [1972Gab] measured a H2 solubility in TiAl 3 at 500 and 600°C extrapolating data which were obtained for H2 pressures of 0.4 to 0.6 kbar to a H2 pressure of 1.01 bar. Only the solubilities at 1.01 bar were given. For both temperatures the solubilities were found to be 1.4 to 1.6 ml H2 per 100 g of alloy. [1977Rud] investigated the solubility and the thermodynamics of solution of H in Ti3Al from 450 to 800°C and for H2 pressures lower than 1.333 bar. The hydrogen solubilities at room temperature under hydrogen pressure of 5 MPa were measured by [2001Has, 2002Has, 2002Ito] around the composition Ti3Al, as well as the temperature at which 50 % hydrogen is desorbed in the whole interval of compositions. [1972Sch] studied the influence of the temperature on the rate of thermal decomposition of hydridation products for Ti alloys with 1.2, 3.0 and 5.9 mass% Al. [2000Sor2] studied the temperatures and details of kinetics of thermal decomposition of two hydrides, obtained by hydridation of Ti3Al under H2 pressure of 3.8 MPa at room temperature. [2002Ito] studied hydrogen adsorption isotherms for another hydridation conditions (127°C, 0.001 to 10 MPa). [1978Rud] studied the interactions of Ti3Al with H 2 at higher H2 pressures than [1977Rud] but using the same specimen and found three metastable phases analogous to hydrides of Ti. The phases exist up to 150°C and decompose at 200°C giving TiH2. Based on metallographic investigations [1975Buk] suggested the existence of a hydride distinguishable from TiH2 after slow cooling from 800°C to room temperature, the Al content being more than 4 to 5 at.%. [1981Kol] confirmed this by X-ray studies. A fragment of the diffraction pattern of the two-phase mixture of TiH2 and a new phase “TiAlHx” is given, but no structural data were extracted; the real composition of that phase is also not known. [1991Sch] obtained a Ti3AlH compound by a reaction of Ti3Al with H2 gas at pressure of 0.1 MPa and a temperature of 600°C, and determined its structure using neutron diffraction. [1999Mae] used the same technique on a product of interaction of Ti3Al with D2 gas at p = 9.2 bar and 200°C; they determined the composition of a higher
Landolt-Börnstein New Series IV/11A3
MSIT ®
72
Al–H–Ti
hydride to be Ti3AlH8-z (z 0.8) and determined its crystal structure. The latter phase was identified with fcc hydride of [1978Rud], but no bcc phase was detected under that conditions. A structural study of the reaction products of Ti3Al with H2 gas at 127°C was performed by [2002Ito], who found and investigated by X-ray diffraction and electronography both phases discovered by [1978Rud]. The metal sublattice of “bcc” hydride called as “H” hydride proved to be an orthorhombic superlattice to bcc. Two modifications of “fcc” phase of [1978Rud] called “ H1” and “ H2” were discovered; the metal sublattices of both have bcc superlattices, close to fcc. The positions of H atoms were not determined but the H2 phase was identified as Ti3AlH 8-z [1999Mae]. The decomposition of higher hydrides gives TiH2 and some Al enriched product. For alloys with 30 % Al and more [1999Mae] found amorphization under H2 treatment. The results of [1999Mae] and [2002Ito] generally confirm those of [1978Rud] and refine the structural data. [2002Ito] also suggested a possible mechanism of mutual transformation of these phases. These authors correlated relative stability of different hydride phases with the kinetics of hydrogen desorption. The H solubility in two-phase samples (Ti3Al+TiAl) from 450 to 570°C was measured by [1995Tak]. [1976Gri] investigated the solubility and thermodynamics of solution of H in liquid Al-Ti alloys up to 8.7 mass% Al between 1700 and 2100°C. The starting materials were Al(A999) and Ti sponge with main impurities of 0.04 mass% Fe, 0.01% Mn, 0.002% Si, 0.004% C, 0.04% O and 0.01% N. The specimens obtained were analyzed yielding 0.03 to 0.4 mass% O and 0.01 mass% N. To prevent contamination, the H saturation was conducted by electromagnetic levitation with subsequent quenching. H content was measured by vacuum extraction. The solubility of H in Ti3Al is theoretically analyzed using a geometrical model [1985Mro]. [1994Bel] performed investigation of H influence on ordering in the Ti3Al phase using GBW model. Ab initio calculation of electronic structure, chemical bonding and hydrogen site preferences in two modifications of Ti3Al and Ti3AlH phase was performed by [2000Sor1]. Binary Systems For the Al-H and Al-Ti binary systems the updated versions [2002Per, 2003Sch] are accepted. The H-Ti edge is believed to be correct as described by [Mas2]. Solid Phases The ternary hydride phases are stable under hydrogen pressure. For instance, at 127°C, Ti3AlH8-z is stable above 10 kPa [2002Ito], and their appearance strongly depends on the conditions of preparation. The crystallographic data for all reported phases, including metastable ternary hydrides, are given in Table 1. For the H phase only the structure of the metal sublattice is known. The H1 and H2 phases are claimed to differ only by H content [2002Ito]; no direct structural data for the former seem to exist. The H composition in H1 is not reported; it was estimated by the present author from its position in hydridation sequence after H and claims of [2002Ito] that it contains less hydrogen than H2. The identity of the latter with fcc hydrides of [1978Rud] and [1999Mae] are accepted, though the real structure may be more complex than determined by [1999Mae]. Isothermal Sections Only [1981Buk] and [1974Sch1, 1974Sch2] tried to present true phase equilibria. The latter data are preferred, mainly because in the former work the H2 pressure in declared three-phase field was not constant. Figures 1 and 2 display the sections at 800 and 900°C after [1974Sch1, 1974Sch2]; in addition these authors give an isothermal section at 850°C which is very similar to that at 800°C and not reproduced here. Thermodynamics A selection of isoactivity lines of H [1974Sch2] at 800 and 900°C is presented in Figs. 1 and 2. A Wagner expansion of the activity coefficient [1976Gri] fits the experimental data within their scatter between 1700 and 2100°C up to 10% Al: log10 ((%H)/p(H2)1/2) = 2323/T - 2.043 - (92.2/T - 0.03) (%Al) MSIT®
Landolt-Börnstein New Series IV/11A3
Al–H–Ti
73
where (%H), (%Al) are in mass%, p(H2) in bar, T in K. The interaction parameters of Al in Ti and Ti have been calculated by [1974Sch2]: in Ti: JH(Al) = dln H / dx(Al) = +5.51 in Ti: JH(Al) = dln H / dx(Al) = +6.04. The positive values of J show that Al dissolved in Ti decreases the hydrogen solubility. The result was experimentally confirmed by [1975Buk] with solid Ti, then by [1976Gri] with liquid Ti. The solubility of H in Ti3Al has been investigated at various temperatures and pressures, up to 200°C and up to 10 MPa [1978Rud]. The 150°C isotherm presents a plateau from H-Ti3AlH2 to H1-Ti3AlH 3 under 1 MPa. The hydrogen uptake goes up to Ti3AlH 4 under 10 MPa H2. The same plateau is estimated under 0.1 MPa at 100°C and under 0.01 MPa at 50°C. The hydrogen pressure at equilibrium H- H1 is given by: RTln(pH2/bar) = -47280 + 127.2T This relation agrees with measurements made later by [2002Ito] which propose a plateau at 127°C and 0.2 MPa. The solubility of H under 1 bar in TiAl has been experimentally determined between 450 and 570°C [1995Tak]. It is given by the following expressions: for Ti50Al50 c/ppm = 1.12 # 104exp(-4380/T), for Ti55Al40 c/ppm = 1.53 # 106exp(-7010/T). These alloys show endothermic uptake of hydrogen. Only the Ti47Al53 alloy takes up hydrogen exothermically. Notes on Materials Properties and Applications The use of Ti3Al for hydrogen storage is discussed from technical point of view in [1981Ive], [1995Tak]. At room temperature under 1 MPa H2 Ti3Al may absorb hydrogen up to the composition Ti3AlH5.6 (H/Me=1.4), under 5 MPa the hydride obtained is Ti3AlH6 (H/Me = 1.5). The hydrogen capacity decreases with off-stoichiometry. For instance, under 5 MPa H2, Ti0.7Al0.3 alloy absorbs hydrogen up to the composition Ti0.7Al0.3H. The desorption of hydrogen reaches 50% by heating at 600°C; it reaches 100% by heating at 800°C [2001Has]. Careful investigation has been carried out by [2002Ito] at higher temperature (127°C). The pressure composition curve of Fig. 3 shows a plateau with hysteresis. On the absorbing edge a plateau is observed at 0.2 MPa for the transition Ti3AlH2 (H/Me=0.5) to Ti3AlH4 (H/Me=1). On the desorbing edge the plateau (narrower and less well defined) is observed at 8 kPa. The position of the plateaus does not change significantly with the preparation of the samples (single crystalline, homogenized, pulverized and as arc-melted). Miscellaneous [1954Ram] suggested as a preparative method to obtain Ti hydride with low O and N content the saturation of an Al-10Ti (mass%) alloy with H2 at 1000°C. This suggestion, however, seems to contradict all other data, especially [1972Gab] who did not find any decomposition of TiAl3 with H2 up to 0.4 to 0.6 kbar of the latter. It may be correlated to some degree with [1972Sch] and [1978Rud], although an observation of [1972Sch] of the decomposition of alloys with only 1 to 6 mass% Al to give free Al (!) seems to be quite surprising. Nevertheless the possible formation of pure Al during decomposition of H2 phase was discussed by both [1999Mae] and [2002Ito], although none of authors could detect it. The suggestion of [1999Mae] on the formation of TiAl (possibly in nanocrystalline or amorphous state) in addition to TiH2 under that conditions seems to be more realistic. The substitution of Ti by Zr or Hf decreases slightly the hydrogen storage capacity of Ti3Al; the substitution of one atom of Ti in Ti3Al by one atom of Mn, Ni, Cu, V or Co decreases hydrogen storage capacity by a factor of 3. The alloys Ti2CrAl or Ti2FeAl had no hydrogen storage capacity at all [2001Ish].
Landolt-Börnstein New Series IV/11A3
MSIT ®
74
Al–H–Ti
References [1954Ram] [1958Ber] [1971Pat] [1972Gab]
[1972Sch]
[1974Sch1]
[1974Sch2]
[1975Buk]
[1976Che]
[1976Gri]
[1977Rud] [1978Rud]
[1981Buk]
[1981Ive] [1981Kol]
[1985Mro]
[1987Ere]
MSIT®
Ramamurthi, S., “Formation of Titanium Hydride in Aluminium- Titanium Alloys”, J. Sci. Ind. Research (India), 13B, 306-307 (1954) (Experimental, 3) Berger, L.W., Williams, D.H., Jaffe, R.J., “Hydrogen in Titanium-Aluminium Alloys”, Trans. Met. Soc. AIME, 212, 509-513 (1958) (Experimental, 7) Paton, N.E., Hickman, B.S., Leslie, D.H., “Behavior of Hydrogen in a Phase Ti-Al Alloys”, Metall. Trans., 2, 2791-2796 (1971) (Experimental, *, 16) Gabidullin, R.M., Shevtsov, I.N., Kolachev, B.A., Archakov, Yu.I., “Solubility of H in Al Intermetallics with Mg, Cu, Mn, Ti and Zr” (in Russian), Stroenie Svoistva i Primenenie Metall., (Publ. 1974), 188-190 (1972) (Experimental, 2) Schekhotsov, M.G., Kolomytsky, F.M., Rubtsov, A.N., “Investigation of Thermal Stability of Titanium Hydride and Hydridated Titanium Based Alloys” (in Russian), Stroenie Svoistva i Primenenie Metall., (Publ. 1974), 185-188 (1972) (Experimental, 4) Schuermann, E., Kootz, T., Preisendranz, H., Schueller, P., Kauder, G., “On the Hydrogen Solubility in the Ti-Al-H, Ti-V-H and Ti-Al-V-H in the Temperature Range 800 to 1000°C at H 2 Pressures 0.1 to 250 mbar. Part 1: Theoretical Basis and Experimental Data” (in German), Z. Metallkd., 65, 167-172 (1974) (Experimental, Thermodyn., *, 32) Schuermann, E., Kootz, T., Preisendranz, H., Schueller, P., Kauder, G., “On the Hydrogen Solubility in the Ti-Al-H, Ti-V-H and Ti-Al-V-H in the Temperature Range 800 to 1000°C at H 2 Pressures 0.1 to 250 mbar. Part 2: Thermodynamic Evaluation” (in German), Z. Metallkd., 65, 249-255 (1974) (Equi. Diagram, Thermodyn., #, *, 3) Bukhanova, A.A., Kolachev, B.A., Nazimov, O.Z., Seregina, E.V., “On the Influence of Al to H Solubility in Ti” (in Russian), Tekhnol. Legk. Splavov, (8), 48-53 (1975) (Experimental, 7) Chernetsov, V.I., Tseiger, E.N., “On the Solubility of H in Aluminium-Bearing Titanium Alloys”, Sov. J. Non-Ferrous Met., (5), 69 (1976), translated from Tsvetn. Met., (5), 67 (1976) (Experimental, 0) Grigorenko, G.M., Lakomskii, V.I., Korzhov, M.P., Tetyukhin, V.V., Konstantoniv, V.S., Kalinyuk, N.M., Gontchar, V.Ya., Solomentsev, A.N., “The Influence of Al to H Activity in Molten Ti” (in Russian), Probl. Spets. Elektrometall., (5), 88-93 (1976) (Thermodyn., Experimental, 12) Rudman, P.S., Reilly, J.J., Wiswall, R.H., “Hydrogen Absorption in Ti3Al”, Ber. Bunsen-Ges. Phys. Chem., 31, 71-80 (1977) (Experimental, 10) Rudman, P.S., Reilly, J.J., Wiswall, R.H., “The Formation of Metastable Hydrides Ti 0.75Al0.25Hx with x < 1.5”, J. Less-Common Met., 58, 231-240 (1978) (Experimental, Crys. Structure, 10) Bukhanova, A.A., Kolachev, “On the Phase Diagram of the Ti-Al-H System between 500 to 800°C” (in Russian), Fazovje Ravnovesija v Metallicheskych Splavach, Publ. Nauka, Moscow, 127-131 (1981) (Equi. Diagram, 3) Ivey, D.G., Northwood, D.O., “Metal Hydrides for Energy Storage”, Can. Metall. Quart., 20, 397-405 (1981) (Review, 40) Kolachev, B.A., Gontchar, V.Ya., Liskovitsch, V.A., “Phase Composition of the Hydrogenation Products of Titanium Alloys”, Inorg. Mater., 17, 1527-1530 (1982), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 17, 2048-2052 (1981) (Experimental, 10) Mrowietz, M., Weiss, A., “Solubility of Hydrogen in Titanium Alloys. II. Blocking Models and Hole Size Considerations”, Ber. Bunsen-Ges. Phys. Chem., 89, 362-371 (1985) (Thermodyn., Theory, 82) Eremenko, V.N., Tretyachenko, L.A., “Physico-Chemical Properties of Titanium”, in “Ternary Systems of Titanium with Transition Metals of IV-VI Groups” (in Russian), Naukova Dumka, Kiev, 5-6 (1987) (Equi. Diagram, Crys. Structure, Review, 14) Landolt-Börnstein New Series IV/11A3
Al–H–Ti [1989Ili]
[1990Sch]
[1991Sch]
[1992Kat]
[1994Bel]
[1995Tak]
[1999Mae]
[2000Sor1]
[2000Sor2]
[2001Bra]
[2001Has] [2001Ish]
[2002Ito]
[2002Has]
[2002Per]
[2003Sch]
Landolt-Börnstein New Series IV/11A3
75
Il'in, A.A., Mamonov, A.M., Mikhailov, Yu.V., “The Phase Diagrams of H Alloyed Ti Alloys” (in Russian), Abstr. 5th All-Union Conf. Phase Diagrams of Metallic Systems, 162 (1989) (Equi. Diagram, Abstract, 0) Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”, Z. Metallkd., 81, 389-396 (1990) (Crys. Structure, Equi. Diagram, Experimental, Review, #, 33) Schwartz, D.S., Yelon, W.B., Berliner R.B., Lederich, R.J., Sastry, S.M., “A Novel Hydride Phase in Hydrogen Charged Ti3Al”, Acta Met. Mater., 39, 2799-2803 (1991) (Crys. Structure, Experimental, *, 8) 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, Equi. Diagram, Thermodyn., #, *, 51) Belov, S.P., Il'in, A.A., Mamonov, A.M., Aleksandrova, A.V., “Theoretical Analysis of Ordering in Ti3Al-Base Alloys. II. Effect of Hydrogen on Stability of Ti3Al Intermetallic Compound”, Russ. Metall., (2), 52-55 (1994), translated from Izv. Ross. Akad. Nauk. Met., (2), 76-78 (1994) (Crys. Structure, Theory, 13) Takasaki, A., Furuya, Y., Ojima, K., Taneda, Y., “Hydrogen Solubility of Two-Phase (Ti3Al+TiAl) Titanium Aluminides”, Scr. Metall. Mater., 32, 1759-1764 (1995) (Phys. Prop., Experimental, 12) Maeland, A.J., Hauback, B., Fjellvag, H., Sorby, M., “The Structure of Hydride Phases in the Ti 3Al/H System”, Int. J. Hydrogen Energy, 24, 163-168 (1999) (Crys. Structure, Experimental, *, 12) Sornadurai, D., Panigrahi, B., Ramani, “Electronic Structure, Hydrogen Site Occupation and Phase Stability of Ti3Al upon Hydrogenation”, J. Alloys Compd., 305, 35-42 (2000) (Crys. Structure, Theory, 22) Sornadurai, D., Panigrahi, B.K., Shashikala, K., Raj, P., Sastry, V.S., Ramani, “X-Ray Diffraction and Differential Scanning Calorimetry Investigations on High-Pressure Hydrogen Gas Charged Ti3Al”, J. Alloys Compd., 312, 251-256 (2000) (Crys. Structure, Kinetics, *, 10) 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, Equi. Diagram, Experimental, #, *, 34) Hashi, K., Ishikawa, K., Aoki, K., “Hydrogen Absorption and Desorption in Ti-Al Alloys”, Met. Mater. Int., 7(2), 175-179 (2001) (Equi. Diagram, Experimental, 8) 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) (Crys. Structure, Kinetics, Experimental, *, 8) Ito, K., Okabe, Y., Zhang, L.T., Yamaguchi, M., “Reversible Hydrogen Absorption/Desorbtion and Related Phase Transformations in a Ti3Al Alloy with Stoichiometry Composition”, Acta Mater., 50, 4901-4912 (2002) (Equi. Diagram, Experimental, *, 18) Hashi, K., Ishikawa, K., Syzuki, K., Aoki, K., “Hydrogen Absorption and Desorption in the Binary Ti-Al System”, J. Alloys Compd., 330/332, 547-550 (2002) (Equi. Diagram, Experimental, 11) Perrot, P., “Al-H (Aluminium-Hydrogen)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14832.1.20, (2002) (Equi. Diagram, Crys. Structure, Assessment, 21) Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 86)
MSIT ®
Al–H–Ti
76 Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 (Ti) 1670 - 882 (Ti) < 882 TiAl3 < 1387
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cI2 Im3m W hP2 P63/mmc Mg tI8 I4/mmm TiAl3
Lattice Parameters Comments/References [pm] a = 404.88
pure Al [V-C]
a = 330.65 a = 328.4
pure Ti at 900°C [V-C] at room temperature, extr. from solid solution [1987Ere] [V-C]
a = 295.2 c = 498.9 a = 384.88 c = 859.82
“Ti2Al5” 1416 - 990 tetragonal superstructure of AuCu-type [2001Bra]
tP28 P4/mmm “Ti2Al5” Ti5Al11 1416 - 1206
tI16 I4/mmm ZrAl3
TiAl2(h) 1433 - 1214
oC12 Cmmm ZrGa2 tI24 I41/amd HfGa2 oP4
TiAl2(r) < 1216 Ti1-xAl1+x ~1445 - 1424 TiAl < 1460
MSIT®
tP4 P4/mmm AuCu(I)
a* = 395.3 c* = 410.4 a* = 391.8 c* = 415.4 a = 390.53 c = 2919.63
a = 392.30 to 393.81 c = 1653.49 to 1649.69 a = 1208.84 b = 394.61 c = 402.95 a = 396.7 c = 2429.68 a = 402.62 b = 396.17 c = 402.62 a = 398.69 c = 405.39
[1990Sch]
chosen stoichiometry [1992Kat] summarizing several phases: Ti5Al11 stable range 1416- 995°C [2001Bra] 66 to 71 at.% Al at 1300°C [2001Bra] (including the stoichiometry Ti2Al5!); [1990Sch] claimed: 68.5 to 70.9 at.% Al and range 1416 -1206°C; at 66 at.% Al [2001Bra] * AuCu subcell only at 71 at.% Al [2001Bra] * AuCu subcell only “Ti2Al5” ~1215 - 985°C [1990Sch]; included in homogeneity region of Ti5Al11 [2001Bra] 29.1 to 31.5 at.% Ti [1990Sch]
33 to 34 at.% Ti [1990Sch]
[1990Sch]
at x = 0.28 [1990Sch] at 38.5 to 52 at.% Ti [1990Sch] at 38.5 at.% Ti, 1000°C
Landolt-Börnstein New Series IV/11A3
Al–H–Ti Phase/ Temperature Range [°C] Ti3Al < 1180
TiHx > 315 JTiHx < 315 * Ti3AlH
* Ti0.75 Al0.25H x
Pearson Symbol/ Space Group/ Prototype hP8 P63/mmc Ni3Sn cF12 Fm3m CaF2 tI6 I4/mmm ThH2 cP5 Pm3m CaTiO3 hP?
Lattice Parameters Comments/References [pm] a = 580.6 c = 465.5 a = 574.6 c = 462.4 a = 445.4
at 78 at.% Ti [L-B]
a = 320.2 c = 427.9
x = 1.72 to 2.0 [Mas, V-C]
a = 408.79
[1991Sch] The parameter is given by [1999Mae] for Ti3AlD metastable, x < 0.2 [1978Rud] decomp. at 200°C, form. at 50 to 150°C at x 0
a = 289 c = 466 * Ti0.75 Al0.25H x
cI? a = 328
* Ti0.75 Al0.25H x
cF? a = 435
* H, Ti3AlH2
oP8 a)
* H1, Ti3AlHx
t??
a 310 a = 390 c = 313
* H2, Ti3AlH8-z
tP12
a = 439.77
* “TiAlyHx”
-
-
a)
77
at 62 at.% Ti [L-B] x = 1.05 to 2.0 [Mas, V-C]
metastable, 0.4 < x < 0.5 (in-reactor state) [1978Rud] decomp. at 200°C, form. at 50 to 150°C at x = 0.35 (estimated) metastable, x > 1.5 [1978Rud] decomp. at 200°C, form. at 50 to 150°C at x = 1.6 (two-phase sample, estimated comp.) [2002Ito]; two-phase sample with H/Me = 0.55; 2×2×1 superstructure to bcc; 2c/a ratio close to 1, varies depending on sample probably identical to cI? phase of [1978Rud] approx. Value for bcc sublattice [2002Ito]; two-phase sample with H/Me = 0.55 (same as previous); x probably between 2 and 7.8 [1999Mae]; sample of gross composition Ti3AlH5.9 contained also 6-7% of TiAl and Ti3AlH; z 0.8; composition of metal sublattice is Ti3(Al0.25Ti 0.75) after [2002Ito] may be bct with c/a ratio close to fcc [1981Kol], the composition of both Al and H is not known
Only metal atoms are counted for Pearson symbol
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–H–Ti
78
Ti Al H
Fig. 1: Al-H-Ti. Partial isothermal section with superimposed (dashed) isoactivity lines (aH) at 800°C. The numbers given are aH=(pH2(bar)/0.981)1/2)
70.00 0.00 30.00
Data / Grid: at.% Axes: at.%
08 0.5
80
20
27 0.3
(β Ti) (αTi)+(β Ti)
90
10
(α Ti)
0.236
0.181
0.073
0.127
0.018 10
Ti
20
Ti Al H
Fig. 2: Al-H-Ti. Partial isothermal section with superimposed (dashed) isoactivity lines (aH) at 900°C. The numbers given are aH=(pH2(bar)/0.981)1/2)
70.00 0.00 30.00
Ti Al H
70.00 30.00 0.00
Ti Al H
70.00 30.00 0.00
Data / Grid: at.% Axes: at.%
26 0.7
80
20
08 0.5
90
0.1 63
0.3 27 (β Ti)
10
0.2 18
0.10 9 (αTi) 0.073 0.018
Ti
MSIT®
10
20
Landolt-Börnstein New Series IV/11A3
Al–H–Ti
79
Hydrogen content (H/Me) 0.8
0.6
1.2
1.0
10
Fig. 3: Al-H-Ti. Absorbtion (A) and desorption (B) isotherms at 127°C for Ti3Al
Pressure (MPa)
1
0.1
(A) 0.01
(B)
0.003 1.0
1.5
2.0
2.5
3.0
Hydrogen content (mass%)
Landolt-Börnstein New Series IV/11A3
MSIT ®
80
Al–Hf–Ni
Aluminium – Hafnium – Nickel Gautam Ghosh Literature Data [1969Mar1] was the first to report the isothermal section of the entire system at 800°C. They prepared about 100 ternary alloys in an arc furnace under Ar atmosphere using the elemental metals Al (99.998 mass%), iodide Hf (99.95 mass%) and electrolytic Ni (99.9 mass%). The alloys were annealed at 800°C for 830 h in evacuated silica tubes followed by quenching into cold water. Phase analysis was performed by microstructural observation and X-ray diffraction techniques. [1981Nas] reported the partial isothermal sections of the Ni corner at 1200 and 1000°C. They prepared 12 ternary alloys containing up to 35 at.% Al and 23 at.% Hf. The alloys were prepared from 99.99 mass% Al, Hf containing about 3 at.% Zr and other impurities of about 0.38 mass%, and 99.99 mass% Ni. The alloy buttons were prepared in an arc furnace under Ar atmosphere. They were placed in alumina crucibles, sealed in silica tubes partially filled with Ar and were homogenized at 1200 and 1000°C for 168 h followed by quenching into water. Phase analysis was carried out by optical microscopy, X-ray diffraction and electron probe microanalysis. [1981Bal] investigated microstructure of two Ni rich ternary alloys, both as-cast and annealed conditions. These results were reviewed by [1991Lee, 1993Gho]. Brief reviews of phase equilibria were presented by [1977Abr, 1990Kum]. Recently, Miura et al. [1999Miu] investigated the solid-liquid phase equilibria of Ni-rich ternary alloys using DTA, XRD and SEM-WDS analysis. They prepared ternary alloys using 99.99 mass% Al, 99.95 mass% Ni, and 95 mass% Hf. [1991Mis] determined the solvus boundary of (Ni) using DTA and SEMEDX analysis. Other recent investigations of the ternary system involve rapid solidification [2002Lou], and very limited thermodynamic measurements [1992Alb]. Binary Systems The Al-Ni binary phase diagram is accepted from [2003Sal], and the Al-Hf binary phase diagram is accepted from [2003Sch]. Recently, Miura et al. [1999Miu, 2001Miu] have determined the liquidus of Ni rich alloys containing up to 13 at.% Al. Unlike Hilpert et al. [1987Hil], Miura et al. [2001Miu] observed a maximum (1466°C) in the liquidus at about 2 at.% Al. Except for [2001Miu], this feature has not been considered in the CALPHAD modeling of the Al-Ni phase diagram [2003Sal]. Unlike [1998Mur], Schuster [2003Sch] did not consider Hf2Al phase in the assessment of Al-Hf equilibrium diagram. This phase was first reported by [1961Now] but subsequent investigations failed to confirm. It is believed that Hf2Al and HfAl3(TiAl 3) might have been stabilized by silicon, and they are not a Al-Hf equilibrium phase [1962Poe1, 1962Poe2, 1964Rie]. The Hf-Ni binary phase diagram is accepted from [1983Nas]. Solid Phases The data of [1981Nas] suggest that the lattice parameter of (Ni) increases more rapidly in the ternary regime than in the binary solid solutions [1985Mis]. In the Hf-Ni system, the rate of increase in the lattice parameter, da/dc, is reported to be 1.0 pm/at.% Hf [1984Och2, 1985Mis]. Figure 1 shows the solubility isotherms of (Ni) [1991Mis]. Ni3Al is reported to dissolve about 8.5 at.% Hf at 1200°C [1981Nas, 1985Mis] and 7 at.% Hf at 1000°C [1983Och], 8 at.% at 1000°C [1981Nas]. On the other hand, [2002Lou] reported a maximum solubility of 11 at.% Hf in Ni3Al in a rapidly solidified Ni74Al15Hf11 alloy with lattice parameter of a = 364.0 pm, even though the alloy contained another metastable cubic phase. Substitution of Al by Hf causes a linear increase in the lattice parameter with increasing Hf content [1984Och1, 1984Och2, 1985Mis]. The rate of increase in the lattice parameter of Ni3Al, da/dc, is reported to be 0.73 pm/at.% Hf [1985Mis]. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Hf–Ni
81
The maximum solid solubility of Hf in NiAl is reported to be about 5 at.% at 1350°C [1990Tak]. Lattice parameters of Ni3Al and NiAl as a function of alloy composition and heat treatment were reported by [1981Nas]. [1981Nas] reported that Hf2Ni7 can dissolve up to about 14 at.% Al with a width of 1.5 at.% Hf at 1200 and 1000°C. On the other hand, [1969Mar1] found that none of the Hf-Ni binary compounds, including Hf2Ni7, can dissolve more than 1 at.% Al at 800°C. Lattice parameters of HfNi3, HfNi5 and Hf 2Ni7 phases as a function of alloy composition and heat treatment were also determined by [1981Nas]. At least ten ternary phases have been reported in this system, of which nine were first reported by Markiv and co-workers [1964Mar, 1966Mar, 1969Mar1, 1969Mar2]. The Hf6Ni8Al15 phase was first reported by [1966Gan1, 1966Gan2] and subsequently confirmed by [1969Mar1]. The ternary phase 2 (Hf10Ni19Al) was reported to be stable above 1000°C, but was not observed by [1981Nas] in the 1200 and 1000°C isothermal sections. This phase was suggested to be an extension of HfNi2 into the ternary region [1972Pet], but it has been disproved [1979Bse]. Also, there is experimental evidence [1979Bse, 1981Nas] suggesting that such a structure is not an equilibrium phase, but most probably stabilized by silica. Incidentally, in Markiv's [1969Mar1] experiment the specimens were in direct contact with silica tubes, whereas [1981Nas] kept their specimens in alumina crucibles during annealing treatments. The structures of the Hf5Ni4Al phase [1969Mar2] and the Hf4Ni16Al5 (3 or L phase) were not determined [1969Mar1]. The latter phase was reported to be present in the isothermal section at 800°C [1969Mar1], but was not observed in the isothermal sections at 1200 and 1000°C [1981Nas]. Accordingly, it has been suggested that the 3 phase forms by a solid state reaction between 1000 and 800°C [1981Nas]. The ternary phase HfNi2Al has been predicted to form by an invariant transition type reaction [1981Nas]. According to [1968Dwi], the structure of the HfNiAl phase can be better described by introducing a slight variation in stacking sequence and by doubling the c-parameter. The details of the crystal structures and lattice parameters of all the solid phases are listed in Table 1. Pseudobinary Systems The section NiAl-HfNi2Al is established as quasibinary using X-ray analysis, metallography and by determining the melting temperatures, but only part of this section up to 30 at.% Hf has been reported [1990Tak]. As shown in Fig. 2, a pseudobinary eutectic reaction LNiAl+HfNi2Al takes place at 1350°C and 15 at.% Hf. Further experiments are necessary to confirm this phase diagram. [1981Bal] observed a eutectic microstructure of NiAl and (Ni,Al)7Hf2 phases embedded in Ni3Al matrix in an as-cast alloy of Ni-20Al-7.5Hf (at.%). This result suggests the possibility of the existence of a pseudobinary eutectic between NiAl and Ni7Hf2. To corroborate this interpretation further experiments are needed. Invariant Equilibria Two transition invariant reactions have been reported [1981Nas] to take place during solidification of the Ni rich alloys: L+Ni3Al(Ni)+Hf2Ni7 (U1) and L+Hf2Ni7(Ni)+HfNi5 (U2). However, the temperatures of occurrence of the above invariant reactions were not reported, but estimated to be between 1275 and 1200°C. Based on the observation of equilibria in the region (Ni)-Ni3Al-Hf2Ni7 between 1200 and 1000°C, [1981Nas] predicted the presence of an invariant U type reaction (Ni)+Hf2Ni7Ni3Al+HfNi5 at some temperature between 1200 and 1000°C. In the same temperature range, [1981Nas] also predicted the possibility of another invariant U type reaction Hf2Ni7+Hf3Ni7HfNi3(r)+HfNi2Al. [1991Lee] reported a speculative invariant reaction scheme for the Ni rich portion of the ternary system. In addition to above mentioned four U type invariant reactions, their speculative reaction scheme includes eight more U type invariant reactions involving the liquid phase. Liquidus Surface Addition of Al in Ni rich Hf-Ni alloys or addition of Hf in Ni rich Al-Ni alloys decrease the liquidus temperature [1999Miu]. Figure 3 shows the probable liquidus surface of the Ni corner [1981Nas]. It is based Landolt-Börnstein New Series IV/11A3
MSIT ®
82
Al–Hf–Ni
on the observation of as-cast microstructures of Ni-(2.5 to 35)Al-(5 to 25)Hf (at.%) alloys. This is in substantial disagreement with the calculated liquidus surface by Kaufman et al. [1974Kau, 1975Kau]. Also, in their calculation [1974Kau, 1975Kau] assumed that the 3 phase (Hf4Ni16 Al5) melts congruently which is not supported by the results of [1981Nas]. Additionally, the calculated liquidus temperatures [1974Kau, 1975Kau] were substantially lower than the measured solidus temperatures [1981Nas]. In other words, the thermodynamic parameters derived by [1974Kau, 1975Kau] certainly overestimate the stability of the liquid phase. Nonetheless, combining the calculated liquidus of [1975Kau] and limited experimental data of [1981Nas], Lee and Nash [1991Lee] proposed a tentative liquidus surface up to 40 at.% Al and 50 at.% Hf. Isothermal Sections Figures 4 and 5 show the partial isothermal sections of the Ni corner at 1200°C [1981Nas, 1985Nas] and 1000°C [1981Nas], respectively. It should be mentioned that the homogeneity ranges of binary Ni3Al at 1200 and 1000°C as reported by [1981Nas, 1985Nas] were considerably higher than those given by the presently accepted binary phase diagram [Mas, 1987Hil, 1988Bre]. Figure 6 shows the isothermal section at 800°C [1969Mar1]. The three-phase fields (Ni)+Ni3Al+HfNi5 and Ni3Al+HfNi5+Hf 2Ni7, as reported in the 800°C isothermal section [1969Mar1], were also found to be present in the 1000°C isothermal section [1981Nas]. These three-phase fields result from an invariant transition type reaction (Ni)+Hf2Ni7Ni3Al+HfNi5 [1981Nas]. However, the calculated isothermal section at 800°C [1974Kau, 1975Kau] showed the presence of a (Ni)+HfNi5+Hf2Ni7 three-phase field, and thus does not take into account the above transition type reaction [1981Nas]. In Figs. 4 to 6, minor adjustments have been made in order to comply with the accepted binary phase diagrams. Since Hf2Al phase is not considered to be an equilibrium phase, previously reported three-phase fields (Hf)+Hf2Al+-5 and Hf2Al+Hf3Al2+-5 in the isothermal section at 800°C [1969Mar1] have been replaced by (Hf)+Hf3Al2+-5. Computer calculated isothermal sections, in the range of Ni-50 at.% (Hf+Al), at 1423 and 1323°C [1974Kau, 1975Kau], at 1223, 1123, 1023°C [1974Kau, 1975Kau, 1976Kau] and at 800°C [1974Kau, 1975Kau] have also been reported. Thermodynamics Experimental thermodynamic data of ternary alloys is very limited. [1992Alb] determined the activity of Hf and Al in (Ni3Al)1-xHfx and Ni0.75Al0.25-xHfx alloys in the temperature range of 1088 and 1407°C. Their data indicate the substitution of Hf for Al in Ni3Al. In fact, thermal conductivity measurement of Ni3(Al,Hf) also corroborate this behavior [2001Ter]. [1999Dar] measured the low-temperature (3.2 to 10.3 K) specific heat of HfNi2Al (-2) using an adiabatic calorimeter, and analyzed the specific heat data in terms of electronic, Debye lattice and Einstein models. The analysis of experimental data yields the Debye temperature D=15°C. They also calculated the electronic structure by tight-binding linearized muffin-tin orbital (TB-LMTO) method. Their results underscore the importance of electron-phonon coupling on the phase stability. Kaufman and Nesor [1974Kau, 1975Kau, 1976Kau] have performed CALPHAD modeling of the ternary system, and calculated several isothermal sections. Notes on Materials Properties and Applications The constitutional equilibria of this ternary system is very important for developing creep resistant hightemperature alloys. [1991Miu] studied the creep behavior of Ni-23.5Al-2Hf (at.%) alloy single crystals oriented close to [001] direction at 850, 900 and 950°C under compressive loads. They observed power-law creep behavior with an exponent of 3.89 and an activation energy of 360 kJ#mol. Hf is a good solid solution strengthener of Ni3Al. Theoretical calculations show that the strengthening effect is related to both the site occupancy and local segregation of Hf at antiphase boundaries [1991Wu]. The microstructure and mechanical properties of melt-spun and bulk Hf1Co9Ni61 Al29 specimens were compared to Al-Co-Ni samples [1990Pan].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Hf–Ni
83
Miscellaneous The solidus temperatures [1981Nas] of some ternary alloys are listed in Table 2. [1997Nag] found that in the presence of boron, the solubility of Al in Ni3Al is increased while the solubility of Ni in NiAl is decreased by about 1.25 at.% at 1130°C. These results suggest that it is easier for Al and Hf to occupy the Ni sites, and it was rationalized in terms of occupancy of interstitial sites by boron atoms. [1991Sas] studied the microstructure of arc-melted (NiAl)99.5Hf0.5 alloy, and did not find any evidence of grain refining effect. Even though they observed the presence of precipitates, the absence of grain refining effect was attributed to the solid-state precipitation. [2002Lou] carried out rapid solidification of Ni78 Al12.5Hf9.5, Ni74Al15Hf11 and Hf 20Ni66Al14 alloys. In the former two alloys they observed a hitherto unknown body-centered cubic phase with lattice parameter a = 220 pm, while the latter alloy has an amorphous structure. It is uncertain if this cubic phase is indeed -4 which also has similar lattice parameter but with face-centered symmetry. This point was not discussed by [2002Lou]. Calorimetric study shows that the crystallization temperature of the amorphous alloy is about 577°C at a heating rate of 1.33 °C/s. Other aspects of crystallization behavior of the rapidly solidified alloys have been discussed by [2002Lou]. References [1961Now] [1962Poe1]
[1962Poe2]
[1964Mar]
[1964Rie]
[1966Gan1] [1966Gan2] [1966Mar]
[1967Kri]
[1968Dwi]
[1969Mar1]
[1969Mar2]
Landolt-Börnstein New Series IV/11A3
Nowotny, H., Schob, O., Benesovsky, F. “The Crystal Structure of Zr2Al and Hf2Al” (in German), Monatsh. Chem., 92, 1300-1303 (1961) (Crys. Structure, Experimental, 10) Poetschke, M., Schubert, K., “On the Constitution of Some Systems Homologous and Quasihomologous to T4 - B3. Part I” (in German), Z. Metallkd., 53, 474-488 (1962) (Experimental, *, 18) Poetschke, M., Schubert, K., On the Constitution of Some Systems Homologous and Quasihomologous to T4 - B 3. Part II” (in German), Z. Metallkd., 53, 548-561 (1962) (Crys. Structure, Experimental, Equi. Diagram, *, 45) Markiv, V.Ya., Voroshilov, Yu.V., Kripyakevich, P.I., Cherkashin, E.E., “New Compounds of MnCu 2Al and MgZn 2 Types Containing Al and Ga”, Sov. Phys.-Crystallogr., 9, 619-620 (1965), transl. from Kristallografiya, 9, 737-738 (1964) (Crys. Structure, Experimental, 4) Rieger, W., Nowotny, H., Benesovsky, F. “Investigations in Systems Transition Metal (T)Boron-Aluminium” (in German), Monatsh. Chem., 95, 1417-1423 (1964) (Crys. Structure, Experimental, 11) Canglberger, E., Nowotny, H., Benesovsky, F., “On Some New G-Phases” (in German), Monatsh. Chem., 97, 219-220 (1966) (Crys. Structure, Experimental, 3) Ganglberger, E., Nowotny, H., Benesovsky, F., “New G-Phases” (in German), Monatsh. Chem., 97, 829-832 (1966) (Crys. Structure, Experimental, 4) Markiv, V.Ya., Kripyakevich, P.I., “Compounds of the Type R(X', X'') 2 in Systems with R = Ti, Zr, Hf; X' = Fe, Co, Ni, Cu; and X'' = Al or Ga and Their Crystal Structure”, Sov. Phys.-Crystallogr., 11, 733-738 (1967), translated from Kristallografiya, 11, 859-865 (1966) (Crys. Structure, Experimental, 25) Kripyakevich, P.I., Markiv, V.Ya., Mel´nik, Ya.V., “Crystal Structure of Zr-Ni-Al, Zr-CuGa and Analogous Compounds” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A8), 750753 (1967) (Crys. Structure, Experimental, 9) Dwight, A.E., Mueller, M.H., Conner, R.A., Downey, J.W., Knott, H., “Ternary Compounds with the Fe2P-Type Structure”, Trans. Metall. Soc. AIME, 242, 2075-2080 (1968) (Crys. Structure, Experimental, 14) Markiv, V.Ya., Burnashova, V.V., “The Hf-Ni-Al System”, Russ. Metall. (Engl. Transl.), (6), 113-115 (1969), translated from Izv. Akad. Nauk SSSR, Met., (6), 181-182 (1969) (Equi. Diagram, Experimental, #, *, 17) Markiv, V.Ya., Burnashova, V.V., “New Ternary Compounds in the (Sc, Ti, Zr, Hf)-(V, Cr, Mn, Fe, Co, Ni, Cu)-(Al, Ga) Systems” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A5), 463-464 (1969) (Crys. Structure, Experimental, 12) MSIT ®
84 [1969Tes] [1972Pet]
[1974Fer]
[1974Kau]
[1975Kau] [1976Kau]
[1977Abr] [1979Bse] [1981Bal]
[1981Fer]
[1981Nas] [1983Nas] [1983Och] [1984Och1] [1984Och2]
[1985Mis]
[1985Nas]
[1987Hil]
[1988Bre]
[1990Kum]
MSIT®
Al–Hf–Ni Teslyuk, M.Yu., Intermetallic Compounds with Structure of Laves Phases (in Russian), Moscow, Nauka, 1969, 1-138 (1969) (Crys. Structure, Equi. Diagram, Review) Pet´kov, V.V., Markiv, V.Ya. Gorsky, V.V., “Compound with the MgCu2-Type of Structure in Alloys of Ni, Zr and Hf” (in Russian), Izv. Akad. Nauk SSSR, Met., (2), 188-192 (1972) (Crys. Structure, Experimental, 10) Ferro, R., Marazza, R., Rambaldi, G., “Equi-Atomic Ternary Phases in the Alloys of the Rare Earths with In and Ni and Pd”, Z. Metallkd., 65, 37-39 (1974) (Crys. Structure, Experimental, 2) Kaufman, L., Nesor, H., “Computer Calculated Phase Diagrams for the Ni-W-Al, Ni-Al-Hf, Ni-Cr-Hf and Co(Cr,Ni)-Ta-C Systems”, NASA Contract No NAS3-17304, National Aeronautics and Space Administration, Washington, D.C. 20546, 1-58 (1974) (Equi. Diagram, Thermodyn., Theory, 28) Kaufman, L., Nesor, H., “Calculation of the Ni-W-Al, Ni-Al-Hf, Ni-Cr-Hf Systems”, Can. Metall. Quart., 14, 221-232 (1975) (Equi. Diagram, Thermodyn., Theory, 22) Kaufman, L., Nesor, H., “Application of Computer Techniques of Prediction of Metastable Transitions in Metallic Systems”, Mater. Sci. Eng., 23, 119-123 (1976) (Equi. Diagram, Theory, 13) Abrikosov, N.Kh., “Phase Diagrams of Al and Mg Alloy Systems” in “Phase Diagrams of Al and Mg Alloy Systems”, Nauka, Moscow, 22-25 (1977) (Crys. Structure, Review, 5) Bsenko, L., “The Hf-Ni and Zr-Ni Systems in the Region 65-80 at.% Ni”, J. Less-Common Met., 63, 171-179 (1979) (Equi. Diagram, Experimental, 13) Baldan, A. and West, D.R.F., “Structural Features of Certain Ni-Al-Ta and Ni-Al-Hf Alloys Containing the ´ and -Phases”, J. Mater. Sci., 16, 24-34 (1981) (Crys. Structure, Experimental, 28) Ferro, R., Marazza, R., “Crystal Structure and Density Data” in “Hafnium: Physicochemical Properties of its Compounds and Alloys”, Atomic Energy Review, Special Issue No.8., K.L. Komarek, Ed., IAEA, Vienna, (8), 121-250 (1981) (Crys. Structure, Review, 645) Nash, P., West, D.R.F., “Phase Equilibria in Ni-Rich Region of the Ni-Al-Hf System”, Met. Sci., 15, 347-352 (1981) (Equi. Diagram, Experimental, #, *, 20) Nash, P., Nash, A., “The Hf-Ni (Hafnium-Nickel) System”, Bull. Alloy Phase Diagrams, 4, 250-253 (1983) (Equi. Diagram, Review, #, *, 23) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data of Ni3Al with Ternary Additions”, Bull. P. M. E., (52), 1-17 (1983) (Equi. Diagram, Experimental, Review, 39) Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32, 289-298 (1984) (Equi. Diagram, Experimental, 90) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions”, Bull. P. M. E., (53), 15-28 (1984) (Crys. Structure, Experimental, 66) Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni( ), Ni3Al( ') and Ni3Ga( ') 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” in “HighTemperature Ordered Intermetallic Alloys”, Koch, C.C., Liu, C.T., Stoloff, N.S., (Eds.), Mat. Res. Soc., Pittsburgh, PA, 423-427 (1985) (Equi. Diagram, Review, #, *, 15) Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremer, F.J., Nickel, H., “Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch., 42a, 1327-1332 (1987) (Equi. Diagram, Experimental, *, 17) Bremar, F.J., Beyss, M., Karthaus, E., Hellwig, A., Schober, T., Welter, J.-M., Wenzl, H., “Experimental Analysis of the Ni-Al Phase Diagram”, J. Cryst. Growth, 87, 185-192 (1988) (Equi. Diagram, Experimental, *, 16) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X=V, Cr,Mn,Fe,Co,Ni,Cu,Zn)”, Int. Mat. Rev., 35, 293-327 (1990) (Equi. Diagram, Review, 158) Landolt-Börnstein New Series IV/11A3
Al–Hf–Ni [1990Pan]
[1990Tak] [1991Lee] [1991Sas]
[1991Mis]
[1991Miu]
[1991Wu]
[1992Alb]
[1993Gho]
[1997Nag] [1998Mur] [1999Dar]
[1999Miu]
[2001Miu]
[2001Ter]
[2002Lou]
[2003Sal]
[2003Sch]
Landolt-Börnstein New Series IV/11A3
85
Pank, D.R., Nathal, M.V., Koss, D.A., “Microstructure and Mechanical Properties of Multiphase NiAl-Based Alloys”, J. Mater. Res., 5, 942-949 (1990) (Experimental, Mechan. Prop., 18) Takeyama, M., Liu, C.T., “Microstructure and Mechanical Properties of NiAl-Ni2AlHf Alloys”, J. Mater. Res., 5, 1189-1196 (1990) (Equi. Diagram, Experimental, #, *, 22) Lee, K.J., Nash, P., “The Al-Hf-Ni System”, J. Phase Equilib., 12, 94-104 (1991) (Equi. Diagram, Crys. Structure, Review, #, 16) 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, Equi. Diagram, Experimental, Mechan. Prop., 10) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Assessment, Equi. Diagram, Experimental, 5) Miura, S., Hayashi, T., Takekawa, M., Mishima, Y., Suzuki, T., “The Compression Creep Behavior of Ni3Al-X Single Crystals”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 623-628 (1991) (Experimental, Phys. Prop., 9) 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) (Calculation, 22) Albers, M., Baba, M.S., Kath, D., Miller, M., Hilper, K., “Chemical Activities in the Solid Solution of Hf in Ni3Al”, Ber. Bunsen-Ges. Phys. Chem., 96(11), 1663-1668 (1992) (Equi. Diagram, Experimental, Thermodyn., 25) Ghosh, G., “Aluminium-Hafnium-Nickel”, in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.12751.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 30) Nagarajan, R.R., Jena, A.K., Ray, R.K., “Phase Equilibria in the ´-Rich Region of the NiAl-Hf System”, Z. Metallkd., 88(1), 87-90 (1997) (Equi. Diagram, Experimental, 16) Murray J.L., McAlister A.J., Kahan D.J., “The Al-Hf (Aluminium-Hafnium) System”, J. Phase Equilib., 19, 376-379 (1998) (Assessment, Crys. Structure, Equi. Diagram, *,14) 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) Miura, S., Hong, Y.-M., Suzuki, T., Mishima, Y., “Liquidus and Solidus Temperatures of Ni-Solid Solution in Ni-Al-X (X: Ti, Zr, and Hf) Ternary Systems”, J. Phase Equilib., 20(3), 193-198 (1999) (Equi. Diagram, Experimental, 11) Miura, S., Unno, H., Yamazaki, T., Takizawa, S., Mohri. T., “Reinvestigation of Ni-Solid Solution/Liquid Equilibria in Ni-Al Binary and Ni-Al-Zr Ternary Systems”, J. Phase Equilib., 22, 457-462 (2001) (Equi. Diagram, Experimental, #, *, 9) 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) Louzguine, D.V., Inoue, A., “Structure and Transformation Behaviour of Rapidly Solidified Ni-Al-Hf Alloys”, J. Alloys Compd., 340, 151-156 (2002) (Crys. Structure, Equi. Diagram, Experimental, 9) Saltykov, P., Cornish, L., Cacciamani, G., “Al-Ni (Aluminum-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 155) Schuster, J.C, “Al-Hf (Aluminium-Hafnium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 39) MSIT ®
Al–Hf–Ni
86 Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) 660.452 (Hf) 2231 - 1743 (Hf) 1743
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cI2 Im3m W hP2 P63/mmc Mg
(Ni) 1455
cF4 Fm3m Cu
Hf2Al < 1160
tI12 I4/mcm CuAl2 tP20 P42/mnm Zr3Al2 hP7 P6/mmm Zr4Al3 oC8 Cmcm CrB oF40 Fdd2 Zr2Al3 hP12 P63/mmc MgZn2 tI8 I4/mmm TiAl3 tI16 I4/mmm ZrAl3 oP16 Pnma NiAl3 hP5 P3m1 Ni2Al3
Hf3Al2 1590 25 Hf4Al3 1200 HfAl 1800 Hf2Al3 1640 25 HfAl2 1650 25 HfAl3(h) 1590 - 700 HfAl3(r) 700 NiAl3 854 Ni2Al3 1138
MSIT®
Lattice Parameters Comments/References [pm] a = 404.88 a = 404.96
[V-C], pure Al at 24°C [Mas2], Al at 25°C
a = 361.5 a = 361.0
[V-C], [Mass2] [2003Sch] dissolves up to 34 at.% Al at 1450°C [V-C], pure Hf at 25°C [Mas2]
a = 319.8 c = 506.1
[2003Sch] dissolves up to 30 at.% Al at 1450°C a = 352.32 [V-C], pure Ni at 20°C a = 353.55 at 0.95 at.% Hf [1985Mis] a = 353.88 at 8.0 at.% Al [1985Mis] a = 352.4 [Mas2] dissolves 21.3 at.%Al at 1372°C [2003Sal] a = 677.6 to 677.9 [1981Fer], in Hf rich two-phase c = 537.2 to 543.3 alloys Si stabilized [20003Sch] a = 753.5 to 754.9 [1981Fer] c = 690.6 to 691.1 a = 513.43 to 533.10 c = 542.2 to 541.4 a = 325.3 b = 1083.1 c = 428.2 a = 952.1 b = 1376.3 c = 552.2 a = 523.0 to 529.0 c = 865.0 to 874.0
[1981Fer] [20003Sch] [1981Fer] [20003Sch] [1981Fer]
[1981Fer]
a = 389.0 to 393.0 [1981Fer] c = 893.0 to 889.0 Si stabilized [20003Sch] a = 398.0 to 401 c = 1714.0 to 1713.0 a = 661.3 b = 736.7 c = 481.1 a = 402.8 c = 489.1
[1981Fer]
[2003Sal] for 37 at.% Al [2003Sal] 59.5 to 63.2 at.% Al
Landolt-Börnstein New Series IV/11A3
Al–Hf–Ni
87
Phase/ Temperature Range [°C] NiAl 1651
Pearson Symbol/ Space Group/ Prototype cP2 Pm3m CsCl
Ni5Al3 700
oC16 Cmmm Pt5Ga3 cP4 Pm3m AuCu3
a = 753.0 b = 661.0 c = 376.0 a = 356.77 to 358.90
cI112 Ia3d Ni3Ga4 tI12 I4/mcm CuAl2 oC8 Cmcm CrB tI40 I4/mcm (or I4/m) Zr9Pt11 cF24 Fd3m Cu2Mg aP20 P1 Hf3Ni7
a = 1140.8 0.1
[2003Sal] solid solubility ranges from 30.8 to 58.0 at.% Al dissolves < 5 at.% Hf at 1350°C [1990Tak] [2003Sal], for 37 at.% Al solid solubility ranges from 32.0 to 37.0 at.% Al [2003Sal], solid solubility ranges from 24.0 to 27.0 at.% Al dissolves 8 at.% Hf at 1000°C [1981Nas], 11 at.% Hf by rapid solidification [2003Sal]
a = 674.3 c = 558.0
[1981Fer] [1983Nas]
a = 322.0 b = 982.0 c = 412.0 a = 979.0 c = 653.0
[1981Fer] [1983Nas]
a = 690.6
[1981Fer] Si stabilized [20003Sch]
a = 651.38 b = 658.9 c = 762.71 = 104.87° = 104.60° = 112.71° a = 912.6 b = 907.8 c = 1227.5 a = 1227.5 b = 907.8 c = 912.6 a = 642.75 b = 800.07 c = 855.4 = 75.18° = 68.14° = 75.61°
[1981Fer] [1983Nas]
Ni3Al 1372
Ni3Al4 < 702 Hf2Ni 1200 HfNi 1530 Hf9Ni11 < 1340
HfNi2 1200 Hf3Ni7 1016 - 1250
Hf7Ni10 1290
Hf8Ni21 1300 - 1175
Landolt-Börnstein New Series IV/11A3
oC68 Aba2 Zr7Ni10 C2ca
aP29 P1 Hf8Ni21
Lattice Parameters Comments/References [pm] a = 286.00 to 288.72
[1981Fer]
[1981Fer]
[1983Nas]
[1981Fer]
MSIT ®
Al–Hf–Ni
88 Phase/ Temperature Range [°C] HfNi3(h) 1350 - 1200
HfNi3(r) 1200 Hf2Ni7 1480
HfNi5 1240 * -1, HfNiAl
* -2, HfNi2Al 1450
Pearson Symbol/ Space Group/ Prototype hR12 R3m BaPb3 hP40 P63/mmc TaRh3 mC36 C2/m Zr2Ni7
cF24 F43m AuBe5 hP9 P62m Fe2P
cF16 Fm3m MnCu2Al
Lattice Parameters Comments/References [pm] a = 527.87 c = 1923.24 a = 525.5 c = 1926.0 a = 527.10 to 528.6 c = 2130.0 to 2139.16 a = 462.0 to 468.0 b = 819.1 to 831.7 c = 1210.2 to 1224.0 = 94.7 to 95.905° a = 668.3 to 669.7
a = 686.0 c = 342.0 a = 684.7 c = 345.9 a = 688.5 c = 683.8 a = 687.3 c = 343.7 a = 608.1
a = 601.8
a = 607.3 a = 606.5
a = 608.2
a = 601.1 a = 607.4 a = 608.1
* -3, Hf3Ni6Al16
MSIT®
tI16 I4/mmm ZrNi2Al5
a = 401.0 c = 1412.0
[1981Fer] [1981Nas] [1981Fer, 1981Nas, 1983Nas]
[1981Fer, 1981Nas, 1983Nas] dissolves up to 11 at.% Al at 1000°C and 14 at.% at 1400°C [1981Fer, 1981Nas]
[1966Mar], annealed at 900°C for 480 h [1967Kri] [1968Dwi], annealed between 700 and 900°C [1974Fer], annealed at 600°C (> 168 h) [1964Mar], at 50 at.% Ni, 25 at.% Al and 25 at.% Hf, annealed at 800°C for 480 h [1981Nas], in an alloy of 60 at.% Ni, 25 at.% Al and 15 at.% Hf, annealed at 1200°C for 168 h [1981Nas], in the same alloy as above but annealed at 1000°C for 168 h [1981Nas], in an alloy of 62 at.% Ni, 15 at.% Al and 23 at.% Hf, annealed at 1200°C for 168 h [1981Nas], in an alloy of 70 at.% Ni, 5 at.% Al and 23 at.% Hf, annealed at 1200°C for 168 h [1981Nas], in an alloy of 61.3 at.% Ni, 20.3 at.% Al and 18.4 at.% Hf, annealed at 1200°C for 168 h [1981Nas], in the same alloy as above but annealed at 1000°C for 168 h [1999Dar] [1969Mar1, 1969Mar2]
Landolt-Börnstein New Series IV/11A3
Al–Hf–Ni Phase/ Temperature Range [°C] * -4, Hf6Ni8Al15
* -5, Hf6NiAl2
* -6, Hf5Ni4Al * 1, Hf5Ni3Al7
* 2, Hf3NiAl5
* 2, Hf10Ni19Al 1000 * 3, Hf4Ni16 Al5
Pearson Symbol/ Space Group/ Prototype cF16 Fm3m Th6Mn23 hP9 P62m Hf6CoAl2 ? hP12 P63/mmc MgZn2 cF24 Fd3m Cu2Mg cF24 Fd3m Cu2Mg -
89
Lattice Parameters Comments/References [pm] a = 1200.0
[1966Gan1, 1966Gan2, 1969Mar1]
a = 783.0 c = 329.0
[1969Mar1, 1969Mar2]
a = 518.0 c = 841.0
[1969Mar2] [1969Mar1]
a = 734.7
[1966Mar, 1969Tes]
a = 690.5
[1969Mar1], possibly stabilized by silica
-
Denoted as L phase by [1969Mar1]
Table 2: Solidus Temperatures as a Function of Alloy Composition [1981Nas] Alloy Composition (at.%) Al Hf 8 9 15 25 13 13 23 15 20 5 16 20 5 5 18.4 20.3 22.5 2.5
Landolt-Börnstein New Series IV/11A3
SolidusTemperature [°C12°C] Ni 83 60 74 62 75 64 70 61.3 75
1237 1233 1262 1227 1262 1237 1233 1233 1233
MSIT ®
Al–Hf–Ni
90
Hf Ni Al
Fig. 1: Al-Hf-Ni. Solubility isotherms of (Ni)
0.00 80.00 20.00
Data / Grid: at.% Axes: at.%
1127°C 1027°C 927°C 827°C
10
10
(Ni)
Hf Ni Al
90
20.00 80.00 0.00
Fig. 2: Al-Hf-Ni. Pseudobinary system NiAl-HfNi2Al
Ni
1750
1651°C
L
1500
Temperature, °C
L+NiAl L+τ2
1350°C NiAl
15% Hf
1250
τ2
NiAl+τ2 1000
Hf Ni Al
MSIT®
0.00 50.00 50.00
10
20
Hf, at.%
Hf Ni Al
30.00 50.00 20.00
Landolt-Börnstein New Series IV/11A3
Al–Hf–Ni Hf Ni Al
Fig. 3: Al-Hf-Ni. Liquidus surface of the Ni corner
91 0.00 70.00 30.00
Data / Grid: at.% Axes: at.%
p1
10
20
Ni3Al
20
10
(Ni) U1
Hf2Ni7 U2 HfNi5 Hf Ni Al
p2 e1
80
30.00 70.00 0.00
Hf Ni Al
Fig. 4: Al-Hf-Ni. Partial isothermal section at 1200°C. The dashed lines represent interpolated boundaries
90
Ni
0.00 50.00 50.00
Data / Grid: at.% Axes: at.%
NiAl 10
τ2 +
NiA l
40
τ2 τ
2 +H
f2 N i7
30
(Ni)+Ni3Al
Ni3Al
20
Hf 2 Ni 7 +N i3 A l
30
NiAl +Ni3Al
Hf2 Ni +τ 7 2 +NiAl
20
40
10
(Ni) (Ni)+HfNi5 L+(Ni) Hf Ni Al Landolt-Börnstein New Series IV/11A3
50.00 50.00 0.00
60
HfNi3(r) 80 Hf8Ni21 Hf2Ni7
70
Hf3Ni7
HfNi5
L
90
Ni
MSIT ®
Al–Hf–Ni
92
Hf Ni Al
Fig. 5: Al-Hf-Ni. Partial isothermal section at 1000°C. The dashed lines represent interpolated boundaries
0.00 50.00 50.00
Data / Grid: at.% Axes: at.%
NiAl 10
40
20
3 Al
NiA l+N i
NiAl+Ni3Al+Hf2Ni7
NiAl+τ 2 +Hf2Ni7
τ2
30
40
30
20
Ni3Al
(Ni)+Ni3Al
(Ni)+HfNi5
τ 2+β HfNi3+Hf2Ni7
10
+Ni3Al
(Ni)
(Ni)+HfNi5 Hf Ni Al
60
50.00 50.00 0.00
70
β HfNi3
80
90
HfNi5
Hf2Ni7
Al Fig. 6: Al-Hf-Ni. Isothermal section at 800°C
Ni
Data / Grid: at.% Axes: at.%
L
HfAl3(r)
20
80
NiAl3
HfAl2 Hf2Al3 HfAl Hf4Al3 Hf3Al2
80
τ3
40
Ni2Al3 60
λ2
τ4 λ1
NiAl
60
40
τ1 τ2
τ5
Ni3Al 20
τ6
λ3 (Ni)
(αHf)
Hf
MSIT®
20
HfNi3 80 HfNi5 Hf2Ni 40 HfNi Hf Ni 60 9 11 Hf7Ni 10 Hf2Ni7
Ni
Landolt-Börnstein New Series IV/11A3
Al–Li–Mg
93
Aluminium – Lithium – Magnesium Gautam Ghosh Literature Data The ternary system contains many technologically important alloys for light weight, high-strength and corrosion resistant applications. Therefore, the phase equilibria of the system are of experimental and theoretical interest. Extensive studies have been carried out on the aging behavior and the structure-property relationship of Al rich alloys. Comprehensive reviews of the phase equilibria have been published by [1990Goe, 1993Gho]. [1948Sha2] was the first to report the entire liquidus surface. Later, [1977Dri, 1979Vos, 1981Sch3] reinvestigated the liquidus surface. Isothermal sections have been investigated several times [1948Sha2, 1954Wei, 1956Lev, 1955Row, 1956Row, 1973Dri1, 1973Dri2, 1976Pad, 1977Dri, 1980Sch4]. Until recent studies on the constitutional equilibria of the entire system by Schürmann and co-workers [1979Vos, 1979Gei, 1980Sch4, 1981Sch3], earlier results were subjected to considerable doubt and inaccuracy in view of high reactivity and volatility of Li and Mg. To overcome this problem, Schürmann et al. [1980Sch1, 1981Sch1] designed a special experimental apparatus to prepare the binary and ternary alloys, and the phases were analyzed by X-ray diffraction, optical metallography and electron probe microanalysis. Accordingly, much of their results are reproduced in this assessment with some amendments. Binary Systems The Al-Li binary phase diagram is taken from the assessment of McAlister [1982McA, Mas]. In his assessment, the experimental results of [1979Vos, 1980Sch2] were not reviewed. Nevertheless, the liquidus data and the invariant reaction temperatures involving liquid, (Al), LiAl, and Li3Al2 phases of [1979Vos, 1980Sch2] agree very well with those of [1982McA]. Also, all these results are in good agreement with the recent thermodynamic assessment of Saunders [1989Sau]. According to [1982McA], the peritectic reaction L+Li3Al2Li9Al4 occurs at 335°C. On the other hand, [1979Vos, 1981Sch2] proposed the peritectic reaction to be L+Li3Al2(Li) at 329°C. But, this was found to be incompatible with the thermodynamic modeling by Saunders [1989Sau]. Voss [1979Vos] reported a eutectoid reaction Li3Al2+(Li)(~Li21Al4) at 242°C, and this feature was also absent in the assessments of [1982McA, 1989Sau]. Also, [1979Vos] reported an unusually high solid solubility (about 13.0 at.%) of Al in (Li). Once again, this feature was found to be incompatible with the thermodynamic modeling of the Al-Li system [1989Sau]. In an earlier assessment [1993Gho] of the Al-Li-Mg system, the Al-Mg binary phase diagram was accepted from the experimental work of [1979Vos, 1980Sch3] which was somewhat different from Murray's assessment [1982Mur, Mas]. The major discrepancy lied in the composition range of 40 to 50 at.% Mg. Schürmann et al. [1979Vos, 1980Sch3] reported two intermediate phases (Mg10 Al11 ) and J(Mg9Al11), which were absent in the assessed phase diagram of [1982Mur]. Also, [1980Sch3] did not observe the R-phase which was reported to exist between 320 to 370°C and at 42 at.% Mg [1982Mur]. Thermodynamically assessed [1990Sau] Al-Mg phase diagram was in excellent agreement with the experimental phase diagram of [1979Vos, 1980Sch3]. Recently, the high-temperature phase equilibria between (Mg2Al3) and (Mg17 Al12 ) phases has been reinvestigated in detail [1997Su, 1998Don, 1998Lia]. Therefore, the Al-Mg phase diagram is accepted from the experimental and thermodynamic calculation of [1998Lia], which was also accepted in the recent evaluation by [2003Luk]. In the composition range of 50 to 60 at.% Al, the phase diagram of [1998Lia] is substantially different from that of [1979Vos, 1980Sch3] and similar to the assessed diagram of [1982Mur]. [1998Lia] found that between (Mg2Al3) and (Mg17 Al12 ) phases, there is only one intermediate phase J(Mg23Al30). Moreover, J phase forms by a peritectoid reaction at 410°C and decomposes by a eutectoid reaction at 210°C. Recent experimental investigations by [1997Su, 1998Don] have shown that the phase reported by Schürmann et al. [1980Sch3, 1981Sch2] in the temperature range of 410 to 452°C does not exist. To account for the additional peaks observed in the X-ray diffraction, [1997Su] assumed the presence of a hypothetical phase having Landolt-Börnstein New Series IV/11A3
MSIT ®
94
Al–Li–Mg
composition between 57 to 58 at.% Al. Donnadieu et al. [1998Don] carried out electron diffraction experiments of several alloys containing 47 to 59 at.% Al which were annealed between 425 to 445°C. They observed modulated microstructure of the phase. The wave vector characterizing the commensurate modulation is temperature and composition dependent. Therefore, the additional peaks observed by [1997Su] in X-ray diffraction could be explained by the commensurate modulation. The Li-Mg binary phase diagram is taken from the recent review and thermodynamic assessment of Nayeb-Hashemi et al. [1984Nay]. Solid Phases Depending on temperature, (Al) can dissolve up to 16.5 at.% Mg [1979Vos, 1980Sch4] and 15.8 at.% Li [1982McA, 1980Sch2]. Solid solubility of (Al) in the ternary regime is shown in Fig. 1, as a function of temperature [1965Fri, 1973Dri2, 1980Sch4]. The results of [1973Dri2, 1980Sch4] agree fairly well, but the results of [1965Fri] indicate that at a given temperature and Li content the solid solubility of Mg in (Al) was less than those reported by [1973Dri2, 1980Sch4]. The solid solubility of Al and Li in (Mg) has been reported by several investigators [1948Sha2, 1952Jon, 1976Pad, 1979Gei, 1980Sch4]. The compositions of (Mg), as a function of temperature, in the (Mg)+(Li)+LiAl() and (Mg)+LiAl()+Mg17Al12( ) three phase fields are listed in Table 1. In general, there is systematic disagreement between the results of [1980Sch4] and those of the others. Figures 2 and 3 show (Mg)/(Mg)+(Li) phase boundaries in vertical sections at 1.0 and 2.0 mass% Al respectively [1952Jon, 1954Wei, 1955Row, 1979Gei, 1980Sch4]. Figures 4, 5 and 6 show the (Mg)/(Mg)+Mg17Al12( ) phase boundaries in vertical sections at 1.0, 2.0 and 4.0 mass% Li respectively [1952Jon, 1954Wei, 1955Row, 1976Pad, 1979Vos, 1980Sch4]. Along these sections, there is significant disagreement between the results of Voss [1979Vos, 1980Sch4] and those due to [1952Jon, 1954Wie, 1955Row, 1976Pad]. In drawing the phase boundaries in Figs. 4 to 6, weightage is given to the results of Voss [1979Vos, 1980Sch4]. In the ternary regime, Mg17Al12( ) dissolves up to about 20 at.% Li, Mg2Al3() dissolves up to about 7 at.% Li, Mg23Al30(J) dissolves about 0.8 at.% Li, and LiAl () dissolves up to about 17 at.% Mg [1979Vos, 1980Sch4]. The lattice parameter of Mg17 Al12( ) decreases with the addition of Li [1956Lev]. Two ternary phases, -1 and -2, have been reported in this system. The -1 phase was first reported by Shamray [1948Sha1, 1948Sha2], and has been confirmed by subsequent investigators [1954Wei, 1955Lev, 1956Lev, 1973Tho, 1976Pad, 1979Vos, 1980Sch4]. Originally, the stoichiometry of the -1 phase was designated as LiMgAl2 [1948Sha1, 1948Sha2, 1954Wei, 1955Lev, 1956Lev]. However, recent results of [1976Pad, 1979Vos] indicate that -1 phase contains 32.0 to 34.2 at.% Li and 13.5 to 14.0 at.% Mg, and this composition is accepted here in drawing the isothermal sections. The Li and Mg contents of -1 phase reported by [1948Sha2, 1954Wei, 1955Lev, 1956Lev] differ significantly as compared to those of [1976Pad, 1979Vos]. The -2 phase, having stoichiometry Li2MgAl and NaTl type of structure, was reported by earlier investigators [1952Jon, 1955Row], but could not be confirmed in subsequent investigations [1954Wei, 1968Pau, 1979Vos]. Rather, it has been reported that -2 is a nonequilibrium transitional phase [1954Wei, 1985Nik]. Accordingly, this phase is not considered in drawing the isothermal sections. The details of the crystal structures and lattice parameters of the equilibrium solid phases are listed in Table 2. Invariant Equilibria Figure 7 shows the reaction scheme associated with the solidification of Al-Li-Mg alloys after [1981Sch3]. However, several modifications are made for consistency with the accepted Al-Mg binary phase diagram. Three pseudo-binary reactions p1, e3 and e4, all of which give rise to a maximum on the liquidus surface, have been reported [1981Sch3]. From the vertical sections reported by Voss [1979Vos, 1981Sch3], the temperatures of the three maxima are estimated to be 545, 485 and 48010°C respectively. The pseudo binary reactions p1 [1948Sha1, 1981Sch3] and e3 [1981Sch3] give rise to the formation of the ternary phase -1, but the latter reaction was originally reported to be occurring at 477°C and peritectic type i.e., L+-1=LiAl() [1948Sha1]. [1981Sch3] reported that three U type reactions U6, U7 and U8 occur at 458, 451 and 449°C, respectively. In this assessment, the U6 invariant reaction of [1981Sch3] is rewritten as a MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Li–Mg
95
ternary peritectic reaction P1. Since the (Mg 10Al11) phase of [1981Sch3] does not exist in the accepted Al-Mg phase diagram and the J(Mg23Al30) phase forms by a solid-state reaction, the U7 and U8 reactions of [1981Sch3] are not accepted. Also, it is doubtful if three invariant reactions, as proposed by Schürmann et al. [1981Sch3], occurring within a temperature interval of 9°C could be firmly established. A new ternary peritectic reaction P2, so far undetected, has been introduced at the Li corner and it is expected to be occurring at 350 20°C. All these amendments are consistent with the experimentally observed isothermal sections, vertical sections and the accepted binary phase diagrams. The compositions of the phases participating in the invariant equilibria [1979Vos, 1981Sch3] are listed in Table 3. Liquidus, Solidus and Solvus Surfaces Figure 8 shows the liquidus surface and the melting grooves separating eleven areas of primary crystallization [1979Vos, 1981Sch3]. Approximate isotherms at 25°C interval are also shown in Fig. 8. There is considerable discrepancy between the liquidus surface reported by Voss [1979Vos, 1981Sch3] and those due to Shamray [1948Sha2] and Drits et al. [1977Dri]. Also the binary phase diagrams accepted by Shamray is quite different from the presently accepted ones. Accordingly, the liquidus surfaces reported by [1948Sha2, 1977Dri] were not considered here. [1986Dub, 1987Dub] employed CALPHAD technique to calculate the liquidus surface of the Al corner. According to their calculation, the temperatures of invariant reactions U1 and U 2 agree very well with those of experimental ones. But, [1987Dub] predicted a ternary eutectic reaction L=(Al)+Mg17 Al12 +Mg2Al3 at 447°C since they assumed no Li-solubility in the Mg2Al3 () phase. Figure 9 shows the solidus surface of the entire ternary system, after [1981Sch3]. The diagram is still incomplete in the Li corner. The solidus temperatures as given in the vertical sections reported by Drits et al. [1973Dri1, 1977Dri] are in reasonably good agreement with [1981Sch3]. Isothermal Sections Partial isothermal sections have been reported several times [1948Sha2, 1952Jon, 1954Wei, 1956Lev, 1956Row, 1956Lev, 1973Dri1, 1973Dri2, 1976Pad, 1977Dri, 1979Gei, 1980Sch4]. Among these, the results of Schürmann et al. [1980Sch4] are considered to be the most accurate. They prepared about 178 ternary alloys in a specially designed vacuum induction furnace [1980Sch1]. The alloys were annealed at 400, 300 and 200°C for 260 h and subsequently quenched in water or oil. The phase analysis was carried out by metallography and electron probe microanalysis. The isothermal sections at 400, 300 and 200°C are shown in Figs. 10, 11 and 12, respectively. In this composition range, the essential feature of the phase fields remain same down to room temperature [1948Sha2, 1954Wei, 1956Row]. Even though several three-phase fields are shown dashed in Figs. 10 to 12, they are consistent with the reaction scheme shown in Fig. 7. Minor adjustments have been made in Figs. 8 to 12 along the binary edges. The partial isothermal sections of [1948Sha2, 1954Wei, 1956Lev, 1956Row, 1973Dri1, 1973Dri2, 1976Pad] agree qualitatively with those of [1980Sch4]. The discrepancies between the results of [1980Sch4] and those of others are primarily due to the fact that the solid solubilities of the binary intermediate phases in the ternary regime were not determined accurately. [1977Sab, 1978Sab] reported the computer calculated isothermal sections in the temperature range of 375 to 500°C. The calculations were done based on the binary solution-phase interaction parameters and compound parameters. Also, binary intermediate phases were assumed to be stoichiometric, and no ternary interaction parameter and ternary phase were taken into account. Accordingly, substantial disagreement between the calculated and the experimental isothermal sections was noticed. However, the isothermal sections of the Al corner calculated by Dubost et al. [1987Dub], agree reasonably well with those experimentally observed. Temperature – Composition Sections [1948Sha1] reported vertical sections at 5, 10, 15, 20, 30, 50 and 60 at.% Li and also along Mg17Al12-LiAl and LiMg2-Al. Among these, the vertical section at 50 at.% Li was reported to be pseudobinary type. Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Li–Mg
96
However, most of his results are incompatible with the accepted binary phase diagram. [1954Wei] reported two vertical sections along Mg-Li5Mg4 and LiMg7-Al. [1973Dri1] determined two isopleths at 30 and 32 mass% Mg. [1977Dri] reported an isopleth at 60 mass% Mg. Thermodynamics [1991Mos] determined the enthalpies of mixing of liquid Al-Li-Mg alloys in the temperature range of 596 to 758°C using an isothermal high temperature mixing calorimeter. Their data indicate the presence of ternary interactions. Miscellaneous Because of structural applications, the decomposition behavior of supersaturated Al-Li-Mg alloys have been studied several times [1950Fro, 1971Fri1, 1971Fri2, 1982Cha, 1983Fri, 1985Nik, 1986Kru, 1987Flo, 1994Kra, 1997Kim, 1998Cho]. The mechanical properties associated with such decomposition process have also been studied a number of occasions [1950Bus, 1956Row, 1965Fri, 1971Fri1, 1982Cha, 1983Fri, 1984Gil, 1985Nik, 1986Kru, 1994Kra, 1997Hwa]. Decomposition of supersaturated Al-(1.5 to 2.0)Li-(4 to 6)Mg (mass%) alloys take place through the formation of a metastable phase [1973Tho]. The structure of this metastable phase has been reported [1980Shc] to be face-centered monoclinic having lattice parameters a = c = 2000.4 pm, b = 1979.7 pm and = 88.83°. [1993Nii, 1994Tsa] reported the formation of a face-centered icosahedral phase in rapidly solidified Li25Mg25Al50 and Li10 Mg40 Al50 alloy, respectively. The electronic origin of such a quasicrystalline phase has been discussed by [1997Del]. It has been predicted [1994Hos] that Mg will occupy the Al sublattice in the metastable phase LiAl3 having L12 structure. References [1948Sha1]
[1948Sha2]
[1950Bus]
[1950Fro] [1952Jon]
[1954Wei]
[1955Row]
[1955Lev]
MSIT®
Shamray, F.I., Kurnakov, N.S., “The Ternary System Aluminium-Magnesium-Lithium. II. State Diagrams of Auxilliary Sections” (in Russian), Bull. Acad. Sci. URSS, Classe Sci. Chim., (1), 83-94 (1948) (Experimental, Equi. Diagram, 0) Shamray, F.I., “Ternary System: Aluminium-Magnesium-Lithium. III. Description of the Ternary System Aluminium-Magnesium-Lithium. Projection of the Liquidus Surface, Isotherms at 400°C and 20°C, and the Process of Crystallisation” (in Russian), Izv. Akad. Nauk SSSR, Otdel Khim. Nauk, (3), 290-301 (1948) (Experimental, Equi. Diagram, *, 0) Busk, R.S., Leman, D.L., Casey, J.J., “The Properties of Some Magnesium-Lithium Alloys Containing Aluminium and Zinc”, Trans. AIME, J. Met., 188, 945-951 (1950) (Experimental, 6) Frost, P.D., Kura, J.G., Eastwood, L.W., “Aging Characteristics of Magnesium-Lithium Base Alloys”, Trans. AIME, J. Met., 188, 1277-1282 (1950) (Experimental, 3) Jones, A., Lennon, J.H., Nash, R.R., W.H. Chang, E.G. Macpeek, “Magnesium Alloy Research Studies”, U. S. At. Energy Comm. Publ., (AF-TR-52-169), 1-130 (1952) (Experimental, Equi. Diagram, #, 16) Weinberg, A.F., Levison, D.W., McPherson, D.J., Rostoker, W., Wolfe, C.P., Humphreys, A., Dvorak, J., Manasevit, H., DuPraw, W., “Phase Relationships in Magnesium - Lithium - Aluminum and Magnesium - Lithium - Zinc Alloys”, Armour Res. Found. Rep., (AD-16567), 1-94 (1954) (Experimental, Equi. Diagram, #, *) Rowland, J.A., Armantrout, Jr.,C.E., Walsh, D.F., “Magnesium-Rich Corner of the Magnesium-Lithium-Aluminum System”, Trans. AIME, J. Met., 203, 355-359 (1955) (Experimental, Equi. Diagram, #, *, 11) Levison, D.W., “Discussion on Magnesium-Rich Corner of the Magnesium - Lithium Aluminum System by Rowland, J.A.,Jr., Armantrout, C.E., Walsh, D.F.”, Trans. AIME, J. Met., 203, 1267 (1955) (Experimental, 1)
Landolt-Börnstein New Series IV/11A3
Al–Li–Mg [1956Lev] [1956Row]
[1965Fri]
[1968Pau]
[1971Fri1]
[1971Fri2]
[1972Sam]
[1973Dri1]
[1973Dri2]
[1973Tho] [1976Pad]
[1977Dri]
[1977Sab] [1978Sab]
[1979Gei]
[1979Vos]
[1980Sch1]
Landolt-Börnstein New Series IV/11A3
97
Levison, D.W., McPherson, D.J., “Phase Relations in Magnesium-Lithium-Aluminum Alloys”, Trans. Am. Soc. Met., 48, 689-697 (1956) (Experimental, Equi. Diagram, #, *, 9) Rowland, J.A., Armantrout, C.E., Walsh, D.F., “Experimental Magnesium Alloys Containing Nickel, Manganese, Lithium and Aluminum”, U. S. Bur. Mines, Rep. Invest., 5250, 1-21 (1956) (Experimental, Equi. Diagram, #, 11) Fridlyander, I.N., Shamray, V.F., Shiryaeva, N.V., “Phase Composition and Mechanical Properties of Alloys of Aluminum with Magnesium and Lithium” (in Russian), Izv. Akad. Nauk SSSR, Met., (2), 153-158 (1965) (Experimental, Equi. Diagram, #, 9) Pauly, H., Weiss, A., Witte, H., “FCC Alloys of Composition Li2MgX with Body-Centred Substructure” (in German), Z. Metallkd., 59, 414-418 (1968) (Crys. Structure, Experimental, *, 15) Fridlyander, I.N., Sandler, V.S., Nikol'skaya, T.I., “Change in the Phase Composition of Aluminum-Magnesium-Lithium Alloy 01420 During Aging” (in Russian), Metall. i Term. Obra. Metallov., (5), 2-5 (1971) (Crys. Structure, Experimental, 7) Fridlyander, I.N., Sandler, V.S., Nikol'skaya, T.I., “Investigation of the Aging of Aluminum-Magnesium-Lithium Alloys”(in Russian), Fiz. Met. Metalloved., 32, 767-774 (1971) (Experimental, 15) Samson, S., “Structural Relationships Among Complex Intermetallic Compounds” (Abstract Only), IXth International Congress of Crystallography, Kyoto, Japan, VII-7, 96 (1972) (Crys. Structure, Experimental, 0) Drits, M.E., Padezhnova, E.M., Guzei, L.S., “On the Question of the Mg-Li-Al System” in “Certain Regularities in the Structure of Phase Diagrams of Metallic Systems”, Baikov Inst. Met., Nauka, Moscow, 147-153 (1973) (Experimental, Equi. Diagram, #, *, 5) Drits, M.E., Kadaner, E.S., Turkina, N.I., Kuz'mina, V.I., “Study of Phase Equilibria in the Solid State in the Al-Corner of the Al-Mg-Li System” (in Russian), Izv. Akad. Nauk SSSR, Met., (2), 225-229 (1973) (Experimental, Equi. Diagram, #, *, 5) Thompson, G.E., Noble, B., “Precipitation Characteristics of Al-Li Alloys Containing Mg”, J. Inst. Met., 101, 111-115 (1973) (Crys. Structure, Experimental, 6) Padezhnova, E.M., Melmik, E.V., Guzei, L.S., Guseva, L.N., “Phase Equilibria in the Magnesium-Lithium-Aluminum System at 300°C” (in Russian), Izv. Akad. Nauk SSSR, Met., (4), 222-226 (1976) (Experimental, Equi. Diagram, #, *, 8) Drits, M.E., Padezhnova, E.M., Guzei, L.S., “Magnesium - Lithium - Aluminum Phase Diagram” (in Russian), Izv. Akad. Nauk SSSR, Met., (2), 205-209 (1977) (Experimental, Equi. Diagram, #, *, 5) Saboungi, M.L., Hsu, C.C., “Computation of Isothermal Sections of the Al-Li-Mg System”, Calphad, 1, 237-251 (1977) (Equi. Diagram, Theory, Thermodyn., 29) Saboungi, M.L., Hsu, C.C., “Estimmation of Isothermal Sections of Ternary Phase Diagrams of Lithium Containing Systems: The Al-Li-Mg System” in “Applications of Phase Diagrams in Metallurgy and Ceramics”, Vol. 2, NBS Special Publ. No 496, Washington, DC, 1109-1138 (1977) (Equi. Diagram, Theory, Thermodyn., 29) Geissler, I., “Phase Equilibria of Al-Li-Mg Alloys at 200, 300 and 400°C and their Hardness in the as Cast State” (in German), Ph. D. Thesis, TU Clausthal (1979) (Experimental, Equi. Diagram, #, *, 44) Voss, H.-J., “Development of an Apparatus for Melting Lithium-Containing Magnesium-Aluminium Alloys and its use for Thermal Analysis” (in German), Ph. D. Thesis, TU Clausthal, 82 pp., (1979) (Experimental, Equi. Diagram, #, *, 14) Schuermann, E., Geissler, I.K., “Phase Equilibria in the Solid Condition of the Aluminum rsp. the Magnesium-rich Corner of the Ternary System of Aluminum-Lithium-Magnesium. Part I. Testing Methods and Design of a Proper Melting Aggregate for Aluminum-Lithium-Magnesium Alloys” (in German), Giessereiforschung, 32(2), 163-164 (1980) (Experimental, Equi. Diagram, #, *, 4)
MSIT ®
98 [1980Sch2]
[1980Sch3]
[1980Sch4]
[1980Shc]
[1981Sch1]
[1981Sch2]
[1981Sch3]
[1982Cha]
[1982McA] [1982Mur] [1983Fri]
[1984Gil]
[1984Nay]
[1985Nik] [1986Dub]
MSIT®
Al–Li–Mg Schuermann, E., Geissler, I.K., “Phase Equilibria in the Solid Condition of the Aluminum rsp. the Magnesium-rich Corner of the Ternary System of Aluminum -Lithium - Magnesium Part II. Phase Equilibria in the Solid Condition of the Aluminium rsp Magnesium Rich Zones of the Binary Systems Aluminium-Lithium and Magnesium-Lithium” (in German), Giessereiforschung, 32(2), 165-167 (1980) (Experimental, Equi. Diagram, #, *, 17) Schuermann, E., Geissler, I.K., “Phase Equilibria in the Solid Condition of the Aluminum rsp. the Magnesium-rich Corner of the Ternary System of Aluminum-Lithium-Magnesium. Part III. Phase Equilibria in the Solid Condition of the Binary System AluminiumMagnesium” (in German), Giessereiforschung, 32(2), 167-170 (1980) (Experimental, Equi. Diagram, #, *, 15) Schuermann, E., Geissler, I.K., “Phase Equilibria in the Solid Condition of the Aluminum rsp. the Magnesium-Rich Corner of the Ternary System of Aluminum-Lithium-Magnesium. Part IV. Phase Equilibria in the Solid Condition of the Ternary System of AluminumLithium-Magnesium” (in German), Giessereiforschung, 32(2), 170-174 (1980) (Experimental, Equi. Diagram, #, *, 4) Shchegoleva, T.V., Rybalko, O.F., “The Structure of the Metastable S'-Phase in an Al-Mg-Li Alloy” (in Russian), Fiz. Met. Metalloved, 50(1), 86-90 (1980) (Crys. Structure, Experimental, 7) Schuermann, E., Voss, H.-J., “Investigation of the Melting Equilibria of the MagnesiumLithium-Aluminum Alloys. Part I. Description of the Melting Equipment and Realization of the Research” (in German), Giessereiforschung, 33(1), 33-35 (1981) (Experimental, Equi. Diagram, *, 5) Schuermann, E., Voss, H.-J., “Investigation of the Melting Equilibria of the Magnesium-Lithium-Aluminum Alloys. Part IV. Melting Equilibria of the Binary System Magnesium - Lithium” (in German), Giessereiforschung, 33(2), 43-46 (1981) (Experimental, Equi. Diagram, #, *, 17) Schuermann, E., Voss, H.-J., “Investigation of the Melting Equilibria of the Magnesium-Lithium-Aluminum Alloys. Part V. Melting Equilibria of the Ternary System of Magnesium - Lithium-Aluminum” (in German), Giessereiforschung, 33(2), 47-53 (1981) (Experimental, Equi. Diagram, #, *, 4) Chanani, G., Narayanan, G. H., Telesman, I.J., “Heat Treatment, Microsrtucture and Mechanical Property Correlations in Al-Li-Cu and Al-Li-Mg P/M Alloys”, “High-Strongth Powder Metallurgy Aluminum Alloys”, Proc. Conf., Dallas, TX, 1982, TMS-AIME, Warrandale, PA, 341-368 (1968) (Crys. Structure, Experimental, 14) McAlister, A.J., “The Al-Li (Aluminum-Lithium) System”, Bull. Alloy Phase Diagrams, 3(2), 177-183 (1982) (Assessment, Equi. Diagram, Thermodyn., #, *, 31) Murray, J.L., “The Al-Mg (Aluminum-Magnesium) System”, Bull. Alloy Phase Diagrams, 3(1), 60-74 (1982) (Equi. Diagram, Review, Thermodyn., #, *, 112) Fridlyader, I.N., Sandler, V.S., Nikol'skaya, T.I., “Characteristics of the Structure and Properties of 1420 Aluminum Alloy” (in Russian), Metall. Term. Obra. Metallov, (7), 20-22 (1983) (Crys. Structure, Experimental, 6) Gilman, P.S., “The Physical Metallurgy of Mechanically Alloyed, Dispersion-Strengthened Al-Li-Mg and Al-Li-Cu Alloys” in “Aluminum-Lithium Alloys II”, Proc. Conf., Monterey, 1984, TMS-AIME, Warrandale, PA, 485-506 (1984) (Crys. Structure, Experimental, 11) Nayeb-Hashemi, A.A., Clark, J.B., Pelton, A.D., “The Li-Mg (Lithium-Magnesium) System”, Bull. Alloy Phase Diagrams, 5(4), 365-374 (1984) (Equi. Diagram, Review, Thermodyn., #, *, 37) Nikulin, L.V., Shevrikuko, S.B., Belozerova, E.V., “Properties and Structure of Cast Mg-Li-Al -Alloys” (in Russian), Tsvetn. Met., (12), 56-59 (1985) (Experimental, 5) Dubost, B., Bompard, P., Ansara I., “Contribution to the Establishment of the Equilibrium Diagram of Phases of the Al-Li-Mg System” (in French), Mem. Etud. Sci. Rev. Metall., 83, 437 (1986) (Experimental, Equi. Diagram, Theory, #, 6) Landolt-Börnstein New Series IV/11A3
Al–Li–Mg [1986Kru]
[1987Dub]
[1987Flo] [1989Sau] [1990Goe] [1990Sau] [1991Mos] [1993Gho]
[1993Nii]
[1994Hos]
[1994Kra] [1994Tsa]
[1997Del]
[1997Hwa]
[1997Kim]
[1997Su]
[1998Cho]
[1998Don]
[1998Lia]
Landolt-Börnstein New Series IV/11A3
99
Kruglov, B.F., Khristoferov, C.M., Sheikman, A.I., “Effect of Natural Aging in an Al-2.2 wt.% Li-5.6 wt.% Mg Alloy” (in Russian), Fiz. Met. Metalloved, 61(1), 190-191 (1986) (Experimental, 11) Dubost, B., Bompard, P., Ansara, I., “Experimental Study and Thermodynamic Calculation of the Al-Li-Mg Equilibrium Phase Diagram”, J. Phys.(France), C3, 473-479 (1987) (Experimental, Equi. Diagram, Theory, Thermodyn., #, 15) Flower, H.M., Gregson, P.J., “Solid State Phase Transformations in Aluminum Alloys Containing Lithium”, Mater. Sci. Technol., 3, 81-90 (1987) (Crys. Structure, Review, 116) Saunders, N., “Calculated Stable and Metastable Phase Equilibria in Al-Li-Zr Alloys”, Z. Metallkd., 80, 894-903 (1989) (Equi. Diagram, Theory, Thermodyn., #, 78) Goel, N.C., Cahoon, J.R., “The Al-Li-Mg System (Aluminum-Lithium-Magnesium)”, Bull. Alloy Phase Diagrams, 11, 528-546 (1990) (Equi. Diagram, Review, #, *, 25) Saunders, N., “A Review of Thermodynamic Assessment of the Al-Mg and Mg-Li Systems”, Calphald, 14, 61-70 (1990) (Equi. Diagram, Theory, Thermodyn., #, 78) Moser, Z., Agarwal, R., Sommer, F., Predel, B., “Calorimetric Studies of Liquid Al-Li-Mg Alloys”, Z. Metallkd., 82, 317-321 (1991) (Experimental, Thermodyn., 9) Ghosh, G., “Aluminium-Lithium-Magnesium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.12175.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 48) Niikura, A., Tsai, A.P., Inoue, A., Masumoto, T., Yamamoto, A., “Novel Face-Centered Icosahedral Phase in Al-Mg-Li System”, Jpn. J. Appl. Phys., 32, L1160-L1163 (1993) (Crys. Structure, Experimental, 9) Hosoda, H., Sato, T., Tezuka, H., Mishima, Y., Kamio, A., “A Substitution Behavior of Additional Elements in the L1 2-Type Al3Li Metastable Phase in Al-Li Alloys” (in Japanese), J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Theory, 26) Kramer, L.S., Langan, T.J., Pickens, J.R., “Development of Al-Mg-Li Alloys for Marine Applications”, J. Mater. Sci., 29, 5826-5832 (1994) (Experimental, Equi. Diagram, 23) Tsai, A.P., Yamamoto, A., Niikura, A., Inoue, A., Masumoto, T., “Structural Model of a Face-Centered Icosahedral Phase in Al-Mg-Li Alloys”, Philos. Mag. Lett., 69, 343-349 (1994) (Crys. Structure, Experimental, 15) Dell'Acqua, G., Krajci, M., Hafner, J., “Face-Centered Al-Mg-Li Alloys: a Free-Electron Quasicrystal”, J. Phys.: Condensed Matter, 9, 10725-10738 (1997) (Crys. Structure, Theory, 46) Hwang, Y.H., Han, C.H., Kim, Y.W., Cho, B.J., Kim, D.H., Hong, C.P., “Effects of Heat Treatment on the Mechanical Properties in Squeeze Cast Mg-Li-Al Alloys” (in Korean), J. Korean Inst. Met. Mater., 35(12), 1653-1659 (1997) (Experimental, 15) Kim, Y.W., Hwang, Y.H., Park, T.W., Kim, D.H., Hong, C.P., “Precipitation Behavior of and During Heat Treatment in Squeeze Cast Mg-Li-Al Alloys” (in Korean), J. Korean Inst. Met. Mater., 35(12), 1609-1615 (1997) (Experimental, 9) Su, H.-L., Harmelin, M., Donnadieu, P., Baetzner, C., Seifert, H.J., Lukas, H.L., Effenberg, G., Aldinger, F., “Experimental Investigation of the Mg-Al Phase Diagram from 47 to 63 at.% Al”, J. Alloys Compd., 247, 57-65 (1997) (Crys. Structure, Experimental, Equi. Diagram, #, *, 20) Cho, B.J., Kim, D.H., Hong, C.P., “Formation and Growth of Widmanstaetten HCP phase in Mg-Li-Al Alloy” (in Korean), J. Korean Inst. Met. Mater., 36(5), 647-654 (1998) (Experimental, 11) Donnadieu, P., Harmelin, M., Seifert, H.J., Aldinger, F., “Commensurately Modulated Stable States Related to the -Phase in Mg-Al Alloys”, Philos. Mag. A, 78, 893-905 (1998) (Crys. Structure, Experimental, *, 21) Liang, P., Sung, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G., Seifert, H.J., Lukas, H.L., Aldinger, F., “Experimental Investigation and Thermodynamic MSIT ®
Al–Li–Mg
100
[2003Luk]
Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89, 536-540 (1998) (Equi. Diagram, Experimental, Thermodyn., #, *, 33) Lukas, H.-L., Lebrun, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49)
Table 1: Temperature Dependence of Solid Solubility of (Mg) in Three-Phase Fields Three-Phase Field
Temperature [°C]
(Mg) + (Li) +
400 300 200
(Mg) + +
100 400 300 200
100
Composition (at.%) Al Li 6.0 20.8 5.2 19.2 2.7 18.7 1.25 17.0 1.3 18.0 0.63 16.9 0.87 16.9 0.28 16.9 11.0 11.7 7.8 11.3 4.7 10.0 3.35 8.7 3.3 8.7 1.45 8.1 2.26 5.7 1.78 3.59
References [1980Sch4] [1977Dri] [1980Sch4] [1976Pad] [1980Sch4] [1977Dri] [1955Row] [1955Row] [1980Sch4] [1977Dri] [1980Sch4] [1976Pad] [1980Sch4] [1977Dri] [1955Row] [1955Row]
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) 660.45 (Li) 180.6 (Mg) 650 , LiAl 700 Li3Al2 520
MSIT®
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cI2 Im3m W hP2 P63/mmc Mg cF16 Fd3m NaTl hR15 R3m Bi2Te3
Lattice Parameters Comments/References [pm] a = 404.88
pure Al at 24°C [V-C]
a = 351.0
pure Li at 25°C [V-C]
a = 320.89 c = 521.01
pure Mg [V-C]
a = 637.0
[V-C, 1982McA], at 50 at.% Li 45 to 55 at.% Li
a = 450.8 c = 1426.0
[V-C, 1982McA] 60 to 61 at.% Li
Landolt-Börnstein New Series IV/11A3
Al–Li–Mg Phase/ Temperature Range [°C]
, Li9Al4, (h) 330 - 275
Pearson Symbol/ Space Group/ Prototype mC26 C2/m Li9Al4
', Li9Al4, (r) 275 , Mg 2Al3 452
-
, Mg17Al12 < 458 J, Mg23Al30 410 - 250 -1, LiMgAl2
cF1168 Fd3m Mg2Al3 cI58 I43m Mn hR159 R3 Mn44Si9 c*456
101
Lattice Parameters Comments/References [pm] a = 1915.51 b = 542.88 c = 449.88 = 107.67° -
[V-C, 1982McA]
[1982McA]
a = 2816 to 2824
a = 1054.38
60-62 at.% Al [2003Luk] 1168 atoms on 1704 sites per unit cell [2003Luk] at 41.4 at.% Al [V-C2] 39.5 to 51.5 at.% Al [2003Luk]
a = 1282.54 c = 2174.78
54.5 to 56.5 at.% Al [2003Luk]
a = 2031.0
[1972Sam, V-C]
Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
L + (Al) + -1
536
U1
L + -1 (Al) +
483
U2
L (Al) -1 L
L + -1 +
464
U3
L + (Al) +
458
P1
L + (Li) (Mg) +
436
U4
L (Mg) + +
418
E1
L + Li3Al2 + (Li)
411
U5
Landolt-Börnstein New Series IV/11A3
-1 (Al) L -1 L (Al) L (Li) (Mg) L (Mg) L Li3Al2 (Li)
Composition (at.%) Al Li 66.0 19.4 53.5 40.7 81. 5 12.0 54.2 34.5 10.8 61.5 31.0 54.5 8.4 79.3 16.4 48.3 39.8 20.1 51.2 34.4 45.5 40.8 42.5 18.6 60.5 6.0 80.7 2.9 51.9 10.7 60.5 7.2 23.9 29.3 0.2 37.5 7.9 20.2 39.5 44.5 19.0 20.6 10.2 12.6 37.7 17.7 41.6 42.5 61.0 12.6 50.5 39.4 67.2 30.5 63.0 0.2
Mg 14.6 5.8 6.5 11.3 27.7 14.5 12.3 35.3 40.1 14.4 13.7 38.7 33.5 16.4 37.4 32.3 46.8 62.3 71.9 16.0 50.4 76.2 44.6 15.9 26.4 10.1 2.3 MSIT ® 36.8
Al–Li–Mg
102
20
Fig. 1: Al-Li-Mg. The solid solubility of (Al) at different temperatures
Mg, mass %
(Al)+(+$
(Al)+J1+( 10
400
°C
430
°C (Al)+J1+0
300°C 200°C (Al) 0
Al
2
1
0
3
Li, mass%
400
Temperature, °C
Fig. 2: Al-Li-Mg. The (Mg)/(Mg)+(Li) phase boundary at a constant Al content of 1.0 mass%
(Mg)+(Li)
(Mg)
300
200
100
0
Li 38.20 Mg 61.10 0.70 Al
MSIT®
70
80
Mg, at.%
90
0.00 Li Mg 99.10 0.90 Al
Landolt-Börnstein New Series IV/11A3
Al–Li–Mg
103
400
Temperature, °C
Fig. 3: Al-Li-Mg. The (Mg)/(Mg)+(Li) phase boundary at a constant Al content of 2.0 mass%
(Mg)+(Li)
(Mg)
300
200
100
0
Li 38.30 Mg 60.40 1.30 Al
70
80
90
Mg, at.%
0.00 Li Mg 98.20 1.80 Al
400
Fig. 4: Al-Li-Mg. The (Mg)/(Mg)+(Li) phase boundary at a constant Li content of 1.0 mass%
300
Temperature, °C
(Mg)
200
(Mg)+Mg17Al12(γ)
100
0
3.40 Li Mg 96.60 0.00 Al
Landolt-Börnstein New Series IV/11A3
10
Al, at.%
3.40 Li Mg 83.20 Al 13.40
MSIT ®
Al–Li–Mg
104
400
Temperature, °C
Fig. 5: Al-Li-Mg. The (Mg)/(Mg) + Mg17Al12 ( ) phase boundary at a constant Li content of 2.0 mass%
300
(Mg)
200
(Mg)+Mg17Al12(γ)
100
0
6.70 Li Mg 93.30 0.00 Al
10
Al, at.%
6.80 Li Mg 80.20 Al 13.00
400
Temperature, °C
Fig. 6: Al-Li-Mg. The (Mg)/(Mg) + Mg17Al12 ( ) phase boundary system at a constant Li content of 4.0 mass%
300
(Mg)
200
(Mg)+Mg17Al12(γ)
100
0
Li 12.70 Mg 87.30 0.00 Al
MSIT®
10
Al, at.%
Li 13.00 Mg 70.30 Al 16.70
Landolt-Börnstein New Series IV/11A3
Landolt-Börnstein New Series IV/11A3
Fig. 7:
η + γ + τ1
464
U5
L+η+γ
P2
(Li) + Li3Al2 + (δ, δ')
ca. 350 L + Li3Al2 + (Li) δ, δ'
536
418
(Mg) + η + γ
L (Mg) + γ + η
L + (Mg) + η
L + (Li) (Mg) + η
L+β+γ
P1
U2
E1
U4
L + (Al) + β
L + (Al) + γ β
(Li) + (Mg) + η
436
(Al) + γ + β
458
L + (Al) + γ
L + τ1 (Al) + γ
U3
U1
(Al) + η + τ1
L + η (Al) + τ1 L + (Al) + τ1
(Al) + γ + τ1
483
L + τ1 η + γ
η + Li3Al2 + (Li)
L + η Li3Al2 + (Li)
L + Li3Al2 + (Li)
411
480 e4 L (Li) + η
485 e3 L τ1 + γ
ca. 545 p1 L + η τ1
Al-Li-Mg
Al-Li-Mg. Reaction scheme for the solidification of Al-Li-Mg alloys
167 e10 l (Li) + δ
335 p5 l + Li3Al2 δ'
520 p2 l + η Li3Al2
600 e1 l (Al) + η
Al-Li
250 e8 εβ+γ
410 p3 β+γε
436 e7 l γ + (Mg)
449.5 e6 lβ+γ
450.5 e5 l (Al) + β
Al-Mg 588 e2 l (Li) + (Mg)
Li-Mg
Al–Li–Mg 105
MSIT ®
Al–Li–Mg
106
Al
Data / Grid: at.%
Fig. 8: Al-Li-Mg. Liquidus surface
Axes: at.%
(Al)
0 60
20
e1
80
0 55 0 50 U2
U1 40
e5 Mg Al (β ) 2 3 60
p1
τ1
LiAl(η)
e3 Mg17Al12(γ )
0 65
60
U4
500
e10
20
Li
500
(Mg)
600
550
450
3 P2 50
E1 20
e4
p5
400
δ
p2
80
e7
450
550
U5
40
U3
0 60 Li3Al2
e6
P1
40
e2 80
60
(Li)
Al
Mg
Data / Grid: at.%
Fig. 9: Al-Li-Mg. Solidus surface
Axes: at.%
(Al) 596
450.5
20
40
53 6, U
1
80
450.5
458, P1
449.5
596
τ1 520
η
β
60
483, U2
449.5
γ
464, U3
436
60
40
418, E1
411, U5 80
20
436 ,U
4
436 (Mg)
(Li)
Li
MSIT®
20
40
60
588
80 588
Mg
Landolt-Börnstein New Series IV/11A3
Al–Li–Mg
107
Al Fig. 10: Al-Li-Mg. Isothermal section at 400°C
Data / Grid: at.% Axes: at.%
(Al) 20
80
(Al)+γ +β
(Al)+τ1+η
β
40
60
τ1 η Li3Al2
(Al)+τ1+γ
η+τ1+γ
γ
60
40
η+γ +(Mg)
l Li 3A 2 Li)+ L+(
80
20
Li3Al2+η+(Li) (Li)+(Mg)+η (Mg)
L
(Li) 20
Li
ε β+γ+ε
40
60
80
Al Fig. 11: Al-Li-Mg. Isothermal section at 300°C
Mg
Data / Grid: at.% Axes: at.%
(Al)
20
80
(Al)+γ +β (Al)+τ 1+η
β
40
η
60
(Al)+τ 1+γ
τ1
ε β+γ+ε
η+τ1+γ
Li3Al2 60
γ 40
η+γ +(Mg)
Li3Al2+δ'+(Li)
δ'
(Li)+L+δ
'
80
(Li)+(Mg)+η
L
Li
Landolt-Börnstein New Series IV/11A3
20
Li3Al2+η+(Li)
(Mg)
(Li) 20
40
60
80
Mg
MSIT ®
Al–Li–Mg
108
Al Fig. 12: Al-Li-Mg. Isothermal section at 200°C
Data / Grid: at.% Axes: at.%
(Al)
20
80
(Al)+γ +β (Al)+τ 1+η
40
τ1 η Li3Al2
β 60
ε
(Al)+τ 1+γ
η+τ +γ 1
β+γ+ε
γ
60
40
η+γ +(Mg)
Li3Al2+δ+(Li)
δ 80
20
Li3Al2+η+(Li) L
Li
MSIT®
η+(Li)+(Mg)
δ+(Li)+L 20
(Mg) 40
(Li)
60
80
Mg
Landolt-Börnstein New Series IV/11A3
Al–Li–Si
109
Aluminium – Lithium – Silicon Oksana Bodak Literature Data The first studies on Al-Li-Si were published in 1926 and the first reviews were made by [1991Goe] and [1995Pav]. Thermal analysis and metallographic techniques have been used to construct a partial liquidus projection for Al rich alloys, Fig. 1, using the data of [1977Dri, 1984Han]. Although topologically similar in the sense that both groups reported the presence of a pseudobinary eutectic reactions L(Al)+-1, and two ternary eutectic reactions L(Al)+(Si)+-1 and L(Al)+LiAl+-1, the results of both groups differ largely in locating the invariant points and in the liquidus isotherms for the primary -1 region. There is also considerable uncertainty with regard to the composition of the ternary compound -1 and to the extension of its homogeneity range. Historically [1926Ass] was the first to study Al-Li-Si alloys with a view to improve their mechanical properties, by ageing between 25 and 525°C. He deduced that the section Al-Li3Si (Li13Si4?) was a pseudobinary section, which is understandable as he was not aware that there is an additional compound, -1. The first report of a ternary compound [1949Boo1] merely stated that the addition of sufficient Li to Al-Si alloys revealed a new phase LixAlySiz. Much more details were revealed by [1949Boo2]. Alloys from 1 to 20 mass% Si were thermally analyzed at cooling rates of 8 K·min-1, remelted under a 50 KCl, 50 LiCl flux with the addition of 1 mass% Li and the thermal analysis repeated. A ternary eutectic reaction was located at 569°C. For hypereutectic Al-Si alloys additions of >1 mass% Li gave a ternary compound as the primary phase. The most significant finding concerned the composition of the ternary compound. An alloy with 7.4Li-11.9Si (at.%) was shown by metallography to contain primary ternary compound. This phase was extracted with hot HCl, the extract was dried and chemically analyzed as 44.1Li-29.6Si (at.%). This composition is close to the formula Li3Al2Si2 for -1. In later work [1976Kad] showed a pseudobinary eutectic e7 L(Al)+Li3Al2Si2, Fig. 1. Using electron probe microanalysis combined with the nuclear microprobe [1987Deg] showed that the primary phase in as cast alloy containing 16.1Li-6.6Si (at.%) was Li3Al2Si2. The crystal structure was not established. However, [1960Now, 1976Sch, 1984Han] refer to the ternary compound as LiAlSi, the lattice parameter of which are very close to the LiAlSi after [1960Now]. The designation of the ternary compound -1 as LiAlSi stems from [1960Now] who prepared about 30 alloys from the elements by heating them in sealed (welded) Fe crucibles at 900-1000°C for 2 h. Practically no attack was observed on the crucible. Examination of the alloys, presumably in their cast state, was solely by X-ray powder diffraction analysis. A cubic phase with a = 594 pm was found at the composition “LiAlSi”. The new phase with a lattice parameter a = 613 pm was detected at the composition "Li2Al2Si". With lower Si contents, on the section “LiAlSi” - LiAl, at a composition of 43.5Li-13Si (at.%), the X-ray examination proved that the alloy was heterogeneous. At the composition “Li2AlSi” the cubic phase had a lattice parameter a = 612 pm. Equilibria in the solid state were studied in alloys containing less then 8.0 at.% of Li and less then 12.0 at.% Si. Aluminum (99.99 mass%), lithium (99.8 mass%), and silicon of semiconductor purity were used as initial materials. [1976Kad] who used thermal analysis and metallographic techniques to study the equilibria in Al-rich alloys showed a wide two-phase region in which (Al)+-1 coexist and therefore a wide homogeneity region for -1. In [1995Pav] it is accepted that the ternary compound -1 is based on the formula Li3Al2Si2, as shown independently by [1949Boo2, 1987Deg], with a homogeneity region that includes the composition “LiAlSi” and “Li2Al2Si”. There is disagreement on the composition “Li2AlSi”; [1960Now] reports it as a cubic phase within the homogeneity region of -1, whereas [1978Ble] regards it as a single phase with cubic structure, different from “LiAlSi”, with a = 606.1 pm and a density of 1.92 g #cm-3. These data were measured from samples prepared under optimum conditions, reacting elements (99.98 Li, 99.999 Al and Si mass%) for 5 d at 500-600°C followed by slow cooling to room temperature. [1978Ble] indicate that a phase with the stoichiometry Li2AlSi did not form. [1992Pav1] studied the system at 200°C and did not detect the Li2AlSi compound and interpreted the ternary compound -1 as LiAlSi with no homogeneity range. All the
Landolt-Börnstein New Series IV/11A3
MSIT ®
110
Al–Li–Si
conflicting data reported by [1960Now, 1974Boc, 1976Sch, 1978Ble, 1992Pav1] rely exclusively on X-ray diffraction analysis of cubic phases with lattice parameters varying from a = 593 pm to a = 612 pm. Experimental difficulties with Li rich alloys have precluded the use of thermal analysis and metallographic techniques. Until results from new techniques are available for these alloys it is concluded that the only ternary phase in alloys containing 50 at.% Li is the compound -1. In the studies of alloys containing >50 at.% Li [1978Ble] reported the presence of the ternary compound Li5.3Al0.7Si2 with 1 formula unit in the elementary cell. This compound showed superlattice reflections, which were ascribed to the presence of a phase with the same composition containing 3 formula units in the elementary cell and having an “a” axis enlarged by 3 . Due to the reactivity of the alloys it was not possible to use high temperature X-ray diffraction analysis to determine whether Li16Al2Si6, with 3 formula units, was a low temperature polymorph to Li5.3Al0.7Si2. [1992Pav1] prepared ternary alloys from 98.2 Li, 99.9998 Al and 99.999 mass% Si by arc-melting in purified Ar atmosphere under 1.01·105 Pa pressure. The alloys were annealed for 240h at 200°C in Ta containers and examined by X-ray diffraction analysis. The ternary compound Li5.3Al0.7Si2 [1978Ble] was confirmed. A ternary compound Li12Al3Si4 was also observed. This compound probably corresponds to a phase called W in [1978Ble]. It has a lattice parameter a of 612 to 615 pm which is about 3 #a the lattice parameter of Li12 Al3Si4, Table 1. Further studies on the phase relations and crystal structures of the compounds were made by [2000Kev, 2001Kev, 2001Gro]. To clarify the relations among the ternary phases [2001Kev] prepared three series of samples made from aluminum powder (99.8 mass%, Alfa), lithium bulk material (99.9 mass%, Chemetall, Frankfurt), and silicon chips (99,9998 mass%, Wacker) as starting materials. The first samples were prepared by arc-melting in purified argon atmosphere. Due to high weight losses (5-10 mass%) by arc-melting, levitation melting under purified argon was performed for most of the alloys. Samples were packed into Ta containers and sealed in silica ampoules. The annealing was carried out at 250°C for up to 1 month. The results for alloys of 15 compositions in the range of 30 to 60 at.% Li, 20 to 50 at.% Si, and 10 to 50 at.% Al are reported by [2001Kev]. Alloys were powdered and investigated using an X-ray powder diffractometer Siemens D-5000 with CoK radiation. The mechanically extracted single crystals of the new ternary phases were also investigated using electron microscope Leitz-AMR 1600T with EDX-detector for the determination of composition. The -1 and -2 phases are confirmed and a new phase of the Al3Li8Si5 composition designated as -3 is found. The other ternary phases reported earlier are assumed to be metastable. The isothermal section at 250°C is presented. [2001Gro] investigated the ternary Al-Li-Si alloys by differential thermal analysis. Melting temperatures were established for the three ternary compounds LiAlSi (-1), Li5.3Al0.7Si2 (-2), and Li8Al3Si5 (-3). Additionally selected ternary alloys were also studied by DTA. These results were combined with the phase relations examined in [2001Kev]. Using these data together with some of the available information from the literature the ternary phase diagram was calculated applying the Calphad method. The thermodynamic model of the ternary system was built by extrapolating the thermodynamic data of the binary subsystems into the ternary. The liquid phase and (Al) were modeled by a simple regular solution model without any ternary interaction parameter. The three experimentally found ternary phases were modeled as stoichiometric phases although there is a homogeneity range confirmed for -1. The phase transformation temperatures found by [2001Kev] were used to fit Gibbs energy functions for the ternary phases. The resulting calculation reproduced the measured DTA data quite well, the model parameters, however were not cited [2001Gro]. This work also presents a number of isothermal sections calculated at 250, 590, 597, 605, 700, 800°C, the liquidus surface and a set of invariant equilibria, but does not give the compositions of the phases. The latest results, published by [2003Spi], puts newly questions on the composition of the compounds in the Al-Li-Si system. The authors synthesized the Li12Al3Si4 compound which according to [1992Pav1, 1992Pav2, 1992Pav3, 1996Dmy] does exist, and which categorically is denied to exist by [2001Kev, 2001Gro]. The alloy was prepared in a tantalum tube weld-sealed under an argon atmosphere. This tube was protected from air by a silica jacket sealed under vacuum. The mixture was heated for 10 h at 950°C in a vertical furnace and shaken several times for homogenization. It was then cooled down at a rate of 6 K#h-1 MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Li–Si
111
for crystal growth. The product of the reaction appeared to be not quite homogeneous, but contained predominantly black and well-crystallized material. A few black crystals were selected and analyzed by atomic absorption flame spectrometry to identify the composition. This analysis led to an Li/Al/Si ratio of 1:0.223(2):0.41(1), corresponding to a mean formula of Li14.63Al3.26Si6. The compound could then be re-prepared following this stoichiometry and obtained in practically 100% yield, as confirmed by X-ray powder pattern (m. p. 824°C). The structure of the Li15Al3Si6 compound, determined by X-ray powder and single crystal analysis, agrees well with data found earlier by [1978Ble], for the compound Li16Al2Si6. Al-Li-Si alloys were investigated in [1994Hos] with the purpose to study the influence of the third component, in this case Si, on the precipitation of metastable phase LiAl3 (`). Binary Systems The binary system Li-Si from [Mas2] is accepted. The binary Al-Si system from [2003Luk] and Al-Li from [2003Gro] are accepted. Solid Phases The -1, -2 and -3 phases of constant composition are the only stable phases in the system according to [2000Kev, 2001Kev, 2001Gro], who worked with high purity initial materials and under well controlled conditions of the experimental environment. The essential differences in composition of compound with cubic structure reported in early works [1949Boo2, 1960Now, 1976Dri, 1976Kad, 1977Dri, 1984Han, 1992Pav1] become understandable after results of [2001Kev]. In this work, in addition to -1 phase, -3 phase, also cubic, but with larger cell parameter and closely-related crystal structure, has been found (Table 1). The resemblance of the crystal structures of -1 and -3 phases and limitations of the film-method used for the determination of crystal structure in early works can be the reason of noticed inaccuracies. All three ternary phases are proposed to melt congruently: -1 at 811°C, -2 at 793°C and -3 at 833°C. The existence of a hexagonal ternary -4 phase, found by [1978Ble, 2003Spi] and the cubic -5 found by [1978Ble, 1992Pav2] need to be confirmed. These phases possibly are stabilized by impurity of other components, contained in the initial metals. Pseudobinary Systems Seven pseudobinary eutectics exist in the system according to computation [2001Gro]. Unfortunately their compositions are not given. The existence of a pseudobinary section extending between the (Al) solid solution and -1 is well established experimentally by [1976Dri, 1976Kad, 1977Dri, 1984Han]. The invariant curve for the liquid phase undergoes a maximum at 635°C for an invariant eutectic reaction according to [1976Kad, 1977Dri], at ~630°C according to [1984Han] and at 657°C according to [2001Gro]. There is disagreement on the reported composition of the eutectic maximum, Fig. 1. [1976Kad, 1977Dri] give a vertical section from the Al corner to 17.5 mass% Li3Al2Si2 with the eutectic composition at 9 mass% Li3Al2Si2 (5.35Li-3.57Si (at.%). [1984Han] quotes a eutectic composition that does not lie on their monovariant curve E2´E3´, Fig. 1 the scaled composition converts to 14Li-4Si (at.%) for e7´. As shown in Fig. 1 the eutectic maximum found by [1984Han] at ~630°C lies very near to the 670°C isotherm given by [1977Dri]. The discrepancies between [1976Kad, 1977Dri, 1984Han] can only be resolved by further investigation. Invariant Equilibria From the eutectic maximum, e7 or e7´ in Fig. 1, monovariant curves descend to a ternary eutectic E2´ or E2´´ and to a ternary eutectic E3´ or E3´´, respectively. According to [1976Kad, 1977Dri] E2 has the composition 28.3Li-1.5Si (at.%) whereas [1984Han] place E´2 at 31.6Li-0.8Si (at.%). The temperature of the reaction was given as 595°C [1976Kad, 1977Dri] and 592°C [1984Han]. The liquidus and the eutectic composition for binary Al-Li alloys given by [1976Kad, 1977Dri] agree more closely with [1989Che] than do the values found by [1984Han]. For example the binary Al-Li eutectic composition is quoted as 25.8 at.% Li Landolt-Börnstein New Series IV/11A3
MSIT ®
112
Al–Li–Si
[1989Che], 27.1 at.% Li [1976Kad, 1977Dri] and 30.2 at.% Li [1984Han]. On this basis the ternary eutectic point E2´´ is preferred above E2´ reported by [1984Han]. The composition of the liquid phase in the second ternary eutectic reaction shows similar differences. [1976Kad, 1977Dri] and [1984Han] place E3 at different compositions; respectively at 0.2Li-11.1Si (at.%) (shown as E3´´ in Fig. 1) and at 5.3Li-12.8Si (at.%) (shown as E3´ in Fig. 1). The ternary eutectic temperature is quoted by [1976Kad, 1977Dri] to be 565°C. [1984Han] locates it at 575°C and according to [2001Gro] the reaction happens at 577°C. The data of [1984Han] for the binary Al-Si eutectic agree well with that given by [2003Luk] whereas that of [1976Kad, 1977Dri] do not. The data of [1984Han] and the composition/temperature of E3´ are preferred above those of [1976Kad, 1977Dri]. On the basis of the assessed experimental data a partial reaction scheme is given in Fig. 2. The calculated invariant equilibria after [2001Gro] are listed in Table 2. The optimization of [2001Gro] revealed a contradiction between the higher melting temperature of -3 (compared to -1) and the eutectic E1: L(Al)+LiAl+-1 reported by [1976Kad, 1976Dri, 1984Han]. A higher melting phase -3 will always result in a tie line between -3 and (Al) at higher temperature. Therefore a eutectic between (Al), LiAl, and -1 will not occur. On the other hand, the four phases, -1, -3, LiAl, and (Al), found in as-cast (not equilibrated) alloys near to -1 and -3 give a hint for an invariant reaction which may change the tie line of -3+(Al) to LiAl+-1. In fact, in the calculation an invariant reaction -3+(Al)LiAl+-1 at 591°C emerges by fitting the parameter for -3 and -1 to the measured melting temperatures and to the experimentally observed phase equilibria at 250°C. This final version of the thermodynamic data set reproduced all experimental results of [2001Kev]. However, the calculated liquidus surface of -1 extends much closer to the Al corner than reported by [1984Han] and somewhat closer than given by [1976Kad]. Figure 3 illustrates the discrepancies between the different reports shown with dashed lines [1976Kad] and dotted lines [1984Han] and the calculation [2001Gro] of the partial liquidus surface shown with solid lines. As discussed above, a eutectic E2´´´: LLiAl+(Al)+-1 does not take place in this calculation. However, at nearly the same temperature as given for E2´´ by [1976Kad] an invariant reaction, E2: LLiAl+(Al), -3 is present in the calculation. It was concluded that the ternary eutectic with -3 instead of -1 describes the correct equilibrium. Liquidus Surface As follows from Fig. 3 there is discrepancy between experimental data of different authors. The calculated liquidus surface given in Fig. 4 after [2001Gro] differs from both experimental series shown in Fig. 3. As follows from three previous chapters additional investigations for liquidus surface are necessary. Isothermal Sections An isothermal section at 550°C was published by [1976Dri] and one at 500°C by [1977Dri]. [1976Kad] has drawn four-phase eutectic planes at 595°C and 565°C and four vertical sections, along 5 mass% Li, along 92 mass% Al, along 2 mass% Si, and one section along Al--1, up to 17.5 mass% Li3Al2Si2. [1976Dri] used an extended annealing schedule, involving 30 h homogenization at 400°C of the cast ingots followed by deformation of 70 % with different annealing, 200 h at 550°C or 200 h at 550°C, plus subsequent 400 h at 500°C; or 200 h at 550°C plus subsequent 1000 h at 200°C. All annealing procedures terminated with water quenching of the samples. Thermodynamically the resulting data for the combined solubility of Li and Si in Al at 550, 500 and 200°C are not feasible. These data were used by [1977Dri] to produce a 500°C isothermal section confined to Al contents above 88 mass%. Plotting in a common scale data from [1976Dri] at 550°C, [1977Dri] at 500°C and [1976Kad] at 595 and 565°C gives an impression of the width of the (Al)+-1 phase region represented by this group of workers. Figure 5 summarizes the data and extrapolates the boundary (Al)--1 tie lines after [1976Kad] to the -1 “composition line” between “LiAlSi” and “Li2AlSi”. See also the discussion in “Introduction”. [1976Kad] did not determine any phase boundaries below the two ternary eutectic temperatures. No check can be made to compare their vertical sections with the isothermal sections at 550 and 500°C. Comparison of the delineation of the 595 and 565°C ternary eutectic planes [1976Kad] with their published vertical sections shows reasonable agreement for the (Al)--1 tie line at 565°C but substantial disagreement for the (Al)--1 tie line at 595°C. In Fig. 5 the tie lines given by [1976Kad] at 595 and 565°C have been preferred to those derived from vertical sections. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Li–Si
113
As can be seen from Fig. 5 there is some difference between the data reported by [1976Dri, 1977Dri, 1976Kad]. The isothermal section at 600°C assessed by [1991Goe] is based on the above mentioned experimental data. The systematic investigations of isothermal sections given in [1992Pav1] and [2001Kev] (Fig. 6) indicate the punctual composition for the compound -1 (LiAlSi). The difference in composition given in works [1976Dri, 1976Kad, 1977Dri] is connected with the existence of -1 and -3 phases with closely-related crystal structures. This is clearly described in [2001Kev]. The equilibria at different constant temperatures, between 590 and 800°C are given in Figs. 7 to 11 as calculated by [2001Gro]. At 800°C (Fig. 11) only two ternary phases, -3 and -1, are present. One hundred degrees lower (Fig. 10) the third ternary phase, -2 appears together with the binary phases Li13Si4 and Li7Si3. The liquid phase extends along the Al-Li edge up to the binary Li-Si eutectic, with little extension into the ternary. At 605°C (Fig. 9) the (Al) solid solution is in equilibrium with -1. The need to reconcile the high melting point of -3 with the solid state LiAl+-1 equilibrium were resolved by [2001Gro] in a series of three nonvariant equilibria: U5: L+-1(Al)+-3 at 600°C E2: LLiAl+(Al)+-3 at 596°C U6: -3+(Al)LiAl+-1 at 591°C. The high melting point of -3 gives a tie line between -3 and an Al-rich liquid, which is also in equilibrium with -1 at 605°C (Fig. 9). The reaction U 5 transforms this tie line, L+-1, into a tie line -3+(Al) shown in Fig. 8. The heat evolution of U5 is suspected to be slow, because a substantial amount of -1 would have to be consumed in this cross-reaction. At 596°C the liquid decomposes by the eutectic reaction E2 to form LiAl+(Al)+-3. At 591°C the (Al)+-3 tie-line transforms into the -1+LiAl equilibrium which is stable down to room temperature, Figs. 7 and 6. As a result from the reaction U6 (Übergangsreaktion) the triangle -1+LiAl+(Al) appears in Fig. 7, describing a three phase field which is well supported by literature data [1976Kad, 1976Dri, 1984Han]. These results, however, would be different if the experimentally found homogeneity range for -1 phase is taken into account. Notes on Materials Properties and Applications Lithium is an important alloying element for weight saving in conventional aluminium alloys. Lithium additions to Al-Si increases the strength and elasticity of alloy, with silicon increasing in particular their hardness [1976Kad]. The improvement of the physical properties by adding silicon to aluminium-lithium alloys is attributed to the formation of lithium silicides. For compositions close to -1 a microhardness of 946 kg·mm-2 has been measured. On quenching, the alloys are in an unstable state, supersaturated with silicides which later, during ageing, appear in a highly dispersed form. Although Al-Li-Si alloys are heat treatable, the improvement in properties is small. The main effect of lithium in Al-Si alloys is the improvement of hardness by the combined effect of Li and Si [1926Ass]. During microprobe study of structure of alloys with composition near E3 the epitaxy between silicon and silicide was observed, leading to the formation of fine silicone structure [1963Boo]. Additions of aluminium to silicides of lithium increase their stability during hydrolyze in dilute H2SO 4 under argon [1974Boc]. References [1926Ass] [1949Boo1] [1949Boo2] [1960Now] [1963Boo]
Landolt-Börnstein New Series IV/11A3
Assmann, P., “The Importance of Si for the Mechanical Improvement of Al by Li or Mg” (in German), Z. Metallkd., 18, 256-260 (1926) (Experimental, 7) Boom, E.A., “A New Phase in the Al-Li-Si System” (in Russian), Dokl. Akad. Nauk SSSR, 66, 645-646 (1949) (Experimental, 3) Boom, E.A., “Physico-Chemical Investigation of Al-Li-Si Alloys” (in Russian), Dokl. Akad. Nauk SSSR, 67, 871-874 (1949) (Experimental, 5) Nowotny, H., Holub, F., “Investigations of Metallic Systems with Fluorspar Phases” (in German), Monatsh. Chem., 91, 877-887 (1960) (Crys. Structure, Experimental, 15) Boom, E.A., “On the Mechanism of the Modification of Silumin” (in Russian), Dokl. Akad. Nauk SSSR, 151, 96-97 (1963) (Experimental, 5) MSIT ®
114 [1974Boc]
[1976Dri]
[1976Kad]
[1976Sch]
[1977Dri]
[1978Ble]
[1984Han] [1987Deg]
[1989Che]
[1991Goe] [1992Pav1]
[1992Pav2]
[1992Pav3]
[1994Hos]
[1995Pav]
[1996Dmy]
[2000Kev]
MSIT®
Al–Li–Si Bockelmann, W., Schuster, H.U., “Crystallographic Aspects of Ternary Phases of Li with 3B and 4B Elements in Ionic and Non-Ionic Compounds” (in German), Anorg. Allg. Chem., 410, 741-750 (1974) (Crys. Structure, Experimental, 5) Drits, M.E., Kadaner, E.S., Kuz’mina, V.I., Turkina, N.I., “Phase Composition of Al-Rich Al-Li-Si Alloys”, Russ. Metall., (5), 177-178 (1976), translated from Izv. Akad. Nauk SSSR, Met., (5), 206-208 (1976) (Equi. Diagram, Experimental, 4) Kadaner, E.S., Turkina, N.I., Kuz’mina, V.I., “Phase Diagram of Al-Li-Si System in the Al-Rich Region”, Russ. Metall., (1), 150-153 (1976), translated from: Izv. Akad. Nauk SSSR, (1), 181-184 (1976) (Equi. Diagram, Experimental, 14) Shuster, H.U., Hinterhauser, H.W., Schäfer, W., Will, G., “Neutron Diffraction Investigations of the Phases LiAlSi and LiAlGe” (in German), Z. Naturforsch., 31B, 1540-1541 (1976) (Crys. Structure, Experimental, 3) Drits, M.E., Bochvar, N.R., Kadaner, E.S., Padezhnova, E.M., Rokhlin, L.l., Sviderskaya, E.A., Turkina, N.I., Phase Diagrams of al and Mg Systems (in Russian), Abrikosov, N.Kh., (Ed.), Nauka, Moscow, 57-58 (1977) (Equi. Diagram, Review, 4) Blessing, J., “Synthesis and Study of Ternary Phases of li with Elements of the 3 and a sub Groups”(in German), Ph. D. Thesis, Univ. Cologne,167 pp. (1978) (Experimental, Crys. Structure, 87) Hanna, M.D., Hellawell, A., “The Liquidus Surface for the Al-Li-Si System from 0 to 20 wt.% Li and Si”, Metall. Trans. A, 15A, 595-597 (1984) (Equi. Diagram, Experimental, 6) Degreve, F., Dubost, B., Dubus, A., Thorne, N. A., Bodart, F., Demortier, G., “Quantitative Analysis of Intermetallic Phases in Al-Li Alloys by Electron, Ion and Nuclear Microprobes”, J. Phys. Colloq., 48, (Suppl. C3), 505-511 (1987) (Experimental, 13) Chen, S.-W., Jan, C.-H., Lin, J.-C., Austin Change, Y., “Phase Equilibria of the Al-Li Binary System”, Metall. Trans. A, 20A, 2247-2258 (1989) (Equi. Diagram, Thermodyn., Experimental, 59) Goel, N.C., Cahoon, J.R., “Tha Al-Li-Si (Aluminium-Lithium-Silicon)”, J. Phase Equilib., 12(2), 225-230 (1991) (Equi. Diagram, Review, 9) Pavlyuk, V.V., Bodak, O.I., Dmytriv, G.S., “Interaction of Components in Li-(Mg, Al)-Si Systems” (in Russian), Ukr. Khim. Zh. (Russ. Ed.), 58, 735-737 (1992) (Equi. Diagram, Experimental, #,6) Pavlyuk, V.V., Bodak, O.I., “The Crystal Structure of Li12Mg3Si4 and Li12 Al3Si4 Compounds” (in Russian), Neorgan. Mater., 28(5), 988-990 (1992) (Crys. Structure, Experimental, 3) Pavlyuk, V.V., Dmytriv, G.S., Starodub, P.K., “Crystal Structure of the Compounds of the Li-M-X (M = Mg, Al; X = Si, Ge, Sn) Systems” (in Russian), VI Conf. Cryst. Chem. Inorg. Coord. Compounds, L’viv (Abstact), 210 (1992) (Crys. Structure, Experimental) Hosoda, H., Sato, T., Tezuka, H., Mishima, Y., Kamio, A., “Substitution Behavior of Additional Elements in the L1(2)-Type Al3Li Metastable Phase in Al-Li Alloys” (in Japanese), J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Crys. Structure, Thermodyn., Theory, 26) Pavlyuk, V., Bodak, O., “Aluminium-Lithium-Silicon”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16694.1.20, (1995) (Crys. Structure, Equi. Diagram, Assessment, 15) Dmytriv, G.S., “Phase Equilibria and Crystal Structure of Compounds in Mg-Li-Si, Ca-Li-{Si, Ge}, Al-Li-{Si, Ge, Sn}, Zn-Li-{Al, Sn}”, Summary of the thesis for kandidate science degree, 1-23 (1996) (Crys. Structure, Equi. Diagram, Experimental, 10) Kevorkov, D., Gröbner, J., Schmid-Fetzer, R., “Experimental Investigations and Thermodynamic Calculation of the Ternary Al-Li-Si Phase Diagram”, Proc. Disc. Meet. Thermodyn. Alloys, 27 (2000) (Thermodyn., Abstract)
Landolt-Börnstein New Series IV/11A3
Al–Li–Si [2001Gro]
[2001Kev]
[2003Spi]
[2003Gro]
[2003Luk]
115
Groebner, J., Kevorkov, D., Schmid-Fetzer, R., “The Al-Li-Si System. 2. Experimental Study and Thermodynamic Calculation of the Polythermal Equilibria”, J. Solid State Chem., 156, 506-511 (2001) (Equi. Diagram, Thermodyn., Experimental, Calculation, 12) Kevorkov, D., Groebner, J., Schmid-Fetzer, R., “The Al-Li-Si System. 1. A New Structure Type Li8Al3Si5 and the Ternary Solid State Phase Equilibria”, J. Solid State Chem., 156, 500-505 (2001) (Crys. Structure, Equi. Diagram, Experimental, 16) Spina, L., Tillard, M., Belin, C., “Li15Al3Si6(Li14.6Al3.4Si6), a Compound Displaying a Heterographite-Like Anionic Framework”, Acta Crystallogr., Sect. C: Cryst. Struct. Commun., C59(2), i9-i10 (2003) (Crys. Structure, Experimental, 9) Groebner, J., “Al-Li (Aluminium-Lithium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.13517.1.20, (2003) (Crys. Structure, Equi. Diagram, Assessment, 21) Lukas, H.L., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 29)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Li) < 180.6
cI2 Im3m W
a = 351.0
pure Li at 25°C [V-C2]
(Al) < 660.45
cF4 Fm3m Cu
a = 404.96
pure Al at 25°C [Mas2] dissolves up to 15 at.% Li and up to 1.5 at.% Si
Li9Al4 ( ) < 347 - 275
mC26 C2/m Li9Al4
a = 1915.51 b = 542.88 c = 449.88 = 107.671°
[2003Gro]
Li9Al4 ( ´) < 275
?
?
[Mas2]
Li3Al2 ( ) < 520
hR15 R3m Li3Al2
a = 450.8 c = 1426
[2003Gro] 60 to 61 at.% Li [Mas2]
LiAl () < 700
cF16 Fd3m NaTl
a = 637
at 50 at.% Li [2003Gro] 45 to 55 at.% Li [Mas2]
LiAl3 (´) < 190 - ~120
cP4 Pm3m Cu3Au
a = 403.8
Metastable [2003Gro]
Li2Si
mC12 C2/m Ge2Os
a = 770 b = 441 c = 656 = 113.4°
Metastable? [V-C2]
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Li–Si
116 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
Li7Si2
oP36 Pbam Li7Si2
a = 799 b = 1521 c = 443
Metastable? [V-C2]
Li7Si3 < 752
hR7 R3m Li7Si3
a = 443.5 c = 1813.4
[Mas2, V-C2]
Li12Si7 < 648
oP152 Pnma Li12Si7
a = 861.0 b = 1973.8 c = 1434.1
[Mas2, V-C2]
Li13Si4 < 722
oP34 Pbam Li13Si4
a = 799 b = 1521 c = 443
[Mas2, V-C2]
Li22Si5 < 628
cF432 F23 Li22Pb5
a = 1875
[Mas2, V-C2]
Li41Si11
cF416 F43m Cu41Sn5
a = 1871
Metastable? [V-C2]
* -1 < 811
cF12 F43m LiAlSi
a = 594
at Li0.33Al0.33Si0.33 (LiAlSi) [1960Now]
a = 593
at Li0.33 Al0.33 Si0.33 'm = 1.95 g#cm-3 'x = 1.97 g#cm-3 [1976Sch] at Li0.33 Al0.33 Si0.33 [1984Han] at Li0.33 Al0.33 Si0.33 [1992Pav1] at Li0.33 Al0.33 Si0.33 [2001Kev]
a = 593 a = 593 a = 592.82 * -2 < 793
hP8 P63/mmc Li5.3Al0.7Si2
a = 435.9 c = 813.6 a = 434.10 c = 810.52
* -3 < 833
cP16 P43m Li8Al3Si5
a = 611.46
a = 613 a = 612
MSIT®
at Li0.66 Al0.09 Si0.25 (Li5.3Al0.7Si2) 'm = 1.35 g#cm-3 'x = 1.38 g#cm-3 [1978Ble] at Li0.66 Al0.09 Si0.25 [2001Kev] at Li0.50 Al0.19 Si031 (Li8Al3Si5) [2001Kev] Li0.42Al0.29Si0.29 (Li3Al2Si2) [1949Boo2, 1976Kad] at Li0.40 Al0.40 Si0.20 (Li2Al2Si) [1960Now] at Li0.50 Al0.25 Si0.25 (Li2AlSi) [1960Now] Li0.42Al0.29Si0.29 (Li3Al2Si2) [1987Deg]
Landolt-Börnstein New Series IV/11A3
Al–Li–Si
117
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
* -4
hP24 P63/m Li15Al3Si6
a = 754.9 c = 809.7 a = 755.0 c = 813.6
* -5
cI76
a = 1062.0
at Li0.63 Al0.13 Si0.24 (Li15Al3Si6) [2003Spi], not shown on the diagram, stability is not confirmed at Li0.67 Al0.08 Si0.25 (Li16Al2Si6) [1978Ble], not shown on the diagram, stability is not confirmed at Li0.63 Al0.16 Si0.21 (Li12Mg3Si4) [1992Pav2], not shown on the diagram, stability is not confirmed
Table 2: Invariant Equilibria Reaction
T [°C]
Type
L -3
832
congruent
L -1 + -3
809
e1 (max)
L -1
810
congruent
L (Si) + -1
802
e2 (max)
L -2
800
congruent
L -3 + -2
798
e3 (max)
L + -1 -3 + (Si)
788
U1
L -2 + Li7Si3
746
e4 (max)
L -2 + Li13Si4
730
e5 (max)
L Li13Si4 + Li7Si3 + -2
727
E1
L + -2 -3 + Li7Si3
718
U2
L -3 + LiAl
686
e6 (max)
L + -3 -2 + LiAl
679
U3
L (Al) + -1
657
e7 (max)
L + Li7Si3 Li12Si7 + -3
630
D1
L + Li13Si4 Li22Si5 + -2
616
U4
L (Si) + Li12Si7 + -3
604
D2
L + -1 -3 + (Al)
600
U5
L LiAl + (Al), -3
596
E2
-3 + (Al) LiAl + -1
591
U6
L (Al) + (Si) + -1
577
E3
L + LiAl Li3Al2 + -2
518
U7
L + Li3Al2 Li9Al4 + -2
334
U8
L (Li) + Li22Si5, -2
180
D3
L (Li) + Li9Al4, -2
175
D4
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Li–Si
118
Li Al Si
Fig. 1: Al-Li-Si. Partial liquidus projection showing the data of [1977Dri] and [1984Han]; numbering of invariant reactions is adapted to [2001Gro]
0.00 65.00 35.00
Data / Grid: at.% Axes: at.%
30
10
[1977Dri, 1976Kad] [1984Han] 20
20
E3´
(Si) E3´´
660 640
30
E2´
620
e7´
E2´´
e7´´
620
610
35.00 LiAl 65.00 0.00
640 650 660 670
τ1
Li Al Si
10
700 680
70
630
80
90
Al-Li-Si
Al-Li
(Al)
Al
A-B-C
Al-Si
635 e7(max) L (Al) + τ1
? 600 e l (Al) + LiAl
L+LiAl+τ1 595
L (Al)+LiAl+τ1
E2
(Al)+LiAl+τ1
577 e l (Al) + (Si)
?
L+(Si)+τ1 575
L (Al)+(Si)+τ1
E3
(Al)+(Si)+τ1 Fig. 2: Al-Li-Si. Partial reaction scheme from assessed experimental data MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Li–Si
119
Li Al Si
Fig. 3: Al-Li-Si. Partial liquidus projection. Comparison between calculation [2001Gro] (solid) and estimations after [1976Kad] (dashed) and [1984Han] (dotted)
0.00 60.00 40.00
Data / Grid: at.% Axes: at.%
[1984Han] [1976Cad] [2001Gro]
(Si) 10
30
20
20
τ1 E3,575 E3,577 E3,565
30
10
τ3 e7,632
e7,635
E2,595
E2,592
e7,657 (Al)
LiAl Li Al Si
70
40.00 60.00 0.00
80
E2,596 U5,600
90
Si
Al
Data / Grid: at.%
Fig. 4: Al-Li-Si. Calculated liquidus surface
Axes: at.%
130 0 20
80
1200 1100 40
Li12Si7
1000
900
U2,718 60
U1,788 40
τ 1 e2,802
τ3 τ2
Li13Si4 U4
60
825
Li7Si3 E1,727
(Si)
e1,809
800
e3,798
80
20
775
Li22Si5
600 500
Li
Landolt-Börnstein New Series IV/11A3
700
20
Li3Al2
E3,577
750 U7,518 40
LiAl
U3,679
e6
60
80
E2,596 U5,600
e7 (Al)
Al
MSIT ®
Al–Li–Si
120
Si
Data / Grid: at.%
Fig. 5: Al-Li-Si. The (Al)+-1 phase region after experimental data
Axes: at.%
20
80
40
500 550 565 595
60
60
τ1
"LiAlSi"
55 50 0 0 56 5
"Li2AlSi"
80
20
Li
20
(Al)+LiAl+τ 1
40
60
80
Si Fig. 6: Al-Li-Si. Experimental isothermal section at 250°C after [2001Kev]
40
(Al)+(Si)+τ 1
(Al)+τ 1
595 500 550
Axes: at.%
20
80
40
60
Li12Si760
40
τ3
Li7Si3
Li22Si5
Al
Data / Grid: at.%
(Si)
Li13Si4
[1977Dri] [1976Dri] [1976Kad] [1976Kad]
τ1
τ2
80
20
(Li)
Li
MSIT®
20
Li9Al4
40
Li3Al2 LiAl
60
80
(Al)
Al
Landolt-Börnstein New Series IV/11A3
Al–Li–Si
121
Si
Data / Grid: at.%
(Si)
Fig. 7: Al-Li-Si. Calculated isothermal section at 590°C
Axes: at.%
20
80
40
Li12Si760
Li22Si5
40
τ3
Li7Si3 Li13Si4
60
τ1
τ2
80
20
L L 20
Li
40
LiAl
60
80
Si
Data / Grid: at.%
(Si)
Fig. 8: Al-Li-Si. Calculated isothermal section at 597°C
Al
(Al)
Axes: at.%
20
80
40
60
Li12Si760
40
τ1
Li7Si3
τ3
Li13Si4
τ2
80
20
Li22Si5
L L
Li
Landolt-Börnstein New Series IV/11A3
20
40
LiAl
60
L
80
(Al)
Al
MSIT ®
Al–Li–Si
122
Si
Data / Grid: at.%
(Si)
Fig. 9: Al-Li-Si. Calculated isothermal section at 605°C
Axes: at.%
20
80
40
60
L Li12Si760
Li13Si4 Li22Si5
40
τ3
Li7Si3
τ1
τ2
80
20
L L 20
Li
40
LiAl
60
L
Si Fig. 10: Al-Li-Si. Calculated isothermal section at 700°C
80
Al
(Al)
Data / Grid: at.%
(Si)
Axes: at.%
20
80
40
60
L 60
τ3
Li7Si3 Li13Si4
40
τ1
τ2
80
20
L
Li
MSIT®
20
40
60
80
Al
Landolt-Börnstein New Series IV/11A3
Al–Li–Si
123
Si
Data / Grid: at.%
(Si)
Fig. 11: Al-Li-Si. Calculated isothermal section at 800°C
Axes: at.%
20
80
40
60
60
40
τ1
τ3
80
20
L
Li
Landolt-Börnstein New Series IV/11A3
20
40
60
80
Al
MSIT ®
124
Al–Li–Zn
Aluminium – Lithium – Zinc Oksana Bodak Literature Data The research on this system started in 1942 when [1942Wei] established the temperature and composition of a ternary eutectic in the Zn corner. One year later [1943Bad] investigated the triangle Al-LiAl-Zn by thermal and microscopic analyses. They found two ternary compounds, -1 and -3, in the LiAl-Zn section and gave eight vertical and two isothermal sections. The -1 phase has been confirmed by [1963Che] and the region of its homogeneity has been determined more accurately; however, these authors found -1 to be in equilibrium with (Zn) neglecting the -3 phase. [1987Dub] and [1989Aud] reported a stable phase to exist in the vicinity of -4, Li3ZnAl5, not far from the Al rich end of the -1 domain. Metastable icosahedral quasicrystals are formed in this system by rapid solidification [1986Cas, 1987Che] or as grain boundary precipitates through solid - solid transformations [1987Cas] with compositions close to the stable -1 phase. The substitution behavior of additional elements in the L12 type metastable compound of Li3Al ( ´ phase) was reported in [1994Hos]. In 1995 a critical review was made inn the MSIT evaluation programs, covering the literature published until 1992, [1995Pav]. Isothermal section of the system at 197°C and crystal structures of compounds were investigated and published in [1993Pav, 1995Dmy, 1996Dmy, 1999Pav]. Alloys of the Al-Li-Zn system were prepared by arc-melting pieces of the pure metals (lithium with a purity 98.2 mass%, zinc with a purity 99.98 mass%, aluminium with a purity 99.99 mass%) under argon atmosphere. The alloys were annealed at 197°C for 400 hours in tantalum containers in evacuated quartz ampoules, quenched in cold water and examined by X-ray diffraction analysis. There are measurements of the enthalpy of mixing of liquid Al-Li-Zn ternary made by high temperature mixing calorimeter in the temperature range 456 - 682°C, [1997Kim]. They used their data in an association model to calculate the thermodynamic mixing functions of the ternary alloys on the basis of the enthalpy of mixing of the binary systems. Aluminium of purity 99.9%), 99.9% pure lithium and zinc of 99.999% were used to prepare the alloy samples for these measurements, executed under pure argon gas at atmospheric pressure. Binary Systems For the Al-Li system phase relations are accepted here as reported by [2003Gro]. For the descriptions of the Al-Zn and Li-Zn phase diagrams the versions given in [Mas2] are accepted. Solid Phases The data for the solid phases are given in Table 1. The quasicrystalline phases are formed by rapid solidification or as grain boundary precipitates by a solid-state reaction in the -1 phase region [1997Kim]. The -1 phase has a high solubility of zinc (16.7-43.3 at.% Zn at 32 at.% Li) and is formed through a peritectic reaction at higher temperature than the -3 and -4 phases [1997Kim]. According to [1993Pav, 1996Dmy, 1999Pav] three ternary compounds are formed in this system: (a) the -1 phase, Li1+xZn0.5-1.5 Al1.5-0.5 with a large homogeneity range which includes the earlier reported composition -1, Li26Al6(Zn1-xAlx)49 (b) the -3 phase, LiZn3Al with an unidentified structure and (c) the -4 phase, Li3ZnAl5. Another compound -2 on the 50 at.% Li section is reported in the work of [1996Dmy]. Pseudobinary Systems The section LiAl-Zn shown in Fig. 1 is pseudobinary [1943Bad]. The solidus and the liquidus of the LiAl phase in Fig. 1 are slightly corrected to agree with the congruent melting point of this phase in the binary system.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Li–Zn
125
The LiAl - Li2Zn3 section has been reported by [1943Bad] as pseudobinary with continuous solid solubility. However this is unlikely because LiAl and Li2Zn3 have different crystal structures, and this contradicts to the existence of the -2 phase proposed by [1996Dmy]. Invariant Equilibria The invariant equilibria established within the triangle Al-LiAl-Zn and in the pseudobinary section LiAl-Zn [1943Bad] are listed in Table 2. An additional four-phase equilibrium at 423°C has been proposed by [1943Bad] following the existence of an intermediate phase in the Al-Zn system. However, in the presently accepted Al-Zn binary this phase does not exist and therefore the invariant reaction at 423°C is eliminated from the reaction scheme and liquidus surface in this evaluation. The reaction scheme is shown in Fig. 2. The temperature and the concentration of the ternary eutectic E1 are reported with some uncertainty, 355°C given by [1943Bad] and 364.25°C. In Fig. 2 and Table 2 the values of [1942Wei] are preferred. Liquidus Surface Figure 3 shows the liquidus surface of the Al-LiAl-Zn partial system the diagram given by [1943Bad]. It had to be amended to match with the accepted binary equilibrium diagrams Al-Zn [Mas2] and Al-Li [1989Che]. The ternary eutectic is incorporated using the data given by [1942Wei]. Isothermal Sections Figure 4 shows the amended partial isothermal section Al-LiAl-Zn at 350°C according to [1943Bad]. The section now is compatible with the accepted binary systems Al-Li [1989Che] and Al-Zn [Mas2] and coherent with [1996Dmy] for which the phase extends in the ternary system along the 50 at.% Li. The homogeneity region of the -1 phase follows [1963Che] and may be expressed by the approximate formula LiZn0.5+xAl1.5+x (0 < x < 0.7). There is no experimental evidence for a large width of the -1 field, so the Li content may be accepted as 34-35 at.% as given by [1943Bad]. [1963Che] found the -1 phase in equilibrium with (Zn) neglecting -3. The isothermal section of the system at 193°C according to [1996Dmy] is shown in Fig. 5. No significant solubilities of Al in binary Li-Zn compounds have been detected. Thermodynamics The values H(xC) of liquid Al-Li-Zn alloys were determined at different temperatures along four sections keeping the concentration ratios of two components constant [1997Kim]: (a) Al0.25Zn0.75-Li, (b) Al0.50 Zn0.50-Li, (c) Al0.70Zn0.30 -Li and (d) Al0.75Li0.25-Zn. They are plotted in Figs. 6 and 7. The H values of the ternary liquid alloys can be obtained by adding the H value of the binary boundary systems: H(xA/xB = const., xC) = (1 - xC) H(xA/xB = const.) + ¦ δ Hi(xC)(x /x = const.) A B
i
For the section Al0.25Zn0.75-Li the agreement between the measured and calculated values is within the experimental error. For other concentration section the experimental H values exhibit more negative values compared with the calculated ones. These deviations could be caused by a negative contribution to the enthalpy of mixing due to the presence of additional ternary interactions or additional ternary associates in the melt which have not been taken into account in the model calculation. The presence of additional ternary interaction in the liquid state is supported by the existence of at least three ternary intermetallic phases in this system [1995Pav]. The difference between measured and calculated values of H is shown in Fig. 8 together with the position of the ternary intermetallic phases. Figure 8 shows that the deviation amounts to - 3.5 kJ#mol-1 in the concentration region where the -1 phase exists, which points to additional ternary interaction in this concentration region. In the region of the ternary -2 and -3 phase the deviation is small in comparison to that in the -1 phase region. This indicates that the ternary interactions in these regions are relatively weak and the influence of the ternary -1 phase is predominant for liquid Al-Li-Zn alloys.
Landolt-Börnstein New Series IV/11A3
MSIT ®
126
Al–Li–Zn
Notes on Materials Properties and Applications Al-Li base alloys have received considerable attention as potential lightweight replacements for conventional Al base alloys in aerospace applications. The addition of 1.8-2.1% Li remarkably alter the precipitation behavior of the Al-Cu-Mg-Zn alloys which are the highest strength aluminum alloys [2000Wei]. References [1942Wei] [1943Bad]
[1963Che]
[1986Cas]
[1987Cas]
[1987Che]
[1987Dub]
[1989Aud]
[1989Che] [1993Pav]
[1994Hos]
[1995Pav]
[1995Dmy]
[1996Dmy]
MSIT®
Weisse, E., Blumenthal, A., Hanemann, H., “Results of an Investigation of Eutectic Zinc Alloys” (in German), Z. Metallkd., 34(9), 221 (1942) (Equi. Diagram, Experimental, 9) Badaeva, T.A., Sal’dau, P.Y., “Physico-Chemical Investigation of Alloys of Aluminium with Zinc and Lithium” (in Russian), Zhur. Obshchey Khimii, 13(9/10), 643-660 (1943) (Equi. Diagram, Experimental, 23) Cherkashin, E.E., Kripyakevich, P.I., Oleksiv, G.I., “Crystal Structures of Ternary Compounds in Li-Cu-Al and Li-Zn-Al Systems” (in Russian), Sov. Phys., -Crystallogr., 8(6), 681-685 (1964), translated from Kristallografiya, 8(6), 846-851 (1963) (Crys. Structure, Experimental, 11) Cassada, W.A., Shen, Y., Poon, S.J., Shiflet, G.J., “Mg 32(Zn,Al)49-Type Icosahedral Quasicrystals Formed by Solid-State Reaction and Rapid Solidification”, Phys. Rev. B: Solid State, 34(10), 7413-7416 (1986) (Experimental, 17) Cassada, W.A., Shiflet, G.J., Poon, S.J., “Quasicrystalline Grain Boundary Precipitates in Al Alloys Through Solid-Solid Transformations”, J. Microsc., 146(3), 323-335 (1987) (Experimental, 26) Chen, H.S., Phillips, J.C., Villars, P., Kortan, A.R., Inoue, A., “New Quasicrystals of Alloys Containing s, p and d Elements”, Phys. Rev. B, Cond. Matter, 35B(17), 9326-9329 (1987) (Crys. Structure, Experimental, 18) Dubost, B., Audier, M., Jeanmurt, P., Lang, J.M., Sainfort, P., “Structure of Stable Intermetallic Compounds of the AlLiCu(Mg) and AlLiZn(Cu) Systems”, J. Phys., Colloq., 48C3(9), 497-504 (1987) (Crys. Structure, Experimental, 16) Audier, M., Janot, C., De Boissieu, M., Dubost, B., “Structural Relationships in Intermetallic Compounds of the Al-Li-(Cu, Mg, Zn) System”, Philos. Mag. B, 60(4), 437-466 (1989) (Crys. Structure, Experimental, 34) Chen, S.-W., Jan, C.- H., Lin, J.-C., Chang, Y. A., “Phase Equilibria of the Al-Li Binary System”, Metall. Trans., 20A(11), 2247-2258 (1989) (Equi. Diagram, Experimental, #, 59) Pavlyuk, V.V., “Synthesis and Crystal Chemistry of Lithium Intermetallic Compounds”, Doct. Thesis, Univ. L’viv, 1-35 (1993) (Equi. Diagram, Crys. Structure, Experimental, Review, 49) Hosoda, H., Sato, T., Tezuka, H., Mishima, Y., Kamio, A., “Substitution Behavior of Additional Elements in the L1(2)-Type Al3Li Metastable Phase in Al-Li Alloys” (in Japanese), J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Crys. Structure, Thermodyn., Theory, 26) Pavlyuk, V., Bodak, O., MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16727.1.20, (1995) (Crys. Structure, Equi. Diagram, Assessment, 9) Dmytriv, G.S., “Isothermal Section of the Phase Diagram of the System Li-Zn-Al at 470 K” (in Ukrainian), Lvivski Khimichni Chytannya Naukova-Praktychna Konferentsiya, LDU, 108 (1995) (Equi. Diagram, Experimental, 0) Dmytriv, G.S., “Phase Equilibria and Crystal Structure of Compounds in Mg-Li-Si, Ca-Li-{Si, Ge}, Al-Li-{Si, Ge, Sn}, Zn-Li-{Al, Sn}” (in Ukrainian), Summary of the Thesis for Candidate Science Degree, Lviv, 1-23 (1996) (Crys. Structure, Equi. Diagram, Experimental, 10)
Landolt-Börnstein New Series IV/11A3
Al–Li–Zn [1997Kim]
[1999Pav]
[2000Wei]
[2003Gro]
127
Kim, Y.B., Sommer, F., “Calorimetric Measurement of Liquid Aluminium-Lithium-Zinc Alloys”, Thermochim. Acta, 291, 27-34 (1997) (Equi. Diagram, Thermodyn., Experimental, 16) Pavlyuk, V.V., Dmytriv, G.S., Bodak, O.I., Stepien-Damm, J., “New Variant of the Structure of the Li1+xZn 0.5-1.5Al1.5-0.5 Intermetallic Compound”, Materials Structure, 6(2), 146-148 (1999) (Crys. Structure, Experimental, 4) Wei, B.C., Chen, C.Q., Huang, Z., Zhang, Y.G., “Aging Behavior of Li Containing Al-Zn-Mg-Cu Alloys”, Mat. Sci. Eng. A, 280(1), 161-167 (2000) (Mechan. Prop., Experimental, 9) Gröbner, J., “Al-Li (Aluminium-Lithium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.13517.1.20, (2003) (Equi. Diagram, Crys. Structure, Assessment, 29)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Li) < 180.6
cI2 Im3m W
a = 351.0
pure Li at 25°C [V-C2]
(Zn) < 419.58
hP2 P63/mmc Mg
a = 266.50 c = 494.70
at 25°C [Mas2]
(Al) < 660.45
cF4 Fm3m Cu
a = 404.96
pure Al at 25°C [Mas2] Dissolves up to 15 at.% Li
, Li9Al4 < 347 - 275
mC26 C2/m Li9Al4
a = 1915.51 b = 542.88 c = 449.88 = 107.671°
[2003Gro]
´, Li9Al4 < 275
?
?
[Mas2]
Li3Al2 < 520
hR15 R3m Li3Al2
a = 450.8 c = 1426
[2003Gro] 60 to 61 at.% Li [Mas2]
, LiAl < 700
cF16 Fd3m NaTl
a = 637
at 50 at.% Li [2003Gro] 45 to 55 at.% Li [Mas2]
´, LiAl3 < 190 - ~120
cP4 Pm3m Cu3Au
a = 403.8
Metastable [2003Gro]
LiZn 4 < 245
hP2 P63/mmc Mg
a = 278.8 c = 439.4
[V-C2], [Mas2]
LiZn4 481 - 65
hP2 P63/mmc
-
[Mas2]
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Li–Zn
128 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
Li2Zn5 < 268
hP*
a = 437.0 c = 251.5
[V-C2], [Mas2]
Li2Zn5 502 - 168
-
-
[Mas2]
LiZn2 < 93
-
-
[Mas2]
Li2Zn3 < 174
cP5
a = 427
[V-C2], [Mas2]
Li2Zn3 520 - 160
-
-
[Mas2]
LiZn < 177
cF16 Fd3m NaTl
a = 623.2
[V-C2], [Mas2]
cI160 * -1, Li1+JZn0.5-1. 3Al1.5-0.7 Im3 LiCuSi
a = 1401.7 0.3 to [1999Pav] a = 1390.4 0.3 single crystal data
* -2, LiZn0.6-0.8 Al0.4-0.2
cF16 Fd3m NaTl
a = 625.7 to a = 621.3
[1996Dmy]
* -3, LiZn3Al < 490
-
-
[1943Bad], [1996Dmy] not found by [1963Che], [1996Dmy]
a = 1391 c = 8205 a = 1390 c = 8245
[1987Dub]
* -4, Li3ZnAl5 P42/mmc
sample composition Li0.33 Zn 0.11Al0.56 [1989Aud]
Table 2: Invariant Equilibria Reaction L + -1 + (Al)
T [°C] 452
Type U1
Phase
Composition (at.%) Al
Li
Zn
(Al)
33.2 < 41.5 < 35 < 86
17.5 39.5 35 7
49.3 19.0> 30 > 7>
L
-1 L + -1 -3 + (Al)
368
U2
L -1 -3 (Al)
15.1 < 33.3 < 20 < 88
9.3 33.3 20 4
75.6 33.3 > 60 > 8>
L (Al) + (Zn) + -3
355 a)
E1
L (Al) (Zn) -3
13.0 < 88.0 < 3.0 < 16.8
8.2 2.0 2.0 16.8
78.8 10.0 > 95.0 > 66.4 >
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Li–Zn T [°C]
Reaction
Type
129 Phase
Composition (at.%) Al
Li
Zn
L + -1
580
p1
L -1
32.2 < 41 < 35
32.2 41 35
35.6 18.0 > 30 >
L + -1 -3
490
p2
L -1 -3
18.6 < 33.3 < 20
18.6 33.3 20
62.8 33.3 > 60 >
L -3 + (Zn)
369
e3
L -3 (Zn)
11 < 16.8 < 2.5
11 16.8 2.5
78 66.4 > 95.0 >
Note: values in brackets < > are estimated. a)
Value given by [1943Bad], 364°C after [1942Wei].
Fig. 1: Al-Li-Zn. The pseudobinary system Zn - LiAl
700°C
700
L
Temperature, °C
600
580°C
500
490°C
β 419.58°C 400
(Zn)
369°C
τ1
τ3 300
Zn
10
20
30
Al, at.%
Landolt-Börnstein New Series IV/11A3
40
Li 50.00 Zn 0.00 Al 50.00
MSIT ®
MSIT®
277 e4 (Al)´´ (Al)´ + (Zn)
381 e2 l (Al) + (Zn)
Al-Zn
Fig. 2: Al-Li-Zn. Reaction scheme
600 e1 l (Al) + β
Al-LiAl
452
355
τ1+τ3+(Al)
275
(Al)´´+(Zn)+τ3
E2
U2
L+τ1+τ3
(Al)´´ (Al)´+(Zn)+τ3
(Al)+(Zn)+τ3
E1
L + τ1 τ3 + (Al)
β + τ1 + (Al)
U1
L (Al) + (Zn) + τ3
L+(Al)+τ3
368
L+(Al)+τ1
L + β τ1 + (Al)
Al-LiAl-Zn
369 e3 L (Zn) + τ3
490 p2 L + τ1 τ3
580 p1 L + β τ1
LiAl-Zn
130 Al–Li–Zn
Landolt-Börnstein New Series IV/11A3
Al–Li–Zn
131
Al
Data / Grid: at.%
Fig. 3: Al-Li-Zn. Partial liquidus surface
Axes: at.%
(Al) 20
e1
550
40
600
700
60
650
LiAl
80
600
500
β
60
p1
40
470
U1
450
τ1
420
430
80
U2
p2
20
Li
40
τ3
60
e2 e3
400 (Zn)
380
80
Al Fig. 4: Al-Li-Zn. Partial isothermal section at 350°C
20
E1
Zn
Data / Grid: at.% Axes: at.%
(Al)´ 20
(Al)´+ τ
1
80
40
β +(Al)´
60
(Al)´´
β
τ1
60
(Al)´+τ 1+τ 3
40
n) (Z τ 3+ ´+ l)´ (A
80
20
τ3 (Zn)
Li
Landolt-Börnstein New Series IV/11A3
20
40
60
80
Zn
MSIT ®
Al–Li–Zn
132
Al Fig. 5: Al-Li-Zn. The isothermal section at 193°C
Data / Grid: at.% Axes: at.%
(Al)
20
80
40
60
β Li3Al2
60
40
τ1
δ´ τ3
80
τ2
20
Li
Fig. 6: Al-Li-Zn. Experimental enthalpy of mixing for ternary undercooled liquid alloys
40
60
β Li2Zn3
20
80 αLi2Zn5 β Li2Zn5αLiZn4 β LiZn4
(Zn)
Zn
5
∆mixH, kJ·mol-1
0
-5 x=30 (610°C) x=75 (518°C) -10
x=50 (554°C) -15 0
Al 100-x x Zn 0.00 Li
MSIT®
20
60
40
80
100
Li Li, at.%
Landolt-Börnstein New Series IV/11A3
Al–Li–Zn
Fig. 7: Al-Li-Zn. Enthalpy of mixing of (Al0.75Li0.25)1-xZnx ternary liquid and undercooled liquid alloys at 682°C
133
∆mixH, kJ·mol-1
5
0
-5
-10 0
Al Zn Li
40
20
75.00 00.00 25.00
60
100
80
Li
Li, at.%
Al Fig. 8: Al-Li-Zn. Difference (in kJ#mol-1) between the experimental and the calculated enthalpy of mixing of Al-Li-Zn ternary liquid and undercooled liquid alloys at 682°C using the association model
Data / Grid: at.% Axes: at.%
experimental 20
calculated
80
40
60
τ2
-1.5 -1.0
9
60
40
τ1 -3.5
τ3
80
Li
Landolt-Börnstein New Series IV/11A3
20
40
60
20
80
Zn
MSIT ®
134
Al–Li–Zr
Aluminium – Lithium – Zirconium Oksana Bodak Literature Data Most investigations on the Al-Li-Zr system concern the metastable phases ´, LiAl3 and ´, ZrAl3. [1985Mak] studied the recrystallization behavior of an 3Li-0.12Zr-Al (mass%) alloy in comparison to that of the binary alloys 2.5Li-Al (mass%), 3Li-Al (mass%) and 0.13Zr-Al (mass%). [1984Gay] prepared an 2.34Zr-Al (mass%) alloy by rapid solidification and observed after aging at 190°C a discontinuous precipitation behavior: ZrAl3 was precipitated as aligned rods or as discrete spheres. The ZrAl3/(Al) interface served as a nucleation site for ´, LiAl 3. The resulting “composite” precipitate contained a core of ZrAl3 and an envelope of ´, LiAl3. [1986Gay1] found in the same alloy a ternary phase between LiAl3 and ZrAl3, expressed by the formula (LixZr 1-x)Al3 with 0.45 < x < 0.8, see Table 1. Physical and thermodynamic properties of this phase were investigated by [1986Gay2]. The metastable phase (LixZr1-x)Al3 is also given in Table 1 because of its technical importance [1986Gay1, 1986Gay2, 1986Sak]. In an alloy 3Zr-Al (mass%), [1986Sak] observed the ´ and ´ phases as distinct phases by a time-of-flight atom-probe field-ion microscopy (ToF atom-probe FIM). The nucleation of ´ on ´ as a substrate was studied theoretically by [1987Tos]. The precipitation of ´ in several ternary and quaternary alloys was reviewed by [1987Flo]. By adiabatic scanning calorimetry [1988Eun] examined precipitation and dissolution reactions. Partial vapor pressures of Li over binary and ternary aluminium melts at 927°C were calculated using an interaction parameter for Zr as a third element [1986Lee]. [1989Sau] calculated phase diagrams for stable as well as for metastable phase equilibria in the Al-Li-Zr system. The effects of mechanical alloying, a low temperature isothermal processing method, and the effect of ternary addition of lithium on the phase stability of the ZrAl3 phase with metastable cubic L1 2 structure were studied in [1991Des]. At 750°C it was found that adding lithium increases the stability of the L12 phase. The literature until 1989 was compiled and critically reviewed by [1995Pav]. The results of an investigation of the isothermal section of Al-Li-Zr at 197°C and data of the crystal structure of the compounds are reported in [2002Zat]. The alloys were prepared by arc melting in purified argon atmosphere under a pressure of ~1.01#105 Pa from a mixture of the pure metals (Zr of 99.98% mass purity, Li of 99.0 mass% purity, and Al of 99.99 mass% purity). The alloy compositions were checked by weight comparison of the initial mixtures and the alloys. The alloys were annealed at 197°C for 400 h in tantalum containers in evacuated quartz ampoules and quenched in cold water. The X-ray powder method was used for the phase analysis and structural investigation. Binary Systems The Al-Li system reported by [2003Gro] is accepted. The Al-Zr phase diagram presented by [2003Sch] shows more likely features than the those given in the diagram by [Mas2], in which all the liquidus lines are drawn tentatively. The Li-Zr system is accepted from [Mas2]. The extremely small solubility of Zr in liquid Li was calculated by [1989Sau]. In the range of 7.5 at.% Li, the stable solid phases are (Al) and , LiAl. However, a metastable LiAl3 occurs and creates order hardening in the alloys. The metastable solvus (Al)/LiAl3 has been experimentally determined by [1998Nob]; Zr additions up to 0.05 at.% do not affect the position of the metastable boundary. Solid Phases The ternary compounds ZrLi2Al has a narrow range of homogeneity and Zr5-xLix+yAl3 (x = 0.2 - 1.0, y = 0 - 1) exhibits a relatively wide homogeneity range, see Table 1.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Li–Zr
135
Liquidus Surface The calculated liquidus surface in the composition range 0 to 10 at.% Zr and 0 to 50 at.% Li is shown in Fig. 1 [1989Sau]. The ternary invariant reactions at 595.4°C with a composition of the melt of 24.65 at.% Li and 5.1#10 -5 at.% Zr cannot be reproduced in Fig. 1 because of the low Zr content in the melt. Isothermal Sections Partial isothermal sections at 500, 300 and 100°C were given for the composition range 0 to 25 at.% Zr and 0 to 50 at.% Li [1989Sau]. Since they are rather similar, only the section at 300°C is presented in Fig. 2. Beyond this the ordering of the ´, (LixZr1-x)Al3 phase was thermodynamically described [1989Sau]. The isothermal section of the Al-Li-Zr phase diagram at 197°C is shown in Fig. 3. The results show good agreement between the experimental data of [2002Zat] and the calculated part of isothermal section [1989Sau]. The formation process of Zr5-xLix+yAl3 ternary intermetallic is realized by a partial substitution of Li atoms by Zr in the 4(d) position and insertion of lithium atoms in holes at the 2(b): 000 position. The change of the lattice parameters in the Zr5-xLix+yAl3 homogeneity range is presented in Fig. 4 after [2002Zat]. The Zr5Al4 binary compound (Ti5Ga4 structure type) has not been found in the Al-Li-Zr system at 197°C. It is stable in the temperature range from 990 to 1530°C. The Zr5-xLix+yAl3 ternary compound, apparently, is a remainder of the high temperature substitution-limited solid solution of the Zr5Al4 binary compound or of the substitution- and insertion-limited solid solution of Zr5Al3 (Mn5Si3 structure type). The characteristic feature of the Al-Li-Zr ternary system is the binary immiscibility region of Li-Zr extending up to ~10 at.% of the third component. Limited solid solutions of the binary compounds of the Al-Zr system were observed in the Al-Li-Zr system. Largest solubility of the third component is found in ZrAl3 (5 at.%), ZrAl2 (10 at.%) and Zr2Al3 (15 at.%). For the ZrLixAl3-x solid solution the change of the lattice parameters vs Li-concentration is presented in Fig. 5 after [2002Zat]. Notes on Materials Properties and Applications In cast aluminium alloys Zr is typically added to achieve grain refinements and to inhibit the recrystallization of wrought structures. This behavior is associated with the formation of coherent ZrAl3 particles of metastable cubic form [1987Flo] which is stabilized by Li [1987Vec]. In addition, Zr is used to impart superplasticity, or improve strength and toughness of rapidly solidified Al-Li alloys. In the ternary system, Zr precipitates in a supersaturated solid solution via a normal nucleation and growth mechanism as coherent spherical or filamentary particles, depending on the heat treatment as (Li,Zr)Al3, metastable, Cu3Au type phase [1989Gay, 1994Hos]. References [1984Gay] [1985Mak]
[1986Gay1] [1986Gay2]
[1986Lee]
Landolt-Börnstein New Series IV/11A3
Gayle, F.W., Vander Sande, J.B., “’Composite’ Precipitates in an Al-Li-Zr Alloy”, Scr. Metall., 18, 473-478 (1984) (Experimental, 13) Makin, P. L., Stobbs, W.M., “Comparison of the Recrystallization Behaviour of an Al-Li-Zr Alloy with Related Binary Systems”, The Institute of Metals, London, Accession Number: 86(8), 72-312; 392-401 (1986) (Experimental, 11) Gayle, F.W., Vander Sande, J.B., “Al3Li Precipitate Modification in an Al-Li-Zr Alloy”, ASTM, Proc. Pennsylvania, 1984, 137-152 (Publ. 1986) (Crys. Structure, Experimental, 16) Gayle, F.W., Vander Sande, J.B., “Al3(Li, Zr), or ´ Phase in Al-Li-Zr System”, The Institute of Metals, London, accession Number, 86(8), 72-312, 376-384 (1986) (Crys. Structure, Experimental, 17) Lee, J.J., Sommer, F., “Thermodynamic Properties of Lithium in Liquid Aluminum Alloys” (in Korean), Taehan Kumsok Hakhoechi, 24(10), 1185-1189 (1986) (Thermodyn., Experimental, 19)
MSIT ®
136 [1986Sak]
[1987Flo] [1987Tos] [1987Vec]
[1988Eun]
[1989Che]
[1989Gay] [1989Sau] [1991Des]
[1994Hos]
[1995Pav]
[1998Nob]
[2002Zat]
[2003Gro]
[2003Sch]
MSIT®
Al–Li–Zr Sakurai, T., Kobayashi, A., Hasegawa, J., Sakai, A., Pickering, H.W., “Atomistic Study of Metastable Phases in Al - 3 wt.% Zr Alloy”, Scr. Metall., 20, 1131-1136 (1986) (Crys. Structure, Experimental, 18) Flower, H.M., Gregson, P.J., “Solid State Phase Transformations in Aluminium Alloys Containing Lithium”, Mater. Sci. Technol., 3(2), 81-90 (1987) (Review, 116) Tosten, M. H., Galbraith, J. M., Howell, P. R., “Nucleation of ´ (Al3Zr) in Al-Li-Zr and Al-Li-Cu-Zr Alloys”, J. Mater. Sci. Lett., 6(2), 51-53 (1987) (Experimental, 10) Vecchio, K.S., Williams, D.B., “Convergent Beam Electron Diffraction Study of Al3Zr in Al-Zr and Al-Li-Zr Alloys”, Acta Metall., 35(12), 2959-2970 (1987) (Crys. Structure, Experimental, 19) Eun, I.-S., Woo, K.-D., Cho, H.K., “The Formation of Precursor Phase During Precipitation in Al-Li-Zr Alloy” (in Korean), J. Korean Inst. Met., 26(11), 1007-1012 (1988) (Thermodyn., Experimental, 10) Chen, S.W., Tan, C.-H., Lin, T.-C., Chang, Y.A., “Phase Equilibria of the Al-Li Binary System”, Metall. Trans. A, 20A(11), 2247-2258 (1989) (Equi. Diagram, Experimental, Thermodyn., #, 59) Gayle, F.W., Vandersande, B., “Phase Transformations in the Al-Li-Zr System”, Acta Metall., 37(4), 1033-1046 (1989) (Crys. Structure, Experimental, Thermodyn., 28) Saunders, N., “Calculated Stable and Metastable Phase Equilibria in Al-Li-Zr Alloys”, Z. Metallkd., 80(12), 894-903 (1989) (Equi. Diagram, Thermodyn., Theory, #, *, 78) Desch, P.B., Schwarz, R.B., Nash, P., “Formation of Metastable L12 Phases in Al3Zr and Al-12.5% X-25 % Zr (X = Li, Cr, Fe, Ni, Cu)”, J. Less-Common Met., 168, 69-80 (1991) (Crys. Structure, Experimental, 25) Hosoda, H., Sato, T., Tezuka, H., Mishima, Y., Kamio, A., “Substitution Behavior of Additional Elements in the L1 2-Type Al3Li Metastable Phase in Al-Li Alloys” (in Japanese), J. Jpn. Inst. Met., 58(8), 865-871 (1994) (Crys. Structure, Theory, Thermodyn., 26) Pavlyuk, V., Bodak, O., “Aluminium-Lithium-Zirconium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.14883.1.20, (1995) (Crys. Structure, Equi. Diagram, Assessment, 15) Noble, B., Bray, S.E., “On the (A1)/ ´(Al3Li) Metastable Solvus in Aluminium-Lithium Alloys”, Acta Mater., 46(17), 6163-6171 (1998) (Calculation, Experimental, Phys. Prop., Thermodyn., 41) Zatorska, G.M., Pavlyuk, V.V., Davydov, V.M., “Phase Equilibria and Crystal Structure of Compounds in the Zr-Li-Al System at 470 K”, J. Alloys Compd., 333, 138-142 (2002) (Equi. Diagram, Crys. Structure, Experimental, #, 11) Groebner, J., “Al-Li (Aluminium-Lithium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.13517.1.20, (2003) (Crys. Structure, Equi. Diagram, Assessment, 21) Schuster, J.C., “Al-Zr (Aluminium-Zirconium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services, GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram, Assessment, 103)
Landolt-Börnstein New Series IV/11A3
Al–Li–Zr
137
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Li) < 180.6 (Al) < 660.45 (Zr) 1855 - 863 (Zr) < 863 Li9Al4 347 - 275
Li9Al4 ( ´) < 275 Li3Al2 ( ) < 520 LiAl () < 700 LiAl3 (´) 400 Zr3Al < 1019 Zr2Al < 1215 Zr5Al3 (r) 1000 Zr5Al3 (h) 1400 - 1000 Zr3Al2 < 1480
Landolt-Börnstein New Series IV/11A3
Pearson Symbol/ Space Group/ Prototype cI2 Im3m W cF4 Fm3m Cu cI2 Im3m W hP2 P63/mmc Mg mC26 C2/m Li9Al4 ? hR15 R3m Li3Al2 cF16 Fd3m NaTl cP4 Pm3m Cu3Au cP4 Pm3m Cu3Au hP6 P63/mmc Ni2In hP16 P63/mcm Mn5Si3 tI32 I4/mcm W5Si3 tP20 P42/mnm Zr3Al2
Lattice Parameters Comments/References [pm] a = 351.0
pure Li at 25°C [V-C2]
a = 404.96
a = 360.90
pure Al at 25°C [Mas2] dissolves up to 15 at.% Li [Mas2]
a = 323.16 c = 514.75
[2003Sch] dissolves up to 8.3 at.% Al at 910°C
a = 1915.51 b = 542.88 c = 449.88 = 107.671° ?
[2003Gro]
a = 450.8 c = 1426
[2003Gro] 60 to 61 at.% Li [Mas2] at 50 at.% Li [2003Gro] 45 to 55 at.% Li [Mas2]
a = 637
[Mas2]
a = 403.8
Metastable [1989Che, 2003Gro]
a = 439.17
[V-C2, Mas2]
a = 489.39 c = 592.83
[2003Sch]
a = 818.4 c = 570.2
[2003Sch]
a = 1104.4 c = 539.1
[2003Sch]
a = 763.0 c = 699.8
[2003Sch]
MSIT ®
Al–Li–Zr
138 Phase/ Temperature Range [°C] Zr4Al3 < 1030 Zr5Al4 1550 - 1000 ZrAl < 1275 Zr2Al3 < 1590 ZrAl2 < 1660 ZrLixAl3-x ZrAl3 < 1580 ZrAl3
* -1, Li2ZrAl
* -2, Lix+yZr5-xAl3
MSIT®
Pearson Symbol/ Space Group/ Prototype hP7 P6/mmm Zr4Al3 hP18 P63/mcm Ti5Ga4 oC8 Cmcm CrB oF40 Fdd2 Zr2Al3 hP12 P63/mmc MgZn2 tI16 I4/mmm ZrAl3
cP4 Pm3m Cu3Au cF12 F43m CuHg2Ti hP18 P63/mcm Ti5Ga4
Lattice Parameters Comments/References [pm] a = 543.0 c = 539.0
[2003Sch]
a = 844.8 c = 580.5
[2003Sch]
a = 335.3 b = 1086.6 c = 426.6 a = 960.1 b = 1390.6 c = 557.4 a = 528.24 c = 874.82
[2003Sch]
a = 400.9 c = 1728.2
x = 0.2 (Li0.2Al2.8Zr) [2002Zat]
a = 401.4 c = 1727.7 a = 408
x = 0 (ZrAl3)
a = 663.3
[2002Zat]
a = 813.36 c = 570.29
Li0.2Zr 4.8Al3 (x = 0.2, y = 0)
a = 817.57 c = 569.09
LiZr4Al3 (x = 1, y = 0) [2002Zat]
[2003Sch]
[2003Sch]
Metastable, stabilized by Li [1987Vec, 1989Gay]
Landolt-Börnstein New Series IV/11A3
Al–Li–Zr
139
10
Fig. 1: Al-Li-Zr. Calculated partial liquidus surface
8
ZrAl3
15
Zr, at.%
6
00
4
14
00
13
00
2
120 0 110 0 900
0
Al
0
800
10
20
30
40
50
Li, at.%
Al Fig. 2: Al-Li-Zr. Calculated partial isothermal section at 300°C
Data / Grid: at.% Axes: at.%
(Al)
10
90
20
80
(Al)+LiAl+ZrAl3
ZrAl3
30
70
40
60
LiAl Li Zr Al Landolt-Börnstein New Series IV/11A3
50.00 0.00 50.00
10
20
30
40
Li Zr Al
0.00 50.00 50.00
MSIT ®
Al–Li–Zr
140
Al Fig. 3: Al-Li-Zr. Isothermal section at 197°C
Data / Grid: at.% Axes: at.%
(Al)
20
80
ZrAl3 ZrAl2 40
60
LiAl Li3Al2
Zr2Al3 ZrAl Zr4Al3
60
Li9Al4
40
τ2
τ1
Zr3Al2 Zr2Al Zr3Al
80
20
L
(Zr) 20
Li
60
80
Zr
329.0
V, pm3 ×10-6
Fig. 4: Al-Li-Zr. Change of lattice parameters for Zr5-xLix+yAl3
40
328.0 327.0
V
326.0
Lattice parameter, pm
570.2 569.8 569.4
c
569.0 818 817
a
816 815 814 813 0
10
20
30
Li, at.%
MSIT®
Landolt-Börnstein New Series IV/11A3
Fig. 5: Al-Li-Zr. Change of lattice parameters for the ZrLixAl3-x solid solution
V, pm3 ×10-6
Al–Li–Zr
141
277.80
V 277.76 277.72 277.68
Lattice parameter, pm
1728.2 1728.0
c
1727.8 1727.6 401.4 401.2
a 401.0 400.8 0
5
10
15
Li, at.%
Landolt-Börnstein New Series IV/11A3
MSIT ®
142
Al–Mg–Mn
Aluminium – Magnesium – Manganese Qingsheng Ran, updated by Joachim Gröbner Literature Data The system has been mainly investigated on the Al-Mg side. [1952Han, 1980Bra] reviewed information about the system, but the scopes covered by each are quite limited. The Al rich corner: with about 150 alloys [1938Lee] studied the system in the range 0-35.5 mass% Mg and 0-12 mass% Mn, by means of metallography and thermal analysis and in some cases supplemented by annealing experiments and X-ray crystallography. Alloys were prepared from aluminium of 99.991% purity, magnesium of 99.996% purity, and an aluminium-manganese hardener containing 13-25 mass% Mn prepared directly from aluminium and dehydrated MnCl2. The results were drawn mainly from microscopical observation of cast alloys. A short version of this work was given by [1938Han]. By X-ray diffraction (Debye-Scherrer and rotating crystal methods) of a 16.28 mass% Mg and 4.26 mass% Mn alloy, [1938Hof] reported the solid phases formed by a metastable eutectic in the Al corner to be Al solid solution, Mg2Al3 and MnAl 4. With about 100 alloys, prepared from Mg, 99.99% pure Al and a high purity Al-Mn alloy, [1940Fah] investigated the joint solubilities of Mg and Mn in (Al) at 500 to 650°C by electrical resistance measurements. To obtain more detailed information about the liquidus surface of the Al corner, [1943But] studied 20 alloys with up to 5 mass% Mg and 2 mass% Mn. Aluminium of super-purity grade and aluminium-manganese master alloys in the same purity degree and magnesium of 99.95% purity were used for preparing the alloys for the determination of cooling curves. Nine alloys were studied by [1943Lit] for determining the effect of Mg on the solubility of Mn at 500°C by microstructure observation. The materials and experimental procedure used by [1943Lit] were the same as those of [1943But]. [1943Lit] stated in addition that no new phases appear in the Al-5Mg-2Mn (mass%) range at 400°C. [1943Mon] drew equilibrium diagrams for the Al corner from data by [1938Lee] and own values, but did not give any details on the results and procedure of their own experiments. [1945But] continued the work of the constitution of the Al corner and determined the solidus isotherms by observation of incipient melting and microstructure. An isothermal section at 630°C was also presented. Considering the limited composition range or nonequilibrium condition, [1948Wak] carried out microstructural observation of 45 samples for determining the phase relationships in the region of aluminium with up to 40 mass% Mg and 25 mass% Mn at 400°C. High purity aluminium and magnesium metals and aluminium-manganese master alloys were melted, cast and annealed at 400°C. In some cases, slowly-cooled alloys were also examined. X-ray diffraction was used for identifying the phases. Using 99.99% Al, 99.9% Mg and 99.9% Mn [1973Ohn1] prepared 40 alloys. After melting, the samples were cast and then annealed at 450°C for 20 days and at 400°C for 40 days, respectively. The quenched samples were investigated by metallography and X-ray diffraction analysis. Isothermal sections of the aluminium side with up to 15 mass% Mg and 6 mass% Mn at 450 and 400°C were established. The structure of a ternary phase was determined. In a work primarily on the quaternary system Al-Cr-Mg-Mn [1973Ohn2] 6 alloys were examined for studying the constitution of the Al corner of the Al-Mg-Mn system at 550°C. Most results of the above investigations are consistent with each other. However, the isothermal sections at 435 and 400°C, established by [1938Lee, 1948Wak, 1973Ohn1], respectively, are inconsistent. The reason might be that the equilibrium state was not achieved by [1938Lee]. The liquidus surface from [1938Lee] is accepted, but more experiment in this region is necessary. The Mg rich corner: [1938Ima] investigated the Mg-35Al-6Mn (mass%) region with 17 samples. The starting materials were 99.8% pure Mg and Al, metallic Mn, an Al-19.8%Mn master alloy and MnCl2 for preparing ternary alloys. Thermal analysis and microscopic examination were used. [1944Bee] determined several solubility curves of Mn and Al in Mg at different temperatures. [1948Age] studied the Mg corner with up to 40 mass% Al and 10 mass% Mn by thermal analysis, metallography and X-ray diffraction. A liquidus surface and some invariant equilibria are presented, but these do not agree with [1938Ima]. [1957Mir] prepared samples from metallic Mg (~99.9%), Al (~99.99%) and electrolytic Mn, from which
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Mn
143
first Mg-Mn and Al-Mn master alloys were made. The liquidus surface of the Al-Mg side with up to ~65 mass% Al was constructed from microstructural analysis on about 20 cast alloys. 75 alloys were examined by metallography and microhardness for determining phase equilibria in the Mg corner at 200 to 400°C. The samples were heat treated in evacuated silica ampules for 9 to 55 days and water quenched. Results were given in two partial isothermal sections. [1986Obe, 1988Sim1, 1988Sim2] reported phases existing in equilibrium with liquid at temperatures between 660 and 760°C in Mg alloys with up to 10 mass% Al and 1.5 mass% Mn. MnCl2 was added to the melts at 780°C to saturate the alloys with Mn. After thorough stirring, the melts were held for 1 to 2 h at 750, 710 and 670°C, respectively. The samples were made either by ordinary casting or rapid quench against a spinning, water-cooled wheel and examined by microstructural, X-ray diffraction and microprobe analysis. Phases in equilibrium with liquid and the single phase region of melt for the temperatures 750, 710 and 670°C were determined. [1992Ars] prepared samples in the Al rich corner with constant 10 mass% Mg by rapid quenching in water. They report a calculated metastable vertical section which is not in equilibria with the ternary phase T. Binary Systems The binary system Al-Mg was updated by [2003Luk]. This version is accepted. The Al-Mn system is accepted from [2003Pis] and Mg-Mn is taken from [Mas]. Solid Phases [1948Wak] revealed a ternary phase T by metallographic observation and X-ray diffraction. The composition of this phase is near MnMg2Al10. The structure was determined by [1973Ohn1, 1994Fun] who suggested the composition of the phase to be Mn2Mg3Al18. A Mn rich phase X was proposed by [1948Age] without giving details on structure or composition. It is quite probably the same phase as X in [1957Mir] who concluded that X should be an Al-Mn binary phase. The ternary phase T and the binary solid phases present in the compiled phase diagrams are listed in Table 1. Invariant Equilibria Some four-phase equilibria were reported. The reactions listed in Table 2 are based on [1938Lee] (the first three) and [1948Age, 1938Ima] (the last two). It should be noted that all these reactions are not certain. According to [1948Wak, 1973Ohn1] the reactions given by [1938Lee] might be metastable. The region of the primary solidification of the ternary compound T reported by [1957Mir] makes the reactions according to [1938Lee] also doubtful. These reactions therefore need further investigation. Liquidus Surface A liquidus surface projection on the Al-Mg rich side is constructed using data from different investigations, Fig. 1. Because of the different opinions on some reactions (see section Invariant Equilibria) and the incomplete determination of other reactions, this liquidus projection has to be considered as tentative. Isothermal Sections Isothermal sections of the Al corner at 630°C [1945But] and 400°C [1948Wak, 1973Ohn1] are given in Figs. 2 and 3, respectively. An isothermal section at a temperature just after the end of crystallization was proposed by [1938Lee], but is contradictory to [1948Wak] and [1973Ohn1], who studied the topic more carefully. The joint solubility of Mg and Mn in solid (Al) is given in Fig. 4; the data are mainly from [1940Fah]. Isothermal sections of the Mg corner at 400°C and 200°C are plotted in Figs. 5 and 6, respectively.
Landolt-Börnstein New Series IV/11A3
MSIT ®
144
Al–Mg–Mn
References [1938Han]
[1938Hof] [1938Ima] [1938Lee] [1940Fah] [1943But]
[1943Lit]
[1943Mon] [1944Bee]
[1945But]
[1948Age]
[1948Wak]
[1952Han] [1957Mir]
[1973Ohn1]
[1973Ohn2]
[1980Bra] [1986Obe]
MSIT®
Hanemann, H., Schrader, A., “On Some Ternary Systems of Aluminium, I Aluminium Iron - Magnesium, Aluminium-Magnesium-Manganese, Aluminium-Manganese-Silicon” (in German), Z. Metallkd., 30, 383-386 (1938) (Equi. Diagram, Experimental, #, 11) Hofmann, W., “X-Ray Methods on Investigation of Aluminium Alloys” (in German), Aluminium, 865-872 (1938) (Crys. Structure, Experimental, 19) Imaki, A., “On the Equilibrium Diagram of Mg-Al-Mn Alloy System” (in Japanese), Trans. Min. Met. Alumi. Assoc., 9, 665-668 (1938) (Equi. Diagram, Experimental, 1) Leemann, W.G., “The Ternary System Aluminium-Magnesium-Manganese” (in German), Aluminium Arch., 9, 6-17 (1938) (Equi. Diagram, Experimental, 7) Fahrenhorst, E., Hofman, W., “The Solubility of Manganese in Aluminium with up to 2 % Mg” (in German), Metallwirtschaft, 19, 891-893 (1940) (Equi. Diagram, Experimental, 3) Butchers, E., Raynor, G.V., Hume-Rothery, W., “The Constitution of Magnesium-Manganese-Zinc-Aluminium Alloys in the Range 0-5 % Magnesium, 0-2 % Manganese, 0-8 % Zinc, I-The Liquidus”, J. Inst. Met., 69, 209-228 (1943) (Equi. Diagram, Experimental, 9) Little, A.T., Raynor, G.V., Hume-Rothery, W., “The Constitution of Magnesium Manganese - Zinc - Aluminium Alloys in the Range 0-5 % Magnesium, 0-2 % Manganese and 0-8 % Zinc, III-The 500C and 400C Isothermals”, J. Inst. Met., 69, 423-440 (1943) (Equi. Diagram, Experimental, 8) Mondolfo, L.F., “Metallography of Aluminium Alloys”, John Wiley and Sons, Inc., New York, 100-101 (1943) (Equi. Diagram, Review, 1) Beerwald, A., “On the Solubility of Iron and Manganese in Magnesium and in Magnesium-Aluminium Alloys” (in German), Metallwirtschaft, 23, 404-407 (1944) (Equi. Diagram, Experimental, 10) Butchers, E., Hume-Rothery, W., “On the Constitution of Aluminium - Magnesium Manganese - Zinc Alloys: The Solidus”, J. Inst. Met., 71, 291-311 (1945) (Equi. Diagram, Experimental, #, 8) Ageev, N.V., Kornilov, I.I., Khlapova, A.N., “Magnesium-Rich Alloy of the System Magnesium-Aluminium-Manganese” (in Russian), Izv. Inst. Fiz.-Khim. Anal., Inst. Obshcheii Neorg. Khim., Akad. Nauk SSSR, 14, 130-143 (1948) (Equi. Diagram, Experimental, #, 11) Wakeman, D.W., Raynor, G.V., “The Constitution of Aluminium-Manganese-Magnesium and Aluminium-Manganese-Silver Alloys, with Special Reference to Ternary Compound Formation”, J. Inst. Met., 75, 131-150 (1948) (Equi. Diagram, Experimental, *, 27) Hanemann, H., Schrader, A., “Ternary Alloys of Aluminium” (in German), Verlag Stahleisen m.b.H., Dusseldorf, 116-120 (1952) (Equi. Diagram, Review, 3) Mirgalovskaya, M.S., Matkova, L.N., Komova, E.M., “The System Mg-Al-Mn” (in Russian), Trudy Inst. Met. Im. A.A. Baikova, Akad. Nauk, 2, 139-148 (1957) (Equi. Diagram, Experimental, #, 3) Ohnishi, T., Nakatani, Y., Shimizu, K., “Phase Diagrams and Ternary Compounds of the Al-Mg-Cr and the Al-Mg-Mn Systems in Al-Rich Side” (in Japanese), Light Metals Tokyo, 23, 202-209 (1973) (Crys. Structure, Equi. Diagram, Experimental, *, 16) Ohnishi, T., Nakatani, Y., Shimizu, K., “Phase Diagram in the Al-Rich Side of the Al-Mg-Mn-Cr Quarternary System” (in Japanese), Light Metals Tokyo, 23, 437-443 (1973) (Equi. Diagram, Experimental, 2) Brandes, E.A., Flint, R.F., “Manganese Phase Diagrams”, Manganese Center, 17 Ave. Hoche, 75008 Paris, France, 82 (1980) (Equi. Diagram, Review, 2) Oberlaender, B.C., Simensen, C.J., Svalestuen, J., Thorvaldsen, A., “Phase Diagram of Liquid Magnesium - Aluminium - Manganese Alloys”, Magnesium Technology, Pros. Conf., London, 133-137 (1986) (Experimental, 3) Landolt-Börnstein New Series IV/11A3
Al–Mg–Mn [1988Sim1]
[1988Sim2]
[1992Ars]
[1994Fun] [2003Luk]
[2003Pis]
145
Simensen, C.J., Oberländer, B.C., Svalestuen, J., Thorvaldsen, A., “Determination of the Equilibrium Phases in Molten Mg - 4 wt.% Al-Mn Alloys”, Z. Metallkd., 79, 537-540 (1988) (Experimental, 6) Simensen, C.J., Oberländer, B.C., Svalestuen, J., Thorvaldsen, A., “The Phase Diagram for Magnesium - Aluminium - Manganese above 650°C”, Z. Metallkd., 79, 696-699 (1988) (Experimental, 10) Arsenov, A.A., Goutan, D., Zolotarevskii, V.S., Kuznetsov, G.M., Lugin, D.V., “Study of Decomposition of the (Al)-Solid Solution Heating for Quenching of Cast Alloys Al-10% Mg and Al-6% Zn-15% Mg-1% Cu Containing Manganese” (in Russian), Metally, 6, 80-83 (1992) (Experimental, 5) Fun, H.-K., Lin, H.-C., Lee, T.-J., Yipp, B.-C., “T-Phase Al18Mg3Mn2”, Acta Crystallogr., C50, 661-663 (1994) (Crys. Structure, 5) Lukas, H.-L., Lebrun, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49) Pisch, A., “Al-Mn (Aluminium-Manganese)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 40)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Al) < 660.5
cF4 Fm3m Cu
a = 404.88
[V-C], pure 23°C
(Mg) < 650
hP2 P63/mmc Mg
a = 320.89 c = 521.01
[V-C], pure
(Mn) < 1079
cP20 P4132 Mn
a = 631.5
pure Mn, [V-C]
(Mn) < 710
cI58 I43m Mn
a = 891.39
pure Mn, [V-C]
, Mg 2Al3 452
cF1168 Fd3m Mg2Al3
a = 2816 to 2824
60-62 at.% Al [2003Luk] 1168 atoms on 1704 sites per unit cell [2003Luk]
, Mg17Al12 < 458
cI58 I43m Mn
a = 1054.38
at 41.4 at.% Al [V-C2] 39.5 to 51.5 at.% Al [2003Luk]
J, Mg23Al30 410 - 250
hR159 R3 Mn44Si9
a = 1282.54 c = 2174.78
54.5 to 56.5 at.% Al [2003Luk]
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Mg–Mn
146 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
MnAl6 < 705
oC28 Cmcm MnAl6
a = 754.5 0.2 b = 649.0 0.3 c = 868.1 0.2
[2003Pis]
, MnAl4 < 923
hexagonal
-
[Mas]
Mn4Al11(r) 916
aP30 P1 Mn4Al11
[2003Pis] a = 509.5 0.4 b = 887.9 0.8 c = 505.1 0.4 = 89.35 0.04° = 100.47 0.05° = 105.08 0.06°
* T, Mn2Mg3Al18
cF184 Fd3m Cr2Mg3Al18
a = 1452.9 a = 1451.7
[1973Ohn1] [1994Fun]
Table 2: Invariant Equilibria T [°C]
Reaction
Type
Phase
Composition (at.%) Al
Mg
Mn
L + Mn4Al11 (r) +
-
U1
L Mn4Al11(r)
67.7 73.3 81.5 62.0
30.6 0 0 37.5
1.7 26.7 19.5 0.5
L + MnAl6 +
-
U2
L MnAl6
69.3 81.5 85.7 61.5
29.5 0 0 38.0
1.2 19.5 14.3 0.5
L (Al) + + MnAl6
437
E1
L Al MnAl6
70.7 84.5 61.0 85.7
28.3 15.0 38.5 0
1.0 0.5 0.5 14.3
~437
U3
L
30
69.5
0.5
~430
E2
L
34
64.6
1.4
L + (Mg) + (Mn)(?) L + (Mn)(?) + X a)
a)
X is an Al-Mn binary compound [1957Mir]
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Mn
147
Al
Data / Grid: at.%
(Al)
Fig. 1: Al-Mg-Mn. Liquidus surface on the Al-Mg side
Axes: at.%
MnAl6
µ 20
E1
80
U2 U1
β 40
60
P X
60
40
γ E2 U3 80
20
Mg
β Mn
Mg αMn
20
40
60
Al Fig. 2: Al-Mg-Mn. Isothermal section of the Al corner at 630°C [1945But]
80
Mn
Data / Grid: at.% Axes: at.%
(Al)
(Al)+L
(Al)+MnAl6
L (Al)+MnAl6+L 10
MnAl6+L
90
MnAl6
Mg 20.00 Mn 0.00 Al 80.00 Landolt-Börnstein New Series IV/11A3
10
Mg 0.00 Mn 20.00 Al 80.00
MSIT ®
Al–Mg–Mn
148
Al Fig. 3: Al-Mg-Mn. Isothermal section of the Al corner at 400°C
Data / Grid: at.% Axes: at.%
(Al) (Al)+MnAl6
10
(Al)+MnAl6+T 20
90
MnAl6
(Al)+T 80
(Al)+T+β
T
30
70
T+β +ε
β Mg 40.00 Mn 0.00 Al 60.00
10
20
Al Fig. 4: Al-Mg-Mn. Joint solubility of Mg and Mn in solid (Al)
MSIT®
Mg 0.00 Mn 40.00 Al 60.00 Data / Grid: at.% Axes: at.%
400°C 500°C
Mg 5.00 Mn 0.00 Al 95.00
30
550°C 600°C
Mg 0.00 Mn 5.00 Al 95.00 Landolt-Börnstein New Series IV/11A3
Al–Mg–Mn Mg 80.00 0.00 Mn Al 20.00
Fig. 5: Al-Mg-Mn. Isothermal section of the Mg corner at 400°C, X is a Al-Mn binary compound
149
Data / Grid: at.% Axes: at.%
(Mg)+γ (Mg)+γ +X 90
10
(Mg)+X
(Mg)+(β Mn)(?)+X (Mg) (Mg)+(β Mn)(?) 10
Mg
Mg 90.00 0.00 Mn Al 10.00
Fig. 6: Al-Mg-Mn. Isothermal section of the Mg corner at 200°C, X is an Al-Mn binary compound
Mg 80.00 Mn 20.00 0.00 Al Data / Grid: at.% Axes: at.%
(Mg)+γ
(Mg)+γ +X
(Mg)+X (Mg)+(β Mn)(?)+X (Mg) (Mg)+(β Mn)(?)
Mg
Landolt-Börnstein New Series IV/11A3
Mg 90.00 Mn 10.00 0.00 Al
MSIT ®
150
Al–Mg–Ni
Aluminium – Magnesium – Nickel Elena L. Semenova Literature Data The Al-Mg-Ni system has been examined first in 1924. From the results of thermal analysis and metallography [1924Fus] concluded that the Mg2Al3-NiAl3 section is a quasibinary one. In [1934Fus] Fuss presented a projection of the liquidus surface in the Al-Mg2Al3-NiAl area showing the lines of double saturation on it. An essential conclusion was that a ternary eutectic equilibrium does not exist in the shown part of the phase diagram. However, [1943Mon, 1944Cha, 1952Han] reported that the invariant eutectic equilibrium exists and is reached independently of the heat treatment and the compositions of the phases, except of solid solution of magnesium in aluminium. These conclusions were based on experimental data obtained on as-cast, annealed and rapidly quenched alloys; their liquidus projection is essentially different from the one without the eutectic invariant reaction proposed by [1934Fus]. [1968Var] studied the structure of the Al-Mg-Ni alloys containing 1 at.% Ni in as-cast conditions. The intermetallic phases were separated by high temperature centrifuging and identified by X-ray analysis. As a result, the AlNi3 and Al3Ni2 phases were found to coexist in the alloy 1Ni-15Mg-Al (at.%). The assessment by [1993Pri] took into account the works published up to 1991 and deals with the Al-rich part of the Al-Mg-Ni ternary system Al-Mg2Al3-Ni2Al3. Later experimental investigations of the ternary system were mainly motivated by the search for new hydrogen storage materials [1998Ori, 2000Yua, 2000Aiz, 2001Gua]. From these studies information on new ternary phases was obtained. [1998Ori] examined the crystallization processes of Alx-Mg 1-x-Ni alloys which were mechanically alloyed under an argon atmosphere by planetary ball milling for 4800 min at ambient temperature and 400 rpm. A phase with CsCl type crystal structure was found in alloys with x = 0.3-0.5 and an amorphous phase formed in alloys with x < 0.2. [2000Yua] synthesized Alx-Mg2-x-Ni (x = 0, 0.1, 0.2, 0.3, 0.4, 0.5) samples by a diffusion method. Mixtures of pure Al, Mg and Ni powders were grounded and pressed into pellets under a pressure of 30 MPa. The pellets were annealed at 540-550°C for 4 h and then cooled to room temperature. X-ray diffraction and SEM were applied to investigate their structure. A new phase of cubic crystal structure of Ti2Ni type was observed in the alloys, so that with x = 0.5 only this phase and a trace of magnesium were detected. [2001Gua] studied by X-ray diffraction the Ni2Mg3Al ternary alloy prepared from components of purities better than 99.95 % by compacting their mixtures at 30 MPa and annealing them at 540-550°C for 4 h under 0.5 MPa argon atmosphere. The composition of the alloy prepared coincided actually with the composition of a new ternary phase found in the investigation by [2000Yua]. [2001Gua] confirmed the existence of the new ternary phase with the composition Ni2Mg3Al and studied its crystal structure using more advanced X-ray techniques. As a result, the crystal structure of Ni2Mg3Al is established and described in more detail than by [2000Yua]. [1991Han] addressed some thermodynamic aspects on the effect that aluminium has on magnesium-nickel melts in presence of 3.8-8.6#10-4 mass% O. [2000Aiz] studied the effect that the substitution of aluminium by magnesium has on hydrogen absorption by a material based on Mg2Ni. Binary Systems The Al-Mg and Al-Ni binary phase diagrams are accepted from [2003Luk], [2003Sal], respectively. The Mg-Ni phase diagram is accepted from [1998Jac]. [1998Jac] made a thermodynamic assessment of the Mg-Ni binary system using the experimental characteristics of the Mg-Ni phase diagram from [1934Hau, 1978Bag, 1996Mic]. The calculated phase diagram is in a good agreement with the data from the experimental works.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Ni
151
Solid Phases The data on the relevant binary phases and ternary phases are listed in Table 1. [2001Gua, 2000Yua] found a new ternary phase of the same stoichiometry Ni2Mg3Al; its structural characteristics were determined and described in detail by [2001Gua]. Although the ternary alloys in both works were prepared in similar ways the Ni2Mg3Al alloy contained different phases in addition to the main phase. Therefore, the real composition of the compound discovered may differ slightly from the stoichiometry given. Invariant Equilibria At least one invariant four-phase equilibrium and one three-phase equilibrium exist in the ternary Al-Mg-Ni system, besides those in the adjacent binary systems. They are in the region of aluminium-rich alloys. The four-phase equilibrium is of eutectic type at a temperature of 449°C [1944Cha, 1952Han, 1993Pri]. The temperature of this equilibrium is assumed to be only by a few tenths of a degree lower than that of the binary eutectic reaction L(Al)+Mg2Al3, which is reliably confirmed to be at 450.5°C [2003Luk]. Type and temperature of the three-phase equilibrium however are not firmly established. It is of eutectic nature and takes place at a temperature between 449°C, where the four-phase eutectic equilibrium is, and 552°C the melting temperature of Mg2Al3, [1993Pri]. The characteristics of the three-phase and four-phase invariant equilibria are listed in Table 2 according to [1993Pri] with some correction for (Al) and Mg2Al3 by [2003Luk]. Concentration of the liquid phase in the three-phase invariant equilibrium is not determined exactly, but taking into account its temperature it is reasonable to assume that it is close to the L(Al)+Mg2Al3 eutectic point in the binary Al-Mg system. The reaction scheme for Al-NiAl3-Mg2Al3 region is shown in Fig. 1. Liquidus, Solidus Surfaces The liquidus surface of the Al-Mg-Ni system in Al-NiAl-Mg2Al3 region is shown in Fig. 2. It is a compilation of the [1952Han, 1934Fus] data with some corrections drawn out that the next phase after NiAl3 should be Ni2Al3 [1968Var, 2003Sal], rather than NiAl2, as it was proposed by [1934Fus]. The temperatures of the invariant reactions in the binary systems are also corrected to comply with the today accepted binary descriptions of Al-Mg and Al-Ni [2003Luk, 2003Sal]. The projection of the solidus surface in the Al-Mg2Al3-NiAl3 region is plotted in Fig. 3 based on [1952Han] with correction of the (Al) and Mg2Al3() homogeneity ranges by [2003Luk]. The Ni2Al3 homogeneity range is shown according to [2003Sal]. Temperature – Composition Sections The statement of [1924Fus] that the Mg2Al3-NiAl 3 section is a quasibinary one can not be correct taking into account the Al-Ni phase diagram [2003Sal], where the NiAl3 phase is shown to form by a peritectic reaction from liquid and Ni2Al3. Figure 4 gives the NiAl3-Mg2Al3 temperature-concentration cut constructed using the data of [1952Han, 2003Luk, 2003Sal]. It can be considered as a quasibinary one only below the solidus temperature of the alloys and within the part between Mg2Al3 and the edge of the Ni2Al3 primary crystallization surface including the e3 eutectic point. Thermodynamics [1991Han] showed that activity of magnesium, containing 3.8-8.6·10-4 % O, in nickel melts increases with addition of aluminium. Notes on Materials Properties and Applications NiMg2 base alloys with addition of Al are candidate materials for hydrogen storage [1998Ori]. Electrochemical capacity and live-cycles of NiMg2-xAlx (0 x 0.5) alloys during absorption and desorption of hydrogen increase with increasing Al contents, due to increasing amount of the Ni2Mg3Al Landolt-Börnstein New Series IV/11A3
MSIT ®
152
Al–Mg–Ni
phase in the alloy [2000Yua]. Addition of Al also improves the corrosion resistance of the NiMg2-xAlx alloys to a certain degree because an Al2O3 oxide layer forms on the surface. The corrosion rate of the ternary alloys is lower than that of NiMg2 [2000Yua]. Chemical modification of NiMg2 alloy by aluminium addition to (NiMg1.8Al0.2) is expected not to lead to significant reduction of onset temperature for hydrogen absorbing [2000Aiz]. NiMg1-xAlx phase with CsCl type crystal structure dissolves hydrogen interstitially without any structural transformation [1998Ori]. References [1924Fus] [1934Fus] [1934Hau] [1943Mon]
[1944Cha] [1952Han]
[1968Var]
[1978Bag]
[1991Han]
[1993Pri]
[1996Mic] [1998Jac] [1998Ori]
[2000Aiz] [2000Yua]
MSIT®
Fuss, V., “On the Constitution of Ternary Al Alloys” (in German), Z. Metallkd., 16, 24, (1924) (Equi. Diagram, Experimental, 1) Fuss, V., “Metallography of Al and its Alloys”, Berlin, The Sherwood Press. Inc., Cleveland, 142-143 (1934) (Equi. Diagram, Experimental, 1) Haughton, J.L., Payne, R.I., J. Inst. Met., 54, 275-283 (1934) quoted by [1998Jac] (Thermodyn.) Mondolfo, L., “Al-Mg-Ni, Aluminium-Magnesium Nickel”, in “Metallography of Aluminium Alloys”, John Wiley and Sons, Inc., New-York - London, 101-102 (1943) (Equi. Diagram, Review, 1) Chao, H.L., “On the Ternary System Al-Mg-Ni”, Thesis, Berlin Techn. Hochschule (1944) (Equi. Diagram, Experimental, 1) Hanemann, H., Schrader, A., “Examples for the Crystallization of Ternary Systems” (in German), Atlas Metallographicus, 3(2), 120-122 (1952) (Equi. Diagram, Experimental, #, *) Varich, N.I., Litvin, B.N., “Structure of Phases in the Aluminium-Magnesium System Containing Transition Metals” (in Russian), Izv. Akad. Nauk SSSR, Met., 6, 179-182 (1968) (Experimental, 4) Bagnoud, P., Feschotte, P., “The Binary Systems Magnesium-Copper and Magnesium Nickel, Especially the Nod-Stoechiometry of the MgCu2 and MgNi2 Laves Phases” (in French), Z. Metallkd., 69, 114-120 (1978) (Crys. Structure, Equi. Diagram. Experimental, 24) Han, Q., Wang, C., “Equilibrium of Mg-O and the Effect of Fe, Al and Cr on the Activity of Mg in Molten Nickel”, Beijing Keji Dexue Xuebao, 13(5), 461-466 (1991) (Experimental, Thermodyn., 4) Prima, S., “Aluminium-Magnesium-Nickel”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.19481.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 10) Micke, K., Isper, H., “Thermodynamic Properties of Liquid Magnesium-Nickel Alloys”, Monatsh. Chem., 127, 7-13 (1996) (Equi. Diagram, Experimental, Thermodyn., 18) Jacobs, M.H.G., Spencer, P.I., “A Critical Thermodynamic Evaluation of the System Mg-Ni”, Calphad, 22(4), 519-525 (1998) (Equi. Diagram, Review, Thermodyn., #, *, 30) Orimo, I.S., Ikeda, K., Fujii, H., “B2-Phase Formation and Hydriding Properties of (Mg1-xAlx)Ni (x = 0~0.5)”, J. Alloys Compd., 266, L1-L3 (1998) (Crys. Structure, Experimental, 10) Aizawa, T., “Solid-State Synthesis of Magnesium Base Alloys”, Mater. Sci. Forum, 350-351, 299-310 (2000) (Experimental, 22) Yuan, H.T., Wang, L.B., Cao, R., Wang, Y.J., Zhang, Y., Yan, D.Y., Zhang, W.H., Gong, W.L., “Electrochemical Characteristics of Mg 2-xAlxNi (0<x<0.5) Alloys”, J. Alloys Compd., 309, 208-211 (2000) (Crys. Structure, Electrochem. Prop., Experimental, 8)
Landolt-Börnstein New Series IV/11A3
Al–Mg–Ni [2001Gua]
[2003Luk]
[2003Sal]
153
Guanglie, L., Linshen, C., Lianbang, W., Huantang, Y., “Study on the Phase Composition of Mg2-xMxNi (M = Al, Ti) Alloys”, J. Alloys Compd., 321(1), L1-L4 (2001) (Crys. Structure, Experimental, 8) Lukas, H.L., Lebrum, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49) 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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu
Lattice Parameters Comments/References [pm] a = 404.96
a = 410.5 0.8 (Mg) < 650
hP2 P63/mmc Mg
a = 320.94 c = 521.07
(Ni) < 1455
cF4 Fm3m Cu
a = 352.40
, Mg17Al12 < 458
cI58 I43m Mn cF1168 Fd3m Mg2Al3 hR159 R3 Mn44Si9 oP16 Pnma NiAl3 hP5 P3m1 Ni2Al3 cI112 Ia3d Ni3Ga4
a = 1054.38
, Mg 2Al3 < 452 J, Mg23Al30 410 - 250 NiAl3 < 856 Ni2Al3 < 1138 Ni3Al4 < 702
Landolt-Börnstein New Series IV/11A3
at 25°C [Mas2] dissolves 0.01 at.% Ni at 639.9°C [2003Sal] and 18.6 at.% Mg at 450.5°C [2003Luk] at 445°C in the alloy with 18.6 at.% Mg [1952Han] at 25°C [Mas2] dissolves 11.5 at.% Al at 436°C [2003Luk] and < 0.04 mol% Ni at 500°C [1934Hau] at 25°C [Mas2] dissolves 20.2 at.% Al at 1385°C [2003Sal] < 0.2 mol% Mg at 1100°C [1998Jac] at 41.4 at.% Al [V-C2] 39.5 to 51.7 at.% Al [2003Luk]
a = 2816 - 2824
60 to 62 at.% Al [2003Luk]
a = 1282.54 c = 2174.78
56.3 at.% Al [2003Luk]
a = 661.3 0.1 b = 736.7 0.1 c = 481.1 0.1 a = 402.8 c = 489.1
[2003Sal]
a = 1140.8 0.1
36.1 to 39.8 at.% Ni [2003Sal]
[2003Sal]
MSIT ®
Al–Mg–Ni
154 Phase/ Temperature Range [°C] NiAl < 1651 Ni5Al3 < 723
Ni3Al < 1372 NiMg2 < 759.31 Ni2Mg < 1147.60 * NiMg1-xAlx
* NiMg1-xAlx * Ni2Mg3Al
Pearson Symbol/ Space Group/ Prototype cP2 Pm3m CsCl oC16 Cmmm Pt5Ga3 cP4 Pm3m AuCu3 hP18 P6222 NiMg2 hP24 P63/mmc Ni2Mg cP2 Pm3m CsCl amorphous phase cF96 Fd3 m derived from Ti2Ni
Lattice Parameters Comments/References [pm] 42.1 to 71.3 at.% Ni [2003Sal] a = 286.0 63 to 68 at.% Ni at 63 at.% Ni [2003Sal]
a = 753 b = 661 c = 376 a = 356.77
75.4 to 76.3 at.% Ni [2003Sal]
a = 520.5 0.1 c = 1320 6
[V-C2] [1998Jac]
a = 482.4 0.2 c = 1582.6
66.2 at.% at 759.31°C to 67.34 at.% Ni at 1095.28°C [1998Jac], [V-C2] In alloys with x = 0.3 - 0.5 prepared by mechanical alloying [1998Ori]
In the alloys with x < 0.2 prepared by mechanical alloying [1998Ori] a = 1154.74 0.02 [2001Gua]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
L Mg2Al3 + NiAl3
449 - 552
e3
L Mg2Al3 + NiAl3 + (Al)
449
E
L Mg2Al3 NiAl3 L Mg2Al3 NiAl3 (Al)
MSIT®
Composition (at.%) Al Mg ~60 ~40 61 39 75 0 64.6 34.6 61 39 75 0 81.39 18.6
Ni ? 0 25 0.8 0 25 0.01
Landolt-Börnstein New Series IV/11A3
Al–Mg–Ni
Al-Mg
155
Al-Mg-Ni
A-B-C
Al-Ni 644 e1 l NiAl3+(Al)
450.5 e2 l β + (Al) e3 L β + NiAl3
449
449
Lβ+NiAl3+(Al)
E
β+NiAl3+(Al) Fig. 1: Al-Mg-Ni. Reaction scheme in the partial Al-Mg 2Al3-NiAl3 system
Al Fig. 2: Al-Mg-Ni. Liquidus surface in the Al-rich alloys
e1,644
Data / Grid: at.% Axes: at.%
(Al)
10
p2,856
90
NiAl3
20
80
Ni2Al3
p1,1138 30
70
E
NiAl
e2,450.5°C
β
40
60
e3
Ni 50.00 Mg 0.00 Al 50.00 Landolt-Börnstein New Series IV/11A3
10
20
30
40
Ni 0.00 Mg 50.00 Al 50.00
MSIT ®
Al–Mg–Ni
156
Al Fig. 3: Al-Mg-Ni. Solidus surface of the partial Al-Mg2Al3-NiAl3 system
Data / Grid: at.% Axes: at.%
(Al) 10
90
20
80
NiAl3
β +(Al)+NiAl3 449°C
30
70
Ni2Al3
β
40
10
Ni 50.00 Mg 0.00 Al 50.00
Fig. 4: Al-Mg-Ni. The NiAl3-Mg 2Ni3 vertical section
60
20
30
40
Ni 0.00 Mg 50.00 Al 50.00
1200
~1116°C 1100
1000
Temperature, °C
L 900
856°C
L+Ni2Al3
800
700
600
L+NiAl3+Ni2Al3
500
L+NiAl3
449
Mg 0.00 Ni 25.00 Al 75.00
MSIT®
e3
Mg2Al3+NiAl3 20
10
Ni, at.%
452°C
Mg 40.00 0.00 Ni Al 60.00
Landolt-Börnstein New Series IV/11A3
Al–Mg–Sc
157
Aluminium – Magnesium – Scandium Evgeniya V. Lysova, updated by Rainer Schmid-Fetzer and Alexander Pisch Literature Data [1976Tur] investigated Al rich alloys of the Al-Mg-Sc system in the range up to 26.0 mass% Mg and 3.0 mass% Sc by thermal and metallographic analysis. The starting materials were 99.99% pure Al, 99.91% pure Mg, and scandium which contained 0.03% Cu, 0.01% Fe, 0.01% Ca, and 0.01% Si. The alloys were melted in an electrical resistance furnace in corundum crucibles under a layer of flux composed of 50% LiCl and 50% KCl and cast into thick-walled Cu moulds. The castings were homogenized at 400°C for 30 h, deformed under various conditions, annealed in evacuated ampoules at 640°C, 550°C and 430°C for 100, 360 and 600 h, respectively and finally quenched in water. Etching of the specimens was possible with a solution of 25 ml HNO3+1.5 ml HCl+1 ml HF per 100 ml of water. A partial isothermal section of the Al-Mg-Sc system at 430°C and two vertical sections have been established. There are indications that an invariant equilibrium exists at approximately 447°C. One alloy Al-2.5Mg-0.4Sc (mass%) has been studied by extensive X-ray diffraction after annealing at 150°C (10 h) and 350°C (2 h) and additional irradiation by electrons with energy 2.3 MeV [1984Rep]. [1989Odi, 1991Odi] examined the Al-Mg rich part of the system by differential thermal analysis, micrographic and X-ray analysis. An isothermal section at 400°C as well as a partial liquidus surface has been proposed for this region. The Al-Mg rich part of this system has also been investigated by [1999Gro]. Based on thermodynamic equilibrium calculations, key samples have been defined to determine the isothermal section at 350°C, the invariant reactions in the Mg rich part as well as the liquidus surface. Starting materials were 99.999% pure Al and 99.99% pure Mg and Sc. Two types of samples have been prepared: binary Al-Sc master alloys by levitation or e-beam melting mixed with Mg for solid state reactions and ternary Al-Mg-Sc samples, melted in an induction furnace in sealed Ta crucibles under Ar atmosphere. Both types of samples have been annealed for 2 months to reach equilibrium at 350°C. The samples were analyzed by X-ray diffraction, optical and electronic microscope, electron microprobe analysis as well as differential thermal analysis. Binary Systems For the binary systems the following versions have been adopted: Al-Mg [1998Lia], which is essentially the same as [2003Luk], Al-Sc [1999Cac] and Mg-Sc [1998Pis]. Solid Phases No ternary phases have been detected in the investigated range of concentrations. Although magnesium dissolves a considerable amount of Al or Sc, the ternary solubility is extremely small. There is a small solubility of magnesium in the four intermetallic Al-Sc compounds at 350°C: 4 at.% Mg in Sc2Al, 12 at.% in ScAl, 1-2 at.% in ScAl2 and 5 at.% in ScAl3 [1999Gro]. The solubility of scandium in the binary Al-Mg phases is small. A value of 1 at.% at 350°C has been found for (Mg17Al12) [1999Gro]. No information has been given for the solubility of aluminium in the binary Mg-Sc phases. Crystallographic data on the binary compounds are given in Table 1. Invariant Equilibria Based on thermodynamic calculations by extrapolation of the three binary systems, [1999Gro] identified 14 invariant reactions. The Al-Mg-Sc system is characterized by a liquid miscibility gap in the ternary with associated invariant reactions of the eutectic type and a series of U type reactions in the Mg-rich corner. The measured temperature of 116540°C for the reaction LL´+ScAl+ScAl2 is reproduced by the calculations. Two of the five U type reactions (U3: L+ScAlScAl 2+(Mg) and U 5: L+ScAl2ScAl3+(Mg)) have been measured and the temperatures agree also well with calculated ones (U3: 59010°C; U5: 66010°C). All Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Mg–Sc
158
other ternary invariant reactions including the binary Al-Mg are degenerated and close to the binary values. The reported L(Al)+ScAl3+(Mg2Al3) at 447°C by [1976Tur] is in agreement with [1999Gro]. The values of [1991Odi] have not been considered due to an incorrect isothermal section at 400°C with subsequent erroneous invariant equilibrium reactions. The invariant equilibria according to [1999Gro] are summed up in Table 2. Liquidus Surface The liquidus surface, based on thermodynamic calculations and experimental DTA results [1999Gro] is drawn in Figs. 1a, 1b. The results of [1991Odi] differ considerably and are not considered. One possible explanation for this discrepancy is the use of sealed quartz ampoules by these authors. Quartz is known to easily react with magnesium. Isothermal Sections Figure 2 shows the isothermal section at 350°C as determined by [1999Gro]. (Mg), depending on the mutual solubilities of Sc and Al, is in equilibrium with ScAl, ScAl2 and ScAl3. All binary Al-Mg phases are only in equilibrium with ScAl3. The isothermal section at 430°C for the Al rich region according to [1976Tur] is reproduced in Fig. 3. Additions of scandium substantially decrease the solubility of magnesium in aluminium. The point of ultimate saturation of aluminium with scandium and magnesium was originally observed at 10.5 mass% (11.6 at.%) Mg and 0.01 mass% (0.006 at.%) Sc [1976Tur]. Due to the low Mg solubility in Al, as shown by [1976Tur] at 430°C (~11 mass% Mg), the (Al)+ScAl3+(Mg2Al3) vertex has been shifted to 12.5 mass% Mg and 0.01 mass% Sc in order to meet the accepted binary value of 13 mass% Mg in Al given by [1981Sch]. The two phase (Al)+(Mg2Al3) region is very narrow and closely adjoins the Al-Mg side, whereas the two phase (Al)+ScAl3 region is quite wide. The proposition of [1984Rep] does not fit with Fig. 2. In both the irradiated and non-irradiated alloys faint X-ray reflections were observed in addition to those of the (Al) matrix. These were attributed to phases Mg2Al3, “Mg 3Al4” and “Mg5Al8”. This is a strange proposition since the alloy Al-2.5Mg-0.4Sc (mass%) is located clearly inside the (Al)+ScAl3 two phase field, which was confirmed in over 50 alloys by [1976Tur]. The work of [1984Rep] does not provide preparation details and does not refer to the prior results of [1976Tur]. Temperature – Composition Sections Figures 4 and 5 show two calculated isopleths [1999Gro] in agreement with the experimental sections from [1976Tur]. The two diagrams correspond to the sections between 17Mg-Al (mass%) and 1.0Sc-Al (mass%) and 22.0Mg-Al (mass%) and 2.0Sc-Al (mass%), respectively. Both sections intersect two regions of primary crystallization: of the aluminium solid solution (Al) and of the ScAl3 compound. The calculated ScAl2-Mg section from [1999Gro], confirmed by selected experiments, is drawn in Fig. 6. One notices the steep liquidus in the Mg-rich part of the diagrams. Notes on Materials Properties and Applications Additions of small amounts of scandium has been found to significantly improve yield stress, fatigue strength and resistance against microcrack growth of Al-Mg alloys [1981Dri, 1984Dri, 1990Saw, 1992Ela, 1997Rod] References [1976Tur]
MSIT®
Turkina, N.I., Kuzmina, V.I., “Phase Reactions in Al-Mg-Sc Alloys (up to 26 % Mg and 3 % Sc)”, Russ. Metall., (4), 179-183 (1976), translated from Izv. Akad. Nauk SSSR, Met., (4), 208-212 (1976) (Equi. Diagram, Experimental, #, 9)
Landolt-Börnstein New Series IV/11A3
Al–Mg–Sc [1981Dri]
[1981Sch]
[1984Dri] [1984Rep]
[1985Sch] [1989Gsc] [1989Odi]
[1990Saw] [1991Odi]
[1992Ela] [1997Rod] [1997Su]
[1998Lia]
[1998Pis]
[1999Cac]
[1999Gro]
[2003Luk]
Landolt-Börnstein New Series IV/11A3
159
Dritz, M.E., Pavlenko, S.G., Toropova, L.S., Bykov, Yu.G., Ber, L.B., “Mechanism of Scandium Effect to Increasing Strength and Termal Stability of Al-Mg Alloys”, Dokl. Akad. Nauk SSSR, 257(2), 353-356 (1981) (Crys. Structure, Experimental, 11) Schuermann, E., Voss, H.J., “Investigation of the Liquid Equilibria of Mg-Li-Al Alloys: Part 4. Liquid Equilibria of the Mg-Al Binary System” (in German), Giessereiforschung, 33, 43-46 (1981) (Equi. Diagram, Experimental, #, 17) Dritz, M.E., Ber, L.B., Bykov, Yu.G., Toropova, L.S., Anastaseva, G.K., “Ageing Alloy Al-0.3 at.% Sc”, Phys. Met. Metallogr., 57, 118-126 (1984) (Experimental) Repnikova, Ye.A., Malinenko, I.A., Chudinova, S.A., Toropova, L.S., Ustinovshchikov, V.M., “Influence of Electron Irradiation on Decomposition of Alloy Al-Mg-Sc”, Phys. Met. Metallogr., 57(3), 103-106 (1984) (Crys. Structure, Experimental, 3) Schuster, J.C., Bauer, J., “The Ternary System Sc-Al-N and Y-Al-N”, J. Less-Common Met., 109, 345 (1985) (Experimental, Crys. Structure) Gscheidner K.A., Calderwood, F.M., “ The Al-Sc (Aluminium-Scandium) System”, Bull. Alloy Phase Diagrams, 10, 34-36, (1989) (Crys. Structure, Equi. Diagram, Review, #, 18) Odinaev, K.O., Ganiev, I.N., Kinzhibalo, V.V., Kotur, B.Y., “Phase Diagram of the Aluminum-Magnesium-Scandium System in the 0-33.3 at.% Sc Interval at 673K”, Dokl. Akad. Nauk Tadzh. SSR, 32(1), 37-38 (1989) (Experimental, Equi. Diagram, 4) Sawtell, R.R., Jensen, C.L., “Mechanical Properties and Microstructures of Al-Mg-Sc Allyos”, Met. Trans., 21A, 421-430 (1990) (Crys. Structure, Mechan. Prop.) Odinaev, K.O., Ganiev, I.N., Vakhobov, A.V., “Quasi Binary Sections and the Liquids Surface of the Al-Mg-Sc System”, Dokl. Akad. Nauk SSSR, Met., (4), 195-197 (1991) (Experimetnal, Equi. Diagram, 10) Elagin, V.I., Zakharov, V.V., Rostova, T.D., “Scandium-Alloyed Aluminium Alloys”, Met. Sci. Heat Treat., 34, 37-45 (1992) (Experimental) Roder, O., Wirtz, T., Gysler, A., Luetjering, G., “Fatigue Properties of Al-Mg Alloys with and without Scandium”, Mater. Sci. Eng. A, A234-236, 181-184 (1997) (Experimental, 5) Su, H.-L., Harmelin, M., Donnadieu, P., Baetzner, C., Seifert, H.J., Lukas, H.L., Effenberg, G., Aldinger, F., “Experimental Investigation of the Mg-Al Phase Diagram from 47 to 63 at.% Al”, J. Alloys Compd., 247, 57-65 (1997) (Crys. Structure, Experimental, Equi. Diagram, #, *, 20) Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M., Quivy, G., Ochin, P., Effenberg, G., Seifert, H.J., Lukas, H.-L., Aldinger, F., “Experimental Investigation and Thermodynamic Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89, 536-540 (1998) (Equi. Diagram, Thermodyn., Experimental, Assessment, *, 33) Pisch, A., Schmid-Fetzer, R., Cacciamani, G., Riani, P., Saccone, A., Ferro, R., “Mg-rich Phase Equilibria and Thermodynamic Assessment of the Mg-Sc System”, Z. Metallkd., 89 (7), 474-477 (1998) (Equi. Diagram, Experimental, *,11) Cacciamani, G., Riani, P., Saccone, A., Ferro, R., Pisch, A., Schmid-Fetzer, R., “Thermodynamic Measurements and Assessment of the Al-Sc System”, Intermetallics, 7, 101-108 (1999) (Experimental, Equi. Diagram, Thermodyn., 26) Groebner, J., Schmid-Fetzer, R., Pisch, A., Cacciamani, G., Riani, P., Parodi, N., Borzone, G., Saccone, A., Ferro, R., “Experimental Investigations and Thermodynamic Calculations in the Al-Mg-Sc System”, J. Phase Equilib., 90(II), 872-880 (1999) (Experimental, Calculation, Themodyn., Equi. Diagram, 23) Lukas, H.-L., Lebrun, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49)
MSIT ®
Al–Mg–Sc
160 Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 (Mg) < 650 (Sc) 1541 - 1337 (Sc) < 1337 , Mg 2Al3 452 , Mg17Al12 < 458 J, Mg23Al30 410 - 250 ScAl3 < 1320 ScAl2 < 1370 ScAl < 1300
Sc2Al < 1300
(Mg-Sc system) < 520
MSIT®
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu hP2 P63/mmc Mg cI2 Im3m W hP2 P63/mmc Mg cF1168 Fd3m Mg2Al3 cI58 I43m Mn hR159 R3 Mn44Si9 cP4 Pm3m AuCu3 cF24 Fd3m Cu2Mg oP8 Cmcm CrB cP2 Pm3m CsCl hP6 P63/mmc Ni2In cP2 Pm3m CsCl
Lattice Parameters Comments/References [pm] a = 404.96
pure at 25°C [Mas2]
a = 320.94 c = 521.05
pure at 25°C [Mas, V-C]
a = 373
[1989Gsc]
a = 330.90 c = 527.33
at room temperature [Mas, V-C]
a = 2816 to 2824
60-62 at.% Al [2003Luk] 1168 atoms on 1704 sites per unit cell [2003Luk] at 41.4 at.% Al [V-C2] 39.5 to 51.5 at.% Al [2003Luk]
a = 1054.38
a = 1282.54 c = 2174.78
54.5 to 56.5 at.% Al [2003Luk]
a = 410.3 a = 411.6
[1989Gsc] 3.7 at.% Mg [1999Gro]
a = 758.2 a = 757.8 a = 757.5 a = 398.8 b = 988.2 c = 365.2 a = 345.0 a = 339.1 a = 343.2 a = 344.7 a = 488.8 c = 616.6 a = 488.5 c = 615.7 -
[1989Gsc] [1999Cac] 1.0 at.% Mg [1999Gro] [1985Sch]
[1989Gsc] 4.9 at.% Mg [1999Gro] 9.7 at.% Mg 11.5 at.% Mg [1989Gsc] 4 at.% Mg [1999Gro] [Mas]
Landolt-Börnstein New Series IV/11A3
Al–Mg–Sc
161
Table 2: Invariant Equilibria According to [1999Gro] Invariant Reaction
T [°C]
Type
L' L + ScAl + ScAl2
1165
M1
L' L + ScAl + Sc 2Al
1144 (calc.)
M2
L + Sc2Al ScAl + (Sc)
787 (calc.)
U1
L + (Sc) ScAl + (Mg)
709 (calc)
U2
L + ScAl ScAl2 + (Mg)
660
U3
(Sc) + ScAl Sc2Al + (Mg)
624 (calc.)
U4
L + ScAl2 ScAl3 + (Mg)
590
U5
(Sc) (Mg) + (Sc), Sc2Al
481 (calc.)
D1
L (Al) + (Mg2Al3), ScAl3
450 (calc.)
D2
L (Mg17Al12) + (Mg2Al3), ScAl3
449 (calc.)
D3
L (Mg) + (Mg17Al12), ScAl3
436 (calc.)
D4
Sc2Al + (Mg) ScAl + (ScMg)
410 (calc.)
U6
(Mg17Al12) + (Mg2Al3) J, ScAl3
410 (calc.)
D5
J (Mg17Al12) + (Mg2Al3), ScAl3
250 (calc.)
D6
Sc
Data / Grid: at.%
Fig. 1a: Al-Mg-Sc. Calculated liquidus surface
Axes: at.%
(β Sc)
e320
80
12 20 e2 e1
AlSc2
40
60
AlSc
M 2'
114 0
60
p1
40
M2
M1'
Al2Sc
L, + L
80
20
Al3Sc
M1
1220
U1 AlSc
1140 e4
Al
Landolt-Börnstein New Series IV/11A3
U2 U3
20
D2 40
D3
60
D4
80
U5
Mg
MSIT ®
Al–Mg–Sc
162
Al 0.00 Mg 70.00 Sc 30.00
Fig. 1b: Al-Mg-Sc. Enlarged schematic Mg rich corner of the calculated liquidus surface
Data / Grid: at.% Axes: at.%
AlSc2
10
20
U1 (β Sc) AlSc 20
10
U2
Al3Sc
g) (M
Al2Sc U5
U3 80
Al 30.00 Mg 70.00 Sc 0.00
90
Mg
Sc
Data / Grid: at.%
Fig. 2: Al-Mg-Sc. Partial isothermal section at 350°C
Axes: at.%
20
80
AlSc2 40
60
AlS c+A lSc
AlSc
2 +M
gSc
60
40
AlS c2 +A lSc+ (Mg )
Al2Sc Al3Sc
80
Al S 3 c+A lSc + 2 (Mg )
MSIT®
20
20
ε+A
not investigated
Al
MgSc Al Sc +M gS c+ (M g)
l3 S c+γ
β
ε 40 β+Al3Sc+ε
γ+Al
3 Sc+(M
(Mg)
g)
γ
60
80
Mg
Landolt-Börnstein New Series IV/11A3
Al–Mg–Sc
163
40
Fig. 3: Al-Mg-Sc. Partial isothermal section at 430°C
Mg, at. %
α + Mg2Al3 + ScAl3 α + Mg2Al3
α + ScAl3
α
0
Al
1.5
0
Sc, at. %
800
Temperature, °C
Fig. 4: Al-Mg-Sc. Calculated vertical section from 0.6Sc-99.4Al to 18.5Mg-81.5Al (at.%)
Thermal arrest [1976Tur]
L 700
L + Al3Sc
600
L + Al3Sc + (Al)
500
Al3Sc + (Al) 450
400
Al 99.40 Mg 0.00 Sc 0.60
Landolt-Börnstein New Series IV/11A3
90
Al, at.%
Al 81.50 Mg 18.50 Sc 0.00
MSIT ®
Al–Mg–Sc
164
800
Fig. 5: Al-Mg-Sc. Calculated vertical section from 98.8 Al-1.2 Sc to 76.2 Al-23.8 Mg (at.%)
Thermal arrest [1976Tur]
700
Temperature, °C
L + Al3Sc
600
L + Al3Sc + (Al) 500
Al3Sc + (Al) 450
400
Al 98.80 Mg 0.00 Sc 1.20
Temperature, °C
Fig. 6: Al-Mg-Sc. Calculated vertical section from Al2Sc to the Mg corner
90
Al 76.20 Mg 23.80 Sc 0.00
80
Al, at.%
Thermal arrest [1999Gro] 1250
1000
L + Al2Sc
750
L + Al2Sc + (Mg) 616
500
(Mg) + Al2Sc (Mg) + Al2Sc + Al3Sc 250
Al 66.66 Mg 0.00 Sc 33.33
MSIT®
60
50
40
30
20
10
Mg
Al, at.%
Landolt-Börnstein New Series IV/11A3
Al–Mg–Si
165
Aluminium – Magnesium – Silicon K.C. Hari Kumar, Nirupan Chakraborti, Hans-Leo Lukas, Oksana Bodak, Lazar Rokhlin Literature Data Al-Mg-Si alloys are being increasingly used in automotive and aerospace industries for critical structure applications because of their excellent castability and corrosion resistance and, in particular, good mechanical properties in the heat treated condition. These are known as 4xxx series of wrought alloys and 3xx.0 and 4xx.0 series of casting alloys. In these alloys, Mg is intentionally added to induce age hardening through precipitation of Mg2Si, metastable phases or Guinier-Preston zones. Several studies pertaining to the liquidus of the system are reported in the literature. [1921Han] studied alloys containing 0 to 11 mass% Si by thermal analysis in the Al-rich corner. [1930Ota] studied 24 alloys in the Al-rich corner, up to 20 mass% Si and 15 mass% Mg. [1931Dix] examined Al-rich alloys by thermal analysis and metallography, while [1931Los] reported thermal analysis data for a total of 150 alloys covering the entire composition range. [1935Saw] measured liquidus temperatures of 29 Mg-rich alloys (50 to 100 mass%) with Si contents up to 12 mass%. [1941Phi] examined the liquidus of certain Al-rich alloys. [1958Gul] gave data on 4 alloys with 1 to 7 at.% Si and approximately equal amounts of Mg and Al. [1976Fis, 1977Sch] again studied the liquidus of the whole system, primarily by thermal analysis and 15 vertical sections were reported along with the isotherms superimposed on the complete liquidus surface. More recent investigations of the liquidus are due to [2001Goe, 2001Li, 2001Bar, 2002Bar], employing thermal analysis. The homogeneity range of the (Al) solid solution was measured using metallography by [1921Han, 1931Dix] and [1943Wes], who reported the data as solubilities of Mg2Si in (Al). [1936Kel] also gave solubility values for Mg2Si and excess Mg. [1940Kuz] measured lattice parameters of the (Al) solid solutions along different lines in the Gibbs triangle and deduced the boundaries of the homogeneity range. Using dilatometry [1997Feu] reported two data points on Al solvus corresponding to the Al-Mg2Si section, which is in reasonable agreement with data reported by [1931Dix] and [1940Kuz]. [1941Phi] constructed several isothermal and vertical sections of the Al corner from metallographic measurements. An isothermal section of the Al corner at 460°C was reported by [1948Axo] based on metallographic experiments. [1988Rok] presented the (Mg) corner of the (Mg)+Mg2Si+Mg17 Al12 three-phase field at 430, 400 and 300°C. The purity of the starting materials was 99.8 to 99.99 % for Mg and Al, but the Si in all investigations contained some impurities (1 % Fe in an 80 % Al, 20 % Si master alloy [1921Han], 0.66 % Fe+0.1 % Ti [1931Dix], 0.3 % Fe in Si [1936Kel], 0.28 % Fe+0.17 % Ca [1931Los], 1.5 % Fe [1941Phi], 1 % mainly Fe [1977Sch] or 0.5 % of impurities not specified [1935Saw]). Reports of thermodynamic measurements for this system are rather limited. The activity of Mg in liquid phase was determined employing emf technique by [1979Seb] using alloys of nine different compositions near the Al-rich corner at three temperatures, 700, 750 and 800°C. The enthalpy change for the ternary eutectic reaction at 560°C was determined by [1980Bir] using DSC. [1986Lue] made a thermodynamic optimization of the system and reported a single ternary interaction parameter for the liquid phase, based mainly upon results of [1977Sch]. The data of [1977Sch], however, failed to extrapolate to the currently accepted melting point of Si [1986Bul] due to the limited purity of the Si used in the experiments. The deviation is about 12°C. [1991Zhi] determined the specific heats and enthalpies of the phase transformations of Al-Si and Al-Mg-Si alloys by isothermal calorimetry measurements. The ternary system was recalculated by [1992Cha] using phase stability values for pure elements recommended by [1989Din]. The calculated diagram agrees with the experimental data reasonably [1921Han, 1930Ota, 1931Dix, 1940Kuz, 1941Phi, 1943Wes, 1948Axo, 1958Gul, 1977Sch, 1979Seb]. [1993Rei] determined temperatures of the secondary phase particle formation in two Al-Mg-Si alloys. One of the alloys had a composition corresponding to the section Al-Mg2Si and the other alloy had composition with some excess of Si as compared with Al-Mg2Si. The results of the experiments confirmed the phase diagram presented by [1992Cha]. [1997Feu] carried out experiments in the Al corner of the system and updated the
Landolt-Börnstein New Series IV/11A3
MSIT ®
166
Al–Mg–Si
thermodynamic assessment of the whole system. They also measured enthalpy of formation and fusion as well as the heat capacity of Mg2Si employing calorimetry. Binary Systems Binary systems Al-Si [2003Luk1] and Al-Mg [2003Luk2] are from the MSIT Binary Evaluation Program. The binary system Mg-Si is from [1997Feu]. Solid Phases No stable ternary compounds have been reported. Several metastable phases were reported to form during annealing of quenched supersaturated (Al) solid solutions [1999Mat, 2000Cay, 2001Mar, 2002Der, 2003Mar]. The stable and some metastable phases are listed in Table 1. The solid phase (Si) has only negligible solubility for Al and Mg. Pseudobinary Systems The Al-Mg2Si section is approximately pseudobinary. It is shown in Fig. 1, calculated from the dataset of [1997Feu]. The calculation shows that the (Al) phase in the eutectic maximum contains more than twice as much Mg than Si. Therefore, the section Al-Mg2Si is not exactly pseudobinary and shows an extended three-phase field L+(Al)+Mg2Si [1997Feu, 2001Zha]. In fact it is reported that the real pseudobinary is shifted more towards the Mg-rich region [1997Feu, 2001Bar, 2001Li, 2002Bar], located along the section Mg2Si-Al97.2 Mg2.4 [2001Goe]. The temperature of the so-called pseudobinary eutectic was given as 595°C by [1930Ota, 1931Dix, 1997Feu, 2001Goe, 2001Li], as approximately 593°C by [1977Sch] and as approximately 590°C by [1921Han, 1931Los, 1941Phi], and 597°C [2001Bar, 2002Bar]. Using the thermodynamic model parameters reported by [1997Feu] it is calculated to be 594°C. Invariant Equilibria The reaction scheme (Fig. 2) proposed by [1992Cha] is calculated from the dataset of [1997Feu] with the data of the binary Al-Mg intermediate phases replaced by those of [1998Lia] assuming zero solubilities of Si in these phases. The calculated compositions of the phases in the reactions containing liquid are listed in Table 2. The temperature of E1 was reported as 551°C [1930Ota, 1941Phi], [1931Los] ~557°C, ~550°C [1977Sch], and 560°C [1980Bir]. Calculations by [1992Cha, 1997Feu] indicate E1 to be at 557°C. E4 was reported to be at 437°C by [1935Saw] and 435°C by [1977Sch, 1988Rok], which are nearly identical to the binary L(Mg)+ eutectic (e7). In the experimental investigation by [1997Feu], E1 is reported to be at 558°C. Liquidus Surface The liquidus surface is shown in Fig. 3, calculated from the dataset of [1997Feu]. Only the ternary equilibria e3, E1, E4 are indicated and not E2, E3, which virtually coincide with the binary reactions e7, e8, respectively. In Fig. 4 the Al corner of the liquidus surface is shown, calculated from the dataset of [1997Feu]. Isothermal Sections Figures 5 and 6 show the isothermal sections at 600 and 550°C. At room temperature, all solid phases are in equilibrium with Mg2Si. The calculated solidus and solvus isotherms of the (Al) solid solution are given in Figs. 7 and 8, respectively, after [1997Feu]. The results of [1988Rok] could not be reproduced by the calculation of [1997Feu] as it contradicts Henry’s rule for dilute solutions. Due to the small Si solubility in (Mg) this rule predicts the Al solubility of (Mg) in the three-phase field (Mg)+ +Mg2Si to be very near to that in the binary two-phase field (Mg)+ .
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Si
167
Temperature – Composition Sections Two calculated vertical sections are shown in Figs. 9 and 10, originally reported by [1992Cha], but recalculated here using the thermodynamic data set of [1997Feu]. Figures 11 to 13 show calculated vertical sections in the Al corner as experimentally investigated and calculated by [1997Feu]. Thermodynamics As mentioned in the section “Literature Data” thermodynamic information on the system is limited to the emf studies [1979Seb] in the Al-rich liquid and the enthalpy of the eutectic reaction E1 [1980Bir]. The enthalpy of melting of the eutectic reaction E1 is reported to be +26.4 kJ#mol-1 [1980Bir]. References [1921Han]
[1930Ota] [1931Dix]
[1931Los] [1935Saw]
[1936Kel]
[1940Kuz]
[1941Phi] [1943Wes]
[1948Axo]
[1958Gul]
[1976Fis]
[1977Sch]
[1979Seb]
Landolt-Börnstein New Series IV/11A3
Hanson, D., Gayler, M.L.V., “The Constitution and Age-Hardening of the Alloys of Aluminium with Magnesium and Silicon”, J. Inst. Met., 26, 321-359 (1921) (Equi. Diagram, Experimental, 4) Otani, B., “Siliciumin and its Structure” (in Japanese), Kinzuku no Kenkyu, 7, 666-686 (1930) (Equi. Diagram, Experimental, 10) Dix, E.H., Keller, F., Graham, R.W., “Equilibrium Relations in Aluminium-Magnesium Silicide Alloys of High Purity”, Trans. A.I.M.M.E., 404-420 (1931) (Equi. Diagram, Experimental, 10) Losana, L., “The Ternary System Al-Mg-Si” (in Italian), Metall. Ital., 23, 367-382 (1931) (Equi. Diagram, Experimental, 14) Sawamoto, H., “Equilibrium Diagram of the Magnesium-Rich Magnesium-AluminiumSilicon Ternary System” (in Japanese), Suiyokwai Shi, 8, 713-727 (1935) (Equi. Diagram, Experimental, 25) Keller, F., Craighead, C.M., “Equilibrium Relations in Aluminium-Magnesium Silicide Alloys Containing Excess Magnesium”, Trans. A.I.M.M.E., 122, 315-323 (1936) (Equi. Diagram, Experimental, 4) Kuznetsov, V.G., Makarov, E.S., “X-Ray Investigation of the Structure of Ternary Solid Solutions of Magnesium and Silicon in Aluminium” (in Russian), Izv. Sekt. Fiz.-Khim. Anal., 13, 177-190 (1940) (Equi. Diagram, Experimental, 18) Phillips, H.W.L., “The Constitution of Alloys of Aluminium with Magnesium and Silicon”, J. Inst. Met., 67, 257-273 (1941) (Equi. Diagram, Experimental, 9) Westlinning, H., Klemm, W., “The Solubility of Mg 2Si, Mg2Ge, Mg2Sn and Mg2Pb in Aluminium” (in German), Z. Elektrochem., 49, 198-200 (1943) (Equi. Diagram, Experimental, 4) Axon, H.J., Hume-Rothery W., “The Effect of 1% Silicon on the Constitution of Aluminium Magnesium Manganese Zinc Alloys at 460°C”, J. Inst. Met., 74, 315-329 (1948) (Equi. Diagram, Experimental, 10) Gul'din, I.T., Dokokina, N.V., “The Aluminium-Magnesium-Iron-Silicon System”, Russ. J. Inorg. Chem., 3, 359-379 (1958), translated from Zh. Neorg. Khim., 3, 799-814 (1958) (Equi. Diagram, Experimental, 5) Fischer, A., “Investigations on Equilibria with Liquid and Mechanical Properties of the Binary Al-Mg and Mg-Si Systems as well as of the Aluminium-Magnesium-Silicon Ternary System” (in German), Thesis, University of Clausthal, F.R. Germany (1976) (Equi. Diagram, Experimental, #, *, 66) Schürmann, E., Fischer, A., “Equilibria with Liquid in the Aluminium- Magnesium-Silicon Ternary System, Part 3, Al-Mg-Si System” (in German), Giessereiforschung, 29, 161-165 (1977) (Equi. Diagram, Experimental, #, *, 14) Sebkova, J., Beranek, M., Halamkova, P., “Thermodynamic Properties of Liquid Al-Mg-Si Alloys” (in Czech), Kovove Mater., 17, 137-143 (1979) (Experimental, Thermodyn., 12) MSIT ®
168 [1980Bir] [1986Bul] [1986Lue] [1988Rok]
[1989Din] [1991Zhi]
[1992Cha] [1993Rei] [1997Feu]
[1998Lia]
[1999Mat]
[2000Cay]
[2001Bar]
[2001Goe] [2001Li]
[2001Mar]
[2001Zha]
[2002Bar]
MSIT®
Al–Mg–Si Birchenall, C.E., Riechmann, A.F., “Heat Storage in Eutectic Alloys”, Metall. Trans. A, A11, 1415-1420 (1980) (Thermodyn., Experimental, 13) “Melting Points of the Elements”, Bull. Alloy Phase Diagrams, 7, 602 (1986) (Review, 0) Lüdecke, D., “Phase Diagram and Thermochemistry of the Al-Mg-Si System”, Z. Metallkd., 77, 278-283 (1986) (Equi. Diagram, Theory, Thermodyn., 33) Rokhlin, L.L., Pepelyan, A.G., “Phase Equilibria in the Mg-Al-Si System in the Magnesium Rich Area” (in Russian), Izv. Akad. Nauk SSSR, Met., 176-179 (1988) (Equi. Diagram, Experimental, 4) Dinsdale, A.T., “SGTE Data for Pure Elements”, NPL Report DMA(A) 195, September (1989) (Review, 20) Zhiguang, H., Sinong, X., Guangzhong, W., Shaohua, M., “Measuring Heat of Thermal Storage of Phase Change Metal” (in Chinese), Gongcheng Rewuli Xuebao (J. Eng. Thermophys.), 12(1), 46-49 (1991) (Thermodyn., Experimental, 5) Chakraborti, N., Lukas, H.L., “Thermodynamic Optimization of the Mg-Al-Si Phase Diagram”, Calphad, 16, 79-86 (1992) (Equi. Diagram, Review, Thermodyn., 27) Reiso, O., Ryum, N., Strid, J., “Melting of Secondary-Phase Particles in Al-Mg-Si Alloys”, Metall. Trans. A, 24A, 2629-2641 (1993) (Equi. Diagram, Experimental, 13) Feufel, H., Gödecke, T., Lukas, H.L., Sommer, F., “Investigation of the Al-Mg-Si System by Experiments and Thermodynamic Calculations”, J. Alloys Compd., 247, 31-42 (1997) (Equi Diagram, Thermodyn., Experimental, Theory, 38) Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G., Seifert, H. J., Lukas, H.L., Aldinger, F., “Experimental Investigation and Thermodynamic Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89, 536-540 (1998) (Equi. Diagram, Thermodyn, Experimental, Theory, 33) Matsuda, K., Naoi, T., Fujii, K., Uetani, Y., Sato, T., Kamio, A., Ikeno, S., “Crystal Structure of the ´´ Phase in an Al-1.0 mass% Mg 2Si-0.4 mass% Si Alloy”, Mater. Sci. Eng. A, A262, 232-237 (1999) (Crys. Structure, Experimental, 14) Cayron, C., Buffat, P.A., “Transmission Electron Microscopy Study of the Phase (Al-Mg-Si Alloys) and QC Phase (Al-Cu-Mg-Si Alloys) Ordering Mechanism and Crystallographic Structure”, Acta Mater., 48, 2639-2653 (2000) (Crys. Structure, Experimental, 38) Barabash, O.M., Sulgenko, O.V., Legkaya, T.N., Korzhova, N.P., “Experimental Analysis and Thermodynamic Calculation of the Structural Regularities in the Fusion Diagram of the System of Alloys Al-Mg-Si”, J. Phase Equilib., 22(1), 5-11 (2001) (Calculation, Equi. Diagram, Experimental, 13) Goedecke, T., “Direction of Crystallisations Paths in Ternary As-cast Alloys” (in German), Z. Metallkd., 92(8), 966-978 (2001) (Equi. Diagram, Experimental, 37) Li, S.-P., Zhao, S.-X., Pan, M.-X., Zhao, D.-Q., Chen, X.-C., Barabash, O.M., “Eutectic Reaction and Microstructural Characteristics of Al(Li)-Mg2Si Alloys”, J. Mater. Sci., 36, 1569-1575 (2001) (Equi. Diagram, Experimental, 10) Marioara, C.D., Andersen, S.J., Jansen, J., Zandbergen, H.W., “Atomic Model for GP-Zones in a 6082 Al-Mg-Si System”, Acta Mater., 49, 321-323 (2001) (Crys. Structure, Metastable, Experimental, 12) Zhang, J., Fan, Z., Wang, Y.Q., Zhou, B.L., “Equilibrium Pseudobinary Al-Mg2Si Phase Diagram”, Mater. Sci. Technol., 17, 494-496 (2001) (Calculation, Equi. Diagram, Experimental, 17) Barabash, O.M., Milman, Yu.V., Korzhova, N.P., Legkaya, T.N., Podrezov, Yu.N., “Design of New Cast Aluminium Materials Using Properties of Monovariant Eutectic Transformation L-Al+Mg2Si”, Mater. Sci. Forum, 396-402, 729-734 (2002) (Equi. Diagram, Mechan. Prop., 9)
Landolt-Börnstein New Series IV/11A3
Al–Mg–Si [2002Der]
[2003Luk1]
[2003Luk2]
[2003Mar]
169
Derlet, P.M., Andersen, S.J., Marioara, C.D., Froseth, A., “A First Principles Study of the ”-Phase in Al-Mg-Si Alloys”, J. Phys.: Condens. Matter, 14, 4011-4024 (2002) (Crys. Structure, Theory, 19) Lukas, H.L., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 29) Lukas, H.L., Lebrum, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49) Marioara, C.D., Andersen, S.J., Jansen, J., Zandbergen, H.W., “The Influence of Temperature and Storage Time at RT on Nucleation of the ” Phase in a 6082 Al-Mg-Si Alloy”, Acta Mater., 51, 789-796 (2003) (Crys. Structure, Experimental, 13)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Mg) < 650 (Al) < 660.452 (Si) < 1414 , Mg 2Al3 < 452 J, Mg23Al30 410 - 250 , Mg17Al12 < 458 Mg2Si < 1076 MgAl3Si6 ”
Pearson Symbol/ Space Group/ Prototype hP2 P63/mmc Mg cF4 Fm3m Cu cF8 Fd3m C (diamond) cF1832 Fd3m Mg2Al3 hR53 R3 Mg23Al30 cI58 I43m Mn cF12 Fm3m CaF2 mP* P2/m ?
(Mg,Al)5Si6 ”
mC* C2/m ?
´
hP* P62m
Landolt-Börnstein New Series IV/11A3
Lattice Parameters Comments/References [pm] a = 320.94 c = 521.07
at 25°C [Mas2]
a = 404.96
pure Al, 25°C [Mas2]
a = 543.06
pure Si, 25°C [Mas2] ~0 at.% Al, ~0 at.% Mg [Mas2] 60-62 at.% Al [V-C2]
a = 2823.9
a = 1282.54 c = 2174.78
[V-C2, 1998Lia]
a = 1048.11 a = 1053.05 a = 1057.91 a = 633.8
52.58 at.% Mg [L-B] 56.55 at.% Mg [L-B] 60.49 at.% Mg [L-B] [Mas2, V-C2]
a = 770 20 b = 670 10 c = 203 = 75 0.5° a = 1516 b = 405 c = 674 = 105.3° a = 710 c = 405
[1999Mat] metastable precipitate in (Al), aged at 150°C [2001Mar, 2002Der] metastable precipitate in (Al)
[2000Cay] metastable precipitate in (Al)
MSIT ®
Al–Mg–Si
170 Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
L (Al) + Mg2Si
594
e3
L (Al) + (Si) + Mg2Si
557
E1
L + Mg2Si
462.5
e5
L + Mg2Si
451.2
e6
L (Al) + + Mg2Si
450
E2
L + + Mg 2Si
449
E3
L (Mg) + + Mg2Si
435.6
E4
L (Al) Mg2Si L (Al) (Si) Mg2Si L Mg2Si L Mg2Si L (Al) Mg2Si L Mg2Si L (Mg) Mg2Si
Fig. 1: Al-Mg-Si. Section from Al to Mg2Si
Composition (at. %) Al Mg 10.8 85.3 2.7 97.1 66.7 0 5.4 81.5 0.70 98.0 0 0 66.7 0 53.8 46.1 53.9 46.1 66.7 0 61.0 38.9 61.1 38.9 0 66.7 36.3 64.0 16.5 83.4 38.9 61.1 66.7 0 42.5 57.4 38.9 61.1 48.1 51.9 66.7 0 69.0 30.9 88.4 11.6 60.1 39.9 66.7 0
Si 3.9 0.2 33.3 13.1 1.3 100.0 33.3 0.1 0 33.3 0.1 0 33.3 0.1 4.0#10 -6 0 33.3 0.1 0 0 33.3 0.1 5.5#10 -5 0 33.3 1076°C
1000
Temperature, °C
L
750
660.452°C
L+Mg2Si L+(Al)
(Al)
L+(Al)+Mg2Si
500
Mg2Si
(Al)+Mg 2Si 250
Al
20
40
Mg, at.%
MSIT®
60
Mg 66.67 0.00 Al Si 33.33 Landolt-Börnstein New Series IV/11A3
Landolt-Börnstein New Series IV/11A3
ca.250
D1
E3
557
γ + β + Mg2Si
ε γ + β, Mg2Si
D2
ε + β + Mg2Si
γ β + ε, Mg2Si
β + γ + Mg2Si
L β + γ + Mg2Si
(Al) + β + Mg2Si
Al-Mg-Si
A-B-C
435.6
E2
E1
(Mg) + γ + Mg2Si
L (Mg) + γ + Mg2Si
451 e6 L β, Mg2Si
(Al) + (Si) + Mg2Si
L (Al) + (Si) + Mg2Si
594 e3 L (Al) + Mg2Si
L (Al) + β + Mg2Si
ε + γ + Mg2Si
ca.410
449
450
462 e5 L γ, Mg2Si
Fig. 2: Al-Mg-Si. Reaction scheme
ca.250 e10 εγ+β
ca.410 p1 γ+βε
436 e9 l (Mg) + γ
449.5 e8 lβ+γ
450.5 e7 l (Al) + β
Al-Mg
E4
639 e2 l (Mg) + Mg2Si
941 e1 l (Si) + Mg2Si
Mg-Si
577 e4 l (Al) + (Si)
Al-Si
Al–Mg–Si 171
MSIT ®
Al–Mg–Si
172
Si
Data / Grid: at.%
Fig. 3: Al-Mg-Si. Liquidus surface
Axes: at.%
1300 °C
20
80
1200
(Si) 40
60
1100
e1
1000 60
900
40
800 80
0 100
700
Mg2Si
600 e2
Mg
e9
20
(Mg)
40
80
Mg 0.00 Al 80.00 Si 20.00
Fig. 4: Al-Mg-Si. Calculated liquidus surface in the Al corner
e4
0 60 (Al)
E2
e8 60β e7
γ
E1
e3
E3
E4
20
Al
Data / Grid: at.% Axes: at.%
660
640 (Si)
620 600 580
E1
580 10
10
600
Mg2Si
620 600 580
Mg 20.00 Al 80.00 Si 0.00
MSIT®
(Al)
0 62
e3
0 64
90
Al
Landolt-Börnstein New Series IV/11A3
Al–Mg–Si
173
Si
Data / Grid: at.%
(Si)
Fig. 5: Al-Mg-Si. Isothermal section at 600°C
Axes: at.%
20
80
40
60
L+(Si)+Mg2Si
60
40
Mg2Si
L+(Si)
80
(Mg)+Mg2Si
L+Mg2Si
L+(Mg)+Mg2Si
Mg
L
20
(Mg)
40
60
80
Si
(Al)
Al
Data / Grid: at.%
(Si)
Fig. 6: Al-Mg-Si. Isothermal section at 550°C
Axes: at.%
20
80
40
60
(Al)+Mg2Si+(Si)
60
Mg2Si
(Mg)+Mg2Si L+(Mg)+Mg2Si
20
40
80
20
L+Mg2Si L+(Al)+Mg2Si
Mg
Landolt-Börnstein New Series IV/11A3
(Mg)
20
40
L
60
80
(Al)
Al
MSIT ®
Al–Mg–Si
174
Fig. 7: Al-Mg-Si. Solidus of the (Al) phase
1.6
(Al) at E1 1.4
°C 575
1.2
Si, at.%
°C 600
1.0
0.8
0.6
0.4
C 5° 62
(Al) at e3
0.2
C
0°
65 0.0
Al
0
2
1
3
4
5
6
7
8
9
10
6
7
8
9
10
Mg, at.%
Fig. 8: Al-Mg-Si. Solvus of the (Al) phase
1.6
(Al) at E1
1.4
1.2
Si, at.%
1.0
0.8
0.6
0.4
C 5° C 0° 55
57
(Al) at e3
50 0° C
0.2
C
0° 45 0.0
Al
0
1
2
3
4
5
Mg, at.%
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Si
Fig. 9: Al-Mg-Si. Vertical section from Si to Mg50.9Al49.1
175
1414°C
1250
Temperature, °C
L
1000
L+(Si)
750
(Si) L+Mg2Si
L+Mg2Si+(Si)
L+Mg2Si+γ
557°C L+(Al)+Mg2Si
500
γ+Mg2Si β +γ+Mg2Si γ+ε+Mg2Si ε+β +Mg2Si β +γ+Mg2Si
(Al)+Mg2Si (Al)+Mg2Si+(Si)
β +(Al)+Mg2Si
250
Si
20
Mg 50.90 Al 49.10 0.00 Si
40
Mg, at.%
Fig. 10: Al-Mg-Si. Vertical section from Mg2Si to Al53Si47
1076°C 1029°C 1000
L
L+Mg2Si
Temperature, °C
L+(Si)
750
L+Mg2Si+(Si)
Mg2Si
577°C 557°C L+(Al)+(Si) 500
(Si)+(Al)
(Al)+Mg2Si+(Si)
Mg 66.67 0.00 Al Si 33.33
Landolt-Börnstein New Series IV/11A3
60
40
20
Mg, at.%
Mg 0.00 Al 53.00 Si 47.00
MSIT ®
Al–Mg–Si
176
700
Fig. 11: Al-Mg-Si. Vertical section at 90 mass% Al L
Temperature, °C
610.3°C 600
596°C 577°C
L+(Si)+(Al)
L+(Al) L+(Al)+Mg2Si 557°C
L+(Al)+Mg2Si
(Si)+(Al) 500
(Si)+(Al)+Mg2Si
Mg 0.00 9.64 Si Al 90.36
(Al)
(Al)+Mg2Si
Mg 10.98 0.00 Si Al 89.02
10
Mg, at.%
700
Fig. 12: Al-Mg-Si. Vertical section at 85 mass% Al L
Temperature, °C
L+(Al) 600
L+(Al)
577°C L+(Si)+(Al)
L+(Al)
L+(Al)+Mg2Si
L+(Al)+Mg2Si
557°C (Si)+(Al) (Si)+(Al)+Mg2Si
(Al)+Mg2Si
500
Mg 0.00 Si 14.49 Al 85.51
MSIT®
10
Mg, at.%
Mg 16.38 0.00 Si Al 83.62
Landolt-Börnstein New Series IV/11A3
Al–Mg–Si
177
700
Temperature, °C
Fig. 13: Al-Mg-Si. Vertical section at 80 mass% Al
L
L+Mg2Si
L+(Si) 600
577°C L+(Al)+Mg2Si L+(Al)+(Si) L+(Si)+Mg2Si
L+(Al) L+(Al)+Mg2Si
(Si)+(Al) 500
(Si)+(Al)+Mg2Si
Mg 0.00 Si 19.36 Al 80.64
Landolt-Börnstein New Series IV/11A3
(Al)+Mg2Si
10
Mg, at.%
20
Mg 21.72 0.00 Si Al 78.28
MSIT ®
178
Al–Mg–Sn
Aluminium – Magnesium – Tin Lazar Rokhlin, updated by Hans Leo Lukas Literature Data [1958Bad] investigated the partial equilibrium diagram Al-Mg2Al3()-Mg2Sn-Sn using thermal and microscopic analyses. They studied nine vertical sections, determined the phase equilibria in the solid state, the nature and temperatures of invariant equilibria and constructed the liquidus surface. [1968Kop, 1969Sem, 1973Sem] investigated the magnesium rich corner of the equilibrium diagram limited by the join Mg17Al12( ) to Mg2Sn using thermal and microscopic analyses. The authors determined the joint solubility of Al and Sn in solid Mg at 400 and 200°C and the temperature and compositions of liquid of the ternary eutectic occurring in this part of the system, constructed four vertical sections and a projection of the liquidus surface. [1938Hum, 1943Wes] determined the solubility of the Mg2Sn compound in solid Al by microscopy or combined microscopy and X-ray diffraction, respectively. [1977Ray] reviewed the papers of [1968Kop, 1969Sem] regarding the Mg corner. A progress report [1950Dow] shows the phase diagram of the pseudobinary section Mg17Al12( ) to Mg2Sn. Binary Systems The binary Al-Sn and Mg-Sn systems are accepted from [Mas2]. Al-Mg is taken from [2003Luk] out of the MSIT collection of binary systems. It is based on the assessment of [1998Lia]. Solid Phases No ternary phases have been detected. The stable binary phases are summarized in Table 1. The (Al) solid solution dissolves up to 16.6 at.% Mg [1998Lia]. The solubility of Sn in solid (Al) is retrograde with a maximum of nearly 0.026 at.% Sn at 625°C [Mas2]. The maximum solubility of Sn in solid Mg is 3.35 at.% at the eutectic temperature, 561.2°C [Mas2]. The solubility of Al in solid (Mg) is 11.6 at.% in the binary Al-Mg system [1998Lia]. Sn and Al decrease somewhat the solubility of each other in solid Mg [1968Kop, 1973Sem]. Solid (Sn) dissolves about 1 at.% Al and nearly no Mg [Mas2]. There is no detectable solubility of Al in Mg2Sn [1938Hum]. It has not been established if some Sn is soluble in the phases , J and of the Al-Mg binary system [1958Bad, 1969Sem]. Pseudobinary Systems Three pseudobinary sections have been established: - Mg2Sn [1950Dow, 1969Sem], - Mg2Sn [1958Bad] and (Al) - Mg2Sn [1958Bad]. They are illustrated in Figs. 1 to 3. The two versions of the section - Mg2Sn disagree in the position and temperature of the eutectic, 11.5 mass% Mg2Sn and 450°C [1969Sem] or 2.5 mass% Mg2Sn and 455°C [1950Dow], respectively. [1969Sem] assumes a large vertical part (450 to 600°C) in the liquidus line, which postulates zero heat of reversible solution of Mg2Sn in liquid [1981Goo], which is not very likely along such a large distance as the liquidus becomes very flat not far above, indicating an appreciable reversible heat of solution there. In constructing the - Mg2Sn section [1969Sem] assumed certain mutual solubilities of the terminating solid phases, which have not been confirmed experimentally. After [Mas2, 1938Hum], however, the solubilities of Al as well of Mg in Mg2Sn are negligible. Therefore in the section in Fig. 1 no solubility is assumed in the terminating solid phases. The eutectic is accepted from [1950Dow]. Invariant Equilibria Three invariant four-phase equilibria have been reported in the Al-Mg-Sn system: L(Mg)+ +Mg 2Sn [1969Sem], L(Al)++Mg2Sn [1958Bad] and L(Sn)+(Al)+Mg2Sn [1958Bad]. In addition, three invariant three-phase equilibria take place in the pseudobinary systems: L +Mg2Sn [1969Sem], MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Sn
179
L+Mg2Sn [1958Bad] and L(Al)+Mg2Sn [1958Bad]. To complete the reaction scheme, an additional four-phase equilibrium L+ +Mg2Sn must be assumed. The compositions are given in Table 2. Those of the phases and are taken from the Al-Mg binary system assuming insignificant solubility of Sn in these phases. The equilibria L+Mg2Sn and L(Al)++Mg2Sn are nearly degenerate and the compositions of the liquid are very near to the Al-Mg binary system [1958Bad]. The temperature of L+Mg 2Sn is given 453°C by [1958Bad], but it cannot be higher than that of the maximum L in the binary Al-Mg system. Thus, 451°C is estimated here for this reaction. Also, the equilibrium L+ +Mg2Sn is assumed to be degenerate. For all the degenerate equilibria, the corresponding concentrations of the binary liquid are adopted for Table 2. The composition of (Sn) in L(Sn)+(Al)+Mg2Sn is adopted from the accepted Mg-Sn [Mas2] and Al-Sn [Mas2] binary systems. The three-phase equilibria in solid state of the Al-Mg binary system are expected to form degenerate four-phase equilibria in the ternary system with Mg2Sn as forth (inert) phase, since the Sn-solubility in all intermediate Al-Mg phases is assumed to be negligible. Figure 4 shows the reaction scheme. The concentration range of the triangle --Mg2Sn has not yet been investigated, but, as no ternary phases exist, it can easily be interpolated between the known parts. Liquidus Surface Figure 5 shows the liquidus surface. The isotherms within the region Mg2Sn-(Sn)-(Al)- are drawn according to [1958Bad] with small corrections due to the accepted binary Al-Sn and Mg-Sn systems. The isotherms within the region (Mg) - Mg2Sn - are constructed from the vertical sections given by [1950Dow, 1968Kop], partially from those of [1969Sem] and from the binary Mg-Sn and Al-Mg systems. The very flat part at 600°C in the field of primary crystallization of (Al) indicates the presence of a metastable miscibility gap in liquid just below the liquidus surface. Isothermal Sections Figure 6 shows the isothermal section at 250°C. It is constructed from the extensions of homogeneity regions of the phases liquid, (Al), , , (Mg) and (Sn) in the accepted binary systems, assuming negligible ternary solubilities in the intermediate phases and . Temperature – Composition Sections Several temperature-composition sections are given in literature. Besides the pseudobinary system - Mg2Sn [1969Sem] reported sections at constant Al content of 15 mass% and at constant Sn content of 18 mass%. [1968Kop] constructed a section at constant Mg content of 75 mass%, which is converted into at.% and redrawn in Fig. 7. The field (Mg)+ is corrected, in the original publication it is drawn too large, contradicting an isothermal section at the Mg corner given in the same paper. [1958Bad] reported 9 sections through the Al corner, at high Mg:Sn ratios only partially until the tie line - Mg2Sn. Figure 8 shows the section at constant atomic ratio Mg:Sn = 1:1. Thermodynamics [1983Som] measured the enthalpy of mixing of the liquid along the sections Mg2Sn - Al, Mg50Sn50 - Al, Mg30Sn70 - Al and Mg50Al50 - Sn at 835°C and along Mg50Sn50 - Al and Mg50Al50 - Sn also at 735°C. Complete thermodynamic datasets of the Al-Mg and Al-Sn binary systems were assessed in the COST 507 action [1998Ans]. Al-Sn in [1998Ans] contains a typing error, the parameter °Lfcc Al,Sn must be 45297.84+8.39814#T, not 45297.848.39814#T J#mol, but was assessed together with the original value from [1991Din] for °G fccSn - °GbctSn = 4150 - 5.2#T J#mol-1. Using the updated value for °GfccSn - °G bctSn = 5510 - 8.46#T J#mol-1 given in [1998Ans] the parameter must be corrected to °LfccAl,Sn = 43410.66 + 11.76812 #T J#mol-1. A thermodynamic dataset for Mg-Sn was assessed by [1993Fri]. For the liquid phase in the Al-Sn and Mg-Sn (above 75 at.% Sn) Gibbs energy datasets are also given by [1996Heu]. As the ternary solubilities in all solid phases are small, only the liquid phase needs additional ternary terms for a complete thermodynamic description of the ternary system. A calculation
Landolt-Börnstein New Series IV/11A3
MSIT ®
180
Al–Mg–Sn
using the above mentioned binary descriptions without ternary terms yields fairly good results in the Sn and Mg rich parts, but fails for the Al rich liquidus surface. Notes on Materials Properties and Applications There is information about favour effect of Sn additive on strength properties of Mg base alloys containing Al at elevated (200-250°C) temperatures [1968Kop]. References [1938Hum] [1943Wes] [1950Dow]
[1958Bad]
[1968Kop]
[1969Sem]
[1973Sem]
[1977Ray] [1981Goo]
[1983Som]
[1991Din] [1993Fri] [1996Heu]
[1998Ans]
[1998Lia]
MSIT®
Hume-Rothery, W., Raynor, G.V., “On the Nature of Intermetallic Compounds of the Type Mg2Sn”, Philos. Mag., 25, Ser. VII, 335-339 (1938) (Experimental, 3) Westlinning, H., Klemm, W., “The Solubility of Mg 2Si, Mg2Ge, Mg2Sn and Mg2Pb in Aluminium” (in German), Z. Electrochem., 49, 198-200 (1943) (Experimental, 3) “Liquidus Determinations of Polynary Magnesium Alloys“, Final Status Report No. 15004; Dow Chemical Company, Off. of Naval Res., Contract No. N9 ONR 85900, 1-20 (1950) (2) (Equi. Diagram, Experimental, 2) Badaeva, T.A., Kuznetsova, R.I., “The Structure of Alloys of Aluminium with Magnesium and Tin”, Tr. Inst. Metall. im. A.A. Baikova, (3), USSR Academy of Science, Moscow, 203-215 (1958) (Equi. Diagram, Experimental, 5) Kopetsky, Ch.V., Padezhnova, E.M., Semenova, E.M., “Investigation of the Equilibrium Diagram of the Mg-Al-Sn in the Magnesium-Rich Region” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., (5), 78-82 (1968) (Equi. Diagram, Experimental, 10) Semenova, E.M., “Equilibrium Diagram of the Mg-Al-Sn System in the Magnesium-Rich Region” (in Russian), Dokl. Akad. Nauk SSSR, 188, 1308-1310 (1969) (Equi. Diagram, Experimental, 8) Semenova, E.M., “Phase Composition and Properties of Alloys of the Mg-Al-Sn System” (in Russian), Tr. Inst. Metall. im. A.A. Baikova, Moscow, Nauka, 165-168 (1973) (Experimental, 5) Raynor, G.V., “Constitution of Ternary and More Complex Alloys of Magnesium”, Int. Met. Rev., (5), 65-95 (1977) (Equi. Diagram, Review, 83) Goodman, D.A., Cahn, J.W., Bennettt, L.H., “The Centennial of the Gibbs-Konovalov Rule for Congruent Points”, Bull. Alloy Phase Diagrams, 2, 29-34 (1981) (Equi. Diagram, Theory, 20) Sommer, F., Rupf-Bolz, N., Predel, B., “Investigations on the Temperature Dependence of the Enthalpy of Mixing of Ternary Alloy Melts” (in German), Z. Metallkd., 74, 165-171 (1983) (Experimental, Thermodyn., 15) Dinsdale, A.T., “SGTE Data for Pure Elements”, Calphad, 15, 317-425 (1991) (Thermodyn., Assessment) Fries, S., Lukas, H.L., “Optimisation of the Mg-Sn System”, J. Chim. Phys., 90, 181-187 (1993) (Equi. Diagram, Thermodyn., Assessment, 32) Heuzey, M.-C., Pelton, A.D., “Critical Evaluation and Optimization of the Thermodynamic Properties of Liquid Tin Solutions”, Metall. Mater. Trans. B, 27B, 810-828 (1996) (Equi. Diagram, Thermodyn., Assessment, 156) Ansara, I., Dinsdale, A.T., Rand, M.H., COST 507, Thermochemical Database for Light Metal Alloys, Vol. 2, European Communities, Luxembourg, Vol. 2, Al-Mg: 48-54; Al-Sn: 81-83 (1998) (Equi. Diagram, Thermodyn., Assessment, Crys. Structure, 0) Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G., Seifert, H.J., Lukas, H.L., Aldinger, F., “Experimental Investigation and Thermodynamic Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89, 536-540 (1998) (Equi. Diagram, Thermodyn., Experimental, Assesssment, 33)
Landolt-Börnstein New Series IV/11A3
Al–Mg–Sn [2003Luk]
181
Lukas, H.L., Lebrum, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 (Mg) < 650 Sn(r) 231.97 - 13.05 Sn(l) < 13.05 , Mg 2Al3 452 , Mg17Al12 < 458 J, Mg23Al30 410 - 250 Mg2Sn < 770
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu hP2 P63/mmc Mg tI4 I41/amd Sn cF8 Fd3m C (diamond) cF1168 Fd3m Mg2Al3 cI58 I43m Mn hR159 R3 Mn44Si9 cF12 Fm3m CaF2
Lattice Parameters Comments/References [pm] a = 404.88
pure Al at 25°C [Mas2]
a = 320.89 c = 521.01
pure Mg at 25°C [Mas2]
a = 583.18 c = 318.18
pure Sn at 25°C [Mas2]
a = 648.92
pure Sn [Mas2]
a = 2816 to 2824
60-62 at.% Al [2003Luk] 1168 atoms on 1704 sites per unit cell [2003Luk] at 41.4 at.% Al [V-C2] 39.5 to 51.5 at.% Al [2003Luk]
a = 1054.38
a = 1282.54 c = 2174.78
54.5 to 56.5 at.% Al [2003Luk]
a = 676.5
[V-C]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
L (Al) + + Mg2Sn
448
E1
L + + Mg 2Sn
447
E2
Landolt-Börnstein New Series IV/11A3
Phase L (Al) Mg2Sn L Mg2Sn
Composition (at.%) Al Mg 63 37 83 17 61 39 0 66.7 57 43 61 39 52 48 0 66.7
Sn 0.4 <0.01 0 33.3 0.4 0 0 33.3
MSIT ®
Al–Mg–Sn
182 T [°C]
Reaction
Type
L (Mg) + + Mg2Sn
428
E3
+ J, Mg2Sn J + , Mg2Sn L (Sn) + (Al) + Mg2Sn
410 250 198
D1 D2 E4
L (Al) + Mg2Sn
605
e1
L + Mg2Sn
455
e3
L , Mg 2Sn
451
e4
Phase
Composition (at.%) Al Mg 31.6 66.0 8.3 91.3 (41) (59) 0 66.7 50.6 49.4 46.4 53.6 0.9 8.8 0.7 (0) 100 (0) 0 66.7 86.2 9.8 98 2 0 66.7 (45.6) (53.5) (46) (54) (0) 66.7 60.8 38.8 61.1 38.9 0 66.7
L (Mg) Mg2Sn L (Sn) (Al) Mg2Sn L (Al) Mg2Sn L Mg2Sn L Mg2Sn
Sn 2.4 0.4 (0) 33.3 0 0 90.3 99.3 <0.01 33.3 4.0 0.02 33.3 0.9 (0) 33.3 0.4 (0) 33.3
Note: Values in brackets () are estimated.
800
Fig. 1: Al-Mg-Sn. The pseudobinary system - Mg2Sn
770.5°C L
Temperature, °C
700
600
L+Mg2Sn
500
458°C
455°C
e3
γ+Mg2Sn
L+γ 400
Mg 54.00 Al 46.00 Sn 0.00
MSIT®
10
20
Sn, at.%
30
Mg 66.70 0.00 Al Sn 33.30
Landolt-Börnstein New Series IV/11A3
Al–Mg–Sn
Fig. 2: Al-Mg-Sn. The pseudobinary system - Mg2Sn
183
800
770.5°C L
Temperature, °C
700
600
L+Mg2Sn
500
451°C
β +Mg2Sn 400
Mg 39.20 Al 60.80 Sn 0.00
Fig. 3: Al-Mg-Sn. The pseudobinary system Mg2Sn - Al
10
20
30
Sn, at.%
Mg 66.70 0.00 Al Sn 33.30
800
770.5°C L
Temperature, °C
700
660.452°C
L+Mg2Sn 605°C 600
e1
L+(Al)
Mg2Sn+(Al) 500
400
Mg 66.70 0.00 Al Sn 33.30
Landolt-Börnstein New Series IV/11A3
20
40
60
80
Al
Al, at.%
MSIT ®
MSIT®
250
β+ε+Mg2Sn
410
447
Fig. 4: Al-Mg-Sn. Reaction scheme
e8 250 εβ+γ
410 p1 β+γε
436 e7 l (Mg) + γ
449 e6 lβ+γ
450 e5 l (Al) + β
Al-Mg
β+γ+Mg2Sn
ε β + γ, Mg2Sn
β + γ ε, Mg2Sn
β+γ+Mg2Sn D1
E2
448
D2
451 e4 L β + Mg2Sn
605 e1 L (Al) + Mg2Sn
A-B-C
E1
(Sn)+(Al)+Mg2Sn
L(Sn) + (Al) + Mg2Sn
(Mg)+γ+Mg Sn
L (Mg) + γ + Mg2Sn
(Al)+β+Mg2Sn
L (Al) + β + Mg2Sn
428
198
ε+γ+Mg2Sn
L β + γ + Mg2Sn
455 e3 L γ + Mg2Sn
Al-Mg-Sn
E4
E3
228 e9 l (Sn) + (Al)
Al-Sn
203 e10 l(Sn) + Mg2Sn
561 e2 l (Mg) + Mg2Sn
Mg-Sn
184 Al–Mg–Sn
Landolt-Börnstein New Series IV/11A3
Al–Mg–Sn
185
Sn Fig. 5: Al-Mg-Sn. Liquidus surface
Axes: at.%
e9
(Sn) e10
35 0
E4
Data / Grid: at.%
0 40
450
20
80
500 550
40
60
(Al)
65 0
60
40
6 610 00 750
70 0
80
e2
550
620
6 600 50
Mg2Sn
e7 E3 40 γ e3
20
Mg
e1
E1
500
0 60
(Mg)
E2 e6 60e4 e5
80
β
Sn Fig. 6: Al-Mg-Sn. Isothermal section at 250°C
Axes: at.%
L
80
40
60
(Al)+L+Mg2Sn
60
40
Mg2Sn
80
20
(Mg)+γ +Mg2Sn
Landolt-Börnstein New Series IV/11A3
(Mg)
Al
Data / Grid: at.%
20
Mg
20
20
β+ γ+
M g
2 Sn
40
γ
(Al )+β 60
β
+M gS 2 n 80
(Al)
Al
MSIT ®
Al–Mg–Sn
186
Fig. 7: Al-Mg-Sn. Temperatureconcentration section at 75 mass% Mg
600
Temperature, °C
L
L+(Mg) 500
L+(Mg)+Mg2Sn
L+(Mg)+γ
428°C
(Mg)+Mg2Sn
(Mg)+γ 400
(Mg)+γ+Mg2Sn
Mg 93.61 0.00 Al Sn 6.39
10
Mg 76.90 Al 23.10 Sn 0.00
20
Al, at.%
700
Temperature, °C
Fig. 8: Al-Mg-Sn. 660.452°C Temperatureconcentration section 604°C L+(Al) 600 at the atomic ratio Mg:Sn = 1:1
L
635°C
31.55%Mg L+Mg2Sn
500
L+(Al)+Mg2Sn
400
300
198°C
200
203.5°C
(Al)+(Sn)+Mg2Sn
L+(Sn)+Mg2Sn
100
Al
10
20
30
Sn, at.%
MSIT®
40
Mg 50.00 0.00 Al Sn 50.00
Landolt-Börnstein New Series IV/11A3
Al–Mg–Ti
187
Aluminium – Magnesium – Titanium Frederick H. Hayes, updated by Andy Watson and Tatyana Dobatkina Literature Data There is general agreement that the solubilities of both Mg in Ti aluminides and Ti in liquid and solid Al-Mg alloys are very small [1954Eis, 1968Var, 1971Dil, 1973Kol, 1984Rus, 1989Ker1]. The observation of [1954Eis] that the solubility of Ti in liquid Mg, 0.04 mass% Ti at 750°C, is sharply decreased by Al additions was later confirmed by [1971Dil]. [1971Dil] found that additions of Al to Mg-Ti alloys reduce the solubility of Ti in the liquid phase as follows: at 720°C from 0.028 to 0.003 by 0.5 mass% Al, at 800°C from 0.042 to 0.004 by 1 mass% Al, at 900°C from 0.08 to 0.02 by 2 mass% Al. [1954Eis] reports that the solubility of Ti in Al-Mg alloys decreases with increasing Mg content at 750°C to become vanishingly small at 90 mass% Mg. At 12.5 mass% Mg [1954Eis] gives the solubility of Ti as 0.122 mass% at 750°C in agreement with the later work [1973Kol]. [1968Var] studied intermetallic phases in Al-Mg alloys containing traces of transition metals; TiAl3 was observed to be present in Al-Mg alloys containing up to 38 at.% Mg and 1 at.% Ti. The liquidus contours given by [1973Kol] from 700 to 850°C for the Al rich corner containing up to 12 mass% Mg and 0.8 mass% Ti are given in Fig. 1. [1984Rus] constructed the 427°C isothermal section for the entire composition range using kinetic data from vapor-diffusion and powder-sintering experiments. Electron-microprobe analysis, electron microscopy and X-ray phase-analysis techniques were used to determine phase compositions. Only binary Al-Ti and Al-Mg intermetallic phases were found. In contrast, the 487°C isothermal section of [1989Ker1] contains a ternary phase in the Al-rich corner; reference to a ternary Al-Mg-Ti phase is also made in the report of [1948Fel]. [1989Ker1] gives the single-phase composition range of the ternary phase as 78.5 to 80.5 at.% Al, 11.4 to 12.8 at.% Mg and 8.7 to 9.1 at.% Ti. In other respects the sections of [1984Rus, 1989Ker1] are in good agreement. No further information on solid-liquid equilibria or invariant reactions is available for this system. Binary Systems The binary Al-Mg system is accepted from [2003Luk], and the Al-Ti system from [1995Hay]. Both differ from those given by [Mas2] in the solid-solid equilibria. The binary Mg-Ti system is accepted from [Mas2]. Solid Phases A ternary phase is reported by [1948Fel, 1987Ker, 1989Ker1, 1989Ker2] but was not mentioned by [1968Var, 1984Rus]. [1987Ker] gave the lattice parameter and suggested it to be of the Cr2Mg3Al18 type. The crystallographic data of all solid phases are given in Table 1. Isothermal Sections Figure 2 gives the 427°C isothermal section based on [1984Rus, 1989Ker1]; it is consistent with the accepted binary phase diagrams. Thermodynamics From a thermodynamic analysis of the Gibbs energy of formation of binary phases in equilibrium with Ti2Mg3Al18 [1989Ker2] concluded that the Gibbs energy of formation of the ternary compound was approximately -15 kJ·mol-1.
Landolt-Börnstein New Series IV/11A3
MSIT ®
188
Al–Mg–Ti
Miscellaneous [1996Set] used X-ray electron probe analysis to define titanium concentrations throughout dendritic cells of a cast aluminium solid solution of Al-Mg-Ti. [1996Set] indicated that Mg levels within the limits of its solubility in aluminium, have no significant effect on titanium intercrystalline segregation. References [1948Fel]
[1954Eis] [1968Var]
[1971Dil]
[1973Kol]
[1984Rus]
[1987Ker]
[1989Ker1]
[1989Ker2] [1990Sch] [1995Hay] [1995Bra] [1996Set] [1998Lia]
[2003Luk]
MSIT®
Feldman, W., Schrader, A., Seemann, J., “Structure of Primary Aluminium and Aluminium Alloys” (in German), FIAT Rev. German Sci., 1939-1946, Non-Ferrous Metallurgy, 1, 153-155 (1948) (Review, 21) Eisenreich, H., Putter, H., “Magnesium-Titanium Ternary Systems” (in German), Metall, 8, 624-625 (1954) (Experimental, 0) Varich, N.I., Litvin, B. N., “Structure of Phases in the Aluminium-Magnesium System Containing Transition Metals” (in Russian), Izv. Akad. Nauk SSSR. Met., (6), 179-182 (1968) (Experimental, 4) Dilov, V.V., Sergeev, V.V., “Effect of Some Elements on Ti Solubility in Liquid Mg” (in Russian), Tr. Vses. N.–I. Proekt. Inst. Alyum. Magn. El., (79), 100-113 (1971) (Experimental, 4) Kolpachev, A.A., Medvedeva, N.D., Samoilova, Yu.A., Titova, I.A., “Solubility of Ti in Al-Mg Alloys” (in Russian), Tekhnol. Legk. Splavov, (8), 15-17, (1973) (Experimental, Equi. Diagram, #, 3) Rusnyak, V.D., Dunaev, S.F., Slyusarenko, E.M., Sokolskii, S.V., Sokolovskaya, E.M., “Study of Phase Equilibria in the Aluminium-Magnesium-Titanium System” (in Russian), Deposited Doc., VINITI, 2189-89, Moscow, 15 pp. (1984) (Experimental, Equi. Diagram, *, 11) Kerimov, K.M., Dunaev, S.F., Slyusarenko, E.M., “Investigation of the Structure of Ternary Phases in Al-Mg-Ti, Al-Mg-V and Al-Mg-Cr Systems”, J. Less-Common Met., 133, 297-302 (1987) (Experimental, Crys. Structure, 9) Kerimov, K.M., Dunaev, S.F., Slyusarenko, E.M., “Study of the Phase Diagrams of the Systems: Aluminium-Magnesium-(Titanium, Zirconium, Hafnium)” (in Russian), Vestn. Mosk. Univ., Ser. 2: Khim., 30(2), 156-161 (1989) (Experimental, Equi. Diagram, *, 8) Kerimov, K.M., Dunaev, S.F., “The M 2Mg3AL18 Phase in Al-Mg Transition Metal Systems”, J. Less-Common Met., 153, 267-273 (1989) (Thermodyn., 10) Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl 3-TiAl”, Z. Metallkd., 81, 389-396 (1990) (Equi. Diagram, Experimental, 33) Hayes, F.H., “The Al-Ti-V (Aluminium-Titanium-Vanadium) System”, J. Phase Equilib., 16(2), 163-176 (1995) (Equi. Diagram, Review) Braun, J., Ellner, M., Predel, B., “Experimental Investigations of the Structure and Stability of the TiAl Phase”, Z. Metallkd., 86(12), 870-876 (1995) (Experimental, Crys. Structure) Setiukov, O.A., Fridlyander, I.N., “Peculiarities of Ti Dendritic Segregation in Aluminium Alloys”, Mater. Sci. Forum, 217-222, 195-200 (1996) (Experimental, 2) Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G., Seifert, H.J., Lukas, H.-L., Aldinger, F., “Experimental Investigation and Thermodynamic Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89, 536-540 (1998) (Equi. Diagram, Thermodyn., Experimental, Theory, *, 33) Lukas, H.-L., Lebrun, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49)
Landolt-Börnstein New Series IV/11A3
Al–Mg–Ti
189
Table 1: Crystallographic Data of Solid Phases Phases/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Al) < 660.45
cF4 Fm3m Cu
a = 404.96
pure Al at 25°C [Mas2] dissolves ~17.0 at.% Mg
(Mg) < 650
hP2 P63/mmc Mg
a = 320.94 c = 521.07
pure Mg at 25°C [Mas2] dissolves ~0.12 at.% Ti, ~12 at.% Al
,(Ti) (h) 1670 - ~865
cI2 Im3m W
a = 330.65
pure Ti at 25°C [Mas2] dissolves 45 at.%Al, ~2.4 at.%Mg
,(Ti) (r) 1490
hP2 P63/mmc Mg
a = 295.06 c = 468.35
pure Ti at 25°C [Mas2] dissolves 51.5 at.% Al, ~1.6 % Mg
, Mg 2Al3 452
cF1168 Fd3m Mg2Al3
a = 2816 to 2824
60-62 at.% Al [2003Luk] 1168 atoms on 1704 sites per unit cell [2003Luk]
, Mg17Al12 < 458
cI58 I43m Mn
a = 1054.38
at 41.4 at.% Al [V-C2] 39.5 to 51.5 at.% Al [2003Luk]
J, Mg23Al30 410 - 250
hR159 R3 Mn44Si9
a = 1282.54 c = 2174.78
54.5 to 56.5 at.% Al [2003Luk]
TiAl3 < 1387
tI8 I4/mmm TiAl3
a = 384.88 c = 859.82
[1990Sch]
TiAl2(r) < 1216
tI24 I41/amd HfGa2
a = 396.7 c = 2429.68
[1990Sch]
TiAl < 1460
tP4 P4/mmm AuCu
2, Ti3Al 1180
* -, Ti2Mg3Al18
Landolt-Börnstein New Series IV/11A3
a = 398.69 c = 405.39 a = 398.8 c = 408.2
38.5 to 52 at.% Ti [1990Sch] at 38.5 at.% Ti, 1000°C [1990Sch] at 45 at.% Ti, 20°C [1995Bra]
hP8 P63/mmc Ni3Sn
a = 580.6 c = 465.5 a = 574.6 c = 462.4
at 78 at.% Ti [L-B]
cF184 Fd3m Cr2Mg3Al18
a = 1477
[1987Ker]
at 62 at.% Ti [L-B]
MSIT ®
Al–Mg–Ti
190
Al
Fig. 1: Al-Mg-Ti. Liquidus contours in the Al-rich corner [1973Kol]
Al
800
750 750 800
700
95
850
Al, at.% Al, at.%
700
95
850
90
90
Mg Ti Al Mg Ti Al
15.00 0.00 85.00 15.00 0.00 85.00 i 1
10
Mg 0.00 1.50 Ti Al 98.50 Mg 0.00 1.50 Ti Al 98.50
5
Mg, at.% 10
5
Mg, at.%
i id
i
h Al i h
[19 3
l]
Al
Fig. 2: Al-Mg-Ti. Isothermal section at 427°C
Data / Grid: at.% Axes: at.%
(Al)
20
τ
TiAl3+τ +(Al)
Mg2Al3+τ +(Al) Mg2Al3+TiAl3+τ
TiAl3
Mg2Al3
g) (M + l3 iA +2 T l A Ti
40
60
TiAl
17 A l 12 +
40
α2+TiAl+(Mg)
(M g) +M g
60
TiAl2
Ti A l 3
Mg2Al3+Mg17Al12+TiAl3 Mg17Al12
80
α2
80
20
(α Ti)+α2+(Mg) (Mg)
Mg
MSIT®
(α Ti) 20
40
60
80
Ti
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn
191
Aluminium – Magnesium – Zinc Dmitriy Petrov, Andy Watson, Joachim Gröbner, Peter Rogl, Jean-Claude Tedenac, Marina Bulanova, Volodymyr Turkevich, updated by Hans Leo Lukas Literature Data The Al-Mg-Zn system has a relatively complex equilibrium diagram. The first investigation of the entire system was carried out by [1913Ege]. By thermal analysis and metallography he determined the liquidus surface and some invariant equilibria. He also detected the first ternary phase, having a large homogeneity range between the binary phases Mg17Al12 and MgZn2. Many investigations were done in the following time by thermal analysis and metallography [1926San, 1936Ham1, 1936Ham2, 1936Koe1, 1936Koe2, 1936Koe3, 1940Ura, 1943But, 1945But, 1945Mik, 1949Sal, 1961Cla, 1962Ale, 1985Kuz1, 1985Kuz2], by electric conductivity measurements [1936Fin], by X-ray diffraction [1935Lav, 1936Fin, 1936Ham1, 1936Ham2, 1936Rie, 1957Ber, 1961Cla, 1985Kuz1, 1985Kuz2, 1995Tak, 1997Don, 2000Lee, 2000Sun] or by TEM [1997Don, 2001Bou1]. Thermodynamic datasets of the ternary system were assessed by [1997Lia, 1998Lia2]. Binary Systems The Al-Mg system is adopted from [2003Luk] which is based on the thermodynamic assessment of the COST 507 action [1998Ans] modified in the central part by [1998Lia1] based on new experimental data. The Al-Zn and Mg-Zn binary systems are accepted from the COST 507 action [1998Ans]. Solid Phases Two ternary phases are known since long time [1913Ege, 1961Cla]. Recently a stable ternary quasicrystalline phase and another stable crystalline phase were detected [1995Tak]. The first ternary phase (-1 in Table 1) was called Mg7Zn6Al3 by [1913Ege], Mg30Zn 25Al20 by [1929Ish, 1930Ish]. [1935Lav] determined a cubic unit cell with a = 1416 pm at a composition Mg3Zn 3Al2. [1936Rie] measured the lattice parameter at different compositions and found 1429 to 1471 pm along the line from Mg3Zn3Al2 to Mg2Al3 and 1429 to 1460 pm along the line from Mg3Zn3Al2 to Mg17Al12. [1952Ber, 1957Ber] determined the complete crystal structure using single crystal X-ray diffraction. The ideal formula is Mg32 (Zn,Al)49 with 162 atoms per unit cell. [2000Sun] and independently [2000Lee] refined the crystal structure stating the position 2(a) of space group Im3 to be empty, where [1952Ber, 1957Ber] assumed Al occupation. For seven of the eight remaining positions of [1952Ber, 1957Ber] both papers agree in having 100% Mg occupation at the sites 16(f), 24(g), 12(e) and Al,Zn mixed occupation on a 48(h) and two different 24(g) positions. [2000Sun], like [1952Ber, 1957Ber] assume also 100% Mg occupation at another 12(e) position, where [2000Lee] assume mixed occupation of about 8 atoms Mg+Zn and about 4 sites to be empty. This agrees with the experimental homogeneity range having significant extension also perpendicular to Al-Zn exchange. The phase is closely related to the quasicrystalline phases and characterized as 1/1 crystalline approximant of these phases. [1959Cla, 1961Cla] established the existence of another ternary phase at 40Mg-40Zn-20Al (mass%) (54.9Mg-20.4Zn-24.7Al (at.%)) and designated this phase (1in Table 1 and in [2001Bou1]). The phase is in equilibrium with (Mg) at 335, 204°C and probably at room temperature. [1961Wri] also discovered the presence of this ternary phase. From transmission electron diffraction data [1997Don] derived an orthorhombic unit cell of this ternary phase and successfully indexed the X-ray powder diagram (except for two reflections) with lattice parameters a = 897.9 pm, b = 1698.8 pm and c = 1934 pm. [1997Don] gave the solubility range of the this phase as (53 to 55)Mg-(18 to 29)Al-(17 to 28)Zn (at.%), which are in good agreement with [1961Cla]. [2001Bou1] derived a model for the crystal structure of this phase from electron diffraction patterns obtained in a transmission electron microscope: Space group Pbcm, Mg84(Al,Zn)68
Landolt-Börnstein New Series IV/11A3
MSIT ®
192
Al–Mg–Zn
(4 times Mg21(Al,Zn)17) on 23 different Wyckoff-positions, 13 of them occupied by Mg and 10 ones occupied by Al+Zn. The lattice parameters were adopted from [1997Don]. Icosahedral quasicrystalline phases can be prepared by ultra rapid quenching of melts of compositions inside the stability range of the -1 phase. [1995Tak] found the quasicrystalline phase to be stable even up to the melting temperature of about 380°C, but only in a very narrow composition range at Mg44Zn41 Al15 . A sample of this composition is still in the quasicrystalline state after 5h annealing at 360°C, whereas a sample of composition Mg45Zn40 Al15 shows small amounts of (Mg) precipitates in a quasicrystalline matrix after 1 h annealing at 360°C. The sample 1 mol% opposite, Mg43Zn 42Al15, after the same annealing conditions transforms into a cubic phase with a = 2291 pm, which also is stable in a very narrow composition range and which is characterized as the 2/1 crystalline approximant of the icosahedral quasicrystalline phase. This phase was confirmed and characterized as package of 8 Bergman atom clusters in the cubic unit cell [2002Hir, 2002Sug], the same clusters, two of which are in the unit cell of -1. The space group is Pa3. Lattice parameter (2310 pm) and composition (Mg46Zn37Al17) differ slightly from those given by [1995Tak]. [2001Bou2] confirmed the appearance of a quasicrystalline phase in Mg cast alloys containing 4 mass% Al and 8 mass% Zn. [2000Bok] reported a reversible transformation from a quasicrystalline phase to the 2/1 approximant in Mg44 Zn41Al15 at 340°C and a subsequent irreversible decomposition of the 2/1 approximant above 420°C. [1961Cla] indicated additionally the probability of the existence of a further ternary phase of undetermined composition near the Mg-Zn boundary in the region of the phases MgZn, Mg2Zn3 and MgZn2. This may be taken as an earlier hint to the two phases identified by [1995Tak]. Metastable precipitates are formed during low temperature annealing of supersaturated (Al) solid solutions, which were formed by quenching from temperatures with higher solubility of Mg and Zn in (Al). Guinier-Preston zones are formed at and slightly above room temperature. At somewhat higher temperatures a metastable phase, ´, is formed coherently in the Al matrix. Several models for its structure were proposed. [2001Wol] calculated the Gibbs energy of three models from first principles and proposed the model of [1974Aul] to be the most likely one. At even higher temperatures (above about 200°C) also the stable phases and -1 may precipitate. The crystallographic data of the stable solid phases are summarized in Table 1. Pseudobinary Systems Data are available in the literature on the following sections Al-MgZn2 [1926San], MgZn2-Mg2Al3 [1936Ham1], MgZn2--1 [1936Koe1], Al--1 [1936Koe1] and -1-Mg17Al12 [1936Koe2]. Essentially, there are three reports on sections from MgZn2 through the ternary -1 phase to the Al-Mg side, in two of these works, [1913Ege, 1936Koe2], the section ends at the Al-Mg side in Mg17Al12 whereas that of [1936Ham1] ends at Mg2Al3. In [1913Ege], only one of the phases, Mg17Al12, is shown in the Al-Mg system. It should be noted that the composition of the -1 phase which is associated with the pseudobinary equilibria at p 1 (L+MgZn2-1), e1 (L(Al)+-1), e3 (L +-1) and e2 (L+-1) varies considerably. This variation in composition reflects the very wide homogeneity range of the -1 phase. The section Mg2Al3-MgZn2 (-MgZn2) is nearly a true quasibinary section with the peritectic point p 1, the eutectic point e2 and the compositions of -1 associated with these two invariant equilibria lying very near to the vertical plane of section [1936Ham1]. Fig. 1 shows a calculation of this section using the dataset of [1998Lia2]. The (MgZn2) homogeneity range is very narrow along constant Mg content and thus extends outside the plane of section. Therefore the corner of the three-phase field L++-1 in Fig. 1 markedly deviates from true quasibinary behavior. However, if the plane of section near this area would be shifted a few tenth of at.% towards more Mg, it would look truly quasibinary. The section from Mg17Al12 to MgZn 2 ( to MgZn 2), although regarded as a quasibinary section by [1913Ege], was shown by [1936Koe1, 1936Koe2] not to be a quasibinary. The tie line joining Mg17Al12-e3--1 at 450°C deviates significantly from the section Mg17Al12-MgZn2 [1936Koe3]. The calculation of this section using the dataset of [1998Lia2] is shown in Fig. 2. A special problem arises in connection with the Al-MgZn2 section. The section was first established in [1913Ege] with the eutectic point at 63Zn-11.6Mg-25.4Al (mass%) and at 473°C. [1923San] observed the MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn
193
same section when studying the Al-Mg-Zn system for the presence of quasibinary sections. In an alloy with approximately 84 mol% MgZn2 on the join Al-MgZn 2 the authors observed a clearly pronounced purely eutectic structure between Al and MgZn2 and thus confirmed the quasibinary nature of the system. [1926San] gave a picture of this quasibinary system with the eutectic point at approximately 63Zn-12Mg-25Al (mass%) and 475°C. [1936Ham1], when summarizing the results of his extensive work on the Al-Mg-Zn system, indicated that the Al-MgZn2 system can be treated as an independent binary system. [1936Koe1] established by microscopic studies the presence of the four-phase transition reaction L+-1(Al)+MgZn2. The quadrangle representing the four-phase region, however, is very nearly a triangle with the liquid phase lying close to the (Al)-MgZn2 join. According to [1986Luk] this type of reaction may be considered as a ternary degenerate equilibrium combining the four-phase equilibrium L+-1+(Al)+MgZn2 with the maximum of the three-phase equilibrium L+-1+(Al). [1943Sie] carried out metallographic and X-ray diffraction investigations and determined the solubility of MgZn2 in (Al). He found a maximum solubility of 17 mass% MgZn2 at 475°C, thus rejecting the result of 30% MgZn2 obtained by [1938Sal]. [1979Sti] gave results of measurements relating to the reaction L+-1(Al)+MgZn2. The four-phase region was represented in this work, as in [1936Koe1], by a triangle with the liquid phase apex on the side Al-MgZn2. The temperature of the reaction is the same, i.e. 475°C. (A misprint occurs in the text of [1936Koe1] and the temperature is given as 375°C). The composition of (Al) in the reaction is shifted substantially towards the Al side in disagreement with the results given in [1936Koe1]. In all the publications mentioned, [1926San], [1938Sal, 1943Sie], except [1913Ege], the temperature of the quasibinary reaction coincides with the temperature of the four-phase equilibrium L+-1(Al)+MgZn2 [1936Koe1, 1979Sti]. The calculation with the dataset of [1998Lia2] shows the liquid phase 2.1 at.% outside the triangle -1-(Al)-MgZn 2, therefore, contrary to [1993Pet], here this reaction is labelled as U type, not as degenerate. The calculated section Al-MgZn2 is shown in Fig. 3. Except near the MgZn2 phase the approximation of it as quasibinary is quite good. At lower temperatures down to about 410°C the three-phase equilibrium -1+(Al)+MgZn 2 crosses the section as its (Al) corner is slightly shifted from the plane of section to the more Zn rich side. The ternary solid solution of MgZn2 extends outside the section into the more Mg-rich side, therefore near MgZn2 the section cuts the three-phase equilibrium MgZn2+(Al)+(Al,Zn), going from the (Al)-miscibility gap of the binary Al-Zn system to MgZn2. Invariant Equilibria Three reaction schemes can be found in the literature. They relate respectively to the Zn corner [1936Koe1], the Mg corner [1936Koe3], and the Mg corner for compositions greater than 50 at.% Mg [1961Cla]. The first and the third schemes served as the basis of the scheme of invariant equilibria in [1986Des] however, the second scheme cannot be accepted because the later discovered ternary phase, 1 [1961Cla], is missing there. This phase over an appreciable range of composition prevents equilibrium between the ternary -1 phase and (Mg). [1988Ito] presented invariant equilibria based on isothermal sections in the Al-rich corner. Since these equilibria used old binary data they were rejected in this assessment. The thermodynamic assessments of two different groups, H. Liang et al. [1997Lia] and P. Liang et al. [1998Lia2], enable the calculation of a complete comprehensive set of invariant equilibria. In both assessments the stable quasicrystalline phase and their 2/1 approximant (-2) detected by [1995Tak] are missing. Below in the section “Thermodynamics” an attempt is described to incorporate these two phases into the dataset of [1998Lia2]. The invariant equilibria calculated from this updated dataset are listed in Table 2 and the corresponding reaction scheme is given in Figs. 4a and 4b. As the stability ranges of the quasicrystalline phase and -2 are known approximative only the reaction scheme has to be taken as partially tentative, indicated by dashed boxes for the reactions. Liquidus Surface [1913Ege] was the first to construct the liquidus surface for the entire field of the Al-Mg-Zn system. Besides (Al), (Mg) and (Zn), he detected only three more solid phases in the system: (Mg17Al12), MgZn2, and a ternary phase Mg7Zn6Al3. [1936Ham1] used thermal, microscopic and X-ray analysis and gave a completely different ternary homogeneity range than [1913Ege], shown in the form of a relatively narrow Landolt-Börnstein New Series IV/11A3
MSIT ®
194
Al–Mg–Zn
band displaced generally towards the Al-Mg system from the field given by [1913Ege]. In [1936Ham2], the MgZn2 field occupies an excessively wide region in the system. The ternary phase was confirmed in general in later works [1934Fus, 1936Koe1, 1936Koe2, 1936Koe3, 1936Fin, 1937Fin, 1940Ura, 1943But, 1943Lit1, 1943Lit2, 1945Mik, 1949Sal], but the ternary homogeneity range has been widened noticeably compared with [1913Ege, 1936Ham1, 1936Ham2] at the expense of the MgZn2 field. [1959Cla, 1961Cla] established the existence of another ternary phase, 1, and connected a new invariant transformation with this phase: L+ +-11 at 393°C; the nature of this transformation was determined by [1961Cla] as peritectic. The author claimed this as the first report of a ternary peritectic reaction in literature on alloy constitution. [1962Ale] investigated in detail the liquidus of the system in the Zn corner within 10 mass% Al and 8 mass% Mg. [1961Yue] used the method of zone melting to determine eutectic compositions in complex metallic systems and established the existence of a ternary eutectic of the composition 50Zn-3Al-47Mg (mass%) at 338°C in the Al-Mg-Zn system. The authors repeated this composition in another work, [1970Yue]. It should be noted, however, that the schematic ternary eutectic equilibrium L(Mg)+ +MgZn suggested by [1970Yue] was earlier given by [1961Cla] as L(Mg)+MgZn+-1. [1973Wil], in a brief but rather detailed review, gave the liquidus surface of the whole Al-Mg-Zn system. The data of the reviews by [1971Mon, 1976Mon] generally agree with those of [1973Wil]. The liquidus surface of [1973Wil] has been adopted virtually without changes in [1986Des]. [1985Kuz1] investigated by thermal, microscopic and X-ray methods the Al-Zn region of the system and corrected to a certain extent the liquidus region of (Mg2Al3). It is narrowed somewhat at the Al-Mg side and spreads to 15 mass% Zn as compared with the results of [1973Wil]. [1986Kuz] studied the portion of the liquidus surface in the Al corner of the system up to 30 mass% Mg and 30 mass% Zn and repeated to a certain extent the results of [1985Kuz1]. The liquidus surface calculated by [1998Lia2] differs in some details from that constructed in the review of [1993Pet]: the line of double saturation of the liquid with -1 and MgZn2 is markedly curved towards the MgZn2 phase, the equilibria between liquid, MgZn 2, -1, Mg2Zn3 and MgZn are interchanged and due to the updated Al-Mg system [1998Lia1] the Al-Mg side is simplified. Taking into account the stability of the quasicrystalline phase q (Mg44Zn 41Al15) as described below in section “Thermodynamics” also a small field of primary solidification of q appears. The liquidus surface calculated from this updated dataset is shown in Fig. 5. Solidus and Solvus Surfaces [1936Koe1, 1936Koe2, 1936Koe3] presented the solidus polytherm of the Al corner. The solidus polytherm in [1952Han] is displaced sharply to substantially lower concentrations of Mg along its whole length as compared with the polytherm in [1936Koe3]. [1945But] constructed a series of solidus isotherms in the temperature interval 630 to 500°C. [1973Wil] gave a generalized solidus surface of the Al corner of the system, which was adopted in the review of [1993Pet]. The solidus isotherms of [1979Sti, 1985Kuz1, 1985Kuz2] as well as the calculated ones of [1998Lia2] virtually coincide with those of [1973Wil]. Figure 6 shows the solidus isotherms of the (Al) phase calculated from the dataset of [1998Lia2]. The solubility of Zn and Mg in (Al) was studied by a number of authors [1933Boc, 1945But, 1947Str, 1955Zam, 1961Sal, 1971Mon]. Their results are compared and presented in the solvus surface of [1973Wil], which was adopted in the review of [1993Pet]. The solvus of the (Al) phase, calculated from the dataset of [1998Lia2] (Fig. 7) agrees very well with that of [1993Pet]. The dataset of [1998Lia2] enables also a calculation of solidus and solvus isotherms of the (Mg) phase. They are shown in Fig. 8. To allow different scaling of the composition axes of Al and Zn, in this figure not Gibbs triangular, but Cartesian rectangular coordinates are chosen. The existence of equilibria between the (Mg) solid solution and the stable quasicrystalline phase is supported by [1995Tak] and [2001Bou2]. The extension of the fields of these equilibria, however, must be taken as tentative.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn
195
Isothermal Sections [1936Fin, 1937Fin] used microscopic and X-ray methods and measurements of electrical resistance to construct a series of isothermal sections in the Al corner in the range up to 20 mass% Mg and 20 mass% Zn over the temperature interval 400 to 200°C, and an isothermal section adjoining the Al-Zn side of the system in the range up to 40 mass% Mg. Similar work was carried out by [1943Lit1, 1943Lit2] who constructed a series of isothermal sections in the Al corner in the range up to 12 mass% Mg and 12 mass% Zn. [1971Mon] gave isothermal sections in the Al corner in the range up to 16 mass% Mg and 16 mass% Zn and compared it with the results of [1936Fin, 1937Fin]. Of the highest interest is the work of [1961Cla] who constructed the isothermal section of the whole ternary system at 335°C using metallographic and X-ray methods and thermal analysis on heating and cooling. This temperature is only 3°C below the lowest solidus temperature of the system. The ternary 1 phase was discovered in the same study. [1961Cla] studied the nature of the invariant equilibria associated with this ternary phase and constructed another two schematic isothermal sections at 394 and 374°C which enabled him to establish the equilibrium L+ +-11 at 393°C. A number of other equilibria involving the 1 phase were investigated so that a reaction scheme could be constructed, as mentioned in the section “Invariant Equilibria”. [1973Wil] in his review reproduced completely the isothermal section at 335°C taken from [1961Cla]. [1986Des] proposed some corrections for the isothermal section of [1961Cla]. [1998Lia2] investigated 34 samples annealed at 335°C mainly to fix the maximum ternary solubilities of the binary phases. These alloys were prepared by induction melting in a graphite crucible placed inside a copper crucible. The samples were re-melted twice to ensure homogeneity, and precautions were taken to reduce Zn and Mg loss. The phase compositions as measured by EPMA are largely in agreement with previous work [1986Des] apart from the Al solubilities in MgZn2, Mg2Zn3 and MgZn, which are of the same magnitude, whereas in the review of [1993Pet] a much lower Al solubility was assumed for Mg2Zn 3. The 335°C isothermal section in Fig. 9 is based on the calculated one of [1998Lia2] with corrections to include the phases q and -2. Thermodynamics Kim et al. [1997Kim] determined the enthalpies of mixing of the liquid phase over the temperature range 610-660°C using high temperature calorimetry. From the results thermodynamic functions of the liquid were assessed using the associated solution model. A thermodynamic description of the solidus and liquidus surfaces in the Al-rich corner has been evaluated by [1990Kuz]. Thermodynamic datasets covering the whole ternary system were assessed by two groups, [1997Lia, 1998Lia2]. The first group used more simplified models for the description of the phases. Ternary solubilities in binary phases are considered only for liquid, (Al), (Mg), (Zn) and the Laves phase MgZn2, 1 is described as a stoichiometric phase and -1 as a line compound with constant Mg content of 39.5 at.%, Mg32(Zn,Al)49. Apart from the restrictions due to these simplifications both assessments do not deviate severely. [1998Lia2] used own measurements to adjust the Al solubilities in Mg2Zn11, MgZn2, Mg2Zn3 and MgZn, as well as the Zn solubilities in the , and J phases of the Al-Mg system. These phases, except MgZn2 and are described as line compounds Mg x(Zn,Al)y, 1 is also described as a line compound, -1 is done by the model Mg26(Mg,Al)6(Mg,Zn,Al)48 Al1, reproducing fairly well the experimental homogeneity range. In the present assessment an attempt was made to incorporate the stable quasicrystal (q) and the 2/1 approximant (-2) into the dataset of [1998Lia2], both described as stoichiometric phases. The Gibbs energies per mole of atoms are expressed as: Gq - 0.15#GAlfcc - 0.44#GMghcp - 0.41#GZnhcp = -7760. -1.#T (J#mol-1) and G-2 - 0.15#GAlfcc - 0.43#GMg hcp - 0.42#GZnhcp = -7900. -1.#T (J#mol-1) Also the description of the 1 phase was updated to satisfy the formula Mg21(Zn,Al) 17 given by [2001Bou1]: GMg:Al1 - 17 #GAlfcc - 21#GMghcp = 38# (-1380. -1.5#T) (J#mol-1) GMg:Zn1 - 21#GMghcp - 17#GZnhcp = 38# (-7150. +1.9#T) (J#mol-1) LMg:Zn,Al1 = 38# (-2100 +1.#T) (J#mol-1)
Landolt-Börnstein New Series IV/11A3
MSIT ®
196
Al–Mg–Zn
This updated dataset was used to calculate all figures and tables of this assessment. With respect to the phases q and -2 it has to be considered as tentative. Miscellaneous [1968Tha] investigated, by electron diffraction, the process of decomposition on long term ageing of quenched Al-6Zn-2Mg (mass%) alloy and discovered for the first time the formation of an intermediate -' phase with a hexagonal cell and lattice parameters a = 267 pm and c = 490 pm. [1976Aul] investigated, by X-ray diffraction, single crystals of the alloy Al-4.0Zn-2.9Mg (mass%) after quenching from 490°C followed by ageing at 155°C for 24 h and also demonstrated the formation of an intermediate phase called -' with a hexagonal cell and the lattice parameters a = 1388 pm and c = 2752 pm. Long holding at 230°C (for 200 d) of an Al-Mg-Zn alloy led to the formation of the exclusively equilibrium -1 phase. The influence of predeformation on the precipitation and resulting mechanical properties were studied by [1997Des] for the alloy Al-6.1Zn-2.35Mg (mass%). The formation and decomposition of metastable quasicrystalline phases were investigated in various rapidly solidified alloys [1986Cas, 1986Raj, 1986Sas, 1988Cha, 2000Bok, 2000Miz, 2000Tak, 2001Bou2]. The resulting internal melting of several Al-Mg-Zn alloys was studied by rapid quenching in a salt bath [1994Dro]. References [1913Ege] [1923San] [1926San] [1929Ish] [1930Ish] [1933Boc]
[1934Fus] [1935Lav]
[1936Fin] [1936Ham1] [1936Ham2]
[1936Koe1]
[1936Koe2]
MSIT®
Eger, G., “Studies on the Constitution of the Ternary Mg-Al-Zn Alloys” (in German), Int. Z. Metallographie, 4, 50-128 (1913) (Equi. Diagram, Experimental, *, 9) Sandner, W., Meissner, K.L., “Equilibrium Studies in the Al-Mg-Si-Zn Quarternary System”(in German), Z. Metallkd., 15, 160-183 (1923) (Equi. Diagram, Experimental, 11) Sandner, W., “The Effect of the Compound MgZn2 on the Improvement of Al Alloys” (in German), Z. Anorg. Allg. Chem., 154, 144-151 (1926) (Equi. Diagram, Experimental, 11) Ishida, S., “On Some Magnesium Base Light Alloys” (in Japanese), Nippon Kogyo Kwai Shi, 45, 256-268, 611-621, 786-790 (1929) (Equi. Diagram, Experimental, 9) Ishida, S., “On Some Magnesium Base Light Alloys” (in Japanese), Nippon Kogyo Kwai Shi, 46, 245 (1930) (Equi. Diagram, Experimental, 9) Bochvar, A.A., Kuznetzov, M.D., “The Transformations in Solid Alloys of Aluminium with up to 30% Zinc and 12% Magnesium” (in Russian), Metallurg, 8, 7-14 (1933) (Equi. Diagram, Experimental, 5) Fuss, V., “Metallography of Al and Its Alloys” (in German), Verlag Julius Springer, Berlin, 154-155 (1934) (Equi. Diagram, Review, 300) Laves, F., Lohberg, K., Witte, H., “On the Isomorphy of the Ternary Compounds Mg3Zn3Al2 and Mg4CuAl6” (in German), Metall-Wirtschaft, 14, 793-794 (1935) (Crys. Structure, Experimental, 6) Fink, W.L., Willey, L.A., “Equilibrium Relations in Aluminium-Magnesium-Zinc Alloys of High Purity”, Met. Technol., 8, 3-32 (1936) (Equi. Diagram, Experimental, 21) Hamasumi, M., “The Complete Equilibrium Diagram of the System Al-Mg-Zn”, Sci. Rep. Tohoku Imp. Univ., 748-776 (1936) (Equi. Diagram, Experimental, *, 8) Hamasumi, M., “Ternary Diagram of the Aluminium-Magnesium-Zinc System” (in Japanese), Tetsu to Hagane (J. Iron Steel Inst. Jpn.), 22, 258-271 (1936) (Equi. Diagram, Experimental, *, 25) Koester, W., Wolf, W., “The Al-Mg-Zn Ternary System. I. The Partial Region Al-Al2Mg3Zn3-MgZn2-Zn” (in German), Z. Metallkd., 28, 155-158 (1936) (Equi. Diagram, Experimental, *, 11) Koester, W., Dullenkopf, W., “The Al-Mg-Zn Ternary System. II. The Partial Region Al-Al3Mg4-Al2Mg3Zn 3-Al” (in German), Z. Metallkd., 28, 309-312 (1936) (Equi. Diagram, Experimental, *, 6)
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn [1936Koe3]
[1936Rie] [1937Fin]
[1938Sal] [1940Ura]
[1943But]
[1943Lit1]
[1943Lit2]
[1943Sie] [1945But]
[1945Mik] [1947Str]
[1949Sal] [1952Ber]
[1952Han]
[1955Zam] [1957Ber]
[1959Cla] [1961Cla] [1961Sal] [1961Wri]
Landolt-Börnstein New Series IV/11A3
197
Koester, W., Dullenkopf, W., “The Al-Mg-Zn Ternary System. III. The Partial Region Mg-Al3Mg4-Al2Mg3Zn3-MgZn2-Mg” (in German), Z. Metallkd., 28, 363-367 (1936) (Equi. Diagram, Experimental, *, 1) Riederer, K., “The Al-Mg-Zn System” (in German), Z. Metallkd., 28, 312-317 (1936) (Crys. Structure, Experimental, *, 12) Fink, W.L., Willey, L.A., “Equilibrium Relations in Aluminium-Magnesium-Zinc Alloys of High Purity”, Trans. Met. Soc. AIME, 124, 78-110 (1937) (Equi. Diagram, Experimental, 21) Saldau, P.Ya., Zamotorin, M.J., “Aging Phenomena in Aluminium-MgZn 2 Alloys” (in Russian), Izv. Sekt. Fiz. -Khim. Anal., 11, 27-36 (1938) (Equi. Diagram, Experimental, 2) Urasov, G.G., Filin, N.A., Shashin, A.B., “The Liquidus of the System Aluminium-Magnesium-Zinc” (in Russian), Metallurg, 15, 3-11 (1940) (Equi. Diagram, Experimental, 7) Butchers, E., Raynor, G.V., Hume-Rothery, W., “The Constitution of Mg-Mn-Zn-Al in the Range 0-5% Mg, 0-2% Mn, 0-8% Zn. I. The Liquidus”, J. Inst. Met., 69, 209-228 (1943) (Equi. Diagram, Experimental, 9) Little, D.T., Raynor, G.V., Hume-Rothery, W., “The Constitution of Mg-Mn-Zn-Al Alloys in the Range 0-5% Mg, 0-2% Mn and 0-8% Zn. III. The 500°C and 400°C Isothermals”, J. Inst. Met., 69, 423-440 (1943) (Crys. Structure, Equi. Diagram, Experimental, 8) Little, D.T., Raynor, G.V., Hume-Rothery, W., “The Constitution of Mg-Mn-Zn-Al Alloys in the Range 0-5% Mg, 0-2% Mn and 0-8% Zn. IV. The Equilibrium Diagram Below 400°C”, J. Inst. Met., 69, 467-484 (1943) (Crys. Structure, Equi. Diagram, Experimental, 9) Siebel, G., “Determination of the Solubility of MgZn2 in Al” (in German), Z. Elektrochem., 49, 218-221 (1943) (Equi. Diagram, Experimental, 7) Butchers, E., Hume-Rothery, W., “The Constitution of Aluminum-MagnesiumManganese-Zinc Alloys. The Solidus”, J. Inst. Met., 71, 291-311 (1945) (Equi. Diagram, Experimental, 8) Mikheeva, V.I., Kryukova, O.S., “The Liquidus of the Al-Mg-Zn System” (in Russian), Dokl. Akad. Nauk SSSR, 50, 234-247 (1945) (Equi. Diagram, Experimental, 9) Strawbridge, D.I., Hume-Rothery, W., Little, A.T., “The Constitution of Al-Cu-Mg-Zn Alloys at 460°C”, J. Inst. Met., 74, 191-225 (1947) (Crys. Structure, Equi. Diagram, Experimental, 11) Saldau, P.Y., “Equilibrium Diagram of the System Al-Mg-Zn” (in Russian), Izv. Sekt. Fiz.-Khim. Anal., 19, 487-496 (1949) (Equi. Diagram, Review, 22) Bergman, G., Waugh, L.T., Pauling, L., “Crystal Structure of the Intermetallic Compound Mg32 (Al,Zn)49 and Related Phases”, Nature, 169, 1057-1058 (1952) (Crys. Structure, Experimental, *, 4) Haneman, H., Schrader, A., “Ternary Al Alloys” (in German), Atlas Metallographicus, Verlag Stahleisen M.B.H. Düsseldorf, Vol. 3, Part 2, 133-149 (1952) (Crys. Structure, Equi. Diagram, Review, 24) Zamotorin, M.I., “Mutual Solubility of Mg and Zn in Al in the Solid State” (in Russian), Tr. Leningrad. Politekhn. Inst., 180, 38-43 (1955) (Equi. Diagram, Experimental, 10) Bergman, G., Waugh, L.T., Pauling, L., “The Crystal Structure of the Metallic Phase Mg32 (Al,Zn)49”, Acta Crystallogr., 10, 254-259 (1957) (Crys. Structure, Experimental, *, 20) Clark, J.B., Rhines, F.N., “Diffusion Layer Formation in the Ternary System Al-Mg-Zn”, Trans. ASM, 51, 199-221 (1959) (Crys. Structure, Equi. Diagram, Experimental, *, 10) Clark, J.B., “Phase Relations in the Mg-Rich Region of the Mg-Al-Zn Phase Diagram”, Trans. ASM, 53, 295-306 (1961) (Crys. Structure, Equi. Diagram, Experimental, #, *, 11) Saldau, P.Ya., “The Mutual Solubility of Mg and Zn in Al” (in Russian), Publishing House (Defence), Moscow, 5-8 (1961) (Equi. Diagram, Review, 10) Wright, E.H., Private Communication to [1961Cla], Aluminium Research Laboratories, Aluminum Company of America (1961) (Crys. Structure, Experimental, *) MSIT ®
198 [1961Yue]
[1962Ale] [1968Tha] [1970Yue] [1971Mon] [1973Wil] [1974Aul] [1976Aul] [1976Mon]
[1979Sti]
[1985Kuz1]
[1985Kuz2]
[1986Cas]
[1986Des]
[1986Kuz]
[1986Luk] [1986Raj]
[1986Sas] [1988Cha]
[1988Ito]
MSIT®
Al–Mg–Zn Yue, A.S., Clark, J.B., “The Determination of the Eutectic Composition by the Zone-Melting Method”, Trans. Met. Soc. AIME, 22, 383-389 (1961) (Equi. Diagram, Experimental, 9) Aleksakhin, I.A., Semionow, A.A., “Zn-Corner of the Zn-Al-Mg System” (in Russian), Metall. Term. Obra. Met., 4, 40-45 (1962) (Equi. Diagram, Experimental, 8) Thackery, P.A., “The Nature and Morphology of Precipitate in Al-Zn-Mg Alloys”, J. Inst. Met., 96, 228-235 (1968) (Crys. Structure, Experimental, *, 16) Yue, A.S., “Determination of Eutectic Compositions in Complex Metall Systems”, Metall. Trans., 1, 19-22 (1970) (Equi. Diagram, Experimental, 13) Mondolfo, L.F., “Structure of Aluminum: Magnesium: Zinc Alloys”, Met. Rev., 16, 95-124 (1971) (Equi. Diagram, Experimental, Review, 621) Willey, L.D., “Al-Mg-Zn (Aluminium-Magnesium-Zinc)”, Metals Handbook, American Society for Metals, Metals Park, Ohio, 8, 397-399 (1973) (Equi. Diagram, Review, #, *, 15) Auld, J.H., Cousland, S. McK., “The Structure of the Metastable ' Phase in Al-Zn-Mg Alloys”, J. Australian Inst. Met., 19, 194-199 (1974) (Crys. Structure, Experimental) Auld, J.H., Cousland, S. McK., “The Metastable T' Phase in Al-Zn-Mg and Al-Ag-Mg Alloys”, Met. Sci., 73, 445-448 (1976) (Crys. Structure, Experimental, 10) Mondolfo, L.F., “Aluminium - Magnesium - Zinc System”, in “Aluminium Alloys; Structure and Properties”, Butterworth and Co., London, 575-590 (1976) (Equi. Diagram, Review, 170) Stiller, W., Hoffmeister, H., “Determination of the Liquid/Solid Phase Equilibria of Al-Mg-Zn Alloys” (in German), Z. Metallkd., 70, 817-824 (1979) (Equi. Diagram, Experimental, 35) Kuznetsov, G.M., Barsukov, A.D., Krivosheeva, G.B., Istomin- Kastorovski, V.V., “A Study of Phase Equilibria and Solidification Processes in Al-Zn-Mg Alloys” (in Russian), Izv. Vyss. Uchebn. Zaved. Tsvetn. Metall., (1), 88-93 (1985) (Equi. Diagram, Experimental, Thermodyn., 13) Kuznetsov, G.M., Barsukov, A.D., Krivosheeva, G.B., Bashashkina, E.V., “Study of Aluminium-Zinc-Magnesium Alloys” (in Russian), Izv. Vyss. Uchebn. Zaved. Tsvetn. Metall., (2), 91-95 (1985) (Equi. Diagram, Experimental, Thermodyn., 13) Cassada, W.A., Shen, Y., Poon, S.J., Shiflet, G.J., “Mg 32(Zn,Al)49-Type Icosahedral Quasicrystals Formed by Solid-State Reaction and Rapid Solidification”, Phys. Rev. B, B34, 7413-7416 (1986) (Experimental, Crys. Structure, 17) Despande, N.U., Ray, K.K., Mallik, A.K., “The Aluminium-Magnesium-Zinc System”, J. Alloy Phase Diagrams, 2, 108-130 (1986) (Crys. Structure, Equi. Diagram, Review, Thermodyn., *, 40) Kuznetsov, G.M., Barsukov, A.D., Krivosheeva, G.B., Dieva, E.G., “Phase Equilibria in Al-Zn-Mg Alloys” (in Russian), Izv. Akad. Nauk SSSR, Met., (4), 198-200 (1986) (Equi. Diagram, Experimental, Thermodyn., 7) Lukas, H.L., Henig, E.T., Petzow, G., “50 Years Reaction Scheme after Erich Scheil” (in German), Z. Metallkd., 77, 360-367 (1986) (Equi. Diagram, Theory, 7) Rajasekharan, T., Akhtar, D.A., Copalan, R., Muraledharan, K., “The Quasi-Crystalline Phase in the Mg-Al-Zn System”, Nature, 322, 528-530 (1986) (Crys. Structure, Experimental, 7) Sastry, G.V.S., Ramaxhandrarao, P., “A Study of the Icosahedral Phase Mg 32(Al,Zn)49”, J. Mater. Res., 1, 247-250 (1986) (Experimental, Crys. Structure, 11) Chandra, S., Suryanarayana, C., “Quasicrystalline to Crystalline Transformation in Rapidly Solidified Mg32(Al,Zn)49”, Philos. Mag., 58, 185-202 (1988) (Crys. Structure, Experimental, 30) Itoh, G., Eto, T., Miyagi, Y., Kanno, M., “Al-Zn-Mg Alloys” (in Japanese), J. Jpn. Inst. Light Met., 38, 818-839 (1988) (Equi. Diagram, Phys. Prop., Experimental, 178)
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn [1990Kuz]
[1993Pet]
[1994Dro] [1995Tak]
[1997Des]
[1997Don]
[1997Lia]
[1997Kim]
[1998Ans]
[1998Lia1]
[1998Lia2]
[2000Bok]
[2000Lee]
[2000Miz]
[2000Sun]
Landolt-Börnstein New Series IV/11A3
199
Kuznetsov, G.M., Ramazanov, S.M., Krivosheeva, G.B., “Phase Equilibria in Alloys of the Aluminium-Zinc-Magnesium System”, Izv. Vyss. Uchebn. Zaved. Tsvetn. Metall., 3, 97-101 (1990) (Thermodyn., Experimental, 5) Petrov, D.A., “Aluminium-Magnesium-Zink”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.11491.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 53) Droenen, P.-E., Ryum, N., “Local Melting in Al-Mg-Zn Alloys”, Metall. Mater. Trans. A, 25A, 521-530 (1994) (Equi. Diagram, Experimental, 23) Takeuchi, T., Mizutani, U., “Electronic Structure, Electron Transport Properties and Relative Stability of Icosahedral Quasicrystals and their 1/1 and 2/1 Approximants in the Al-Mg-Zn Alloy System”, Phys. Rev. B, 52B, 9300-9309 (1995) (Crys. Structure, Equi. Diagram, Experimental, 27) Deschaps, A., Brechet, Y., Guyot, P., Livet, F., “On the Influence of Dislocations on Precipitation in an Al-Zn-Mg Alloy”, Z. Metallkd., 88(8), 601-606 (1997) (Experimental, Equi. Diagram, Review, 24) Donnadieu, P., Quivy, A., Tarfa, T., Ochin, P., Dezellus, A., Harmelin, M.G., Liang, P., Lukas, H.L., Seifert, H.J., Aldinger, F., Effenberg, G., “On the Crystal Structure and Solubility Range of the Ternary 1 Phase in the Mg-Al-Zn System”, Z. Metallkd., 88(12), 911-916 (1997) (Crys. Structure, Experimental, 16) Liang, H., Chen, S.-L., Chang, Y.A., “A Thermodynamic Description of the Al-Mg-Zn System”, Met. Mater. Trans., 28A, 1725-1734 (1997) (Equi. Diagram, Thermodyn., Assessment, 75) Kim, Y.B., Sommer, F., Predel, B., “Calorimetric Investigation of Liquid Aluminium-Magnesium-Zinc Alloys”, J. Alloys Compd., 247, 43-51 (1997) (Thermodyn., Theory, Experimental, 20) Liang, P., Lukas, H.-L., “System Al-Mg-Zn” in ”COST 507, Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H. (Eds.), European Communities, Luxembourg, Vol. 2, 329-332 (1998) (Equi. Diagram, Thermodyn., Assessment, 0) Liang, P., Su, H.-L., Donnadieu, P., Harmelin, M.G., Quivy, A., Ochin, P., Effenberg, G., Seifert, H.J., Lukas, H.-L., Aldinger, F., “Experimental Investigation and Thermodynamic Calculation of the Central Part of the Mg-Al Phase Diagram”, Z. Metallkd., 89, 536-540 (1998) (Equi. Diagram, Thermodyn., Experimental, Assesssment, *, #, 33) Liang, P., Tarfa, T., Robinson, J. A., Wagner, S., Ochin, P., Harmelin, M.G., Seifert, H.J., Lukas, H.L., Aldinger, F., “Experimental Investigation and Thermodynamic Calculation of the Al-Mg-Zn System”, Thermochim. Acta, 314, 87-110 (1998) (Equi. Diagram, Thermodyn., Experimental, Assessment, *, #, 69) Bokhonov, B.B., Ivanov, E.Y., Tolochko, B.P., Sharaphutdinov, M.P., “In Situ Study of Structural Transformations of Mg44Al15Zn41 Quasicrystals under Heating”, Mater. Sci. Eng. A, A278, 236-241 (2000) (Crys. Structure, Experimental, 8) Lee, C.-S., Miller, G.J., “Where are the Elements in Complex Aluminides? An Experimental and Theoretical Investigation of the Quasicrystalline Approximant, Mg2-y(ZnxAl1-x)3+y”, J. Am. Chem. Soc., 122, 4937-4947 (2000) (Crys. Structure, Experimental, Theory, 72) Mizutani, U., “Electron Transport Mechanismin the Pseudogap System: Quasicrystals, Approximants and Amorphous Alloys”, Mater. Sci. Eng. A, A294-296, 464-469 (2000) (Crys. Structure, Experimental, Theory, 23) Sun, W., Lincoln, F.J., Sugiyama, K., Hiraga, K., “Structure Refinement of (Al,Zn) 49Mg32-Type Phases by Single-Crystal X-Ray Diffraction”, Mater. Sci. Eng. A, A294-296, 327-330 (2000) (Crys. Structure, Experimental, 9)
MSIT ®
Al–Mg–Zn
200 [2000Tak]
[2001Bou1]
[2001Bou2]
[2001Wol]
[2002Hir]
[2002Sug]
[2003Luk]
Takeuchi, T., Mizuno, T., Banno, E., Mizutani, U., “Magic Number of Electron Concentration in the Icosahedral Cluster of AlxMg40X60-x (X = Zn, Cu, Ag, and Pd) 1/1 Cubic Approximants”, Mater. Sci. Eng. A, A294-296, 522-526 (2000) (Crys. Structure, Experimental, Theory, 14) Bourgeois, L., Muddle, B.C., Nie, J.F., “The Crystal Structure of the Equilibrium 1 Phase in Mg-Zn-Al Casting Alloys”, Acta Mater., 49, 2701-2711 (2001) (Crys. Structure, Experimental, Theory, 72) Bourgeois, L., Mendis, C.L., Muddle, B.C., Nie, J.F., “Characterization of Quasicrystalline Primary Intermetallic Particles in Mg-8wt% Zn-4 wt% Al Casting Alloy”, Philos. Mag. Lett., 81, 709-718 (2001) (Crys. Structure, Experimental, 33) Wolverton, C., “Crystal Structure and Stability of Complex Precipitate Phases in Al-Cu-Mg-(Si) and Al-Zn-Mg Alloys”, Acta Mater., 49, 3129-3142 (2001) (Crys. Structure, Theory, 64) Hiraga, K., Sugiyama, K., Ishi, Y., “Arrangement of Atomic Clusters in a 2/1 Cubic Approximant in the Al-Zn-Mg Alloy System”, Philos. Mag. Lett., 82, 341-347 (2002) (Crys. Structure, Experimental, 21) Sugiyama, K., Sun, W., Hiraga, K., “Crystal Structure of a Cubic Al17Zn 37Mg46; a 2/1 Rational Approximant Structure for the Al-Zn-Mg Icosahedral Phase”, J. Alloys Compd., 342, 139-142 (2002) (Crys. Structure, Experimental, 8) Lukas, H.L., Lebrum, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] , (Al) < 660.45 (Mg) < 650 (Zn) < 419.58 , Mg 2Al3 452 , Mg17Al12 < 458 J, Mg23Al30 410 - 250
, Mg 51Zn20 342 - 325
MSIT®
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu hP2 P63/mmc Mg hP2 P63/mmc Mg cF1168 Fd3m Mg2Al3 cI58 I43m Mn hR159 R3 Mn44Si9 oI142 Immm Mg51Zn20
Lattice Parameters Comments/References [pm] a = 404.96
at 25°C [Mas2]
a = 320.94 c = 521.07
at 25°C [Mas2]
a = 266.50 c = 494.70
at 25°C [Mas2]
a = 2816 to 2824
60-62 at.% Al [2003Luk] 1168 atoms on 1704 sites per unit cell [2003Luk] at 41.4 at.% Al [V-C2] 39.5 to 51.5 at.% Al [2003Luk]
a = 1054.38
a = 1282.54 c = 2174.78
54.5 to 56.5 at.% Al [2003Luk]
a = 1408.3 b = 1448.6 c = 1402.5
[Mas2], called Mg7Zn 3, lattice parameters for Mg72Zn28 [V-C]
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn Phase/ Temperature Range [°C] MgZn < 347 , Mg2Zn3 < 416 , MgZn2 < 590
Pearson Symbol/ Space Group/ Prototype -
mC110 B2/m hP12 P63/mmc MgZn2 cP39 , Mg2Zn11 < 381 Pm3 Mg2Zn11 * -1, Mg32 (Zn,Al)48 cI160 Im3 Mg32(Zn,Al)48 * 1, Mg21(Zn,Al)17 oP152 Pbcm Mg21(Zn,Al)17 * -2, Mg43Zn42Al15 or cP640 (?) Mg46Zn37Al17 Pa3 * q, Mg44Zn41 Al15
201
Lattice Parameters Comments/References [pm] -
[Mas2]
-
[Mas2]
a = 522.1 c = 856.7
[Mas2, V-C]
a = 855.2
[Mas2, V-C]
a = 1413 - 1471
[1957Ber] (gave cI162) [2000Lee, 2000Sun] 1/1 approximant of icosahedral phase [2001Bou1] lattice parameters from [1997Don]
a = 897.9 b = 1698.8 c = 1934 a = 2291 a = 2310
quasicrystalline, icosahedral
[1995Tak] [2002Hir, 2002Sug] 2/1 approximant of icosahedral phase [1995Tak] stable quasicrystalline phase
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
L + MgZn2 -1
530
p1 (max)
L (Al) + -1
480
e1 (max)
L + -1 (Al) + MgZn 2
476
U1
L + -1
451
e2 (max)
L + -1
451
e3 (max)
L + + -1
449
E1
L MgZn 2 -1 L (Al) -1 L -1 (Al) MgZn 2 L -1 L -1 L
Landolt-Börnstein New Series IV/11A3
Composition (at.%) Al Mg Zn 21.8 37.0 41.2 7.0 33.5 59.5 18.2 36.2 45.6 51.1 21.4 27.5 90.2 4.2 5.6 23.2 33.8 43.0 5.1 18.7 36.2 33.0 47.5 19.5 89.6 2.9 7.5 7.5 33.3 59.2 55.4 39.2 5.4 56.6 38.9 4.5 48.1 41.0 10.9 46.6 47.0 6.4 46.5 49.0 4.5 46.0 42.4 11.6 52.0 43.1 4.9 59.5 38.9 1.6 50.6 47.6 1.8
MSIT ®
Al–Mg–Zn
202 Reaction
T [°C]
Type
Phase
L (Al) + + -1
447
E2
L +MgZn2 + -1 Mg 2Zn3
434
P1
L + + -1 1
388
P2
L + -1 q
380
p5 (max)
q + -1 -2
377
p6 (max)
L + (Mg) + 1
368
U2
L + -1 q + 1
362
U3
L + -1 Mg2Zn3+ q
362
U4
q + -1 Mg2Zn3 +-2
356
U5
L + MgZn2 Mg2Zn11 + (Al)
355
U6
L + q + Mg2Zn3 MgZn
353
P3
L (Al) -1 L MgZn 2 -1 Mg2Zn3 L -1 1 L q q -1 -2 L (Mg) 1 L -1 1 L -1 Mg2Zn3 -1 Mg2Zn3 L MgZn 2 Mg2Zn11 (Al) L
Mg2Zn3 L + 1 (Mg) + q
345
U7
L (Al,Zn) + Mg2Zn11 + (Zn)
344
E3
MSIT®
MgZn L 1 (Mg) L (Al,Zn) Mg2Zn11 (Zn)
Composition (at.%) Al Mg Zn 60.1 34.3 5.6 85.8 13.2 1.0 4.7 56.4 38.9 47.7 40.1 12.2 5.6 60.6 33.8 3.6 33.8 62.6 49.0 10.7 40.3 5.0 40.0 55.0 17.4 65.8 16.8 34.1 58.1 7.8 25.8 30.1 44.1 12.0 56.0 32.0 7.2 66.7 26.1 15.8 41.8 42.4 15.0 44.0 41.0 15.0 44.0 41.0 15.0 41.7 43.3 15.0 43.0 42.0 14.0 69.8 16.2 6.0 92.2 1.8 32.1 60.1 7.8 19.3 25.5 55.2 8.7 68.4 22.9 19.7 42.7 37.6 19.9 55.3 24.8 4.7 67.9 27.4 12.1 41.5 46.4 5.2 40.0 54.8 12.3 41.3 46.4 5.4 40.0 54.6 10.9 7.6 81.5 2.0 33.0 65.0 3.1 15.4 81.5 45.0 0.3 54.7 4.4 68.7 26.9 5.1 40.0 54.9 5.1 48.0 46.9 6.5 70.1 23.4 18.3 55.3 26.4 3.0 94.6 2.4 8.7 6.0 85.3 34.8 0.1 65.1 2.4 15.4 82.2 2.5 0.2 97.3
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn
203
Reaction
T [°C]
Type
Phase
L + Mg51Zn20 (Mg)+MgZn
340
U8
L q + (Mg) + MgZn
339
E4
L (Mg) MgZn L (Mg) MgZn MgZn 2 (Al,Zn) (Al) Mg2Zn11 (Al,Zn) (Al) (Zn) Mg2Zn11 Mg2Zn3 -1 MgZn 2 -2 q (Mg) 1 MgZn
MgZn2 + (Al,Zn) (Al) + Mg 2Zn 11 331
U9
(Al,Zn) (Al)+(Zn), Mg2Zn11
277
D1
Mg2Zn 3 + -1 MgZn 2 + -2
214
U10
q + (Mg) 1 + MgZn
84
U11
Fig. 1: Al-Mg-Zn. Calculated section from (Mg38.5Al62.8) to MgZn2
600
500
L+β
Temperature, °C
Composition (at.%) Al Mg Zn 0.1 71.0 28.9 0.05 97.0 2.95 0.1 48.0 51.9 4.4 70.1 25.5 2.0 95.3 2.7 5.3 48.0 46.7 1.8 33.1 65.1 52.3 0.2 47.5 23.2 76.7 0.1 2.9 15.4 81.7 41.0 0.06 58.94 85.8 0.03 14.17 98.3 1.6 0.1 2.1 15.4 82.5 6.1 40.0 53.9 39.0 45.2 15.8 3.8 33.4 62.8 15.0 43.0 42.0 15.0 44.0 41.0 0.4 99.3 0.3 19.3 55.3 25.4 47.6 4.4 48.0
590°C
L+τ1+β
L+(Al)+β
L
L+η
L+τ1
L+τ1+η
η
L+τ1+(Al)
400
(Al)+β
τ1+β +(Al)
τ1+(Al)
τ1
τ1+η
300
η+ζ τ1+η+ζ η+q τ1+η+q η+ζ +q
200
Mg 38.00 Zn 0.00 Al 62.00
Landolt-Börnstein New Series IV/11A3
10
20
30
Zn, at.%
40
50
60
Mg 33.33 Zn 66.67 0.00 Al
MSIT ®
Al–Mg–Zn
204
Fig. 2: Al-Mg-Zn. Calculated section from (Mg17Al12) to MgZn2
600
590°C
L+η
L 500
Temperature, °C
L+τ1
L+τ1+η
L+γ
τ1+τ2+q
L+γ+τ1
400
η
τ1+η
τ1
L+τ1+q
γ+τ1
L+η +ζ
τ1+ζ
τ1+τ2
ζ +η
τ1+η+ζ
γ 300
γ+φ +τ1
τ1+φ +q τ1+φ
τ1+τ2+ζ
200
τ1+τ2+η τ2+η
γ+φ Mg 58.62 Zn 0.00 Al 41.38
Fig. 3: Al-Mg-Zn. Calculated section from MgZn2 to Al
10
20
30
η+ζ +τ2 40
50
Mg 33.33 Zn 66.67 0.00 Al
60
Zn, at.%
700
660°C L 600
Temperature, °C
590°C
L+(Al) L+η
L+(Al)+τ1
500
η
(Al)+τ1+η
L+(Al)+η
(Al)
(Al)+τ1
400
L+η +θ
(Al)+η
(Al,Zn)+θ +η (Al,Zn)+(Al)+η 300
θ +η (Al)+θ +η
200
Mg 33.33 Zn 66.67 0.00 Al
MSIT®
20
40
60
80
Al
Al, at.%
Landolt-Börnstein New Series IV/11A3
Landolt-Börnstein New Series IV/11A3
U7 L+q+ζ e13 P3
E2
449
434
τ1+(Al)+η
L+τ1+φ
q+τ1+τ2
377 p6 q + τ 1 τ2
380 p5 L + τ1 q
τ1+η+ζ
P1
L + τ1 q + ζ U4 362 L + τ1 φ + q U3 τ1+q+ζ L+q+φ τ1+φ+q 356 q + τ1 τ2 + ζ U5 τ1+τ2+ζ U10 U q+τ2+ζ U7 U6 6 E3 E3 U10
γ+(Mg)+φ
U2
P2
L+τ1+ζ
L + τ1 + η ζ
τ1+β+γ
E1
L+(Al)+η 451 e3 L τ1 + γ
U1
530 p1 L + η τ1
A-B-C
L τ1 + β + γ
L + τ1 (Al) + η
Al-Mg-Zn
L + τ1 + γ φ L+γ+φ
L + γ (Mg) + φ
τ1+γ+φ
388
(Al)+τ1+β
L (Al) + τ1 + β
L+(Mg)+φ
362
368
447
450 e5 Lβ+γ
451 e2 L β + τ1
476
480 e1 L (Al) + τ1
Fig. 4a: Al-Mg-Zn. Reaction scheme, part 1
410 p3 β+γε
436 e7 l (Mg) + γ
449 e6 lβ+γ
450 e4 l (Al) + β
Al-Mg
360 e9 l (Zn) + θ
381 p4 l+ηθ
416 p2 l+ηζ
Mg-Zn
380 e8 l (Al) + (Zn)
Al-Zn
Al–Mg–Zn 205
MSIT ®
MSIT®
84
L + δ (Mg) + MgZn
331
q+φ+MgZn
U11
(Mg)+φ+MgZn
q + (Mg) φ + MgZn
τ1+τ2+ζ
E4 η+(Al,Zn)+(Al)
214
η+ζ+τ2
U10
τ1+η+ζ
D1
τ1 + ζ η + τ2
(Al)+(Zn)+θ
τ1+η+τ2
P1 U1 p4
(Zn)+(Al,Zn)+θ
(Al,Zn) (Al) + (Zn),θ
(Al,Zn)+(Al)+θ
U9
η+θ+(Al,Zn)
U8
E3
η + (Al,Zn) θ + (Al)
277
η+(Al)+θ
q+(Mg)+MgZn
L q + (Mg) + MgZn
q+φ+(Mg)
339
U7
L+φ+q
L+(Al,Zn)+θ
L θ + (Zn) + (Al,Zn)
L+(Mg)+MgZn
340
344
P3
U6
U3
A-B-C L+(Al)+η
L + η (Al) + θ
Al-Mg-Zn
L + φ (Mg) + q
q+ζ+MgZn
L+q+(Mg)
345
355
U5
L + q + ζ MgZn
L+q+ζ
τ1+τ2+ζ
L+(Mg)+φ
353
U4
L+q+MgZn
p3 U2
Fig. 4b: Al-Mg-Zn. Reaction scheme, part 2
250 e13 εβ+γ
Al-Mg e9 e8
325 e11 δ (Mg) + MgZn
341 e10 l MgZn + δ
341.1 p8 l + (Mg) δ
347 p7 l + ζ MgZn
Mg-Zn
277 e12 (Al,Zn) (Al)+(Zn)
Al-Zn
206 Al–Mg–Zn
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn
207
Al Fig. 5: Al-Mg-Zn. Liquidus surface, calculated using the data of [1998Lia2]
Data / Grid: at.% Axes: at.%
600 20
80
(Al) 55 0 e4 40
e6
E2
β
e1
e3
γ
60
60
50 0
e2 E1
U1 40
e7
τ1 450
400
550
η
p e10 ζ 40 MgZn 8 U8 p7 p2
60
) (Mg
600
P1
U7
5
q
U4
400 U6
80
Al Fig. 6: Al-Mg-Zn. Calculated solidus isotherms of the (Al) phase
20
e8 E3 (Zn)
P3
E4 20
Mg
0 50
φ Up3
p1
450
500
U2
P2
500
80
p4 θ e9
Zn
Data / Grid: at.% Axes: at.%
640° C
620 600 580
560 540 520 10
500
0 48
90
e1 U1
0 46 E2 Mg 15.00 Zn 0.00 Al 85.00 Landolt-Börnstein New Series IV/11A3
10
Mg 0.00 Zn 15.00 Al 85.00
MSIT ®
Al–Mg–Zn
208
Al
Data / Grid: at.%
Fig. 7: Al-Mg-Zn. Calculated solvus isotherms of the (Al) phase
Axes: at.%
340 0 36 8 3 0 400
τ1
420
η
440 460°C 10
90
e1 U1
β E2 10
Mg 15.00 Zn 0.00 Al 85.00 3
Mg 0.00 Zn 15.00 Al 85.00
U8
Fig. 8: Al-Mg-Zn. Calculated solidus and solvus isotherms of the (Mg) phase
solidus isotherms MgZn 40 0° C 2
E4
solvus isotherms univariant equilibria
U7 q
45
U2
30
Zn, at.%
0
35
0 0
50
0
0
25
1
20 0
15 0
55
0
60
0
0
Mg
0
2
e7 4
6
8
10
12
Al, at.%
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zn
209
Al Fig. 9: Al-Mg-Zn. Isothermal section at 335°C, calculated from the data of [1998Lia2]
Data / Grid: at.% Axes: at.%
(Al)
20
80
(Al)+η
β 40
ε
60
η+(Al)
+(Al,Zn)
60
γ
γ +φ +τ 1 80
φ +γ +φ g) φ (M g)+ φ (M )+q+ (Mg
(Mg)
(Al)+τ 1
γ +τ 1
(Al,Zn) 40
φ +τ 1
τ1 20
q τ2
θ+(Zn) +(Al,Zn)
η+θ +(Al,Zn)
(Mg)+MgZn+q (Mg)+MgZn
Mg
Landolt-Börnstein New Series IV/11A3
20
40
MgZn
60
ζ
η
80
θ
(Zn)
Zn
MSIT ®
210
Al–Mg–Zr
Aluminium – Magnesium – Zirconium Natalia R. Bochvar, updated by Marina V. Bulanova Literature Data [1958Ich, 1959Ich, 1968Bab] investigated the solubility of Zr and Al in liquid Mg by the chemical analysis of alloys quenched in water from 700 to 800°C. The solubility of zirconium decrease sharply on the addition of even minor quantities of aluminium. It was also found that the ZrAl3 compound is in equilibrium with liquid magnesium. [1960Sch] found, by very careful chemical analysis of the Mg rich liquid separated at 740°C from the solid equilibrium phases and by X-ray investigation of those, that ZrAl3 did not appear up to at least 0.0506 mass% Al and 0.0372 mass% Zr. Below that Al content the precipitates were identified as ZrAl2, Zr2Al3, Zr4Al3 and a more Zr-rich phase. [1964Cro] found by metallographic and X-ray inspection that ZrAl2 and Mg17Al12 were present in cast Mg rich alloys with 3 to 10 mass% Al and 0.01 to 1.6 mass% Zr. ZrAl2 was present even in the sample with 0.01 mass% Zr, indicating a very low joint solid solubility in (Mg). [1969Dri1, 1969Dri2] studied the aluminium corner of the system by metallographic and differential thermal analyses and reported the existence of a ternary phase -(ZryMgyAlx) in equilibrium with (Al) though the composition and crystal structure of this phase have not been determined. In following investigations of [1989Ker, 1992Fri], however, the existence of the ternary compound in the Al rich corner of the system was not confirmed. [1989Ker] studied the interaction of Al-Mg alloys with Zr by the diffusion-couple technique and constructed the isothermal section at 400°C in the whole concentration range; transition zones were analyzed by the electron microprobe technique. Using metallographic analysis, X-ray diffraction and electron microprobe techniques, [1992Fri] investigated the alloys annealed at 400°C with the constant Mg content of 6 mass%. According to [1989Ker, 1992Fri], only the phases which belong to the corresponding binaries exist in the system. [1977Asa1, 1977Asa2] investigated the magnesium corner of the system by measurements of electrical conductivity, thermo-emf and calculations based on general thermodynamic relations. They supposed the existence of two additional ternary compounds in the system, “Zr3Mg8Al9” and “ZrMg6Al3”. The vertical sections, however, have been constructed by the authors with violations of the phase rule. Their interpretation of the obtained results is considered unreliable and these works will not be further discussed. Binary Systems The following binary systems are accepted: Al-Mg [2003Luk], Al-Zr [Mas2, 1992Per] in the Zr-rich part and Mg-Zr [Mas2]. Solid Phases The solid phases are given in Table 1. No ternary phase is accepted. Invariant Equilibria It may be supposed that in the Al rich corner of the system, a ternary eutectic (Al)+, Mg2Al3+ZrAl3 exists at 450°C. This supposition is based on the absence of a ternary phase [1989Ker, 1992Fri] and on the absence of any arrests on the thermocurves except that at 450°C [1969Dri2]. Isothermal Sections Figure 1 shows the isothermal section at 400°C [1989Ker]. Miscellaneous Figure 2 shows the solubility limit at 740°C in the magnesium corner from 0.001 to 0.05 mass% Al [1960Sch]. The most Al poor precipitate could not be identified. With increasing Al content Zr4Al3, Zr2Al3 MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zr
211
and ZrAl2 were found in this careful investigation [1960Sch]. The solubility data of 0.4 to 0.04 mass% Zr and 0.05 to 0.015 mass% Al are much lower compared with earlier studies [1958Ich, 1959Ich]. Data from those earlier studies are given in Fig. 3 as isotherms of solubility at 700 and 800°C in the magnesium corner of the system according to [1958Ich, 1959Ich], where only ZrAl3 is mentioned as the precipitated phase. [1984Kai] reported the possibility of obtaining a zirconium-supersaturated solid solution of aluminium in an Al-2 Mg-0.16Zr (at.%) alloy by quenching at a high rate. [1991Lav] studied the alloy Mg-8.4Al-0.2Zr (mass%) (Mg-7.64Al-0.04Zr (at.%)) after it was spray atomized and the deposited. The research methods were SEM, EDAX, X-ray diffraction. The deposited residual contained (Mg) and eutectic (Mg)+ , Mg17Al12 . The deposit contained (Al) as well, that proves nonequilibrium process. References [1958Ich]
[1959Ich]
[1960Sch] [1964Cro] [1968Bab] [1969Dri1]
[1969Dri2]
[1977Asa1]
[1977Asa2]
[1984Kai]
[1989Ker]
[1991Lav]
[1992Fri]
[1992Per]
Landolt-Börnstein New Series IV/11A3
Ichikawa, R., “Solubility of Zr in Mg and its Alloys in the Liquid State. II. Alloys Containing Al, Fe, Mn and Si” (in Japanese), Nagoya Kogo Daigaku Gakuho, 10, 197-203 (1958) (Equi. Diagram, Experimental, #, 8) Ichikawa, R., “Al, Fe, Mn and Si as Impurities in Mg-Zr Alloys. I. Solubility of Zr in Molten Mg. II. Intermetallic Compounds Formed by Impurities and Zr” (in Japanese), Nippon Kinzoku Gakkaishi, 23, 192-194 (1959) (Equi. Diagram, Experimental, #, 5) Schneider, A., Stendel, J., “Precipitation of Intermetallic Phases from a Liquid Solvent Metal” (in German), Z. Anorg. Allg. Chem., 303, 227-246 (1960) (Experimental, 28) Crosby, R.L., Higley, L.W., “Intermetallic Compounds in Mg-Rich Mg-Al-Zr Alloys”, U.S. Bur. Mines, Rep. Invest., 1-23 (1964) (Equi. Diagram, Experimental, 7) Babkin, V.M., “Solubility of Zr in Molten Mg and ML5 Alloy” (in Russian), Metalloved. Term. Obrab. Met., 3, 61-64 (1968) (Equi. Diagram, Experimental, #, 4) Drits, M.E., Kadaner, E.S., Kuzmina, V.I., “Phase Diagram of the Al-Mg-Zr System in Al-Rich Region”, Russ. Metall., translated from Izv. Akad. Nauk SSSR, Met., 5, 170-173 (1969) (Equi. Diagram, Experimental, #, 6) Drits, M.E., Kadaner, E.S., Kuzmina, V.I., “Interaction of Components in Ternary Systems” (in Russian), in a Collection of Papers “Aluminium Alloys”, Metallurgiya, Moscow, 6, 146-149 (1969) (Equi. Diagram, Experimental, #, 2) Asanovich, V.Ya., Sryvalin, I.T., Korpachev, V.G., “An Electrometric Study of the Aluminium-Magnesium Zirconium System” (in Russian), Nauchn. Tr. Kuban. Univ., 3, 42-48 (1977) (Equi. Diagram, Experimental, #, 6) Asanovich, V.Ya., “Phase Diagram of the Aluminium- Magnesium- Zirconium System”, Russ. Metall., (4), 169-171 (1977), translated from Izv. Akad. Nauk SSSR, Met., (4), 208-210 (1977) (Equi. Diagram, Experimental, #, 5) Kaibyshev, O.A., Valiev, R.Z., Tsenev, N.K., “Influence of the Grain Boundary State on the Superplastic Flow”, Sov. Phys. -Dokl., 29, 752-754 (1984), translated from Dokl. Akad. Nauk SSSR, 278, 93-97 (1984) (Experimental, 12) Kerimov, K.M., Dunaev, S.F., Slusarenko, E.M., “Investigations on Phase Equilibria in Aluminium - Magnesium- (Titanium, Zirconium, Hafnium) Systems” (in Russian), Vestn. Mosk. Univ. Ser. 2: Khim., 30, 156-161 (1989) (Experimental, Equi. Diagram, 8) Lavernia, E.J., Baram, J., Gutierrez, E., “Precipitation and Excess Solid Solubility in Mg-Al-Zr and Mg-Zn-Zr Processed by Spray Atomization and Deposition”, Mat. Sci. Eng. A, A132, 119-133 (1991) (Experimental, Crys. Structure) Fridman, A.S., Dobatkina, T.V., Muratova, E.V., “Section of Isothermic Tetrahedron of the Al-Rich Portion of the Al-Mg-Sc-Zr System at 500°C” (in Russian), Izv. Akad. Nauk SSSR, Met., (1), 234-236 (1992) (Equi. Diagram, Experimental, 4) Peruzzi, A., “Reinvestigation of the Zr-Rich End of the Zr-Al Equilibrium Phase Diagram”, J. Nucl. Mater., 186, 89-99 (1992) (Equi. Diagram, Experimental, 17)
MSIT ®
Al–Mg–Zr
212 [2003Luk]
Lukas, H.-L., Lebrun, N., “Al-Mg (Aluminium-Magnesium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 49)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Al) < 660.45
cF4 Fm3m Cu
a = 404.96
pure Al at 25°C [Mas2]
(Mg) < 650
hP2 P63/mmc Mg
a = 320.94 c = 521.05
pure Mg at 25°C [Mas2]
(Zr) 1855 - 863
cI2 Im3m W
a = 360.99
[Mas2]
(Zr) < 863
hP2 P63/mmc Mg
a = 323.12 c = 514.77
pure Zr at 25°C [Mas2]
, Mg 2Al3 452
cF1168 Fd3m Mg2Al3
a = 2816 to 2824
60-62 at.% Al [2003Luk] 1168 atoms on 1704 sites per unit cell [2003Luk]
, Mg17Al12 < 458
cI58 I43m Mn
a = 1054.38
at 41.4 at.% Al [V-C2] 39.5 to 51.5 at.% Al [2003Luk]
J, Mg23Al30 410 - 250
hR159 R3 Mn44Si9
a = 1282.54 c = 2174.78
54.5 to 56.5 at.% Al [2003Luk]
ZrAl3 < 1580
tI16 I4/mmm ZrAl3
a = 401.4 c = 1732.0
[V-C, Mas2]
ZrAl2 < 1645
hP12 P63/mmc Zn2Mg
a = 528.24 c = 874.82
[V-C, Mas2]
Zr2Al3 < 1595
oF40 Fdd2 Zr2Al3
a = 960.1 b = 1390.6 c = 557.4
[V-C, Mas2]
ZrAl < 1275
oC8 Cmcm CrB
a = 335.3 b = 1086.6 c = 426.6
[V-C, Mas2]
Zr4Al3 < 1030
hP7 P6 Zr4Al3
a = 543.3 c = 539.0
[V-C, Mas2]
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mg–Zr
213
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
Zr3Al2 < 1480
tP20 P42/mnm Zr3Al2
a = 763.0 c = 699.8
[V-C, Mas2]
Zr5Al3 1395 to ~1000
tI32 I4/mcm W5Si3
a = 1104.9 c = 539.6
[V-C, Mas2]
Zr2Al < 1215
hP6 P63/mmc Ni2In
a = 489.39 c = 592.83
[V-C, 1992Per]
Zr3Al < 1019
cP4 Pm3m AuCu3
a = 437.2
[V-C, 1992Per]
Al
Data / Grid: at.% Axes: at.%
(Al)
Fig. 1: Al-Mg-Zr. Isothermal section at 400°C 20
80
ZrAl3 ZrAl2 Zr2Al3
40
β
60
ZrAl Zr4Al3 Zr3Al2
γ 60
40
Zr2Al Zr3Al 80
20
(Mg)
αZr
Mg
Landolt-Börnstein New Series IV/11A3
20
40
60
80
Zr
MSIT ®
Al–Mg–Zr
214
10-0.5
Fig. 2: Al-Mg-Zr. Liquid solubility in the Mg corner at 740°C (log scale) [1960Sch]
L+?
Zr, at.%
10-1.0
L+Zr4Al3
10-1.5
L+Zr2Al3 L
L+ZrAl2
10-2.0
10-2.5 10-2.5
10-3.0
10-2.0
10-1.0
10-1.5
Al, at.%
Mg 98.00 0.00 Zr 2.00 Al
Fig. 3: Al-Mg-Zr. Liquidus solubility limits in the Mg corner at 700 and 800°C suggested by [1958Ich, 1959Ich]
Axes: at.%
700°C
L
Mg
MSIT®
Data / Grid: at.%
800°C L+ZrAl3 Mg 98.00 2.00 Zr 0.00 Al Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
215
Aluminium – Manganese – Palladium Oleksandr Dovbenko, Tamara Velikanova, Sergiy Balanetskyy Literature Data The alloys of the Al-Mn-Pd system have been investigated in many works, however the equilibrium diagram has not been determined for the whole composition range yet. In accordance with the assessment by [1993Ran], based on the works of [1968Web, 1981Sol1, 1981Sol2], the MnPd2Al Heusler phase exists in the system. The MnPd2Al alloy has bcc W type structure at high temperature, CsCl type below 1010°C and the structure of MnCu2Al type below 710°C [1981Sol2]. Many reports with data on ternary quasicrystalline and crystalline phases in Al-rich part of the phase diagram (up to 50 at.% Al) have been published later. [1968Web] determined the MnPd2Al structure as CsCl type in samples prepared by arc melting from pure metals and annealed in vacuum at 800°C for 24 h, followed either by quenching or slow cooling over 36 h. [1981Sol2] studied the solid state transition of the MnPd2Al alloy by X-ray diffraction. Components (99.999% purity) were mixed and then melted in an induction furnace. The ingots were powdered and pressed into tablets of 10 mm diameter. These were sealed in silica ampoules under vacuum of 0.1 to 0.01 Pa, homogenized at 800°C for 60 h, and quenched. The temperatures of phase transformations were determined by diffractometry and high-temperature photography in the range of 30 to 1230°C with a heating rate of 600K#h-1. The order-disorder transition was also treated theoretically by [1981Sol1]. The icosahedral quasicrystalline phase in the composition range 5-25 at.% Pd and 10-20 at.% Mn was first found in the rapidly solidified samples by [1990Tsa1] and in samples prepared by arc melting and then annealed between 850 and 900°C (10h) in vacuum by [1990Tsa2]. The XRD, TEM [1990Tsa1, 1990Tsa2] and DTA, DSC, SEM and optical microscopy (OM) [1990Tsa2] methods were used for investigation. The structure of alloys and phase equilibria in the Al-rich part of the system were investigated by [1991Yok, 1991Tsa, 1992Was, 1992Yok, 1993Aud, 1995Goe1, 1995Goe2, 1998Sim, 1999Gru, 1999Yok, 2000Gru, 2000Kle1, 2001Goe, 2002Ant, 2002Yur]. [1991Yok, 1992Yok] studied the ternary alloys in the composition range (14-26)Pd-(3-15)Mn (at.%), prepared from pure metals (Pd 99.996 mass% and Al, Mn 99.999 mass%) by arc melting under argon atmosphere. Ingots were annealed for 12 h at 850 and 870°C in vacuum and quenched in water. In addition the pre-alloyed ingots, after heating in an argon atmosphere for 1 h at 1130°C, were examined by DTA at a cooling rate of 0.033 K#s-1. EDX, XRD, SEM, SSM and OM analysis were used, too. The temperature-composition section at Al80-xPd20Mnx, (with x = 0 to 20) and equilibria of the icosahedral phase with liquid at 850 and 870°C were reported. The area of formation (Pd+Mn=20-30 at.%) and the composition dependence of the lattice parameter (a) of the supercooled icosahedral phase have been determined by XRD [1999Yok]. The samples were prepared from master alloys by the zone melting process in ultrahigh vacuum using metals of the purity as in the works of [1991Yok, 1992Yok]. Czochralski and Bridgman techniques with a flux were used to produce icosahedral single phase samples. The composition range in which the icosahedral phase forms was also studied by [1991Tsa, 1992Was, 2000Gru] and a partial liquidus projection in the Al corner was drafted by [1993Aud]. The authors investigated samples prepared from pure metals by induction melting under argon atmosphere. The single quasicrystalline samples were prepared by Bridgman and Czochralski technique. SEM-XEDS, DTA, EPMA, TEM, HREM, X-ray wavelength dispersive spectroscopy (XWDS) were used to analyze the structure of the phases and for investigating the phase composition in the area of quasicrystalline phases. The liquidus projection and the isothermal sections at 894, 875°C were constructed by [1995Goe1, 1995Goe2, 2001Goe]. The liquidus surface results from DTA at usual cooling and heating rates, the isothermal sections are based on samples annealed at 894 and 875°C for 4 d, at 840°C for 6 d and 600°C for 18 d, in a range of 60-100 at.% Al. Similarly the vertical sections were constructed from thermal analysis data. Approximately 70 samples of different compositions were investigated prepared from high purity metals (Mn 99.985 %, Al and Pd better than 99.998 %) by induction melting. The components were placed in a corundum crucible and this in turn was closed in a silica ampule, evacuated and filled with 650 hPa
Landolt-Börnstein New Series IV/11A3
MSIT ®
216
Al–Mn–Pd
argon. Silica ampoules with extremely thin walls were used for DTA with heating and cooling rates of 2-10°C/min. Several alloys have been investigated by measuring magnetic susceptibility versus temperature (MTA) and some samples were studied by EDX analysis. Based on a limited number of samples annealed for 45 min [2000Kle1] determined liquid-solid equilibria at 685, 730, 840, 920, 945 and 952°C and proposed a projection of the liquidus surface in the region 50-100 at.% Al. The alloys were prepared from 99.999 mass% Al, 99.9 mass% Pd and Mn by induction melting and some of the ingots were re-melted in alumina crucibles under argon atmosphere in order to produce large single grains by the Bridgman technique. The crystallization sequences were deduced from SEM-XEDS, XWDS and TEM methods and from DTA samples of about 0.15 g, applying cooling and heating rates of 5°C/min. In the Al-rich part of the system isothermal sections at 880, 870, 850 and 790°C were reported by [1999Gru, 2002Yur] in the vicinity of the quasicrystalline phases. The 33 ingots were prepared by induction (levitation) melting in a cold crucible under argon atmosphere. Part of the ingots were annealed at 880, 870 and 850°C for 65-70 h, at 790°C for 590 h and at 710°C for 1450 h and quenched in water. The samples were examined by OM, SEM-EDX and XRD methods. The bulk composition of some single-phase samples was measured by inductively coupled plasma optical emission spectroscopy, ICP-OES. Single-phase samples whose compositions were measured by ICP-OES were used as standards for the correction of the EDX data. Selected samples were studied by DTA at heating and cooling rates of 20°C/min and by TEM. The TEM examinations were carried out on the powdered materials, which were spread on copper grids with carbon film. The composition of the liquid phase coexisting with icosahedral and compositions of adjacent ternary crystalline phases were determined by [1998Sim]. Five ternary alloys (Mn7.2Pd20.7Al72.1, Mn5Pd18Al77, Mn3.6Pd16.6Al79.8, Mn3.5Pd20Al76.5 and Mn1.7Pd19.2Al79.1) and one binary alloy (Pd19Al81) were investigated, prepared from Al 99.999 mass%, Pd and Mn 99.9 mass% by induction melting in a cold crucible under argon atmosphere. DTA, chemical composition analysis, structure analysis by X-ray or electron diffraction and neutron scattering were applied to the as-cast samples. All investigated ternary solid phases were found to melt incongruently. Thin films were deposited by simultaneous evaporation of the metals from separate sources on carbon substrates or on glass plates at temperatures up to 500°C and the structures and compositions of ternary films were investigated by TEM, electron diffraction, and EDX [2002Ant]. Icosahedral order was observed for aluminium contents above 75 at.% and a phase diagram for thin films at a deposition temperature of 475°C was constructed. There is a number of experimental and theoretical works devoted to the crystallographic investigation of the quasicrystalline and periodic crystalline ternary phases in the Al-rich part of this system [1991Bee, 1991Bou, 1991Don, 1992Bou, 1992Was, 1993Aud, 1993Bou, 1993Dau, 1993Hir1, 1993Ste, 1993Sun, 1993Tsa, 1993Was, 1994Aud, 1994Bee, 1994Li1, 1994Li2, 1995Bee1, 1995Bee2, 1995Boi, 1995Ish, 1996Bou, 1996Yam, 1997Ama, 1997Hae, 1997Kle, 1997Kra, 1997Mat, 1997Son, 1997Zur, 1998Ber, 1998Boi, 1998Mat, 1998Wan, 1999Aud, 1999Cap, 1999Fis, 2000Bee1, 2000Bee2, 2000Dun, 2000Fra, 2000Fre, 2000Gwo, 2000Hir, 2000Jan, 2000Jac, 2000Kaj, 2000Kle3, 2000Let, 2000Nic, 2000Qua, 2000Sch1, 2000Shr, 2000Sta, 2000Ste1, 2000Ste2, 2000Uch, 2000Yam, 2001Nau, 2002Hir, 2002Lei, 2002Shr, 2002Yam, 2002Yan, 2002Zha1, 2002Zha2]. [1992Bou, 1993Bou] examined single grain samples of the icosahedral phase by XRD and neutron diffraction. [1993Sun] investigated Al70Pd20Mn10 samples rapidly solidified and annealed at 800°C by TEM, XRD, HREM, EDXA. TEM and XRD methods were used by [1991Bee, 1992Was, 1993Was, 1995Ish, 1998Boi, 1999Aud, 1994Aud]. The modifications of the icosahedral phase have been studied by [2000Hir, 1998Boi, 1995Ish, 2000Let, 1999Aud, 2002Yam]. [1995Ish] investigated samples of composition MnxPd29-xAl71, (6.5 < x < 9.5), prepared from pure elements (Al 99.999 mass%, Pd 99.95 mass%, Mn 99.99 mass%) in a plasma-jet furnace. The samples, put into a graphite crucible and sealed in silica tubes, were annealed for 50 h at 803 4 and 48-400 h 602 2°C, quenched into water and subsequently into liquid nitrogen. An alloy of Mn8Pd21 Al71 composition annealed at different temperatures was examined by powder X-ray diffraction using CuK radiation at room temperature by [2000Hir], whereas [1998Boi] performed in-situ heating experiments using synchrotron light source. [2000Let] examined ingots of composition Mn8.8Pd21.4 Al69.8 by X-ray and in-situ neutron MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
217
diffractions. [2002Yam] investigated a sample at Mn8Pd21Al71 by XRD with an IP-Weissenberg camera and carried out a structure refinement of the icosahedral quasicrystal. A two-axis diffractometer was used by [2000Shr] to measure the diffuse background for two different single grained samples: icosahedral (Mn8.2Pd21.4Al70.4) and !´ (Mn4Pd22Al74) phases. The T and R orthorhombic ternary phases in this system were found by [1993Aud, 1994Aud, 1997Kle, 1997Mat, 1998Ber]. [1997Kle] obtained single crystal T-Al3(Pd,Mn) phase by Bridgman technique from an ingot of nominal composition Mn19 Pd7.1Al73.9, prepared by induction melting in a cold crucible from the pure elements (99.999 mass% Al, 99.9 mass% Pd, 99.9 mass% Mn). [1997Mat] performed XRD analysis of the single crystal of nominal composition Mn25Pd5 Al70 prepared by arc melting pure metals under Ar atmosphere. [1998Ber] examined the sample of the Mn3Pd9.3 Al87.7 composition by XWDS, TEM and HREM and observed the structural defects near the R-! interfaces and within both phases. Single crystals of the ternary !´ phase were obtained by [1996Bou] using the Bridgman technique from ingot with starting nominal composition of Mn3.5Pd19 Al77.5 . [2002Shr] observed this phase in samples of the composition around 4.5Mn-22.9Pd-72.6Al (at.%), (which corresponds to the nominal composition of the !´ phase, prepared by the Bridgman crystal growth method. Powder and single crystal XRD experiments were carried out. The electronic structure and electronic densities of the decagonal phase Mn17Pd13Al70 and related crystalline phases have been calculated by [1997Kra] and a structure-induced pseudogap in the Al band was shown to exist in decagonal as well as in related ternary periodic phases of similar composition. The stabilization by the Hume-Rothery-like band-gap was found to decrease in the sequence: crystalline (Al6Mn) - icosahedral (Mn8Pd22Al70) - decagonal (Mn17Pd13 Al70 ). The results of the photoemission spectroscopy on the electronic structure of quasicrystals have been reviewed by [2000Sta] and the existence of the theoretically predicted pseudogap at the Fermi level being confirmed. Binary Systems The Al-Pd and Al-Mn systems are accepted from [2003Bal] and [1997Oka, 2003Pis], respectively. Data concerning the Mn-Pd system are from [Mas2]. Solid Phases Crystallographic data on the known unary and binary phases as well as recently reported ternary ones are listed in Table 1. A peculiar feature of the system is the formation the two stable quasicrystalline phases: decagonal -2, usually labeled as “D”, and icosahedral -3, usually labeled as “I”. Another peculiarity is the close crystallographic relationship between the solid solutions based on binary phases such as Mn4Al11 (h), and J, MnAl4 and other periodic phases, and the quasicrystalline phases. The mutual solid solubility of the isostructural , MnAl and ( Mn) on the one hand, with the , PdAl (h) and , MnPd (h) phases on the other are reported by [1995Goe1, 1995Goe2, 2000Kle1]. However, the authors did not give details about the phase relationships in the range where disordered bcc , MnAl of the W type and the ordered CsCl type , PdAl do coexist. At 1010°C the cubic W type phase transforms into the cubic CsCl type phase [1981Sol2] which at 710°C transforms into the -1 MnPd2Al Heusler phase. The composition of the decagonal -2 quasicrystalline phase is very close to Mn18.1 Pd12.1Al69.8, its melting point is at 896°C according to [1995Goe1, 1995Goe2]. A crystalline “pseudo-decagonal” phase of the same composition identified as DH, a high temperature modification of D, with a B-centered orthorhombic cell is reported by [1995Bee1, 1995Bee2, 1995Goe2] above 864°C. The homogeneity range of the icosahedral phase -3 is Mn8-10.2Pd20.3-23.2Al68-69.5 according to [2000Kle1] and in temperature range from 880 to 710°C it is Mn6-10Pd19.2-24.5Al69.5-70.8, according to [1999Gru, 2000Gru, 2002Yur]. Three modifications of the icosahedral phase were found by [1995Ish, 1998Boi, 2000Hir, 2000Let] at different temperatures and in close compositional vicinity, i.e. Mn9.2Pd22.0Al68.8, Mn8.8Pd21.4Al69.8, and Mn8.7Pd22.0Al69.3, labeled as F, F2 and F2M, respectively. The high temperature F phase has a 6 dimensional reciprocal primitive cubic cell with strong chemical order. The F2 phase is considered as a superstructure of the F-phase and could be described as a P type 6D hypercubic lattice with Landolt-Börnstein New Series IV/11A3
MSIT ®
218
Al–Mn–Pd
parameter aP = -aF (aP = 2088.1 pm), or a 6 dimensional diamond-type structure with aF2 = 2aF (aF2 = 2581.0 pm) [1995Ish, 2000Hir]. According to [2000Let] the F2 phase is not stable and corresponds to a transient state in the process of the transformation of the icosahedral F phase to the F2M phase. The F2M phase has a domain structure with a cubic but non-periodic symmetry [1998Boi] and forms from high temperature F phase at 715°C [2000Let]. According to [1993Aud, 1994Aud, 1996Sun, 1997Kle, 1998Ber, 1996Bou, 2000Shr, 2000Kle1, 2002Shr, 2002Tex] four additional orthorhombic ternary phases designated as R, T, !´ and ! exist. The R, T and !´ phases form from the liquid by peritectic reactions. According to [2000Kle1] the composition range of the R and T phases on solidus overlap (3.5-6.6 at.% Pd, 16-25.5 at.% Mn), but separated in the liquidus surfaces of the system. The T phase is reported to be isostructural to the MnAl3 phase where Pd substitutes Al only in one position and this position shifts along the b axis. The R phase is reported to be isostructural to Mn11Ni4Al60 [1994Aud, 2000Kle1, 1997Kle]. According to [1997Kle] the T phase might transform into an R phase with aR = -aT, b R = bT, cR = --1cT. From HREM data both the R and T phases are pointed to exhibit different tilings in their ac plane but built from the same flattened hexagon. The orthorhombic !' phase has a composition of Mn5.0-4.6Pd22.1-22.4Al73.3 at 730°C according to [2000Kle1], and exists in several modifications [1996Kle, 2002Shr]. According to [1996Bou] the structure of this phase is very close to J6 the “PdAl3” phase in which the Mn atoms substitute only two Pd sites of high symmetry among 12 Pd sites. The additional two modifications, called !´_3 and !´_5 by [2002Shr], have a periodicity defined as a!´_n = a!´, b!´_n = b!´ and c!´_n = c!´´ (n+-), where n equals 3 or 5, respectively and - is the golden mean. The structure of the !´_5 phase is different from the structure of the !´ phase, but not very much. Also, a modification !´_2 with c!´_2 = c!´ (2+-) may exist in this system as well [2002Shr]. According to [2000Kle1] the T, R, ! and !´ phases are ternary compounds and have to be distinguished from the binary phases. One of the reasons for this conclusion was that in the samples Mn1.7Pd19.2Al79.1 (obtained by Bridgman) and Mn11.4Pd1.3Al87.3 (not annealed sample after DTA experiment) interfaces between the ternary and binary phases have been observed, i.e. between “PdAl3” and !´ phases, or Mn4Al11 and T phases, respectively. These, however, also could result from segregation and nonequilibrium conditions of the samples. The reported composition ranges of the R, T and !´ phases are close to the binary MnAl4, Mn4Al11(h), J6 and J28 phases and their crystal structures are very similar; R phase is isostructural to metastable %-MnAl4 phase, T phase to Mn4Al11 (h) or MnAl3, and !´ is isostructural to the J6 phase. According to [1995Goe1, 1995Goe2, 1999Gru, 2002Yur] the H solid solution, based on the Mn4Al11 (h) binary phase, and “J” solid solution exist in the composition range reported for the T and !´ phases, respectively, where “J” is considered to be the ternary extension of the J6 and J28 binary phases. According to [2002Yur] both J6 and J28 binary phases were observed in the ternary overall “J” field and, in addition, J22 and J34 ternary phases were observed by TEM. However the SEM/EDX analysis did not reveal any compositional inhomogeneities typical for such a multiphase sample. The J6, J22 , J28, J34 phases can be correlated with !´, !´_2, !´_3 and !´_4 respectively, according to designation by [2002Shr]. The coexistence of the J6, J22, J28 , J34 phases in the “J” continuous range of the ternary Al-Mn-Pd system is unclear. Their stability at temperatures under investigation is consistent with long-term annealing and the “J-phase” is stable in the ternary system up to the melting temperature of 845°C according to DTA data [2002Yur]. Thus the question wether Jl (!´), R and T phases are solid solutions based on binary phases or individual ternary phases is open and needs additional investigations. All data for the R, T and !´ phases are given in Table 1 together with data for %, MnAl4, Mn4Al11(h) and J22, J34 phases, respectively. The stability range of the cubic and orthorhombic phases reported by [1992Was, 1993Aud, 1994Aud, 1994Li2, 1998Ber, 1999Hip, 2000Kle1], based on XRD and TEM data, is not clear. An orthorhombic phase, labeled as -2-R, with space group Amm2 and lattice parameters a = 1243, b = 2030, c = 6250 pm, or a = bR, b = -2cR, c = -2aR, according to [2000Kle1] is reported by [1993Aud, 1994Aud]. The periodic cubic approximant (2/1) with lattice parameter a = 2030 pm was observed by [1992Was] in the Mn4Pd26Al70 sample after heat treatment at 750°C and by [1994Li2] in the Mn10Pd20Al70 sample annealed at 800°C for 3 d. The authors of [1994Li2] also found in this sample the cubic phase (1/1) with lattice a parameter a = 1240 pm. Four orthorhombic phases: the (1/1,1/1) with lattice parameters a = 1260, b= 1240, c = 1480 pm; (2/1,5/3) with a = 1920, b = 1240 c = 6140 pm; (5/3,3/2) with a = 5050, b = 1240, c = 3780 pm and MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
219
the (8/5,5/3) with a = 8400, b = 1240, c = 6200 pm were found by [1994Li2] in a Mn15Pd15Al70 sample which was annealed for 3 d at 800°C and quenched in water. Invariant Equilibria A partial reaction scheme is presented in Fig. 1a for the Al-rich part, according to [1995Goe1] with some corrections made according to the accepted binary diagrams and data of [1999Gru, 2002Yur] on additional invariant ternary transition equilibrium -3+-2+ (U 4) at 860°C. The thermal DTA effects obtained by [1995Goe1, 1995Goe2] at 864°C are shown as dotted lines in Figs. 9-11, 14. These lines are very close to the above-mentioned temperature reported by [1999Gru, 2002Yur], and interpreted by [1995Goe2] as DHD transition temperature. It is not clear wether these temperatures really correspond to different processes. The temperature of the invariant equilibrium L+-3+ (U 3) that is 867°C according to [1995Goe1] and 870 < T < 880°C according to [1999Gru, 2002Yur], was accepted from the last work (~875°C). The temperature of the reaction P3 (L+ +-3J) is accepted from [2002Yur]. According to [1995Goe1, 1995Goe2] the decagonal quasicrystalline phase can form in two ways. According to the equilibrium diagram, the decagonal phase crystallizes from the liquid by peritectic reaction P1: L+Mn4Al11 +-2 at 896°C (Figs. 1a, 2a) before the icosahedral phase. In a metastable state, for example after rapid solidification, -2 phase is formed in the solid state at ~770°C [1995Goe1] by the reaction H++-3-2 (Figs. 1b, 2b). The icosahedral phase is formed from the liquid by peritectic reaction at approximately 893°C [1995Goe1, 1995Goe2, 1999Gru, 2002Yur]. As mentioned in the chapter Solid Phases, according to [2000Kle1] the T, !´ and R phases are ternary compounds and form from the liquid by peritectic reactions. The authors [2000Kle1] however noted nonequilibrium state of the samples therefore this data were not used in the discussion of the reaction scheme and this question needs additional investigations. The ternary eutectic reaction L(Al)+!´+R at a temperature 6175°C is reported by [1998Sim]. According to [1995Goe1] the equilibrium U9 and binary eutectic e4 occur at the same temperature. Thus, the nature of this transition remains undefined. Liquidus and Solidus Surfaces The projection of the liquidus surface in the Al-rich part of the system in Fig. 2a is given after the data by [1995Goe1, 1995Goe2, 2001Goe] and exhibits surfaces of the binary phases and only two stable ternary compounds (decagonal and icosahedral phases) firmly determined. The surface with the stable liquid phases and the primary solidification of -2 and the corresponding nonvariant equilibria were extrapolated in [1995Goe1] from the long term heating curves of the DTA and MTA plots and from the heat treatment experiments on alloys of 18 at.% Pd and 10-18 at.% Mn. A continuous ternary (MnPdAl) liquidus surface, extending from Al-Mn (55.1 to 71.7 at.% Al) to Al-Pd (38.3 to 70.5 at.% Al) binary boundary liquidus, is proposed by [1995Goe1, 1995Goe2, 2000Kle1] notwithstanding that the different crystal structures (W and CsCl types) of the binary compounds form corresponding mutual ternary solid solutions. It can not be excluded that the peculiar shape of the (MnPdAl) liquidus reflects this fact. But mutual transformations of cubic phases - W and CsCl type are not investigated. The doted lines in Fig. 2 correspond to the expected phase transition of cubic (MnPdAl) solid solution and to the monovariant line l+J28+J6 that starts from peritectic reaction p6 in the Al-Pd system. The metastable liquidus surface is projected in Fig. 2b for the Al-rich part of the phase diagram after [1995Goe1, 1995Goe2]. A different partial metastable liquidus surface projection was proposed by [2000Kle1]. The extensions of the liquidus phase fields was determined from the solidification sequences of DTA samples and the liquid-solid equilibria in samples annealed for 45 min at 952, 945, 920, 840, 730, 685°C and subsequently quenched. The three additional ternary compounds (T, R, !´) on the solidus supposed by [2000Kle1] and the separate liquidus fields corresponding to these ternary phases are accepted to exist in addition to the one quasicrystalline phase (icosahedral): for T and R phases in the H phase liquidus surface (after [1995Goe1]) and for the !´ phase in “J” region. The temperature limits of the monovariant reactions obtained by [2000Kle1] are in good agreement with the data of [1995Goe1, 1995Goe2] if one supposes that the phase fields H and “J” [1995Goe1, 1995Goe2] correspond to the phase fields R, T and !´ [2000Kle1], respectively. Landolt-Börnstein New Series IV/11A3
MSIT ®
220
Al–Mn–Pd
The partial solidus surface projection corresponding to the liquidus presented in Fig. 2 is given in Fig. 3 according to data [1995Goe1]. Isothermal Sections The partial isothermal section (in the Al-rich part of the system) at 894°C is given in Fig. 4 after work [1995Goe1]. The sections at 875, 840, 710 and 600°C presented in Figs. 5-8 result from evaluating all the data presented by [1995Goe1, 1999Gru, 2002Yur]. Mutual transformations of cubic phases - W and CsCl type, which are to take place in the ternary system, are not investigated. In isothermal sections above 840°C which is the temperature of eutectoid decomposition of the bcc phase in the Al-Mn system, the phase fields of (W) and (CsCl type) are separated in the same way as in the solidus projection in Fig. 3. According to [1995Goe1] at 894°C (Fig. 4) the decagonal -2 phase coexists in equilibrium with liquid and solid solutions based on the binary cubic phase and H phase, which is a ternary extension of the binary Mn4Al11 (h) phase. According to [1995Goe1, 1999Gru, 2002Yur] the last phase is stabilized by Pd additions in the ternary alloys down to temperatures that are lower then its decomposition temperature in the binary system (stable down to 600°C [1995Goe1]). Two ternary compounds -2 (decagonal) and -3 (icosahedral) in the isothermal sections at 875, 840, 710 and 600°C are shown. At 840°C (Fig. 6) the overall “J” phase field is a ternary extension (solid solution) of the binary J phases, which includes J6, J22, J34 and J28 phases stabilized by Mn to higher temperature than in the binary Al-Pd system and at 600°C this phase field joins the edge boundary system. There is the nonvariant four-phase equilibrium L+-3+ (transition type) at 867°C according to [1995Goe1], but according to more precise data of [1999Gru, 2002Yur] it should be in the temperature range 880-870°C (Fig. 5) (accepted here ~875°C). The authors [1995Goe1] reported the equilibrium between -2 phase and solid solution at 875, 840 and 600°C. However, according to the detailed investigation of [1999Gru, 2002Yur] at 850 and 710°C the -2 phase is in equilibrium with the binary phase as a result of the transition type reaction -3+-2+ which was determined at 860°C (from the DTA data for Al-20Pd-12Mn alloy annealed at 850°C). The isothermal sections at 840, 710 and 600°C are shown in Figs. 6-8 taking into account the above mentioned data of [1999Gru, 2002Yur]. The phase equilibria at 790°C are similar to those for 710°C, [1999Gru, 2002Yur]. At 875°C a broad band of liquidus extends from the Al-Pd binary system to 10 at.% Mn and the -3 phase was found to be homogeneous from 7.5 to 10 at.% Mn, [1995Goe1]. The homogeneity range of the -3 phase at 880°C extends from 70.2 to 71.2 at.% Al and from ~8.2 to 10.4 at.% Mn [2002Yur]. The homogeneity range of the -3 phase at 870°C spans from 69.6 to 71.6 at.% Al and about 8 to 10.5 at.% Mn [2000Gru, 2002Yur]. The composition of the H phase in equilibrium with -2 and -3 is 71-73 at.% Al and 6-7 at.% Mn. The solubility of the Mn in the solid solution based on the phase reaches up to 2.0 at.% at 850°C and 1.6 at.% at 710°C. The homogeneity range of the icosahedral phase extends from 70.0 to 71.6 at.% Al and 6.7 to 10 at.% Mn at 850°C and from 70.0 to 71.0 at.% Al. and 5.6 to 8.5 at.% Mn. The homogeneity range of the -2 phase at 710°C spans from 69.4 to 70.2 at.% Al and from 14.8 to 17.3 at.% Mn. The partial isothermal sections at 850 and 870°C reported by [1992Yok] expose the liquid-solid equilibria of the -3 phase. These data are in satisfactory agreement with the above given data. Temperature – Composition Sections The temperature - composition sections across 10 at.% Pd, 20 at.% Pd, 70 at.% Al and 6 at.% Mn are given in Figs. 9-12 according to [1995Goe1]. The sections Mn32.8Al67.2 - Pd27Al73, Mn31.2 Al68.8 - Pd29Al71 and Pd3Al97 - Mn13Pd30Al57 (Figs. 13-15) are given according to [1995Goe2] taking into account the above mentioned corrections from the isothermal sections. The latter section was obviously mislabeled as Pd3Al97 - Mn22.8Pd30Al47.2 by [1995Goe2].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
221
Thermodynamics The heat capacity of a Mn10Pd20 Al70 alloy was measured by [1997Ina] in the 1-350 K temperature range with an adiabatic calorimeter (above 5 K) and an isoperibolic one installed in a helium-3 cryostat (below 7 K). The samples were obtained by arc melting pure metals and from their melts single grains of the icosahedral phase were grown at 1200°C at a growth rate of 1mm/h from melt (Bridgman method). The Cp/T vs T2 plot from these data is presented in Fig. 16. The specific heat of the icosahedral phase at constant pressure (Cp) and at constant volume (Cv) was investigated by [2000Eda]. Below 427°C (Fig. 17a) Cp shows the typical behavior of ordinary solids, it appears to approach 3 kB asymptotically with increasing temperature, obeying the Dulong-Petit’s law. Above 427°C, however, it increases dramatically and reaches approximately 5 kB at 807°C. The temperature dependence of Cv has been calculated in the same work (Fig. 17b). The sample Mn9Pd20Al71 was prepared from elemental constituents by arc melting under argon atmosphere and annealed at 750°C for 72 h. XRD, DSC measurements with heating rate 10°C/min were applied in the temperature range of 277-807°C and XRD measurements at higher temperatures (7-527°C). For the same phase the specific heat Cp(T) was measured by [1998Wae] in the 1.6-280 K temperature range using two different techniques (Fig. 18). A standard relaxation technique was adopted at 1.6-30 K, and adiabatic continuous-heating calorimeter was used in the 15-300 K range. [1998Wae] grew single grain with final composition Mn9Pd22.8Al68.2 by Czochralski method. The polygrain Mn9Pd21Al70 was synthesized using Al (99.997 mass%), Pd (99.9 mass%) and Mn (99.94 mass%). For homogeneity the sample was arc melted several times and subsequently quenched into water from 800°C. [2000Kaj] concluded the thermal expansion of the Mn9Pd20 Al71 alloy in the temperature range between -263 and 427°C from X-ray diffractometry. The alloy was prepared by arc melting and annealed at 750°C for 72 h. The linear thermal expansion coefficient (T) of the icosahedral phase is about half of that for the pure aluminium phase at room temperature and does not show negative thermal expansion at low temperature. For the -3 phase, the same author estimated Cv(T) above 77°C by the Debye approximation using the reported Debye temperature (a limiting value as T of 188°C). Generalized vibrational densities of states (GVDOS) at T = 23, 327, 527, 727°C have been measured by [2002Sch]. The sample with composition Mn4Pd22Al74 (!´ phase) was prepared from pure elements (Al 99.999%, Pd 99.95% and Mn 99.99%) arc melted and annealed for 6 d at 590°C, 2 d at 825°C. Subsequently it was cooled down to room temperature with a rate of 5°C/min. Neutron scattering experiments under vacuum have been made and the temperature dependence of the heat capacity was calculated from GVDOS data between 2 and 32 K. The specific heat Cp (Fig. 19) was measured in the 2-80 K temperature range by relaxation-type method and the sound carrying velocities was calculated. Notes on Materials Properties and Applications The diffusion coefficient of 63Ni in Mn9Pd21Al70 alloy [2000Zum] and those of 65Zn and 114In in single icosahedral quasicrystals of undefined composition [2000Gal] were measured by radioactive tracers. The activation enthalpies of diffusion Q in icosahedral phase was found to be 209.0 kJ#mol-1 for Ni, 121.31.3 kJ#mol-1 for Zn and 165.95 kJ#mol-1 for In. The diffusion of 103Pd and 195Au in icosahedral quasicrystal (Mn8.5Pd21.3Al70.2) under proton irradiation was investigated by [2000Blu]. Paramagnetic Curie (-21530 K) and Néel (240 K) temperatures were determined by [1968Web] for the MnPd2Al Heusler phase. The magnetic properties of the quasicrystalline and related crystalline phases have been studied in several works [1998Sim, 1999Fis, 1999Hip, 1999Yok, 2000Lai, 2000Sch2, 2000Sim, 2002Miz, 2002Mot]. The neutron scattering experiments on several Al-Mn-Pd liquid alloys with Mn content between 3.5 and 7.2 at.% were carried out by [1998Sim]. The comparison between polarized neutron scattering experiments and magnetic susceptibility suggested that the magnetic moments are present in the liquid state but not in the solid. Temperature dependence of the magnetic susceptibility has been measured in the liquid state and during the solid-liquid transformation. [1998Sim] found that all investigated samples are diamagnetic at room temperature. It is found that the appearance of a magnetic moment on a Mn atom strongly depends on its position in the crystalline lattice [2000Lai]. [1999Hip] investigated the magnetic properties of the Mn6Pd24Al70 alloy, produced by planar flow casting and Landolt-Börnstein New Series IV/11A3
MSIT ®
222
Al–Mn–Pd
annealed at 815°C for 2 h, and three orthorhombic phases: !´-Mn4.6Pd23.5Al71.9, as single crystal obtained by Czochralski growth, T-Mn21.7Pd5.2Al73.1 and T-Mn16.6Pd4.9Al78.5 obtained by Bridgman growth. No magnetic Mn atoms are presented in the Mn6Pd24Al70 alloy, none in the !´ phase nor in the Mn-poor T phase. However, a fraction of the Mn atoms carries magnetic moments in the Mn-rich T phase in which a spin-glass transition occurs at -259°C. The temperature dependence of the dc magnetic susceptibility of these phases has been studied. A strong decrease of the total magnetic moment after annealing was observed by [2000Sch2] using a vibrating sample magnetometer to study an annealed Mn9Pd20Al71 sample which was produced from the melt in a vacuum furnace and cast in a cold Cu mould, annealed at 825°C for 3 h and quenched in water; annealed again at 625°C for 0.5 h and cooled down to room temperature with a rate of 1 K#min-1. The number and the magnitude of the magnetic moments were determined by measuring the field and temperature dependence of the total magnetic moment. [2002Mot] measured the temperature dependence of the magnetic susceptibility of the F2M phase (Mn8Pd21Al71 sample) annealed at different temperatures in the range of -271 to 327°C and found that the susceptibility decreases with increasing temperature from -273 < T < 73°C, it increases with increasing temperature from 73 < T < 327°C. The data could be fitted to the Curie-Weiss law with an additional term proportional to the square of the temperature. The additional term indicates a pseudogap in the electronic density of states at the Fermi energy. The magnetic properties of Mn8.2Pd20.3Al71.5 as-cast samples and those of samples annealed at 727°C for 50 h, have been investigated by [1999Yok] from -268.8 to RT when an external magnetic fields up to 10 kOe is applied. The magnetic order at low temperature was found to be typical for the canonical spin glass phenomenon. The value of diamagnetic component susceptibility (30) was estimated to be about 3.3-4 # 10-7 emu/gOe. The results of electrical properties measurements can be found in [1997Son, 1999Fis, 1999Yok, 2000Tho, 2002Bil, 2002Dem, 2002Miz, 2003Ban, 2003Cap]. In particular, [2002Bil] measured the thermoelectric properties of polygrained icosahedral quasicrystal Mn8.5Pd19.5Al72. They used a kind of self-flux technique, where the ternary melt is first slowly cooled, and then the remaining melt is decanted in the temperature range from -263°C to RT. The electrical resistivity is 'RT = 1.2 m6#cm-1 at RT and increases with decreasing temperature showing a maximum at -153°C. The thermoelectric power is positive in the whole temperature range and at room temperature S = 70 V#K-1. Thermal conductivity at RT is K = 3.4 W/mK and shows a maximum at -243°C and a broad minimum around 148°C. According to [1999Yok] the electrical resistivity in the -268 to -3°C temperature range shows anisotropic dependence on the different symmetrical axes (2-, 3-, 5-fold directions) in the Mn10Pd20Al70 as-grown sample prepared by Czochralski method. After annealing at 627°C for 50 h differences among the electric resistivities along the various directions become smaller than that in the as-grown state. The electric resistivity of Mn10Pd20Al70 thin films was measured from -185 to 12°C by [1997Son] and showed strong negative temperature dependence. The thin films were prepared by laser ablation on fused silica at different deposition temperatures (-196, RT, 165, 350°C) and examined by XRD. Several works were dedicated to the investigation of the mechanical properties of the icosahedral alloys [1996Tan, 1999Yok, 2000Bar, 2000Bru1, 2000Bru2, 2000Feu, 2000Kaj, 2000Mes, 2000Sch3, 2000Sch4, 2002Duq, 2002Lei, 2002Kab, 2002Tak, 2002Tex]. [1996Tan] measured the elastic constants for a Mn6Pd24Al70 alloy (density 5150 kg#m-3) over a temperature range from -269 to 800°C by the rectangular parallelepiped resonance method. At room temperature the results were: Lamé constants (c12) = 74.9, (c44) = 72.4, Young modulus 182 GPa, bulk modulus 123 GPa, Poisson ratio 0.254. The temperature dependence of the above mentioned properties have been determined. The Vickers hardness of the Al70Pd20Mn10 sample presented in Fig. 20 is noticeably different along the 2-fold, 3-fold and 5-fold directions [1999Yok]. The plastic deformation of icosahedral Al-Mn-Pd single quasicrystals and !´ phase has been investigated by [2000Bar, 2000Bru2, 2000Feu, 2000Kle2, 2000Mes, 2000Wan, 2002Kab, 2002Tex]. The oxidation of the Mn9Pd20Al71 icosahedral quasicrystals (sample made by hot isostatically pressing) and Mn8.5Pd21Al70.5 (obtained by arc melting) at 800°C is strongly influenced by the evaporation of Mn, according to [2000Weh]. The reflectivity of the icosahedral phase (Mn8.4Pd21.2Al70.4 sample prepared by induction melting) was found to be very high at low frequency in the far-infrared range, and then it decreases suddenly [2002Dem]. [2000Lan] examined the phase structure and evolution of the quasicrystalline coatings, thermal diffusivity, hardness, and friction coefficient. They found that the quasicrystalline coatings with composition Mn10Pd20 Al70 (prepared by a plasma spray process from gases MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
223
of atomized powders) show cracking after heat treatment at 700 and 800°C. In the as-sprayed condition the measured thermal diffusivity of the coatings is low, but heat treatment increases the thermal diffusivity with increasing heat treatment temperature. The thermal diffusivity of the coatings thus increases as the volume fraction of icosahedral phase decreases and the decagonal and crystalline phases increase. The coefficients of friction of the coatings are reasonably low in the as sprayed condition. Little change is observed after heat treatment up to 600°C, but after heat treatment at 700 and 800°C the values are considerably reduced. [2000Bel] reviewed the main results on transport properties of quasicrystals and introduced the necessary mathematical background. Some questions about practical applications of the quasicrystals are discussed in [2000Cyr, 2000Dub]. Miscellaneous The concentration regime and lattice parameter changes for the supercooled icosahedral phase are reported by [1999Yok]. The temperature dependence of the (quasi-)lattice constants measured by high temperature X-ray diffraction experiments are given in Fig. 21 according to [2000Kaj]. The high-temperature solution growth of large single-grain crystals and quasicrystals of the Al-Mn-Pd system are discussed in [1999Fis, 2000Fis, 2001Can]. Such kind of methods as low-energy electron diffraction (LEED), X-ray photoemission spectroscopy (XPS), scanning tunneling microscopy (STM), ultraviolet-photoemission spectroscopy (UPS), X-ray photoelectron diffraction (XPD), secondary-electron imaging and Auger electron spectroscopy (AES) are used to investigate the atomic and electronic structure, properties, decomposition at elevated temperature, the surface structural phase transitions, voids in the as-grown and annealing single quasicrystalline icosahedral crystal, etc. The results can be found in the publications [1997Gie, 1997She, 1998Bol, 1998Gie, 1998Was, 2000Bol, 2000Cap1, 2000Cap2, 2000Klu, 2000Led, 2000Nau, 2000Ros, 2000Sch5, 2000Sch1, 2002Klu, 2000Klu, 2002Pap, 2003Ebe]. The structural perfection of icosahedral phase has been studied by mechanical spectroscopy [2000Dam] as well as by combined synchrotron X-ray diffractometry and imaging technique [2000Man] and by means of positron annihilation spectroscopy and time-differential dilatometry [2000Bai]. References [1968Web]
[1981Sol1]
[1981Sol2]
[1990Ell]
[1990Tsa1]
[1990Tsa2]
[1991Bee] [1991Bou]
Landolt-Börnstein New Series IV/11A3
Webster, P.J., Tebble, R.S., “Magnetic and Chemical Order in Pd2MnAl in Relation to Order in the Heusler Alloys Pd2MnIn, Pd2MnSn and Pd 2MnSb”, J. Appl. Phys., 39, 471 (1968) (Magn. Prop., Crys. Structure, *, 4) Soltys, J., “X-Ray Diffraction Research of the Order-Disorder Transitions in the Ternary Heusler Alloys B2MnAl (B = Cu, Ni, Co, Pa, R)”, Phys. Status Solidi A, 66(2), 485-491 (1981) (Crys. Structure, Experimental, *, 18) Soltys, J., Kozubski, R., “A Simple Model of the Order-Disorder Phase Transition in Ternary Alloys and Its Application to Several Selected Heusler Alloys”, Phys. Status Solidi, 1, 35-44 (1981) (Crys. Structure, Experimental, *, 23) Ellner, M., “The Structure of the High-Temperature Phase MnAl(h) and the Displacive Transformation from MnAl(h) into Mn 5Al8”, Metall. Trans., A, 21, 1669-1672 (1990) (Crys. Structure, Experimental, 18) Tsai, A.P., Inoue, A., Yokoyama, Y., Masumoto, T., “New Icosahedral Alloys with Superlattice Order in the Al-Pd-Mn System Prepared by Rapid Solidification”, Philos. Mag. Lett., 61(1), 9-14 (1990) (Crys. Structure, Experimental, 10) Tsai, A.P., Inoue, A., Yokoyama, Y., Masumoto, T., “Stable Icosahedral Al-Pd-Mn and Al-Pd-Re Alloys”, Mater. Trans., JIM, 31(2), 98-103 (1990) (Crys. Structure, Experimental, 19) Beeli, C., Nissen, H.-U., Robadey, J., “Stable Al-Mn-Pd Quasicrystals”, Philos. Mag. Lett., 63(2), 87-95 (1991) (Equi. Diagram, Experimental, 24) Boudart, M., de Boissieu, M., Janot, C., Dubois, J.M., Dong, C., “The Structure of the Icosahedral AlPdMn Quasicrystal”, Philos. Mag. Lett., 64(4), 197-206 (1991) (Crys. Structure, Experimental, 18) MSIT ®
224 [1991Don]
[1991Tsa]
[1991Yok]
[1992Bou]
[1992Li]
[1992Was]
[1992Yok]
[1993Aud]
[1993Bou]
[1993Dau] [1993Hir1]
[1993Hir2]
[1993Ran]
[1993Ste] [1993Sun]
[1993Tsa]
[1993Was]
MSIT®
Al–Mn–Pd Dong, C., Dubois, J.M., Boissieu, M., Boundard, M., Janot, C., “Growth of Stable Al-Pd-Mn Icosahedral Phase”, J. Mater. Res., 6(12), 2637-2645 (1991) (Crys. Structure, Experimental, 29) Tsari, A.-P., Yokoyama, Y., Inoue, A., Masumoto, T., “Formation, Mictostructure, Chemical Long-Range Order, and Stability of Quasicrystals in Al-Pd-Mn Alloys”, J. Mater. Res., 6(12), 2646-2652 (1991) (Crys. Structure, Equi. Diagram, Experimental, 19) Yokoyama, Y., Tsai, A.-P., Inoue, A., Masumoto, T., “Production of the Quasicrystalline Al-Pd-Mn Alloys with Large Single Domain Size”, Mater. Trans., JIM, 32(12), 1089-1097 (1991) (Abstract, Crys. Structure, Magn. Prop., 15) Boudard, M., de Boissieu, M., Janot, C., Heger, G., Beeli, C., Nissen, H.-U., Vincent, H., Ibberson, R., Audier, M., Dubois, J.M., “Neutron and X-Ray Single-Crystal Study of the AlPdMn Icosahedral Phase”, J. Phys.: Condens. Matter, 4, 10149-10168, (1992) (Crys. Structure, Experimental, 48) Li, X.Z., Kuo, K.H., “The Structural Model of Al-Mn Decagonal Quasicrystal Based on a New Al-Mn Approximant”, Philos. Mag. B, B65(3), 525-533 (1992) (Crys. Structure, Experimental, 9) Waseda, A., Morioka, H., Kimura, K., Ino, H., “An Icosahedral Quasicrystal and Its Cubic Approximant in the Al-Pd-Mn System”, Philos. Mag. Lett., 65(1), 25-32 (1992) (Crys. Structure, Experimental, *, 20) Yokoyama, Y., Miura, T., Tsai, A., Inoue, A., Masumoto, T., “Preparation of a Large Al70Pd20Mn10 Single-Quasicrystal by the Czochralski Method and Its Electrical Resistivity”, Mater. Trans. Jpn. Inst. Metals, 33(2), 97-101 (1992) (Equi. Diagram, Crys. Structure, Electr. Prop., Experimental, 15) Audier M., Durand-Charre M., De Boissieu M., “Aluminum-Palladium-Manganese Phase Diagram in The Region of Quasicrystalline Phases”, Philos. Mag. B, 68(5), 605-618 (1993) (Equi. Diagram, Crys. Structure, Experimental, *, 20) Boudard, M., de Boissieu, M., Janot, C., Heger, G., Beeli, C., Nissen, H.-U., Vincent, H., Audier, M., Dubois, J.M., “Atomic Structure of the Al-Pd-Mn Icosahedral Phase”, J. Non-Cryst. Solids, 153&154, 5-9 (1993) (Crys. Structure, Experimental, *, 21) Daulton, T.L., Kelton, K.F., “The Orthrhombic (Al11Mn4)-Pd Decagonal Approximant”, Philos. Mag. B, 68(5), 697-711 (1993) (Crys. Structure, Experimental, 13) Hiraga, K., Sun, W., “The Atomic Arrangement of an Al-Pd-Mn Decagonal Quasicrystal Studied by High-Resolution Electron Microscopy”, Philos. Mag. Lett., 67(2), 117-123 (1993) (Crys. Structure, Experimental, 7) Hiraga, K., Kaneko, M., Matsuo, Y., Hashimoto, S., “The Structure of Al3Mn: Close Relashionship to Decagonal Quasicrystals”, Philos. Mag. B, B67(2), 193-205 (1993) (Crys. Structure, Experimental, 12) Ran, Q., “Aluminium-Manganese-Palladium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16728.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 3) Steurer, W., “Comparative Structure Analysis of Several Decagonal Phases”, J. Non-Cryst. Solids, 153-154, 92-97 (1993) (Crys. Structure, Experimental, 16) Sun, W., Hiraga, K., “Interface Structure Between Decagonal and Icosahedral Quasicrystals in Al-Pd-Mn Alloy”, Philos. Mag. Lett., 67(3), 159-164 (1993) (Crys. Structure, Experimental, 14) Tsai, A.-P., Yokoyama, Y., Inoue, A., Masumoto, T., “Chemically Driven Structural Change in Quasicrystalline Al-Pd-Mn Alloys”, Met. Abstr. Light Metals and Alloys, 26, 32 (1993) (Equi. Diagram, Experimental) Waseda, A., Araki, K., Kimura, K., Ino, H., “Quasicrystals and Approximants in the Al-Co-(Fe, Ru) and Al-Pd-Mn Systems”, J. Non-Cryst. Solids, 153-154, 635-639 (1993) (Crys. Structure, Equi. Diagram, Experimental, *, 19) Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd [1994Aud]
[1994Bee]
[1994Li1]
[1994Li2] [1994Shi]
[1995Bee1]
[1995Bee2]
[1995Boi]
[1995Goe1]
[1995Goe2]
[1995Ish] [1996Bou]
[1996Kle] [1996Sun]
[1996Tan]
[1996Yam] [1997Ama]
Landolt-Börnstein New Series IV/11A3
225
Audier, M., Duneau, M., Vacher, M., “Structural Approach of the Decagonal and Approximant Phases in the Al-Pd-Mn System: an Application of the Linear Phason Strain Field Theory”, Advances in Physical Metallurgy, Gordon & Breach, 61-72 (1994) (Crys. Structure, *, 9) Beeli, C., Horiuchi, S., “The Structure and its Reconstruction in the Decagonal Al70Mn17 Pd13 Quasicrystal”, Philos. Mag. B, B70(2), 215-240 (1994) (Crys. Structure, Experimental, 33) Li, X.Z., Dubois, J.M., “Structural Sub-Units of the Al-Mn-Pd Decagonal Quasicrystal Derived from the Structure of the T3 Al-Mn-Zn Phase”, J. Phys.: Condens. Matter, 6, 1653-1662 (1994) (Crys. Structure, Theory, 22) Li, H.L., Kuo, K.H., “Some New Crystaline Approximantas of Al-Pd-Mn Quasicrystals”, Philos. Mag. Lett., 70(1), 55-62 (1994) (Crys. Structure, Experimental, 32) Shi, N.C., Li, X.Z., Ma, Z.S., Kuo, K.H., “Crystalline Phases Related to a Decagonal Quasicrystal. I. A Single-Crystal X-Ray Diffraction Study of the Orthorhombic Al3Mn Phase”, Acta Crystallogr., Sect. B: Struct. Crystallogr. Crys. Chem., B50, 22-30 (1993) (Crys. Structure, Experimental, 24) Beeli, C., Stadelmann, P., Gödecke, T., Lück, R., “The Decagonal Al-Mn-Pd Phase and its Modification”, Proc. Intern. Conf. on Aperiodic Crystals (Aperiodic 94), Chapius, G., Paciorek, W., (Eds.), World Scientific Publ., Singapore 1995, 361-365 (1995) (Crys. Structure, Experimental, *, 7) Beeli, C., Stadelmann, P., Lueck, R., Gödecke, T., “Decagonal Al-Mn-Pd Quasicrystals Free of Linear Phason Strain”, Proc. 5th Intern. Conf. on Quasicrystals, Janot, C., Mosseri, R. (Eds.), World Scientific. Publ.; Singapore 1995, 680-683 (1995) (Crys. Structure, Experimental, *, 12) de Boissieu, M., Boudard, M., Hennion, B., Bellisent, R., Kycia, S., Goldman, A.I., Janot, C., Audier, M., “Diffuse Scattering and Phason Elasticity in the AlPdMn Icosahedral Phase”, Phys. Rev. Lett., 75(1), 89-92 (1995) (Crys. Structure, Experimental, 24) Goedecke, T., Lueck, R., “The Aluminium-Palladium-Manganese System in the Range from 60 to 100 % Al”, Z. Metallkd., 86(2), 109-121 (1995) (Equi. Diagram, Experimental, #, *, 30) Goedecke, T., Lueck, R., Beeli, C., “The Formation of Quasicrystalline Alloys from the Melt in the Aluminium-Palladium-Manganese System”, Proc. 5th Int. Conf. Quasicryst., 644-647 (1995) (Equi. Diagram, Experimental, #, *, 14) Ishimasa, T., “Superlattice Ordering in the Lowe-Temperature Icosahedral Phase of Al-Pd-Mn”, Philos. Mag. Lett., 71(1), 65-73 (1995) (Crys. Structure, *, 14) Boudard, M., Klein, H., de Boissieu, M., Audier, M., “Structure of Quasicyrstalline Approximant Phase in the Al-Pd-Mn System”, Philos. Mag. A, 74(4), 939-956 (1996) (Crys. Structure, *, 31) Klein, H., Audier, M., Boudard, M., De Boissieu, M., “Phason Defects in Al-Pd-Mn Approximant Phases”, Philos. Mag. A, 73(2), 309-331 (1996) (Crys. Structure, 31) Sun, W., Hiraga, K., “High-Resolution Transmission Electron Microscopy of the Al-Pd-Mn Decagonal Quasicrystal with 1-6nm Periodicity and its Crystalline Approximants”, Philos. Mag. A, 73(4), 951-971 (1996) (Crys. Structure, Experimental, 19) Tanaka, K., Mitarai, Y., Koiwa, M., “Elastic Constants of Al-Based Icosahedral Quasicrystals”, Philos. Mag. A, 73(6), 1715-1723 (1996) Crys. Structure, Mechan. Prop., Experimental, 18) Yamamoto, A., “Crystallography of Quasiperiodic Crystals”, Acta Crystallogr., Sect. A: Found. Crystallogr., 52, 509-560 (1996) (Calculation, Crys. Structure, Review, 211) Amazit, Y., Perrin, B., Fischer, M., Itie, J.P., Polian, A., “X-Ray Diffraction Measurements in Icosahedral Al-Pd-Mn up to 40 GPa”, Philos. Mag. A, 75(6), 1677-1688 (1997) (Crys. Structure, Experimental, 23)
MSIT ®
226 [1997Gie]
[1997Hae]
[1997Ina]
[1997Kle]
[1997Kra]
[1997Mat]
[1997Oka] [1997She]
[1997Son]
[1997Zur]
[1998Ber]
[1998Boi]
[1998Bol]
[1998Gie]
[1998Mat]
MSIT®
Al–Mn–Pd Gierer, M., van Hove, M.A., Goldman, A.I., Shen, Z., Chang, S.-L., Jenks, C.J., Zhang, C.-M., Thiel, P.A., “Structural Analysis of the Fivefold Symmetric Surface of the Al70Pd21Mn9 Quasicrystal by Low Energy Electron Diffraction”, Phys. Rev. Lett., 78(3), 467-470 (1997) (Crys. Structure, Experimental, 17) Häussler, D., Beeli, C., Nissen, H.-U., “One-Dimensionally Modulated Quasicrystal Phase Related to Icosahedral Al-Mn-Pd”, Philos. Mag. Lett., 75(2), 117-124 (1997) (Crys. Structure, Experimental, 14) Inaba, A., Tsai, A.P., Shibata, K., “Vibrational Properties of Quasicrystals of Al-Cu-Ru, Al-Pd-Re and Al-Pd-Mn Deduced from Heat Capacities”, Proc. of the 6th International Conf. on Quasicrystals, Takeuchi, S., Fujiwara, T. (Eds.) (ICQ6), World Scientific, Singapore, 1997, p.443-450 (Crys. Structure, Experimental, 10) Klein, H., Boudard, M., Audier, M., de Boissieu, M., Vincent, H., Beraha, L., Duneau, M., “The T-Al3(Mn, Pd) Quasicrystalline Approximant: Chemical Oreder and Phason Defect”, Philos. Mag. A, 75(4), 197-208 (1997) (Crys. Structure, Experimental, *, 21) Krajci, M., Hafner, J., Mihalkovic, M., “Atomic and Electronic Structure of Decagonal Al-Pd-Mn Alloys and Approximant Phases”, Phys. Rev. B, 55(2), 843-855 (1997) (Crys. Structure, Experimental, 54) Matsuo, Y., Kaneko, M., Yamanoi, T., Kaji, N., Sugiyama, K., Hiraga, K., “The Structure of an Al3Mn-Type Al3(Mn, Pd) Crystal Studied by Single-Crystal X-Ray Diffraction Analysis”, Philos. Mag. Lett., 76(5), 357-362 (1997) (Crys. Structure, Experimental, *, 9) Okamoto, H., “Al-Mn (Aluminum-Manganese)”, J. Phase Equilib., 18(4), 398-399 (1997) (Crys. Structure, Equi. Diagram, Review, 8) Shen, Z., Jenks, C.J., Anderegg, J., Delaney, D.V., Kograsso, T.A., Thiel, P.A., Goldman, A.I., “Structure and Stability of the Twofold Surface of Icosahedral Al-Pd-Mn by Low-Energy Electron Diffraction and X-ray Photoemission Spectroscopy”, Phys. Rev. Lett., 78(6), 1050-1053 (1997) (Crys. Structure, Experimental, 22) Sonsky, J., Jelinek, M., Jastrabik, L., Studnicka, V., Chvostova, D., Bryknar, Z., “Study of Quasicrystalline Thin Films Based on Al-Pd-Mn and Al-Cu-Fe Prepared by PLD”, Czechoslov. J. Phys., 47(10), 1019-1024 (1997) (Crys. Structure, Electr. Prop., Experimental, 16) Zurkirch, M., Crescenzi, M.D., Erbudak, M., Hochstrasser, M., “Comparison of the Structure of AlPd and Al 70Pd20Mn10”, Phys. Rev. B, 55(14), 8808-8811 (1997) (Crys. Structure, Experimental, 21) Beraha, L., Duneau, M., Klein, H., Audier, M., “Phason Defects in Al-Pd-Mn Approximant Phases: Another Example”, Philos. Mag. A, 78(2), 345-372 (1998) (Crys. Structure, Experimental, *, 23) de Boissieu, M., Boudard, M., Ishimasa, T., Elkaim, E., Lauriat, J.P., Letoublon, A., Audier, M., Duneau, M., Davroski, A., “Reversible Transformation Between an Icosahedral Al-Pd-Mn Phase and a Modulated Structure of Cubic Symmetry”, Philos. Mag. A, 78(2), 305-326 (1998) (Crys. Structure, Experimental, *,39) Bolliger, B., Erbudak, M., Vvedensky, D.D., Zurkirch, M., “Surface Strcutural Transitions on the Icosahedral Quasicrystal Al70Pd20Mn10”, Phys. Rev. Lett., 80(24), 5369-5372 (1998) (Crys. Structure, Experimental, 25) Gierer, M., Van Hove, M.A., Goldman, A.I., Shen, Z., Chang, S.-L., Pinhero, P.J., Jenks, C.J., Anderegg, J.W., Zhang, C.-M., Thiel, P.A., “Fivefold Surface of Quasicrystalline AlPdMn: Strcuture Determination Using Low-Energy-Electron Diffraction”, Phys. Rev. B: Condens. Matter, 57(13), 7628-7641 (1998) (Crys. Structure, Experimental, 55) Matsuo, Y., Yamamoto, Y., Ishii, Y., “Investigation of Phason Strains in Decagonal Al-Pd-Mn Single Qasicryustals by Means of X-ray Diffraction”, J. Phys.: Condens. Matter., 10, 983-994 (1998) (Crys. Structure, Experimental, 12)
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd [1998Sim]
[1998Wae]
[1998Wan]
[1998Was]
[1999Aud]
[1999Cap]
[1999Fis]
[1999Gru] [1999Hip]
[1999Yok]
[2000Bai]
[2000Bar]
[2000Bee1]
[2000Bee2] [2000Bel]
[2000Blu]
[2000Bol]
Landolt-Börnstein New Series IV/11A3
227
Simonet, V., Hippert, F., Klein, H., Audier, M., Bellissent, R., Fischer, H., Murani, A.P., Boursier, D., “Local Order and Magnetism in Liquid Al-Pd-Mn Alloys”, Phys. Rev. B, 58(10), 6273-6286 (1998) (Crys. Structure, Experimental, Magn. Prop., 63) Waelti, C., Felder, E., Chernikov, M.A., Ott, H.R., de Boissieu, M., Janot, C., “Lattice Excitations in Icosahedral Al-Mn-Pd and Al-Re-Pd”, Phys. Rev. B: Condens. Matter, 57(17), 10504-10511 (1998) (Experimental, Thermodyn., 46) Wang, R., Feuerbacher, M., Yang, W., Urban, K., “Stacking Faults in High-Temperature-Deformed Al-Pd-Mn Icosahedral Quasicrystals”, Philos. Mag. A, 78(2), 273-284 (1998) (Experimental, 19) Waseda, Y., Suzuki, S., Urban, K., “Novel Morphology of Voids in Single-Quasicrystalline Icosahedral Al70.5 Pd21.0Mn8.5”, Z. Naturforsch. A, 53A, 679-683 (1998) (Crys. Structure, Experimental, 16) Audier, M., Duneau, M., de Boissieu, M., Boudard, M., Letoublon, A., “Superlattice Ordering of Cubic Symmetry in an Icosahedral Al-Pd-Mn Phase”, Philos. Mag. A, 79(2), 255-270 (1999) (Crys. Structure, Experimental, *, 9) Capitan, M.J., Calvayrac, Y., Quivy, A., Joulaud, J.L., Lefebvre, S., Gratias, D. “X-Ray Diffuse Scattering from Icosahedral Al-Pd-Mn Quasicrystals”, Phys. Rev. B, 60(9), 6398-6404 (1999) (Crys. Structure, Experimental, 23) Fisher, I.R., Kramer, M.J., Wiener, T.A., Islam, Z., Ross, A.R., Lograsso, T.A., Kracher, A., Goldman, A.I., Canfield, P.C., “On the Growth of Icosahedral Al-Pd-Mn Quasicrystals from the Ternary Melt”, Philos. Mag. B, 79(10), 1673-1684 (1999) (Experimental, Phys. Prop., 20) Grushko, B., Yurechko, M., Tamura, N., “A Contribution to the Al-Pd-Mn Phase Diagram”, J. Alloys Compd., 290, 164-171 (1999) (Equi. Diagram, Experimental, #, *, 24) Hippert, F., Simonet, V., Trambly de Laissardiere, G., Audier, M., Calvayarac, Y., “Magnetic Properties of AlPdMn Appximant Phases”, J. Phys.: Condens. Matter, 11, 10419-10435 (1999) (Crys. Structure, Experimental, Magn. Prop., 48) Yokoyama, Y., Yamada, Y., Fukaura, K., Sunada, H., Note, R., Inoue, A., Sugiyama, K., Hiraga, K., “Strain Affected Properties of Icosahedral Al-Pd-Mn Single Ingot”, Jpn. J. Appl. Phys., 38(1)(3A), 1495-1499 (1999) (Crys. Structure, Equi. Diagram, Electr. Prop., Magn. Prop., Experimental, *, 16) Baier, F., Mueller, M.A., Grushko, B., Schaefer, H.-E., “Atomic Defects in Quasicrystals: an Approach with Positron Annihilation Spectroscopy and Time-Differential Dilatometry”, Mater. Sci. Eng. A, 294-296, 650-653 (2000) (Crys. Structure, Experimental, 13) Bartsch, M., Geyer, B., Haeussler, D., Feuerbacher, M., Urban, K., Masserschmidt, U., “Plastic Properties of Icosahedral Al-Pd-Mn Single Quasicrystals”, Mater. Sci. Eng. A, 294-296, 761-764 (2000) (Crys. Structure, Experimental, Phys. Prop., 14) Beeli, C., Soltmann, C., Poon, S.J., “Relationship of Phason Strain and Electronic Properties in Icosahedral Al-Pd-(Re,Mn) and Al-Cu-Os”, Mater. Sci. Eng. A, 294-296, 531-534 (2000) (Crys. Structure, Experimental, 14) Beeli, C., “High-Resolution Electron Microscopy of Quasicrystals”, Mater. Sci. Eng. A, 294-296, 23-28 (2000) (Crys. Structure, Experimental, 40) Bellissard, J., “Anomalous Transport: Results, Conjectures and Applications to Quasicrystals”, Mater. Sci. Eng. A, 294-296, 450-457 (2000) (Crys. Structure, Phys. Prop., Review, 56) Blueher, R., Frank, W., Grushko, B., “Diffusion of 103Pd and 195Au in Icosahedral Al70.2Pd21.3Mn8.5 under Proton Irradiation”, Mater. Sci. Eng. A, 294-296, 689-692 (2000) (Crys. Structure, Experimental, Phys. Prop., 10) Bolliger, B., Erbudak, M., Hensch, A., Vvedensky, D.D., “Surface Structural Phase Transitions on Icosahedral Al-Pd-Mn”, Mater. Sci. Eng. A, 294-296, 859-862 (2000) (Crys. Structure, Experimental, 10)
MSIT ®
228 [2000Bru1]
[2000Bru2]
[2000Cap1]
[2000Cap2]
[2000Cyr]
[2000Dam]
[2000Dub]
[2000Dun]
[2000Eda]
[2000Feu]
[2000Fis]
[2000Fra]
[2000Fre] [2000Gal]
[2000Gru]
[2000Gwo]
MSIT®
Al–Mn–Pd Brunet, P., Zhang, L.M., Sordelet, D.J., Besser, M., Dubois, J-M., “Comparative Study of Microstructural and Tribological Properties of Sintered, Bulk Icosahedral Samples”, Mater. Sci. Eng. A, 294-296, 74-78 (2000) (Mechan. Prop., Experimental, 6) Brunner, D., Plachke, D., Carstanjen, H.D., “The Strain-Softering Phenomenon of Icosahedral Al-Pd-Mn Single Quasicrystals”, Mater. Sci. Eng. A, 294-296, 773-776 (2000) (Mechan. Prop., Experimental, 19) Cappello, G., Schmithuesen, F., Chevrier, J., Comin, F., Stierle, A., Formoso, V., Boissieu, M., Boudard, M., Lograsso, T.A., Jenks, C., Delaney, D., “Characterization of Surface Morphologies at the Al-Pd-Mn Fivefold Surface”, Mater. Sci. Eng. A, 294-296, 822-825 (2000) (Crys. Structure, Experimental, 14) Cappello, G., Dechelette, A., Schmithuesen, F., Decossas, S., Chevrier, J., Comin, F., Formoso, V., Boissieu, M., Jach, T., Colella, R., Lograsso, T.A., Jenks, C., Delaney, D., “Bulk and Surface Evidence for the Long-Range Spatial Modulation of X-Ray Absorption in the Al-Pd-Mn Quasicrystal at Bragg Incidence”, Mater. Sci. Eng. A, 294-296, 863-866 (2000) (Crys. Structure, Experimental, 11) Cyron-Lackmann, F., “Quasicrystals as Potential for Thermoelectric Materials”, Mater. Sci. Eng. A, 294-296, 611-612 (2000) (Calculation, Crys. Structure, Electr. Prop., Phys. Prop., Thermal Conduct., 15) Damson, B., Weller, M., Feuerbacher, M., Grushko, B., Urban, K., “Mechanical Spectroscopy of i-Al-Pd-Mn and d-Al-Ni-Co”, Mater. Sci. Eng. A, 294-296, 806-809 (2000) (Crys. Structure, Experimental, 18) Dubois, J-M., “New Prospects from Potential Applications of Quasicrystalline Materials”, Mater. Sci. Eng. A, 294-296, 4-9 (2000) (Crys. Structure, Experimental, Phys. Prop., Review, 38) Duneau, M., “Covering Clusters in the Katz-Gratias Model of Icosahedral Quasicrystals”, Mater. Sci. Eng. A, 294-296, 192-198 (2000) (Calculation, Crys. Structure, Experimental, 34) Edagawa, K., Kajiyama, K., “High Temperature Specific Heat of Al-Pd-Mn and Al-Cu-Co Quasicrystals”, Mater. Sci. Eng. A, 294-296, 646-649 (2000) (Crys. Structure, Thermodyn., Experimental, *, 21) Feuerbacher, M., Klein, H., Bartsch, M., Messerschmidt, U., Urban, K., “A Comparative Study of the Plastic Behavior of Icosahedral and pri-Al-Pd-Mn”, Mater. Sci. Eng. A, 294-296, 736-741 (2000) (Crys. Structure, Experimental, Phys. Prop., 23) Fisher, I.R., Kramer, M.J., Islam, Z., Wiener, T.A., Kracher, A., Ross, A.R., Lograsso, T.A., Goldman, A.I., Canfield, P.C., “Growth of Large Single-Grain Quasicrystals from High-Temperature Metallic Solutions”, Mater. Sci. Eng. A, 294-296, 10-16 (2000) (Crys. Structure, Experimental, 22) Fradkin, M.A., “Finite-Resolution Correction to the Diffraction Intensity in Icosahedral Quasicrystals”, Mater. Sci. Eng. A, 294-296, 319-322 (2000) (Calculation, Crys. Structure, Experimental, 4) Frey, F., “Disorder Diffuse Scattering of Decagonal Quasicrystals”, Mater. Sci. Eng. A, 294-296, 178-185 (2000) (Crys. Structure, Experimental, 15) Galler, R., Mehrer, H., “Diffusion in Icosahedral Al-Pd-Mn Quasicrystals: Temperature and Pressure Dependence”, Mater. Sci. Eng. A, 294-296, 693-696 (2000) (Crys. Structure, Experimental, Phys. Prop., 17) Grushko, B., “Composition and Presipitation Behavior of Icosahedral Al-Pd-Mn Quasicrystals”, Mater. Sci. Eng. A, 294-296, 45-48 (2000) (Crys. Structure, Experimental, *, #, 16) Gwozdz, J., Grushko, B., Surowiec, M., “Mosaic Structure of Single Al-Pd-Mn Icosahedral Quasi-Crystals”, Mater. Sci. Eng. A, 294-296, 49-52 (2000) (Crys. Structure, Experimental, 7)
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd [2000Hir]
[2000Jac] [2000Jan] [2000Kaj]
[2000Kle1]
[2000Kle2]
[2000Kle3]
[2000Klu]
[2000Lai]
[2000Lan]
[2000Led]
[2000Let]
[2000Man]
[2000Mes]
[2000Nau]
[2000Nic]
[2000Qua]
Landolt-Börnstein New Series IV/11A3
229
Hirai, I., Ishimasa, T., Letoublon, A., Boudard, M., Boissieu, M., “Formation Conditions of Two Quasiperiodic Modifications of Al-Pd-Mn Icosahedral Phase Studied by Annealing Method”, Mater. Sci. Eng. A, 294-296, 33-36 (2000) (Crys. Structure, Experimental, *, 9) Jach, T., “Quasicrystal Element Correlations from X-Ray Standing Waves”, Mater. Sci. Eng. A, 294-296, 315-318 (2000) (Calculation, Crys. Structure, Theory, 13) Janot, C., Loreto, L., Farinato, R., “Clusters in Quasicrystals: Tiling Versus Covering and Porosity”, Mater. Sci. Eng. A, 294-296, 405-408 (2000) (Assessment, Crys. Structure, 19) Kajiyama, K., Edagawa, K., Suzuki, T., Takeuchi, S., “Thermal Expansion of Icosahedral Al-Pd-Mn and Decagonal Al-Cu-Co Quasicrystals”, Philos. Mag. Lett., 80(1), 49-56 (2000) (Crys. Structure, Experimental, Mechan. Prop., 19) Klein, H., Durand-Charre, M., Audier, M., “Liquid-Solid Equilibria in the Quasicrystalline Regions of the Al-Pd-Mn Phase Diagram”, J. Alloys Compd., 296, 128-137 (2000) (Equi. Diagram, Experimental, *, 41) Klein, H., Feuerbacher, M., Schall, P., Urban, K., “Bending Experiments on the 3´-(Al-Pd-Mn) Quasicrystal Approximant”, Philos. Mag. Lett., 80(1), 11-18 (2000) (Crys. Structure, Experimental, 12) Klein, H., Feuerbacher, M., Urban, K., “Dislocation in Al-Pd-Mn Approximants: a High Resolution Electron Microscopy Study”, Mater. Sci. Eng. A, 294-296, 769-772 (2000) (Crys. Structure, Experimental, 7) Kluge, F., Ebert, P., Grushko, B., Urban, K., “Influence of Grown-in Voids on the Structure of Cleaved Icosahedral Al-Pd-Mn Quasicrystal Surfaces”, Mater. Sci. Eng. A, 294-296, 874-877 (2000) (Crys. Structure, Experimental, 22) Laissardiere, G.T., Mayou, D., “Conditions on the Occuerence of Magnetic Moments in Quasicrystals and Related Phases”, Mater. Sci. Eng. A, 294-296, 621-624 (2000) (Calculation, Crys. Structure, Magn. Prop., 18) Lang, C.I., Sordelet, D.J., Besser, M.F., Shechtman, D., Biancaniello, F.S., Gonzales, E.J., “Quasicrystalline Coatings: Thermal Evolution of Structure and Properties”, J. Mater. Res., 15(9), 1894-1904 (2000) (Experimental, Mechan. Prop., Phys. Prop., 41) Ledieu, J., Muryn, C.A., Thornton, G., Cappello, G., Chevrier, J., Diehl, R.D., Lograsso, T.A., Delaney, D., McGrath, R., “Decomposition of the Five-Fold Surface of Al70Pd21Mn9 at Elevated Temperature”, Mater. Sci. Eng. A, 294-296, 871-873 (2000) (Crys. Structure, Experimental, 22) Letoublon, A., Ishimasa, T., de Boissieu, M., Boudard, M., Hennion, B., Mori, M., “Stability of the F2-(Al-Pd-Mn) Phase”, Philos. Mag. Lett., 80(4), 205-213 (2000) (Crys. Structure, Equi. Diagram, Experimental, *, 14) Mancini, L., Letoublon, A., Agliozzo, S., Wang, J., Gastaldi, J., Boissieu, M., Haertwig, J., Klein, H., “Effect of Annealing on the Structural Perfection of Al-Pd-Mn Icosahedral Quasicrystal Grains”, Mater. Sci. Eng. A, 294-296, 57-60 (2000) (Crys. Structure, Experimental, 20) Messerschmidt, U., Haeussler, D., Bartsch, M., Geyer, B., Feuerbacher, M., Urban, K., “Microprocesses of the Plastic Deformation of Icosahedral Al-Pd-Mn Single Quasicrystals”, Mater. Sci. Eng. A, 294-296, 757-760 (2000) (Crys. Structure, Experimental, 15) Naumovic, D., Aebi, P., Schlapbach, L., Beeli, C., “Atomic and Electronic Structure of Five-Fold i-Al-Pd-Mn Surfaces”, Mater. Sci. Eng. A, 294-296, 882-885 (2000) (Crys. Structure, Experimental, Phys. Prop., 28) Nicula, R., Jianu, A., Grigoriu, C., Barfels, T., Burkel, E., “Laser Ablation Synthesis of Al-Based Icosahedral Powders”, Mater. Sci. Eng. A, 294-296, 86-89 (2000) (Crys. Structure, Experimental, Mechan. Prop., 12) Quandt, A., Elser, V., Kresse, G., Hafner, J., “An Ab Initio Based Structure Model of i(Al-Pd-Mn)”, Mater. Sci. Eng. A, 294-296, 351-354 (2000) (Calculation, Crys. Structure, 19) MSIT ®
230 [2000Ros]
[2000Sch1]
[2000Sch2]
[2000Sch3]
[2000Sch4]
[2000Sch5]
[2000Shr]
[2000Sim]
[2000Sta] [2000Ste1] [2000Ste2] [2000Tho]
[2000Uch]
[2000Wan]
[2000Weh]
[2000Yam]
[2000Zum]
MSIT®
Al–Mn–Pd Ross, A.R., Wiener, T.A., Fisher, I.R., Canfield, P.C., Lograsso, T.A., “Formation and Morphological Development of Porosity in Icosahedral Al-Pd-Mn Alloys”, Mater. Sci. Eng. A, 294-296, 53-56 (2000) (Crys. Structure, Experimental, 11) Schaub, T., Delahaye, J., Berger, C., Grenet, T., Guyot, H., Belkhou, R., Taleb-Ibrahimi, A., Prejean, J.J., Calvayrac, Y., “High Resolution Experiment on the Electronic Density of States in Icosahedral-Al-Pd-Mn”, Mater. Sci. Eng. A, 294-296, 512-515 (2000) (Crys. Structure, Experimental, 24) Scheffer, M., Suck, J.-B., “Influence of Vacancies on the Magnetic Properties of Icosahedral Al7.10 Pd20.0Mn9.0”, Mater. Sci. Eng. A, 294-296, 629-632 (2000) (Crys. Structure, Experimental, Magn. Prop., 14) Schurack, F., Eckert, J., Schultz, L., “Quasicrystalline Al-Alloys with High Strength and Good Ductility”, Mater. Sci. Eng. A, 294-296, 164-167 (2000) (Crys. Structure, Experimental, Mechan. Prop., 6) Schall, P., Feuerbacher, M., Bartsch, M., Messerschmidt, U., Urban, K., “Dislocation Arrangement and Density in Deformed Al-Pd-Mn Single-Quasicrystals”, Mater. Sci. Eng. A, 294-296, 765-768 (2000) (Calculation, Crys. Structure, Experimental, 9) Schmithuesen, F., Boissieu, M., Boudard, M., Chevrier, J., Comin, F., “Electron Energy Loss Spectroscopy Investigation of Volume and Surface Plasmonts at the Al-Pd-Mn Fivefold Surface”, Mater. Sci. Eng. A, 294-296, 867-870 (2000) (Crys. Structure, Experimental, 15) Shramchenko, N., Klein, H., Caudron, R., Bellissent, R., “Comparison of Local Order in Icosahedral Al-Pd-Mn Quasicrystal and in Approximant Phase by Thermal Neutron Scattering”, Mater. Sci. Eng. A, 294-296, 335-339 (2000) (Crys. Structure, Experimental, 10) Simonet, V., Hippert, F., Audier, M., Calvayras, Y., “Magnetism of Approximants in the Al-Mn and Al-Pd-Mn Systems”, Mater. Sci. Eng. A, 294-296, 625-628 (2000) (Crys. Structure, Experimental, Magn. Prop., 24) Stadnic, Z.M., “Photoemission Studies of Qusicrystals”, Mater. Sci. Eng. A, 294-296, 470-474 (2000) (Crys. Structure, Electr. Prop., Experimental, 26) Steurer, W., “The Quasicrystal-to-Crystal Transformation. I. Geometrical Principles”, Z. Kristallogr., 215, 323-334 (2000) (Calculation, Crys. Structure, 44) Steurer, W., “Geometry of Quasicrystal-to-Crystal Transformations”, Mater. Sci. Eng. A, 294-296, 268-271 (2000) (Assessment, Crys. Structure, 12) Thompson, E., Vu, P.D., Pohl, R.O., “Glasslike Lattice Vibrations in the Quasicrystal Al72.1Pd20.7Mn7.2”, Phys. Rev. B, 62(17), 11437-11443 (2000) (Crys. Structure, Experimental, Phys. Prop., Thermal Conduct., 49) Uchiyama, H., Takahashi, Y., Sato, K., Kanazawa, I., Kimura, K., Komori, F., Suzuki, R., Ohdaira, T., Tamura, R., Takeuchi, S., “Stable Quasicrystals Studied by Means of the Slow Positron Beam”, Nucl. Instrum. Methods Phys. Res./B, B171, 245-250 (2000) (Crys. Structure, Experimental, 21) Wang, R., Yang, W., Gui, J., Urban, K., “Dislocation Mechanism of High-Temperature Plastic Deformation of Al-Cu-Fe and Al-Pd-Mn Icosahedral Quasicrystals”, Mater. Sci. Eng. A, 294-296, 742-747 (2000) (Crys. Structure, Experimental, 18) Wehner, B.I., Koster, U., Rudiger, A., Pieper, A., Sordelet, D.J., “Oxidation of Al-Cu-Fe and Al-Pd-Mn Quasicrystals”, Mater. Sci. Eng. A, 294-296, 830-833 (2000) (Crys. Structure, Experimental, 16) Yamamoto, A., Hiraga, K., “Six-Dimensional Model of an i-Al-Pd-Mn Quasicrystal Compatible with its 2/1 Approximant”, Mater. Sci. Eng. A, 294-296, 228-231 (2000) (Crys. Structure, Experimental, Review, 3) Zumkley, T., Guo, J.Q., Tsai, A.P., Nakajima, H., “Diffusion in Quasicrystalline Al-Ni-Co and Al-Pd-Mn”, Mater. Sci. Eng. A, 294-296, 702-705 (2000) (Crys. Structure, Experimental, Phys. Prop., 18) Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd [2001Can]
[2001Goe] [2001Nau]
[2002Ant]
[2002Bil]
[2002Dem]
[2002Duq] [2002Hir]
[2002Kab]
[2002Klu]
[2002Lei]
[2002Miz]
[2002Mot]
[2002Pap]
[2002Sch]
[2002Shr] [2002Tak]
Landolt-Börnstein New Series IV/11A3
231
Canfield, P.C., Fisher, R., “Hihg-Temperature Solution Growth of Intermetallic Single Crystal and Quasicrystal”, J. Cryst. Growth, 225, 155-161 (2001) (Crys. Structure, Experimental, Magn. Prop., 11) Gödecke, T., “Ableitung des Kristallizationpfades in Ternaeren Gusslegierungen” (in German), Z. Metallkd., 92(8), 966-978 (2001) (Equi. Diagram, Experimental, *, 37) Naumovic, D., Aebi, P., Schlapbach, L., Beeli, C., Kunze, K., Lograsso, T.A., Delaney, D.W., “Formation of Stable Decagonal Quasicrystalline Al-Pd-Mn Surface Layer”, Phys. Rev. Lett., 87(19), 195506-1-195506-4 (2001) (Crys. Structure, Experimental, 35) Anton, R., Kreutzer, P., “Growth and Electrical and Optical Properties of Al(PdMn) Alloy Thin Films Prodused by Simultaneous Vapor Deposition of the Components”, J. Alloys Compd., 342(1-2), 464-468 (2002) (Crys. Structure, Equi. Diagram, Experimental 10) Bilusic, A., Bodrovic, Z., Smontara, A., Dolinsek, J., Canfield, P.C., Fisher, I.R., “Transport Properties of Icosahedral Quasicrystal Al72Pd19.5 Mn8.5”, J. Alloys Compd., 342(1-2), 413-415 (2002) (Electr. Prop., Experimental, Phys. Prop., Thermal Conduct., 23) Demange, V., Milandri, A., Weerd, M.C., Machizaud, F., Jeandel, G., Dubois, J.M., “Optical Conductivity of Al-Cr-Fe Approximant Compounds”, Phys. Rev. B, 65, 144205-1 -144205-11 (2002) (Calculation, Crys. Structure, Experimental, Optical Prop., 39) Duquesne, J.-Y., Perrin, B., “Elastic Wave Interaction in Icosahedral AlPdMn”, Physica B, 316-317, 317-320 (2002) (Experimental, Mechan. Prop., 9) Hiraga, K., “The Structure of Quasicrystals Studied by Atomic-Scale Observations of Transmission Electron Microscopy”, Adv. Imag. Electr. Phys., 122, 1-86 (2002) (Assessment, Crys. Structure, 99) Kabutoya, E., Edagawa, K., Tamura, R., Takeuchi, S., Guo, J.Q., Tsai, A.-P., “Plastic Deformation of Icosahedral Al-Pd-Mn Single Qusicrystals to large Strains I. Experiments”, Philos. Mag. A, 82(2), 369-377 (2002) (Experimental, Mechan. Prop., 15) Kluge, F., Yurechko, M., Urban, K., Ebert, Ph., “Influence of Growth Kinetics and Chemical Composition on the Shape of Voids in Quasi-Crystal”, Surf. Sci., 519, 33-39 (2002) (Crys. Structure, Experimental, 15) Lei, J., Wang, R., Yin, J., Duan, X., “Diffuse Electron Scattering Determination of Elastic Constants of Al-Pd-Mn Icosahedral Quasicrystal”, J. Alloys Compd., 342(1-2), 326-329 (2002) (Experimental, Mechan. Prop., 8) Mizutani, T., Nakano, H., Kashimoto, S., Takatani, Y., Mori, M., Ishimasa, T., Matsuo, S., “Ten-Fold-Like Magnetic Anisotropy in Electrical Conductivity of Al-Pd-Mn Icosahedral Quasicrystal”, J. Alloys Compd., 342(1-2), 360-364 (2002) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop., 7) Motomura, S., Ishimasa, T., Hirai, I., Kashimoto, S., Nakano, H., Matsuo, S., “Magnetic Properties of F2 M-Type Al-Pd-Mn Quasicrystals”, J. Alloys Compd., 342(1-2), 393-396 (2002) (Crys. Structure, Experimental, Magn. Prop., 14) Papadopolos, Z., Kasner, G., Ledieu, J., Cox, E. J., Richardson, N.V., Chen, E. J., Diehl, R.D., Lograsso, T. A., Ross, A. R., McGrath, R., “Bulk Termination of the Quasicrystalline Fivefold Surface of Al70 Pd21Mn9”, Phys. Rev. B, 66(18), 184207-1-184207-13 (2002) (Crys. Structure, Experimental, 47) Scheffer, M., Suck, J-B., “Inelastic Neutron Scattreing Study of the Dynamics of Al74Pd22Mn4 (`)”, J. Alloys Compd., 342, 310-313 (2002) (Calculation, Experimental, Thermodyn., 13) Shramchenko, N., Denoyer, F., “The Al-Pd-Mn Quasicrystalline Approximant (3)`-Phase Revisited”, Eur. Phys. J. B, 29(1), 51-59 (2002) (Crys. Structure, Experimental, *, 27) Takeuchi, S., Tamura, R., Kabutoya, E., Edagawa, K., “Plastic Deformation of Icosahedral Al-Pd-Mn Single Quasicrystals to Large Strains II. Deformation Mechanism”, Philos. Mag. A, 82(2), 379-385 (2002) (Experimental, Mechan. Prop., 9)
MSIT ®
Al–Mn–Pd
232 [2002Tex]
[2002Yam]
[2002Yan]
[2002Yur]
[2002Zha1]
[2002Zha2]
[2003Ban]
[2003Bal]
[2003Cap]
[2003Ebe]
[2003Pis]
Texier, M., Proult, A., Bonneville, J., Rabier, J., Baluc, N., Cordiers, P., “Microstructure of Icosahedral Al-Pd-Mn Quasicrystals Deformed at Room Temperature in an Anisotropic Confining Medium”, Philos. Mag. Lett., 82(12), 659-669 (2002) (Crys. Structure, Experimental, Mechan. Prop., 28) Yamamoto, A., Takakura, H., Tsai, A.P., “Structure Refinement of i-Al-Pd-Mn Quasicrystals by IP-Weissenberg Camera Data”, J. Alloys Compd., 342, 159-163 (2002) (Crys. Structure, Experimental, 11) Yang, W., Feuerbacher, M., Urban, K., “Cluster Structure and Low-Energy Planes in Icosahedral Al-Pd-Mn Quasicrystals”, J. Alloys Compd., 342(1-2), 164-168 (2002) (Crys. Structure, Experimental, 13) Yurechko, M., “Phase Equilibria in Ternary Systems on the Aluminum Basis, which Contain Quasiperiodic and Related Periodic Phases”, PhD Thesis, The Taras Shevchenko Kiev National University, Kiev (2002) (Equi. Diagram, Experimental, #, *, 160) Zhang, Y., Colella, R., Kycia, S., Goldman, A.I., “Absolute Structure-Factor Measurements of an Al-Pd-Mn Quasicrystal”, Acta Crystallogr., Sect. A: Found. Crystallogr., 58, 385-390 (2002) (Crys. Structure, Experimental, 18) Zhang, Y., Ehrlich, S.N., Colella, R., Kopecky, M., Widom, M., “X-Ray Diffuse Scattering in the Icosahedral Quasicrystal Al-Pd-Mn”, Phys. Rev. B, 66, 104202-1-104202-7 (2002) (Crys. Structure, Experimental, Theory, 18) Banerjee, G.N., Banerjee, S., Goswami, R., “Point Contact Spectroscopy of Al70 Pd30-xMnx Quasicrystals”, J. Phys.: Condens. Matter, 15(14), 2317-26 (2003) (Crys. Structure, Experimental, Electr. Prop., 22) Balanetskyy, S., Grushko, B., “Al-Pd (Aluminium - Palladium)”, MSIT Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, to be published, (2003) (Equi. Diagram, Crys. Structure, Assessment, 26) Capitan, M.J., Alvarez, J., Prejean, J.J., Berger, C., “Conductivity and Superlattice Ordering in an Icosahedral Al-Pd-Mn Phase”, Phys. Rev. B, 68(6), 064203-1-9 (2003), (Crys. Structure, Experimental, Electr. Prop., 32) Ebert, Ph., Yurechko, M., Kluge, F., Cai, T., Grushko, B., Thiel, P.A., Urban K., “Surface Structure of Al-Pd-Mn Quasicrystals: Existence of Supersaturated Bulk Vacancy Concentrations”, Phys. Rev. B: Condens. Matter, 67(2), 24208-1-8 (2003) (Crys. Structure, Experimental, Phys. Prop., 35) Pisch, A., “Al-Mn (Aluminium-Manganese)”, MSIT Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, to be published, (2003) (Equi. Diagram, Crys. Structure, Assessment, 40)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 ( Mn) 1246 - 1140
MSIT®
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cI2 Im3m W
Lattice Parameters Comments/References [pm] a = 404.96
a = 308.0
at 25°C [Mas2] dissolves ~0.2 at.% Pd [2003Bal] and ~0.62 at.% Mn [1997Oka] [Mas2] dissolves ~39.5 at.% Al [1997Oka] and ~4 at.% Pd [Mas2]
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd Phase/ Temperature Range [°C] ( Mn)
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype cF4 a = 386.0 Fm3m Cu
(Mn) 1100 - 727
cP20 P4132 Mn cI58 I43m Mn cF4 Fm3m Cu
(Mn) < 727 (Pd) < 1555
, (MnPdAl) 1400 - 840 , MnAl 1177 - 840 1, MnxAl1-x 1048 - 957 2, Mn5Al8 < 991 , MnxAl1-x 1275 - 870
Landolt-Börnstein New Series IV/11A3
cI2 Im3m W
a = 631.52
a = 891.26
a = 389.0
a = 308.3 (?) a = 306.3 0.3
233 Comments/References
[Mas2] dissolves ~9.1 at.% Al at 1073°C [1997Oka] and ~22.6 at.% Pd at 1147°C [Mas2] [Mas2] dissolves ~40.5 at.% Al [1997Oka] and ~3 at.% Pd [Mas2] at 25°C [Mas2] dissolves ~2 at.% Al [1997Oka] and ~2 at.% Pd [Mas2] [Mas2] dissolves ~20 at.% Al at 1055°C [2003Pis] and ~30.5 at.% Mn at 1350°C [Mas2] MnPd2Al at >1010°C [1981Sol2] 34.5 to 51.3 at.% Mn [1997Oka] Mn45 Al55 at 957°C [1990Ell] 30 to 38.2 at.% Mn [1997Oka]
hR26 R3m Cr5Al8 hP2 P63/mmc Mg
a = 1273.9 c = 1586.1 a = 270.5 - 270.5 c = 436.1 - 438
31.4 to 50.01 at.% Mn [2003Pis] at 42 at.% Mn [V-C2] 53.2 to 60 at.% Mn [1997Oka] usually called J
MSIT ®
234
Al–Mn–Pd
Pearson Symbol/ Lattice Parameters Comments/References Phase/ Temperature Range Space Group/ [pm] Prototype [°C] Labeled as “H” by [1995Goe1, H, (MnPdAl)4(MnPdAl)11 1995Goe2, 1999Gru, 2002Yur], 6 to 7 at.% Pd, 71 to 73 at.% Al at 880-870°C 7 at.% Pd, 70.6 to 71.6 at.% Al at 850°C ~6.5 at.% Pd, 73.5 to 76.8 at.% Al at 710°C [2002Yur] Labeled as “T” by [1993Aud, 1994Aud, 1997Kle, 1997Mat, 2000Kle1, 2002Tex] Mn4Al11 (h) 3.5 to 6.6 at.% Pd, 16 to 25.5 at.% Mn 1002 - 895 oP160 [2000Kle1] Pnma 25 to 28.7 at.% Mn [1997Oka] MnAl3 MnAl3 [1992Li] a = 1479 b = 1242 c = 1259 a = 1483 MnAl3, as-cast [1993Hir2] b = 1243 c = 1251 a = 1483.7 0.4 MnAl3 (Pn21a?) [1994Shi] b = 1245.7 0.2 c = 1250.5 0.2 a = 1471.7 Mn24.5 Pd3.2Al72.3 single crystal b = 1251.0 obtained by Bridgman technique c = 1259.4 [1997Kle]. a = 1476 Mn23 Pd6Al71, obtained by Bridgman or b = 1243 Czochralski technique [1993Aud, c = 1256 1994Aud] a = 1417 in Mn11.85Pd21.88Al66.27 rapid b = 1251 quenched sample, together with -3 and c = 1259.4 (Al-Pd) [2002Tex] a = 1472.7 0.3 Mn20.9 Pd4.4Al74.7 single crystal, b = 1250.9 0.3 as-cast [1997Mat] c = 1260.0 0.3 25 to 28.7 at.% Mn [1997Oka] Mn4Al11 (l) aP30 [V-C2] < 916 P1 a = 509.5 0.4 b = 887.9 0.8 Mn4Al11 c = 505.1 0.4 = 89.35 0.04° = 100.47 0.05° = 105.08 0.06° , MnAl4 a = 1998 19 to 20.8 at.% Mn [1997Oka] hP574 c = 2467.3 < 923 P63/mmc MnAl4
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd Phase/ Temperature Range [°C] , MnAl 4 < 693
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype hP586 P63/m a = 2838.2 c = 1238.9
%, (Mn,Pd,Al)1(Mn,Pd,Al)4
orthorhombic Cmcm Mn11Ni4Al60
%, MnAl4
MnAl12 < 500 1, Mn3Al10 < 860 MnxAl1-x
i-MnAl d-MnAl PdAl4 < 604 Pd8Al21 640
Landolt-Börnstein New Series IV/11A3
oC28 Cmcm MnAl6
cI26 Im3 WAl12 hP26 P63/mmc Co2Al5 tP2 P4/mmm CuAu icosahedral m35 decagonal D3 P6322 PtAl4 tI116 I41/a Pt8Al21
Comments/References
16.8 to19 at.% Mn [Mas2] [2003Pis] space group does not fit 100%, probably P63 Phase of composition 3.5-6.6 at.% Pd, 16-25.5 at.% Mn [2000Kle1], labeled as “R” by [1993Aud, 1994Aud, 1998Ber, 2000Kle1]
a = 2360 b = 1240 c = 770
in rapidly solidified MnAl4 alloy after heating to 600°C [1992Li], metastable
a = 2388 b = 1243 c = 776
Mn15.6 Pd5.7Al78.6 (EPMA), in samples obtained by induction melting and Bridgman technique [1993Aud, 1994Aud] Mn15.7 Pd5.7Al78.6 (EPMA) [1998Ber]
a = 2388 b = 1243 c = 778 MnAl6 < 705
235
a = 755.51 b = 649.94 c = 887.24 a = 754.5 0.2 b = 649.0 0.3 c = 868.1 0.2
14.2 at.% Mn [1997Oka] [V-C2]
[2000Yam]
a = 747
7.7 at.% Mn [1997Oka], [V-C2]
a = 754.6 0.3 c = 289.5 0.2
[2003Pis], metastable
a = 278 - 279 c = 356 - 357
55.8 to 55.1 at.% Al, metastable [2003Pis]
a = 1308.6 c = 963.1
~20 at.% Mn [2003Pis], quasicrystal, metastable ~22 at.% Mn, [2003Pis] quasicrystal, metastable 20 at.% Pd [2003Bal], usually labeled as “”
a = 1299.8 c = 1072.9
27.5 at.% Pd [2003Bal], usually labeled as “ ”
a = 1240.0
MSIT ®
236 Phase/ Temperature Range [°C] “J-phase”
J6 782 - 579
J28 792 - ?
J22
J34 Jx
J6*
MSIT®
Al–Mn–Pd Pearson Symbol/ Lattice Parameters Comments/References [pm] Space Group/ Prototype Overall ternary extension of the binary J phases, 0 to 5 at.% Mn at 710°C [2002Yur] ~25.4 to 26.9 at.% Pd [2003Bal] orthorhombic a = 2350.0 Pnma b = 1680.0 c = 1230.0 phase labeled as !´ by [1996Bou, 2000Kle1, 2002Shr], composition Mn5.0-4.6Pd22.1-22.4Al73.3 at 730°C [2000Kle1] Mn4.1Pd22.4Al73.5, single-crystal Pnma a = 2354.1 obtained by Bridgman method, XRD b = 1656.6 [1996Bou] c = 1233.9 Pnma or P21ma a = 2354.3 0.5 Mn4.5Pd22.9Al72.6 sample obtained by b = 1664.3 0.5 Bridgman method [2002Shr] c = 1237.4 0.4 a = 2350.0 orthorhombic 28.1 to < 26.9 at.% Pd [2003Bal] b = 1680.0 B2mm c = 5700.0 a = 2357.4 orthorhombic Mn4.5Pd22.9Al72.6 sample obtained by b = 1661.0 Bridgman method, labeled as !´_3 c = 5712.0 [2002Shr] a = 2350 orthorhombic 3.1 to 4.6 at.% Mn at ~750-~800°C b = 1680 [2002Yur], c = 4490 according to [2002Shr] could be named !´_2 a = 2350 orthorhombic 1.2 to 3.1 at.% Mn at 710-790°C b = 1680 B2221 ? [2002Yur], according to [2002Shr] c = 7010 could be named !´_4 a = 2353.1 Mn4.5Pd22.9Al72.6, labeled as !´_5, c = b = 1658.0 (c!´ (5+-)) [2002Shr] c = 8183.2 a = 2354.7 0.5 single crystal? [2002Shr] orthorhombic b = 1670 3 c = 8261 37 a = (4700.0) ordered J6 [2003Bal] b = 3360.0 c = 2460.0
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd Phase/ Temperature Range [°C] !
Pearson Symbol/ Space Group/ Prototype orthorhombic Cmcm
Bmmb
, (Mn,Pd,Al)2(Mn,Pd,Al)3
, (MnPdAl)
, PdAl (h) 1645 - 545 , MnPd 1515 - 540 ', PdAl (l) < 850 < 740 , Pd5Al3 1315 - 615 ', Pd2Al < 1418
Pd5Al2 < 980 Pd3Al < 775
Landolt-Börnstein New Series IV/11A3
Lattice Parameters Comments/References [pm] a = 2032 b = 1650 c = 1476 a = 2032 b = 1657 c = 1475 a = -#c!´ b = b!´ c = --1#a!´ a = 2032 b = 1650 c = 1476
hP5 P3m1 Ni2Al3
, Pd2Al3 < 952
237
Mn4Pd23Al73, Mn 4Pd21.6 Al74 (EMPA) [1993Aud] Mn4Pd23Al73 [1998Ber]
[2000Kle1]
in as-cast Mn5Pd20Al75 alloy [1996Sun] 0 to 2 at.% Mn at 850°C, 0 to 1.6 at.% Mn at 710°C [2002Yur]
a = 422.7 c = 516.7 cP8 Pm3m CsCl a = 308.3 a = 303.6 a = 273.3
38 to ~42.2 at.% Pd [2003Bal] at 40 at.% Pd [2003Bal] ~5 to 12 at.% Mn, ~57 at.% Al at 880-870°C in equilibrium with -3 [2002Yur] MnPd2Al, at 1010-710°C [1968Web] 43.5 to ~56 at.% Pd [2003Bal] 34 to ~62 at.% Mn [Mas2] HT [V-C2] ~48.5 to ~52.2 at.% Pd [2003Bal]
hR78 R3
a = 1565.9 c = 525.1
CP8 P213 FeSi oP16 Pbam Rh5Ge3 oP12 Pnma Co2Si
a = 486.2
~48 to ~49 at.% Pd [2003Bal]
a = 1047.1 b = 537.3 c = 503.5
62.5 at.% Pd [2003Bal]
oP28 Pnma Pd5Ga2 orthorhombic P21ma -
a = 541.0 b = 405.5 c = 776.0 a = 540.0 b = 403.4 c = 1840.5 a = 540.7 b = 403.2 c = 1580.2
~65 to ~76 at.% Pd [2003Bal] at 66.1 at.% Pd [2003Bal]
~70.7 to ~71.7 at.% Pd [2003Bal] Usually labeled as “-” 75 at.% Pd [2003Bal], usually labeled as “1”
MSIT ®
Al–Mn–Pd
238 Phase/ Temperature Range [°C] Pd5Al 660 1, MnPd 1200 MnPd3 < 750 Mn3Pd5 < 500
Mn11Pd21 < 197 * -1, MnPd2Al < 710 * -2 (l)
Pearson Symbol/ Space Group/ Prototype Pnma tP4 P4/mmm AuCu tI16 I4mm Au3Cd oC16 Cmmm Ga3Pt5 tP32 P4/mmm Mn11Pd21 cF16 MnCu2Al decagonal D3
< 864 decagonal P105/mmc
* -2 (h) 896 - 864
MSIT®
B-centered orthorhombic
Lattice Parameters Comments/References [pm] a = 1070.0 b = 400.0 c = 807.4 a = 406.9 c = 358.5 a = 391.3 b = 1549.6 a = 807.2 b = 727.9 c = 404.4 a = 806.1 c = 733.0 a = 618.2
83 at.% Pd [2003Bal], sometimes called “%” 31 to ~53 at.% Mn [Mas2] low T [V-C2] 32.5 to ~35.5 at.% Mn [Mas2] [V-C2] ~37.5 at.% Mn [Mas2] [V-C2]
~34.5 at.% Mn [V-C2] RT, Mn25.2Pd49.7Al25.1 [1981Sol2]
quasicrystal, usually labeled as “D” Mn17.5-17.9Pd13.5-12.1Al69-70 at 880-710°C [2002Yur] Mn17 Pd13Al70 annealed at 800°C 4d a = 1200 [1993Hir1] a5 = 1255.7 0.1 Mn16.5 Pd13Al70.5 single crystal V5D = 1030.1#10-10 pm5 [1993Ste] a = 1240 Mn16.5 Pd13Al70.5 (SEM-XMA) in sample Mn2Pd2Al9 annealed at 780°C 2.5d and quenched in water [1991Bee] a = 1250 very close to Mn18.1Pd12.1Al69.8, samples annealed at 855 (5d), 830 (14d), 600°C (60d) [1995Bee1, 1995Bee2] very close to Mn18.1Pd12.1Al69.8, labeled as DH by [1995Bee1, 1995Bee2, 1995Goe2] a = 2030 in samples annealed above 865°C b = 1250 [1995Bee1, 1995Bee2, 1995Goe2] c = 6250
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd Phase/ Temperature Range [°C] * -3 < 893
Pearson Symbol/ Lattice Parameters Comments/References [pm] Space Group/ Prototype quasicrystal, usually labeled as “I” icosahedral Mn6-10Pd24.5-19.2Al69.5-70.8 [2002Yur] Mn8-10.2Pd20.3-23.2Al68-69.5 [2000Kle1] 6D face-centered a6D = 645.990.03 hypercubic lattice Pm35 a6D = 645.1
aF = 1290.1
aF = 1290.1
superstructure of the F-phase
P-type
< 715 [2000Let]
Landolt-Börnstein New Series IV/11A3
239
(a6D)F2 = 2-a6D aP = 2091.1 aP = 2088.1 aP = -aF
“diamond” type aF2 = 2581.0 aF2 = 2aF cubic symmetry aF = 1292.6 approximately m3, superstructure of the F2-phase aP = 646.3 0.5
sometims labeled as “F” Mn10 Pd19Al71, spinning and subsequent annealing, neutron diffraction [1991Bou] Mn9.6Pd21.7Al68.7 sample obtained by Bridgman method (EPMA, XRD). Density 5.1 0.2 g#cm-3. Mn9Pd21Al70, obtained by plasma jet melting and annealed at 800°C 3d subsequently quenched in liquid nitrogen. Neutron diffraction [1992Bou]; Mn8.5Pd21Al70.5 single grains, grown using conventional casting procedures, X-ray and neutron diffraction [1993Bou] in Mn8Pd21Al71 sample annealed at different temperatures (XRD) [2000Hir] in Mn7.5Pd21.5Al71 sample obtained by plasma jet melting and annealed at 803 4°C 50h, quenched into water and subsequently into liquid nitrogen (XRD) [1995Ish] sometimes labeled as F2, not stable, corresponds to a transient state in the process of the transformation F to F2M [2000Let] [2000Hir] [2003Cap] Mn7.5Pd21.5Al71 alloy obtained by plasma jet melting and annealed at 6022°C 48-400 h, quenched into water and subsequently into liquid nitrogen. XRD [1995Ish] Mn8.8Pd21.4Al69.8 (EPMA) at RT [2000Let] sometims labeled as F2M
at RT, Mn8.7Pd22.0Al69.3 (EPMA, XRD) [1998Boi] MSIT ®
Al–Mn–Pd
240 Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Al
Mn
Pd
L + 1 H +
952
U1
L
66
12
22
1 H + + 2
923
E1
1
-
-
-
L+ H + -2
896
P1
L
70.5
19
10.5
L + -2 + -3
893
P2
L
71
19.5
9.5
L + -2 H + -3
887
U2
L
72
18.5
9.5
L + -3 +
~875
U3
L
73.5
21
5.5
-3 + -2 +
860
U4
-
-
-
-
L + + -3 J
845
P3
L
76
20.5
3.5
L + -3 H + J
832
U5
L
73
15.5
6.5
H + - 2 + 2
755
U6
-
-
-
-
L + H + MnAl6
647
U7
L
96
2
2
L + MnAl6 (Al) + H
626
U8
L
95
4
1
L + H + (Al)
618
U9
L
92.5
7
0.5
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
Al-Mn
Al-Mn-Pd
1048 p1 l + γ γ1 1002
241
Al-Pd
γβ
p2
l+γ1Mn4Al11(h)
991 e1 γ1 + γ γ2 957
e2
γ1Mn4Al11(h)+γ
923
γβ L + γ1 β + H
952
γβ
U1
952 p3 l+βδ
β+γ1+Η
p4
923
γ1 β + γ2 + H
L+H+β
E1
l+Mn4Al11(h)µ
γ+β2+H
896
L + H + β τ2
τ2+H+β
P1
L + τ2 + β τ3
893
L+H+τ2
L+β+τ2
L+β+τ3
L+τ2+τ3
L + τ2 H + τ3
887
U2
L + β τ3 + δ
~875
τ2+β+τ3
τ2+H+τ3
τ3 + β τ2 + δ L+H+τ3 β+τ2+δ τ3+τ2+δ 845 L+τ3+ε
L + τ3 H + ε
U5
L+τ2+δ
U4
L + δ + τ3 ε δ+τ3+ε
755
658.5 e3 l (Al)+MnAl6
H + β τ 2 + γ2 δ+τ3+γ
P3
L+δ+ε 792 p5 l + δ ε28
τ3+H+ε
705 p7 l + µ MnAl6
U3
β+τ3+δ
860
832
P2
782 p6 l + ε28 ε6
U6
H+τ2+γ2 L+H+ε
647
L + µ H + MnAl6 U7
L+H+MnAl6 626
µ+H+MnAl6
L+MnAl6 (Al)+H U8
MnAl6+(Al)+H
L+(Al)+H 618
L + H (Al) + ε
U9
H+ε+(Al) L+ε+(Al)
616 e4 l (Al) + ε6
Fig. 1a: Al-Mn-Pd. Reaction scheme in the Al-rich part of the system
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Mn–Pd
242
Al-Mn
Al-Mn-Pd
Al-Pd
γβ
1048 p1 l + γ γ1 1002 p2 l+γ1Mn4Al11(h) 991 e γ1 + γ γ2
γβ
957 e γ1Mn4Al11(h)+γ
γβ L + γ1 H + β
952
952 p3 l+βδ
U1
β+γ1+H γ1β + γ2 + H
923
U3
L+H+β
E1
β+γ2+H
894,876
U6
L+H+τ3
L + H + β τ3
L+β+τ3
τ3+H+β
U3
U5 H + τ3 + β τ2
770
H+β+τ2
H+τ2+τ3
P1
P2 τ2+β+τ3
U6
Fig. 1b: Al-Mn-Pd. Metastable reaction scheme in the Al-Mn-Pd system according to [1995Geo1]
Al
Data / Grid: at.%
e3,658.5°C MnAl6 p7,705°C
Axes: at.%
U7
µ
10
(αAl) U8 e4,616°C U9
0 75
Fig. 2: Al-Mn-Pd. Partial liquidus surface projection: a) - stable, b) - metastable
P1
p2,1002°C
00 10
τ3
H U2
γ1
p1,1048°C
τ2
0 105
U1
1025
00 11
p,1177°C
δ
p6,782°C p5,792°C p3,952°C
P1
30
40
U3
P3
950 0 100
70
P2
00 11 0 0 12 0 0 13
γ (MnPdAl)
δ
b)
80
U5
P3
p5
U3
ε
5 87 0 90 5 92 0 95
20
τ2
90
0 80 0 85
p4,923°C
p6
U5
60
β ((MnPdAl)
MnxAl1-x Mn 50.00 Pd 0.00 Al 50.00
MSIT®
10
20
30
a)
40
Mn 0.00 Pd 50.00 Al 50.00 Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
243
Al Fig. 3: Al-Mn-Pd. Solidus surface projection
Data / Grid: at.%
(α Al)
70 5
Axes: at.%
10
90
626 MnAl6
(α Al)+H+ε
20
80
92 3
µ
6 61
618°C
647
H
1002 30
γ1
1048
887
H+γ 1
ε
832
89 τ 2+τ 3+β 6 893
τ3
70
84 5
2 79
H+γ 1+β 952
τ2
H+τ 3
β +τ 3
δ
875
40
60
2 95
β +δ
γ (MnPdAl) β (MnPdAl) 10
Mn 50.00 Pd 0.00 Al 50.00
20
30
40
Al
Mn 0.00 Pd 50.00 Al 50.00
Data / Grid: at.%
Fig. 4: Al-Mn-Pd. Isothermal section in the Al-rich part of the system at 894°C
Axes: at.%
10
90
L
µ
20
80
L+H Mn4Al11(l) 30
H
τ2 γ 2+β +H
γ2
70
L+δ
τ 2+β +L L+β
H+β
δ
40
60
γ 2+β β (MnPdAl) Mn 50.00 Pd 0.00 Al 50.00 Landolt-Börnstein New Series IV/11A3
10
γ (MnPdAl)
20
30
40
Mn 0.00 Pd 50.00 Al 50.00
MSIT ®
Al–Mn–Pd
244
Al
Data / Grid: at.%
Fig. 5: Al-Mn-Pd. Isothermal section in the Al-rich part of the system at 875°C
Axes: at.%
10
90
L
µ
20
80
H+L
Mn4Al11(l) H
30
τ2 γ 2+H+β 40
τ 2+τ 3+β
H+β
70
L+δ
τ3 τ 3+β
γ2
δ 60
β +δ
γ 2+β β (MnPdAl) 10
Mn 50.00 Pd 0.00 Al 50.00
γ (MnPdAl)
20
30
40
Al
Mn 0.00 Pd 50.00 Al 50.00
Data / Grid: at.%
Fig. 6: Al-Mn-Pd. Isothermal section in the Al-rich part of the system at 840°C
Axes: at.%
10
90
L
µ
20
H+L 80
Mn4Al11(l)
L+ε
L+H+τ 3
H
H+τ 3
30
τ2
ε τ 2+τ 3+δ
70
τ3
γ 2+H+β 40
H+β
γ2
δ
τ 2+β
60
δ +β γ 2+β β (MnPdAl) Mn 50.00 Pd 0.00 Al 50.00
MSIT®
10
20
30
40
Mn 0.00 Pd 50.00 Al 50.00 Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
245
Al Fig. 7: Al-Mg-Pd. Isothermal section in the Al-rich part of the system at 710°C
Data / Grid: at.% Axes: at.%
L 10
90
µ +L+H µ
20
80
L+H+ε
Mn4Al11(l)
H
H+τ 3
30
γ 2+τ 2+H
γ 2+H
40
L+ε
τ3
τ2
ε
ε6 ε28 ε+δ 70
τ 2+δ+τ 3
γ 2+τ 2+β
γ2
δ
60
τ 2+β
β +δ
γ 2+β 10
Mn 50.00 Pd 0.00 Al 50.00
β (MnPdAl) 20
30
40
Al
Data / Grid: at.%
(αAl)
Fig. 8: Al-Mn-Pd. Isothermal section in the Al-rich part of the system at 600°C
10
(α Al)+H+ε
MnAl6
µ
Axes: at.%
90
PdAl4
20
80
Mn4Al11(l) 30
Mn 0.00 Pd 50.00 Al 50.00
ε H γ 2+H+τ 2
70
τ3
τ2
Pd8Al21
δ
40
γ2
60
τ 2+γ 2+β
τ 2+β
γ 2+β
50
µ
50
β´
(β Mn) Mn 60.00 Pd 0.00 Al 40.00 Landolt-Börnstein New Series IV/11A3
(β Mn)+β 10
β (MnPdAl) 20
30
40
50
Mn 0.00 Pd 60.00 Al 40.00
MSIT ®
Al–Mn–Pd
246
1100
Fig. 9: Al-Mn-Pd. Partial vertical section at 10 at.% Pd
L+β L+γ1
1000
Temperature, °C
L
L+γ1+β 952
L+γ1+H L+H+β 887
900
864
832 H+τ 3
L+H
923 H+β +γ1 β +γ1+γ2
896
L+H+τ3
800
H+β H+β +γ2
~755
H+ε H+τ3+ε
τ2+β +γ2
700
(Al)+ε+L 600
(Al)+PdAl4+ε
τ2+γ2
L+H+ε H+τ2+γ2
618
(Al)+ε
β +γ2
H+τ2 H+τ2+β
ε+L
β
γ1+β
(Al)+H+ε 10
Mn 0.00 Pd 10.00 Al 90.00
20
30
Mn 32.00 Pd 10.00 Al 58.00
Mn, at.%
1000
Fig. 10: Al-Mn-Pd. Partial vertical section at 20 at.% Pd
L 900
L+ε+τ3
L+H+β
τ3+β L+β +τ3
L+H+τ3
Temperature, °C
γ1+β
952
τ3
L+τ3
L+β
τ2+τ3+β
896 893
860
β
923 864 H+β
832 800
H+τ3
L+ε
τ2+τ3 τ2+H+β
τ2+δ
H+γ2+β ~755
β +γ2
L+H+ε 700
τ2+β τ2+β +γ2
(Al)+ε+L (Al)+ε
(Al)+ε+H H+τ3+ε 618 H+ε
τ2+δ +β
600
(Al)+PdAl4+ε
Mn 0.00 Pd 20.00 Al 80.00
MSIT®
τ2+τ3+H 10 τ2+τ3+δ
Mn, at.%
20
Mn 28.00 Pd 20.00 Al 52.00
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
Fig. 11: Al-Mn-Pd. Partial vertical section at 70 at.% Al
1002°C 1000
975°C
Temperature, °C
L+β
1048°C
900
247
L+γ1+β
L+γ1
γ1+H γ2+H
L L+γ1+H 952 923
895°C
L+τ2
L+H
L+τ3+β
896 864
H+β
L+τ2+τ3
887
L+τ3
H+τ3
H+τ2+β
τ3
800
~875 τ3+β 845
H+τ2+τ3
~755
792°C
γ2+Mn4Al11(l)
δ +ε
τ2+τ3
τ2+H
640°C Pd8Al21+δ
τ2+τ3+δ
γ2+Mn4Al11(l)+H
L+δ +ε
τ3+δ +ε
H+τ2+γ2
600
L+δ L+τ3+δ
τ3+δ
H+β +γ2
700
952°C
L+β
893 L+τ2+β
Pd8Al21+δ +ε 10
Mn 30.00 Pd 0.00 Al 70.00
20
Mn 0.00 Pd 30.00 Al 70.00
Pd, at.%
1000
Fig. 12: Al-Mn-Pd. Partial vertical section at 6 at.% Mn
L+β L L+τ3+β
Temperature, °C
900
L+τ3+δ L+τ3 L+τ3+ε
L+β +δ
τ3+β +δ 845
800
860
τ3+ε
L+H L+H+µ
705°C
β
τ3+δ +ε
832
L+µ
~875
L+H+ε
τ2+β +δ
τ3+δ
700
L+MnAl6 L+µ+MnAl6 658.8°C
626 H+(Al)+MnAl6 600
(Al)+MnAl6
β +δ
τ2+τ3+δ
L+(Al)+H 618
τ2+δ
(Al)+H+ε
Mn 6.00 (Al)+H Pd 0.00 Al 94.00
Landolt-Börnstein New Series IV/11A3
H+τ3+ε L+H+MnAl6 H+ε
647
10
20
Pd, at.%
30
Mn 6.00 Pd 40.00 Al 54.00
MSIT ®
Al–Mn–Pd
248
920
Fig. 13: Al-Mn-Pd. Partial vertical section from Mn32.8Al67.2 to Pd27Al73
Temperature, °C
L+H+β
L L+H
H+β
900
P1,896 H+β +γ2
H+τ2+L
τ2+L
P2,893
H+γ2
U2,887
γ2
L+δ
τ2+τ3+L L+β
τ3+β +L
L+β +δ L+β +τ3
880
τ2+τ3+β
H+β +τ2
U3,~875
τ3+L
τ2(h?)
L+τ3+δ
τ3+β 864
τ2
860 10
Mn 32.80 Pd 0.00 Al 67.20
Mn 0.00 Pd 27.00 Al 73.00
Pd, at.%
1000
Fig. 14: Al-Mn-Pd. Partial vertical section from Mn31.2Al68.8 to Pd29Al71
20
L+γ1
γ1
L L+γ1+H
957°C
γ1+H
U1,952
Temperature, °C
L+H
γ1+γ2+H β +γ1+γ2 900
895°C
E1,923
L+β
L+τ2+β
H+β +γ2 H+γ2
L+H+β
H+β
P1,896 P2,893
L+τ2+τ3
L+τ3+δ
τ2
Mn4Al11(l)+H+γ2
H+τ2+β
L+δ +β U3,~875
U2,887
τ2(h?)
L+δ
L+τ3+β L+τ3
P3,845
τ2+τ3
τ3+δ +ε
τ3
Mn4Al11(l)+γ2
L+ε+δ
δ +ε
800
Mn 31.20 Pd 0.00 Al 68.80
MSIT®
10
20
Pd, at.%
τ3+δ
Mn 0.00 Pd 29.00 Al 71.00
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
249
1000
Fig. 15: Al-Mn-Pd. Partial vertical section from Pd3Al98 to Mn13Pd30Al57
L+β L
τ2+τ3+β τ3+β L+τ3+β
900
Temperature, °C
L+τ3 U5,832
τ3
H+τ3+ε
β
U4,860
τ3+δ
800
τ2+β τ2+δ L+H+ε H+τ3
L+H 700
τ2+τ3+δ
H+ε L+(Al) L+(Al)+H L+(Al)+ε
τ2+τ3 H+τ2+τ3
U9,618 (Al)+ε
600
Mn 0.00 Pd 3.00 Al 97.00
τ2+δ +β
(Al)+ε+H 10
20
Mn 13.00 Pd 30.00 Al 57.00
Pd, at.%
8
Fig. 16: Al-Mn-Pd. Low-temperature heat capacities plotted in the form of Cp/T vs T2 for the Mn10Pd20Al70 quasicrystal [1997Ina]
Cp/T, mJ/molK2
6
4
2
0
50
100 2
150
2
Temperature , K
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Mn–Pd
250
6
5
b)
4
Cp/KB
Fig. 17: Al-Mn-Pd. Temperature dependence of specific heat per atom at constant pressure (a) and at constant volume (b) measured for Mn9Pd20Al71 quasicrystal [2000Eda]
a)
3
2
[1997Ina] 1
0 0
200
400
600
800
1000
Temperature, K
101
2.0 10
10-1 1.0 10-2 0.5
(Cp-Cv)/Cv, %
1.5
Cp, J·g-atom-1·K-1
Fig. 18: Al-Mn-Pd. Low temperature specific heat Cp(T) of icosahedral Al68.2Mn 9Pd22.8 as a function of temperature between 1.6 and 280 K. The Cp(T) and Cv(T) due to thermal-expansion effects, as a function of temperature [1998Wae]
10-3 0
100
200
300
Temperature, K 10-4 1
10
100
Temperature, K
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Pd
251
10
Fig. 19: Al-Mn-Pd. Temperature dependence of the specific heat Cp of Mn4Pd22Al74, measured with a heat relaxation system [2002Sch] Cp, J·mol-1·K-1
1
0.1
0.01
0.001 1
10
100
Temperature, K
900
800
Hardness, HV/kg·mm2
Fig. 20: Al-Mn-Pd. Vickers hardness as a function of temperature for several symmetrical atomic surfaces in as-grown single icosahedral Mn10Pd20Al70 ingots with 3-fold growth direction by the Czochralski method [1999Yok]
600
400
2-fold 3-fold 5-fold 200
0 300
400
500
600
700
800
Temperature, K
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Mn–Pd
252
9
Fig. 21: Al-Mn-Pd. Temperature dependence of the quasilattice constant of the icosahedral phase [2000Kaj]: a = a(T)-a0, where a0 = a(0 K)
8
7
6
a/a0 (103)
5
4
3
2
1
0 0
100
200
300
400
500
600
700
Temperature, K
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Ti
253
Aluminium – Manganese – Titanium Andy Watson Literature Data A number of phase diagram studies have been made of the ternary system. [1954Dom, 1955Dom] investigated some 100 ternary alloys ranging from pure titanium to the 60 at.% Al and 40 at.% Mn binary limits. Metallographic and incipient melting techniques were employed. Samples prepared from iodide titanium, high purity aluminium and electrolytically-refined manganese were homogenized by repeated arc melting and annealing at 1000°C for 24 h. Phase equilibria were determined by the metallographic examination of specimens quenched after annealing under argon or in vacuum at temperatures between 700 and 1200°C. Corresponding annealing times ranged from 17 to 4 d; no X-ray examination or tests for equilibrium were reported. Six vertical sections and isothermal sections at 750, 800, 900, 1000, 1100 and 1200°C were constructed. The major variation in form occurs between 900 and 800°C, with a (Ti)+TiAl high temperature equilibrium being superseded at lower temperatures by an (Ti)+TiMn2 phase field as a result of a four-phase invariant reaction, (Ti)+TiAl(Ti)+TiMn2, which is reported to occur at 865°C. A second four-phase reaction (Ti)+TiMn2(Ti)+TiMn was inferred to occur between 700 and 550°C. The results of [1954Dom, 1955Dom] are summarized in [1974Zwi]. The isothermal sections reported by [1954Dom, 1955Dom] are not completely consistent with more recent binary data in that they fail to take into account the existence of the phases Ti3Al and TiMn and they assume that the Laves phase TiMn 2 is a purely binary phase of invariant composition. Electrical resistivity measurements and metallographic observations on samples annealed at 600, 800, 1000 and 1200°C were used by [1960Sat] to construct a partial phase diagram for alloys containing up to 49 at.% Al and 52 at.% Mn. Since this investigation assumes the intervention of a high-temperature phase Ti3Al2 and associated reactions, L+(Ti)Ti3Al2 (1620°C), L+Ti3Al2TiAl (1460°C) and (Ti)+Ti3Al2Ti3Al (1400 to 1300°C), their proposed ternary equilibria are generally incompatible with other studies. As in [1954Dom, 1955Dom], an invariant reaction was proposed at 860 10°C, but [1960Sat] suggests that it takes the form (Ti)+Ti 3Al(Ti)+TiMn. A limited study of ternary equilibria around the TiAl phase has been made by [1988Has] using optical microscopy, electron probe microanalysis and X-ray diffraction. Extension of the Al-rich boundary of the TiAl phase field to approximately 4.5Mn-46.7Al (at.%) at 1000°C was reported, but it should be noted that in a second publication of these results [1989Tsu] the authors emphasize that equilibrium as not achieved by the heat treatment used (7 d). The manganese-rich regions of the system (0 to 40 at.% Al, 0 to 40 at.% Ti) have been investigated by [1977Cha]. Samples were argon-arc melted from electrolytic or high purity grade materials. After annealing under argon at 1000°C for 7 d the alloys were quenched and the phases present were determined using optical metallography and X-ray powder diffraction techniques. Observations on the binary Al-Mn alloys conflict with generally accepted information since they indicate the (Mn) field as extending to only 39 at.% Al. The Laves phase was shown to exhibit an extensive ternary homogeneity range projecting along lines of approximately constant Ti content. Its limits were not completely established by [1977Cha] but X-ray studies by [1974Dwi] and [1978Jac] show that it extends at least to the equiatomic ternary composition. The presence of the hexagonal phase at the composition 37.5Ti-25Mn (at.%) has been reported by [1988Has]. Aluminium-rich alloys containing up to 1.2 at.% Ti and 2.3 at.% Mn have been investigated by [1958Mal]. Using differential thermal analysis, liquidus surfaces and the corresponding lines of the secondary reactions L+TiAl3(Al), L(Al)+MnAl6 and LTiAl3+MnAl6 were determined and an invariant four-phase reaction point, L+TiAl3(Al)+MnAl6 was located at 0.07Ti-0.94Mn (at.%) and 663°C. Solid state isothermal equilibria were investigated by means of metallographic and X-ray examination of alloys annealed at 650°C (121 h), 600°C (136 h), 500°C (228 h) and 400°C (168 h). The ternary sections given by [1954Dom] and [1955Dom] did not include the Al rich portion of the diagram and thus the ternary L12 phase was not seen. A number of studies have focussed on equilibria involving this phase. The extent of the L12 phase field was investigated by [1991Nic], who arc melted Al, Mn and Ti
Landolt-Börnstein New Series IV/11A3
MSIT ®
254
Al–Mn–Ti
(>99.9 mass%) under argon followed by a HIP treatment at 1200°C and 172 MPa for 2 h. Microstructures were examined optically and by SEM. EDS and WDS analyses were also carried out. Later, [1993Nak] studied the Al rich part of the diagram using arc melted samples (material purities: Al >99.999 mass%, Ti >99.9 mass%, Mn >99.99 mass%) which had been homogenized for 2 d at 1000°C, and powdered metals (purities: Al 99.9 mass%, Mn 99.9 mass%, Ti 99.5 mass%) which were sintered at 1000°C. They found that the ternary phase field centered at Ti25Mn9Al66 . [1996Mab, 1998Mab] used similar techniques and material purities in their investigations at 1150°C. Also, they discovered that the ternary phase melts incongruently. Equilibria involving the (Ti)(hcp), (Ti)(bcc) and the TiAl (L1 0) phases were studied experimentally by [2000Kai] over the temperature range 1000-1300°C. Arc melted samples using Al, Mn and Ti of purity 99.99, 99.9 and 99.7 mass% respectively, were wrapped in Mo foil and sealed in Ar filled silica capsules for annealing. Heat treatment times of 504 h, 168 h and 24 h were used for temperatures of 1000, 1200 and 1300°C, respectively. Microstructural and EPMA analysis revealed equilibria that were qualitatively in agreement with [1954Dom] and [1955Dom], but the composition of the (Ti) phase was substantially different. [1996Che] studied arc-melted and annealed samples by SEM, EDS and EPMA. An alloy composition of Ti-42Al-10Mn was chosen, and annealing temperatures of 1000 and 800°C were used. At the higher temperature, an equilibrium between the (Ti), TiAl and the TiMn2 Laves phase containing a considerable amount of Al, given as Ti(AlMn)2 was found. The appearance of this Laves phase contradicts the work of [1988Has] who suggested the presence of a ternary compound with the formula Ti3Mn2Al3. The data for the equilibrium between the (Ti) and the TiAl phase are consistent with the data of [2000Kai]. At the lower temperature, a three phase equilibrium between (Ti), TiAl and the Laves phase was found suggesting a four phase invariant just above 800°C, as was presented by [1955Dom]. While [1955Dom] and [1960Sat] both indicated the presence of a four-phase solid transition reaction at approximately 860 10°C, they disagree on the phases involved. Considering the extent of observed stability ranges of (Ti) [1955Dom, 2000Kai] and Ti(Mn,Al)2 [1974Dwi, 1977Cha, 1978Jac, 1996Che], the reaction suggested by [1960Sat] involving (Ti) and TiMn perhaps appears least probable. Taking into account currently accepted binary Al-Ti equilibria suggests, however, that TiAl3 rather than (Ti) is likely to be a product in the reaction proposed by [1955Dom]. However, [1996Che] suggest that the nature of the reaction is eutectoid, (Ti) (TiAl)+Ti(Mn,Al)2, but unfortunately, this refers to unpublished work. In a series of articles, [1997But, 1998But, 1999But] investigated the solidification behavior and phase transition sequences in alloys with a Ti:Al ratio of 1:14, plus 5, 10, 20 and 30 at.% Mn (for [1999But] the alloy used was Ti-20Mn-37Al (at.%). DTA measurements and microstructural observation were used and the results were compared with thermodynamic calculations. Only partial agreement was obtained, but the thermodynamic data used for the calculations were only extrapolations from the edge binary systems. Some agreement between the experimental results and those of [2000Kai] were found at 1300°C for the 5 at.% Mn composition. CVM calculations of the locations of the (Ti)/Ti3Al and Ti3Al/TiAl phase boundaries were conducted by [2001Kan]. The results agreed well with the experimental data of [1988Has]. The crystallography of the ternary L12 phase has been discussed by a number of authors [1990Kum, 1991Dur, 1991Mae, 1992Mor, 1996Mab, 2001Mil]. Most authors suggest that the ternary compound is based on the addition of Mn to the binary TiAl3 compound. However, [1991Dur] argues that following geometrical considerations the compound should be based on Ti5Al11. The site selectivity of Mn in TiAl has been studied by [1991Bab]. They conclude that at x < 1.85 in Ti50-xMnxAl50, Mn substitutes for Ti. At higher Mn concentrations, some Al sites are also occupied by Mn. [1993Ers] conducted LMTO-ASA calculations to study the effects of Mn substitution on Al and Ti sites in TiAl. The calculations were made for compositions of TiMnAl2 and Ti2MnAl but they argue that it is possible to extrapolate their results to small percentage additions of Mn to the binary compound. They found that Mn substitutes for Al which is in broad agreement with the earlier work of [1991Bab]. This work was confirmed by [1999Hao], who used atom location channelling enhanced microanalysis. They determined the substitutional sites by Mn (1-5 at.%) in TiAl and (1-2 at.%) in Ti3Al. They found Mn substitution for Ti in Ti3Al.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mn–Ti
255
Binary Systems The Al-Ti and Al-Mn binary systems produced by the MSIT binary evaluation programme have been accepted [2003Pis, 2003Sch]. The Mn-Ti phase diagram is taken from [Mas2]. Solid Phases The stable phases are given in Table 1. There is only one ternary phase which occurs at a composition of Ti25Mn9Al66. It was found that the range of homogeneity increased with increasing temperature [1996Mab]. Allusions had been made to two more ternary compounds (TiMnAl [1974Dwi] and Ti3Mn2Al3 [1988Has]) but these are more likely to be the Ti(Mn,Al)2 Laves phase owing to the extensive solubility of Al in the binary TiMn2 compound [1996Che]. Invariant Equilibria The only ternary invariant reaction established with certainty is the U type liquid transition reaction reported in aluminium-rich alloys at 663°C. A partial reaction scheme is given in Fig. 1. The temperature and composition of the liquid and (Al) solid phases given in Table 2 were read from diagrams given in [1958Mal]; negligible ternary solution in TiAl3 and MnAl6 was indicated. The presence of a four phase reaction has been suggested at approximately 860 10°C by [1955Dom, 1960Sat, 1996Che]. Considering the stability ranges of the phases involved, it is most likely to be the transition reaction, (Ti)+TiAlTi 3Al+Ti(Mn,Al)2. However, owing to uncertainties in the locations of the phase boundaries of the phases involved, it is not possible to give their compositions. Liquidus, Solidus and Solvus Surfaces Investigations of ternary liquidus surfaces are confined to those in the extremely aluminium-rich corner of the system. Figure 2 shows the isotherms of the surfaces of primary crystallization of (Al), TiAl3 and MnAl6, the secondary reaction lines, the tertiary point and the solid limits of the tertiary reaction plane as reported by [1958Mal]. It was necessary to make some adjustments to the location of the isotherms to ensure agreement with the accepted Al-Ti binary phase diagram. Isothermal Sections Figures 3 and 4 show partial isothermal sections for equilibria involving the (Ti), (Ti) and TiAl phases at 1300 and 1200°C taken from [2000Kai]. Very minor adjustments have been made to ensure consistency with the phase boundaries of the accepted Al-Ti phase diagram. Figure 5 presents the equilibria surrounding the L12 ternary phase based on the work of [1998Mab]. However, the information in the original article referring to equilibria between the ternary phase and the Mn8Al5 phase have been ignored here. This binary phase is unstable above 991°C, and moreover, the liquid phase is stable at Al contents up to around 60 at.% in the binary system at this temperature. No equilibria involving the liquid phase are given in the original article. Phase boundaries drawn in Fig. 5 involving the Al-Mn binary should therefore be taken as very tentative. Figure 6 shows the isothermal section for 1000°C. It was constructed by combining the results of [1958Mal, 1974Dwi, 1993Nak, 1996Che] and [2000Kai]. Modifications to the original phase boundaries were made to maintain consistency with the accepted binary phase diagrams. This was particularly important with respect to equilibria between the ternary phase and the Al-Mn binary edge as in Fig. 5. The work of [1993Nak] claimed equilibrium between the ternary phase and the Mn5Al8 ( 2) phase at 1000°C. This is unlikely as this phase is unstable at this temperature in the binary system as indicated above. Therefore, tentative equilibria between the ternary compound and the and 1 phases of the binary Al-Mn system have been added to maintain consistency with the accepted binary diagram. The composition range of the -1 phase increases with increasing temperature which also results in a slight shift of the phase field in Fig. 6 with respect to Fig. 5. The work of [1996Che] and [2000Kai] suggest a higher solubility of Al and Mn in the (Ti) phase than the earlier work. Figure 7 shows the partial isothermal section for 800°C given by [1955Dom] with minor alterations to allow consistency with the binary edges. Figure 8 is a composite Landolt-Börnstein New Series IV/11A3
MSIT ®
256
Al–Mn–Ti
diagram of partial isothermal sections for Al-rich alloys for temperatures 650-400°C taken from [1958Mal] with some minor adjustments to the phase boundaries coincident with the accepted binary phase diagrams. Notes on Materials Properties and Applications Both TiAl and TiAl3 have been identified as possible materials for aerospace applications owing to their high temperature stability and their low density. However, they suffer from poor room temperature ductility and poor workability, even at high temperatures. Hence there has been much interest in studying the mechanical properties of materials consisting of TiAl or TiAl3 alloyed with a third component [1991Kum, 1991Mae, 1991Nic, 1992Win, 1992Mor, 1993Has, 1996Mab, 1999Has, 2000Jin, 2001Mil]. It was found that the addition of Mn to either of the binary compounds improved ductility at room temperature. The increased ductility correlates with a lowering of the antiphase boundary energy allowing formation of partial super dislocations [1992Mor]. Studies of the hydrogen absorption/desorption of Ti3Al found that absorption properties were improved with the addition of Mn to the compound [2001Ish]. Mn substituted Ti3Al showed a reduction in the hydrogen desorption temperature. Nanocrystallites of the L12 phase have been prepared by ball-milling [1998Var]. A Ti25.6 Mn9.4Al65 alloy, homogenized for 100 h at 1000°C was ball milled for 386 h. It was found that the crystallite size of ~20-30 m produced was unchanged after 200 h of milling. References [1954Dom] [1955Dom] [1958Mal]
[1960Sat]
[1974Dwi]
[1974Zwi] [1977Cha] [1978Jac]
[1988Has]
[1989Tsu]
[1990Kum]
[1990Sch]
MSIT®
Domagala, R.F., Rostoker, W., “The System Titanium - Aluminium - Manganese”, Trans. Am. Soc. Met., Reprint No. 4 (1954) (Equi. Diagram, Experimental, #, *, 6) Domagala, R.F., Rostoker, W., “The System Titanium - Aluminium - Manganese”, Trans. Am. Soc. Met., 47, 565-577 (1955) (Equi. Diagram, Experimental, #, *, 6) Mal`tsev, M.V., Van Bok, Y., “Investigation of the Equilibrium Diagram of the Aluminium-Manganese-Titanium System” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., (2), 130-142 (1958) (Experimental, Equi. Diagram, #, *, 3) Sato, T., Huang, Y.-C., Kondo, Y, “A Metallographic Study on Titanium - Aluminium Manganese Alloys”, Sumitomo Light-Metal Technical Reports, 1, 36-44 (1960) (Equi. Diagram, Experimental, 12) Dwight, A.E., “Alloying Behaviour of Zirconium, Hafnium and the Actinides in Several Series of Isostructural Compounds”, J. Less-Common Met., 34, 279-284 (1974) (Experimental, Crys. Structure, 6) Zwicker, U., “Titanium and Titanium Alloys” in “Pure and Applied Metallurgy in Individual Descriptions” (in German), 21, 576-585 (1974) (Equi. Diagram, Review, 22) Chakrabarti, D.J., “Phase Stability in Ternary Systems of Transition Elements with Aluminium”, Metall. Trans. B., 8B, 121-123 (1977) (Experimental, Equi. Diagram, #, *, 13) Jacob, I., Shaltiev, D., “A Note on the Influence of Aluminium on the Hydrogen Sorption Properties of Ti(AlxB1-x)2 (B = Cr, Mn, Fe, Co)”, Mater. Res. Bull., 13, 1193-1198 (1978) (Experimental, Crys. Structure, 10) Hashimoto, K., Doi, H., Kasahara, K., Tsujimoto, T., Suzuki, T., “Effects of Third Elements on the Structures of TiAl-Based Phases”, J. Jpn. Inst. Met., 52, 816-825 (1988) (Experimental, Equi. Diagram, Crys. Structure, 31) Tsujimoto, T., Hashimoto, K., “Structure and Properties of TiAl-Base Alloys Containing Manganese”, Mater. Res. Soc. Symp. Proc., 133, 391-396 (1989) (Experimental, Crys. Structure, 9) 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, Equi. Diagram, Review, 158) Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”, Z. Metallkd., 81(6), 389-396 (1990) (Crys. Structure, Equi. Diagram, Experimental, Review, #, 33)
Landolt-Börnstein New Series IV/11A3
Al–Mn–Ti [1991Bab] [1991Dur]
[1991Kum]
[1991Mae]
[1991Nic]
[1992Kat]
[1992Mor]
[1992Win]
[1993Ers]
[1993Has]
[1993Nak] [1996Che] [1996Mab]
[1997But]
[1997Sah]
[1998But]
[1998Mab]
[1998Var]
[1999But]
Landolt-Börnstein New Series IV/11A3
257
Babu, S.V., Seehra, M.S., “Site Selectivity of Mn Atoms in -TiAl Alloys Determined by X-Ray Scattering”, J. Mater. Res., 6(2), 339-342 (1991) (Crys. Structure, Experimental, 10) 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) Kumar, K.S., Brown, S.A., Whittenberger, J.D., “Compression, Bend and Tension Studies on Forged Al67Ti25Cr8 and Al66Ti25Mn9 L1(2) Compounds”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 481-486 (1991) (Mechan. Prop., Experimental, 11) 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) Nic, J.P., Zhang, S., Mikkola, D.E., “Alloying of Al 3Ti with Mn and Cr to Form Cubic L1(2) Phases”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 697-702 (1991) (Crys. Structure, Equi. Diagram, Experimental, 12) 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, Equi. Diagram, Thermodyn., #, *, 51) Morris, D.G., Gunter, S., “Ordering Ternary Atom Location and Ageing in Ll2 Trialuminide Alloys”, Acta Metall. Mat., 40(11), 3065-3073 (1992) (Crys. Structure, Mechan. Prop., Experimental, 23) Winnicka, M.B., Varin, R.A., “Microstructure and Ordering of L12 Titanium Trialuminides”, Metall. Trans. A, 23A(11), 2963-2972 (1992) (Crys. Structure, Mechan. Prop., Experimental, 24) Erschbaumer, H., Podloucky, R., Rogl, P., Temnitschka, G., Wagner, R., “Atomic Modelling of Nb, V, Cr and Mn Substitutions in -TiAl. I: c/a Ratio and Site Preference”, Intermetallics, 1, 99-106 (1993) (Calculation, Crys. Structure, 31) Hashimoto, K., Masao, K., “Effects of Third Element Addition on Mechanical Properties of TiAl”, Struct. Intermet.: 1st Int. Symp. Struct. Intermetallics, Champion Pa. Sept., 309-318 (1993) (Equi. Diagram, Experimental, Mechan. Prop., 18) Nakayama, Y., Mabuchi, H., “Formation of Ternary L1(2) Compounds in Al3Ti-Base Alloys”, Intermetallics, 1, 41-48 (1993) (Crys. Structure, Equi. Diagram, Experimental, 40) Chen, Z., Jones, I.P., Small, C.J., “Laves Phase in Ti-42Al-10Mn Alloy”, Scr. Mater., 35(1), 23-27 (1996) (Equi. Diagram, Experimental, *, 14) Mabuchi, H., Kito, A., Nakamoto, M., Tsuda, H., Nakayama, Y., “Effects of Manganese on the L12 Compound Formation n Al3Ti-Based Alloys”, Intermetallics, 4, S193-S199 (1996) (Experimental, Equi. Diagram, 34) Butler, C.J., McCartney, D.G., Small, C.J., Horrocks, F.J., Saunders, N., “Solidification Microstructures and Calculated Phase Equilibria in the Ti-Al-Mn System”, Acta Mater., 45(7), 2931-2947 (1997) (Calculation, Equi. Diagram, Experimental, 26) 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) Butler, C.J., McCartney, D.G., “An Experimental Study of Phase Transformations and a Comparison with Calculated Phase Equilibria in Ti-Al-Mn Alloys”, Acta Mater., 46(6), 1875-1886 (1998) (Calculation, Crys. Structure, Equi. Diagram, Experimental, 31) Mabuchi, H., Tsuda, H., Tateno, T., Morii, K., “Phase Equilibrium and the Formation of a Graded Diffusion Layer by Bonding L10- and L12- Alloys in the Ti-Al-Mn System” (in Japanese), J. Jpn. Inst. Met., 62(11), 999-1005 (1998) (Equi. Diagram, Experimental, 13) Varin, R.A., Wexler, D., Calka, A., Tbroniek, L., “Formation of Nanocrystalline Cubic (L1(2)) Titanium Trialuminide by Controlled Ball Milling”, Intermetallics, 6, 547-557 (1998) (Calculation, Crys. Structure, Experimental, Mechan. Prop., 26) Butler, C.J., McCartney, D.G., “Phase Transformations and Phase Equilibria in a Ti-37% Al-20% Mn Alloy”, Intermetallics, 7, 663-669 (1999) (Calculation, Equi. Diagram, Experimental, 14) MSIT ®
Al–Mn–Ti
258 [1999Hao]
[1999Has]
[2000Dub]
[2000Jin]
[2000Kai]
[2001Bra]
[2001Ish]
[2001Kan]
[2001Mil]
[2003Pis]
[2003Sch]
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) Hashimoto, K., Yamamoto, Y., Kimura, T., Nobuki, M., “Effect of Vanadium on Residual Strain in L1 2-Type (AlMn) 3Ti(V) Alloy Powders and Bend Ductility of Pre-Milling Alloys”, Mater. Trans., JIM, 40, 400-403 (1999) (Crys. Structure, Experimental, Phys. Prop., 14) Dubrovinskaia, N., Dubrovinsky, L., Vennstrom, M., Anderson, Y., Abrikosov, I., Eriksson, O., “High-Pressure, High-Temperature In-Situ Study of Alloys: Ti3Al”, Proc. Disc. Meet. Thermodyn. Alloys, 23 (2000) (Thermodyn.) Jinxu, Z., Gengxiang, H., Jiansheng, W., “Electron Structure and Bonding Characteristics of Al3Ti Intermetallic Alloys”, J. Mater. Sci. Lett., 19(18), 1685-1686 (2000) (Equi. Diagram, Experimental, Phys. Prop., 7) Kainuma, R., Fujita, Y., Mitsui, H., Ishida, K., “Phase Equilibria Among Alfa (hcp), (bcc) and (L1(0)) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855-867 (2000) (Crys. Structure, Equi. Diagram, Experimental, #, *, 29) 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, Equi. Diagram, Experimental, #, *, 34) 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) (Phys. Prop., Experimental, 9) 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, 15) 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 L1sub/2/ Phases on Its Basis”, Intermetallics, 9, 839-845 (2001) (Crys. Structure, Experimental, Mechan. Prop., 36) Pisch, A., “Al-Mn (Aluminium-Manganese)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 40) Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 86)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al)
(Al) < 660.452 ( Mn) 1246 - 1138 ( Mn) 1138 - 1100
MSIT®
Pearson Symbol/ Space Group/ Prototype hP2 P63/mmc Mg cF4 Fm3m Cu cI2 Im3m W cF4 Fm3m Cu
Lattice Parameters Comments/References [pm] a = 269.3 c = 439.8
at 25°C, 20.5 GPa [Mas2]
a = 404.96
at 25°C [Mas2]
a = 308.0
[Mas2]
a = 386.0
[Mas2]
Landolt-Börnstein New Series IV/11A3
Al–Mn–Ti Phase/ Temperature Range [°C] (Mn) 1100 - 727 (Mn) < 727 (7Ti)
(Ti) 1670 - 882 (Ti) < 882 MnAl6 < 705 Mn4Al11 (HT) 916 - 1002 < 1177 1 < 1048 2, Mn5Al8 < 991 J, Mn3Al2 < 1312 TiMn 950 TiMn 1200 Ti(Mn,Al)2
TiMn2 < 1325 TiMn3 ~1250 - 950
Landolt-Börnstein New Series IV/11A3
Pearson Symbol/ Space Group/ Prototype cP20 P4132 Mn cI58 I43m Mn hP3 P6/mmm 7Ti cI2 Im3m W hP2 P63/mmc Mg oC28 Cmcm MnAl6 oP160 Pnma cI2 Im3m W
259
Lattice Parameters Comments/References [pm] a = 631.52
[Mas2]
a = 891.26
at 25°C [Mas2]
a = 462.5 c = 281.3
at 25°C, HP 1 atm [Mas2]
a = 330.65
[Mas2]
a = 295.06 c = 468.35
at 25°C [Mas2]
a = 755.51 b = 649.94 c = 887.24 ?
[V-C2]
a = 306.3
[V-C2] Also designated Mn55Al45
hR26 R3m Cr5Al8 hP2 P63/mmc Mg tP30 P42/mnm CrFe -
a = 1273.9 c = 1586.1
at 58 at.% Al [V-C2]
a = 270.5 - 270.5 c = 436.1 - 438
44.2 - 44.9 at.% Al [2003Pis]
a = 888 c = 454.2
[Mas2], [V-C2]
-
[Mas2]
hP12 P63/mmc MgZn2
a = 495.3 c = 805.8 a = 497.8 c = 815.1 a = 499.7 c = 890.8 a = 483.33 0.09 c = 793.84 0.11 -
at 17 at.% Al and 50 at.% Mn [1978Jac]
o**?
[2003Pis]
at 33.3 at.% Al and 33.3at.% Mn [1974Dwi] at 37.5 at.% Al and 25 at.% Mn [1988Has] [Mas2], [V-C2]. 60-70 at.% Mn [Mas2]
MSIT ®
Al–Mn–Ti
260 Phase/ Temperature Range [°C] TiMn4
Ti3Al < 1164 (up to 10 GPa at RT)
Ti3Al (I) 15 to > 41 GPa TiAl < 1463
Pearson Symbol/ Space Group/ Prototype hR53 R3m Co5Cr2Mo3 hP8 P63/mmc Ni3Sn
hP16 P63/mmc TiNi3 tP4 P4/mmm AuCu
Lattice Parameters Comments/References [pm] a = 1100.3 c = 1944.6
a = 580.6 c = 465.5 a = 574.6 c = 462.4 a = 531.2 c = 960.4
a = 400.0 c = 407.5 a = 398.4 c = 406.0
at 38 at.% Al [L-B] [1997Sah] at 16 GPa, not confirmed by [2000Dub] (0-35 GPa, 25-2250°C) 46.7 to 66.5 at.% Al [1992Kat] 50 to 62 at.% Al at 1200°C [2001Bra] L10 ordered phase (“ -TiAl”) at 50.0 at.% Al, [2001Bra] at 62.0 at.% Al, [2001Bra]
oC12 Cmmm ZrGa2
a = 1208.84 b = 394.61 c = 402.95
tP4 P4/mmm AuCu
a = 403.0 c = 395.5
tI24 I41/amd HfGa2 tP32 P4/mbm Ti3Al5
a = 397.0 c = 2497.0
chosen stoichiometry [1992Kat] summarizing several phases: metastable modification of TiAl2, only observed in as-cast alloys [2001Bra]; listed as TiAl2(h) by [1990Sch] (66 to 67 at.% Al, 1433-1214°C) Ti1-xAl1+x; 63 to 65 at.% Al at 1250°C, stable range 1445-1170°C [2001Bra]; listed as orthorhombic, Pmmm, with pseudotetragonal cell by [1990Sch] (range ~1445-1424°C). at 1300°C [2001Bra] stable structure of TiAl2 <1216 [2001Bra];
a = 1129.3 c = 403.8
listed as TiAl2(r) by [1990Sch] Ti3Al5, stable below 810°C [2001Bra]
TiAl2 < 1199
MSIT®
[V-C2], [Mas2]. Phase referred to as TiMn5 in [V-C2]. Prototype given as ~ (Mo,Ni) in [Mas2] ~20 to 38.2 at.% Al [1992Kat] D019 ordered phase (“2-Ti3Al”) [1997Sah]. at 22 at.% Al [L-B]
Landolt-Börnstein New Series IV/11A3
Al–Mn–Ti Phase/ Temperature Range [°C] “Ti2Al5” 1416 - 990
Pearson Symbol/ Space Group/ Prototype
261
Lattice Parameters Comments/References [pm] chosen stoichiometry [1992Kat] summarizing several phases: Ti5Al11 stable range 1416 - 995°C [2001Bra] 66 to 71 at.% Al at 1300°C [2001Bra] (including the stoichiometry Ti2Al5!); [1990Sch] claimed: 68.5 to 70.9 at.% Al and range 1416 - 1206°C; at 66 at.% Al [2001Bra] * AuCu subcell only at 71 at.% Al [2001Bra] * AuCu subcell only “Ti2Al5” ~1215 - 985°C [1990Sch]; included in homogeneity region of Ti5Al11 [2001Bra] 74.2 to 75.0 at.% Al [1992Kat] 74.5 to 75 at.% Al at 1200°C [2001Bra] D022 ordered phase stable above 735°C (Al-rich) [2001Bra]
tetragonal superstructure of AuCu-type [2001Bra]
tP28 P4/mmm “Ti2Al5”
a* = 395.3 c* = 410.4 a* = 391.8 c* = 415.4 a = 390.53 c = 2919.63
TiAl3 (h) < 1393
TiAl3 (l) < 950 (Ti-rich) * -1, Ti25Mn9Al66
tI32 I4/mmm TiAl3 (l) cP4 Pm3m AuCu3
a = 384.9 c = 860.9 a = 387.7 c = 3382.8
74.5 to 75 at.% Al [2001Bra]
a = 395.8 a = 395.9
Ti25Mn8Al67 [1991Nic] Ti43Mn11Al66 [2001Mil]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
(Ti) + TiAl Ti3Al + Ti(Mn,Al)2 L + TiAl3 (Al) + MnAl6
~865
U1
-
Composition (at.%) Al Mn Ti -
663
U2
L TiAl3 (Al) MnAl6
98.99 75.0 99.37 85.7
Landolt-Börnstein New Series IV/11A3
0.94 0 0.09 14.3
0.07 25.0 0.54 0
MSIT ®
Al–Mn–Ti
262
Al-Mn-Ti
Al-Ti
Al-Mn
664 L+TiAl3+MnAl6
663
p
l+TiAl3(Al)
L+TiAl3 (Al)+MnAl6
U2
(Al)+TiAl3+MnAl6
658.5
e
l(Al)+MnAl6
Fig. 1: Al-Mn-Ti. Partial reaction scheme
Al Fig. 2: Al-Mn-Ti. Partial liquidus surface
Data / Grid: at.%
p (664°C)
Axes: at.%
(αAl)
U2 (663°C) e (658.5°C)
MnAl6
0°C 90
°C 700 °C 0 75 °C 0 80 0°C 85
TiAl3
C 0° 95
Ti 2.30 Mn 0.00 Al 97.70
MSIT®
Ti 0.00 Mn 2.30 Al 97.70 Landolt-Börnstein New Series IV/11A3
Al–Mn–Ti
263
Ti 40.00 Mn 0.00 Al 60.00
Fig. 3: Al-Mn-Ti. Partial isothermal section at 1300°C
Data / Grid: at.% Axes: at.%
TiAl 50
50
(α Ti) (β Ti)
60
40
10
Ti 70.00 Mn 0.00 Al 30.00
20
Ti 40.00 Mn 0.00 Al 60.00
Fig. 4: Al-Mn-Ti. Partial isothermal section at 1200°C
Ti 40.00 Mn 30.00 Al 30.00 Data / Grid: at.% Axes: at.%
TiAl 50
60
50
(αTi)
40
(β Ti)
70
Ti 75.00 Mn 0.00 Al 25.00 Landolt-Börnstein New Series IV/11A3
30
10
20
30
Ti 40.00 Mn 35.00 Al 25.00
MSIT ®
Al–Mn–Ti
264
Al Fig. 5: Al-Mn-Ti. Partial isothermal section at 1150°C (Al-rich part)
Data / Grid: at.% Axes: at.%
L
20
80
TiAl3 Ti2Al5 TiAl2
τ1
40
60
TiAl
γ 60
40
ε
Ti(Mn,Al)2 20
Ti 70.00 Mn 0.00 Al 30.00
40
60
Al
Data / Grid: at.%
L
Fig. 6: Al-Mn-Ti. Isothermal section at 1000°C TiAl3 Ti2Al5
Axes: at.%
20
80
Mn4Al11(h)
TiAl2
γ1
τ1
40
60
TiAl
γ
60
ε
Ti3Al Ti3Al+(β Ti) 80
(α Ti)
(β Ti)
Ti
MSIT®
20
Ti 0.00 Mn 70.00 Al 30.00
(β Ti)+Ti(Mn,Al)2
40
Ti(Mn,Al)2
β TiMn
60
Ti(Mn,Al)2 +(β Mn)
40
(β Mn)
TiMn380 TiMn4
20
(αMn)
Mn
Landolt-Börnstein New Series IV/11A3
Al–Mn–Ti
265
30.00 Ti 0.00 Mn Al 70.00
Fig. 7: Al-Mn-Ti. Partial isothermal section at 800°C (Ti-rich part)
Data / Grid: at.% Axes: at.%
TiAl2 40
60
TiAl
TiAl+Ti(Mn,Al)2+τ 1 60
40
Ti3Al
80
20
Ti3Al+(β Ti)+Ti(Mn,Al)2 (αTi) (β Ti)
(β Ti)+α TiMn 20
Ti
40
Al Fig. 8: Al-Mn-Ti. Isothermals 650 to 400°C (Al-rich part)
αTiAl
60
30.00 Ti Mn 70.00 0.00 Al
Data / Grid: at.% Axes: at.%
400°C 500°C 550°C
600°C
(αAl)
400 °C 50 0°C 55 0°C
650°C
65 0° C
60 0° C
(αAl)+MnAl6
Ti 1.00 Mn 0.00 Al 99.00 Landolt-Börnstein New Series IV/11A3
(αAl)+TiAl3
(αAl)+TiAl3+MnAl6
Ti 0.00 Mn 1.00 Al 99.00
MSIT ®
266
Al–Mo–Ni
Aluminium – Molybdenum – Nickel Kostyantyn Korniyenko, Vasyl Kublii Literature Data 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 [1969Mar, 1969Vir, 1971Pry]. Binary Systems The Al-Mo, Al-Ni and Mo-Ni systems are accepted from [2003Sch], [2003Sal] 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 5 of composition Mo7,5Ni58,0 Al34,5 at 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.%) Mo50 Ni25 Al25 , Mo43Ni31Al26 and Mo9Ni53Al38 , the latter being fairly close to the composition of the reported 5 phase [1959Gua2]. Thus, [1969Vir] concluded that the 5 phase in fact did not belong to the Al-Mo-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 5 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, -1, Mo(NixAl1-x)3 [1969Mar, 1969Vir, 1971Pry] and -2. The composition of -2 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 -1 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 / ´ boundary may be obtained from [1989Gai] and [1989Hon1]. A model based on X-ray measurements to show the effect of Mo on the ´ structure has been suggested by [1977Aig]. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni
267
Pseudobinary Systems On the basis of X-ray, DTA as well as optical microscopy data [1986Mas2] plotted the phase diagram of the partial pseudobinary NiAl-Mo system. The temperature of the L+ equilibrium is equal to 16007°C; the maximum solubility of molybdenum in the 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 phase at 1651°C [2003Sal], 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 phase and the reaction temperature of U1 (1340°C), the higher pseudobinary eutectic temperature (1600°C) is recommended. The vertical section Mo-Ni3Al, according to the data of [1971Pry], demonstrates a peritectoid reaction + ´+, 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 observed in the system. [1977Pea] reported the equilibrium L++ at 1300°C, but later it was established by [1977Aig, 1983Nas] and [1983Wak] that instead of the phase the ´ phase takes part in the eutectic reaction L+ + ´, 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 , , , ´ and 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 + 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 , and phases participating in the equilibrium LU2+ + at 1310°C. [1974Tho, 1976Nes, 1976Hen] and [1976Spr] discovered by directional solidification studies the existence of eutectic reactions L+ ´, L+ and L + . 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 [2003Sal]. 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].
Landolt-Börnstein New Series IV/11A3
MSIT ®
268
Al–Mo–Ni
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 ´ phase is much smaller than experimental data, and also the position of the ´+ 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 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 / ´ 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 (Ni) solid solution area increases with rising temperature. The position of the nickel-rich boundary of the ´ 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 and 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 ´ phase particles in ternary Al-Mo-Ni alloys was used. In the calculation, the alloy composition was adjusted in order to have the same fraction of the ´ phase as experimentally observed. This gave a small difference in composition compared with the experimental data.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni
269
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, and ´. The 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 ´ 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 ´ based alloys molybdenum substitutes preferentially for aluminium. The microstructure and chemical characteristics of the nanocrystalline phase are studied by [2002Alb]. It was established that the addition of molybdenum tends to slightly refine the grain size of the phase based alloy. The specimens with 2, 4 or 6 at.% Mo are polycrystalline containing, at the same time, the 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 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 / ´ matrix reinforced with faceted Mo fibres, which are primarily square in cross section and with the following orientation relationship: [001] ´//[001], (010) //(010) ´//(110), (100) ´//(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 (, ´ and 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 solidified ´, 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 ´´ and ´ phases to be parallel: [100] ´//[110] ´, (010) ´´//(111) ´ and [001] //[112] ´. [2001Kai] studied the effect of molybdenum on the morphological stability of the interface between the ´ and 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. References [1925Pfa] [1933Roe] [1951Ham]
Landolt-Börnstein New Series IV/11A3
Pfautsch, H., “The System Aluminium-Molybdenum-Nickel” (in German), Z. Metallkd., 19(4), 125-127 (1925) (Experimental, Equi. 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 [2003Sch]
MSIT ®
270 [1952Tay]
[1954Ada] [1958Woo]
[1959Gua1] [1959Gua2] [1960Bag] [1960Vig]
[1962For] [1963Arb]
[1964Lea] [1965Ram]
[1967Bel]
[1969Mar]
[1969Rec]
[1969Vir] [1970Cli] [1971Cli]
[1971Nal]
[1971Pry]
[1971Rex]
MSIT®
Al–Mo–Ni 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, Equi. 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., “-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 ( ´ phase)”, Trans. Met. Soc. AIME, 215, 807-814 (1959) (Equi. Diagram, Experimental, 27) Guard, R.W., Smith, E.A., “Constitution of Nickel-Base Ternary Alloys”, J. Inst. Met., 88, 283-287 (1959-1960) (Equi. 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, Equi. 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, Equi. Diagram, 16) Forsyth, J.B., Gran, G., “The Structure of the Intermetallic Phase (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) (Equi. 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 - (U)-Related Phases”, Z. Metallkd., 60, 594-600 (1969) (Equi. 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) (Equi. 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) Pryakhina, L.I., Myasnikova, K.P., Markiv, V.Ya., Burnasheva, V.V., “Investigation of the Molybdenum-Nickel-Aluminium Ternary System”, in “Phase Diagrams of Metal Systems” (in Russian), Nauka, Moscow, 112-116 (1971) (Equi. 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, Equi. Diagram, 23) Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni [1973Wal] [1974Kau] [1974Tho] [1976Hen] [1976Jac]
[1976Mon] [1976Nes]
[1976Spr] [1977Aig]
[1977Pea]
[1978Gul]
[1980Fer]
[1981Sch]
[1982Shi]
[1983Nas] [1983Och] [1983Wak]
[1984Kov1]
Landolt-Börnstein New Series IV/11A3
271
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) (Equi. 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 ' in - (Ni-Al-Mo) Eutectics”, Scr. Metall., 10, 955-957 (1976) (Equi. 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, Equi. 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) (Equi. 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) (Equi. 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) (Equi. Diagram, Crys. Structure, Experimental, Review, 20) as quotet by [1993Kub] Pearson, D.D., Lemkey, F.D., “Solidification and Properties of / ´-Mo Ductile/Ductile Eutectic Superalloy”, Proc. Conf. Solidification and Casting of Metals, Metals Soc. London, Sheffield, U.K., 526-532 (1977) (Publ. 1979) (Equi. 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, Equi. 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 / '-(Mo) Eutectic Alloy”, Conf. Mat. Eng., Freund Publ. House Tel Aviv, 1981, 10-13 (1981) (Equi. 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) (Equi. 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) (Equi. 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 / ' -Mo Aligned Eutectics”, Acta Metall., (11), 1937-1944 (1983) (Equi. Diagram, Experimental, 19) as quoted by [1993Kub] Kovacova, K., Kristin, J., “Morphological Properties of / '-Mo Eutectic Composite Material” (in Czech), Kovove Mater., 22(3), 347-356 (1984) (Equi. Diagram, Experimental, 18) as quoted by [1993Kub]
MSIT ®
272 [1984Kov2] [1984Mir]
[1984Och1] [1984Och2]
[1984Sch]
[1985Mis]
[1985Nas]
[1986Hor]
[1986Hua] [1986Kau]
[1986Mas1]
[1986Mas2]
[1987Kha] [1987San]
[1987Sve]
[1988Li]
[1988Mas]
[1989Ell]
MSIT®
Al–Mo–Ni Kovacova, K., “Undirectional Solidification of Ni-Al-Mo Alloy”, J. Cryst. Growth, 66, 426-430 (1984) (Equi. 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) (Equi. Diagram, Experimental, #, 12) Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behavior of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32, 289-298 (1984) (Equi. Diagram, Experimental, 90) Ochiai, S., Mishima, Y., Suzuki, T.S., “Lattice Parameter Data of Ni( ), Ni3Al ( ') and Ni3Ga ( ') 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. Mat. Sci. Lett., (3), 363-366 (1984) (Experimental, Crys. Structure, 6) Mishima, S., Ochiai, S., Suzuki, T.Y., “Lattice Parameters of Ni ( ), Ni3Al ( ') and Ni3Ga ( ') 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”, Mat. Res. Soc. Symposia Proceedings, Kreh, C.C., Liu, C.T., Stoloff, N.S., (Eds.), MRS-Pennsylvania, Boston, Massachusetts (1984), 39, 423-427 (1985) (Equi. Diagram, Review, 15) Horita, Z., Sano, T., Nemoto, M., “Identification of Fine Particles in Unidirectionally Solidified Ni-Al-Mo 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 Nickel-Aluminium-Molybdenum System at 1300-2000°C”, Russ. Metall., (1), 203-209 (1986), translated from Izv. Akad. Nauk SSSR, Met., (1), 198-205 (1986) (Equi. 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) (Equi. 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 + + ) Line in the Ni-Mo-Al System and Plane Front Solidification in / ´ - Alloys”, Russ. Metall., (6), 186-192 (1987), translated from Izv. Akad. Nauk SSSR, Met., (6), 183-189 (1987) (Equi. 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) (Equi. 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) Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni [1989Gai]
[1989Hon1]
[1989Hon2]
[1989Mas]
[1990Yam] [1991Eno]
[1991Kim]
[1991Mis]
[1991Sas]
[1991Sch]
[1992Mur]
[1993Kha]
[1993Kub]
[1994Mur] [1995Gri] [1996Pau]
[1996Vik]
[1997Bou]
Landolt-Börnstein New Series IV/11A3
273
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 Solvus Surface in Ni-Al-X (X = Cr, Mo and W) Ternary Systems”, ISIJ International, 29(1), 78-84 (1989) (Equi. Diagram, Experimental, #, 25) Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of ' Solvus in Ni-Al-X Ternary Systems”, Mat. Res. Soc. Symp. Proc., 733, 431-438 (1989) (Equi. 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) (Equi. Diagram, Crys. Structure, #, 12) Yamamoto, M., Iada, J., Nenno, S., “The Microstructure of a Two-Phase Mixture in a Ni75Mo20 Al5 Alloy”, J. Mat. Sci. Lett., 9, 34-35 (1990) (Crys. Structure, Experimental, 5) Enomoto, M., Harada, H., Yamazaki, M., “Calculation of ´/ Equilibrium Phase Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15(2), 143-158 (1991) (Assessment, Calculation, Equi. 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 [2003Sal] Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Assessment, Equi. 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, Equi. Diagram, Experimental, Mechan. Prop., 10) Schuster, J.C., Ipser, H., “The Al-Al8Mo3 Section of the Binary System AluminiumMolydenum”, Met. Trans., A22, 1729-1736 (1991) (Experimental, Crys. Structure, Equi. Diagram, 20) Murakami, Y., Otsuka, K., Hanada, S., Watanabe, S., “Crystallography of Stress-Induced B27R 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) (Equi. 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, Equi. 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 Mo 4Al17 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) MSIT ®
274 [1997Jin]
[1997Poh] [1997Pot]
[1998Mur] [1998Rav]
[1998Sim]
[1999Kau]
[1999Lu]
[2000Bor]
[2001Ter]
[2001Kai]
[2002Alb]
[2002Gru]
[2002Guo]
[2003Sal]
[2003Sch]
MSIT®
Al–Mo–Ni Jin, Y., Chaturvedi, M.C., Han, Y.F., Zhang, Y.G., “Crystal Structure of -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, Equi. 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) (Equi. 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.) 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 '/ Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168-175 (2001) (Equi. 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, Equi. 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; to be published, (2003) (Crys. Structure, Equi. 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; submitted for publication, (2003) (Crys. Structure, Equi. Diagram, Assessment, 61)
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni
275
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 MoxNiyAl1-x-y
, (Mo) < 2623 ,(Mo1-x-yNiyAlx)
, (Ni) < 1455 , (MoxNi1-x)
, (Ni1-xAlx)
, MoNi4 < 870 ´´, MoNi3 < 910
Landolt-Börnstein New Series IV/11A3
Pearson Symbol/ Lattice Parameters Comments/References Space Group/ [pm] Prototype cF4 a = 404.88 pure Al at 24°C [V-C] Fm3m Cu x = 0, y = 0 to 0.004 [2003Sal] 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] a = 314.7 pure Mo, at 25°C [V-C] cI2 Im3m W a = 314.6 x = 0, y = 0.004 [1980Fer] a = 314.5 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] pure Ni at 25°C [1984Och2, Mas2] a = 352.40 cF4 pure Ni at 20°C [V-C] a = 352.32 Fm3m quenched from 800°C [V-C] Cu a = 355.8 x = 0.03, quenched from 1000°C [1984Och1, a = 353.9 1984Och2, 1985Mis] x = 0.06, quenched from 1000°C [1984Och1, a = 355.2 1984Och2, 1985Mis] x = 0.09, quenched from 1000°C [1984Och1, a = 356.5 1984Och2, 1985Mis] x = 0.097 [1980Fer] a = 356.3 x = 0.218 [1980Fer] a = 361.0 x = 0 to 0.2 [2003Sal] x = 0.2 at 1372°C [2003Sal] x = 0.025 Slowly cooled alloy [1952Tay] a = 352.8 x = 0.05 Slowly cooled alloy [1952Tay] a = 353.2 a = 572.0 [V-C] tI10 c = 356.4 I4/m MoNi4 a = 506.4 [V-C] oP8 b = 422.2 Pmmn c = 444.8 TiCu3
MSIT ®
Al–Mo–Ni
276 Phase/ Temperature Range [°C]
, MoNi < 1362
Pearson Symbol/ Space Group/ Prototype oP112 Cmcm MoNi
(Mo,Ni,Al)1 (Ni,Mo,Al)1
MoAl12 < 712 MoAl5 (h2) 846 to 800 - 750
MoAl5 (h1) 800 - 750 to ~648 MoAl5 (r)
648
Mo5Al22 964 to 831
cI26 Im3 WAl12 hP12 P6 3 WAl5 hP60 P3 MoAl5 (h1) hP36 R3c MoAl5 (r) oF216 Fdd2 Mo5Al22
Mo4Al17 < 1034
mC84 C2 Mo4Al17
MoAl4 1177 to 942
mC30 Cm WAl4
MSIT®
Lattice Parameters Comments/References [pm]
a = 910.8 b = 910.8 c = 885.2 a = 455 b = 1663 c = 873
a = 757.3 a = 758.15 a = 491.2 c = 886.0 a = 489 c = 880 a = 493.3 c = 4398 a = 495.1 c = 2623
a = 7382 3 b = 916.1 0.3 c = 493.2 0.2
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] 92.4 at.% Al [1991Sch] [1954Ada] [1980Fer] 83.8 at.% Al [1991Sch] [1980Fer] at 83.3 at.% Al [1991Sch]
at 83.3 at.% Al [1991Sch]
81.7 at.% Al [1991Sch] [1995Gri]
80.9 at.% Al [1991Sch] a = 915.8 0.1 [1995Gri] b = 493.23 0.08 c = 2893.5 0.5 = 96.71 0.01 79 to 80 at.% Al [1991Sch] [1964Lea] a = 525.5 0.5 b = 1776.8 0.5 c = 522.5 0.5 = 100.88 0.06° [1991Sch] a = 525.5 b = 1176.8 c = 522.5 = 100.7°
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni Phase/ Temperature Range [°C] Mo1-xAl3+x 1260 to 1154 MoAl3 1222 to ~818
Pearson Symbol/ Space Group/ Prototype cP8 Pm3n Cr3Si mC32 C2/m MoAl3
Mo3Al8 < 1555 10
mC22 Cm Mo3Al8
Mo2Al3 1570 to 1490 MoAl 1750 to 1470
-
Mo3Al
cP8 Pm3n Cr3Si
2150 (Mo,Ni,Al)3 (Al,Mo,Ni)1 J, NiAl3 < 856
Ni2Al3 < 1138
´, Ni3Al4 < 702
Landolt-Börnstein New Series IV/11A3
cP2 Pm3m CsCl
oP16 Pnma NiAl3
oP16 Pnma Fe3C hP5 P3m1 Ni2Al3
cI112 Ia3d Ni3Ga4
277
Lattice Parameters Comments/References [pm]
a = 494.5 0.1
76 to 79 at.% Al [1991Sch] [1991Sch]
a = 1639.6 0.1 at 75 at.% Al [1991Sch] b = 359.4 0.1 c = 838.6 0.4 = 101.88 0.07 72 to 75 at.% Al [Mas2] a = 920.8 0.3 [1962For] b = 363.78 0.03 c = 1006.5 0.3 = 100.78 0.05° Called “1” (h) [1971Rex] 46 to 51.7 at.% Al [Mas2] Called “2” (h) [1971Rex] a = 309.8 [1971Rex] a = 309.8 to 309.9 [1980Fer] 22 to 27 at.% Al [Mas2] a = 495 [1958Woo] a = 487.6 at 6 at.% Ni, 75 at.% Mo [1969Vir] 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
[L-B]
[1996Vik]
[1997Bou, V-C]
36.8 to 40.5 at.% Al [Mas2] [L-B] [1997Bou, V-C] [1989Ell, V-C]
MSIT ®
Al–Mo–Ni
278 Phase/ Temperature Range [°C] , NiAl < 1651
Pearson Symbol/ Space Group/ Prototype cP2 Pm3m CsCl
Lattice Parameters Comments/References [pm]
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
´, Ni3Al < 1372
cP4 Pm3m Cu3Au
Ni3(Al1-xMox)
a = 753 b = 661 c = 376 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
MSIT®
mP22 P2 1/c Ni2Al9
a = 868.5 0.6 b = 623.2 0.4 c = 618.5 0.4 = 96.50 0.05°
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]
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 1000C [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]
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni Phase/ Temperature Range [°C] NixAl1-x 0.60 < x < 0.68
Pearson Symbol/ Space Group/ Prototype tP4 P4/mmm AuCu
m**
Ni2Al
hP3 P3m1 CdI2 aP126 P1
D1 (Al-Ni) decagonal D4 (Al-Ni) decagonal * -1, Mo(NixAl1-x)3 tI8 I4/mmm TiAl3
superstructure? c = 2c0 * -2, Mo11 Ni14 Al75 orthorhombic
Landolt-Börnstein New Series IV/11A3
279
Lattice Parameters Comments/References [pm]
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 = 94.3° a = 407 c = 499 a 1252 b 802 c 1526 90° 109.7° 90° 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
a = 1005.4 0.4 b = 1528.8 0.4 c = 851.9 0.2
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]
Metastable [1988Li] Metastable [1988Li] [1971Pry], at 3 to 8 at.% Ni 25 at.% Mo, 600°C at 1.6 to 6.0 at.% Ni [2002Gru] Called “Mo 2NiAl5” [1965Ram], from a three phase 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]
MSIT ®
Al–Mo–Ni
280 Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Al
Mo
Ni
L+
1600
e1 (max)
L
45 1 48.5
10 97.5 2.5
45 1.5 49
L + + ´
1340
U1
L ´
20.5 <4 <33 <25
13 95.5 1.5 2
66.5 >0.5 >65.5 >73
L+ +
1310
U2
L
8.58 <0.5 <2.5 <8
27.22 97.5 50.5 20
64.2 >2.0 >47 >72
L + + ´
1300
E
L ´
18 <2 <17 <20.5
16 96.0 10 5.5
66 >2 >73 >74
+ ´ +
1130
U3
´
1 2 10 20
98.2 49 14 5
0.8 49 76 75
+ ´+ ´´
890
U4
´ ´´
1 5.5 20.5 2.5
51 14 5.5 22.5
48 80.5 74 75
+ ´´ ´ +
730
U5
´ ´´
5.5 20 3 1
11.5 5 22 19
83 75 75 80
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni
Fig. 1: Al-Mo-Ni. Partial pseudobinary system Mo-NiAl
281
1900
1850°C
L
1800
L+α
1700
Temperature, °C
1651°C L+β
1600±7°C
1600
e1max
β
1500
1400
α +β 1300
1200
1100
Mo 20.00 Ni 40.00 Al 40.00
Mo 0.00 Ni 50.00 Al 50.00
10
Mo, at.%
Mo-Ni
Al-Mo-Ni
Al-Ni
1600 e1 Lα+β
1372 p1 l + γ γ´
1362 p2 l+αδ 1317 e3 lγ+δ
1340
L + β α + γ´
1369 e2 l β + γ´
U1
α+β+γ´ 1310
L+δα+γ
U2
L+α+γ´
L+α+γ δ+α+γ
ca.1300
L α + γ + γ´
E
α+γ+γ´ ca.1130 910 p3 δ + γ γ´´
γ + α γ´ + δ
U3 δ+γ+γ´
α+δ+γ´
ca.890
870 p4 γ + γ´´ θ
δ + γ γ´ + γ´´
γ+γ´+γ´´ ca.730
γ + γ´´ γ´ + θ
θ+γ´+γ´´
U4
δ+γ´+γ´´
U5
θ+γ+γ´
Fig. 2: Al-Mo-Ni. Reaction scheme of the partial Mo-NiAl-Ni system Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Mo–Ni
282
Al
Data / Grid: at.%
Fig. 3: Al-Mo-Ni. Liquidus surface in the Ni-rich region
Axes: at.%
20
80
40
60
e1max 60
40
β e2
U1
γ'
1369
80
160 0
α
E
δ 20
Mo
1445
60p e 2 3
40
1415 1425
γ
U2 80
Al
Ni
Data / Grid: at.%
Fig. 4: Al-Mo-Ni. Partial isothermal section at 1260°C
Axes: at.%
20
80
40
60
60
40
γ'
80
20
α+γ +γ '
Mo
γ
α+δ+γ
α
MSIT®
p1
20
20
40
δ
60
80
Ni
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni
283
Al Fig. 5: Al-Mo-Ni. Isothermal section at 1200°C, calculated by [1974Kau]
Data / Grid: at.% Axes: at.%
L
20
80
Mo3Al8
L+β +Mo3Al8 40
60
Mo3Al+Mo3Al8+β
β
60
40
Mo3Al 80
α+γ '+β
γ'
α+Mo3Al+β
20
γ '+γ +α
α
γ
α +γ +δ 20
Mo
40
δ
60
80
Al
Ni
Data / Grid: at.%
Fig. 6: Al-Mo-Ni. Partial isothermal section at 1100°C
Axes: at.%
20
80
40
60
β 60
40
γ'
α+β +γ '
80
20
δ+γ +γ ' α+δ+γ '
γ
α
Mo
Landolt-Börnstein New Series IV/11A3
20
40
δ
60
80
Ni MSIT ®
Al–Mo–Ni
284
Al
Data / Grid: at.%
Fig. 7: Al-Mo-Ni. Partial isothermal section at 1000°C
Axes: at.%
10
90
L 20
80
MoAl3
τ2
Mo3Al8
τ1
30
τ 2+L+Ni2Al3
τ 1+τ 2+Ni2Al3
70
Ni2Al3 10
Mo 40.00 Ni 0.00 Al 60.00
20
30
Al 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
Mo 0.00 Ni 40.00 Al 60.00 Data / Grid: at.% Axes: at.%
MoAl4 MoAl3 Mo3Al8
L
20
L+MoAl3 80
L+τ 1
τ1
L+τ 1+Ni2Al3
40
Ni2Al3 60
60
40
β γ´
80
20
α+β +γ ´ δ+γ +γ ´
α+γ ´+δ
γ
α
Mo
MSIT®
20
40
δ
60
80
Ni
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ni
285
Al 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
Data / Grid: at.% Axes: at.%
L MoAl4 MoAl3 Mo3Al8
20
80
τ2
τ1
Ni2Al3
40
60
60
40
β
δ +γ +γ ´
α +δ +γ ´
α 20
Mo
γ´
α+β +γ ´
80
40
δ
60
20
γ 80
Al
Ni
Data / Grid: at.%
Fig. 10: Al-Mo-Ni. Partial isothermal section at 880°C
Axes: at.%
20
80
40
60
β 60
40
α+β +γ '
80
γ'
20
γ ´´+γ +γ ' α+γ '+δ
δ+γ '+γ ´´
α
Mo
Landolt-Börnstein New Series IV/11A3
20
40
δ
60
γ ´´ 80
γ
Ni
MSIT ®
Al–Mo–Ni
286
Al
Data / Grid: at.%
Fig. 11: Al-Mo-Ni. Partial isothermal section at 700°C
Axes: at.%
20
80
40
60
β 60
40
Ni5Al3
α+β +γ '
80
α+γ '+δ
Mo α
20
40
γ'
20
δ+γ '+γ ´´
δ
γ
γ ´´ 80 θ
60
γ +γ '+γ ´´ γ +γ '+θ
Ni
1600
Fig. 12: Al-Mo-Ni. Partial isopleth at 14 at.% Al
1500
L 1400
Temperature, °C
1300
L+γ
L+α
L+γ'+α 1300
L+γ+α
1200
L+γ+γ'
γ
γ+γ'+α
1130 1100
γ+γ'
γ+γ'+δ 1000
γ'+δ +α 890
900
δ +γ'
800
γ´´+γ'+δ
γ+γ'+γ´´ 730
700
γ+γ'+θ
γ'+γ´´
600
Mo 28.00 Ni 58.00 Al 14.00
MSIT®
60
γ'+γ´´+θ
70
Ni, at.%
80
Mo 0.00 Ni 86.00 Al 14.00
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti
287
Aluminium – Molybdenum – Titanium Ludmila Tretyachenko Literature Data The phase equilibria in this system were assessed by [1993Bud] based on results published up to 1990. Experimental data in these investigations have been interpreted from earlier versions of the Al-Ti phase diagram, which essentially differs from those accepted at present. So, the earliest studies of the Ti rich alloys did not take into account the Ti3Al based phase 2 [1958Boe, 1962Ere, 1962Ge, 1963Ge1, 1963Ge2, 1963Ge3]. Phase equilibria in the Ti rich alloys involving the 2 phase have been studied by [1969Cro, 1969Kor, 1969Nar, 1969Fed, 1978Ban, 1980Ban1, 1980Ban2]. However, the phase equilibria with phase at temperatures higher than ~1200°C were not determined by [1958Boe, 1963Ge1, 1963Ge2, 1963Ge3, 1969Cro, 1969Kor, 1969Fed, 1978Ban, 1980Ban2, 1987Ere]. Phase equilibria involving phases, which later were found to exist between TiAl and TiAl3 [1990Sch, 2001Bra], have not been considered by [1962Ge, 1970Han, 1987Ere] and could not be shown in the isothermal sections at 1300 and 900°C given by [1993Bud] in those days. The following information on the Al-Mo-Ti system was available before 1990: an existence of a wide region of bcc solid solutions and ordering of bcc solid solution that resulted in formation of the 2 CsCl type phase with a composition of Ti2MoAl; significant solubilities of third component in some binary phases, e.g. up to 20 at.% Mo in TiAl3 (J1), 8 at.% Mo in TiAl ( ), more than 20 at.% in Mo3Al (') and small solubilities of Mo in Ti3Al (2) and (Ti). The ternary ) phase Ti1.5Mo1.5Al2 was found in alloys annealed at 925°C for a week [1970Han]; the four-phase invariant equilibrium +2+2 was suggested to exist at 550°C [1972Ham, 1973Ham, 1975Ham]. The partial ternary phase diagram in the Ti rich corner presented by [1993Bud] has taken into account the coexistence of these four phases. A similar version of the phase equilibria was used by [1981Tre] to describe phase transformations in Ti rich alloys. The appearance of the 2++ phase region [1980Ban2] has been discussed taking into account a version of the Al-Ti phase diagram in which the 2 phase exists at high temperatures up to the melt. The isothermal sections at 1600°C [1988Ere1] and 1300°C [1987Ere] were determined. An additional investigation of the alloys resulted in an refining of some elements of the isothermal section at 1300°C, the construction of the isothermal section at 1000°C, a preliminary version of the solidus surface and a reaction scheme in the Al-Mo-Ti system up to 75 at.% Al. The reaction scheme takes into account new information on the binary systems Al-Ti [1996Tre1] and Al-Mo [1991Sch]. The study was made using optical microscopy (OM), later also electron microprobe (EMPA), X-ray diffraction (XRD) and differential thermal (DTA) analyses [1988Ere2, 1990Ere, 1996Tre2]. The phase equilibria in the Al rich region of the Al-Mo-Ti system (> 65 at.% Al) have been studied by [1994Sok] using OM and XRD and published as a partial isothermal section at 500°C. Crystallization of the (Ti1-xMox)Al3 aluminides from dilute melts containing less than 0.5 at.% (Ti+Mo) was studied by [1990Abd], who cooled very slowly from 1000 down to 700°C and then let the samples cool down to room temperature inside a furnace. Most of the investigations performed after the review by [1993Bud] concerned phase transformations and microstructures of alloys adjacent to the Ti-Al side of the ternary phase diagram. The alloys based on Ti3Al were studied by [1991Dja, 1992Dja1, 1992Dja2]. The alloys have been prepared by arc melting and, after various heat treatments, were studied by means of OM, transmission electron microscopy (TEM), scanning electron microscopy (SEM), selected area diffraction (SAD), anomalous small-angle X-ray scattering (ASAXS). Mechanical properties were determined as well. The continuous cooling transformation diagrams, from 1100°C down to room temperature were determined for different cooling rates and the phase and structure transformations have been analyzed. The (2), 7at, 2´, 2 phases were observed. Numerous investigations of AlTi based alloys have been carried out to obtain an information useful in development of titanium aluminide alloys with improved mechanical properties and structural stability. Such alloys have been studied using OM, XRD, and EMPA of arc melted, annealed at 1300°C for 5 h,
Landolt-Börnstein New Series IV/11A3
MSIT ®
288
Al–Mo–Ti
1200°C for 48 h, 1100°C for 120 h and quenched alloys. Moreover in situ XRD at temperatures up to 1400°C [1992Kim] and studies of diffusion couples [1993Has, 1998Kim] have been made. The partial isothermal sections were calculated [1998Kim] using the ThermoCalc program. The isothermal sections were published for the region of the ++ phase field at 1200 and 1300°C [1993Has, 1998Has, 1998Kim]. Earlier the thermodynamic calculations together with experimental studies of phase boundaries of the + region were performed by [1975Zan, 1977Zan, 1986Gro, 1988Gro]. High temperature phase equilibria were studied by [1993Das1, 1993Das2] using OM, SEM, XRD, EMPA, DTA and TEM of the Ti-50Al-5Mo and Ti-45Al-3Mo alloys; here and further compositions of alloys and phases are given in at.%, if not stated differently. The location of the 2+ and +2+ phase fields at 1175°C were determined using EMPA of the above alloys annealed at 1300°C for 3 d and then at 1175°C for 6 d. The microstructures of the Ti-48Al alloys containing 0.5 or 2 at.% Mo were studied as cast (plasma melted) and quenched from temperatures between 1000 and 1350°C, by OM, SEM and TEM [1993Li]. Crystallographic analysis of the solidification microstructure of the Ti-48Al-2Mo alloy was used to investigate high-temperature phase equilibria by OM, SEM, TEM, EMPA [1995Nak]. The mechanism of phase transformations of the phase was studied on continuous cooling experiments. The 2+2+ alloy, Ti-44Al-2Mo, prepared by plasma melting was studied by OM, TEM, SEM and mechanical testing [1994Li, 1994Mor] on samples as cast, as HIPped (Hot Isostatic Pressed) at 1250°C, 150 MPa and as heat-treated at 1200 and 900°C for 120 and 500 h. 12 alloys containing 44 to 50 at.% Al and 2 to 6 at.% Mo were studied as cast and annealed in the temperature range 1100 to 1400°C by means of TEM, XRD and EMPA [1997Sin1, 1997Sin2]. Solidification paths and postsolidification transformations were analyzed. Phases present after heat treatments were determined and their compositions established. Partial Ti-rich isothermal sections at 1400, 1300 and 1200 - 1100°C were developed and projections of the liquidus and solidus surfaces involving , , and L phases were proposed. A projection of the partial liquidus surface near Al-Ti side was constructed from microstructural analysis of arc melted ingots of Ti alloys containing 45 to 60 at.% Al and 2 to 7.5 at.% Mo using OM, SEM [1998Joh]. The experimental data for the liquidus surface have been employed to calculate thermodynamically a solidification path. There is a calculated partial isothermal section at 1500°C and a discussion on directional solidification in the literature. The partial liquidus surface in the regions of primary solidification of the and phases and directional solidification of alloys have been analyzed by [2002Jun] too. Two- and three-phase equilibria involving , 2, (2) and phases have been studied by [2000Kai] who arc melted alloys, annealed them at 1000°C for 168 or 504 h, at 1200°C for 168 h and at 1300°C for 24 h and characterized them by OM and EMPA. So partial phase diagrams at 1000, 1200 and 1300°C were established. Similar phase relations were addressed by [1998Tak]. A detailed study of the Ti-50Al-15Mo alloy was made by [1997Che] using OM, XRD, SEM, TEM and EMPA. The alloy was plasma arc melt and annealed at 1400°C for 1.5 h and at 1350°C for 2 h. The latter samples were annealed additionally at 1200, 1000 or 800°C for 96, 144 and 504 h respectively and water quenched after each of the heat treatments. The resulting phases, their compositions and crystal structure were determined. In addition to the well known phases /2 (hcp), /2 (bcc/B2), (L10) and the phase with D022 structure on the base of TiAl3, three new phases were reported and designated as L60, ´ and ´´. The results by [1997Che] were used in the review by [1999Flo]. In Ti-(5.5-15)Mo-(2-7)Al (mass%) alloys, which were quenched from 1000°C, [1972Luz] studied transformations during aging at 200 to 500°C and examined the influence of these transformations on the mechanical properties. [1980Sas] studied the crystal structures of martensites in Ti-(0-17)Mo-3Al (mass%) alloys quenched from 1000°C. The Ti-(0-30)Mo-3Al (mass%) were researched by [1971Kho] with respect to the transformation temperature and the mechanical properties. Mechanical properties were also studied by [1975Hid] together with the structure of the Ti-7Mo-(16,19)Al alloys, quenched from 960°C and aged at 600 and 400°C. Physical properties and phase transformations were studied for the Ti3Al-1 % Mo alloy by [1976Zel]. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti
289
Site substitution behavior of Ti3Al and TiAl was calculated theoretically [1990Nan, 1993Rub, 1998Woo, 2000Yan, 2001Kan] and determined experimentally by [1999Hao] using the atom location channelling enhanced microanalysis (ALCHEMI). The site occupancy of the alloying element in the and 2 phases was used to estimate +2 phase equilibrium and /2 and 2/ phase boundaries [1999Yan, 2000Yan, 2001Kan]. The sublattice occupancy in B2 phases in the ALCHEMI experiments was analyzed by [1995Che]. The local atomic order in the Ti2MoAl phase was determined from the EXAFS (Extended X-Ray Absorption Fine Structure) study which revealed that this alloy has a pseudo-B2 structure, in which Mo and Al atoms occupy one sublattice and Ti atoms the other one [1996Sik]. The relative stability of different structures in the Ti50Mo25Al25 alloy was calculated theoretically by [2000Alo]. The stability of the aluminides Ti3Al, TiAl and the B2 phase in Ti2MoAl base alloys, has been considered by [1992Nak, 1997Nak] as an information, which can be useful in developing Al-Mo-Ti based materials for structural applications. Binary Systems The accepted Al-Mo system assessed by [2003Sch2] is based on the data of [1971Rex] for the Mo-Mo3Al8 region and on the results of [1991Sch] for the Al rich part. The Al-Ti phase diagram is accepted from the assessment of [2003Sch1], who has proposed a version based on the results by [1992Kat, 1997Zha]. The TiAl-TiAl3 region shown by [1992Kat] summarizes complicated phase relations in this concentration range as shown by [1990Sch] and recently reinvestigated by [2001Bra]. The data by [1996Tre1] are in good agreement with the results of recent studies, particularly as for the Ti5Al11 phase. The Mo-Ti system is accepted as described by [Mas2]. Solid Phases Data on solid phases observed in the ternary and relative binary systems are given in Table 1. The bcc solid solutions existing in a wide range of compositions are the high temperature phase at the Al-Ti side of the ternary phase diagram. They undergo a number of phase transformations as the temperature decreases, 2, giving rise to a variety of microstructures depending on the temperature of the heat treatment and on the cooling rate. Molybdenum is a strong stabilizer and its addition stabilizes the bcc structure down to the room temperature [1991Dja]. Ordering of bcc solid solutions to ordered CsCl type phase (2) was discovered by [1958Boe] and confirmed in the works of [1972Ham, 1991Dja, 1992Dja1, 1992Dja2, 1993Das1, 1993Das2, 1993Li, 1994Li, 1994Mor, 1995Nak, 1997Che, 1997Sin2]. Ordering takes place in a wide range of compositions. The temperature of ordering depends on the composition of the phase and is supposed to be the highest at ~1400°C, for the composition Ti2MoAl. An XRD study often is unable to recognize the ordered 2 phase owing to very weak superstructure reflections. Therefore an electron diffraction analysis was used to identify the 2 phase [1993Das1]. The ternary ) phase detected by [1970Han] was confirmed by [1988Ere2, 1990Ere, 1996Tre2]. This ) phase forms through a peritectoid reaction at ~1250°C. The wide homogeneity range of the J phase based on the binary TiAl3 compound earlier found by [1970Han] was confirmed by [1987Ere, 1990Abd, 1990Ere, 1996Tre2]. The homogeneity range of TiAl3, which is not more than ~1 at.% in the binary Al-Ti system, was found to extend up to ~22 at.% Mo at 75 at.% Al and up to ~16 at.% Mo along the 25 at.% Ti isopleths. The substitution of both Ti and Al atoms by Mo atoms results in decreasing lattice parameters of the J phase. The c/a ratio decreases insignificantly, from 2.234 for TiAl3 to 2.214 for Ti3Mo22Al75, but the substitution of Al by Mo makes the c/a ratio decrease to ~2.12. The Mo solubility in the TiAl based phase increases with increasing Al content and reaches ~9 at.% at ~60 at.% Al. The lattice parameters of the phase were observed to decrease with c/a ratio increasing from 1.015 to ~1.035 with increasing Mo content [1990Ere, 1996Tre3]. [1997Che] observed a 1 phase with D022 type structure in the Ti-50Al-15Mo alloy which was annealed at 1200 - 800°C and water quenched. This work suggests that a transformation of the high temperature (L1 0) phase to the 1 phase takes place, which can not be suppressed. Towards high Mo-contents in the phase
Landolt-Börnstein New Series IV/11A3
MSIT ®
290
Al–Mo–Ti
region [1990Ere, 1996Tre3] observed a similar phase in alloys as cast and such annealed at 1300 and 1000°C. The crystal structure of the ´ phase was supposed to be characterized by Mo and excess-Al ordering in Ti layers in TiAl [1997Che]. The Mo solubilities in and phases were found to be small, about ~2 at.% [1990Ere]. The ordered ´´ phase was observed by [1997Che] in the Ti-50Al-15Mo alloy after prolonged aging at 800°C. The crystal structure of the ´´ phase was found to be similar to that of TiAl3 (D022) and different only by its sublattice. The proposed model of the ´´ phase is consistent with the chemical formula of (Ti, Mo)3Al5. The ´´ phase was suggested to form from 2 (B2) phase or between 2 (B2) and ´ (D022). The martensite phases ´, ´´ were observed in an alloy close to the Mo-Ti side of the ternary phase diagram [1980Sas]. The metastable 7 phase was reported by [1971Wil, 1972Ham, 1972Luz, 1980Sas] and also observed in the research of [1991Dja, 1992Dja1, 1992Dja2] in Al-Ti base alloys of ~20 - 25 at.% Al and 3 - 4 at.% Mo, where also the 2´ martensite phase was observed which is based on Ti3Al. Additions of Al to Mo-Ti alloys were found to suppress the formation of the 7 phase [1972Luz]. Invariant Equilibria The reaction scheme shown in Fig. 1a is based on results obtained by [1990Ere, 1996Tre2] mainly for the Ti-TiAl3-MoAl3-Mo region. Temperatures of phase transformations were determined by DTA. Because of the large losses of Al during heating at temperatures above ~1600°C, even for the time of an DTA experiment, the temperature of the invariant equilibrium '+L+2 was developed from the Al-Mo binary data and from temperatures determined on alloys of the nearest regions. As phase transformations in alloys along the Al-Mo side could not be suppressed during cooling, only the phases existing at lower temperatures have been observed. So, the equilibria involving the 1 and 2 phases were concluded to exist tentatively from the analysis of results obtained from DTA, XRD and OM in as cast and annealed alloys. The reactions in the region between the and J phase fields is shown simplified because phase relations between the phase (L10) and ´ (D0 22) are not determined. The + ´+J (or +J+ / ´), +J+), +'+) and J+'+) phase fields were found to exist at 1000°C [1990Ere], but the + +) and +J+) phase fields were observed at 925°C by [1970Han]. So, the invariant equilibria +J)+ ´ and + ´ +) (or summarized as +J / ´+)) were supposed to exist at temperatures in the range of 925-1000°C. According to [1972Ham, 1975Ham] the invariant equilibrium +2+2 exists at 550°C in the Ti rich region of the ternary system. However, the new version of the binary Mo-Ti phase diagram with a monotectoid reaction +´ existing at 675°C will lead to a three-phase region ++´ in the ternary system. It can be supposed that at lower temperature this three phase region and the +2+ one will give rise to the invariant four-phase equilibrium of +2+´ rather than that proposed by [1972Ham, 1975Ham]. The phase taking part in this equilibrium may have an ordered B2 crystal structure. Nevertheless, the invariant equilibrium +2+2 suggested by [1972Ham, 1975Ham] takes place but at a temperature between 675 and 850°C, which are the temperatures of the monotectoid reaction and the maximum point of the binodal curve +´ in the binary Mo-Ti system. One of the preceding three-phase equilibria, 2++2 may emerge from a contact of two-phase regions, 2+ (2) and +´ based on the +´ phase field in the Mo-Ti system (one of the phases may have the ordered B2 structure). One of the equilibria succeeding the invariant equilibrium, ++´, must move towards the binary Mo-Ti system down to monotectoid line ´ at 675°C. The eutectoid reaction +2+´ may be considered as one more version of the invariant phase equilibrium in the Ti rich region of the ternary system. Based on the data for the binary systems Al-Mo [1991Sch] and Al-Ti [2003Sch1] a tentative reaction scheme for the Al rich region of the Al-Mo-Ti system is shown in Fig. 1b. Liquidus and Solidus Surfaces The solidus surface projected on the Ti-TiAl3-MoAl3-Mo region of the ternary system is shown in Fig. 2. It mainly results from [1990Ere] and integrates additional data for the binary Al-Mo system by [1991Sch], which were reported also by [1997Smi] and accepted by [2003Sch2]. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti
291
The liquidus surface has been determined only near the Al-Ti side. The projections of the boundaries separating the fields of the primary crystallization of , and phases were constructed by [1997Sin1, 1998Joh]. An increasing stability of the phase or both the and phases with Mo addition was found. It might be supposed that maxima on the boundary liquid curves corresponding to invariant three-phase reactions L+ and L+ do exist. The boundary / liquid curve was established also by [2002Jun] from analyzing the dendrite morphology of directionally solidified alloys. There is a good agreement between the results obtained in the above studies. The partial liquidus surface projection is shown in Fig. 3 with a maximum on the curve of the liquid involving in the reaction L+. Earlier the liquidus surface was calculated by [1982Dan] using subregular solution approximation. The calculation was performed without taking into account the existence of several Al-Ti and Al-Mo binary phases. Isothermal Sections The isothermal section at 1600°C is shown in Fig. 4 [1988Ere1, 1988Ere2, 1990Ere, 1996Tre2]. Figure 5 shows the estimated partial section at 1500°C [1998Joh]. The tentative partial isothermal section at 1400°C is given by Fig. 6 [1997Sin2]. The section was constructed from a study of 12 alloys annealed at 1400°C for 1 h and quenched. Earlier the phase equilibria at 1400°C in the Ti rich region (Ti content > 50 mass%) were reported by [1980Ban1], who has obtained similar results. Some discrepancies in phase boundaries can be attributed to a different purity of alloys. The phase equilibria at 1300°C are shown in Fig. 7 [1990Ere, 1996Tre2] and those between the , and phases have been reported also by [1980Ban1, 1993Has, 1998Kim, 2000Kai]. A good agreement is observed between obtained results. Phase equilibria in the region between the and J phase were not ascertained definitely. The Mo solubility in (Ti5Al11) was found to be not more than ~1 at.%. The +J+ phase field was found to exist in a narrow range at ~2 at.% Mo. The +J equilibrium existing at higher Mo contents was observed to be replaced by J being in equilibrium with another phase. The crystal structure of this phase seems to be the same as that of the J phase, the D022 type, but with the c/a ratio close to 1.05, for a sublattice. A similar phase was observed by [1997Che]. The phase relations involving this phase designated as ´ were not firmly established and they are shown in Fig. 7 tentatively. The phase equilibria between the , and phases at 1200°C were presented by [1980Ban1, 1993Has, 1997Sin2, 1998Has, 1998Kim, 2000Kai]. [1998Has, 1998Kim] have attempted to assess experimental results by means of thermodynamic calculation. The partial phase diagram obtained for this region is shown in Fig. 8 [2000Kai]. Similar phase diagrams were presented by [1993Has, 1998Has, 1998Kim] but another location of apices of the ++ phase triangle was proposed by [1997Sin2], especially for the and phases. An ordered phase has not been detected by [1980Ban1, 1987Ere, 1990Ere], while the more recent works have shown the ordered modification 2 of bcc solid solution [1992Kim, 1993Has, 1994Mor, 1995Nak, 1997Che, 1997Sin2, 1998Kim, 1998Tak]. The location of the three-phase +2+ triangle at 1175°C established by [1993Das1, 1993Das2] is consistent with that presented by [1993Has, 1998Has, 1998Kim, 2000Kai] for 1200°C. The phase equilibria in the Ti rich alloys, i.e. with Ti content > 50 mass% have been presented by [1963Ge1, 1963Ge2, 1980Ban2, 1997Sin2]. The 2 phase was not identified in the earlier works, so the phase instead of 2 was shown to coexist with the phase [1963Ge1, 1963Ge2]. The and 2 phases have not been separated by [1980Ban2]. [1997Sin2] has reported that the Ti based phase which coexists with the and 2+ phases is an ordered 2 phase. A good agreement is observed as to the phase composition of the studied alloys and compositions of the (2) and phases but there is a great difference between the composition of the phase reported by [1997Sin2] and that shown by [1963Ge1, 1963Ge2, 1980Ban2]. Figure 9 shows the partial isothermal section at 1100°C developed mainly from that shown by [1963Ge1, 1963Ge2] and the accepted in this evaluation Al-Ti binary system. The data by [1963Ge1, 1963Ge2] were preferred because a large number of alloys annealed at 1100°C for 100 h and water quenched were investigated, while [1997Sin2] studied 12 alloys in the narrow composition range (44 to 50 at.% Al, 2 to Landolt-Börnstein New Series IV/11A3
MSIT ®
292
Al–Mo–Ti
6 at.% Mo) and annealed only 6 h at 1100°C and [1980Ban2] examined only 4 alloys of the same composition range. The phase equilibria at 1000°C are shown in Fig. 10 [1988Ere2, 1990Ere, 1996Tre2]. A main peculiarity of these phase equilibria is the ternary ) phase with a composition close to Ti3Mo3Al4. In the region between the and J phase fields, a coexistence of the J phase and another phase with the D0 22 type crystal structure ( ´) was observed at a higher contents of molybdenum in the alloys. Phase relations between this ´ phase and the and phase were not established. The coordinates of the 2+/2+ phase field agree well with those determined by [2000Kai] as well as with the data by [1997Che] for the /2 and phases in the Ti-50Al-15Mo alloy. The region of the ordered 2 phase is shown mainly by [1958Boe], whose data are in good agreement with those of [1993Das1, 1994Li, 1994Mor, 1997Sin2, 1997Che, 1997Nak]. The phase equilibria at 925°C are shown in Fig. 11 mainly from [1970Han] with corrections due to recent data on the binary Al-Ti and Al-Mo and the ternary systems. The phase relations involving the ) phase, which was discovered by [1970Han], are distinguished from those found at 1000°C. So, the invariant reaction /2+J +) is supposed to take place at a temperature between 1000 and 925°C. The existence of an +2 phase field seems to be hardly probable as the ordering transformation 2 is believed to be of second order. The region of the 2 phase is shown tentatively, the two-phase +2 phase field is omitted in Fig. 11. At lower temperatures the two-phase +2 field would be possible, if attributed to a miscibility gap. The phase equilibria at 800°C have been presented in the Ti rich part of the phase diagram by [1963Ge1, 1963Ge2]. The structure of alloys in the region of Ti-Al-(Ti ~30Mo) annealed at 800°C for 200 to 220 h have been investigated by [1990Ere, 1996Tre2]. The partial section at 800°C shown in Fig. 12 was constructed from the above works and information reported by [1958Boe, 1963Luz, 1972Ham, 1978Ban]. [1997Che] has reported the equilibria of the Ti-50Al-15Mo alloy annealed at 1350°C for 2 h, then at 800°C for 504 h and water cooled. The alloy was found to consist of the 2+ ´+ ´´ phases, an information which is not consistent with the isothermal section at 925°C shown above, because the equilibrium 2+ ´(TiAl)+ ´´(TiAl3) excludes the existence of the + +) phase field, which was found earlier by [1970Han] at 925°C. The phase equilibria in the Ti rich corner at 700 and 600°C are similar to those at 800°C as it is shown in Figs. 13 and 14 which incorporate compatibly data from [1972Ham, 1962Ge, 1963Ge1, 1963Ge2] respectively, the binary Al-Ti and Mo-Ti phase diagrams, data by [1958Boe] for the /2 boundary and data by [1990Ere]. [1994Sok] studied the part of the system and published a partial isothermal section at 500°C for the Al rich region; the Mo solubility in TiAl3 were found to be only 2 at.%; TiAl3 was found to coexist with MoAl3, MoAl5 and MoAl12. The Ti solubilities in above aluminides were reported to be 2, 4 and 2 at.%, respectively, but it is unknown what modification of MoAl3 was implied, no information on crystal structures of the phases was reported. The presented data are not consistent with the data by [1990Abd], who obtained the (Ti1-xMox)Al 3 aluminides during very slow cooling, which allowed the equilibrium phase to crystallize from the Al melt containing ~0.5 at.% (Ti+Mo). The (Ti1-xMox)Al3 aluminides with the TiAl3 type structure were obtained up to x = 0.47 (~12 at.% Mo). The phase composition of J+(Al) was found for alloys in the Al-MoAl3-TiAl3 region almost up to x = 0.7 (the alloys were annealed at 600°C for 44 h, solidus temperatures of these alloys were determined to be 650°C) [1990Ere]. Thermodynamics Evaluated thermodynamic parameters used to asses isothermal sections in Al-Mo-Ti system by [1998Kim]. For modelling of individual phases the sublattice concept was applied. The calculated energy of formation and the chemical potentials of elements, including that of Mo, in (TiAl) are given by [1998Woo] for low temperatures and stoichiometric compositions. The energy of formation for the A2 and B2 phases in the Ti50Al25Mo25 composition was evaluated by [2000Alo].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti
293
Notes on Materials Properties and Applications Ti alloys of the Al-Mo-Ti system are characterized by a variety of phase transformations, which can take place during different heat treatments. Depending on a composition and a heat treatment, both equilibrium and metastable phases can occur and result in various microstructures having an influence upon their properties. Mechanical properties of Ti based Al-Mo-Ti alloys have been studied depending on a composition and heat treatments in earlier works. Composition - hardness relations of Ti rich alloys (Ti > 50 mass%) quenched from various temperatures have been determined by [1963Ge1]. The maximum hardness (HV = 400 to 500 kg#mm-2)has been observed for + and + +2 alloys, solid solutions exhibited the minimum hardness, (HV = 250 kg#mm-2), at ~15 mass% Al near Ti3Al. High temperature hardness of Ti-1Mo-(5 to 20)Al (mass%) has been determined by [1962Ge]. [1971Kho] studied the influence of a thermomechanical treatment on tensile properties of Ti-3Al-(0 to 30)Mo (mass%) alloys, the maximum strengthening was obtained for the Ti-3Al-15Mo alloy. [1972Luz] determined mechanical properties by tensile tests and observed that additions of Mo gave rise to increasing strength in quenched alloys and suppressed the formation of the 7 phase. [1973Ham, 1975Hid] again observed a correlation between microstructures and mechanical properties of Ti-(7-19)Al-7Mo. A study of Ti3Al based alloys containing up to 32 mass% Mo was made by [1969Kor]. The TiAl and Ti3Al aluminides were a subject of recent investigations because they were found to have potential use for high temperature applications in aerospace engines. These aluminides combine low density, high specific strength, good resistance to oxidation, but they have low ductility at room temperature. Molybdenum was found to be an alloying addition, which can have a favorable influence on the properties of intermetallic alloys based on the Ti aluminides. [1991Mae] has found that Ti rich TiAl modified by Mo exhibited higher tensile ductility at room temperature and improved creep strength. Room temperature tensile tests have been carried out also by [1994Li, 1994Mor]. The high strength obtained at room temperature for TiAl based alloys has been attributed to the presence of the ordered 2 phase. Hardness measurements were carried out on the individual phases. The hardness values were measured to be H2 = 394 15 kg#mm-2, H2 = 430 20 kg#mm-2 and H = 273 10 kg#mm-2. Also )0.2 the stress values at 0.2 % strain, the ductility J, the maximum flow stress )max were measured from tensile tests [1994Mor]. Mechanical properties of TiAl based alloy at temperatures ranging from 77 to 1473 K were examined by [1993Has]. The mechanical properties of TiAl can be greatly improved by control of microstructure and morphology of secondary phases, which can be changed with Mo additions affecting the stability of the phases. Tensile properties of Ti3Al based alloys with Mo at room temperature have been examined on samples thermomechanically processed (TMP) and heat treated (HT) [1992Dja2]. It has been shown that the tensile properties of Al-Mo-Ti aluminides may be optimized by specific TMP and HT. Electrical conductivity and a coefficient of thermal expansion in the temperature range from 20 to 1000°C, hardness at 20 to 800°C, a modulus of elasticity and internal friction were measured on the Ti3Al-1 mass% Mo alloy by [1976Zel]. An abrupt change of physical properties with a heat absorption has been observed at 1080°C. Calorimetric studies of superconducting (Ti0.75 Mo0.25)1-xAlx alloys with x = 0 to 0.06 have revealed that the superconducting transition temperature Tc decreases linearly from 3.9 0.1 K at x = 0 with a rate of approximately 0.3 K per at.% Al [1985Ho]. Miscellaneous Mo atoms tend to Ti sites [1999Hao, 2000Yan] in Ti3Al alloys. The Mo atoms were shown to occupy both sublattices in TiAl [1990Nan, 1998Woo, 2000Yan, 2001Kan] and show different site preference of Mo in TiAl alloys than in Ti3Al [1999Hao]. The formation of Ti3Al phase was shown to obey the electron concentration rule. The experimental boundary of the 2 phase was found to agree with that calculated using an electron model with N = 2.12 [1984Li]. [1995Che] studied two alloys, Ti-42Al-7.5Mo and Ti-50Al-15Mo, which were annealed at 1350°C for 2 h and WQ, then the latter alloy was annealed at 800°C for 504 h and WQ. Compositions of three 2 phases Landolt-Börnstein New Series IV/11A3
MSIT ®
294
Al–Mo–Ti
of these alloys were determined (the first alloy was single phase 2) and sublattice occupancies were established using ALCHEMI. The 2 phases were found to be Ti-48Al-15Mo, Ti-41.6Al-7.3Mo and Ti-37.1Al-22.2Mo. The first two ones were described as (Ti,Mo)52Al48, Ti51(Al,Mo)49. In the third 2 phase, Mo was suggested to be distributed randomly on both sublattices, (Ti,Mo)50(Al,Mo)50. In all cases 2 contained more Al than Ti2AlMo [1958Boe]. Applying CVM, the cluster variation method [1993Rub] calculated from experimental binary data an isothermal section at 1000°C and found a miscibility gap in the inner part of the section besides of bcc () and B2 (2) fields. A comparison between the energy of formation of the A2 () and B2 (2) phases of the same composition Ti2AlMo calculated from first principles has shown the B2 phase to be more stable than the A2 one [2000Alo]. [1980Sas] reported martensite phases in Ti-3Al-(0 to 17)Mo (mass%) alloys quenched from 1000°C (the field). The crystal structure of the martensite at low Mo-content (4 mass%) was found to be hcp (´), at 7 12 mass% (3.5 - 6.2 at.%) Mo it was orthorhombic (´´). No martensite was observed at Mo contents higher than 13 mass% (6.8 at.%). However, slight deformation caused orthorhombic martensite to occur at 13 to 17 mass% Mo. A distorted bcc phase was observed at 12 mass% Mo. [1971Wil] studied a decomposition of a metastable phase in the alloys Ti-(3, 6)Al-20Mo (mass%) quenched from 1000°C and has found that Mo additions reduced the volume fraction and time of stability of the 7 phase. The influence of Mo additions on the occurrence of the 7 phase in the alloys containing 4 to 8 at.% Mo and 0 to 3 at.% Al was studied by [1993Cui]. It was shown that formation of the 7 phase obeys the electron concentration rule. The -7 boundary was calculated and determined experimentally (at the valence electron number 4.10, from ~4.5 at.% Mo to ~6 at.% Mo at 0 and 3 at.% Al). A formation of an athermal 7 phase (“tweed microstructure”) has been observed in Ti3Al based alloys containing 3.4 and 4.4 at.% Mo quenched from the field [1991Dja, 1992Dja2]. [1991Dja, 1992Dja2] have presented continuous cooling transformation diagrams for Ti3Al based alloys with different Al and Mo contents, which have been annealed in the region and cooled with rates varying from 80 to 0.1°C#s-1. The Ti-50Al-5Mo alloy was found to be single phase at 1400°C and to exhibit during cooling a sequence of phase transformations ++ + +2. The + + phases were found in the alloy annealed at 1240°C for 150 h. The 2+ phase composition was established in the alloy annealed at 1175°C for 6 h. In the Ti-45Al-5Mo alloy, the + structure observed in the alloy annealed at 1300°C for 3 d was found to be changed to 2+ after annealing at 1175°C for 6 h. The 2+ alloys were found to be stable to a high temperature exposure at 1240°C for 150 h, but some modifications took place at longer time. The partitioning tendency of Mo into different phases (, /2 and ) was found to be as follows: > > [2000Kai]. Sintering of elemental powders at 1150°C to obtain a ternary intermetallic compound of the L12 type has resulted in the phase composition D022+(TiAl2) in the Ti-67Al-8Mo alloy [1993Nak]. Sulfidation properties of the TiAl-2Mo alloy at 900°C and 1.3 Pa sulphur pressure have been studied by [2000Izu]. References [1958Boe]
[1962Ere] [1962Ge]
[1963Ge1]
MSIT®
Boehm, H., Loehberg, K., “A Superstructure CsCl Type Phase in the Titanium Molybdenum - Aluminium System” (in German), Z. Metallkd., 49, 173-178 (1958) (Crys. Structure, Experimental, #, 10) Eremenko, V.N., Mnogokomponentnyye Splavy Titana (in Russian), Izd. Akad. Nauk Ukr. SSR, Kiev, 27-29 (1962) (Equi. Diagram, Review, 8) Ge Chzhi Min, Kornilov, I.I., Pylayeva, E.N., “Investigation of Structure and Properties of Alloys of the Titanium - Molybdenum System” (in Russian), Izv. Acad. Nauk SSSR, Otd. Tekh. Nauk, Metall. i Toplivo, (4) 114-118 (1962), translated in Russ. Metallurgy and Fuels, (4) 86-98 (1962) (Equi. Diagram, Experimental, #, 14) Ge Chzhi Min, Kornilov, I.I., Pylayeva, E.N., “Investigation of the Ti-Al-Mo Phase Diagram in the Region of Ti-Rich Alloys” (in Russian), Zh. Neorg. Khim., 8, 366-372 Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti
[1963Ge2]
[1963Ge3]
[1963Luz]
[1969Cro]
[1969Kor]
[1969Nar] [1969Fed]
[1970Han]
[1971Kho]
[1971Rex]
[1971Wil]
[1972Ham]
[1972Kam]
[1972Luz]
[1973Ham]
[1975Ham]
Landolt-Börnstein New Series IV/11A3
295
(1963), translated in Russ. J. Inorg. Chem., 8, 189-193 (1963) (Equi. Diagram, Experimental, #, 7) Ge Chzhi Min, Pylayeva, E.N., “Investigation of a Phase Equilibrium in the Ti-Al-Mo System” (in Russian), “Titan I Yego Splavy”, (10), AN SSSR, Moskva, 14-21 (1963), translated in “Titanium and its Alloys”, 10, 11-18 (1966) (Equi. Diagram, Experimental, #, 7) Ge Chzhi Min, Pylayeva, E.N., “Investigation of Phase Transformation in the Ti-Mo-Al System” (in Russian), “Titan I Yego Splavy”, (10), AN SSSR, Moskva, 22-26 (1963), translated in “Titanium and Its Alloys”, 10, 19-23 (1966) (Equi. Diagram, Experimental, 8) Luzhnikov, L.P., Novikova, V.M., Mareyev, A.P., “Solubility of -Stabilizers in -Ti” (in Russian), Metalloved. Term. Obrab. Met., (2) 13-16 (1963) (Equi. Diagram, Experimental, 4) Crossley, F.A., “Effects of the Ternary Additions: O, Sn, Zr, Cb, Mo, and V on the /+Ti3Al Boundary of Ti-Al Base Alloys”, Trans. Metall. Soc. AIME, 245, 1963-1968 (1969) (Equi. Diagram, Experimental, 15) Kornilov, I.I., Nartova, T.T., Shirokova, N.I., “Structure and Properties of the Ti3Al Aluminide Containing Molybdenum” (in Russian), Metalloved. Term. Obrab. Met., (8) 40-42 (1969) (Equi. Diagram, Experimental, 4) Nartova, T.T., Shirokova, N.I., “Phase Equilibrium in a Part of the Ti-Al-Mo System” (in Russian), Izv. Akad. Nauk SSSR, Met., (6) 163-166 (1969) (Equi. Diagram, Experimental, 9) Fedotov, S.G., Ronami, G.N., Konstantinov, K.M., Kuznetsova, S.M., Sinodova, E.P., Starokozhev, B.S., “Composition of an -Solid Solution in Ternary Alloys of Titanium with Aluminium and Molybdenum or Vanadium” (in Russian), Izv. Akad. Nauk SSSR, Met., (6) 167-171 (1969) (Equi. Diagram, Experimental, 7) Hansen, R.C., Raman, A., “Alloy Chemistry of )(U)-Related Phases. III. )-Phases with Non-Transition Elements”, Z. Metallkd., 61, 115-120 (1970) (Crys. Structure, Equi. Diagram, #, 24) Khorev, A.I., Chinenov, A.M., Martynova, M.M., “Mechanical-Thermal Treatment of Alloys of the Ti-Al-Mo System” (in Russian), Metalloved. Term. Obrab. Met., (9) 43-46 (1971) (Equi. Diagram, Experimental, 10) Rexer, J., “Phase Equilibria in the Aluminium - Molybdenum System at Temperatures above 1400°C” (in German), Z. Metallkd., 62, 844-848 (1971) (Crys. Structure, Equi. Diagram, Experimental, 23) Williams, J.C., Hickman, B.S., Leslie, D.H., “The Effect of Ternary Additions on the Decomposition of Metastable Phase Ti Alloys”, Metall. Trans., 2, 477-484 (1971) (Experimental, 20) Hamajima, T., Luetjering, G., Weissman, S., “Microstructure and Phase Relations for Ti-Mo-Al Alloys”, Metall. Trans., 3, 2805-2810 (1972) (Crys. Structure, Equi. Diagram, Experimental, #, 15) Kamei, K., Ninomiya, T., Terauchi, S., “Aluminium - Molybdenum Binary Phase Diagram”, Tech. Rep. Kansai Univ., 13, 93-106 (1972) (Crys. Structure, Equi. Diagram, Experimental, 7) Luzhnikov, L.P., Novikova, V.M., Orlova, I.S., “Transformations during Heat Treatment of Alloys of the Ti-Mo System with Additions of Al, Zr, Sn” (in Russian), Novy Konstr. Mater. Titan, Nauka, Moscow, 41-48 (1972) (Equi. Diagram, Experimental, 3) Hamajima, T., Luetjering, G., Weissman, S., “Importance of Slip Mode for Dispersion-Hardened -Titanium Alloys”, Metall. Trans., 4, 847-856 (1973) (Equi. Diagram, Experimental, 10) Hamajima, T., Weissman, S., “Thermal Equilibria and Mechanical Stability of Ti3Al Phase in Ti-Mo-Al Alloys”, Metall. Trans., 6A, 1535-1539 (1975) (Equi. Diagram, Experimental, 6)
MSIT ®
296 [1975Hid] [1975Zan]
[1976Zel]
[1977Zan]
[1978Ban] [1980Ban1]
[1980Ban2]
[1980Sas]
[1981Kin] [1981Tre]
[1982Dan]
[1984Li]
[1985Ho]
[1986Gro]
[1987Ere]
[1988Ere1]
[1988Ere2]
[1988Gro]
MSIT®
Al–Mo–Ti Hida, M., Weissman, S., “High-Temperature Strength and Ductility Increases in Ti-Mo-Al Alloys by Step Aging”, Metall. Trans., 6A, 1541-1546 (1975) (Experimental, 7) Zangvil, A., Osamura, K., Murakami, Y., “Determination of Phase Equilibrium in the Ti-Rich Ti-Mo-Al Ternary System Using the X-Ray Microanalyzer”, Met. Sci., 9, 27-31 (1975) (Equi. Diagram, Experimental, 15) Zelenkov, I.A., Osokin, E.N., “A Change of Some Physical Properties of the Ti3Al Compound and Hard Alloys on its Base at Phase Transitions” (in Russian), Poroshk. Metall., (2) 44-48 (1976) (Experimental, 12) Zangvil, A., Osamura, K., Murakami, Y., “Determination of Interaction Parameters from EMPA Data in the Ti-Mo-Al Ternary Systems”, Trans. Jpn. Inst. Met., 18, 503-508 (1977) (Equi. Diagram, Theory, Thermodyn., 10) Banerjee, D., Krishnan, R.V., Vasu, K.I., “Transformation Microstructures in a Ti-31Al-13Mo Alloy”, Scr. Metall., 12, 27-30 (1978) (Experimental, 9) Banerjee, D., Arunachalam, V.S., “The 2 Transformation in Ti-Al-Mo Alloys”, “Titanium´80. Science and Technology”, Proc. 4 Int. Conf., Kyoto, N.Y., 4, 2959-2969 (1980) (Experimental, 28) Banerjee, D., Krishnan, D.V., Vasu, K.I., “A Reconsideration of Phase Relations in the Ti-Al-Mo and Ti-Mo Systems”, Metall. Trans., 11A, 1095-1105 (1980) (Crys. Structure, Equi. Diagram, Experimental, 24) Sasano, H., Suzuki, T., Nakano, O., Kimura, H., “Crystal Structures of Martensites in Ti-Mo-Al Alloys”, “Titanium´80. Science and Technology”, Proc. 4 Int. Conf., Kyoto, (1980), N.Y., 4, 717-724 (1980) (Crys. Structure, Experimental, 16) King, H.W., “Crystal Structure of the Elements at 25°C”, Bull. Alloy Phase Diagrams, 2, 401-402, (1981) Crys. Structure, Review, 5) Tretyachenko, L.A., “On the Phase Diagrams of the Ti-Mo-Al System. Boundaries of the + Region” (in Russian), “Vliyaniye Termich. Obrab. na Svoistva Titan. Splavov”, Proc. I Vses. Conf., Dnepropetrovsk, 1980, 113-121 (1981) (Equi. Diagram, Review, 27) Danilenko, V.M., Rubashevsky, A.A., “Calculation of the Liquidus Surface of the Ti-Mo-Al System” (in Russian), Poroshk. Metall., (9) 46-49 (1982) (Equi. Diagram, Thermodyn., Theory, 5) Li, D., Liu, Y., “On the Thermal Stability of Ti Alloys. II. The Behaviour of Transition Elements in Ti3X-Phase Formation” (in Chinese), Acta Metall. Sin.(China), 20, A384-A390 (1984) (Equi. Diagram, Experimental, Theory, 1) Ho, J.C., Majerich, D., Gegel, H.L., “Calorimetric Studies of Superconducting (Ti0.75 Mo0.25)1-xAlx Alloys”, J. Mater. Sci. Lett., 4, 1261-1263 (1985) (Experimental, Thermodyn., 9) Gros, J.P., Ansara, I., Allibert, M., Alheritière, E., “Thermodynamic Study of the Ti-Rich Side of the Ti-Al-Mo System” (in French), Mem. Etud. Sci. Rev. Metall., 83, 448 (1986) (Equi. Diagram, Experimental, Theory, Thermodyn., 1) Eremenko, V.N., Sukhaya, S.A., Tretyachenko, L.A., “Isothermal Section of the Ti-Al-Mo Phase Diagram at 1300°C” (in Russian), Stabil. i Metastabil. Fazy v Mater., IPM, Kiev, 106-114 (1987) (Crys. Structure, Equi. Diagram, Experimental, #, 9) Eremenko, V.N., Sukhaya, S.A., Tretyachenko, L.A., “Phase Equilibria in the Ti-Al-Mo System at 1600°C” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., (4) 97-100 (1988) (Crys. Structure, Equi. Diagram, Experimental, #, 9) Eremenko, V.N., Sukhaya, S.A., Tretyachenko, L.A., Buyanov, Yu.I., “On the Phase Equilibria in the Mo-Ti-Al System at 1600, 1300, 1000°C” (in Russian), VI Vses. Soveshch. po Chim. i Technol. Mo i W, 1988, Nalchik, Abs. Rep., 132 (1988) (Equi. Diagram, Experimental, 0) Gros, J.P., Ansara, I., Allibert, M., “Prediction of / Equilibria in Titanium-Based Alloys Containing Al, Mo, Zr, Cr (Part II), Sixth World Conf. on Titanium, III, Cannes, France, 1559-1564 (1988) (Equi. Diagram, Experimental, Theory, 0) Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti [1990Abd]
[1990Nan] [1990Ere]
[1990Sch]
[1991Dja]
[1991Mae] [1991Sch]
[1992Dja1]
[1992Dja2]
[1992Kat]
[1992Kim]
[1992Nak]
[1993Bud]
[1993Cui] [1993Das1]
[1993Das2]
Landolt-Börnstein New Series IV/11A3
297
Abdel-Hamid, A.A., “Crystallization of Complex Aluminide Compounds from Dilute Al-Ti Metals Containing One or Two Transition Metals of IVB to VIB Groups”, Z. Metallkd., 81, 601-605 (1990) (Equi. Diagram, Experimental, 16) Nandy, T.K., Banerjee, D., Gogia, A.K., “Site Substitution Behaviour of TiAl Intermetallics”, Scr. Metall. Mater., 24, 2019-2022 (1990) (Crys. Structure, Theory, 8) Eremenko, V.N., Tretyachenko, L.A., Sukhaya, S.A., Petukh, V.M., “Investigation of the Structure of Alloys of the Ti-Mo-Al System” (in Russian), “Physico-Chemical Investigation of Binary and Ternary Systems of Transition Metals of IV-VIII Groups of the Periodic System and Development of Principles for Control of Mechanical Properties of Alloys on Their Base (Theme 2.26.30, Final Report, State Regist. No. 01 86 0 060682)”, Akad. Nauk Ukr. SSR, IPM, Kiev, 83-135, 141-143 (1990) (Crys. Structure, Equi. Diagram, Experimental, #, 24) Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”, Z. Metallkd., 81, 389-396 (1990) (Crys. Structure, Equi. Diagram, Experimental, Review, 33) Djanarthany, S., Servant, S., Penelle, R., “Phase Transformations in Ti3Al and Ti3Al+Mo Aluminides”, J. Mater. Res., 6, 969-986 (1991) (Crys. Structure, Equi. Diagram, Experimental, 24) Maeda, T., Okada, M., Shida, Y., “Ductility and Strength in Mo Modified TiAl”, Mat. Res. Soc. Symp. Proc., 213, 556-560 (1991) (Crys. Structure, Experimental, 15) Schuster, J., Ipser, H., “The Al-Al8Mo3 Section of the Binary System AluminiumMolybdenum”, Metall. Trans., 22A, 1729-1736 (1991) (Crys. Structure, Equi. Diagram, Experimental, 20) Djanarthany, S., Servant, S., Lyon, O., “Phase Separation in a Ti-Al-Mo Alloy Studied by Anomalous Small-Angle X-Ray Scattering. A Synchrotron Radiation Experiment”, Philos. Mag., 66A, 575-590 (1992) (Crys. Structure, Equi. Diagram, Experimental, Theory, 14) Djanarthany, S., Servant, S., Penelle, R., “Influence of an Increasing Content of Molybdenum on Phase Transformations of Ti-Al-Mo Aluminides - Relation with Mechanical Properties”, Mater. Sci. Eng., A152, 48-53 (1992) (Equi. Diagram, Experimental, 14) Kattner, U.R., Lin, J.C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the Ti-Al System”, Metall. Trans., 23A, 2081-2090 (1992) (Equi. Diagram, Review, Theory, Thermodyn., 51) Kimura, M., Hashimoto, K., Morikawa, H., “Study on Phase Stability in Ti-Al-X Systems at High Temperatures”, Mater. Sci. Eng., A152, 54-59 (1992) (Crys. Structure, Equi. Diagram, Experimental, #, 12) Naka, S., Thomas, M., Khan, T., “Potential and Prospects of Some Intermetallic Compound for Structural Applications”, Mater. Sci. Technol., 8, 291-298 (1992) (Equi. Diagram, Experimental, Review, 26) Budberg, P., Schmid-Fetzer, R., “Aluminium - Molybdenum - Titanium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.17143.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 27) Cui, Y., Li, D., Wan, X., “7 Phase Formation in Ti Alloys” (in Chinese), Acta Metall. Sin. (China), 29, A61-A67 (1993) (Crys. Structure, Experimental, Theory, 9) Das, S., Mishurda, J.C., Allen, W.P., Perepezko, J.H., Chumbley, L.S., “Development of a ( +0) Lamellar Microstructure in a Ti45Al50Mo5 Alloy”, Scr. Metall. Mater., 28, 489-494 (1993) (Crys. Structure, Equi. Diagram, Experimental, 17) Das, S., Jewett, T.J., Perepezko, J.H., “High Temperature Phase Equilibria of Some Ternary Titanium Aluminides”, in “Structural Intermetallics”, Darolia, R., Lewandowski, J.J., Liu, C.T., Martin, P.L., Miracle, D.B., Nathal, M.V., (Eds.), Min., Met., Mater. Soc., 420
MSIT ®
298
[1993Gam]
[1993Has]
[1993Li] [1993Nak] [1993Oka] [1993Rub]
[1994Kai] [1994Li]
[1994Mor]
[1994Sok]
[1995Che] [1995Gri] [1995Nak]
[1996Sik]
[1996Tre1]
[1996Tre2]
[1996Tre3] [1997Bul]
MSIT®
Al–Mo–Ti Commonwealth Dr., Warrendale, Pens. 15086, 35-43 (1993) (Equi. Diagram, Experimental, Review, 48) Gama, S., “Aluminium - Niobium - Titanium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16070.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 22) Hashimoto, K., Kimura, M., “Effects of Third Element Additions on Mechanical Properties of TiAl”, in “Structural Intermetallics”, Darolia, R., Lewandowski, J.J., Liu, C.T., Martin, P.L., Miracle, D.B., Nathal, M.V., (Eds.), Min., Met., Mater. Soc., 420 Commonwealth Dr., Warrendale, Pens. 15086, 309-318 (1993) (Equi. Diagram, Experimental, 18) Li, Y.G., Loretto, M.H., “Antiphase Boundaries in Ti-48Al2Mo”, Acta Metall. Mater., 41, 3413-3419 (1993) (Equi. Diagram, Experimental, 11) Nakayama, Y., Mabuchi, H., “Formation of Ternary L1 2 Compounds in Al3Ti-Base Alloys”, Intermetallics, 1, 41-48 (1993) (Equi. Diagram, Experimental, 40) Okamoto, H., “Al-Ti (Aluminium - Titanium)”, J. Phase Equilib., 14, 120-121 (1993) (Crys. Structure, Equi. Diagram, Review, 16) Rubin, G., Finel, A., “Calculation of Phase Diagrams of Ternary Systems with Cluster-Variation-Method Entropy”, J. Phys.: Condens. Matter., 5, 9105-9120 (1993) (Equi. Diagram, Theory, Thermodyn., 34) Kainuma, R., Palm, M., Inden, G., “Solid-Phase Equilibria in the Ti-Rich Part of Ti-Al System”, Intermetallics, 2, 321-332, (1994) (Equi. Diagram, Experimental, 35) Li, Y.G., Loretto, M..H., “Microstructure and Fracture Behaviour of Ti-44Al-xM Derivatives”, Acta Metall. Mater., 42, 2913-2919 (1994) (Crys. Structure, Equi. Diagram, Experimental, 12) Morris, M.A., Li, Y.G., Leboeuf, M., “Variation of the Phase Distribution in a Ti-44Al-2Mo Alloy by Annealing: Influence on its Strength and Ductility”, Scr. Metall. Mater., 31, 449-454 (1994) (Crys. Structure, Equi. Diagram, Experimental, 11) Sokolovskaya, E.M., Kazakova, E.F., Poddyakova, E.I., Portnoy, V.K., Temirbayeva, A.A., “Isothermal Section of the Al-Mo-Ti System at 770 K” (in Russian), Vestn. Mosk. Univ., Ser. 2: Khim., 35, 95-97 (1994) (Equi. Diagram, Experimental, 6) Chen, Z., Jones, I.P., “Sublattice Occupancy in Three Ti-Al-Mo B2 Phase”, Scr. Metall. Mater., 32, 553-557 (1994) (Crys. Structure, Experimental, 5) Grin, Yu.N., Ellner, M., “The Crystal Structures of Mo4Al17 and Mo5Al22”, Z. Kristallogr., 210, 96-99 (1995) (Crys. Structure, Experimental, 11) Nakai, K., Ono, T., Ohtsubo, H., Ohmori, Y., “Phase Stability and Decomposition Processes in Ti-Al Based Intermetallics”, Mater. Sci. Eng., A192, 922-929, (1995) (Equi. Diagram, Experimental, 21) Sikora, T., Hug., G., Jaouen, M., Flank, A.-M., “EXAFS Study of the Local Atomic Order in Ti2AlX (X = Nb, Mo) B2 Intermetallic Compounds”, J. Phys. IV, 6, C2-15-C2-20 (1996) (Crys. Structure, Experimental, 8) Tretyachenko, L.A., “On the Ti-Al System”, Phase Diagrams in Material Science, Fifth International School, 1996, Katsyveli, Crimea, Ukraine, 118 (1996) (Equi. Diagram, Experimental, #, 0) Tretyachenko, L.A., “Phase Equilibria in the Ti-Mo-Al System”, Phase Diagrams in Materials Science, Fifth International School, 1996, Katsyveli, Crimea, Ukraine, 119 (1996) (Equi. Diagram, Experimental, 0) Tretyachenko. L.A., unpublished data Bulanova, M., Tretyachenko, L., Golovkova, M., “Phase Equilibria in Ti-Rich Corner of the Ti-Si-Al System”, Z. Metallkd., 88, 256-267 (1997) (Crys. Structure, Equi. Diagram, Experimental, #, 15)
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti [1997Che]
[1997Nak]
[1997Sau] [1997Sin1]
[1997Sin2]
[1997Smi] [1997Zha] [1998Has]
[1998Joh]
[1998Kim]
[1998Tak]
[1998Woo]
[1999Hao]
[1999Flo] [1999Yan]
[2000Alo]
[2000Izu]
[2000Kai]
[2000Oka]
Landolt-Börnstein New Series IV/11A3
299
Chen, Z., Jones, I.P., Small, C.J., “The Structure of the Alloy Ti-50Al-15Mo between 800 and 1400°C”, Acta Mater., 45, 3801-3815 (1997) (Crys. Structure, Equi. Diagram, Experimental, 18) Naka, S., Khan, T., “Designing Novel Multicomponent Intermetallics: Contribution of Modern Alloy Theory in Developing Engineering Materials”, J. Phase Equilib., 18, 635-649 (1997) (Equi. Diagram, Review, 17) Saunders, N., “The Al-Mo System (Aluminium - Molybdenum)”, J. Phase Equilib., 18, 370-376 (1997) (Crys. Structure, Equi. Diagram, Review, Thermodyn., 40) Singh, A.K., Banerjee, D., “Transformations in 2+ Titanium Aluminide Alloys Containing Molybdenum: Part I. Solidification Behavior”, Metall. Mater. Trans., 28A, 1735-1741 (1997) (Equi. Diagram, Experimental, 13) Singh, A.K., Banerjee, D., “Transformations in 2+ Titanium Aluminide Alloys Containing Molybdenum: Part II. Heat Treatment”, Metall. Mater. Trans., 28A, 1745-1753 (1997) (Equi. Diagram, Experimental, 7) Smith, J.F., “Appendix” to [1997Sau], J. Phase Equilib., 18, 376-378 (1997) (Crys. Structure, Equi. Diagram, Review, 1) Zhang, F., Chen, S.L., Chang, Y.A., Kattner, U.R., “A Thermodynamic Description of the Ti-Al System”, Intermetallics, 5, 471-482 (1997) (Equi. Diagram, Theory, Thermodyn., 45) Hashimoto, K., Kimura, M., Mizuhara, Y., “Alloy Design of Gamma Titanium Aluminides Based on Phase Diagrams”, Intermetallics, 6, 667-672 (1998) (Equi. Diagram, Experimental, Theory, 14) Johnson, D.R., Chihara, K., Inui, H., Yamaguchi, M., “Microstructural Control of TiAl-M-B Alloys by Directional Solidification”, Acta Mater., 18, 6529-6540 (1998) (Equi. Diagram, Experimental, Theory, Thermodyn., 33) Kimura, M., Hashimoto, K., “High-Temperature Phase Equilibria in Ti-Al-Mo System”, J. Phase Equilib., 20, 224-230 (1998) (Equi. Diagram, Experimental, Theory, Thermodyn., #, 19) Takeyama, M., Ohmura, Y., Kikuchi, M., Matsuo, T., “Phase Equilibria and Microstructural Control of TiAl Based Alloys”, Intermetallics, 6, 643-646, (1998) (Equi. Diagram, Experimental, Theory, Thermodyn., #, 33) Woodward, C., Kajihara, S., “Site Preferences and Formation Energies of Substitutional Si, Nb, Mo, Ta and W Solid Solutions in L1 0 Ti-Al”, Phys. Rev. B, 57, 13459-13470 (1998) (Crys. Structure, Theory, 45) 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, 1129-1139 (1999) (Crys. Structure, Experimental, 41) Flower, H.M., Christodoulou, J., “Phase Equilibria and Transformations in Titanium Aluminides”, Mater. Sci. Technol., 15, 45-52 (Crys. Structure, Equi. Diagram, Review, 46) Yang, R., Hao, Y.L., “Estimation of ( +2) Equilibrium in Two Phase Ti-Al-X Alloys by Means of Sublattice Site Occupancies of X in TiAl and Ti3Al”, Scr. Mater., 41, 341-346 (1999) (Equi. Diagram, Theory, 13) Alonso, P.R., Rubiolo, G.H., “Relative Stability of bcc Structures in Ternary Alloys with Ti 50Al25Mo25 Composition”, Phys. Rev. B, 62, 237-242 (2000) (Crys. Structure, Equi. Diagram, Theory, 19) Izumi, T., Yoshika, T., Hayashi, S., Narita, T., “Sulfidation Properties of TiAl-2 at.% X (X = V, Fe, Co, Cu, Mo, Nb, Ag and W) Alloys at 1173 K and 1.3 Pa Sulfur Pressure in an H 2S-H2 Gas Mixture”, Intermetallics, 8, 891-901 (2000) (Experimental, 42) Kainuma, R., Fujita, Y., Mitsui, H., Ohnuma, I., Ishida, K., “Phase Equilibria Around (hcp), (bcc) and (L10) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855-867 (2000) (Equi. Diagram, Experimental, #, 29) Okamoto, H., “Al - Ti (Aluminium - Titanium)”, J. Phase Equilib., 21, 311 (2000) (Equi. Diagram, Review, 2) MSIT ®
Al–Mo–Ti
300 [2000Yan]
[2001Bra]
[2001Kan]
[2002Jun]
[2003Kar]
[2003Sch1]
[2003Sch2]
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, 296-301 (2000) (Crys. Structure, Equi. Diagram, Review, 38) Braun, J., Ellner, M., “Phase Equilibria Investigations on the Aluminium-Rich Part of the Binary System Ti-Al”, Metall. Mater. Trans., 32A, 1037-1047 (2001) (Crys. Structure, Equi. Diagram, Experimental, Review, 34) Kang, S.-Y., Onodera, H., “Analyses of HCP/D0 19 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) (Equi. Diagram, Theory, 15) Jung, I.S., Jang, H.S., Oh, M.H., Lee, J.H., Wee, D.H., “Microstructure Control of TiAl Alloys Containing Stabilizers by Directional Solidification”, Mater. Sci. Eng., A329-331, 13-18 (2002) (Equi. Diagram, Experimental, 19) Karpets, M.V., Milman, Yu.V., Barabash, O.M., Korzhova, N.P., Senkov, O.N., Miracle, D.B., Legkaya, T.N., Voskoboynik, I.V., “The Influence of Zr Alloying on the Structure and Properties of Al 3Ti”, Intermetallics, 11, 241-249 (2003) (Crys. Structure, Experimental, 16) Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 85) Schuster, J.C., “Al-Mo (Aluminium - Molybdenum)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 61)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 664 < 660.452 , (Ti1-x-yMoxAly) (Ti)(h) 1670 - 882
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cI2 Im3m W
a = 404.96 a = 330.65
a = 314.70 a = 314.2 a = 317.8
(Mo) < 2623 * 2
Lattice Parameters Comments/References [pm]
cP2 Pm3m CsCl
a = 321
a = 320.1 a = 321.5
MSIT®
0 to 0.6 at.% Ti [1992Kat, 2003Sch1] 0 to < 0.01 or 0.03 at.% Mo [2003Sch2] pure Al at 25°C [1981Kin, Mas2] 0 x 1 [Mas2] pure Ti at 900°C; dissolves up to 44.8 at.% Al at x = 0 [1992Kat, 1993Oka, 2003Sch1] dissolves up to 20.5 at.% Al pure Mo [1981Kin, Mas2] for Mo - ~20 at.% Al [1972Kam] in Ti-50Al-5Mo annealed at 1240°C for 150 h ( +2) [1993Das1] ordered form of bcc (Ti,Mo,Al) solid solution [1958Boe, 1972Ham, 1975Ham, 1991Dja, 1992Dja1, 1992Dja2, 1992Nak, 1993Das1, 1993Das2, 1994Li, 1994Mor, 1995Che, 1996Sik, 1997Che, 1997Nak, 1997Sin2, 1998Joh] in the Ti-44Al-2Mo alloy (2+2+ ) HIPped at 1250°C, 150 MPa for 4 h [1994Li] in as HIPped Ti-44Al-2Mo alloy [1994Mor] in the Ti-44Al-2Mo alloy annealed at 1200°C and 900°C [1994Mor]
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti Phase/ Temperature Range [°C] , (Ti1-x-yMoxAly)
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype hP2 P6 3/mmc Mg
a = 295.03 c = 468.36 a = 294.9 c = 467.6 a = 293.1 c = 464.3
(Ti)(r) < 882
', Mo3Al 2150
cP8 Pm3n Cr3Si a = 495 a = 496 a = 497
a = 498.7 a = 498.7 2, MoAl(h) ~1750 - 1470
1, Mo 37Al63(h) 1570 - 1490
Landolt-Börnstein New Series IV/11A3
cP2 Pm3m CsCl cI2 Im3m W
301 Comments/References
47.3 to 51.4 at.% Al at x = 0 at solidus temperatures 1490 to 1463°C [1992Kat, 1997Zha, 2003Sch1] ~48 to 51 at.% Al at solidus temperatures 1520 to 1485°C [1996Tre1, 1997Bul] pure Ti at 25°C [1981Kin, Mas2] dissolves up to ~0.4 at.% Mo [Mas2] for the single phase Ti-2.5Al-2.5Mo alloy annealed at 800°C/222 h [1990Ere] for (+) alloy (Ti-5Al-5Mo) annealed at 800°C/222 h [1990Ere] ~23-28.5 at.% Al [2003Sch2] dissolves up to ~14 at.% Ti at 1600°C, ~22 at.% Ti at 1300 and 1000°C [1987Ere, 1988Ere1, 1990Ere] [V-C2] in the Ti-40Al-50Mo (J+') alloy annealed at 1000°C for 200 h [1990Ere] in the Ti-40Al-40Mo alloy ('+)+J) annealed at 1000°C for 200 h [1990Ere] in the Ti-19Al-55Mo alloy (+') annealed at 1600°C/53 h+1300°C/101 h [1990Ere] in the Ti-39Al-37Mo alloy (+') annealed at 1300°C for 101 h [1990Ere] ~46 to 52 at.% Al [2003Sch2]
a = 308.9 to 309.8 [1971Rex]
[1971Rex, Mas2, 1997Sau, 2003Sch2]
MSIT ®
Al–Mo–Ti
302 Phase/ Temperature Range [°C]
, Mo 3Al8 < 1555 10
Pearson Symbol/ Space Group/ Prototype mC22 c2/m Mo3Al8
MoAl3(h) 1222 - 818
mC32 Cm MoAl3
Mo1-xAl3+x(h) 1154 - 1260
cP8 Pm3n Cr3Si mC30 Cm WAl4
MoAl4(h) 1177 - 942
Mo4Al17 < 1034
MSIT®
mC84 C2 Mo4Al17
Lattice Parameters Comments/References [pm] a = 920.8 b = 363.8 c = 1006.5 = 100.78° a = 916.4 b = 363.9 c = 1004.0 = 100.50° a = 920.7 0.3 b = 364.1 0.1 c = 1006.0 0.5 = 100.78 0.09° a = 919 b = 363 c = 1008 = 101° a = 913 b = 354 c = 1009 = 100.33° a = 913 b = 362 c = 1002 = 100.62° a = 916.2 b = 363.8 c = 1000.3 = 100.47° a = 910 b = 364 c = 1005 = 100.82° a = 1639.6 b = 359.4 0.1 c = 838.6 0.4 = 101.88° a = 494.5
[V-C2] 72.7 at.% Al [2003Sch2]
[1991Sch]
[1990Ere], in the Ti-60Al-30Mo (+J+ ) alloy annealed at 1300°C for 63 h
[1990Ere], in the Ti-55Al-40Mo (+'+ ) alloy annealed at 1300°C for 107 h
[1990Ere], in the Ti-47Al-51Mo ('+ ) alloy annealed at 1000°C
[1990Ere], in the Ti-75Al-23Mo (J+ ) alloy annealed at 1000°C
[1990Ere], in the Ti-55Al-40Mo (J+ +') alloy annealed at 1000°C
[1991Sch, 1997Smi, 2003Sch2]
76 to 79 at.% Al [1991Sch, 1997Smi, 2003Sch2]
79 to 80 at.% [1991Sch] [V-C2] a = 525.5 b = 1776.8 c = 522.5 = 100.88° [1991Sch] [1995Gri] a = 915.8 0.1 b = 493.23 0.08 c = 2893.5 0.5 = 96.71 0.01° Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti Phase/ Temperature Range [°C] Mo5Al22(h) 964 - 831 MoAl5(h 2) 846 (750 < T < 800)
Pearson Symbol/ Space Group/ Prototype oF216 Fdd2 Mo5Al22 hP12 P6 3 WAl5
hP60 P3 hP36 R3c cI26 Im3 WAl12 tI18 J, (Ti1-xMox)1+yAl3-y I4/mmm TiAl3(h)
MoAl5(h 1) (850 - 750) - 648 MoAl5(r) < 650 MoAl12 712
Lattice Parameters Comments/References [pm] a = 7382 3 b = 916.1 0.3 c = 493.2 0.2 a = 491.2 0.2 c = 886.0 0.4 a = 493.7 c = 924.3 a = 493.3 0.1 c = 4398 9 a = 495.1 0.1 c = 2623 1 a = 758.15 a = 758.77
TiAl3 a = 384.9 c = 860.9 a = 385.3 c = 858.7 a = 384 c = 859 a = 380.7 c = 839.2 a = 379.8 c = 836.7 a = 380 c = 841 a = 384.0 c = 830.8 a = 387.1 c = 831.8 a = 389 c = 829 a = 383 c = 849 a = 387.4 c = 830.3 a = 390 c = 825 a = 386.5 0.3 c = 843.9 0.1
Landolt-Börnstein New Series IV/11A3
303
[1991Sch] [1995Gri]
[1991Sch] [V-C2] [1991Sch] [1991Sch] [V-C2] [1991Sch] [V-C2], D022 ordered phase 0 x 0.88; 0 y ~0.21 [1970Han, 1987Ere, 1990Ere, 1996Tre2, 1990Abd] 72.4 to 75.0 at.% Al [2003Sch1] < 1425°C [1999Tre1, 1997Bul] 1385 to 735°C, 74.5-75 at.% Al at 1200°C [2001Bra] melting temperature 1408°C [2003Kar] [1970Han] Ti-75Al-12.5Mo [1970Han] Ti-76Al-16Mo [1970Han] Ti-75Al-20Mo annealed at 1000°C for 121 h [1990Ere] Ti-68Al-16Mo [1970Han] Ti-64Al-10Mo [1970Han] Ti-62.5Al-12.5Mo annealed at 1000°C [1990Ere] Ti-70Al-13Mo annealed at 1000°C for 100 h [1990Ere] Ti-65Al-15Mo annealed at 1300°C/50 h + 1000°C/147 h [1990Ere] Ti-60Al-15Mo annealed at 1300°C/50 h + 1000°C/ 147 h [1990Ere] Ti-67Al-10Mo, 1300°C [1996Tre3]
MSIT ®
Al–Mo–Ti
304 Phase/ Temperature Range [°C] TiAl3(l) < 950
Pearson Symbol/ Space Group/ Prototype tI32 I4/mmm TiAl3(l)
Lattice Parameters Comments/References [pm] a = 387.7 c = 3382.8
tetragonal superstructure of AuCu type a* = 395.3 c* = 410.4 a* = 391.8 c* = 415.4 tI16 I4/mmm ZrAl3
tP28 P4/mmm Ti2Al5
a = 398.81 to 392.3 c = 1646.69 to 1653.49 a = 399.1 1.3 c = 1646.6 0.5 a = 392.8 0.6 c = 1656.3 1.5 a = 390.53 c = 2919.63
, TiAl2 < 1199 tP4 P4/mmm AuCu
orthorhombic, Pmmm, with pseudotetragonal cell tI24 I41/amd HfGa2 oC12 Cmmm ZrGa2
tP32 P4/mbm Ti3Al5 MSIT®
a = 403.0 c = 395.5 a = 402.62 b = 396.17 c = 402.62
74.5 to 75 at.% Al [2001Bra]
summarizes several phases [2003Sch1] Ti5Al11 [2001Bra] stable in the range 1416-995°C, 66 to 71 at.% Al at 1300°C [2001Bra] (including the stoichiometry Ti 2Al5) at 66 at.% Al, * AuCu subcell only [2001Bra] at 71 at.% Al, * AuCu subcell only [2001Bra] Ti5Al11, D023 type [V-C] 65.8 to 70.9 at.% Al, 1416-1206°C [1990Sch] 69 to 71 at.% Al, 1450-~990°C [1996Tre1, 1997Bul] in the as cast Ti-68Al-2Mo alloy (+J) [1996Tre3] in the Ti-70Al-2Mo alloy (+J) annealed at 1300°C for 24 h [1996Tre3] Ti2Al5, 1416-990°C [1992Kat], ~1215-985°C [1990Sch]; included in the homogeneity range of Ti5Al11 [2001Bra] chosen stoichiometry [1992Kat] summarizes several phases [2003Sch1]: Ti1-xAl1+x, 63 to 65 at.% Al at 1300°C, stable in the range 1445-1170°C [2001Bra] for Ti36Al64 at 1300°C [2001Bra] 1445-1424°C [1990Sch] for as arc melted Ti36 Al64 [1990Sch]
a = 397.0 c = 2430.9
stable structure of TiAl 2 < 1216°C, 66 to 67 at.% Al at 1000°C [2001Bra];
a = 396.7 c = 2429.68 a = 1208.84 b = 394.61 c = 402.98 a = 1209.44 b = 395.91 c = 403.15 a = 1129.3 c = 403.8
shown as TiAl2(r) < 1214°C [1900Sch] metastable modification of TiAl2 observed only in as cast alloys [2001Bra] TiAl2(h), 66 to 67 at.% Al, 1433-1214°C [1990Sch] Ti3Al5, stable below 810°C [2001Bra]
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti Phase/ Temperature Range [°C] , TiAl < 1463
Pearson Symbol/ Space Group/ Prototype tP4 P4/mmm AuCu
Lattice Parameters Comments/References [pm] a = 400.5 c = 407.0
a = 400.0 0.1 c = 407.5 0.1
[V-C], L1 0 ordered phase 46.7 to 66.5 at.% Al [1992Kat, 1993Oka]; 50 to 62 at.% Al at 1200°C [2001Bra] ~52 to 65 at.% Al at solidus temperatures, ~50 to 60 at.% Al at 1000°C [1996Tre1, 1997Bul] at 50 at.% Al [2001Bra]
a = 398.4 0.1 c = 406.0 0.1
at 62 at.% Al [2001Bra]
a = 398.1 c = 407.5
in Ti-50Al-5Mo ( +2) alloy annealed at 1240°C for 150 h [1993Das1]
a = 397 c = 408
Ti-55Al-5Mo annealed at 1300°C/ 111 h + 1000°C/ 150 h [1990Ere]
a = 399.2 c = 405.6
in Ti-65Al-5Mo ( +J) alloy annealed at 1300°C /13 h + 1000°C/ 26 h + 800°C/ 205 h [1996Tre3] the same ( +J) alloy annealed at 1300°C/ 150 h [1990Ere]
a = 396 c = 408
Landolt-Börnstein New Series IV/11A3
305
a = 396.0 1.1 c = 407.5 0.2
in Ti-50Al-10Mo (+ ) alloy annealed at 1300°C/ 13 h + 1000°C 26 h + 800°C/ 205 h [1996Tre3]
a = 400.6 c = 405.7
in Ti-34.5Al-1.5Mo(mass%) alloy (Ti50.8Mo0.6Al48.6) annealed at 1000°C for 1 h [1991Mae]
a = 401.3
c/a = 1.008, 1.015 or 1.013 in the Ti-44Al-2Mo alloy as HIPpped (High Isostatic Pressed), annealed at 900°C or 1200°C, respectively [1994Mor]
MSIT ®
306 Phase/ Temperature Range [°C] 2, (Ti3Al) Ti3Al < 1164
* ), ~Ti3Mo3Al4
L60
MSIT®
Al–Mo–Ti Pearson Symbol/ Lattice Parameters Comments/References Space Group/ [pm] Prototype D019 ordered phase hP8 P6 3/mmc ~20 to 38.2 at.% Al, maximum at 30.9 at.% Al Ni3Sn [1992Kat, 1993Oka, 2003Sch1] < 1180°C [1993Gam] maximum at 32.5 at.% Al, ~1200°C [1996Tre1, 1997Bul] < 1210°C (+2) [1994Kai, 2000Oka] [V-C] a = 578.2 c = 468.9 at 28 at.% Al [L-B] a = 580.6 c = 465.5 at 28 at.% Al [L-B] a = 574.6 c = 462.4 at 32 at.% Al [1997Bul] a = 577.5 0.4 c = 463.7 0.5 at 25 at.% Al, annealed at 1300°C/40 h a = 579.5 + 1000°C/90 h + 800°C/222 h [1990Ere] c = 464.1 metastable 2 phase a = 567 c = 451 (Ti-54.2Al-13.0Mo) in the Ti-50Al-15Mo alloy annealed at 1400°C/2 h and water quenched (WQ) [1997Che] a = 606 in the Ti-21.6Al-3.4Mo alloy aged at 450°C c = 495 [1992Dja1] a = 576.2 in the Ti-44Al-2Mo alloy (2+2+ ) c = 461.9 [1994Mor] [V-C2], single phase (Ti26 Al41 Mo33 tP30 P4 2/mnm at 925°C) [1970Han] in the Ti-42Al-25Mo alloy annealed at 925°C )CrFe a = 966.7 [1970Han] c = 501.8 in the Ti-42Al-33Mo alloy annealed at 925°C a = 965.1 [1970Han] c = 501.8 in the Ti-42Al-36Mo alloy annealed at 925°C a = 963.6 [1970Han] c = 499.7 in the Ti-48Al-26Mo alloy annealed at 925°C a = 959.1 [1970Han] c = 496.6 ) (Ti~28Al~40Mo~32) a = 966 c = 502 in the Ti-40Al-30Mo (+'+)) alloy annealed at 1300°C/63.5 h+1000°C/200 h [1990Ere] a = 966 in the Ti-40Al-40Mo alloy (J+'+)) annealed c = 501 at 1000°C [1990Ere] tP4 a = 395 intermediate phase observed in the P4/mmm b = 403 Ti-50Al-15Mo alloy annealed at 1400°C for b/a = 1.020 1.5 h and WQ (composition of the phase Ti-57.4Al-9.4Mo) [1997Che]
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti Phase/ Temperature Range [°C] ´
Pearson Symbol/ Space Group/ Prototype tI18 I4/mmm
´´
tP18 P4/mmm
´
hP2 P6 3/mmc Mg oC4 P2221 U hP3 P6/mmm 7TiCr
´´
7
Landolt-Börnstein New Series IV/11A3
307
Lattice Parameters Comments/References [pm] a = 397 c = 815 c/a = 2.05
the ordered D022 type form of the phase formed by diffusionless way, observed in the Ti-50Al-15Mo alloy quenched from temperatures in the range of 1400-800°C [1997Che] the ordered phase observed in the Ti-50Al-15Mo alloy after prolonged aging at 800°C, supposed to be formed as a result of further ordering of the ´ (D022) [1997Che] martensite phase in Ti-xMo-3Al alloys (0 x 4) [1980Sas] martensite phase in Ti-xMo-3Al alloys (7 x 12) [1980Sas] metastable phase, appeared during quenching of /2 phases (7 athermal) or aging of metastable (quenched) /2 phases (7 isothermal) [1971Wil, 1972Luz, 1980Sas, 1991Dja, 1992Dja1, 1992Dja2, 1993Cui, 1997Che]
MSIT ®
MSIT®
α2+β+γ
<1125
1145
ε+δ+ρ
1325
1400
β+γ/γ´+σ
<550
α+β+β´
α+α2+β
β+ρ+σ
α + β α2 + β´ α+α2+β´
U10
>1125 e5 α α2 + β
β+ε+δσ
P
E2
U5
U3
U11
ε+ρ+σ
β+γ/γ´+ε
L β + γ/γ´ + ε
γ/γ´+ε+L
β+δ+ε
U2
β+ρ+ζ2
Lβ+ε
δ+Lβ+ε
Lβ+δ
L + ζ1 β + δ
β+δ+ζ1
β+ε+σ
1250
β+ε+ρ
1370
E1
1470
β + ε γ/γ´ + σ
E3
γ/γ´+ε+σ
1000>T>925
α α2 + β + γ
γ+ε+η
U9
U8
U7
α+α2+β
γ+ε+ζ
γ+ζε+η
β+δε+ρ
β+δ+ρ
L+ζγ+ε
L+β+γ
U6
ζ2 β + δ + ρ
>1475 p5 L+ δ ε
U4
1500
L+β+ζ1
U1
A-B-C
L + ζ2 β + ζ1
L + ρ β + ζ2
L+β+ζ2
1550
β + ζ1 δ + ζ2
β+ζ1+ζ2
α+Lβ+γ
1455
β+δ+ζ2
α+β+γ
1440
p3
1495
β+lα
1600>T>1500
Al-Mo-Ti
1750>T>1600
Fig. 1a: Al-Mo-Ti. Reaction scheme up to 75 at.% Al
990 e7 ζε+η
1118 e6 α α2 + γ
p9 1199 γ+ζη
1416 p7 l+γζ 1393 p8 lζ+ε
1463 p6 α+lγ
1490 p4 β+lα
Al-Ti
1470 e4 ζ2 δ + ρ
1490 e3 ζ 1 δ + ζ2
2150 p1 l+βρ 1720 e1 l ρ + ζ2 1570 p2 l + ζ2 ζ1 1535 e2 l δ + ζ1
Al-Mo
675 e8 β α + β´
Mo-Ti
308 Al–Mo–Ti
Landolt-Börnstein New Series IV/11A3
Landolt-Börnstein New Series IV/11A3
ε+Mo8Al22+MoAl5
L+ε+Mo8Al22
U7
650 ε+MoAl12+(Al)
Lε+MoAl12+(Al)
L+ε+MoAl12
ε+Mo4Al17+MoAl5
E
L+ε+MoAl5
ε+MoAl12+Mo4Al17
MoAl5+Mo4Al17+MoAl12
U11
U8
U6
ε+MoAl5Mo4Al17+MoAl12 U12
ε+MoAl5+MoAl12
U4
U2
ε+δ+MoAl3
L+MoAl5 ε+MoAl12
ε+δ+Mo4Al17
U10
L+Mo8Al22ε+MoAl5
ε+MoAl3+Mo4Al17
ε+MoAl4MoAl3+Mo4Al17
ε+MoAl3+MoAl4
ε+Mo1-xAl3+xMoAl4+MoAl3
ε+MoAl3δ+Mo4Al17
ε+Mo8Al22Mo4Al17+MoAl5 U9
ε+Mo4Al17+Mo8Al22
U5
ε+MoAl4+Mo4Al17
L+Mo4Al17ε+Mo8Al22
L+ε+Mo4Al17
L+MoAl4 ε+Mo4Al17
ε+Mo1-xAl3+x+MoAl4
L+Mo1-xAl3+xε+MoAl4
L+ε+MoAl4
MoAl3+Mo1-xAl3+x+ε
A-B-C
δ+Mo1-xAl3+xε+MoAl3 U3
U1
δ+ε+Mo1-xAl3+x
L+δ ε+Mo1-xAl3+x
Al-Mo-Ti
L+ε+Mo1-xAl3+x
L+δ+ε
Fig. 1b: Al-Mo-Ti. Reaction scheme for the Al-rich part
664 p8 l + ε (Al)
Al-Ti
e1
e2
e3
e6
MoAl5Mo4Al17+MoAl12
648
660 e5 lMoAl12+(Al)
712 p7 l+MoAl5MoAl12
818 e4 MoAl3 δ+Mo4Al17
Mo8Al22Mo4Al17+MoAl5
831
846 p6 l +Mo8Al22MoAl5
MoAl4MoAl3+Mo4Al17
942
964 p5 l+Mo4Al17Mo8Al22
1034 p4 l+MoAl4Mo4Al17
Mo1-xAl3+xMoAl3+MoAl4
1154
1177 p3 l+Mo1-xAl3+xMoAl4
1222 p2 δ+Mo1-xAl3+xMoAl3
1260 p1 l + δ Mo1-xAl3+x
Al-Mo
Mo-Ti
Al–Mo–Ti 309
MSIT ®
Al–Mo–Ti
310
Al
Data / Grid: at.%
Fig. 2: Al-Mo-Ti. Projection of the partial solidus surface
Axes: at.%
20
80
ζ
~1400 >1415
40
1470
Mo1-xAl3+x
δ
ε
γ
~1500 60
1370 ~1440
α
ζ1
~1550
ζ2
60
40
β ρ 80
20
20
Ti
40
Mo
Data / Grid: at.% Axes: at.%
γ
40
p6
p4
MSIT®
80
Ti 30.00 Mo 0.00 Al 70.00
Fig. 3: Al-Mo-Ti. Partial liquidus surface projection
Ti 60.00 Mo 0.00 Al 40.00
60
50
U6
60
α
p3
50
β
10
20
Ti 30.00 Mo 30.00 Al 40.00 Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti
311
Al
Data / Grid: at.%
Fig. 4: Al-Mo-Ti. Isothermal section at 1600°C
Axes: at.%
20
80
L 40
60
ζ2
L+β 60
40
β
ρ
80
20
20
Ti
40
60
80
Ti 30.00 Mo 0.00 Al 70.00
Fig. 5: Al-Mo-Ti. Partial isothermal section at 1500°C
Mo
Data / Grid: at.% Axes: at.%
40
60
L L+α
α
50
L+β 50
α +β L+β
β
Ti 60.00 Mo 0.00 Al 40.00 Landolt-Börnstein New Series IV/11A3
10
20
Ti 30.00 Mo 30.00 Al 40.00
MSIT ®
Al–Mo–Ti
312
Ti 35.00 Mo 0.00 Al 65.00
Fig. 6: Al-Mo-Ti. Tentative partial section at 1400°C
Data / Grid: at.% Axes: at.%
40
60
γ γ +β
α+γ
50
50
α α+β
β
10
Ti 60.00 Mo 0.00 Al 40.00
20
Al Fig. 7: Al-Mo-Ti. Isothermal section at 1300°C
Ti 35.00 Mo 25.00 Al 40.00
Data / Grid: at.% Axes: at.%
L
20
80
ζ
δ
ε 40
γ' 60
γ
α 60
40
β2 ρ 80
20
β
Ti
MSIT®
20
40
60
80
Mo
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti
313
Ti 40.00 Mo 0.00 Al 60.00
Fig. 8: Al-Mo-Ti. Partial isothermal section at 1200°C
Data / Grid: at.% Axes: at.%
γ
50
60
50
β +γ
α
40
β
70
30
10
Ti 80.00 Mo 0.00 Al 20.00
20
30
Ti 40.00 Mo 0.00 Al 60.00
Fig.9: Al-Mo-Ti. Partial isothermal section at 1100°C
Ti 40.00 Mo 40.00 Al 20.00 Data / Grid: at.% Axes: at.%
50
γ
50
60
40
α2 70
30
α 2+β α 80
20
β
Ti 90.00 Mo 0.00 Al 10.00 Landolt-Börnstein New Series IV/11A3
10
20
30
40
Ti 40.00 Mo 50.00 Al 10.00
MSIT ®
Al–Mo–Ti
314
Al
Data / Grid: at.%
L
Fig. 10: Al-Mo-Ti. Isothermal section at 1000°C
Axes: at.%
Mo4Al17 MoAl4(h)
20
η 40
80
MoAl3(h)
δ
ζ ε γ1
60
γ
α2
α
σ
60
40
ρ
β2
80
20
β 20
Ti
40
60
80
Al
Data / Grid: at.%
L
Fig. 11: Al-Mo-Ti. Isothermal section at 925°C
Mo
Axes: at.%
Mo5Al22(h) Mo4Al17
20
80
η
MoAl3(h)
δ
ε
40
60
γ
α2
α
σ
60
40
ρ
β2
80
20
β
Ti
MSIT®
20
40
60
80
Mo
Landolt-Börnstein New Series IV/11A3
Al–Mo–Ti
315
Al
Data / Grid: at.%
Fig. 12: Al-Mo-Ti. Partial isothermal section at 800°C
Axes: at.%
20
80
ε
η 40
60
γ
60
40
α2
80
20
α
β2 β β +β ' 20
Ti
40
60
80
40.00 Ti 0.00 Mo Al 60.00
Fig. 13: Al-Mo-Ti. Partial isothermal section at 700°C
Mo
Data / Grid: at.% Axes: at.%
γ
60
40
α2
80
20
β2 α
Ti
Landolt-Börnstein New Series IV/11A3
β'
β
20
40
40.00 Ti Mo 60.00 0.00 Al
MSIT ®
Al–Mo–Ti
316
Al Fig. 14: Al-Mo-Ti. Isothermal section at 600°C
Data / Grid: at.%
(Al)
Axes: at.%
MoAl12 MoAl5(r) Mo4Al17
20
80
δ
ε
η 40
60
γ σ
60
40
α2
ρ
80
20
α
Ti
MSIT®
β 20
β2 β' 40
60
80
Mo
Landolt-Börnstein New Series IV/11A3
Al–N–Si
317
Aluminium – Nitrogen – Silicon Hans Leo Lukas Literature Data During the investigation of the quaternary Al-Si-N-O system [1975Gau] found a phase Al5+xSi3-xN 9-xOx, which exists in the range 0 x 3. [1978Sch, 1980Sch] however, assumed this phase to be unstable in the ternary Al-N-Si system, as it needs some oxygen to be stabilized. Other investigations regarding quaternary systems with AlN-Si3N4 as boundary system [1978Lan, 1983Hua, 1986Hua, 1988Fuk, 1990Wei] do not mention this phase and assume AlN to be in equilibrium with Si3N4. [1992Hil] thermodynamically calculated the Al-N-Si system, assuming the ionic liquid model with ideal solution behavior for a nitride liquid. These authors did not consider Al5Si3N9 to be a stable phase. Thus at 1 bar pressure the only stable phases taken into account are liquid, the solid metals (Al) and Si, solid AlN and solid Si3N4. All these phases have only small ranges of homogeneity, which for AlN and Si3N4 were neglected by [1992Hil] in their calculation. [2001Kas] synthesized Al1-xSixN solid solutions up to x = 0.12 by metalorganic vapor-phase epitaxial growth. From the thermodynamic point of view it is very likely, that this solid solution has to be considered as metastable supersaturated although the crystal quality is very perfect, measured by the full width at half maximum of 100 arcsec ( = 0.028°) of an X-ray rocking curve (single crystal rotation technique focussed on a single X-ray peak). The temperature during preparation (900°C) may be far too low to enable equilibration. The solid solution was characterized as substitutional, one Si atom replaces one Al atom. The same authors [2001Tan] reported lattice parameter measurements of Al1-xSixN in dependence of x, extrapolated from the epitaxial layer to zero residual strain. [2002Wu] prepared Al containing solid solutions of Si3N4 by Al ion implantation in order to study the influence of Al on the oxidation behavior of Si3N4. No structural details of the solid solution were reported. Binary Systems The Al-Si system is accepted from [2003Luk]; it is based on the thermodynamic assessment of [1997Feu]. The N-Si and Al-N systems are accepted from the thermodynamic assessments of [1991Hil1] and [1991Hil2], respectively. The calculation of the ternary system by [1992Hil] used the latter two binary assessments and an older assessment of the Al-Si system without any ternary excess term. The calculated results, except near the eutectic of the binary Al-Si system, do not show a visible dependence on the selection of the binary Al-Si assessment. Solid Phases Stable binary phases are AlN and Si3N4. Pure Si3N4 is metastable but formed as the main product during reaction of Si with N2. It is stabilized, however, by large cations, e.g. rare earth oxides. The phase Al 5Si3N9 possibly exists only in the oxygen stabilized form Al5+xSi3-xN9-xOx with x > 0. All solid phases are summarized in Table 1. Invariant Equilibria At 1 bar pressure the only four-phase equilibria are: (i) Gas+LSi3N4+AlN at 1839.5°C, which is nearly degenerated and very near to the quasibinary three-phase equilibrium Gas+LSi3N4 at 1840.5°C; (ii) L(Al)+(Si), AlN at 577°C, which is totally degenerated and identical to the binary Al-Si eutectic. Isothermal Sections Figures 1 and 2 show the isothermal sections at 1 bar and 2400 or 1800°C, calculated by [1992Hil].
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–N–Si
318 Temperature – Composition Sections
The AlN-Si3N4 section, calculated for 1 bar [1992Hil], is shown in Fig. 3. Notes on Materials Properties and Applications Epitaxially grown Al1-xSixN layers are promising candidates as materials for flat panel displays (FE-displays), as the Si content in AlN decreases the electric field necessary for field emission (FE). As part of the Al-Si-N-O system, Al-Si-N is interesting for high temperature materials based on SIALON. References [1975Gau]
[1976Jac] [1978Lan]
[1978Sch]
[1980Sch] [1983Hua]
[1986Hua]
[1988Fuk]
[1990Wei]
[1991Hil1] [1991Hil2] [1992Hil] [1997Feu]
[2001Kas]
[2001Tan]
MSIT®
Gauckler, L.J., Lukas, H.L., Petzow, G., “Contribution to the Phase Diagram Si3N4-AlN-Al2O3-SiO 2”, J. Am. Ceram. Soc., 58, 366-367 (1975) (Experimental, Equi. Diagram, 10) Jack, K.H., “Review: Sialon and Related Nitrogen Ceramics”, J. Mater. Sci., 11, 1135-1158 (1976) (Review. Equi. Diagram, Crys. Structure, 41) Land, P.L., Wimmer, J.M., Barns, R.W., Choudhury, N.S., “Compounds and Properties of the System Si-Al-O-N”, J. Am. Ceram. Soc., 61, 56-60 (1978) (Experimental, Equi. Diagram, 25) Schneider, G., “Equilibrium Investigations in the Si, Al, Be/C, N System” (in German), Thesis, University of Stuttgart, Germany (1978) (Experimental, Equi. Diagram, Crys. Structure, 71) Schneider, G., Gauckler, L.J., Petzow, G., “Phase Equilibria in the System AlN - Si3N4 Be3N2”, J. Am. Ceram. Soc., 63, 32-35 (1980) (Experimental, Equi. Diagram, 7) Huang, Z.K., Greil, P., Petzow, G., “Formation of -Si3N4 Solid Solutions in the System Si3N4-AlN-Y2O3”, J. Am. Ceram. Soc., 66, C-96-C-97 (1983) (Experimental, Equi. Diagram, 5) Huang, Z.K., Tien, T.-Y., Yen, T.-S., “Subsolidus Phase Relationships in Si3N4-AlN-Rare Earth Oxide Systems”, J. Am. Ceram. Soc., 69, C-241-C-242 (1986) (Experimental, Equi. Diagram, 5) Fukuhara, M., “Phase Relationships in the Si3N 4 Rich Portion of the Si3N4-AlN-Al2O3-Y 2O3 System”, J. Am. Ceram. Soc., 71, C359-361 (1988) (Experimental, Equi. Diagram, 10) Weitzer, F., RemsChnig, K., Schuster, J.C., Rogl, P., “Phase Equilibria and Structural Chemistry in the Ternary Systems M-Si-N and M-B-N (M = Al, Cu, Zn, Ag, Cd, In, Sn, Sb, Au, Tl, Pb, Bi)”, J. Mater. Res., 5, 2152-2159 (1990) (Experimental, Equi. Diagram, Crys. Structure, 39) Hillert, M., Jonsson, S., “Report, Trita-Mac-465”, Royal Inst. of Technology, Stockholm, Sweden, (1991) (Thermodyn., Equi. Diagram, Assessment, 0) Hillert, M., Jonsson, S., “Report, Trita-Mac-466”, Royal Inst. of Technology, Stockholm, Sweden, (1991) (Thermodyn., Equi. Diagram, Assessment, 0) Hillert, M., Jonsson, S., “Prediction of the Al-Si-N System”, Calphad, 16, 199-205 (1992) (Thermodyn., Equi. Diagram, Assessment, 11) Feufel, H., Gödecke, T., Lukas, H.L., Sommer, F., “Investigation of the Al-Mg-Si System by Experiments and Thermodynamic Calculations”, J. Alloys Comp., 247, 31-42 (1997) (Experimental, Assessment, Thermodyn., Equi. Diagram, 38) Kasu, M., Taniyasu, Y., Kobayashi, N., “Formation of Solid Solution of Al1-xSixN (0x12%) Ternary Alloy”, Jpn. J. Appl. Phys. 2, 40(10A), L1048-L1050 (2001) (Experimental, 12) Taniyasu, Y., Kasu, M., Kobayashi, N., “Lattice Parameters of Wurtzite Al1-xSixN Ternary Alloys”, Appl. Phys. Lett., 79(26), 4351-4353 (2001) (Experimental, Crys. Structure, 14)
Landolt-Börnstein New Series IV/11A3
Al–N–Si [2002Wu]
[2003Luk]
319
Wu, J., Wang, Y., Ye, J., Du, H.H., “The Cyclic and Continuous Oxidation of with and without Aluminum Implantation”, Key Eng. Mater., 224-226, 803-806 (2002) (Experimental, Corrosion, 14) Lukas, H.L., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 29)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Al) < 660.45
cF4 Fm3m Cu
a = 404.93
at 23°C [V-C2]
(Si) < 1414
cF8 Fd3m C (diamond)
a = 543.06
at 30°C [V-C2]
Al1-xSixN < 2800 50
hP4 P6 3mc ZnS (wurtzite)
a = 311.15 at 17°C, x = 0 [V-C2] c = 497.98 a = 311.13 - 14.12x 0 x 0.12 [2001Tan] c = 498.18 - 22.99x metastable ?
Si3N4
hP14 Be2SiO 4
a = 760.8 c = 291.1
[V-C2]
Si3N4
hP28 Si3N4
a = 775 to 782 c = 562 to 559
metastable, stabilized by rare earth oxides three sets of parameters [V-C2]
Al5+xSi3-xN9-xOx
hexagonal
a = 307.9 c = 530
0(?) x 3 [1975Gau] possibly not stable at x = 0 [1978Sch, 1980Sch]; parameters from [1976Jac] for Si3Al7N11 formula
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–N–Si
320
N Fig. 1: Al-N-Si. Isothermal section at 2400°C
Data / Grid: at.% Axes: at.%
Gas 20
80
40
60
AlN 60
L+Gas +AlN
40
L+Gas 80
20
L+AlN
20
Al
40
L
60
80
N
Si
Data / Grid: at.%
Fig. 2: Al-N-Si. Isothermal section at 1800°C
Axes: at.%
20
80
Gas+AlN +Si3N4 40
60
Si3N4
AlN 60
40
L+AlN +Si3N4 L+AlN
80
Al MSIT®
20
40
20
L
60
80
Si
Landolt-Börnstein New Series IV/11A3
Al–N–Si
Fig. 3: Al-N-Si. Section from AlN to Si3N4
321
3000
Gas
Temperature, °C
2750
L+Gas 2500
2250
L+Gas+AlN 2000
1839.6 1750
Al Si N
Landolt-Börnstein New Series IV/11A3
50.00 0.00 50.00
10
1840.6°C
Si3N4+L+Gas
Si3N4+AlN 20
Si, at.%
30
40
Al Si N
0.00 42.86 57.14
MSIT ®
322
Al–N–Ti
Aluminium – Nitrogen – Titanium Vasyl Tomashik and Pierre Perrot Literature Data A critical assessment of the Al-N-Ti ternary system has been published by [1993Jeh], which included the literature data up to the year 1991. Thermodynamic data appearing up to 1997 are included in the thermodynamic assessment made by [1998Che]. Subsequently this system was investigated in different experimental approaches and for different temperatures. The present evaluation takes care of all data, from the first publication to the present. The investigations in this ternary system are concerned with (a) phase diagram studies, (b) preparation and characterization of the ternary compounds and (c) the formation of metastable solid solutions in the AlN-TiN pseudobinary system. The equilibria in the Ti-rich part of the ternary system have been determined by [1954Thy] for 0 to 10 mass% Al and 0 to 1 mass% N. This study applied micrograph analysis and X-ray diffraction of samples annealed at 600 to 1250°C, for 576 to 6 h. These samples were prepared from high purity arc molten alloys. The obtained results are given as vertical sections for constant N content. Annealing of Al-TiN bilayers on SiO2 for 15 h at 645°C leads to the formation of AlN and Al3Ti, as the data of [1982Wit] show. These phases are also formed by reaction sintering of powder mixtures Al+TiN, containing 10, 15, 20 and 30 mol% TiN [1992Koy]. Titanium specimens with embedded AlN particles, as well as AlN-Ti and AlN-TiN diffusion couples were annealed at 900 to 1000°C up to 40 h by [2000Par]. It was shown that in AlN-TiAl diffusion couples a ternary Ti2AlN phase is formed at the interface. A more complex AlN-TiN-Ti3AlN-Ti3Al-Ti-Ti reaction zone was observed at the AlN-Ti interface. Thermodynamic calculations give the same sequence of expected layers between AlN and pure Ti [1998Lee] (the composition of Ti at the Ti3AlN/Ti interface is close to the composition of Ti3Al). AlN never is in contact with Ti3AlN [2000Par]. Nitriding the intermetallic TiAl3 in nitrogen and ammonium flow was studied by [1983Psh] in a temperature range of 600 to 1200°C. This work states that Al and Ti are nitrated in fact simultaneously, which results in formation of a heterogeneous mixture of practically not interacting binary nitrides. Experimental results imply that AlN-TiN, TiAl3-AlN and TiAl3-TiN are stable tie lines in the Al-N-Ti ternary system at low temperatures [1984Bey]. Phase equilibria in this ternary system were investigated at 1000 and 1300°C using previously prepared Al-Ti alloys, AlN, TiN and Ti powders [1984Sch]. About 30 ternary alloys were cold-pressed and sintered at the following conditions: 1000°C for 240 to 800 h in BN crucibles sealed in evacuated quartz tubes, 1200°C for 60 h in Mo crucibles under dynamic vacuum, 1300°C for 60 h in Mo crucibles under dynamic vacuum or for 50 h in BN (Mo) crucibles under argon and 1400°C in Mo crucibles under dynamic vacuum. As the alloys sintered at 1000°C were initially not in equilibrium they were powderized again, cold-pressed and sintered again. These two isothermal sections were included in the reviews [1985Sch, 1992Sch, 1993Jeh, 1998Che]. The isothermal section at 900°C was constructed by [1997Dur] which was supported by the thermochemical calculations. Based on such calculations the 1000°C isotherm is expected to be virtually not altered with respect to the 900°C isotherm, which disagrees with [1984Sch]. It was concluded by [1997Dur] that the samples of [1984Sch] were not heat treated sufficiently long to reach equilibrium at 1000°C. For the 850°C isotherm the thermochemical calculations predict a three-phase field AlN+TiN+TiAl2 rather than AlN+TiN+TiAl3 [1997Dur]. The 1325°C equilibrium isothermal section of the Al-N-Ti ternary system with accounting of Ti4AlN3-x formation was constructed by [2000Pro2]. The isothermal sections of the Al-N-Ti system at 1200, 1400, 1580, 1600, 1900 and 2500°C were calculated thermodynamically by [1998Che] but in these calculations the existence of the Ti3Al2N2 was taken into account. As the new investigations indicate that the more probable composition of Ti3Al2N2 in this system is Ti4AlN 3-x these isothermal sections must be recalculated and the Al-N-Ti ternary system needs a revised thermodynamic assessment.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–N–Ti
323
The phase diagram of the AlN-TiN pseudobinary system was calculated by [1988Hol, 1989Hol]. Unfortunately, they used a melting point of 2930°C for TiN instead of 3290°C. It was shown that the eutectic temperature and eutectic composition depend on the size of both AlN and TiN particles [1991And]. Mixtures of AlN and TiN containing from 30 to 90 mass% AlN did not show evidence of a reaction between these two materials [1976Kuz]. Similarly no reaction was found in annealing TiN powder and AlN plates up to 2000°C for 6 h. AlN-TiN composite materials were prepared by pressureless sintering in N2 atmosphere at 1870°C for 6 h [2002Tan]. Ti1-xAlxN metastable solid solutions (0 x < 0.7) can be obtained in the AlN-TiN pseudobinary system using cathodic arc plasma depositing process [1981Bee, 1986Mue, 1988Pen, 1988Ran, 1991Ike, 1992Tan, 1993Tan], or reactive dc and radio-frequency magnetron sputtering [1986Jeh, 1986Kno, 1987Hak, 1987Ina, 1987Kno, 1988Jeh, 1990McI, 1991Adi, 1993Pet, 1993Wah]. Such films could be prepared onto polished flat high speed steel surfaces [1986Jeh, 1986Kno, 1988Jeh], or stainless-steel substrates [1988Pen, 1990McI] or stellite surfaces [1986Kno], or MgO(001) substrates [1991Adi, 1991Hul, 1993Adi, 1993Pet, 1993Wah], or oxidized silicon surfaces [1991Hul, 1993Adi, 1993Wah], or Si and WC-Co substrates [1992Tan, 1993Tan]. These solid solutions based on TiN1-x phase crystallize in a cubic structure [1986Jeh, 1988Jeh] and the lattice parameter of the Ti1-xAlxN films linearly decreases with increasing Al content [1986Kno, 1987Ina, 1987Kno, 1993Adi, 1993Tan, 1993Wah]. According to the data of [1986Jeh, 1988Jeh] another phase was found in coatings deposited at low nitrogen pressures and in pure Ar atmosphere. Although Ti0.5Al0.5N is thermodynamically metastable it exhibits a good high-temperature stability during post annealing [1991Hul]. Such alloys deposited at 400°C were stable up to 1.5 h at 900°C [1990McI]). The films which contain more than 70 mol% AlN crystallize in the wurtzite structure [1991Hul, 1992Tan, 1993Tan, 1993Wah]. According to the data of [1981Bee] the amorphous Ti1-xAlxN films can be obtained when the N2 content in Ar-N 2 atmosphere is greater than 20%. The existing experimental results and thermodynamic calculations lead to a so-called vapor deposition phase diagram representing the range of metastable phases which were established by [1988Hol, 1989Hol] and then refined by [2001Spe]. The composition at which the structural transition takes place was experimentally verified at about 63 and 69 mol% AlN [2001Spe]. Binary Systems Al-N: The solubility of nitrogen in Al(s) and Al(l) is very small. Only one compound AlN exists in the Al-N binary system. The decomposition temperature of AlN under 0.1 MPa nitrogen pressure is 2437.4°C [2003Fer]. AlN undergoes a congruent melting point towards 2800 50°C under a nitrogen pressure of 10 MPa [1984Jon]. On increasing nitrogen pressure above 1GPa, AlN undergoes a transition from the wurtzite type to the rock salt type structure. Al-Ti: Three ordered phases Ti3Al, TiAl and TiAl3 are stable in this system [2003Sch]. The composition range between the phases TiAl and TiAl3, however, is still controversial, especially at temperatures above 1200°C because of the large number of long period structures. In total of five phases were suggested for this region, some occurring in narrow temperature ranges only and/or with a range of solubility. These five phases were subsumed in a simplified version by two stoichiometric compounds, TiAl2 and Ti2Al5 [2003Sch]. N-Ti: The solubility of nitrogen both in (Ti) and (Ti) is significant. The congruently melting TiN1-x compound with wide homogeneity range and incongruently melting Ti2N compound exist in this binary system [Mas2]. According to the data of [1992Rog] the new phases Ti3N2-x and Ti4N3-x are also formed in the N-Ti system. Solid Phases Three compounds (-1, Ti2AlN, -2, Ti3AlN and -3, Ti4AlN3) are formed in this system among which -1, Ti2AlN is the most stable [1995Wu] and belongs to the group of H phases [1964Now]. An excellent agreement exists between the various determinations of the lattice parameters [1963Jei, 1976Ivc2, 1977Ivc, 1984Sch, 1985Sch, 1986Kau, 1995Wu, 1999Far, 2000Bar2, 2000Gam, 2001Per]. It has been observed to exist over the temperature range from 700 to 1600°C and being deficient in nitrogen above 1300°C Landolt-Börnstein New Series IV/11A3
MSIT ®
324
Al–N–Ti
[1992Sch]. Its melting point is above at least 1950°C [1998Che], probably above 2500°C [1992Mab]. -1, Ti2AlN is easily obtained by numerous ways: (I) from Al and Ti powders which react exothermically in gaseous nitrogen to form Ti2AlN particles in a matrix of TiAl [1992Mab], (II) by hot pressing powder mixtures of Al, Ti and TiN and homogenizing the samples at 850°C for 200 h [1963Jei], (III) from AlN and Ti metal powder by sintering above 1500°C [1976Ivc2, 1977Ivc], (IV) by mixing elemental or binary powders, followed by cold pressing, then hot pressing in sealed evacuated containers at temperatures from 1275 to 1600°C and pressures of up to 1100 MPa for up to 24 h [1999Far], (V) by nitriding Al-Ti alloys at 1000°C [1999Mag], (VI) by heating 2Ti+AlN mixtures at 1400°C for 48 h under a pressure of 40 MPa [2000Bar2] and (VII) by reactive sintering AlN and Ti for 16 h under a vacuum of 10-3 Pa [2000Gam]. -2, Ti3AlN which exhibits a negligible range of homogeneity [1984Sch] has a cubic structure with a lattice parameter that varies only within experimental errors [1984Sch, 1985Sch, 1986Kau, 1992Sch]. This compound becomes nitrogen deficient above 1300°C and melts incongruently at 1590 10°C, decomposing presumably into either L+ TiN1-x+Ti 2AlN or L+ TiN1-x [1998Che]. -3, Ti4AlN3 is stable between ~1250 and 1500°C under Ar, but decomposes in air at 1400°C to form TiN [2000Pro1]. It tends to be deficient in nitrogen Ti4AlN3- (where 0 < < 0.1) [1999Bar, 1999Ho, 2000Bar1, 2000Fin, 2000Pro2, 2000Raw]. The formulae Ti3Al2N2 [1984Sch, 1985Sch, 1992Sch, 1998Bar] and Ti3Al1-xN2 [1997Lee] were initially accepted for this compound; however chemical analysis using energy dispersive spectroscopy (EDS) unequivocally proved a stoichiometry of Ti4AlN3 [1999Bar]. Fully dense polycrystalline samples of Ti4AlN 3- were processed by mixing TiH2, TiN and AlN to the desired stoichiometry [1999Bar, 1999Ho, 2000Bar1, 2000Pro1, 2000Pro2]. The mixed powders were cold-pressed at ~200 MPa, sealed in evacuated borosilicate tubes and hot isostatically pressed at 1275°C for 24 h under a pressure of ~70 MPa. To complete the reaction such samples were annealed further at a temperature of 1325°C for 168 h under an Ar atmosphere. The solubility limit for nitrogen in TiAl alloys are lower than 0.1 at.%, because the precipitation of nitrides occurs even at the smallest content of N in these alloys [1991Kaw]. Nitrogen solubility in Ti3Al should be higher than 2.32 at.% [2001Per] and can be as high as 3.5 at.% [1997Dur]. Solid solution based on aluminium does not hold detectable amount Ti, and TiN1-x dissolves very small amounts of Al [1984Sch]. Details of crystal structure of all solid phases are given in Table 1. Pseudobinary Systems The phase diagram of the AlN-TiN pseudobinary sub-system, has been calculated using the model of regular solutions for the solid phases and that of an ideal solution for the liquid phase [1988Hol, 1989Hol]. Figure 1 shows the calculated diagram modified to take into account the accepted melting point of TiN (3290°C instead of 2930°C). The eutectic temperature and eutectic composition depend experimentally on the dimension of both AlN and TiN particles [1991And] because of the possible formation of Ti1-xAlxN metastable solid solutions. Hard coatings prepared by the cathodic arc ion plating method allow to form a cubic solid solution Ti1-xAlxN (0 < x < 0.7) and a wurtzite type solid solution Ti1-xAlxN (0.8 < x < 1) [1991Ike, 1992Tan]. The existing experimental results and thermodynamic calculations lead to the so-called vapor deposition phase diagram, Fig. 2, [2001Spe]. Isothermal Sections According to the calculations of [1984Bey] AlN-TiN, TiAl3-AlN and TiAl3-TiN are stable tie lines in the Al-N-Ti ternary system at low temperatures. The sintering of Al with TiN powders leads to a hardening of the alloy due to the formation of AlN and TiAl3 during sintering [1992Koy]. Figures 3 and 4 show isothermal sections of the Al-N-Ti diagram at 900 and 1325°C, respectively. These sections were constructed using the experimental data and accounting for the formation of Ti4AlN3 [1997Dur, 2000Pro2]. JTi2N does not coexist with any of the ternary compounds [1984Sch].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–N–Ti
325
Temperature – Composition Sections Substantial solid solubility of nitrogen in (Ti) and (Ti) solid solutions has been reported by [1954Thy]. Unfortunately, these ternary phase boundary data do not match with the currently established phase boundaries in the Al-Ti and N-Ti binaries. Nitrogen raises the /(+) phase boundaries toward higher temperatures and widens the + field of the Al-Ti system. Thermodynamics Heat capacity of the Ti4AlN3- compound was measured between 2 and 10 K using a standard adiabatic calorimeter in a liquid helium cryostat [1999Ho]. It was determined that Cp = 0.00812T + 0.033#10 -3T3 J#mol-1#K-1 and the characteristic Debye temperature (D) equals 506°C. According to the data of [2000Bar1] Cp = 232 - 24350T-1 from 25 to 1030°C and D = 498°C (D = 489°C [2000Fin]. The molar heat capacity at room temperature is 150 J#mol-1#K-1 and increases monotonically with increasing temperature, reaching a plateau at 220 J#mol-1#K-1 at 1030°C. The Gibbs energy of formation of Ti2AlN at 850°C equals -135.5 kJ#mol-1 of atoms [1997Dur]. Notes on Materials Properties and Applications Ti2AlN is more wear-resistant than carbides of transition metals [1976Ivc1, 1977Ivc] and its abrasive ability gives up only on diamond, B4C, B and BN. Composites containing 30 vol.% Ti 2AlN and 70 vol.% TiAl have a high strength at both room and elevated temperatures and show some intrinsic compressive ductility at room temperature [1992Mab]. The yield strength and fracture stress increase with increasing nitrogen content in the TiAl phase [1991Kaw]. At room temperature Young’s (ERT) and shear ( RT) moduli and Poisson’s ratio of Ti4AlN 3- are 310 2, 127 2 GPa and 0.2 respectively [2000Fin, 2000Pro1]. This ternary nitride is relatively soft (Vickers hardness 2.5 GPa), lightweight (4.58 g#cm-3) and machinable [2000Pro1]. Increasing the Al content in the Ti1-xAlxN metastable films leads to an increase coating roughness and a change in color from gold to black-purple when the Al content increases from 13 to 27 mass% [1991Col]. Because of differences in chemical composition, the sputtered Ti1-xAlxN coatings show colors changing from metallic silver for low nitrogen coatings to a very dark blue for layers with high nitrogen contents [1986Jeh, 1988Jeh]; [1987Ina] indicates that these solid solutions in the composition range of 0.13 x 0.58 were greenish brown in color. These films have good decorative properties and excellent wear as well, [1986Kno, 1987Kno, 1988Ran, 1992Tan]. The incorporation of Al into the nitride films improves the oxidation resistance as well as the cutting performance of Ti1-xAlxN coated drills [1986Mue, 2001Spe]. It has been noted by [1990McI, 1991Ike, 1992Tan] that metastable single-phase polycrystalline Ti0.5 Al 0.5N alloy films exhibit much better high-temperature (750 - 900°C) oxidation resistance than polycrystalline TiN1-x films grown under similar conditions. It was found that Ti1-xAlxN films upon oxidation in air at 1000°C formed two-phase mixtures of TiO2 and Al 2O3 [1991Ike, 2001Hug]. The thickness of the oxide layer grown on these films decreases with increasing Al content in the films [2001Hug]. The electric resistivity of Ti1-xAlxN metastable solid solutions raised with increasing Al content [1987Ina]. Based upon resistivity and elevated-temperature interfacial reaction measurements, Ti1-xAlxN appears to be a promising candidate for improved diffusion-barrier layers between Al and Si [1993Pet]. Metastable Ti1-xAlxN coatings with the cubic NaCl structure are already being produced commercially for cutting tool applications [2001Spe]. When the amount of TiN particles was increased [2002Tan] the AlN-TiN composite materials showed an increasing Vickers hardness (adding 21 vol.% TiN to AlN-ceramics increased the hardness more than 15%), a decreasing fracture strength (20%) and a slightly increasing Young’s modulus (6%). Such composites with high content of AlN (> 20 vol.%) have a great thermic stability against cyclic heating and cooling in gas environments and in water [1976Kuz].
Landolt-Börnstein New Series IV/11A3
MSIT ®
326
Al–N–Ti
References [1954Thy]
[1963Jei] [1964Now]
[1976Ivc1]
[1976Ivc2]
[1976Kuz]
[1977Ivc]
[1981Bee]
[1982Wit] [1983Psh]
[1984Bey] [1984Jon] [1984Sch] [1985Sch]
[1986Jeh]
[1986Kau]
MSIT®
van Thyne, R.J., Kesler, H.D., “Influence of Oxygen, Nitrogen and Carbon on the Phase Relationships in the Ti-Al System”, Trans. AIME, J. Met., (2), 193-199 (1954) (Experimental, Equi. Diagram, 7) Jeitschko, W., Nowotny, H., Benesovsky, F., “Ti2AlN, a Nitrogen Containing H-Phase” (in German), Monatsh. Chem., 94(6), 1198-1200 (1963) (Experimental, Crys. Structure, 2) Nowotny, H., Jeitschko, W., Benesovsky, F., “Novel Complex Carbides and Nitrides and Their Relation to Phases of Hard Substances” (in German), Planseeber. Pulvermetall., 12, 31-43 (1964) (Experimental, Equi. Diagram, 18) Ivchenko, V.I., Kosolapova, T.Ya., “Investigation of Abrasive Properties of Ternary Compounda in the Systems Ti-Al-C and Ti-Al-N” (in Russian), Poroshk. Metall., (8), 56-59 (1976) (Experimental, Crys. Structure, Phys. Prop., 6) Ivchenko, V.I., Lesnaya, M.I., Nemchenko, V.F., Kosolapova, T.Ya., “Study of Preparation Conditions and Certain Physical Properties of the Ternary Compound TI2AlN” (in Russian), Poroshk. Metall., (4), 60-63 (1976) (Experimental, Crys. Structure, Phys. Prop., 6) Kuzenkova, M. A., Kislyi, P. S., Pshenichnaya, O. V., “The Structure and Properties of Composite Materials Based on the Nitrides of Ti, Zr and Al” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 12(3), 430-434 (1976) (Experimental, Equi. Diagram, Mechan. Prop., Phys. Prop., 8) Ivchenko, V. I., Kosolapova, T. Y., “Study of Preparation Conditions and Some Properties of Ternary Compounds in the Ti-Al-C and Ti-Al-N Systems” (in Russian), Nauchn. Trudy Moskov. Inst. Stali i Splavov, (99), 86-90 (1977) (Experimental, Crys. Structure, Phys. Prop., 11) Beensh-Marchwicka, G., Kròl-Stpniewska, L., Posadowski, W., “Structure of Thin Films Prepared by the Cosputtering of Titanium and Aluminium or Titanium and Silicon”, Thin Solid Films, 82(4), 313-320 (1981) (Experimental, Equi. Diagram, 10) Wittmer, M., “Interfacial Reactions Between Aluminium and Transition-Metal Nitride and Carbide Films”, J. Appl. Phys., 53(2), 1007-1012 (1982) (Experimental, Equi. Diagram, 16) Pshenichnaya, O.V, Verkhovodov, P.A., Kislyi, P.S., Kuzenkova, M.A., Goncharuk, A.B., “Test Methods and Properties of Powder Metallurgical Materials. Nitriding of the Intermetallic Compound TiAl3”, Sov. Powder Metall. Met. Ceram., (10), 851-855 (1983), transl. from Poroshk. Metall., (10), 76-80, 1983 (Experimental, Equi. Diagram, 9) Beyers, R., Sinclair, R., Thomas, M. E., “Phase Equilibria in Thin-Film Metallizations”, J. Vac. Sci. Technol., B2(4), 781-784 (1984) (Calculation, Equi. Diagram, 15) Jones, R.D., Rose, K., “Liquidus Calculations for III-N Semiconductors”, Calphad, 8(3), 343-354, (1984) (Equi. Diagram, Calculation, #, 28) Schuster, J.C., Bauer, J., “The Ternary System Titanium-Aluminium-Nitrogen”, J. Solid State Chem., 53, 260-265 (1984) (Experimental, Equi. Diagram, Crys. Structure, 24) Schuster, J.C., Bauer, J., Nowotny, H., “Applications to Materials Science of Phase Diagrams and Crystal Structures in the Ternary Systems Transition Metal-Aluminium-Nitrogen”, Rev. Chim. Miner., 22(4), 546-554 (1985) (Experimental, Equi. Diagram, Crys. Structure, 20) Jehn, H.A., Hofmann, S., Rueckborn, V.-E., Muenz, W.-D., “Morphology and Properties of Magnetron-Sputtered (Ti,Al)N Layers on High Speed Steel Substrates as a Function of Deposition Temperatures and Sputtering Atmosphere”, J. Vac. Sci. Technol., A4(6), 2701-2704 (1986) (Experimental, Crys. Structure, 22) Kaufman, M.J., Konitzer, D.G., Shull, R.D., Fraser, H.L.,“An Analytical Electron Microscopy Study of the Recently Reported ’Ti2Al Phase’ in -TiAl Alloys”, Scr. Metall., 20(1), 103-108 (1986) (Experimental, Crys. Structure, 13)
Landolt-Börnstein New Series IV/11A3
Al–N–Ti [1986Kno] [1986Mue] [1987Hak]
[1987Ina] [1987Kno] [1988Hol] [1988Jeh] [1988Pen]
[1988Ran]
[1989Hol] [1990McI]
[1991Adi]
[1991And]
[1991Col] [1991Hul]
[1991Ike]
[1991Kaw]
[1992Koy]
Landolt-Börnstein New Series IV/11A3
327
Knotek, O., Boehmer, M., Leyendecker, T, “On Structure and Properties of Sputtered Ti and Al Based Hard Compound Film”, J. Vac. Sci. Technol., A4(6), 2695-2700 (1986) Muenz, W. D., “Titanium Aluminium Nitride Films: A New Alternative to TiN Coatings”, J. Vac. Sci. Technol., A4, 2717-2721 (1986) (Experimental, Crys. Structure, 26) Håkansson, G., Sundaren, J.-E., McIntyre, D., Greene, J.E., “Microstructure and Physical Properties of Polycrystalline Metastable Ti0.5Al0.5N Alloys Grown by D.C. Magnetron Sputter Deposition”, Thin Solid Films, 153(1-3), 55-65 (1987) (Experimental, Crys. Structure, 17) Inamura, S., Nobugai, K., Kanamaru, F., “The Preparation of NaCl-type Ti1-xAlxN Solid Solutions”, J. Solid State Chem., 68(1), 124-127 (1987) (Experimental, Crys. Structure, 3) Knotek., O., Leyendecker, T, “On the Structure of (Ti,Al)N-PVD Coatings”, J. Solid State Chem., 70(2), 318-322 (1987) (Experimental, Crys. Structure, 12) Holleck, H., “Metastable Coatings - Prediction of Composition and Structure”, Surf. Coat. Technol., 36, 151-159 (1988) (Calculation, Equi. Diagram, 8) Jehn, H., Hofmann, S, Muenz, W.-D., “(Ti,Al)N Coatings-an Example of ’Ternary’ Nitride Hard Coatings”, Metall, 42(7), 658-669 (1988) (Experimental, Crys. Structure, 30) Penttinen, I., Molarius, J.M., Korhonen, A. S., Lappalainen, R., “Structure and Composition of ZrN and (Ti,Al)N Coatings”, J. Vac. Sci. Technol., 6(3), 2158-2161 (1988) (Experimental, Crys. Structure, 9) Randhawa, H., Johnson, P.C., Cunningham, R., “Deposition and Characterization of Ternary Nitrides”, J. Vac. Sci. Technol., 6(3), 2136-2139 (1988) (Experimental, Mechan. Prop., 6) Holleck, H., “Advanced Concepts in PVD Hard Coatings” (in German), Metall, 43(7), 614-624 (1989) (Experimental, Crys. Structure, 23) McIntyre, D., Greene, J.E., Hakansson, G., Sundaren, J.-E., Muenz, W. D., “Oxidation of Metastable Single-Phase Polycrystalline Ti0.5Ai0.5N Films: Kinetics and Mechanisms”, J. Appl. Phys., 67(3), 1542-1553 (1990) (Experimental, Crys. Structure, 38) Adibi, F., Petrov, I., Hultman, L., Wahlstroem, U., Shimizu, T., McIntyre., D., Green., J.E., Sundgren, J.-E., “Defect Structure and Phase Transitions in Epitaxial Metastable Cubic Ti 0.5Al0.5N Alloys Grown on MgO(001) by Ultra-High-Vacuum Magnetron Sputter Deposition”, J. Appl. Phys., 69(9), 6437-6450 (1991) (Experimental, Crys. Structure, 34) Andrievskii, R.A., Anisimova, N.A., “Phase Diagram Calculations for Titanium Nitride-Based Pseudobinary Nitride Systems”, Inorg. Mat., 27(7), 1220-1223 (1991), transl. from Izv. Akad. Nauk SSSR, Neorg. Mater., 27(7), 1450-1453 (1991) (Calculation, Equi. Diagram, 17) Coll, B.F., Fontana, R., Gates, A., Sathrum, P., “(Ti-Al)N Advanced Films Prepared by Arc Process’, Mater. Sci. Eng., A140, 816-824 (1991) (Experimental, Mechan. Prop., 12) Hultman, L., Hakansson, G., Wahlstroem, U., Sundaren, J.-E., Petrov, I., Adibi, F., Green, J.E., “Transmission Electron Microscopy Studies of Microstructural Evolution, Defect Structure and Phase Transitions in Polycrystalline and Epitaxial Ti1-xAlxN and TiN Films Grown be Reactive Magnetron Sputter”, Thin Solid Films, 205(2), 153-164 (1991) (Experimental, Crys. Structure, 45) Ikeda, T., Satoh, H., “Phase Formation and Characterization of Hard Coatings in the Ti-Al-N System Prepared by the Cathodic Arc Ion Platting Method”, Thin Solid Films, 195(1-2), 99-110 (1991) (Experimental, Crys. Structure, 17) Kawabata, T., Tadano, M., Izumi, O., “Effect of Carbon and Nitrogen on Mechanical Properties of TiAl Alloys”, ISIJ International, 31(10), 1161-1167 (1991) (Experimental, Equi. Diagram, Mechan. Prop., 45) Koyama, K., Morishita, M., Suzuki, K., Yagi, S., “A New Ternary Al-Ti-N Alloy Prepared by the Reaction Sintering Process” (in Japanese), J. Japan. Soc. Powder Powder Metall., 39(10), 823-829 (1992) (Experimental, Equi. Diagram, 12)
MSIT ®
328 [1992Mab]
[1992Rog]
[1992Sch] [1992Tan]
[1993Adi]
[1993Jeh]
[1993Pet]
[1993Tan]
[1993Wah]
[1995Wu] [1997Dur]
[1997Lee] [1998Bar] [1998Che] [1998Lee] [1999Bar]
[1999Far]
MSIT®
Al–N–Ti Mabuchi, H., Tsuda, H., Nakayama, Y., “Processing of TiAl-Ti2AlN Composites and their Compressive Properties”, J. Mater. Res., 7(4), 894-900 (1992) (Experimental, Mechan. Prop., 21) Rogl, P., Schuster, J.C., “Ti-B-N (Titanium - Boron - Nitrogen)” in “Phase Diagrams of Ternary Boron Nitride and Silicon Nitride Systems” (Monogr. Ser. of Alloy Phase Diag.), Materials Park, Ohio: Materials Informations Soc., 103-106 (1992) (Review, Equi. Diagram., Crys. Structure, Thermodyn., #, *, 19) Schuster, J.C., “System Aluminium - Nitrogen - Titanium: Summary of Constitution Data”, Int. Report, (1992) (Review, Equi. Diagram, 18) Tanaka, Y., Guer, T.M., Kelly, M., Hagstrom, S.B., Ikeda, T., Wakihira, K., Satoh, H., “Properties of (Ti1-xAlx)N Coating Tools Prepared by the Cathodic Arc Ion Plating Method”, J. Vac. Sci. Technol., A10(4), 1749-1756 (1992) (Experimental, Crys. Structure, 21) Adibi, F., Petrov, I., Green., J.E., Wahlstroem, U., Sundaren, J.-E., “Design and Characterization of a Compact Two-Target Ultrahigh Vacuum Magnetron Sputter Deposition System: Application to the Growth of Epitaxial Ti1-xAlxN Alloys and TiN/Ti1-xAlxN Superlattices”, J. Vac. Sci. Technol., A11(1), 136-142 (1993) (Experimental, Crys. Structure, 25) Jehn, H.A., “Aluminium-Nitrogen-Titanium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.13521.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 29) Petrov I., Mojab, E., Adibi, F., Greene, J.E., Hultman, L., Sundgren, J.-E., “Interfacial Reactions in Epitaxial Al/Ti1-xAlxN (0 x 0.2) Model Diffusion-Barrier Structure”, J. Vac. Sci. Technol., A11(1), 11-17 (1993) (Experimental, Crys. Structure, 25) Tanaka, Y., Guer, T.M., Kelly, M., Hagstrom, S.B., Ikeda, T., “Strusture and Properties of (Ti1-xAlx)N Films Prepared by Reactive Sputtering”, Thin Solid Films, 228(1-2), 238-241 (1993) (Experimental, Crys. Structure, 13) Wahlstroem, U., Hultman, L., Sundgren, J.-E., Adibi, F., Petrov, I., Greene, J.E., “Crystal Growth and Microstructure of Polycrystalline Ti(1-x)AlxN Alloy Films Deposited by Ultra-High-Vacuum Dual-Target Magnetron Sputtering”, Thin Solid Films, 235(1-2), 62-70 (1993) (Experimental, Crys. Structure, 32) Wu, Z.L., Pope, D.P., Vitek, V., “Ti2NAl in L12 Al3Ti-Base Alloys”, Metall. Mater. Trans., A26(3), 521-524 (1995) (Experimental, Crys. Structure, 15) Durlu, N., Gruber, U., Pietzka, M.A., Schmidt, H., Schuster, J.C., “Phases and Phase Equilibria in the Quaternary System Ti-Cu-Al-N at 850°C”, Z. Metallkd., 97(5), 390-400 (1997) (Experimental, Review, Crys. Structure, Equi. Diagram, 32) Lee, H.D., Petuskey, W.T., “New Ternary Nitride in Ti-Al-N System”, J. Am. Ceram. Soc., 80(3), 604-608 (1997) (Experimental, Crys. Structure, 8) Barsoum, M.W., Schuster, J.C., “Comment on “New Ternary Nitride in Ti-Al-N System”, J. Am. Ceram. Soc., 81(3), 785-786 (1998) (Experimental, Crys. Structure, 10) Chen, G., Sundman, B., “Thermodynamic Assessment of the Ti-Al-N System”, J. Phase Equilib., 19(2), 146-160 (1998) (Assessment, Equi. Diagram, Thermodyn., 42) Lee, B.-J., “Predictive Analysis of Ti/AlN Interfacial Reaction Using Diffusion Simulation”, Scr. Mater., 38(3), 499-507 (1998) (Calculation, Equi. Diagram, 15) Barsoum, M.W., Farber, L., Levin, I., Procopio, A., El-Raghy, T., Berner, A., “High-Resolution Transmission Electron Microscopy of Ti 4AlN 3, or Ti3Al2N2 Revisited”, J. Am. Ceram. Soc., 82(9), 2545-2547 (1999) (Experimental, Crys. Structure, 23) Farber, L., Levin, I., Barsoum, M.W., El-Raghy, T., Tzenov, T., “High-Resolution Transmission Electron Microscopy of Some Tin+1 AXn Compounds (n = 1, 2; A = Al or Si; X = C or N)”, J. Appl. Phys., 86(5), 2540-2543 (1999) (Experimental, Crys. Structure, 23)
Landolt-Börnstein New Series IV/11A3
Al–N–Ti [1999Ho]
[1999Mag] [2000Bar1]
[2000Bar2]
[2000Fin]
[2000Gam] [2000Par]
[2000Pro1] [2000Pro2]
[2000Raw]
[2001Hug]
[2001Per]
[2001Spe]
[2002Tan]
[2003Fer]
[2003Sch]
Landolt-Börnstein New Series IV/11A3
329
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) Magnan, J., Weatherly, G.C., Cheynet, M.-C., “The Nitriding Behavior of Ti-Al Alloys at 1000°C”, Metall. Mater. Trans. A, 30A(1), 19-29 (1999) (Experimental, Equi. Diagram, 27) Barsoum, M.W. Rawn, C.J., El-Raghy, T., Procopio, A.T., Porter, W.D., Wang, H., Hubbard, C.R., “Thermal Properties of Ti4AlN3”, J. Appl. Phys., 87(12), 8407-8414 (2000) (Experimental, Crys. Structure, Phys. Prop., 33) Barsoum, M.W., Ali, M., El-Raghy, T., “Processing and Characterization of Ti2AlC, Ti 2AlN and Ti2AlC0.5N0.5”, Metall. Trans. A, 31A(7), 1857-1865 (2000) (Experimental, Crys. Structure, Phys. Prop., 36) Finkel, P., Barsoum, M.W., El-Raghy, T., “Low Temperature Dependencies of the Elastic Properties of Ti4AlN3, Ti3Al1.1C1.8, and Ti3SiC2”, J. Appl. Phys., 87(4), 1701-1703 (2000) (Experimental, Mechan. Prop., 22) Gamarnik, M.Y., Barsoum, M.W., El-Raghy, T., “Improved X-Ray Powder Diffraction Data for Ti2AlN”, Powder Diffr., 15(4), 241-242 (2000) (Experimental, Crys. Structure, 7) Paransky, Y., Gotman, I., Gutmanas, E.Y., “Reactive Phase Formation at AlN-Ti and AlN-TiAl Interfaces”, Mater. Sci. Eng. A, A277, 83-94 (2000) (Experimental, Equi. Diagram, 28) Procopio, A.T., Barsoum, M.W., El-Ragny, T., “Characterization of Ti4AlN3”, Metall. Mater. Trans. A, 31A(2), 333-337 (2000) (Experimental, Crys. Structure, Phys. Prop., 24) Procopio, A.T., El-Raghy, T., Barsoum, M.W., “Synthesis of Ti4AlN3 and Phase Equilibria in the Ti-Al-N System”, Metall. Mater. Trans. A, 31A(2), 373-378 (2000) (Experimental, Equi. Diagram, Crys. Structure, 24) Rawn, C.J., Barsoum, M.W., El-Raghy, T., Procopio, A., Hoffmann, C.M., Hubbard, C.R., “Structure of Ti4AlN3 - A Layered Mn+1 AXn Nitride”, Mater. Res. Bull., 35, 1785-1796 (2000) (Experimental, Crys. Structure, 14) Hugon, M.C., Varniere, F., Letendu, F., Agius, B., Vickridge, I., Kingon, A.I., “18O Study of the Oxidation of Reactively Sputtered Ti1-xAlxN Barrier”, J. Mater. Res., 16(9), 2591-2599 (2001) (Experimental, Crys. Structure, Phys. Prop., 24) Perdix, F., Trichet, M.-F., Bonnentien, J.-L., Cornet, M., Bigot, J., “Influence of Nitrogen on the Microstructure and Mechanical Properties of Ti-48Al Alloy”, Intermetallics, 9, 147-155 (2001) (Experimental, Equi. Diagram, 19) Spencer, P.J., “Computational Thermochemistry: from its Early Calphad Days to a Cost-Effective Role in Materials Development and Processing”, Calphad, 25(2), 163-174 (2001) (Calculation, Equi. Diagram, 31) Tangen, I.-L., Grande, T., Yu, Y.D., Hoier, R., Einarsrud, M.-A., “Preparation and Mechanical Characterisation of Aluminium Nitride-Titanium Nitride and Aluminium Nitride-Silicon Carbide Composites”, Key Eng. Mater., 206-213, 1153-1156 (2002) (Experimental, Mechan. Prop., 2) Ferro, R., Bochvar, N., Sheftel, E., Ding, J.J., “Al-N (Aluminum-Nitrogen)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 33) Schmid-Fetzer, R., “Al-Ti (Aluminium-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 85)
MSIT ®
Al–N–Ti
330 Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Al) < 660.452
cF4 Fm3m Cu
a = 404.96
at 25°C [Mas2]
(N) < -237.54
cP8 Pa3 N
a = 566.1
[Mas2]
(Ti) 1670 - 882
cI2 Im3m W
a = 330.65
[Mas2] dissolves up to 6.2 at.% N at 2020°C dissolves up to 44.8 at.% Al at 1490°C
(Ti) < 882
hP2 P63/mmc Mg
a = 295.06 c = 468.35
at 25°C [Mas2] dissolves up to 23 at.% N at 1050°C dissolves up to 51.8 at.% Al at 1463°C
AlN < 2434.7
hP4 P63mc ZnS (wurtzite)
a = 311.14 c = 497.92
at 25°C [2003Fer]
Ti2N < 1100
tP6 P42/mnm TiO2
a = 494.52 c = 303.42
at 33 to 34 at.% N [V-C2]
, TiN 1-x < 3290
cF8 Fm3m NaCl
a = 423.9 0.1
[V-C2] From 28 at.% N at 2350°C to > 50 at.% N
Ti3N2-x 1103 - 1066
hR2 ? VTa2C2
a = 297.95 c = 2896.5
at 29 at.% N [1992Rog]
Ti4N3-x 1291 - 1078
hR2 ? V4C3
a = 298.09 c = 2166.42
at 31.5 at.% N [1992Rog]
Ti3Al 1164
hP8 P63/mmc Ni3Sn
a = 580.6 c = 465.5 a = 574.6 c = 462.4
at 22 at.% Al [2003Sch]
tP4 P4/mmm AuCu
a = 400.0 c = 407.5 a = 398.4 c = 406.0
at 50.0 at.% Al, [2003Sch]
TiAl2 < 1199
tI24 I41/amd HfGa2
a = 397.0 c = 2497.0
[2003Sch]
“Ti2Al5” 1416 - 990
tP28 P4/mmm “Ti2Al5”
a = 390.53 c = 2919.63
[2003Sch]
TiAl < 1463
MSIT®
at 38 at.% Al [2003Sch]
at 62.0 at.% Al, [2003Sch]
Landolt-Börnstein New Series IV/11A3
Al–N–Ti
331
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
TiAl3(h) < 1393
tI8 I4/mmm TiAl3(h)
a = 384.9 c = 860.9
[2003Sch]
TiAl3(l) < 950 (Ti-rich)
tI32 I4/mmm TiAl3 (l)
a = 387.7 c = 3382.8
[2003Sch]
* -1, Ti3AlN
cP5 ? CaTiO3
a = 411.20 a = 411.70 0.07
[1984Sch, 1985Sch] [1992Sch]
* -2, Ti2AlN
hP8 P63/mmc Cr2AlC
a = 298.9 c = 1361.4 a = 299.9 c = 1365.0 a = 300.9 c = 1365.0
at 25°C [2000Bar2]
* -3, Ti4AlN 3
* Ti1-xAlxN metastable
Landolt-Börnstein New Series IV/11A3
hP16 P63/mmc Ti4AlN 3
cF8 Fm3m NaCl
a = 299.05 0.01 c = 2338.0 0.1 a = 300.45 0.02 c = 2348.1 0.2 a = 302.22 0.02 c = 2360.8 0.2 a = 298.80 0.02 c = 2337.2 0.2 a = 299.10 0.02 c = 2339.6 0.1 a = 424 a = 422.6 a = 420.6 a = 419.9 a = 416.9 a = 416
at 400°C [2000Bar2] at 800°C [2000Bar2] Nitrogen deficient Ti4AlN 3-x [1999Bar] at 25°C [2000Bar1] at 570°C [2000Bar1] at 1094°C [2000Bar1] Ti4AlN2.78 , neutron powder diffraction [2000Raw] X-ray powder diffraction [2000Raw] at x = 0.1 [1993Pet] at x = 0.2 [1993Pet] at x = 0.3 [1993Tan] at x = 0.42 [1993Tan] at x = 0.5 [1993Tan] at x = 0.7 [1991Ike]
MSIT ®
Al–N–Ti
332
3290°C 3250
Fig. 1: Al-N-Ti. Calculated phase diagram of the AlN - TiN pseudobinary system
3000
L
TiN+L
2800°C
2750
L+AlN 2500
2500°C
TiN
Temperature, °C
2250
AlN
2000 1750 1500
TiN+AlN 1250 1000 750 500 250
Ti Al N
50.00 0.00 50.00
10
20
30
40
Al, at.%
Ti Al N
0.00 50.00 50.00
Ti Al N
0.00 50.00 50.00
900
Fig. 2: Al-N-Ti. Metastable TiN - AIN phase diagram
800
cubic+hexagonal 700
Temperature, °C
600
500
(Ti,Al)N cubic
(Al,Ti)N hexagonal
400
300
200
100
Ti Al N
MSIT®
0
50.00 0.00 50.00
10
20
30
Al, at.%
40
Landolt-Börnstein New Series IV/11A3
Al–N–Ti
333
N
Data / Grid: at.%
Fig. 3: Al-N-Ti. Isothermal section at 900°C
Axes: at.%
20
80
40
60
AlN TiN1-x 60
40
Ti2N
τ2
80
20
τ1
(αTi) 20
Ti (βTi)
40
Ti3Al
TiAl
60
TiAl2
L
TiAl3 80
N
Al
Data / Grid: at.%
Fig. 4: Al-N-Ti. Isothermal section at 1325°C
Axes: at.%
20
80
40
60
AlN
60
TiN1-x
τ3
40
τ2 80
20
τ1
(α Ti)
Ti
Landolt-Börnstein New Series IV/11A3
(β Ti) 20
40
TiAl
60
Ti2Al5
80
TiAl3
L
Al
MSIT ®
334
Al–Nb–Ti
Aluminium – Niobium – Titanium Ludmila Tretyachenko Literature Data Titanium aluminide based alloys are candidate materials for high temperature structural applications; among alloying elements particularly niobium is expected to exert a favorable influence on low temperature ductility. Data on phase equilibria in the Al-Nb-Ti system are a prerequisite to promote the development of appropriate alloys. The status of investigations in the Al-Nb-Ti system was summarized by [1993Gam] in a critical assessment comprising all constitution-relevant literature data up to 1990. As a result the experimental data by [1990Hel] and [1990Per] were chosen for the liquidus projection as well as for the isothermal sections at 1200 and 1000°C. Furthermore, a liquidus projection and several isothermal sections calculated by [1992Kat1] were given. Phase relations at that time were characterized by a series of equilibrium phases: (i) a wide region of the bcc disordered solid solution (Ti,Nb,Al) [1962Pop, 1970Nar, 1972Nar, 1974Nar, 1983Tro, 1984Zak, 1989Jew, 1989Kal, 1990Hel, 1990Per], which transforms to an ordered ternary solid solution phase (B2 or 0) in a wide range of compositions [1987Ban, 1989Ben, 1990Hel, 1990Per]; (ii) extended solid solution phases on the base of binary compounds TiAl ( ), Nb3Al ( ) and Nb2Al ()), (iii) a continuous solid solution between the binary boundary phases TiAl3 and NbAl3 (from now on designated as J), (iv) solid solutions based on the low-temperature modification of titanium Ti () and the Ti3Al based phase (2) as well as (v) the ternary compounds Ti2NbAl (so-called O phase, discovered by [1988Ban]) [1990Moz, 1990Mur, 1990Wey, 1991Ben] and Ti4NbAl3 with the Ni2In type (B82) [1990Ben1, 1990Ben2, 1991Ben]. Although the general features of the phase relations remained unchanged, new investigations refined various details in the constitution of the ternary system and furthermore solved a series of controversies, which essentially concerned (a) the stability of ternary phases and (b) the extension of solid solution phases. A listing of recent and some earlier experiments and the techniques used is presented in Table 1. One of the problems is linked to the two ternary phases, T1 (Ti-18Nb-34Al) and T2 (Ti-11Nb-44Al), reported by [1989Jew] in an isothermal section at 1200°C, which turned out to be part of ternary solutions: T1 was shown to have the structure of the ordered bcc phase (B2 or 0 in this assessment [1990Per]), whilst the T2 phase was supposed to be an isolated region of the same phase. The authors of [1990Per, 1990Ben2, 1990Kno, 1990Mis, 1990Wey] meanwhile agree that due to numerous phase transformations alloys in the area of T1 and T2 are very sensitive to composition, temperature and the cooling rate. The second problem is related with the so-called 1 phase. Although the TiNbAl3 ( 1) phase was reported in the Ti-NbAl3 section by [1962Pop, 1983Tro, 1984Zak], which in the review by [1984Arg] was assumed to be pseudobinary, the 1 phase was, however, not observed by the authors of [1989Jew, 1989Kal, 1990Per]. A study of diffusion couples at 1000°C [1990Hao1, 1990Hao2] again was interpreted in terms of two ternary compounds, TiNbAl3 ( 1) and Ti5NbAl2 each with a large solubility range. Whilst the second phase is to be identified with the O phase, Ti2NbAl [1988Ban, 1989Kes], the existence of the 1 phase was denied in an investigation of partial isothermal sections at 1100, 900 and 800°C [1991Smi, 1992Smi, 1991Zak, 1992Zak, 1992Pav1, 1992Pav2]. Nevertheless, claim for the existence of the 1 phase was again raised by [1993Zha] and [1994Che1] and a model of its crystal structure was reported by [1994Che2, 1994Wan]. Furthermore 1 phase fields were shown in the isothermal sections at 1000, 1150 and 1400°C by [1996Che]. Despite [1997Jew] studied in detail the alloy Ti-23Nb-51Al (which was prepared by arc melting, annealed at 1200°C for 180 h and then at 1150°C for 50 h and water quenched) by backscattered electron imaging (BSEI), energy dispersive X-ray analysis (EDX) and XRD, and could not confirm the existence of 1, the previous authors did not agree with the comment of [1997Jew] and again presented (i) the 1 phase in their isothermal section at 1400°C [1998Wan], (ii) in refined versions of the sections at 1000 and 1150°C [1998Din] and (iii) the crystal structure of the 1 phase [1998Che]. However, in a detailed reinvestigation MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
335
of the isothermal sections at 1200 and 1000°C by [1998Hel] employing optical microscopy, EMPA, TEM and XRD on diffusion couples and bulk samples neither TiNbAl3 ( 1) nor the phases T1 or T2 could be traced. No other ternary compounds were observed. Considerable solid solubility of the third element in the most of the binary phases were confirmed and refined. A separate area of the ordered B2 phase was detected at 1000°C. The third problem area covers (i) the so-called (orthorhombic) O phases near the composition Ti2NbAl and (ii) the Ti4NbAl3 phase. As these problems are related to the crystallography of the phases mentioned, a detailed discussion is included in the section “Solid Phases”. Vaporization of solid alloys has been studied by Knudsen-effusion mass spectroscopy in the temperature range between 897 and 1362°C to derive Ti, Al partial pressures and thermodynamic activities of Ti and Al, partial enthalpies and entropies of mixing at 1200°C. Data on the phase compositions of 25 alloys in the range adjoining to the Al-Ti side and containing up to ~30 at.% Nb are given for 1200, 1100 and 1000°C [1999Eck]. Based on the experimental phase equilibria and thermodynamic data, thermodynamic assessments of the Al-Nb-Ti system were performed by [1998Ser] using the Redlich - Kister polynomial to describe the excess Gibbs energies of liquid, bcc and hcp phases. The intermetallic compounds, which exhibit a homogeneity range, were modeled using two or three sublattices. The sublattice model was also used to describe the order - disorder transformations D019 - hcp and A2 - B2. Both O1 and O2 forms were modeled as separate phases with two and three sublattices, respectively. As a result a liquidus projection has been calculated, as well as partial isothermal sections of the Nb rich corner at 700, 900 and 1200°C (in weight fractions), isothermal sections at 700, 800, 900, 1000, 1020, 1060, 1100, 1150, 1175, 1200, 1400 and 1650°C (in at. fractions) the isopleth at 27.5 at.% Al up to 35 at.% Nb. The representation of the thermodynamic properties of two states of the orthorhombic phase, ordered O1 and disordered O2, with a unique function was proposed by [2001Ser]. Two other models were proposed for thermodynamic modeling of the orthorhombic phase. The two sets of thermodynamic parameters obtained according to both models were used to calculate the isothermal sections at 990 and 700°C. Fields of B2 and bcc phase stability in the isothermal section at 1000°C were calculated using the CPA-GPM (coherent potential approximation - generalized perturbation method) within the cluster variation method (CVM) [1993Rub] and with application of linear muffin-tin orbitals (LMTO) [1995Rub]. The CVM in the irregular tetrahedron approximation was furthermore used to calculate the limits of the B2 phase field at 800, 1000, 1200 and 1400°C [1996Jac, 1999Cha1] and in the vertical section at 50 at.% Ti and 50 at.% Nb [1996Jac]. The results obtained were proven by experimental studies [1999Cha2]. The results of [1993Rub] and [1999Cha1] were included in a review by [2001Col] and used for the mixed CVM-CALPHAD method to calculate the phase equilibria in ternary system (isothermal section at 1000°C). [2001Kan] applied the CVM in the octahedron and tetrahedron approximation to calculate the /2 and /2 phase equilibria at 1000°C. The grand potential approach was applied to obtain thermodynamic parameters used to calculate the / and / phase equilibria at 1150 and 1400°C [2001Li1]. Binary Systems The Al-Nb and Nb-Ti systems are accepted from [Mas2] and [1987Mur], respectively. A critical assessment of the Al-Ti phase diagram is due to [2003Sch]. The version accepted therein and in [1993Oka1] is primarily based on the work of [1992Kat2], which is in essential agreement with recent data by [1996Tre]. However, the Ti5Al11 stoichiometry was shown in the latter phase diagram. Recently the Al-rich part of the system has been reinvestigated by [2001Bra], who also has shown the Ti5Al11 phase to exist. Solid Phases Data of the solid phases in the Al-Nb-Ti system are given in Table 2. The bcc solid solution () exists in a wide range of composition up to 40 at.% Al [1995Zdz, 1998Hel, 2000Leo2, 2002Leo1]. The transformation of the disordered (A2) phase to ordered 0 (B2) has been observed by many research groups [1987Ban, 1989Ben, 1992Men, 1994Hou, 1996Men, 1996Vas, 1998Rho, 1999Cha2] and others. The transition temperatures were shown to be sensitive to composition [1999Cha2] with the highest ordering temperatures Landolt-Börnstein New Series IV/11A3
MSIT ®
336
Al–Nb–Ti
(>1600°C, close to melting temperature) found for alloys in the vicinity of Ti2NbAl [1996Vas]. There is still some discrepancy on the ordering temperature for the Ti2NbAl alloy for which [1989Ben] reported a temperature higher than 1400°C, [1990Hel] estimated 1000 to 1200°C but [1999Cha2] from in situ neutron diffraction recorded only 1182 5°C. The A2B2 transformation temperature decreases from Ti2NbAl toward the Nb corner. The ordered 0 phase can be obtained in metastable form on quenching from the high temperature field and decomposes at aging. There are two well established ternary phases with the stoichiometries Ti2NbAl [1988Ban, 1990Moz] and Ti4NbAl3 [1990Ben1, 1990Ben2, 1992Ben, 1996Sad, 2000Sad]. The orthorhombic Ti2NbAl based phase O arises from the phase as a result of a sequence of phase transformations [1992Mur1, 1994Ben2, 1995Mur2, 1999Boe, 2001Sad, 2003Sad]. The formation of the O phase was suggested to occur immediately from the B2 (0) phase [1989Kes, 1991Ben, 1994Ben2] as well as through the peritectoid reaction 0+2O [1992Mur1, 1992Mur2, 1995Mur1, 1999Boe, 2003Sad] below ~1000°C. The orthorhombically distorted phase was observed at the 2/0-interface with the same composition and site occupancy as the 2 phase; as a similar structure has been obtained in hydrogenated Ti3Al-Nb alloys, the authors [1990Mur] concluded that the O phase appears as a result of hydrogen absorption during thin foil preparation in the acid-containing electrolyte. The homogeneity range of the O phase extends preferably at constant Al content of 26 - 27 at.%. The orthorhombic phase was shown to exist in two forms, O1 and O2, with crystal lattice of the same space group and lattice parameters, but with different site occupancies [1990Mur, 1990Wey, 1992Mur2, 1995Mur1, 1995Mur2, 2002Wu]. In the O1 form, which was observed to exist at higher temperatures from ~1000 down to ~900°C, Ti and Nb atoms randomly occupy the same sublattice (as in hexagonal 2), while Nb atoms occupy a distinctive sublattice in the O2 form detected at temperatures below 900°C. The transformation between these two forms was shown to be reversible. A first order transition was suggested for the O1O2 reaction [1995Mur2]. A very “weak” first order transition was predicted using the Bragg Williams model [2002Wu]. As to the O1 phase, it was suggested that the disordered orthorhombic martensite observed in the binary Nb-Ti system can be stabilized into an equilibrium phase at certain Al and Nb contents in the ternary Al-Nb-Ti system around the Ti2NbAl composition [1995Mur1]. [1994Ben1, 1994Ben2] have outlined possible paths for the constant composition coherent transformation of bcc Ti2NbAl high temperature phases to the hcp or orthorhombic low temperature phases employing crystallographic group-subgroup relations. The Ni2In (B82) type phase Ti4NbAl3 is formed from the CsCl (B2) type phase at ~900°C. This phase was found in the Ti-20Nb-30Al alloy annealed at 900°C by [1992Ben], however, in an in situ neutron diffraction of the Ti-12.9Nb-36.5Al by [2000Sad] it was only revealed at temperatures at or below 800°C. A thermodynamic calculation of the phase transformation in the Ti-10.8Nb-36.9Al alloy yielded Ti4NbAl3 below 1060°C [1996Sad]. The transformation of the B2 phase to Ti4NbAl3 involves the formation of metastable 7´´ with trigonal structure [1990Ben1, 1990Ben2, 1990Sho, 1996Sad]. From TEM-analysis [1990Ben2] reported also a new phase with a tripled hexagonal lattice for which he assumed further substitutional ordering of the B82 type phase in terms of either a possible Ti5Ga4 type phase with 18 atoms per unit cell and (Ti3Al3)(AlNb2) stoichiometry or in terms of the Mn5Si3 type structure (16 atoms/u.c.) with (Ti3Al3)Nb2 stoichiometry. A phase with hexagonal structure (a = 579 pm, c = 1409 pm) was found in the as cast alloy Ti4NbAl4 by means of TEM [1995Zdz]. It was supposed to be a superstructure of 2. Formation of metastable phases ´ and ´´ during rapid cooling was observed in Ti3Al-Nb alloys containing up to 5 at.% Nb [1988Str, 1990Wey, 1995Xu]. At higher Nb contents various metastable 7 related phases, both athermal and isothermal, have been detected in alloys rapidly cooled from high temperatures or aged at ~350-550°C [1978Zak, 1982Str, 1988Str, 1991Li, 1992Hsi1, 1992Sur, 2000Leo2, 2000Sad, 2001Sad] and others, as well as in Nb-Ti alloys with low Al content [1992Voz, 1996Men]. The 7´ and 7´´ phases are described by [1990Ben1] as two configurations of the same trigonal P3ml phase. They are related to the ordered B2 type phase and are distinguished by site occupancies. The 7´ modification is considered as the idealized state with the B2 chemical order inherited in a diffusionless transition. The chemical order in the 7´´ configuration is changed but the space group is the same. This configuration is more stable. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
337
The 1 phase (TiNbAl 3), which earlier was reported by [1962Pop, 1983Tro, 1984Arg, 1984Zak] and not confirmed by [1989Kal, 1989Jew, 1990Hel, 1990Per], was again reported by [1990Hao1, 1990Hao2, 1994Che2, 1994Wan, 1996Che, 1998Wan] and even later by [1998Che, 1998Din] in spite of the fact that [1997Jew] once more disproved the existence of this phase. The crystal structure of 1 was identified as tetragonal with a = 558 to 584 pm, c = 815 to 845 pm and has been considered as a superstructure of the L10 structure of (TiAl). The transformation (L1 0) 1 was suggested to be a continuous ordering process taking place with increasing Nb content in the (TiAl) phase. The ordering process has been presumed to proceed as a first order transition at 1000 and 1150°C but as a second order transition at 1400°C. The ordered 1 phase has been considered to be formed at the Nb content of 18 at.%, when Nb atoms occupy a specific sublattice. A possible relation of 1 with Ti2Al5 (tP32, P4/mbm) [1982Mii] was suggested. [1993Jac] detected in the Ti-20.3Nb-42.9Al alloy a high temperature phase with a Ti3Cu (L6 0) type lattice, a modification of the CuAu (L10) type lattice. The L6 0 phase (tP4, space group P4/mmm) was found to differ from the L10 phase in the site occupancy and was suggested to be an intermediate phase between the high temperature phase and the lower temperature and or 2 phases. Besides major amounts of the (Ti,Nb)Al3 phase, [1991Spa] claimed the formation of a cubic Cu 3Au (L1 2) type phase (composition Ti27.8Nb12.3 Al60.9 , a = 397.8 pm) using XRD, SEM and EMPA on the Ti-12Nb-63Al alloy arc melted and annealed at 1200°C for 16 h. However, [1993Nak2] from optical microscopy, XRD and SEM analyses did not confirm the L12 phase in the Ti-8Nb-67Al alloy sintered at 1150°C for 24 h. A metastable ordered tetragonal transition phase T with a composition of Ti5NbAl2 arising during the B2 to 2 transition in a plasma sprayed Ti-11Nb-24Al alloy after aging for 10 min at 650°C was reported by [1992Hsi1, 1992Hsi2, 1992Hsi3]. On prolonged aging the T phase transforms to an ordered O phase and further to 2. The phase was detected by means of XRD, SEM and TEM. The crystal structure of the T phase was found to be similar to the D03 type structure but with a tetragonal distortion (P4/mmm, a = 650 10 pm, c/a 1.02) and structural relationships and habit plane between T and O phases were established. [1994Ban] analyzed the diffraction patterns obtained by [1992Hsi2, 1992Hsi3] and found that they were not consistent with the proposed structure, but can be attributed to the structure of the metastable O phase proposed by [1990Moz]. Recently a new phase evolution path during aging at 650°C was proposed to be B2B19O´´O´2, with O´´ and O´ phases instead of T and O (Ti2AlNb) phases involved in the previous phase evolution path: B2TO2 [1995Hsi]. The phases taking part in the newly proposed phase transition sequence were the following: B2 (Pm3m), a = 325; B19 (Pmmm), a = 325, b = c = 460; O´´ (Cmcm, previously T), a = 660, b = 920, c = 460; O´ (Cmcm, previously O), a = 605 b = 980, c = 473; 2 (P63/mmc), a = 580, c = 465 (the lattice parameters in pm). The B19 and O´´´ phases can only be resolved with difficulties owing to overlapping peaks and weak reflection intensity, however a tetragonal distortion of the B2 phase was detected. A novel tetragonal phase, designated as , was observed by [2000Leo1] in Ti-Nb-40Al alloys (Ti from 24 to 36 at.%) aged below 1000°C for times up to 3600 h followed by water quenching. Phase identification was performed by XRD and EMPA. Convergent beam electron diffraction yielded a bct cell and space group I41/amd, a = 510.6 pm, c = 2816.8 pm, on the basis of which indexing of the X-ray powder pattern was satisfactory. The composition was evaluated as ~25Ti-45Nb-30Al and orientation relationships between and (TiAl) were determined. From the low concentration of elements (O, N, C) interstitial contamination was ruled out. The phase was reported to be thermodynamically stable. A hydride phase with the same crystal structure and nearly the same lattice parameters as the phase was observed to replace the 2 phase in a Ti-48Al-2Cr-2Nb duplex alloy at hydrogen charging for 60 h at 12.8 MPa and 800°C. However, there is no reason to suppose a high content of H in the studied samples [2000Leo1], in particular, for crushed powder samples analyzed by means of XRD. Precipitates, which occurred in the single phase alloy Ti-5Nb-54Al containing < 900 ppm O 2, were shown by SEM and EDS analysis to be a cubic ternary Al-O-Ti compound with a = 690 pm [2001Cao]. A stress induced orthorhombic 9R phase was observed at incoherent twin or incoherent pseudotwin boundaries of the phase in the Ti-10Nb-45Al alloy, which was hot-forged at 1050°C [1997Wan]. The lattice parameters of the 9R phase were obtained from HRTEM as follows: a = 490 pm, b = 282 pm, c = 2080 pm. Landolt-Börnstein New Series IV/11A3
MSIT ®
338
Al–Nb–Ti
Formation of a martensite type fcc phase with a = 437 pm was observed at electrical polishing of thin foils [1978Zak]. Pseudobinary Systems Continuous solid solubility between TiAl3 and NbAl3 was confirmed by experimental investigations [1996Che, 1995Zdz, 1998Din, 1998Wan] and was accepted in the thermodynamic assessment of the Al-Nb-Ti system by [1998Ser]. No new experimental data were reported on the melting temperatures within the (Ti,Nb)Al3 solid solution. The TiAl3-NbAl3 section is shown in Fig. 1 taking into account calculated liquidus temperatures reported by [1998Ser] for the Ti rich solid solutions. They are somewhat higher than those shown earlier by [1990Per] and [1992Kat1], and are more reasonable for the phase with stoichiometric composition. The highest values of solidus temperatures shown by [1990Per] are used to draw the solidus. Invariant Equilibria There are five invariant equilibria with a liquid phase, but the type of only one of them is well established: L+ +) (U type). Various types have been proposed for other invariant equilibria (Table 3). In addition the temperatures of the invariant equilibria are not well established and compositions of phases participating in equilibria are not known. Different stoichiometry for the third aluminide, Ti5Al11, Ti2Al5 or Ti9Al23, has been accepted in equilibria including the phases on the base of Al-rich titanium aluminides TiAl ( ) and TiAl3 (J). The existence of a three phase invariant equilibrium, L +J, was shown by [1990Per] and [1995Zdz] but the temperature of this equilibrium was not established and the position of the maximum point on the liquidus curve is different in [1990Per] and [1995Zdz]. The existence of an invariant equilibrium ++J at ~1100°C was reported by [1989Kal]. [2002Leo1] suggested that the four-phase equilibrium + +)+O occurs at 900°C from a convergence of the + +) and +O+) phase fields in the alloy Ti-37.5Nb-25Al. The existence of the invariant peritectoid reaction 2+0O at about 1000°C was proposed by [1995Mur1], whilst an eutectoid reaction 0O+) was considered by [2001Mis]. Liquidus Surface The liquidus surface was presented earlier by [1989Kal, 1990Per] from experimental studies. A thermodynamic calculation was performed by [1992Kat1]. The liquidus surface, shown in Fig. 2, was constructed by [1995Zdz] and is similar to that of [1990Per]. The liquidus surface presented by [1992Pav1] has not been constructed for the part of the phase diagram adjoining to the Al-Ti side. A peritectic reaction L++) was proposed. Recently [2000Leo1] reinvestigated the liquidus surface and has found that the field of primary crystallization of the phase is wider than earlier reported. Figure 3 shows the liquidus surface projection calculated by [1998Ser]. There are four maximum points, which indicate the existence of three-phase pseudobinary reactions. Isothermal Sections Figure 4 shows the calculated section at 1650°C [1998Ser]. Experimental data [1992Men, 1996Men] for Nb rich alloys show good agreement with calculated boundaries for the + region. The isothermal section at 1400°C was presented by [1996Che, 1998Wan] (Fig. 5) from results of an experimental study and was calculated by [1992Kat1] (shown also in [1993Gam]) and [1998Ser] (Fig. 6). The calculated versions are in good agreement with each other, the existence of the ordered 0 phase is shown in the latter version. The field of the questionable 1 phase is shown in Fig. 5. Boundaries / and / calculated by [2001Li1] are in better agreement with the data of [1996Che, 1998Wan] than with the boundaries calculated by [1998Ser].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
339
It should be noted that the (Ti,Nb)Al3 solid solutions (J) have to remain in solid state in the whole homogeneity range at 1400°C taking into account recent data on the melting temperature of TiAl3 (1408°C [2003Kar], 1425°C [1996Tre, 1997Bul]). The calculated isothermal section at 1300°C [2001Sad] is shown in Fig. 7. The structure of two alloys, Ti-21.8Nb-29.7Al and Ti-31.7Nb-23.4Al, studied by [2001Sad], is consistent with the calculated section. The experimental region in the vicinity of the phase field [2000Kai] (with insignificant corrections to adjust to the accepted binary Al-Ti system) is presented in Fig. 8. There is agreement between calculated and experimental Nb solubility in the phase. The isothermal section at 1260°C [2001Sad] is similar to that at 1300°C. The isothermal section at 1200°C shown in Fig. 9 was taken from [1995Zdz] with minor changes to comply with the accepted binary Al-Ti system. This shows a good agreement with the calculated section (Fig. 10) [1998Ser]. [1992Pav1] presented the isothermal section at 1200°C for the Nb rich side, where the Ti solubility in Nb3Al and Nb2Al were found to be less than shown formerly and confirmed later [1989Jew, 1989Kal, 1990Per, 1990Hel, 1992Sur, 1998Hel, 1993Ebr, 1998Che, 2002Leo1]. Satisfactory agreement also exists between phase compositions of alloys investigated at 1200°C by [1992Jac, 1993Ebr, 1993Nak1, 1994Che1, 1999Eck] and the phase equilibria proposed by [1995Zdz]. [2000Kai] suggested a slightly different configuration of the phase field and adjacent phase fields, but the same Nb solubility in the phase, ~10 at.%. The isothermal section at 1150°C was constructed from results of a diffusion couple study [1996Che, 1998Din] and calculated by [1998Ser] (Fig. 11). Unlike the predicted phase equilibria shown in Fig. 11, those obtained by [1996Che] and modified by [1998Din] are characterized with an existence of equilibrium between and ) phases and + +) and ++) phase fields as well as a separate region of the questionable 1 phase coexisting with the ), and J phases. The data by [1998Yu] on the ++ phase field are consistent with the prediction of [1998Ser]. The isothermal section at 1100°C has been represented by [1991Smi, 1992Smi, 1996Che, 1998Din] and calculated by [1992Kat1, 1998Ser]. Opposite results were obtained for the phase equilibria in the Ti rich part of the system by [1992Kat1] and [1998Ser]. The coexistence of the 2 and ) phases was shown by [1992Kat1], while according to [1998Ser] (Fig. 12) the 2 and ) phase fields are separated by the 0 phase field and the 0 phase coexists with the phase. [1991Smi, 1992Smi], who studied the Nb rich part of the phase diagram, reported the existence of a + +) field though none of the studied alloys was in this region and directions of tie-lines show better correlation with the version by [1992Kat1] rather than [1998Ser]. The version proposed by [1992Che, 1998Din] satisfactorily agrees with [1992Kat1] in the part adjacent to the Nb-Ti side. The above mentioned 1 phase also was shown at higher Al contents [1992Che, 1998Din]. However, numerous results obtained for certain alloys are in agreement with the version by [1998Ser, 1989Ben, 1989Mur, 1990Ben1, 1990Ben2, 1991Ben, 1992Qua, 1994Hou, 1994Ben2, 1999Boe, 1999Eck, 2001Mis, 2002Leo1]. The calculated isothermal section at 1020°C [1998Ser] shown in Fig. 13 is consistent with the experimental section at 1000°C [1998Hel] (Fig. 14). These versions were found to be more reliable than those reported by [1990Hao1, 1994Kum, 1996Che, 1998Din]. [1990Hao1, 1996Che, 1998Din] have shown the not well established 1 phase. [1994Kum] reported only a small part of the section including the O phase, what can be explained by the temperature of formation of O slightly below 1000°C. The ternary phase in the region of the existence of the O phase also was shown by [1990Hao1]. It should be noted that the coexistence of the 2+) phases shown in Figs. 13 and 14 has been observed by [2001Sad], but this has not been reported by a majority of researchers who studied phase transformations and structures of alloys in the appropriate region. The occurrence of the disordered O phase is shown in the calculated isothermal section at 990°C [2001Ser] (Fig. 15). Phase equilibria in the Nb rich part at 900°C have been studied by [1991Smi, 1992Pav1, 1992Pav2, 1992Smi, 1992Zak] and have been presented as a partial isothermal section. The Ti solubility in Nb3Al and Nb2Al found seems to be too low. The calculated isothermal section at 900°C [1998Ser] is shown in Fig. 16. [2003Sad] reported the content of Al to be ~22 at.% in the O phase coexisting with the 0 phase. According
Landolt-Börnstein New Series IV/11A3
MSIT ®
340
Al–Nb–Ti
to [1995Mur1], the largest extension of the homogeneity range of the O phase is along the isopleth at ~27.5 at.% Al. An occurrence of the Ti4NbAl3 ternary compound, designated here as -, is shown in the calculated section at 800°C [1998Ser] (Fig. 17). The experimental partial isothermal sections were presented in [1991Smi, 1992Smi]. The calculated isothermal section at 700°C is accepted from [2001Ser] (Fig. 18). The calculated Nb rich corner at 600°C is shown in Fig. 19 after [1998Ser]. Temperature – Composition Sections The experimental isopleth for 27.5 at.% Al is shown in Fig. 20 on the base of that earlier proposed by [1995Mur1]. The calculated version of the same isopleth [1998Ser] is presented in Fig. 21. The latter version proceeds from the existence of the 2+) equilibrium, while the first one does not suggest this equilibrium. The presented versions of the isopleth also differ by the position of the homogeneity range of the O phase, which according to [1998Ser] is supposed to be located at an Al content less than 27.5 at.%. The existence of the O phase in two forms is shown in both versions. Thermodynamics The thermodynamic activities of Ti and Al, as well as partial enthalpies and entropies of mixing were evaluated from measurements of Ti and Al partial pressures using Knudsen effusion mass spectrometry [1999Eck]. Among the twenty four Al-Nb-Ti alloys studied, more attention was paid to those within the , 2+ and 2 phase fields. The measurements were carried out in the temperature range between 897 and 1362°C and thermodynamic properties were evaluated for the mean temperature of 1200°C. Figures 22 to 24 summarize the thermodynamic activities, partial enthalpies of mixing and partial entropies of mixing for the TiAl based alloys. Additional data for the alloy series (Ti0.48-xNbxAl0.52, Ti0.35NbxAl0.65-x, (Ti0.8Al0.2)1-xNbx, (Ti0.7Al0.3)1-xNbx and Ti0.67NbxAl0.33-x) are given in [1999Eck]. Thermodynamic activities of Al and Ti were calculated using a two-sublattice quasi-subregular solution model for based alloys (Ti0.32Al0.08)1-xNbx (0 < x < 0.2), Ti0.48-xNbxAl0.52 and Ti0.44-xNbxAl0.56 (0 < x < 0.15) [2001Wan]. The Gibbs free energies of the , and phases were described by a subregular solution model; interaction parameters were calculated and used to calculate / and / phase equilibria at 1150 and 1400°C by a grand potential approach [2001Li1]. The Gibbs energy of formation of phases in the Al-Nb-Ti system were derived by [1998Ser] from an optimization procedure using all the available experimental data on thermodynamics and phase equilibria. A multi-sublattice model was used to describe the ordered compounds, whilst solution phases were described by means of Redlich - Kister polynomials. The thermodynamic modeling of the orthorhombic phase was reanalyzed by [2001Ser]. A representation of the thermodynamic properties of ordered and disordered states with a continuous function was applied. Two different models of the orthorhombic phase were performed. The thermodynamic parameters used to model the order/disorder transformation in the orthorhombic O phase were reported. Notes on Materials Properties and Applications The increased interest in titanium aluminides is due to their promising properties, which make them attractive for potential application as aerospace materials, in particular, for jet engine components. These intermetallics are characterized with low density, good strength at elevated temperatures, high resistance to oxidation, good creep properties. However, they exhibit poor ductility at room temperature and low fracture toughness, which can both be significantly improved by additions of niobium. An increased high temperature strength was reported for Nb additions to Ti3Al [1970And, 1972And], but the variation of high temperature strength versus composition was found to exhibit a maximum at 3 mass% Nb and a minimum at 15 mass% Nb. A Ti3Al based alloy with ~5 at.% Nb at 760°C after various heat treatments exhibited a fine acircular Widmanstaetten structure yielding a very high mechanical strength [1977Sas]. However, this structure is unstable at high temperatures and the strength decreases with time. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
341
Nevertheless, the alloy revealed higher strength and ductility than Ti3Al. A similar structure was proposed for Ti based alloys containing 13.5 - 15.3 % Al and 23.4 - 30 % Nb [1981Bla]. These alloys were suggested for application up to 750°C. Tensile tests at room and elevated temperatures and a study of the creep behavior at 650°C were carried out by [1991Row1, 1991Row2] on Ti2NbAl based alloys: the best heat resistance was found after heat treatment in the field. Significant strengthening and resistance to fracture have been achieved in alloys with a Widmanstaetten O+0 structure. Similar results were obtained by [1990Kno] for Ti-11Nb-24Al alloys. [1991Row3] proposed the Ti3Al based alloys Ti-(18-34)Nb-(18-30)Al, which were reported to exhibit an elevated heat resistance and a hot stamping ability, for gas turbine components. [1997Nak] reported a large tensile elongation (~16 - 28 %) at room temperature for Ti3Al based alloys. Elongations up to 810 % have been achieved for a Ti-10Nb-25Al alloy after a deformation rate of 5#10-5 s-1 at 980°C [1992Yan]. The dislocation structure and deformation behavior of the O and 2 phases at RT and at 650°C were examined as a function of the Nb concentration in the alloys Ti-21Nb-26Al and Ti-16Nb-25Al. The O phase was found to deform on all slip systems observed in 2 in spite of the lower (orthorhombic) symmetry [1991Ban, 1995Ban]. The alloys of Ti-11Nb-(24-26)Al are the most studied. Creep testing at 650°C was carried out to evaluate the influence of cooling rate from the field on the steady state strain rate and time to rupture [1990Mis]. Deformation and fracture processes were examined by [1991Akk]. [1992Aco] studied microstructure and microhardness of spot welds. Phase transformations resulting from laser and gas-tungsten-arc welding and solid state processing have been characterized to optimize mechanical properties [1990Cie]. Plasticity of the Ti-25Nb-25Al alloy was improved after rapid quenching and disappeared after annealing [1991Cha]. Slow cooling from the region followed by aging in the 2+ phase field resulted in the formation of relatively stable Widmanstaetten structure and a good balance of compressive and tensile properties of the forged Ti-11Nb-24Al alloy produced by powder metallurgy [1993Sob]. Dynamic material modeling (DMM) was used to analyze the mechanical behavior of the Ti-11Nb-25Al alloy [1993Lon]. Unstable and stable flow zones were predicted by DMM and attributed to the O2 transformation. Data of hot compression tests have been used to construct instability maps for Ti-11Nb-25Al [1994Sag] and Ti-15Nb-25Al alloys [1998Sag]. [2000Mur] determined regimes of unstable material flow during hot deformation of the Ti-15Nb-25Al alloy. [1995Sem] reported on microstructure evolution during rolling of sheets of Ti-23Nb-22Al. A significant increase of hardness (from ~270 VHN to ~440-470 VHN) was observed in the solution treated Ti-22.8Nb-11.1Al alloy as a result of precipitation hardening [1992Qua]. The age hardening occurred in the disordered matrix in the temperature range of 575 - 675°C due to the formation of lath-like 2 precipitates. A similar increase of hardness as observed for the quenched Ti-60Nb-8Al alloy annealed at 600°C has been attributed to the precipitating O phase [1992Voz]. Alloys on the base of (TiAl) have been discussed by [1989Kim] (phase relations, microstructure, processing, mechanical properties, deformation and fracture, factors affecting ductility). A possibility to improve the oxidation resistance of based alloys has been reported earlier by [1962Pop]. [1993Zha] reported two heat resistant alloys Ti-10Nb-45Al and Ti-8Nb-48Al, which were developed for high temperature application. The specific strength of these alloys at 800 - 1100°C was found to be higher than that of TiAl and superalloys (the compressive yield strength was about ~700 MPa at 800°C, 350 MPa at 1100°C, the density was ~4.3 g#cm-3). The alloys showed some ductility at room temperature and oxidation resistance better than that of TiAl and Ti3Al. The 2 phase transformation, which occurred at grain boundaries during high temperature stress rupture deformation, has been studied by [2000Che]. Internal friction at high temperature and creep measurements were carried out for a Ti-4Nb-46.5Al alloy [2000Wel]. Planar fault energies and sessile dislocation configurations were studied in (Ti1-xAlx)1-yNby alloys, 0.48 < x < 0.51, 0 < y < 0.02 [1996Woo]. High temperature strength (compression testing up to 1100°C) and oxidation behavior (at 900 - 1200°C in air) of alloys in a wide composition range (Ti3Al - TiAl 3 - NbAl 3 - TiNbAl3) have been investigated by [1992Che]. The alloys with 55 - 64 at.% Al and a Ti:Nb ratio of 2 to 5 yielded the highest oxidation resistance besides high tensile strength. Landolt-Börnstein New Series IV/11A3
MSIT ®
342
Al–Nb–Ti
Fracture toughness measurements and fractographic analysis were carried out to evaluate the toughening mechanism of the ) phase with particles of and phases [1993Ebr]. Superconducting properties of solid solutions on the base of Nb3Al were studied by [1975Pan, 1975Sha, 1977Ale]. The critical temperature Tc of the superconducting transition was found to decrease down to ~9 K with increasing Ti content up to ~13 at.% [1975Sha]. [1981Ish] investigated the influence of Al additions on the critical current density in superconducting Nb-60Ti alloys. Electrical resistivity and its temperature dependence in the range of 20 to 220°C, as well as emf in a couple with Cu have been studied for Ti alloys containing up to 50 mass% Nb and 10 mass% Al. Aluminium additions to Nb-Ti alloys resulted in a decrease of heat conductivity [1965Kal]. Electrical resistivity, hardness and density of Ti3Al-Nb alloys (up to 50 mass% Nb) have been studied by [1970And]. Temperature dependence of the 0.2 % proof stress at a compression rate 10-4 s-1 for Ti0.25Nb0.75Al3 with the D022 structure was presented by [1990Sau]. Miscellaneous The effect of Nb on the phase equilibria and transformation behavior in Al-Ti alloys based on /0, 2 and phases has been discussed for development of advanced high temperature materials [1999Flo]. It was pointed out that data reported on the phase transformation in the appropriate field of the phase diagram are fragmentary and often they are mutually incompatible. This may be due to limitations in experimental techniques or interstitial contamination. It can be added that elements of phase diagrams often contradict requirements of the phase equilibria theory. CCT - Curves-(Continuous Cooling Transformation) Schematic curves of continuous cooling transformations were derived from a study of microstructure occurring in the Ti-11Nb-24Al alloy during continuous cooling from 1230°C ( field) down to room temperature by immersing a wedge-shaped specimen with a narrow end into ice water [1990Wey]. [1995Lon] used DTA (600 - 1300°C) and in situ high temperature XRD (600 - 1300°C) to investigate phase stability during continuous heating/cooling of Ti-11Nb-25Al alloy. The sequence of the phase fields 2++O2++ was established at heating, the same fields were identified at cooling. The alloy was in the 2++O region up to 850°C, the field was found to exist above 1200°C. [2001Sad] constructed CCT diagrams for Ti-21.8Nb-27.9Al and Ti-31.7Nb-23.4Al alloys from samples, which had been cooled from 1260°C with the rates from 100 to 0.25 K#s-1, using dilatometry, DTA, XRD, SEM, TEM and microhardness measurements. Out-of-equilibrium phase transformations were observed for fast cooling, while quasi-equilibrium transformations were detected for lower cooling rates. The sequence of transformations at a cooling rate of 0.25 K#s-1 was established to be 0)+02+0+)O+0+)-+)+O for Ti-21.8Nb-27.9Al and 0)+02+0+)O+0+) for Ti-31.7Nb-23.4Al. The CCT diagrams for the alloys Ti-21.8Nb-27.9Al and Ti-31.7Nb-23.4Al, respectively, are shown in Figs. 25, 26. Three non-equilibrium phases, 7´, 7´´ and O m (a massive orthorhombic phase, which formed by a diffusion-less mechanism and had the chemical composition of the parent B2 phase) were observed. According to an in-situ neutron diffraction study [2000Sad] the transition from 0+2+ to 2+ +- in Ti-12.9Nb-36.5Al occurs between 800 and 960°C. Atomic Structure and Electronic Structure The electronic structure and the total energy of Ti2NbAl in B2 (0), D019 (2) and O structure were calculated with the self-consistent tight binding linear muffin-tin orbital method [1999Rav]. The obtained results were used to study the phase stability and cohesive properties of these phases. The B2 phase was shown to be the most stable one. The presence of all these phases in equilibrium over a range of temperature is possible because they are close in energy. The heats of formation H were calculated to be -0.239, -0.208 and -0.036 (eV/atom) for the B2, D019 and O phases, respectively. The linear muffin-tin orbital method was also employed to elucidate the atom site distribution in ordered (TiAl) compounds (L10), TiXAl2 and Ti2AlX (X = transition metal) via calculation of the electronic structure and total energies from first principles [1993Ers]. Niobium was found to preferentially substitute on Ti sites thereby increasing c/a. Accordingly, preferential Nb substitution for Ti in TiAl was established MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
343
experimentally [1986Kon, 1991Moh, 1999Hao] and has been predicted using a thermodynamic approach based on a Bragg - Williams model [1990Nan], a plane-wave pseudopotential method [1996Woo, 1998Woo] and CVM [2001Kan]. Ti/Nb substitution is also supported from the partial entropy of mixing for the (Ti0.38Al0.62)1-xNbx (0 < x < 0.2 alloy series [1999Eck]. Ti substitution for Nb on Nb sublattice sites was determined for the ordered 0 phase in five Nb rich alloys [2002Leo2]. The local atomic order in 0 phase of Ti2NbAl composition has furthermore been studied using the Extended X-Ray Absorption Fine Structure (EXAFS) [1996Sik]. The site composition was shown to be written as (Ti1.5Nb0.5)A(Ti0.5Nb0.5Al) B. CVM in the irregular tetrahedron approximation, used to calculate the 800, 1000, 1200 and 1400°C isothermal partial sections, revealed a Heusler type phase instead of the CsCl (B2) phase for the Ti rich region with a miscibility gap between the ordered Heusler phase and the disordered W(A2) type phase [1996Jac, 1999Cha1]. CVM was furthermore employed by [1999Cha1] to calculate the sublattice occupation of the 0 phase by Nb, Ti and Al atoms. It was shown that Ti atoms occupy one sublattice, Al atoms on the other but Nb atoms prefer one or both sublattices depending on the composition. In agreement with the CVM calculation, a neutron diffraction study [1999Cha2] at room temperature and in situ up to 1600°C has shown that sublattice occupation of the 0 phase is sensitive to the concentration. [1987Ban] determined the site occupancy in the ordered solid solution phase 0 in the Ti-10Nb-25Al alloy using ALCHEMI. It was shown that Ti atoms occupy one of two possible sublattices whilst Al and Nb atoms are found in the other one. Theoretical and experimental investigations of sublattice substitution of Nb in (TiAl) and 2 (Ti3Al) based alloys [1999Hao, 1999Yan] comparing binding energy data and the Bragg-Williams model with ALCHEMI (Atom Location Channeling Enhanced Microanalysis) measurements were summarized by [2000Yan]. ALCHEMI data prompted a strong preference of Nb atoms to substitute for Ti in both TiAl and Ti3Al [1999Hao]. [1986Kon] confirmed the Nb/Ti substitution in the Ti3Al lattice. The ordering tie-line (OTL) approach to represent sublattice occupations was adopted by [2000Ama]: the OTLs were determined via the ALCHEMI method. It was suggested that the order-disorder transformation is a second-order phase transformation. Studies of corrosion An addition up to 15 mass% Al to Nb alloys containing 20-40 mass% Ti significantly decreases the oxidation rate at 1100°C [1991Pav]. Oxidation kinetics of a Ti-25Nb-50Al alloy was studied using thermogravimetry in air, pure O2 and their mixture at 1300°C at the pressure of 100 kPa [1992Bra]. A study of cyclic oxidation of a Ti-24Nb-14Al alloy by [1988Sub] demonstrated the benefits of a protecting TiAl3 coating. Stress corrosion cracking (SCC) was shown to occur for a Ti-11Nb-24Al (2+) alloy in methanol and aqueous solutions and needs to be taken into account in developing and applying Ti3Al-Nb alloys [1992Zha]. Electro-spark deposition (ESD) was used to produce crack-free TiAl3 aluminide coating on a Ti3Al-Nb alloy (Ti-10.8Nb-24.1Al) to improve its high temperature oxidation resistance [2001Li2]. An Al plate was used as an electrode material. Isothermal oxidation tests at 800 and 900°C in air proved the low oxidation rate of the coating. The use of Ti hydride instead of pure Ti for the synthesis of O phase based alloys by ball-milling resulted in a reduced contamination with oxygen and nitrogen, in considerable particle refinement and it accelerated the amorphization of the powders [2002Bou]. [1989Shi] investigated the hydrogenation behavior in Ti3Al observing a “hydride” phase in the Ti-11Nb-24Al alloy. The crystal structure of this phase was not established but an orthorhombic distortion of the hexagonal base structure was reported. [1992Roz] studied the influence of hydrogen on phase transformations in Ti-11Nb-24Al. Cathode charging hydrogen resulted in the formation of a Ti3Al-H hydride in a thin surface layer and induced cracking. Temperature and pressure dependencies of hydrogen solubility in a Ti-11Nb-24Al alloy were reported by [1992Chu] and the hydrogenization behavior of three alloys with compositions in the vicinity of Ti2NbAl was investigated by [2001Zha]; a beneficial effect of the O phase on the hydrogenization properties was established, i.e. Hf becomes more negative with increasing volume fraction of the O phase. Landolt-Börnstein New Series IV/11A3
MSIT ®
344
Al–Nb–Ti
[2002Hod] studied the behavior of Ti-7Nb-6Al (mass%) alloy under simulated biological conditions (specific ions, pH, temperature) i.e. the electrochemical characterization by impedance spectroscopy and photoelectrochemistry of the passive film. An investigation of the sulfidation process of TiAl-2Nb (at.%) alloy was undertaken in order to find out the alloying element, which would improve oxidation resistance [2000Izu]. The sulfidation amount was found to be close to that for binary TiAl. Disordering of the phase with tetragonal lattice and a new phase formation with a smaller c/a ratio were observed at a neutron irradiation treatment of a TiAl-Nb alloy [1986Ibr]. The diffusivity in the phase was estimated at 1200 and 1400°C using the diffusion couple method [1996Ebr]: Ti seems to be the fastest species, Al having a mobility close to Ti and Nb being the slowest species. A bulk Ti-19.9Nb-14.6Al nanophase material with the structure of the O phase was synthesized and consolidated from powders with structure produced by ball milling [1991Chr]. The grain size of the consolidated material was ~10 nm, the density was 4.48 g#cm-3 and Vicker´s hardness was 498 VHN. References [1962Pop] [1965Kal]
[1970And]
[1970Nar]
[1972And]
[1972Nar]
[1974Nar]
[1975Fed]
[1975Pan] [1975Sha]
[1977Ale]
[1977Sas]
MSIT®
Popov, I.A., Rabezova, V.I., “Investigation of the Phase Diagram of the Al-Nb-Ti System” (in Russian), Zh. Neorg. Khim., 7, 436-439 (1962) (Equi. Diagram, Experimental, 9) Kalinin, G.R., Elyutin, O.P., Mamontovskaya, L.Y., “Physical and Mechanical Properties of Alloys of the Ti-Nb-Al System” (in Russian), Izv. Akad. Nauk SSSR, Met., (3), 146-150 (1965) (Experimental, 8) Andreyev, O.N., “Phase Structure and High-Temperature Strength of Ti3Al-Nb Alloys” (in Russian), Izv. Akad. Nauk SSSR, Met., (1), 193-196 (1970) (Equi. Diagram, Experimental, 10) Nartova, T.T., Sopochkin, G.G., “Investigation of the Phase Structure of Ti3Al-Nb Alloys” (in Russian), Izv. Akad. Nauk SSSR, Met., (2), 220-223 (1970) (Equi. Diagram, Experimental, 9) Andreyev, O.N., Kornilov, I.I., “Study of Effect of Some Elements on High-Temperature Strength of Ti3Al” (in Russian), in “Nov. Konstr. Mater. Titan”, Nauka, Moscow, 101-164 (1972) (Experimental, 5) Nartova, T.T., Sopochkin, G.G., “Phase Equilibrium Study of Alloys of the Ti-Al-Nb System” (in Russian), in “Nov. Konstr. Mater. Titan”, Nauka, Moscow, 19-23 (1972) (Equi. Diagram, Experimental, 4) Nartova, T.T., Sopochkin G.G., “Reaction of Titanium Aluminide Ti3Al with Niobium and Molybdenum” (in Russian), in “Stroyeniye, Svoistva i Primeneniye Metallidov”, Nauka, Moscow, 80-83 (1974) (Equi. Diagram, Experimental, 9) Fedorova, M.A., Turchinskaya, M.I., Sokolovskaya, E.M., “Effect of Group IVB Elements on the Structure and Superconductivity of the Nb 3Al Intermetallic” (in Russian), Vestn. Mosk. Univ., Ser. 2: Khim., 16, 238-239 (1975) (Experimental, 4) Pan, V.M., Latysheva, V.I., “Superconductivity of Nb-Al-Ti Alloys” (in Russian), Metallofizika, (57), 74-77 (1975) (Equi. Diagram, Experimental, 6) Shamrai, V.F., Postnikov, A.M., “Investigation of Some Ternary Solid Solutions Based on the Nb3Al Compound” (in Russian), Dokl. Akad. Nauk SSSR, 224, 1130-1133 (1975) (Crys. Structure, Equi. Diagram, Experimental, 8) Alekseyevskiy, N.Y., Ageev, N.V., Shamrai, V.F., “Superconductivity of Some Ternary Solid Solutions Based on the Nb 3Al Compound” (in Russian), Fiz. Met. Metalloved., 43, 38-44 (1977) (Crys. Structure, Equi. Diagram, Experimental, 14) Sastry, S.M.L., Lipsitt, H.A., “Ordering Transformations and Mechanical Properties of Ti 3Al and Ti3Al-Nb Alloys”, Metall. Trans., 8A, 1543-1552 (1977) (Crys. Structure, Equi. Diagram, Experimental, 22)
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti [1978Zak]
[1980Jor]
[1981Bla] [1981Ish] [1981Kin] [1982Mii]
[1982Str]
[1983Tro]
[1984Arg] [1984Zak]
[1986Ibr]
[1986Kon]
[1987Ban]
[1987Mur] [1988Ban] [1988Has]
[1988Str]
[1988Sub] [1989Ben]
Landolt-Börnstein New Series IV/11A3
345
Zakharova, M.I., Khatanova, N.A., Kozlovskaya, N.A., “An Investigation of Supersaturated Solid Solution Decomposition in a Ti-Nb-Al Alloy” (in Russian), Vestn. MGU Fiz., Astron., 19, 121-124 (1978) (Crys. Structure, Experimental, 11) Jorda, J.L., Fluekiger, R.J., Mueller, “A New Metallurgical Investigation of the Niobium-Aluminium System”, J. Less-Common Met., 75, 227-239 (1980) (Crys. Structure, Equi. Diagram, Experimental, 20) Blackburn, M.J., “Titanium Alloys of the Ti3Al Type”, Pat. 60264 USA, Cit. by Ref. J. Metallurgiya, (7), Abs. 10I449P (1982) (in Russian) Ishida, F., “Influence of Third Element Additions on the Critical Current Density of Nb-60 at.% Ti Alloys” (in Japanese), J. Jpn. Inst. Met., 45, 517-524 (1981) (Experimental, 8) King, H.W., “Crystal Structure of the Elements at 25°C”, Bull. Alloy Phase Diagrams, 2, 401-42, (1981) (Crys. Structure, Review, 5) Miida, R., Hashimoto, S., Watanabe, D., “New Type of A5B3 Structure in Al-Ti and Ga-Ti Systems; Al 5Ti3 and Ga5Ti3”, Japan. J. Appl. Phys., 21, L59-L61 (1982) (Crys. Structure, Experimental, 10) Strychor, R., Williams, J.C., “Phase Transformations in Ti-Al-Nb Alloys”, Proc. Int. Conf. Solid-Solid Phase Transformations, Pittsburgh, 1981, Warrandale, 249-253 (1982) (Crys. Structure, Equi. Diagram, Experimental, 10) Troitskii, B.S., Zakharov, A.M., Karsanov, G.V., Vergasova, L.L., “Polythermal Sections of the Nb-Ti-Al System” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., (3), 77-80 (1983); translated from Sov. Non-Ferrous Met. Res., 11, 231-232 (1983) (Crys. Structure, Equi. Diagram, Experimental, 8) Argent, B.B., “Phase Diagrams of Alloys Based on Niobium”, Met. Soc. (AIME), Acc. No. 84(7), 72-486, 325-415 (1984) (Crys. Structure, Equi. Diagram, Review, 11) Zakharov, A.M., Karsanov, G.V., Troitskii, B.S., Vergasova, L.L., “Isothermal Sections of the System Nb-Ti-Al at 1200-600ºC” (in Russian), Izv. Akad. Nauk SSSR, Met., 1, 200-202 (1984) (Crys. Structure, Equi. Diagram, Experimental, 7) Ibragimov, S.S., Kofanov, B.A., Melikhov, V.D., “Change of Intermetallic Phases Structure in Three-Phase TiAl-Nb Alloy at Neutron Irradiation” (in Russian), Izv. Akad. Nauk Kazakh. SSR, Ser. Fiz. Mat., (2), 29-33 (1986) (Experimental, 8) Konitzer, D.G., Jones, I.P., Fraser H.L., “Site Occupancy In Solid Solutions of Nb in the Intermetallic Compounds TiAl And Ti3Al”, Scr. Metall., 20, 265-268 (1986) (Crys. Structure, Experimental, 9) Banerjee, D., Nandy, T.K., Gogia, A.K., “Site Occupation in the Ordered -Phase of Ternary Ti-Al-Nb Alloys”, Scr. Metall., 21, 597-600 (1987) (Crys. Structure, Experimental, 21) Murray, J.L., “Nb-Ti (Niobium - Titanium)”, in “Phase Diagrams of Binary Titanium Alloys”, ASM Publication, 188-194 (1987) (Crys. Structure, Equi. Diagram, Review, 44) Banerjee, D., Gogia, A.K., Nandi, T.K., Joshi, V.A., “A New Ordered Orthorhombic Phase in a Ti3Al-Nb Alloy”, Acta Metall., 36, 871-882 (1988) (Crys. Structure, Experimental, 22) Hashimoto, K., Doi, H., Kasahara, K., Tsujimoto, T., Suzuki, T., “Effects of the Third Elements on the Structures of TiAl-Based Alloys” (in Japanese), J. Jpn. Inst. Met., 52, 816-825 (1988) (Crys. Structure, Equi. Diagram, Experimental, 31) Strychor, R., Williams, J.C., Soffa, W.A., “Phase Transformations and Modulated Microstructures in Ti-Al-Nb Alloys”, Metall. Trans., 19A, 225-234, (1988) (Equi. Diagram, Experimental, 44) Subrahmanyam, J., “Cyclic Oxidation of Aluminized Ti-14Al-24Nb Alloy”, J. Mater. Sci., 23, 1906-1010 (1988) (Experimental, 7) Bendersky, L.A., Boettinger, W.J., “Investigation of B2 and Related Phases in the Ti-Al-Nb Ternary System”, Mater. Res. Soc. Symp. Proc., 133, 45-50 (1989) (Equi. Diagram, Experimental, 6)
MSIT ®
346 [1989Jew]
[1989Kal]
[1989Kes]
[1989Kim] [1989Mur] [1989Shi] [1990Ben1]
[1990Ben2]
[1990Cie]
[1990Hao1]
[1990Hao2]
[1990Hel]
[1990Kno] [1990Mis] [1990Moz]
[1990Mur]
[1990Nan] [1990Per]
MSIT®
Al–Nb–Ti Jewett, T.J., Lin, J.C., Bonda, N.R., Seitzman, L.E., Hsieh, K.C., Chang, A.Y., Perepezko, J.H., “Experimental Determination of the Titanium-Niobium-Aluminum Phase Diagram at 1200°C”, Mater. Res. Soc. Symp. Proc., 133, 69-74 (1989) (Equi. Diagram, Experimental, 8) Kaltenbach, K., Gama, S., Pinatti, D.G., Schulze, K., Henig, E.-T., “A Contribution to the Ternary System Al-Nb-Ti”, Z. Metallkd., 80, 535-539 (1989) (Equi. Diagram, Experimental, 13) Kestner-Weykamp, H.T., Ward, C.H., Broderick, T.F., Kaufman, M.J., “Microstructures and Phase Relationships in the Ti 3Al+Nb System”, Scr. Metall., 23, 1697-1702 (1989) (Crys. Structure, Equi. Diagram, Experimental, 13) Kim, Y.-W., “Intermetallic Alloys Based on Titanium Aluminide”, JOM, 41, 24-30 (1989) (Review, 61) Muraleedharan, K., Banerjee, D., “Alloy Partitioning in Ti-24Al-11Nb Analytical Electron Microscopy”, Metall. Trans., 20A, 1139-1142 (1989) (Equi. Diagram, Experimental, 10) Shih, D.S., Scarr, G.K., Wasielewski, G.E., “On Hydrogen Behavior in Ti3Al”, Scr. Metall., 23, 973-978 (1989) (Experimental, 13) Bendersky, L.A., Boettinger, W.J., Burton, B.P., Biancaniello, F.S., “The Formation of Ordered 7-Related Phases in Alloys of Composition Ti4Al3Nb”, Acta Metall. Mater., 38, 931-943 (1990) (Crys. Structure, Equi. Diagram, Experimental, 24) Bendersky, L.A., Burton, B.P., Boettinger, W.J., Biancaniello, F.S., “Ordered 7-Derivatives in a Ti-37.5Al-12.5Nb (at.%) Alloy”, Scr. Metall. Mater., 24, 1541-1546 (1990) (Crys. Structure, Equi. Diagram, Experimental, 6) Cieslak, M.J., Headly, T.J., Baeslack III, “Effect of Thermal Processing of the Microstructure if Ti-26Al-11Nb: Application to Fusion Welding”, Metall. Trans., 21A, 1273-1286 (1990) (Experimental, 27) Hao, S., Zhao, Q., “Investigation of the 1000°C Isothermal Section of Ti-Al-Nb Ternary Phase Diagram” (in Chinese), Proc.: 6 th National Symp. Phase Diagrams, Shenyang, China, 1990, 141-143 (1990) (Equi. Diagram, Experimental, 4) Hao, S., Zhao, Q., “A New Ternary Intermetallic Compound in Ti-Al-Nb System”, Proc.: 6th National Symp. Phase Diagrams, Shenyang, China, 1990, 144-145, 149 (1990) (Crys. Structure, Experimental, 3) Helwig, A., “Experimental Study About the Constitution of the Aluminium - Titanium Niobium System” (in German), Ph.D. Thesis, University of Dortmund (1990) (Experimental, 68) as quoted by [1993Gam] Knorr, D.B., Stoloff, N.S., “Effect of Heat Treatment on Microstructure and Texture in Ti-24 at.% Al-11at.% Nb”, Mater. Sci. Eng., A123, 81-87 (1990) (Experimental, 23) Misra, R.S., Banerjee, D., “On the Influence of Cooling Rate in Solution Treatment for a Ti-25Al-11Nb Alloy”, Scr. Metall. Mater., 24, 1477-1482 (1990) (Experimental, 18) Mozer, B., Bendersky, L.A., Boettinger, W.J., “Neutron Powder Diffraction Study of the Orthorhombic Ti2AlNb Phase”, Scr. Metall. Mater., 24, 2363-2368 (1990) (Crys. Structure, Experimental, 10) Muraleedharan, K., Naidu, C.V.N., Banerjee, D., “Orthorhombic Distortion of the 2 Phase in Ti3Al-Nb Alloys: Artifacts and Facts”, Scr. Metall. Mater., 24, 27-32 (1990) (Crys. Structure, Experimental, 7) Nandy, T.K., Banerjee, D., Gogia, A.K., “Site Substitution of TiAl Intermetallic”, Scr. Metall. Mater., 24, 2019-2022 (1990) (Crys. Structure, Theory, Thermodyn., 13) Perepezko, J.H., Chang, Y.A., Seitzman, L.E., Lin J.C., Bonda, N.R., Jewett, T.J., Mishurda. J.C., “High Temperature Phase Stability in the Ti-Al-Nb System”, in “High Temperature Aluminides and Intermetallics”, Wang, S.H., Liu, C.T., Pope, D.P., Stiegler, J.O., (Eds.), The Minerals, Metals and Materials Society, 19-47 (1990) (Equi. Diagram, Experimental, 20)
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti [1990Sau] [1990Sch]
[1990Sho]
[1990Wey]
[1991Akk]
[1991Ban]
[1991Ben]
[1991Cha] [1991Chr] [1991Li] [1991Moh] [1991Pav]
[1991Row1]
[1991Row2]
[1991Row3]
[1991Smi]
[1991Spa]
[1991Zak]
Landolt-Börnstein New Series IV/11A3
347
Sauthoff, G., “Intermetallic Alloys-Overview on New Materials Developments for Applications in West Germany”, Z. Metallkd., 81, 855-861 (1990) (Review, 36) Schuster, J.C., Ipser, H., “Phases and Phase Relations in the Partial System TiAl3-TiAl”, Z. Metallkd., 81, 389-396 (1990) (Crys. Structure, Equi. Diagram, Experimental, Review, 33) Shoemaker, C.B., Shoemaker, D.P., Bendersky, L.A., “Structure of 7-Ti3Al2.25Nb0.75”, Acta Crystallogr., Sect. C: Cryst. Struct. Commun., C46(3), 374-377 (1990) (Crys. Structure, Experimental, 9) Weykamp, H.T., Baker, D.R., Paxton, D.M., Kaufman, M.J., “Continuous Cooling Transformations in Ti 3Al+Nb Alloys”, Scr. Metall. Mater., 24, 445-450 (1990), (Crys. Structure, Experimental, 13) Akkurt, A.S., Liu, G., Bond, G.M., “Micromechanisms of Deformation and Fracture in a Ti 3Al-Nb Alloy”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 455-460 (1991) (Crys. Structure, Experimental, 11) Banerjee, D., Rowe, R.G., Hall, E.L., “Deformation of the Orthorhombic Phase in Ti-Al-Nb Alloys”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc.: Johnson, L.A., Pope, D.P., Stiegler, J.O., (Eds.), 213, 285-290 (1991) (Crys. Structure, Experimental, 16) Bendersky, L.A., Boettinger, W.J., Roytburd, A., “Coherent Precipitates in the B.C.C./Orthorhombic Two-Phase Field of the Ti-Al-Nb System”, Acta Metall. Mater., 39, 1959-1969 (1991) (Crys. Structure, Equi. Diagram, Experimental, 23) Chang, C.P., Loretto, M.H., “The Decomposition Process of Rapidly Solidified Ti-25 at.% Al-25 at.% Nb”, Philos. Mag. A, 63, 389-406 (1991) (Crys. Structure, Experimental, 23) Christman, T., Jain, M., “Processing and Consolidation of Bulk Nanocrystalline Titanium Aluminide”, Scr. Metall. Mater., 25, 767-772 (1991) (Crys. Structure, Experimental, 32) Li, D., Zhou, J., Chang, X., Guan, S., “On the Ordering Transformations in Ti3Al-Nb Alloy”, Acta Metall. Sin. (China), 4A(3), 204-208 (1991) (Equi. Diagram, Experimental, 6) Mohandas, E., Beaven, P.A., “Site Occupation of Nb, V, Mn and Cr in -TiAl”, Scr. Metall. Mater., 25, 2023-2027 (1991) (Crys. Structure, Experimental, 15) Pavlov, A.V., Zakharov, A.M., Karsanov, G.V., Vergasova, L.L., “An Influence of Al and Si upon Heat Resistivity of Nb-Ti Alloys at 1100°C” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., (5), 89-94 (1991) (Experimental, 17) Rowe, R.G., Hall, E.L., “Stress-Assisted Discontinuous Precipitation during Creep of Ti 3Al-Nb Alloys”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc., 213, 449-454 (1991) (Experimental, 11) Rowe, R.G., Konitzer, D.G., Woodfield, A.P., Chesnutt, J.C., “Tensile and Creep Behavior of Ordered Orthorhombic Ti2AlNb-Based Alloys”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc., 213, 703-708 (1991) (Experimental, 10) Rowe, R.G., “Tri-Titanium Aluminide Alloys Containing at Least Eighteen Atom Percent Niobium”, Pat. 5032357 USA, Cit. by Ref. J. Metallurgiya, (10), Abs. 10I449P (1992) (in Russian) Smirnova, T.R., Zakharov, A.M., Oleinikova, S.V., Filipyeva, O.A., “Phase Composition of Alloys in the Nb-Ti-Al System with 0-20 % Al and Ti:Nb1 at 1100-800°C” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., 4, 93-100 (1991) (Crys. Structure, Equi. Diagram, Experimental, 4) Sparks, C.J., Porter, W.D., Schneibel, J.H., Oliver, W.C., Golec, C.G., “Formation of Cubic L1 2 Phases from Al3Ti and Al3Zr by Transition Metal Substitutions for Al”, Mater. Res. Soc. Symp. Proc., 186, 175-180 (1991) (Crys. Structure, Experimental, 15) Zakharov, A.M., Pavlov, A.V., Kachanova, T.L., “The Molybdenum Influence on the Phase Composition of the Nb-Ti-Al Alloys at 1400 - 1600°C” (in Russian), Izv. Akad. Nauk SSSR, Met., (3), 102-106 (1991) (Equi. Diagram, Experimental, 10)
MSIT ®
348 [1992Aco]
[1992Ben]
[1992Bra]
[1992Che] [1992Chu] [1993Gam]
[1992Hsi1]
[1992Hsi2] [1992Hsi3]
[1992Hsi4] [1992Jac] [1992Kat1] [1992Kat2]
[1992Kim]
[1992Men] [1992Mur1]
[1992Mur2]
[1992Pav1]
[1992Pav2]
MSIT®
Al–Nb–Ti Acoff, V.L., Thompson, R.G., Griffin, R.D., Radhakrishnan, B., “Effect of Heat Treatment on Microstructure and Microhardness of Spot Welds in Ti-26Al-11Nb”, Mater. Sci. Eng., A152, 304-309 (1992) (Abstract) (Experimental, 5) Bendersky, L.A., Boettinger, W.J., Biancaniello, F.S., “Intermetallic Ti-Al-Nb Alloys Based on Strengthening of the Orthorhombic Phase by 7-Type Phases”, Mater. Sci. Eng., 152A, 41-47 (1992) (Experimental, 15) Brady, M.P., Nanrahan, R.J. (Jr.), Elder, R.S.P., Verink, E.D. (Jr.), “The Effect of Nitrogen on the Oxidation Behavior of 25Nb-25Ti-50Al”, Scr. Metall. Mater., 26, 767-770 (1992) (Experimental, 6) Chen, G., Sun, Z., Xhou, X., “Oxidation and Mechanical Behavior of Intermetallic Alloys in the Ti-Nb-Al Ternary System”, Mater. Sci. Eng., 153, 597-601 (1992) (Experimental, 6) Chu, W.-Y., Thompson, A.W., Williams, J.C., “Hydrogen Solubility in a Titanium Aluminide Alloy”, Acta Metall. Mater., 40, 455-462 (1992) (Experimental, 38) Gama, S., “Aluminium - Niobium - Titanium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16070.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 22) Hsiung, L.M., Cai., W., Wadley, H.N.G., “Microstructure and Phase Evolution in Rapidly-Solidified Ti-24Al-11Nb”, Mater. Sci. Eng., 152, 295-303 (1992) (Experimental, 14) Hsiung, L.M., Wadley, H.N.G., “A New Ordered Tetragonal Phase in the Ti3Al+Nb System”, Scr. Metall. Mater., 26, 35-40 (1992) (Crys. Structure, Experimental, 10) Hsiung, L.M., Wadley, H.N.G., “Structural Relationships between the T and O Phases in Ti-24Al-11Nb”, Scr. Metall. Mater., 26, 1071-1076 (1992) (Crys. Structure, Experimental, Theory, 7) Hsiung, L.M., Wadley, H.N.G., “Stability of the Ordered Orthorhombic Phase in Ti-24Al-11Nb”, Scr. Metall. Mater., 27, 605-610 (1992) (Crys. Structure, Experimental, 9) Jackson, A.G., Lee, D.S., “Characterization of the Phases Present in a Ti-45 at.% Al-10 at.% Nb Alloy”, Scr. Metall. Mater., 26, 1575-1579 (1992) (Crys. Structure, Experimental, 8) Kattner, U.R., Boettinger, W.J., “Thermodynamic Calculation of the Ternary Ti-Al-Nb System”, Mater. Sci. Eng., A152, 9-17 (1992) (Equi. Diagram, Thermodyn., #, 20) Kattner, U.R., Lin, J.-C., Chang, Y.A., “Thermodynamic Assessment and Calculation of the Ti-Al System”, Metall. Trans., 23A, 2081-2090 (1992) (Equi. Diagram, Review, Theory, Thermodyn., #, 51) Kimura, M., Hashimoto, K., Morikawa, H., “Study on Phase Stability in Ti-Al-X Systems at High Temperatures”, Mater. Sci. Eng., A152, 54-59 (1992) (Equi. Diagram, Experimental, 12) Menon, E.S.K., Subramanian, P.R., Dimiduk, D.M., “Phase Equilibria in Niobium Rich Nb-Al-Ti Alloys”, Scr. Metall. Mat., 27, 265-270 (1992) (Equi. Diagram, Experimental, 22) Muraleedharan, K., Gogia, A.K., Nandy, T.K., Banerjee, D., Lele, S., “Transformation in a Ti-24Al-15Nb Alloy: Part I. Phase Equilibria and Microstructure”, Metall. Trans., 23A, 401-415 (1992) (Equi. Diagram, Experimental, 28) Muraleedharan, K., Gogia, A.K., Nandy, T.K., Banerjee, D., Lele, S., “Transformation in a Ti-24Al-15Nb Alloy: Part II. A Composition Invariant 0O Transformation”, Metall. Trans., 23A, 417-431 (1992) (Crys. Structure, Experimental, 20) Pavlov, A.V., Zakharov, A.M., “Phase Equilibria in the Nb-Ti-Al System” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., (1-2), 98-104 (1992) (Crys. Structure, Equi. Diagram, Experimental, 24) Pavlov, A.V., Zakharov, A.M., Karsanov, G.V., Vergasova, L.L., “Isothermal Sections of the Nb-Ti-Al System at 900 and 600°C” (in Russian), Russ. Akad. Nauk, Metally, (5), 117-119 (1992) (Equi. Diagram, Experimental, 10)
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti [1992Qua] [1992Roz] [1992Smi]
[1992Sur] [1992Tre]
[1992Voz]
[1992Yan] [1992Zak]
[1992Zha]
[1993Ebr]
[1993Ers]
[1993Jac] [1993Lon] [1993Mur] [1993Nak1]
[1993Nak2] [1993Oka1] [1993Oka2] [1993Rub]
[1993Sob]
Landolt-Börnstein New Series IV/11A3
349
Quatrocchi, L.S., Koss, D.A., Scarr, G., “Precipitation Hardening of Beta Titanium Alloy by the Alpha-Two Phase”, Scr. Metall. Mater., 26, 267-272 (1992) (Experimental, 10) Rozenak, P., Dangur, M., “Effects of Hydrogen on the Hydride Transformation in Ti-24Al-11Nb Alloys”, J. Mater. Sci., 27, 2273-2278 (1992) (Experimental, 13) Smirnova, T.P., Zakharov, A.M., Oleinikov, S.V., Filipyeva, O.A., “Phase Composition of Alloys of the Nb-Ti-Al System with 0-20% Al And Ti:Nb Ratio 1 at 1100-800°C” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., (1-2), 91-98 (1992) (Crys. Structure, Equi. Diagram, Experimental, 3) Surynarayana, C., Lee, D.S., “Phase Relations in Ti-Al-Nb Alloys at 1200°C”, Scr. Metall. Mater., 26, 919-924 (1992) (Crys. Structure, Equi. Diagram, Experimental, 16) Trenogina, T.L., Vozilkin, V.A., Volkova, S.B., “On the Stability of Ordered Orthorhombic 0-Phase in Ti60 Nb8Al Alloy” (in Russian), Fiz. Met. Metalloved, (12), 96-98 (1992) (Crys. Structure, Equi. Diagram, Experimental, 6) Vozilkin, V.A., Trenogina, T.L., Volkova, S.B., “Influence of Aluminium on the Structure and Properties of Ti-60% Nb Alloy” (in Russian), Fiz. Met. Metalloved., (11), 108-113 (1992) (Crys. Structure, Equi. Diagram, Experimental, 7) Yang, H.S., Jin, P., Mukherjee, A.K., “Superplastic Behavior of Regular 2 and Super 2 Titanium Aluminides”, Mater. Sci. Eng., A153, 457-464 (1992) (Experimental, 22) Zakharov, A.M., Oleinikova, S.V., Smirnova, T.R., “Phase Equilibria in the Nb-Ti-Al System in the Concentration Range 25-40% Ti and 0-20% Al”, Russ. Metall., (5), 102-105 (1992), transl.: Russ. Akad. Nauk, Metally, (5), 112-116, 1992 (Crys. Structure, Equi. Diagram, Experimental, 3) Zhang, Y., Wang, Y.-B., Chu, W.-Y., Hsiao, C.-M., Thompson, A.W., “Stress Corrosion Cracking of Titanium Aluminide Alloys in Aqueous Solutions and Methanol”, Scr. Metall. Mater., 26, 925-928 (1992) (Experimental, 8) Ebrahimi, F., Hoelzer, D.T., Castillo-Gomez, J.R., “Fracture Toughness of )+x Microstructure in the Nb-Ti-Al System”, Mater. Sci. Eng., A171, 35-45 (1993) (Equi. Diagram, Experimental, 16) Erschbaumer, H., Podloucky, R., Rogl, P., Temnitschka, G., Wagner, R., “Atomic Modelling of Nb, V, Cr and Mn Substitutions in TiAl. I: c/a Ratio and Site Preference”, Intermetallics, 1, 99-106 (1993) (Crys. Structure, Theory, 31) Jackson, A.G., “Identification of the L60 Phase in a -Ti-Al-Nb Alloy”, Scr. Metall. Mater., 28, 673-675 (1993) (Crys. Structure, Experimental, 4) Long, M., Rack, H.J., “Thermo-Mechanical Stability of Forged Ti-25Al-11Nb (at.%)”, Mater. Sci. Eng., A170, 215-226 (1993) (Experimental, Theory, 30) Muraleedharan, K., Banerjee, D., “Phase Transformations Involving the 2 and O Phases in Ti-Al-Nb Alloys”, Scr. Metall. Mater., 29, 527-532 (1993) (Experimental, 16) Nakamura, H., Takeyama, M., Yamabe, Y., Kikuchi, M., “Phase Equilibria in TiAl Alloys Containing 10 and 20 at.% Nb at 1473 K.”, Scr. Metall. Mater., 28, 997-1002 (1993) (Equi. Diagram, Experimental, 10) Nakayama, Y., Mabuchi, H., “Formation of Ternary L12 Compounds in Al3Ti Base Alloys”, Intermetallics, 1, 41-48 (1993) (Crys. Structure, Experimental, 40) Okamoto, H., “Al-Ti (Aluminium - Titanium)”, J. Phase Equilib., 14, 120-121 (1993) (Crys. Structure, Equi. Diagram, Review, 16) Okamoto, H., “Al-Ti (Aluminium - Titanium)”, J. Phase Equilib., 14, 764 (1993) (Equi. Diagram, Review, 5) Rubin, G., Finel, A., “Calculation of Phase Diagrams of Ternary Systems with Cluster Variation - Method Entropy”, J. Phys.: Condens. Matter, 5, 9105-9120 (1993) (Theory, Thermodyn., 34) Soboyejo, W.O., “An Investigation of the Effect of the Heat Treatment on the Microstructure and Mechanical Behavior of 2+ Forged Ti-24Al-11Nb”, in “Titanium´92:
MSIT ®
350
[1993Zha]
[1994Ban] [1994Ben1]
[1994Ben2]
[1994Che1]
[1994Che2]
[1994Hou]
[1994Kum] [1994Sag] [1994Wan]
[1995Ban] [1995Hsi]
[1995Lon]
[1995Mur1]
[1995Mur2] [1995Rub]
[1995Sem] [1995Xu]
MSIT®
Al–Nb–Ti Science and Technology”, Froes, H.F., Caplan, I., (Eds.), Miner., Met. Mater. Soc., 359-366 (1993) (Experimental, 21) Zhang, W.-J., Chen, Q.-Z., Wang, Y.-D., Sun, Z.-Q., “Characteristics of Heat Resistant Alloys Ti10Nb45 Al and Ti18Nb48Al”, Scr. Metall. Mater., 28, 1113-1118 (1993) (Crys. Structure, Equi. Diagram, Experimental, 12) Banerjee, D., “Is There an Ordered Tetragonal Phase in the Ti3Al-Nb System?”, Scr. Metall. Mater., 30, 855-858 (1994) (Crys. Structure, Theory, 14) Bendersky, L.A., Roytburd, A., Boettinger, W.J., “Phase Transformations in the (Ti, Al) 3Nb Section of the Ti-Al-Nb System. - I. Microstructural Predictions Based on a Subgroup Relation between Phases”, Acta Metall. Mater., 42, 2323-2335 (1994) (Crys. Structure, Theory, 36) Bendersky, L.A., Boettinger, W.J., “Phase Transformations in the (Ti, Nb)3Al Section of the Ti-Al-Nb System. - II. Experimental TEM Study of Microstructures”, Acta Metall. Mater., 42, 2337-2352 (1994) (Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 19) Chen, Z., Jones, I.P., Saunders, N., Small, C.J., “Characterization of Phases in Ti-42Al-8Nb Alloy at 1200°C”, Scr. Metall. Mater., 30, 1403-1408 (1994) (Equi. Diagram, Experimental, 9) Chen, G.L., Wang, J.G., Sun, Z.Q., Ye, H.Q., “Continuous Ordering in the TiAl+Nb System”, Intermetallics, 2, 31-36 (1994) (Crys. Structure, Equi. Diagram, Experimental, 24) Hou, D.H., Yang, S.S., Shyue, J., Fraser, H.L., “Investigation of B2 and Related Phases in Ti-Modified Nb-Al Alloys”, Mater. Res. Soc. Symp. Proc., 322, 437-442 (1994) (Crys. Structure, Equi. Diagram, Experimental, 19) Kumar, S.G., Reddy, R.G., Brewer, L., “Phase Equilibria in Ti3Al-Nb Alloys at 1000°C”, J. Phase Equilib., 15, 279-284 (1994) (Equi. Diagram, Experimental, 16) Sagar, P.K., Banerjee, D., Prasad, Y.V.R.K., “Processing of an -2 Aluminide Alloy, Ti-24Al-11Nb”, Mater. Sci. Eng., A177, 185-197 (1994) (Experimental, 21) Wang, J., Chen, G., Sun, Z., Ye, H., “Structure of a New Ordered Ternary Intermetallic Compound in TiAl+Nb System” (in Chinese), Acta Metall. Sin. (China), 30, A525-A531 (1994) (Crys. Structure, Experimental, 11) Banerjee, D., “Deformation of the O and 2 Phases in the Ti-Al-Nb System”, Philos. Mag. A., 72, 1559-1587 (1995) (Experimental, 39) Hsing, L.M., Wadley, H.N.G., “Time-Temperature Transformation Behavior of Ti-24Al-11Nb”, Mater. Sci. Eng., A192/193, 908-913 (1995) (Crys. Structure, Experimental, 11) Long, M., Rack, H.J., “Phase Stability During Continuous Heating/Cooling of TiAl-(Nb, V, Mo) Titanium Aluminide Alloys”, Mater. Sci. Technol., 11, 150-158 (1995) (Equi. Diagram, Experimental, 31) Muraleedharan, K., Nandy, T.K., Banerjee, D., “Phase Stability and Ordering Behaviour of the O phase in Ti-Al-Nb Alloys”, Intermetallics, 3, 187-199 (1995) (Crys. Structure, Equi. Diagram, Experimental, Theory, #, 30) Muraleedharan, K., Banerjee, D., “The 2-to-O Transformation in Ti-Al-Nb Alloys”, Philos. Mag., 71, 1011-1036 (1995) (Crys. Structure, Experimental, Theory, 24) Rubin, G., Finel, A., “Application of First-Order Principles Methods to Binary and Ternary Alloy Phase Diagram Predictions”, J. Phys.: Condens. Matter, 7, 3139-3152 (1995) (Equi. Diagram, Theory, 30) Semiatin, S.L., Smith, P.R., “Microstructural Evolution During Rolling of Ti-22Al-23Nb Sheet”, Mater. Sci. Eng., A202, 26-35 (1995) (Experimental, 13) Xu, R., Li, D., Cui, Y., Xu, D., Li, Q., Hu, Z., “A New Phase in Rapidly Solidified Ti 3Al-Based Alloys”, Scr. Metall. Mater., 32, 305-308 (1995) (Crys. Structure, Experimental, 6)
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti [1995Zdz]
[1996Che]
[1996Ebr] [1996Jac]
[1996Men]
[1996Sad]
[1996Sik]
[1996Tre]
[1996Vas]
[1996Woo]
[1997Bul]
[1997Jew]
[1997Nak]
[1997Wan]
[1998Che]
[1998Din]
[1998Hel]
Landolt-Börnstein New Series IV/11A3
351
Zdziobek, A., Durand-Charre, M., Driole, J., Durand, F., “Experimental Investigation of High Temperature Phase Equilibria in the Nb-Al-Ti System”, Z. Metallkd., 86, 334-340 (1995) (Crys. Structure, Equi. Diagram, Experimental, #, 23) Chen, G.L., Wang, X.T., Ni, K.Q., Hao, S.M., Cao, J.X., Ding, J.J., Zhang, X., “Investigation on the 1000, 1150 and 1400°C Isothermal Section of the Ti-Al-Nb System”, Intermetallics, 4, 13-22 (1996) (Crys. Structure, Equi. Diagram, Experimental, 27) Ebrahimi, F., Ruiz-Aparicio, J.G.L, “Diffusivity in the Nb-Ti-Al Ternary Solid Solution”, J. Alloys Compd., 245, 1-9 (1996) (Experimental, 16) Jacob, V., Colinet, C., Desre, P., Moret, F., “Calculation of the A2/B2 Phase Boundary in the Nb-Ti-Al System with the Cluster Variation Method” (in French), J. Phys. IV, Col. 2, 6, C2-3-C2-10 (1996) (Equi. Diagram, Theory, 17) Menon, E.S.K., Subramanian, P.R., Dimiduk, D.M., “Phase Transformations in Nb-Al-Ti Alloys”, Metall. Mater. Trans., 27A, 1647-1659 (1996) (Crys. Structure, Equi. Diagram, Experimental, 30) Sadi, F., Servant, C., “Transformation During Continuous Cooling of the Ti4Al3Nb Alloy” (in French), J. Phys. IV, Colloq 2, 6, C2-241-C2-246 (1996) (Equi. Diagram, Experimental, 14) Sikora, T., Hug, G., Jaouen, M., Flank, A.-M., “EXAFS Study of the Local Atomic Order in Ti2AlX (X = Nb, Mo) B2 Intermetallic Compounds”, J. Phys. IV, Colloq. 2, 6, C2-15-C2-30 (1996) (Crys. Structure, Experimental, 8) Tretyachenko, L.A., “On the Ti-Al System”, “Phase Diagrams in Material Science”, Fifth International School, Katsyveli, Crimea, Ukraine, 118 (1996) (Equi. Diagram, Experimental, #, 0) Vasudevan, V.K., Yang, J., Woodfield, A.P., “On the to B2 Ordering Temperature in a Ti-22Al-26Nb Orthorhombic Titanium Aluminide”, Scr. Mater., 35, 1033-1039 (1996) (Crys. Structure, Equi. Diagram, Experimental, 10) Woodward, C., MacLaren, J.M., “Planar Fault Energies and Sessile Dislocation Configurations in Substitutionally Disordered Ti-Al with Nb and Cr Ternary Additions”, Philos. Mag. A, 74, 337-357 (1996) (Crys. Structure, Theory, 27) Bulanova, M., Tretyachenko, L., Golovkova, M., “Phase Equilibria in Ti-Rich Corner of the Ti-Si-Al System”, Z. Metallkd., 88, 256-267 (1997) (Crys. Structure, Equi. Diagram, Experimental, #, 15) Jewett, T.J., “Comment on “Investigation on the 1000, 1150 and 1400°C Isothermal Section of the Ti-Al-Nb System””, Intermetallics, 5, 157-159 (1997) (Crys. Structure, Equi. Diagram, Experimental, 14) Naka, S., Khan, T., “Designed Novel Multiconstituent Intermetallics: Contribution of Modern Alloy Theory in Developing Engineered Materials”, J. Phase Equilib., 18, 635-649 (1997) (Review, 17) Wang, J.G., Zhang, L.C., Chen, G.L., Ye, H.Q., “Formation of Stress-Induced 9R Structure in a Hot-Deformed Ti-45Al-10Nb Alloy”, Scr. Mater., 37, 135-140 (1997) (Crys. Structure, Experimental, 21) Chen, G.L., Wang, J.G., Wang, X.T., Ni, X.D., Hao, S.M., Ding, J.J., “Reply to the “Comment on Investigation on the 1000, 1150 and 1400°C Isothermal Section of the Ti-Al-Nb System”” - Part I. Ordering of Nb in -TiAl and 1-Phase”, Intermetallics, 6, 323-327 (1998) (Crys. Structure, Experimental, 12) Ding, J.-J., Hao, S.-M., “Reply to the “Comment on Investigation on the 1000, 1150 and 1400°C Isothermal Section of the Ti-Al-Nb System”” - Part II. Modification of 1000 and 1150°C Isothermal Sections of the Ti-Al-Nb System”, Intermetallics, 6, 329-334 (1998) (Crys. Structure, Equi. Diagram, Experimental, 15) Hellwig, A., Palm, M., Inden, G., “Phase Equilibria in the Al-Nb-Ti System at High Temperatures”, Intermetallics, 6, 79-94 (1998) (Crys. Structure, Equi. Diagram, Experimental, #, 57) MSIT ®
352 [1998Rho] [1998Sag]
[1998Ser]
[1998Tak]
[1998Wan]
[1998Woo]
[1998Yu] [1999Boe]
[1999Cha1]
[1999Cha2]
[1999Eck]
[1999Flo] [1999Hao]
[1999Rav]
[1999Yan]
[2000Ama] [2000Che]
[2000Izu]
MSIT®
Al–Nb–Ti Rhodes, C.G., “Order/Disorder Temperature of the bcc Phase in Ti-21Al-26Nb”, Scr. Mater., 38, 681-685 (1998) (Equi. Diagram, Experimental, 10) Sagar, P.K., Prasad, Y.V.R.K., “Hot Deformation and Microstructural Evolution in an 2/O Titanium Aluminide Alloy Ti-25Al-15Nb”, Z. Metallkd., 89, 433-441 (1998) (Experimental, 25) Servant, C., Ansara, I., “Thermodynamic Assessment of the Al-Nb-Ti System”, Ber. Bunsenges. Phys. Chem., 102, 1189-1205 (1998) (Equi. Diagram, Review, Thermodyn., #, 76) Takeyama, M., Ohmura, Y., Kikuchi, M., Matsuo, T., “Phase Equilibria and Microstructural Control of TiAl Based Alloys”, Intermetallics, 6, 643-646 (1998) (Equi. Diagram, Review, 20) Wang, X.T., Chen, G.L., Ni, K.Q., Hao, S.M., “The 1400°C Isothermal Section of the Ti-Al-Nb Ternary System”, J. Phase Equilib., 19, 200-205 (1998) (Equi. Diagram, Experimental, #, 18) Woodward, C., Kajihara, S., “Site Preferences and Formation Energies of Substitutional Si, Nb, Mo, Ta, and W Solid Solution in L10 Ti-Al”, Phys. Rev. B, 57, 13459-13470 (1998) (Crys. Structure, Theory, Thermodyn., 45) Yu, T.H., Koo, C.H., “Phase Characterization of a Hot-Rolled Ti-40Al-10Nb Alloy at 1000 to 1200°C”, Scr. Mater., 39, 915-922 (1998) (Equi. Diagram, Experimental, 9) Boehlert, C.J., “The Phase Evolution and Microstructural Stability of an Orthorhombic Ti-23Al-27Nb Alloy”, J. Phase Equilib., 20, 101-108 (1999) (Equi. Diagram, Experimental, 17) Chaumat, V., Colinet, C., Moret, F., “Study of Phase Equilibria in the Nb-Ti-Al System Theoretical Study: CVM Calculation of the Phase Diagram of bcc Nb-Ti-Al”, J. Phase Equilib., 20, 389-398 (1999) (Equi. Diagram, Experimental, Theory, 22) Chaumat, V., Ressouche, E., Ouladdiaf, B., Desre, P., Moret, F., “Experimental Study of Phase Equilibria in the Nb-Ti-Al System”, Scr. Mater., 40, 905-911 (1999) (Crys. Structure, Equi. Diagram, Experimental, 14) Eckert, M., Kath, D., Hilpert, K., “Thermodynamic Activities in the Alloys of the Ti-Al-Nb System”, Metall. Mater. Trans., 30A, 1315-1326 (1999) (Equi. Diagram, Experimental, Thermodyn., #, 44) Flower, H.M., Christodoulou, J., “Phase Equilibria and Transformation in Titanium Aluminides”, Mater. Sci. Technol., 15, 45-52 (1999) (Equi. Diagram, Review, 46) 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, 1129-1139 (1999) (Crys. Structure, Equi. Diagram, Experimental, 41) Ravi, C., Vajeeston, P., Mathijaya, S., Asokamani, R., “Electronic Structure, Phase Stability and Cohesive Properties of Ti2XAl (X = Nb, V, Zr)”, Phys. Rev. B, 60, 15683-15690 (1999) (Crys. Structure, Equi. Diagram, Theory, 32) Yang, R., Hao, Y.L., “Estimation of ( + 2) Equilibrium in Two-Phase Ti-Al-X Alloys by Means of Sublattice Site Occupancies of X in TiAl and Ti3Al”, Scr. Mater., 41, 341-346 (1999) (Equi. Diagram, Theory, 13) Amancherla, S., Banerjee, R., Banerjee, S., Fraser, H.L., “Ordering in Ternary B2 Alloys”, Inter. J. Refract. Met. Hard Mater., 18, 245-252 (2000) (Equi. Diagram, Theory, 23) Cheng, Z.Y., Du, X.W., Zhu, J., Cao, C.X., Sun, F.S., “ 2 Phase Transformation in Fractured High Temperature Stress Rupture Ti-48Al-2Nb (at.%)”, J. Mater. Sci., 35, 4501-4505 (2000) (Experimental, 22) Izumi, T., Yoshioka, T., Hayashi, S., Narita, T., “Sulfidation Properties of TiAl-2 at.% X (X = V, Fe, Co, Cu, Nb, Mo, Ag and W) Alloys at 1173 K and 1.3 Pa Sulfur Pressure in an H 2S-H2 Gas Mixture”, Intermetallics, 8, 891-901 (2000) (Experimental, 42)
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti [2000Kai]
[2000Leo1]
[2000Leo2]
[2000Mur] [2000Oka] [2000Sad]
[2000Wel] [2000Yan]
[2001Bra]
[2001Cao]
[2001Col]
[2001Kan]
[2001Li1]
[2001Li2] [2001Mis]
[2001Sad]
[2001Ser]
[2001Sun]
Landolt-Börnstein New Series IV/11A3
353
Kainuma, R., Fujita, Y., Mitsui, H., Ishida, K., “Phase Equilibria Among (hcp), (bcc) and (L10) Phases in Ti-Al Base Ternary Alloys”, Intermetallics, 8, 855-867 (2000) (Equi. Diagram, Experimental, #, 29) Leonard, K.J., Mishurda, J.C., Molloseau, B., De Graef, M., Vasudevan, V.K., “Identification of a New Tetragonal Phase in the Nb-Ti-Al System”, Philos. Mag. Lett., 80, 295-305 (2000) (Crys. Structure, Experimental, 6) Leonard, K.J., Mishurda, J.C., Vasudevan, V.K., “Examination of Solidification Pathways and the Liquidus Surface in the Nb-Ti-Al System”, Metall. Mater. Trans. B, 31B, 1305-1321 (2000) (Crys. Structure, Equi. Diagram, Experimental, #, 21) Murty, S.V.S., Rao, B.N., Kashyap, B.P., “Development of a Processing Map for the Hot Working of Ti-25Al-15Nb”, Z. Metallkd., 91, 769-774 (2000) (Theory, 8) Okamoto, H., “Al - Ti (Aluminium - Titanium)”, J. Phase Equilib., 21, 311 (2000) (Equi. Diagram, Review, 2) Sadi, F.-A., Servant, C., “In Situ Neutron Diffraction on the Alloy 50.6Ti-36.5Al-12.9Nb (at.%)”, Z. Metallkd., 91, 504-509 (2000) (Crys. Structure, Equi. Diagram, Experimental, 21) Weller, M., Chatterjee, A., Haneczok, G., Clemens, H., “Internal Friction of -TiAl Alloys at High Temperature”, J. Alloys Compd., 310, 134-138 (2000) (Experimental, 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, 296-301 (2000) (Crys. Structure, Equi. Diagram, Review, Theory, 38) Braun, J., Ellner, M., “Phase Equilibria Investigations on the Aluminium-Rich Part of the Binary System Ti-Al”, Metall. Mater. Trans., 32A, 1037-1047 (2001) (Crys. Structure, Equi. Diagram, Experimental, 34) Cao, G.H., Liu, Z.G., Shen, G.J., Liu, J.-M., “Identification of a Cubic Precipitate in -Titanium Aluminides”, J. Alloys Compd., 325, 263-268 (2001) (Crys. Structure, Experimental, 16) Colinet, C., “Applications of the Cluster Variation Method to Empirical Phase Diagram Calculations”, Calphad, 25, 607-623 (2001) (Equi. Diagram, Review, Theory, Thermodyn., 108) 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) (Equi. Diagram, Theory, 15) Li, J., Jiang, M., Hao, S., Li, S., Zhong, Z., “Thermodynamic Calculation of the / and / Phase Equilibria in the Ti-Al-Nb Ternary System” (in Chinese), Acta Metall. Sin. (China), 37, 1064-1068 (2001) (Equi. Diagram, Thermodyn., Theory, 13) Li, Zh., Gao, W., He, Y., “Protection of a Ti3Al-Nb Alloy by Electro-Spark Deposition Coating”, Scr. Mater., 45, 1099-1105 (2001) (Experimental, 23) Mishurda, J.C., Vasudevan, V.K., “An Estimate of the Kinetics of the 0 to Orthorhombic Phase Transformation in the Nb-Ti-Al System”, Scr. Mater., 45, 677-684 (2001) (Equi. Diagram, Experimental, 14) Sadi, F.A., Servant, C., Cizeron, G., “Phase Transformations in Ti-29.7Al-21.8Nb and Ti-23.4Al-31.7Nb (at.%) Alloys”, Mater. Sci. Eng., A311, 185-199 (2001) (Crys. Structure, Equi. Diagram, Experimental, #, 20) Servant, C., Ansara, I., “Thermodynamic Modelling of the Order-Disorder Transformation of the Orthorhombic Phase of the Al-Nb-Ti System”, Calphad, 25, 509-525 (2001) (Equi. Diagram, Theory, Thermodyn., #, 17) Sun, F.-S., Cao, C.-X., Kim, S.-E., Lee, Y.-T., Yan, M.-G., “Alloying Mechanism of Beta Stabilizers in a TiAl Alloy”, Metall. Mater. Trans., 32A, 1573-1589 (2001) (Crys. Structure, Equi. Diagram, Experimental, 37)
MSIT ®
Al–Nb–Ti
354 [2001Wan]
[2001Zha]
[2002Bou]
[2002Hod]
[2002Leo1]
[2002Leo2]
[2002Wu]
[2003Kar]
[2003Sad] [2003Sch]
Wang, X., Chang, H., Lei, M., “Thermodynamic Aspects of Oxidation for Nb Alloying -TiAl Intermetallic Compounds”, Acta Metall. Sin. (China), 37, 810-814 (2001) (Theory, Thermodyn., 20) Zhang, L.T., Ito, K., Vasudevan, V.K., Yamaguchi, M., “Beneficial Effects of O-Phase on the Hydrogen Absorption of Ti-Al-Nb Alloys”, Intermetallics, 9, 1045-1052 (2001) (Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 13) Bououdina, M., Guo, Z.X., “Characterization of Structural Stability of (Ti(H2)+22Al+23Nb) Powder Mixtures During Mechanical Alloying”, Mater. Sci. Eng., A332, 210-222 (2002) (Crys. Structure, Experimental, 20) Hodgson, A.W.E., Mueller, Y., Forster, D., Virtanen, S., “Electrochemical Characterization of Passive Films on Ti Alloys under Simulated Biological Conditions”, Electrochim. Acta, 47, 1913-1923 (2002) (Experimental, 55) Leonard, K.J., Mishurda, J.C., Vasudevan, V.K., “Phase Equilibria at 1100°C in the Nb-Ti-Al System”, Mater. Sci. Eng., A329-331, 282-288 (2002) (Crys. Structure, Equi. Diagram, Experimental, 25) Leonard, K.J., Vasudevan, V.K., “Site Occupancy Preferences in the B2 Ordered Phase in Nb-Rich Nb-Ti-Al Alloys”, Mater. Sci. Eng., A329-331, 461-467 (2002) (Crys. Structure, Equi. Diagram, Experimental, 19) Wu, B., Shen, J., Chu, M., Shang, Sh., Zhang, Z., Peng, D., Liu, S., “The Ordering Behaviour of the O Phase in Ti2AlNb-Based Alloys”, Intermetallics, 10, 979-984 (2002) (Crys. Structure, Theory, Thermodyn., 10) Karpets, M.V., Milman, Yu.V., Barabash, O.M., Korzhova, N.P., Senkov, O.N., Miracle, D.B., Legkaya, T.N., Voskoboynik, I.V., “The Influence of Zr Alloying on the Structure and Properties of Al3Ti”, Intermetallics, 11, 241-249 (2003) (Crys. Structure, Experimental, 16) Sadi, F.A., Servant, C., “On the B2O Phase Transformation in Ti-Al-Nb Alloys”, Mater. Sci. Eng., A346, 19-28 (2003) (Crys. Structure, Equi. Diagram, Experimental, Theory, 28) Schmid-Fetzer, R., “Al - Ti (Aluminium - Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 85)
Table 1: Experimental Investigations after [1993Gam] and some Earlier Works Achievement isothermal section at 1200°C
Sample Preparation and Characterization 14 arc-melted alloys, annealed at 1200°C for two weeks; diffusion couples; LOM, SEM-EMPA, XRD isothermal sections at 1700 and 750°C; 36 arc-melted alloys, annealed at 1700°C for < 30 at.% Ti and 30 at.% Al; 2.5 - 3.0 at.% 25 h and at 750°C for 500 h; LOM, XRD Ti solubility in Nb3Al structure of Nb rich alloys alloys at 25 and 20 at.% Al, containing < 5 at.% Ti; annealed at 800°C, 500 h, water quenched after homogenization at 1700°C for 300 h. LOM, XPD and EMPA Nb3Al based solid solutions; Ti solubility (Nb,Ti)3Al up to ~14 at.% Ti; arc-melted, in Nb3Al at 700°C > 10 at.% homogenized at 1650°C for 3 h, annealed at 700°C for 250 h; XRD, Tc (superconducting transition)
MSIT®
References [1989Jew]
[1975Pan]
[1975Fed]
[1975Sha]
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti Achievement phase equilibria in the Nb rich region up to 40 mass% Al and 40 mass% Ti (~70 at.% Al, 50 at.% Ti); partial liquidus and solidus projections and isothermal sections at 1200, 900, 600°C; invariant equilibria L +
+ ) at 1950°C, L + + ) at 1750°C and L ) + J + at 1470°C composition-temperature section at 6 mass% Al for 25 - 35 mass% Ti (from Ti-55.4Nb-15.5Al to Ti-40Nb-14Al (at.%) from solidus at ~2000°C down to 600°C partial isothermal sections at 1100, 900 and 800°C for the range of 0 - 20 mass% (0 to ~46 at.%) Al and from 25 to 40 mass% Ti (~39 to 56 at.%)
Sample Preparation and Characterization References [1992Pav1, arc melted alloys, step-wise annealed: 1400°C/10 h + 1200°C/50 h + 900°C/100 h + 1992Pav2] 600°C/150 h, water quenched from 1200 600°C; LOM, XRD, EMPA, and solidus temperature measurements
arc melted alloys, step-wise annealed: 1400°C/10 h + 1200°C/50 h + 900°C/100 h + 600°C/150 h, water quenched from 1200 600°C; LOM, XRD, EMPA arc melted alloys, step-wise annealed: (1400°C/5 h + 1300°C/30 h + 1100°C/100 h) + (1100°C/2 h + 900°C/300 h + 800°C/500 h), water quenched from 1100 - 800°C; LOM, XRD, EMPA structure of Ti-60 mass% Nb alloy with 1 - TEM, XRD of alloys quenched from 1150°C, 8 at.% Al (up to ~40 at.% Nb and ~12 at.% aged at 400 - 900°C Al) in the temperature range 1150 - 400°C; ordering of the bcc phase and precipitation of orthorhombic O phase structure of Ti-60Nb-8Al (mass%) alloy TEM, XRD 5 alloys arc melted, annealed at 1650, 1200 boundary of the and phases at 1650, 1200 and 1000°C in the Nb corner; the and 1000°C for 50 h, 14 d and 30 d, respectively; LOM, XRD, SEM, TEM, EMPA phase ordering phase relations in Ti-Nb-15Al alloys up to plasma arc melted alloys; TEM, EMPA and ALCHEMI 40 at.% Ti in the temperature range > 800°C; site occupancy in the ordered (B2) phase ordering and phase transformations in the extruded at ~1232°C, annealed in the and Ti3Al based alloy with ~5 at.% Nb + fields and quenched, then annealed at 700 - 1000°C and again quenched; XRD, TEM phase transformations alloys Ti-(10-20)Nb-25Al on quenching and low-temperature aging at 400 - 500°C; TEM and 2 phase boundaries up to 7.5 at.% Nb containing Ti-(34-38) mass% Al alloys Nb at 1000°C aged at 1000°C for 605 ks; LOM, XRD, EMPA 5 alloys, arc melted, homogenized at 1400°C phase equilibria in the region around Ti2NbAl (transformations of the bcc phase for 3 h, annealed at 1100°C for 4 d; LOM, to the B2 and Ti4NbAl3 phases) TEM, SAD phase transformations (ordering of the bcc Ti-(0-30)Nb-25Al alloys both bulk and melt phase, the O phase formation) spun ribbons heat treated at 700 -800°C; LOM, XRD, SEM, TEM compositions of the 0 and 2 phases in the analytical electron microscopy technique Ti-11Nb-24Al alloy in the temperature range of 1200 - 1020°C phase transformations to Ti4NbAl3 1400 - 700°C; LOM, TEM, SEM
Landolt-Börnstein New Series IV/11A3
355
[1991Zak]
[1991Smi, 1992Smi, 1992Zak]
[1992Voz]
[1992Tre] [1992Men, 1996Men] [1994Hou]
[1977Sas]
[1982Str, 1988Str] [1988Has]
[1989Ben]
[1989Kes]
[1989Mur]
[1990Ben1, 1990Ben2] MSIT ®
356 Achievement ordering, structure of ordered phases (0, O)
Al–Nb–Ti
Sample Preparation and Characterization Ti-10Nb-25Al; Ti-12.5Nb-25Al; channeling enhanced microanalysis, convergent beam electron diffraction (CBED) effect of heat treatment on microstructure Ti-11Nb-24Al hot-rolled sheets annealed at 1000 and 1200°C, WQ or furnace cooled (FC); LOM, XRD, Vicker’s hardness measurements influence of cooling rate on microstructure Ti-11Nb-25Al, solution treated at 1150°C for and creep properties 45 min and cooled with rates from 0.02 K#s-1 to 10 K#s-1 or aged at 750°C for 24 h; LOM, SEM, creep testing continuous cooling transformations Ti-11Nb-24Al, wedge-shaped specimens, heated at 1230°C for 1 h, cooled in ice water; Ti-20Nb-24Al annealed at 1250°C for 8 h, air cooled; LOM, SEM, TEM, hardness measurements Ti-12.5Nb-25Al; extruded and heat treated at microstructure and compositions of the 1040°C for 1 h, aged at 760°C for 1 h, creep phases (/0, 2, O) tested at 650°C; TEM, EMPA behavior of ordering transformation Ti-21Nb-14Al (mass%), arc melted, forged, rolled, annealed at 1060°C for 0.5 h, WQ or air cooled, aged at 700°C for 1 h; TEM, XRD, SAD partial composition-temperature section at arc melted alloys Ti2NbAl and Ti4Nb3Al, 50 at.% Ti homogenized at 1400°C, annealed at 700°C for 26 d; TEM, LOM, SEM microstructure of the Ti-20Nb-3Al alloy arc melted, homogenized at 1400°C, heat treated in the range 1100 - 700°C; TEM, LOM Ti-11Nb-24Al; TEM, SAD, microdiffraction study of microstructure and evolution of (MD); CBED phases; reaction sequence during isothermal aging at 650 and 850°C including a new transition T phase; crystal structure of the T phase and structural relationships between T and O phases transformations during aging at 450 plasma-sprayed alloy Ti-11Nb-24Al, TEM, 850°C involving transition metastable XRD phases other than in [1992Hsi1, 1992Hsi4] phase transformations from the to O Ti-15Nb-24Al alloy, various heat treatments phase: O phase exists in two forms in the temperature range from 1200 to 650°C; LOM, TEM, and EMPA phase transformations in the temperature Ti-13Nb-28.5Al alloy; TEM range from 900 to 400°C vertical section Ti-27.5Al up to 25 at.% Nb six Ti-(12.5-25)Nb-27.5Al alloys, arc melted, from 1200 to 700°C; refined version of the heat treated in the range 1200 to 700°C, water section at 27.5 at.% Al with both forms of cooled; EMPA, TEM and SAD the O phase; formation of the O phase by the peritectoid reaction 2 + 0 O transformation from 2 to O at isothermal alloy Ti-13Nb-28.5Al (at.%); TEM, SAD, aging at 900°C from 15 min to 200 h CBED MSIT®
References [1987Ban, 1988Ban] [1990Kno]
[1990Mis]
[1990Wey]
[1991Akk]
[1991Li]
[1991Ben]
[1992Ben] [1992Hsi1, 1992Hsi4]
[1992Hsi2, 1992Hsi3] [1995Hsi]
[1992Mur1, 1992Mur2] [1993Mur] [1995Mur1]
[1995Mur2]
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti Achievement transformation temperatures from 600 to 1300°C phase equilibria near Ti3Al at 1000°C possible transformation paths from high temperature bcc/B2 to low temperature hcp or O phase fields were predicted schematic pseudobinary Ti3Al-Nb 3Al section up to ~35 at.% Nb the structure of alloys in the vicinity of TiAl, identification of the L60 structure the structure of alloys in the vicinity of TiAl the structure of alloys in the vicinity of TiAl ( + 0 + , 0 + + ))
Sample Preparation and Characterization Ti-11Nb-25Al alloy; calorimetric differential thermal analysis (CDTA); in situ high temperature XRD four sintered alloys, equilibrated for 225 h LOM, XRD and EMPA alloys in the Ti3Al-Nb3Al section
TEM study of three alloys in the Ti3Al-Nb3Al section annealed at 1100 and 700°C Ti-20Nb-43Al at 1200°C; TEM, electron diffraction alloy Ti-2.14Nb-47.2Al, plasma arc melted, annealed at 1050°C for 96 h, LOM, XRD three alloys in a region of Ti-(~10-20)Nb-(~40-45)Al, 1200°C for 33 h; SEM, EMPA the structure of alloys in the vicinity of alloys Ti-10Nb-45Al and Ti-18Nb-48Al, TiAl 1200°C/24 h; TEM, SEM, XRD, electron diffraction vertical sections at 10 at.% Nb, 48 at.% Al; + 0 + phase equilibria (/2 + /0 + ) Ti-10Nb-40Al, from 1200°C/24 h to 1000°C/400 h, XRD, TEM, EMPA phase equilibria between , (0) and at LOM and EMPA 1300 and 1250°C up to ~15 at.% Nb effects of Nb on the microstructure and in alloys (Ti52Al48-xNbx (0 x 6 at.%), arc phase constituents 2, and melted, hot isostatic pressing at 1200°C for 3 h, annealed at 1200°C for 12 h, aged at 900°C for 8 h; LOM, SEM, EMPA, XRD, X-ray photoelectron spectroscopy (XPS) high temperature phase equilibria; liquidus alloys inductively melted, homogenized at projection, isothermal section at 1200°C; 1300°C for 20 h and annealed at 1200°C for invariant reactions of [1990Per] two weeks; SEM, XRD, EMPA and TEM; confirmed; no ternary phases at 1200°C melting temperatures measured with a pyrometer isothermal sections at 1200 and 1000°C; diffusion couples and bulk samples annealed neither TiNbAl3 ( 1) nor T1 [1989Jew] or at 1200°C for 48 h, at 1000°C for 96 h and water quenched; LOM, EMPA, TEM and T2 [1989Jew, 1990Per] were found; no XRD other ternary compounds; separate area of the ordered B2 phase detected at 1000°C; considerable solid solubilities of the third element in most of the binary phases phase relations involving Ti4NbAl3 alloy Ti-20Nb-30Al, 1100°C/24 h, 900 700°C up to 18 d, TEM, SAD
Landolt-Börnstein New Series IV/11A3
357 References [1995Lon]
[1994Kum] [1994Ben1]
[1994Ben2] [1992Jac] [1992Kim] [1993Nak1]
[1993Zha]
[1998Tak] [1998Yu] [2000Kai] [2001Sun]
[1995Zdz]
[1998Hel]
[1992Ben]
MSIT ®
358 Achievement phase relations involving Ti4NbAl3
phase relations involving Ti4NbAl3
liquidus projection by [1995Zdz] was proposed to be changed with respect to the wider field of primary crystallization without changing the nature and direction of the liquid phase reactions; the solid state transformations were considered ( ) + , massive transformation, + ) eutectoid-like transformation)
phase transformations
phase evolutions from B2 to the O phase
effect of cooling rate on the transformations from B2 to the O phase
phase evolutions from B2 to the O phase
phase evolutions from B2 to the O phase
MSIT®
Al–Nb–Ti Sample Preparation and Characterization alloy Ti-10.8Nb-36.9Al, continuous cooling from 1200°C at various rates, optical pyrometry, DTA, dilatometry, thermoresistometry, differential microcalorimetry, electron microscopy, TEM, Vicker’s hardness alloy Ti-12.9Nb-36.5Al, arc melted, annealed at 1300°C/6 h, SEM, in situ neutron diffraction 25 - 960°C 15 alloys in the range 15 to 40 at.% Al with Nb:Ti ratios of 4:1. 2:1, 1.5:1 and 1:1.5; alloys were arc melted, homogenized through hot isostatic pressing (HIP) at 1425 and 1475°C for times up to 7 h at 138 MPa followed by water or oil quenching; LOM, XRD, DTA, BSEI, EMPA, TEM and microhardness measurements; data on phase equilibria in the same 15 alloys annealed at 1100°C for 720 h examined by BSEI, optical microscopy, XRD and EMPA alloys Ti-21.8Nb-27.9Al and Ti-31.7Nb-23.4Al prepared by vacuum arc melting, homogenized at 1300°C during 1 week and annealed at 1260°C for 20 and 70 h, at 1100°C 20 and 75 h, at 900°C 140 and 1000 h, at 700°C 1500 h with ice-water quenching after each heat treatment; microhardness measurements, dilatometry, DTA, XRD, SEM, TEM; continuous cooling of the alloys from 1260°C with rates from 100 to 0.25 K·s-1 alloy Ti-27Nb-23Al rolled sheet, 650 1090°C, up to 450 h; EMPA, SEM, TEM, XRD, DTA alloys Ti-37.5Nb-25Al, Ti-35Nb-30Al, Ti-~44.5Nb-~25.6Al; DTA up to 1500°C, SEM, TEM, electrical resistivity measurements three alloys around Ti2NbAl, arc melted, annealed at 1200°C for 3 h, aged at 600 900°C for 0.5 to 300 h; XRD, LOM, TEM, SEM three alloys around Ti2NbAl, annealed at 1350, 900, 800 and 700°C up to 1500 h and quenched; XRD, SEM, TEM
References [1996Sad]
[2000Sad]
[2000Leo2]
[2001Sad]
[1999Boe]
[2001Mis]
[2001Zha]
[2003Sad]
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
359
Table 2: Crystallographic Data of Solid Phases Phase / Temperature Range [°C] (Al) < 664.2 < 660.452 , (Ti1-x-yNbxAly)
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype cF4 Fm3m Cu a = 404.96 cI2 Im3m W
a = 330.07 a = 327.6 0.3 a = 330.65 a = 328
(Nb) < 2469 (Ti)h 1670 - 882
a = 327 5 a = 327.3 to 337.5 a = 327.3 to 328.5 a = 326 a = 330.3 * 0
cP2 Pm3m CsCl
a = 323.5 a = 324 3 a = 326 3 a = 326.6 a = 325.1 a = 326.9 a = 327.1 a = 327.5 a = 326.8 a = 324.4 a = 326.1 a = 323.05 0.05 a = 322.50 0.05 a = 325 a = 330 a = 328
Landolt-Börnstein New Series IV/11A3
Comments /References
0 to 0.6 at.% Ti [1992Kat2] [V-C2] pure Al at 25°C [1981Kin] 0 x 1 >882°C at y = 0 [Mas2, 1987Mur] 0 y 0.448 at x = 0 [1993Gam, 1993Oka1, 2000Oka, 2003Sch] 0 y 0.46 at x = 0 [1996Tre, 1997Bul] 0 y 0.215 at x + y = 1 [Mas2] pure Nb at 25°C [1981Kin] for Nb-21.5 at.% Al [1980Jor] [Mas2] for Ti-45Nb-10Al in Ti-37.2Nb-12.2Al (at.%) alloy annealed at 700°C for 26 d, [1991Ben] for Ti-60.3Nb-10.8 Al alloy homogenized at 1300°C for 20 h [1999Cha1], 900°C [1984Zak] [1983Tro] Ti-47Nb-6.3Al [1992Pav2] Ti-40.8Nb-17.4Al (2++)) 900°C [1992Pav1] ordered form of the high temperature (Ti,Nb,Al) solid solutions [1989Ben, 1989Kes, 1991Ben, 1991Cha, 1992Voz, 1994Hou, 1995Mur1, 1996Men, 1996Vas, 1998Hel, 1998Rho, 1999Cha2, 1999Boe, 1999Flo, 1999Rav, 2000Leo1, 2001Sad, 2002Leo2, 2003Sad] Ti-25Nb-25Al after rapid quench. [1991Cha] for as cast Ti-20.6Nb-26.7Al [1999Cha2] for Ti-54.3Nb-15.4Al, homogenized at 1300°C for 20 h [1999Cha2] Ti-42.5Nb-15Al, as cast [2000Leo2] Ti-37.5Nb-25Al, as cast [2000Leo2] Ti-51Nb-15Al, as cast [2000Leo2] Ti-56.7Nb-15Al, as cast [2000Leo2] Ti-68Nb-15Al, as cast [2000Leo2] Ti-40.90Nb-15.44Al, annealed at 1100°C for 720 h [2002Leo1] Ti-26.8Nb-21.8Al, annealed at 1350°C [2003Sad] Ti-30.2Nb-19.7Al, annealed at 1350°C [2003Sad] Ti-14.4Nb-30.1Al (1200°C) [1998Hel] Ti-16.8Nb-34.6Al (1200°C) [1998Hel] Ti-25.4Nb-25.1Al (1000°C) [1998Hel] Ti-45Nb-25Al, as cast [1995Zdz] Ti-45Nb-10Al [1991Ben] and Ti-11Nb-25Al [1995Lon] MSIT ®
Al–Nb–Ti
360 Phase / Temperature Range [°C] (Ti1-x-yNbxAly) < 1490
Pearson Symbol/ Lattice Parameters [pm] Space Group/ Prototype hP2 P63/mmc Mg
a = 295.03 c = 468.36
(Ti)r < 882
a = 291 c = 469 (Ti1-xNbx)3Al, 2
hP8 P63/mmc Ni3Sn
Ti3Al < 1164 a = 580.6 c = 465.5
Comments /References
at x = 0 47.3 to 51.4 at.% Al at solidus temperatures 1490 to 1462°C [1993Oka1, 1993Oka2, 2000Oka, 1993Gam, 2003Sch] at x = 0 from ~48 at.% Al at 1520°C to 51 at.% Al at 1485°C [1996Tre, 1997Bul] dissolves up to 10 at.% Nb at 1200°C [1998Hel] pure Ti at 25°C [Mas2, V-C2, 1981Kin] dissolves up to ~2 at.% Nb in the Nb-Ti system [Mas2] Ti-5Nb-40Al annealed at 1400°C for 6 h, WQ [1996Che] ~20 to 38.2 at.% Al D019 ordered phase (“2Ti3Al”); maximum at 30.9 at.% Al [1992Kat2, 1993Oka1, 1993Oka2] < 1180°C [1993Gam] maximum at 32.5 at.% Al and 1200°C [1996Tre, 1997Bul] at 22 at.% Al [L-B]
a = 574.6 c = 462.4
at 38 at.% Al [L-B] [V-C]
a = 576.7 0.4 c = 465.4 0.7
Ti-12.4Nb-30.9Al (1000°C) [1998Hel]
a = 580 10 c = 480 10
Ti-11Nb-24Al [1990Wey]
a = 580 10 c = 460 10
in thin films Ti-11Nb-24Al [1992Hsi1]
a = 580 c = 466
in alloy Ti-11Nb-25Al [1995Lon]
a = 574.3 c = 498.4
Ti-13.8Nb-13.4Al (mass%) [1984Zak]
a = 572.4 to 574.3 [1983Tro] c = 498.4
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti Phase / Temperature Range [°C] TiAl, < 1463
Pearson Symbol/ Space Group/ Prototype tP4 P4/mmm AuCu
361
Lattice Parameters Comments /References [pm] a = 400.5 c = 407.0
L1 0 ordered phase (“ TiAl”) 46.7 to 66.5 at.% Al [1992Kat2, 1993Oka1] ~52 to 65 at.% Al at solidus temperatures, ~50 to 60 at.% Al at 1000°C [1996Tre, 1997Bul] 50 to 62 at.% Al at 1200°C [2001Bra]
Landolt-Börnstein New Series IV/11A3
a = 400.0 0.1 c = 407.5 0.1
at 50 at.% Al [2001Bra]
a = 398.4 0.1 c = 406.0 0.1
at 62 at.% Al [2001Bra]
a = 399 c = 408
at 55.4 and 61.8 at.% Al [1998Hel]
a = 399.4 c = 409.6
(Ti0.70 Nb0.30)Al [1991Smi, 1992Zak]
a = 399.3 c = 410.4
Ti-19Nb-53Al ( in a + ) alloy) [1997Jew]
a = 397.9 c = 412.6
Ti-18.9Nb-55.6Al (1200°C) [1998Hel]
a = 398 c = 419
Ti-10Nb-50Al, annealed at 1400°C for 6 h, WQ [1996Che]
a = 399 c = 407
Ti-15Nb-55Al, annealed at 1400°C for 6 h, WQ [1996Che]
MSIT ®
Al–Nb–Ti
362
Pearson Symbol/ Lattice Parameters Comments /References Phase / [pm] Temperature Range Space Group/ Prototype [°C] chosen stoichiometry [1992Kat2] summarizing TiAl2, several phases [2003Sch]: < 1199 oC12 Cmmm ZrGa2
a = 1208.84 b = 394.61 c = 402.95
tP4 P4/mmm AuCu
a = 403.0 c = 395.5 tI24 I41/amd HfGa2 a = 397.0 c = 2430.9
tP32 P4/mbm Ti3Al5
MSIT®
metastable modification of TiAl2, only observed in as-cast alloys [2001Bra]; listed as TiAl2(h) (66 to 67 at.% Al, 1433-1214°C) by [1990Sch] Ti1-xAl1+x; 63 to 65 at.% Al at 1300°C, stable range 1445 - 1170°C [2001Bra]; listed as orthorhombic, Pmmm, with pseudotetragonal cell by [1990Sch] (range ~1445 - 1424°C) for Ti36Al64 at 1300°C [2001Bra]
stable structure of TiAl2 <1216°C [2001Bra]; 66 to 67 at.% Al at 1000°C [2001Bra]; listed as TiAl2(r) by [1990Sch]; < 1210°C [1996Tre, 1997Bul] [2001Bra]
a = 396.7 c = 2429.68
[1990Sch]
a = 397 c = 2430
stable between 65.5 and 66.9 at.% Al, dissolves ~5 at.% Nb at 1200°C [1998Hel]
a = 394.89 c = 412.36
Ti-4.1Nb-64.6Al ( in a + J alloy at 1200°C) (for the CuAu type subcell) [1998Hel]
a = 397.16 c = 405.92
Ti-3.8Nb-65.8Al ( in a + + J alloy at 1000°C (for the CuAu type subcell) [1998Hel]
a = 1129.3 c = 403.8
Ti3Al5, stable below 810°C [2001Bra]
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
363
Pearson Symbol/ Lattice Parameters Comments /References Phase / [pm] Temperature Range Space Group/ Prototype [°C] summarizes several phases [2003Sch]: Ti5Al11 tetragonal superstructure of stable range 1416 - 995°C [2001Bra] AuCu type 66 to 71 at.% Al at 1300°C [2001Bra] [2001Bra] (including the stoichiometry Ti2Al5); at 66 at.% Al [2001Bra] a* = 395.3 * AuCu subcell only c* = 410.4 at 71 at.% Al [2001Bra] a* = 391.8 * AuCu subcell only c* = 415.4 D023 type [V-C] tI16 I4/mmm a=398.81-392.3 68.5 to 70.9 at.% Al, 1416 - 1206°C [1990Sch] ZrAl3 c=1649.69-1653.49 69-71 at.%Al, 1450-990°C [1996Tre, 1997Bul] for 69.4 at.% Al, accepted as Ti2Al5, stable a = 393 c = 1654 between 69.4 and 71.8 at.% Al at 1200°C, dissolves ~2 at.% Nb [1998Hel] a = 390.53 “Ti2Al5”; 1416 - 990°C [1992Kat2] tP28 c = 2919.63 P4/mmm ~1215 - 985°C [1990Sch]; Ti2Al5 included in hom. region of Ti5Al11 [2001Bra] 0 x 1 [1989Jew, 1989Kal, 1990Per, (Ti1-xNbx)Al 3, J tI8 1992Pav2, 1995Zdz, 1996Che, 1998Din, I4/mmm 1998Hel, 1998Wan] TiAl3 74.2 to 75.0 at.% Al [2003Sch] TiAl3 (h) D022 ordered phase, 1387 - 735°C, < 1393 a = 384.9 c = 860.9 74.5 to 75 at.% Al at 1200°C [2001Bra] homogeneity range 74.4 - 75.3 at.% Al [1998Hel], 74.3 - 75.6 at.% Al [Mas2] [2003Kar] a = 385.3 c = 858.7 [1980Jor] NbAl3 a = 384.1 c = 860.9 < 1680 Ti-7.4Nb-73.6Al (in a + alloy) [1998Hel] a = 385.2 c = 859.9 Ti-12.0Nb-74.9Al [1998Hel] a = 384.6 c = 862.0 Ti-16Nb-72Al, annealed at 1400°C for 6 h, a = 384 WQ [1996Che] c = 865 Ti-18.8Nb-74.6Al [1998Hel] a = 384.6 c = 860.9 Ti-19.0Nb-75.1Al [1998Hel] a = 384.2 c = 861.6 Ti-21.2Nb-72.4Al (in a +J alloy) [1998Hel] a = 385.9 c = 857.6 Ti-23.2Nb-72.5Al (in a +)+J alloy) a = 385.6 [1998Hel] c = 858.6 Ti12Nb16Al72 [1991Spa] in alloy a = 386.2 c = 859.0 Ti-12Nb-63Al
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Nb–Ti
364 Phase / Temperature Range [°C] TiAl3 (l) < 950 (Ti rich) TiAl3
Nb2Al, ) < 1940
Nb3Al, < 2060
MSIT®
Pearson Symbol/ Space Group/ Prototype tI32 I4/mmm TiAl3 (l) cP4 Pm3m AuCu3 tP30 P42/mnm )CrFe
cP8 Pm3n Cr3Si
Lattice Parameters Comments /References [pm] a = 387.7 c = 3382.8
74.5 to 75 at.% Al [2001Bra]
a = 397.2
metastable, obtained at 85 at.% Al from splat cooling [2001Bra]
32 to 42 at.% Al at solidus temperatures, 32 to 35 at.% Al at 1300°C [Mas2, V-C] a = 995.2 to 986.6 [1980Jor] c = 516.8 to 518.7 a = 990.1 c = 517.0
Ti-46.9Nb-41.2Al (in a + ) + J alloy, 1200°C) [1998Hel]
a = 992.64 c = 515.54
Ti-48.9Nb-36.2Al (1000°C) [1998Hel]
a = 991.5 c = 517.3
Ti-54Nb-35.9Al (900°C) [1992Pav1]
18.6 to 25 at.% Al [Mas2] a = 519.7 to 518.0 at 19 to 25 at.% Al [1980Jor] a = 516.97
Ti-45.2Nb-24.1Al (in a +) alloy, 1000°C) [1998Hel]
a = 517.26
Ti-52.7Nb-22.2Al (in a + +) alloy, 1200°C) [1998Hel]
a = 519.0
in the Ti-65.8Nb-26.7Al ( +)) alloy annealed at 900°C for 100 h
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti Phase / Temperature Range [°C] * Ti2NbAl, O 980
Pearson Symbol/ Lattice Parameters [pm] Space Group/ Prototype oC16 Cmcm NaHg
Comments /References
[1988Ban, 1990Moz] exists in two forms, O1(h) (~980 to 900°C) and O2 (r) (below ~900°C) with different site occupancies [1992Mur2, 1995Mur1, 1995Mur2, 2002Wu] a = 608.93 0.02 Ti-25Nb-25Al, annealed at 700°C for 228 h b = 956.94 0.04 (neutron powder diffraction, Rietveld c = 466.66 0.02 refinement) [1990Moz];
a = 605 b = 961 c = 465
Landolt-Börnstein New Series IV/11A3
365
ordered, distorted Ni3Sn type for Ti2NbAl in the Ti-37.2Nb-12.2Al alloy annealed at 700°C, 26 d [1991Ben]
a = 610.6 0.3 b = 955.7 0.3 c = 463.1 0.6
in the Ti-40.1Nb-18.4Al alloy, 1150°C, WQ + 700°C/5 h [1992Tre]
a = 609.5 0.3 b = 956.9 0.3 c = 466.0 0.6
in the same alloy, 1150°C, WQ + 700°C/5 h + 400°C/5 h [1992Tre]
a = 610 10 b = 980 10 c = 470 10
in Ti-20Nb-25Al, cooled from field [1990Wey] and in thin films Ti-11Nb-24Al [1992Hsi4]
a = 608 b = 962 c = 466
in Ti-20Nb-25Al, 800°C [1989Kes]
a = 612 b = 956 c = 466
in Ti-25Nb-25Al, 800°C [1989Kes]
a = 615 b = 953 c = 466
in Ti-30Nb-25Al, 800°C [1989Kes]
a = 596 b = 986 c = 467
for O1 and O2 in Ti-15Nb-24Al [1990Mur]
a = 604 b = 971 c = 464
in Ti-11Nb-25Al [1995Lon]
MSIT ®
Al–Nb–Ti
366 Phase / Temperature Range [°C] * Ti4NbAl3, 900
´
Pearson Symbol/ Space Group/ Prototype hP6 P63/mmc Ni2In
hP2 P63/mmc Mg
´´
oP4 P2221
T
hP3 P6/mmm 7TiCr
7´´
trigonal P3ml
Lattice Parameters Comments /References [pm]
a = 458.0 0.3 c = 552.0 0.4
[1990Ben1, 1990Ben2, 1992Ben, 2000Sad] in the Ti-12.5Nb-37.5Al alloy aged at 700°C for 26 d (Ti 4NbAl3+0+(2)) [1990Ben1]
a = 455.5 c = 554.2
in situ neutron diffraction at RT for the Ti-12.9Nb-36.5Al alloy (0+2+Ti4NbAl3) [2000Sad]
a = 457.6 c = 552.4
the same alloy at 805°C [2000Sad]
a = 580 10 c = 480 10 a = 296.5 b = 492.8 c = 464.6
disordered martensite phase in Ti-Nb-25Al at Nb content < 5 at.% [1988Str] [1990Wey] metastable phase, in rapidly solidified Ti3Al-Nb alloys containing < 2 at.% Nb [1995Xu] [1990Wey] metastable phase, in Ti-5Nb-25Al, aged at 350 - 550°C [1978Zak] [1990Ben1, 1991Li, 1992Sur, 1992Voz, 1996Men, 2000Leo2, 2000Sad, 2001Sad] metastable phase [1990Ben1, 1990Sho, 1992Sur, 1994Che1, 2000Leo2, 2000Sad, 2001Sad]; T - the idealized version of the same phase [1990Ben1]
a = 455.54 0.10 Ti3Nb0.75Al2.25 at 23°C [1990Sho] c = 554.15 0.14
MSIT®
a = 460 c = 580
[1989Ben] Ti4NbAl3
a = 457.5 c = 560.4
[1991Cha] Ti-25Nb-25Al
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
367
1800
Fig. 1: Al-Nb-Ti. The TiAl3-NbAl3 section
1700
1680+-5°C
Temperature, °C
L 1600
L+ε 1500
ζ +L L+ζ +ε
ε
1400
1393°C
1300
1200
Ti 25.00 Nb 0.00 Al 75.00
10
0.00 Ti Nb 25.00 Al 75.00
20
Nb, at.%
Al Fig. 2: Al-Nb-Ti. Liquidus surface projection [1995Zdz]. Dotted lines is the limit of primary crystallization field found by [2000Leo1]
Data / Grid: at.% Axes: at.%
20
80
ε
ζ 40
60
γ α
σ
60
40
δ 80
20
β
Ti
Landolt-Börnstein New Series IV/11A3
20
40
60
80
Nb
MSIT ®
Al–Nb–Ti
368
Al Al
Data // Grid: Grid: at.% at.% Data Axes: at.% at.% Axes:
Fig. 3: Al-Nb-Ti. Calculated liquidus surface projection [1998Ser] 20 20
80 80
ε L 40 40
ζ 60 60
γ L+σ
β0
α L+α
60 60
β 0+σ
σ
α
σ β +β 0
α +β
40 40
β +σ
δ
20 20
β +δ
ββ
20 20
Ti Ti
δ
σ+δ
80 80
40 40
60 60
80 80
Al
Nb Nb
Data / Grid: at.%
Fig. 4: Al-Nb-Ti. Calculated isothermal section at 1650°C [1998Ser]
Axes: at.%
20
80
L 40
60
60
β0
L+α
L+σ
β 0+σ
40
α α +β
σ β +β 0
β +σ
80
20
β +δ
β
Ti
MSIT®
δ
σ+δ
20
40
60
80
Nb
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
369
Al Fig. 5: Al-Nb-Ti. Experimental isothermal section at 1400°C [1996Che, 1998Wan]
Data / Grid: at.% Axes: at.%
L
L+ζ+ε
20
80
ζ
ε ε +γ
40
60
γ1
γ
ε+σ
α+γ
60
40
α
γ +σ
β +α
σ
β +γ
σ+δ
80
β +σ
δ β +δ
20
β 20
Ti
40
60
80
Al Fig. 6: Al-Nb-Ti. Calculated isothermal section at 1400°C [1998Ser]
Data / Grid: at.% Axes: at.%
L
20
80
ε
ζ ε +γ
40
α +γ
60
γ ε+σ
γ +σ
60
40
α σ
β0 β +σ
80
Landolt-Börnstein New Series IV/11A3
σ+δ 20
δ β +δ
β
Ti
Nb
20
40
60
80
Nb
MSIT ®
Al–Nb–Ti
370
Al
Data / Grid: at.%
L
Fig. 7: Al-Nb-Ti. Calculated isothermal section at 1300°C [2001Sad]
Axes: at.%
L+ε 20
80
ε ζ γ +ε
40
60
γ
α +γ
ε+σ
γ +σ
60
40
α
β0
σ
β 0+σ
δ
σ+δ
80
β +δ
β
20
Ti
40
60
Ti Nb Al
Fig. 8: Al-Nb-Ti. Partial isothermal section at 1300°C [2000Kai]
20
80
40.00 0.00 60.00
Nb
Data / Grid: at.% Axes: at.%
γ 50
50
γ +α γ +β α
α +β
60
40
β
Ti Nb Al
MSIT®
70.00 0.00 30.00
10
20
Ti Nb Al
40.00 30.00 30.00
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
371
Al Fig. 9: Al-Nb-Ti. Experimental isothermal section at 1200°C [1995Zdz]
Data / Grid: at.%
L
Axes: at.%
L+ε
20
80
ζ
ε
η 40
60
γ ε +σ
γ +σ
α +γ
60
40
α
σ
β +σ
σ+δ
80
20
β +δ
β
20
Ti
δ
40
60
80
Al Fig. 10: Al-Nb-Ti. Calculated isothermal section at 1200°C [1998Ser]
Nb
Data / Grid: at.% Axes: at.%
L
L+ε 20
80
ε ζ ε +γ
40
60
γ α +γ
60
ε+σ
γ +σ
α
σ
β 0+σ
β0
40
σ +δ
80
δ 20
β +δ β
Ti
Landolt-Börnstein New Series IV/11A3
20
40
60
80
Nb
MSIT ®
Al–Nb–Ti
372
Al
Data / Grid: at.%
Fig. 11: Al-Nb-Ti. Calculated isothermal section at 1100°C [1998Ser]
Axes: at.%
20
η
80
ε
ζ
γ +ε
40
60
γ 60
γ +α
ε +σ
40
α β0
α2
σ
β 0+σ
α
σ+δ
δ
80
20
β +δ
β
20
Ti
40
60
80
Al
Nb
Data / Grid: at.%
Fig. 12: Al-Nb-Ti. Calculated isothermal section at 1100°C [2001Sad]
Axes: at.%
20
80
ζ η
ε ε+γ
40
60
γ 60
α
α 2+γ
α2 β0
80
ε+σ
γ +σ
40
σ
β 0+σ
σ +δ
20
β +δ
β
Ti
MSIT®
20
δ
40
60
80
Nb
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
373
Al
Data / Grid: at.%
Fig. 13: Al-Nb-Ti. Calculated isothermal section at 1020°C [1998Ser]
Axes: at.%
20
80
ζ η
ε γ +ε
40
60
γ 60
α 2+γ
σ β 0+σ
β0
α 2+β
α
40
β0
α2 80
ε+σ
γ +σ
δ
δ +σ
20
β +δ β 20
Ti
40
60
80
Al Fig. 14: Al-Nb-Ti. Experimental isothermal section at 1000°C mainly based on [1998Hel]
Nb
Data / Grid: at.%
L
Axes: at.%
20
80
ζ
ε
η
ε +γ
40
60
γ 60
β0 α2
ε+σ
γ +σ
γ +α2
40
σ
α 2+σ
σ+δ
α 2+β
80
δ
α
20
β +δ β
Ti
Landolt-Börnstein New Series IV/11A3
20
40
60
80
Nb
MSIT ®
Al–Nb–Ti
374
Al
Data / Grid: at.%
Fig. 15: Al-Nb-Ti. Calculated isothermal section at 990°C [2001Ser]
Axes: at.%
20
80
ε
η γ +ε
40
60
γ 60
α 2+γ
β0
40
σ
α2 α
ε +σ
γ +σ
O1
σ+δ
80
δ 20
β /β 0 20
Ti
40
60
80
Al
Nb
Data / Grid: at.%
Fig. 16: Al-Nb-Ti. Calculated isothermal section at 900°C [1998Ser]
Axes: at.%
20
80
ε ε+η
η
ε+γ
40
60
ε+γ +σ γ 60
ε +σ
γ +σ
τ
40
σ α2 80
σ+δ
O
α
δ 20
β +δ β
Ti
MSIT®
20
40
60
80
Nb
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
375
Al
Data / Grid: at.%
Fig. 17: Al-Nb-Ti. Calculated isothermal section at 800°C [1998Ser]
Axes: at.%
20
η
80
ε
η+ε γ +ε
40
60
γ +ε+σ
γ γ +τ
ε+σ
γ +σ
60
40
τ O2
α2
σ
O2+σ
σ +δ
80
α
20
O2+β
α 2+β
δ
β +δ β
20
Ti
40
60
80
Al
Nb
Data / Grid: at.%
Fig. 18: Al-Nb-Ti. Calculated isothermal section at 700°C [2001Ser]
Axes: at.%
20
80
ε η 40
60
γ +ε
γ
γ +ε+σ ε+σ
γ +τ 60
α2 80
Landolt-Börnstein New Series IV/11A3
40
τ +γ +σ O2+σ
α2+O2
O2
σ σ+δ
δ 20
δ+β
O2+β
α α+β
Ti
τ
β 20
40
60
80
Nb
MSIT ®
Al–Nb–Ti
376
Al
Data / Grid: at.%
Fig. 19: Al-Nb-Ti. Calculated partial isothermal section at 600°C [1998Ser]
Axes: at.%
20
80
ε
+ε γ+η γ+ε +ε γ+τ
40
60
τ +σ +ε
σ +ε
60
40
τ +O2+σ
O2+σ
O2
σ
σ +δ
O2+δ+σ
80
δ 20
O2+β +δ
O2+β
β 20
Ti
Fig. 20: Al-Nb-Ti. The partial isopleth along 27.5 at.% Al
40
60
80
Nb
1200
β0
α
Temperature, °C
1100
1000
α2
O1
900
O2 800
700
Ti 72.50 Nb 0.00 Al 27.50
MSIT®
10
20
Nb, at.%
30
Ti 37.50 Nb 35.00 Al 27.50
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
Fig. 21: Al-Nb-Ti. The calculated partial isopleth along 27.5 at.% Al [1998Ser]
1300
1200
1100
β
β +α
α
377
β0
α +β 0
β +α 2
α +α 2
β 0+σ
Temperature, °C
α 2+β 0 1000
α 2+β 0+σ
β 0+σ+O1
α 2+σ α2
α 2+O1+σ
900
σ+O1
α 2+O2+σ
α 2+O1
σ+O2
800
α 2+O2 700
O2+τ
α 2+O2+τ
O2+τ+σ
600
Ti 72.50 Nb 0.00 Al 27.50
10
20
Ti 37.50 Nb 35.00 Al 27.50
30
Nb, at.%
0.15
Fig. 22: Al-Nb-Ti. Thermodynamic activities of Ti and Al in the alloys (Ti0.38Al0.62)1-xNbx at 1200°C [1996Eck]
Ti
α
0.1
0.05
Al
0 0
0.05
0.1
0.15
0.2
x
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Nb–Ti
378
50
Fig. 23: Al-Nb-Ti. Partial enthalpies of mixing of Ti and Al in the alloys (Ti0.38Al0.62)1-xNbx at 1200°C [1999Eck]
30 10
Ti
∆mixH, kJ·mol-1
-10 -30 -50 -70 -90 -110
Al -130 -150 0
0.05
0.1
0.2
0.15
x
50
Fig. 24: Al-Nb-Ti. Partial entropies of mixing of Ti and Al in the alloys (Ti0.38Al0.62)1-xNbx at 1200°C [1999Eck]
Ti 30
∆mixS, J·mol-1· K-1
10
-10
-30
-50
Al
-70 0
0.05
0.1
0.15
0.2
x
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Nb–Ti
379
Thermodynamic calculation 1378
Fig. 25: Al-Nb-Ti. CCT diagram for the Ti-27.9Al-21.8Nb alloy (cooling rate in K#s-1) [2001Sad]
0 1200
0
1000
Temperature, °C
1107
0 2
0+
0Om
0O
800
0+ 2O
2+ +O2
992 915 845
+ 2+O2
0
0
0+ +O1
600
15K/s 400
200
0.5K/s 3K/s 10K/s 2K/s 1K/s 0.25K/s 5K/s
100K/s 50K/s 20K/s 10-1
100
101
102
103
Time, s
Fig. 26: Al-Nb-Ti. CCT diagram for the Ti-23.4Al-31.7Nb alloy [2001Sad]
Thermodynamic calculation 1254 +
1200
1217
0
0+
0 2
Temperature, °C
1000
0+ 2O 0O
800
0+ +O1 +O1 O1 O1+O2 O2
1017 985 940 915 895
0 0Om 600
2K/s 0.5K/s 3K/s 1K/s 0.25K/s
400
200
50K/s 4K/s 30K/s 10K/s 100K/s 20K/s 5K/s 10-1
100
101
102
103
Time, s
Landolt-Börnstein New Series IV/11A3
MSIT ®
380
Al–Nb–Zr
Aluminium – Niobium – Zirconium Lazar Rokhlin, Natalia Bochvar Literature Data In his investigations along the section ZrAl3-NbAl3 [1962Poe] found that at 660°C a significant amount (about 12.5 at.%) of Zr can be dissolved in NbAl3. Nb-rich alloys containing up to 26 at.% Zr and 30 at.% Al were investigated first by [1967Yam] using X-ray diffraction methods and metallography, later Nb rich alloys were investigated in the same way by [1974Fed1, 1974Fed2]. From the latter works a partial isothermal section results at 800°C. The work of [1967Yam] allowed him to construct partial isothermal sections of the phase diagram at 1300 and 1100°C. Detailed microscopy studies of Zr-rich quenched samples allowed [1968Tre] to construct partial isothermal sections at 1000, 900, 800 and 700°C for the Zr corner of the phase diagram. The description presented later by [1970Ali] is merely based of the on [1968Tre]. [1970Han] studied the Al-Nb-Zr system in almost the entire concentration range employing X-ray diffraction method to determine the crystal structures and lattice constants of the phases involved. The conclusions of these experiments were presented as isothermal section at 925°C which confirms the significant solubility of Zr in NbAl3 found earlier by [1962Poe]. In [1977Ale] the solubility of Zr in Nb3Al was established once more by X-ray measurements of lattice parameters, this time after annealing at 700°C. Aiming to establish the boundaries between the phase areas more precisely, [1990Per] studied the Zr corner of the phase diagram once more. These authors employed light metallography, quantitative metallography, X-ray diffractometry and electron microprobe X-ray analysis on samples with controlled oxygen and nitrogen contents. The resulting isothermal sections at 800, 771 and 730°C showed the same phase fields, but compared with [1968Tre] and [1990Per] shifted phase boundaries. [1989Sub] discussed the phase stability of NbAl3 depending on the solubility of Zr using the solubility data from [1970Han]. [1993Hub] presented a review on the Al-Nb-Zr phase diagram basing on investigations conducted by [1970Han, 1968Tre, 1974Fed1, 1974Fed2]. [1993Hub] gave the isothermal section of the phase diagram at 925°C which was constructed according to [1970Han] with addition of the phase areas in Zr corner according to [1968Tre]. The most recent investigation of the Al-Nb-Zr system were done in the Al corner by [1994Sok] using X-ray phase analysis and light metallography. The partial isothermal section at 500°C, constructed by [1994Sok], shows a significantly larger solubility of Nb in ZrAl3 and significantly smaller solubility of Zr in NbAl3 than the earlier work by [1970Han] does. This fact is difficult to explain even if the lower temperature of [1994Sok] is taken into account. In this evaluation the data on solid solubilities in the compounds by [1970Han] are preferred because they were obtained on more exact measurements of the lattice constants. Binary Systems The edge binary systems Al-Zr and Al-Nb are accepted as recently evaluated by [2003Sch] and [2003Vel], respectively. Phase relations in Nb-Zr are accepted as drawn by [1992Oka]. Solid Phases Two ternary compounds exist in the Al-Nb-Zr system according to [1970Han]. The ternary compound -1 has a homogeneity range limited by 12 to 25 at.% Nb and 46 to 54 at.% Al. The -1 homogeneity field in the isothermal section at 925°C has the shape of a deformed ellipse. [1970Han] gave for the -1 the formulae Zr5Nb2Al6-Zr 3Nb3Al7. In the assessment of [1993Hub] for the -1 compound the generalized formulae Zr5-2xNb2+xAl6+x was assumed with 0 x 1. The composition of -2 is about Zr35Nb30Al35. The crystal structure of -2 has not been described yet. MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Nb–Zr
381
Analyzing X-ray diffraction patterns [1970Han] indicates that a third ternary compound with cubic crystal structure of the CsCl type may exist in the middle part of the Zr3Al-Nb3Al section, at elevated temperatures. The homogeneity ranges of the binary compounds ZrAl2, Zr4Al3, Zr5Al3 and Zr2Al extend substantially into the ternary system, up to 19 at.% Nb. ZrAl3 dissolves up to 2 at.% Nb at 925°C. These solubility values from [1970Han] are accepted here, although [1994Sok] reports that at 500°C as much as 8 at.% Nb would be dissolved in ZrAl3. Other Al-Zr compounds, ZrAl, Zr3Al2, Zr 3Al, dissolve only insignificant amounts of Nb. The ) phase (Nb,Zr)2Al has a large homogeneity range in the binary system and extends substantially into the ternary system dissolving up to 15 at.% Zr [1970Han]. For (Nb,Zr)3Al investigations at different temperatures by [1970Han, 1974Fed1, 1977Ale] give a consistent trend for the amount of Zr that can be dissolved in this phase: about 10 at.% at 925°C, 5 at.% at 800°C and 4 at.% at 700°C. The data reported by [1967Yam] i.e. 3 at.% Zr dissolved at 1300 and 1100°C has to be taken with care. Details on crystal structure data of the solid phases are presented in Table 1. Isothermal Sections Figure 1 displays the isothermal section of the Al-Nb-Zr phase diagram at 925°C. It is constructed mainly after [1970Han] with additions of two supposed three-phase fields, Zr 3Al2+(Zr,Nb)5Al3+(Zr,Nb)4Al3, Zr3Al+(Zr,Nb)2Al+(Zr 1-x-yNbx-yAly) and (Nb,Zr)Al3+(Nb,Zr)Al2+Nb 3Al2 which should exist according to the phase rule. The boundaries of the miscibility gap in the (Nb,Zr) continuous solid solution shown at the Nb-Zr side take the miscibility gap in the binary Nb-Zr system into account. The section shows the compound (Zr,Nb)5Al3(h) at 925°C established firmly by [1970Han]. This does not contradict the binary Al-Zr phase diagram by [2003Sch] because the lower limit of existence for that phase is shown by [1970Han] at about 1000°C only tentatively. Unlike in [1970Han] the homogeneity ranges of the Al-Zr compounds are shown as line-compounds taking into consideration that they are very narrow in the binary Al-Zr system. Figures 2, 3, 4 display the partial isothermal sections of the Zr corner of the phase diagram at 800, 771 and 730°C after [1990Per]. The estimated solubility of Al in (Zr) had to be shifted to meet the binary Al-Zr after [2003Sch]. Temperature – Composition Sections Figure 5 shows vertical sections of the surface between (Zr 1-x-yNbx-yAly) solid solution and (Zr)+(Zr1-x-yNbx-yAly) phase areas. The sections correspond to the constant Al contents of 0, 3.3 and 6.7 at.%. The sections were constructed mainly after [1990Per] with some corrections to be consistent with the accepted Al-Zr binary phase diagram. Notes on Materials Properties and Applications Additions of Al and Nb lower the corrosion resistance of Zr in water at elevated temperatures and high pressure [1968Tre]. Adding Nb additive has a favorable effect on high temperature hardness and creep resistance of Zr3Al, as found by [2003Tew]. The superconductivity of the compound Nb3Al with Zr additives was studied in [1975Fed, 1975Sha, 1977Ale]. Addition of Zr to the compound decreased temperature of the superconductivity transition Tc. Miscellaneous [1985Zak] studied structural transformations during decomposition of the (Zr) based solid solution in Zr-rich alloys containing Al and Nb, and described the sequential formation of a number of metastable phases. Similarly [1999Tew] studied the structure transformations in the Zr3Al alloys containing up to 10 mass% Nb. The alloys were rapidly quenched from liquid state and annealed then. Sequence of the solid phase formations was established.
Landolt-Börnstein New Series IV/11A3
MSIT ®
382
Al–Nb–Zr
References [1962Poe]
[1967Yam]
[1968Tre]
[1970Ali] [1970Han]
[1974Fed1]
[1974Fed2]
[1975Fed]
[1975Sha]
[1977Ale]
[1985Zak]
[1989Sub]
[1990Per]
[1992Oka] [1993Bar]
MSIT®
Poetzschke, M., Schubert, K., “On the Constitution of Some Systems Homologous or Quasihomologous to T4-B3. II. The System Ti-Al, Zr-Al, Hf-Al, Mo-Al and Some Ternary Systems” (in German), Z. Metallkd., 53(8), 548-561 (1962) (Equi. Diagram, Crys. Structure, Experimental, 45) Yamamoto, A.S., “The Determination of the Niobium-Rich Region of the Ternary Phase Diagram, Niobium-Aluminium-Zirconium. Phase Equilibria of the Niobium - Tungsten Hafnium and Niobium - Tungsten - Zirconium Alloy Systems”, U.S. At. Energy Comm. Publ., 1-82 (1967) (Equi. Diagram, Experimental, Mechan. Prop., 26) Tregubov, I.A., Kudryavzev, D.L., “The Zr Corner of the Phase Diagram and Properties of Alloys of the Zr-Al-Nb System” (in Russian), in “Fiziko-Khimiya Splavov Zirkoniya”, 14-17 (1968) (Equi. Diagram, Experimental, Corrosion, 3) Alisova, S.P., Budberg, P.B., “Aluminium-Niobium-Zirconium” (in Russian), Diagrammy Sostoyaniya Met. Sistem, 14, 133-133a (1970) (Equi. Diagram, Review, 1) Hansen, R.G., Raman, A., “Alloy Chemistry of )(-U)-Related Phases. III. )-Phases with Non-Transition Elements”, Z. Metallkd., 61, 115-120 (Equi. Diagram, Crys. Structure, Experimental, #, 24) Fedorova, M.A., Burnashova, V.V., Turchinskaya, M.I., Sokolovskaya, E.M., “Phase Composition and Superconductivity in Alloys of the System Nb-Al-Ti {Zr, Hf}” (in Russian), Moskovskii Universitet, Moscow, 2137-74, (1974) (Experimental, 10) (quoted in Alisova, S.P., Budberg, P.B., Diagrammy Sostoyaniya Met. Sistem, 20, 133-134 (1974) (Equi. Diagram, Review, 1) Fedorova, M.A., Burnashova, V.V., Sokolovskaya, E.M., Kripyakevich, P.I., “Ternary Compounds in (Ti, Zr, Hf)-(Nb, Ta)-Al Systems” (in Russian), Tezisy Dokl. Vses. Konf. Kristallokhim. Intermetall. Soedin., 2nd, Lvov, 18 (1974) (Crys. Structure, 0) Fedorova, M.A., Turchinskaya, M.I., Sokolovskaya, E.M., “Influence of Group IVb Elements on the Structure and Superconductive Properties of the Intermetallic Compound Nb3Al”, Phys. Met. Metallogr., 30, 86-87 (1975), translated from Vest. Mosk. Univ., Ser. 2: Khim., 30, 238-240 (1975) (Experimental, 4) Shamrai, V.F., Postnikov, A.M., “Study og Some Ternary Solid Solutions Based on the Compound Nb3Al” (in Russian), Dokl. Akad. Nauk SSSR, 224, 1130-1133 (1975) (Experimental, 8) Alekseevskii, N.Yu., Ageev, N.V., Shamrai, V.V., “Superconductivity of Some Three-Component Solid Solutions Based on the Compound Nb3Al”, Phys. Met. Metallogr., 43(1), 29-35 (1977), translated from Fiz. Met. Metalloved., 43(1), 38-44 (1977) (Experimental, 14) Zakharova, M.I., Badaev, O.P., “Influence of Aluminium on Structure Transformations of the Solid Solution in Alloy Zr-Nb-Al, Phys. Met. Metallogr., 60(1), 188-200 (1985), translated from Fiz. Met. Metalloved., 60(1), 199-201 (1985) (Experimental, 0) Subramanian, P.R., Simmons, J.P., Mendiratta, M.G., Dimiduk, D.M., “Effect of Solutes on Phase Stability in Al3Nb”, Mat. Res. Soc. Symp. Proc., 133(3), 51-56 (1989) (Equi. Diagram, Expermental, 12) Peruzzi, A., Bolcich, J., “Experimental Determination of the Phase Relationships in Zr/2.5-8.0 at.% Nb/0-6.7 at.% Al Alloys with 750 at. ppm O and 250 at. ppm N Between 730-900°C”, J. Nucl. Mater., 174, 1-15 (1990) (Equi. Diagram, Experimental, #, 18) Okamoto, H., “Nb-Zr (Niobium-Zirconium)”, J. Phase Equilib., 13(5), 577 (1992) (Equi. Diagram, Review, 8) Barth, E.P., Sanchez, J.M., “Obersevation of a New Phase in the Niobium-Alumionium System”; Scr. Metall. Mater., 28, 1347-1352 (1993) (Crys. Structure, Equi. Diagram, Experimental, 9)
Landolt-Börnstein New Series IV/11A3
Al–Nb–Zr [1993Hub]
[1994Sok]
[1999Tew]
[2000Tew]
[2003Sch]
[2003Vel]
383
Hubert-Protopopescu, M., Lukas, H.L., Ran, Q., “Aluminium-Niobium-Zirconium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16071.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 13) Sokolovskaya, E.M., Kazakova, E.F., Podd'yakova, E.I., Portnoi, V.K., Tolmachiova, N.Yu., “Isothermal Section of Al-Nb-Zr System at 770 K” (in Russian), Vest. Mosk. Univ., Ser. 2: Khim., 35(4), 342-344 (1994) (Equi. Diagram, Experimental, 6) Tewari, R., Mukhopadhyay, P., Banerjee, S., Bendersky, L.A., “Evolution of Ordered 7 Phases in (Zr3Al)-Nb Alloys”, Acta Mater., 47(4), 1307-1323 (1999) (Crys. Structure, Experimental, 48) Tewari, R., Dey, G.K., Ravi, K., Kutty, T.R.G., Banerjee, S., “Hot Hardness and Indepentation Creep Behaviour of Zr3Al-Nb Alloys”, Trans. Indian Inst. Met., 53(3), 381-389 (2000) (Experimental, Mechan. Prop., 22) Schuster, J.C., “Al-Zr (Aluminium-Zirconium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 168) Velikanova, T., Ilyenko, S., “Al-Nb (Aluminium-Niobium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 84)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Zr1-x-yNbx-yAly)
cI2 Im3m W
(Zr) 1855 - 863
Lattice Parameters Comments/References [pm]
a = 360.90
a = 330.4
(Nb) < 2477
at 0 x 1 and 0 y 0.1 at 925°C [1970Han] at x = 0, y = 0, dissolves up to 25.9 at.% Al at 1350°C [2003Sch] at x = 1, y = 0, dissolves up to 21.5 at.% Al at 2060°C [2003Vel]
hP2 P63/mmc Mg
a = 323.16 c = 514.75
(Al) < 660.452
cF4 Fm3m Cu
a = 404.96
pure Al at 25°C [Mas2]
Zr3Al < 1019
cP4 Pm3m Cu3Au
a = 437.2 0.3
[2003Sch], dissolves small amount of Nb [1970Han]
(Zr) < 863
Landolt-Börnstein New Series IV/11A3
pure Zr at 25°C [Mas2] dissolves up to 8.3 at.% Al at 910°C [2003Sch]
MSIT ®
Al–Nb–Zr
384 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Zr,Nb)2Al < 1215
hP6 P63/mmc Ni2In
(Zr,Nb)5Al3(h) 1400 - ?
tI32 I4/mcm W5Si3
Lattice Parameters Comments/References [pm]
a = 461.0 c = 591.3 a = 489.39 0.05 c = 592.83 0.05 a = 489.4 c = 592.8
a = 1087 c = 529.6 a = 1104.4 c = 539.1 a = 1105 c = 539.6
dissolves up to 19 at.% Nb at 925°C [1970Han], for Zr59Nb20 Al21 [1970Han] for Zr2Al [2003Sch] for Zr2Al [1970Han] dissolves about 16 at.% Nb at 925°C [1970Han], for Zr50Nb17 Al33 [1970Han] for Zr5Al3 (h) [2003Sch] for Zr3Al5 (h) [1970Han]
Zr5Al3(r) ?
hP16 P63/mcm Mn5Si3
a = 817.4 c = 569.8
[2003Sch]
Zr3Al2 < 1480
tP20 P42/mnm Zr3Al2
a = 763.0 0.1 c = 699.8 0.1
[2003Sch], dissolves small amount of Nb [1970Han]
(Zr,Nb)4Al3
hP7 P6/mmm Zr4Al3
Zr4Al3 1030
a = 536.8 c = 533.3 a = 543.3 0.5 c = 539.0 0.5
dissolves about 16 at.% Nb at 925°C [1970Han], for Zr50Nb10 Al40 [1970Han] for Zr4Al3 [2003Sch, 1970Han]
Zr5Al4 1550 - ~1000
hP18 P63/mcm Ti5Ga4
a = 844.8 c = 580.5
[2003Sch]
ZrAl < 1275 25
oC8 Cmcm CrB
a = 335.9 0.1 b = 1088.7 0.3 c = 427.4 0.1
[2003Sch], dissolves small amount of Nb [1970Has]
Zr2Al3 < 1590
oF40 Fdd2 Zr2Al3
a = 960.1 0.2 b = 1390.6 0.2 c = 557.4 0.02
[2003Sch], dissolves up to 1 at.% Nb at 500°C [1994Sok]
(Zr,Nb)Al2 < 1660
hP12 P63/mmc MgZn2
MSIT®
a = 525.4 c = 869.0 a = 528.24 0.05 c = 874.82 0.05 a = 528.2 c = 874.8
dissolves about 15 at.% Nb at 925°C [1970Han], for Zr25Nb15 Al60 [1970Han] for ZrAl2 [2003Sch] for ZrAl2 [1970Han]
Landolt-Börnstein New Series IV/11A3
Al–Nb–Zr Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Zr,Nb)Al3 < 1580
tI16 I4/mmm ZrAl3
(Nb,Zr)3Al
cP8 Pm3m Cr3Si
Nb3Al < 2060 Nb3Al2 1590
tP20 P42/mnm Al2Zr3
(Nb,Zr)2Al () phase) < 1940
tP30 P42/mnm )CrFe
(Nb1-xZr x)Al3
NbAl3 < 1680
Landolt-Börnstein New Series IV/11A3
tI8 I4/mmm TiAl3
385
Lattice Parameters Comments/References [pm]
a = 400.5 c = 1727 a = 399.93 0.05 c = 1728.3 0.2 a = 401.0 c = 1731.5
dissolves about 2 at.% Nb at 925°C [1993Hub], for Zr20Nb5Al75 [1970Han] for ZrAl3 [2003Sch] for ZrAl3 [1962Poe]
a = 519.7
dissolves about 10 at.% Zr at 925°C [1970Han] for Zr2.6Nb72.4Al25 [1977Ale]
a = 518.6 a = 518.7
for Nb3Al [2003Vel] for Nb3Al [1970Han]
a = 707 8 c/a 0.05
[1993Bar] 42.4 at.% Al, equilibria to be checked
a = 995.2 c = 517.4
dissolves about 15 at.% Zr at 925°C [1970Han], for Zr13.5Nb53.5 Al33 [1970Han]
a = 988.7 c = 516.2
for Zr7Nb53Al40 [1970Han]
a = 994.3 c = 518.6
for Nb2Al [1970Han]
a = 387.9 c = 877.1
at 0 x 0.68 at 925°C [1970Han], at x = 0.5 [1970Han]
a = 389 c = 876
at x = 0.5, annealed at 660°C, two-phase alloy [1962Poe]
a = 387 c = 874
at x = 0.24, annealed at 660°C, two-phase alloy [1962Poe]
a = 384.1 c = 860.9
at x = 0 [2003Vel]
a = 384.5 c = 860.1
at x = 0 [1970Han]
a = 384 c = 858
at x = 0 [1962Poe]
MSIT ®
Al–Nb–Zr
386 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
* -1, Zr5-2xNb2+xAl6+x
hR13 R3m W6Fe7
* -2, Zr35Nb30Al35
-
* -3
cP2 Pm3m CsCl
Lattice Parameters Comments/References [pm]
a = 522.7 c = 2830
at 0 x 1 at x = 1 [1970Han]
a = 528.2 c = 2858
at x = 0.5 [1970Han]
a = 529.6 c = 2873
at x = 0.1 [1970Han]
-
[1970Han] Assumed to be stable in the as cast condition, between 25 to 38 at.% Nb along section Zr3Al-Nb 3Al [1970Has]
Al Fig. 1: Al-Nb-Zr. Isothermal section at 925°C
Data / Grid: at.%
(Al)
Axes: at.%
(Zr,Nb)Al3 20
80
(Nb,Zr)Al3
(Zr,Nb)Al2 Zr2Al3 40
60
ZrAl (Zr,Nb)4Al3 Zr3Al2 60 (Zr,Nb)5Al3 (Zr,Nb)2Al
τ1
Nb3Al2 40
τ2
(Nb,Zr)2Al
Zr3Al
(Nb,Zr)3Al
80
20
(β Zr1-x-yNbx-yAly)
Zr MSIT®
20
40
60
80
Nb
Landolt-Börnstein New Series IV/11A3
Al–Nb–Zr
387
Zr Nb Al
Fig. 2: Al-Nb-Zr. Partial isothermal section at 800°C
86.00 0.00 14.00
Data / Grid: at.% Axes: at.%
90
10
(α Zr)+Zr3Al
(α Zr)+(β Zr1-x-yNbx-yAly)+Zr3Al
(αZr) (α Zr)+(β Zr1-x-yNbx-yAly) (β Zr1-x-yNbx-yAly) 10
Zr Zr Nb Al
Fig. 3: Al-Nb-Zr. Partial isothermal section at 771°C
86.00 0.00 14.00
Zr Nb Al
86.00 14.00 0.00
Zr Nb Al
86.00 14.00 0.00
Data / Grid: at.% Axes: at.%
90
10
(α Zr)+(β Zr1-x-yNbx-yAly) (αZr)+(β Zr1-x-yNbx-yAly)+Zr3Al
(α Zr) (αZr)+(β Zr1-x-yNbx-yAly) (β Zr1-x-yNbx-yAly)
Zr
Landolt-Börnstein New Series IV/11A3
10
MSIT ®
Al–Nb–Zr
388
Zr Nb Al
Fig. 4: Al-Nb-Zr. Partial isothermal section at 730°C
86.00 0.00 14.00
Data / Grid: at.% Axes: at.%
90
10
(αZr)+Zr3Al
(α Zr)+(β Zr1-x-yNbx-yAly)+Zr3Al (αZr) (αZr)+(β Zr1-x-yNbx-yAly)
(β Zr1-x-yNbx-yAly) 10
Zr
Zr Nb Al
86.00 14.00 0.00
1000
Fig. 5: Al-Nb-Zr. Partial vertical sections at 0 at.% Al (1), 3.3.at.% Al (2), 6.7 at.% Al (3) Temperature, °C
900
(βZr1-x-yNbx-yAly) 3
800
(βZr1-x-yNbx-yAl2y) 3 (αZr) + (βZr1-x-yNbx-yAly) 2 (αZr)+(βZr1-x-yNbx-yAly) 1
700
1
600 0
4
12
8
16
20
Nb, at.%
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ru
389
Aluminium – Nickel – Ruthenium Benjamin Grushko Literature Data The previous assessment of this system was made in the MSIT Evaluation Program [1993Tre] based on the data from [1980Tsu, 1985Sok] for 550ºC, partial isothermal sections at 1000 and 1250°C in the Al-poor region and on a liquidus projection published by [1985Cha1, 1985Cha2, 1986Cha]. No ternary phases and no significant solubility of the third element in the binary compounds were revealed apart from about 5 at.% Ru that solve in Ni2Al3 [1980Tsu]. Even NiAl and RuAl, both of the CsCl type structure and alike lattice parameters, were found to form a wide miscibility gap at 1250°C [1986Cha] which, according to [1980Tsu], becomes significantly wider at 550°C. However, in [1985Sok] and then in [1998Hor] a continuous range of solid solutions between these phases was concluded. A number of the ternary alloys were investigated in [1997Hor1, 1997Hor2]. But the data obtained are of limited use in the determination of isothermal sections and the homogeneity ranges for the phases because the experimental technique employed resulted in as-cast, thus not equilibrated samples. In [1997Poh, 2000Sun1, 2000Sun2, 2001Sun, 2002Sun, 2002Hir] the formation of the several quasicrystalline phases was observed in the high-Al range of the Al-Ni-Ru system. Only for the decagonal phase the stability was confirmed by [2003Mi1, 2003Mi2, 2004Mi1] with a periodicity of 1.6 nm (D4 phase) together with four stable crystalline phases which were determined in the Al-rich region. The phase equilibria in the temperature range of 700 to 1100°C can be described by partial isothermal sections built on [2003Mi2]. Additional data are reported by [1997Hor1] from investigations on alloys of low Al contents. Binary Systems The description of the Al-Ni phase equilibria has been accepted from [2003Sal]. According to the recent work [2004Mi2] the Al-Ru system contains six intermediate phases: RuAl6, Ru4Al13, RuAl2, Ru2Al5, Ru 2Al3 and RuAl, that have been previously reported in [1996Bon1, 1996Bon2]. RuAl12 reported in [1963Obr] was not confirmed. Different crystal structures were associated with RuAl2 and Ru2Al3. The data accepted in Table 1 originate from [2004Mi2]. Apart from the stable phases a metastable icosahedral phase (I) and decagonal D1 were reported in this system [1990Wan]. Only RuAl exhibits a significant compositional range. By heating up to 2100°C the melting point of RuAl was yet not reached [2004Mi2]. Ru dissolves up to 4 at.% Al [Mas2]. The Ni-Ru system does not contain stable intermediate phases [L-B]. At 1550°C Ni dissolves up to 34.5 at.% Ru while Ru up to 50 at.% Ni. Variations of the lattice parameters of these solid solutions with the compositions are compiled in [L-B]. A metastable phase was reported in [1979Var] in the range of 30 to 40 at.% Ru. Solid Phases At 1500-1600°C the congruent RuAl and NiAl phases form a continuous range of solid solution (Ru,Ni)Al [1998Hor] which naturally separates the high-Al and low-Al ranges. Considering the very high melting temperatures [2001Liu] the constitution of the corresponding alloys in equilibrium is unclear even at 1000°C. At 1100°C the Ru4Al13, RuAl6, and RuAl2 phases dissolve up to 7.0, 0.5 and 0.7 at.% Ni, while Ni2Al3 and NiAl3 dissolve about 0.7 and < 0.5 at.% Ru, respectively [2003Mi2]. The ternary m phase ((Ru,Ni)2Al9) is isostructural to Co2Al9 and forms at almost constant 82 at.% Al between 4.5 and 7.0 at.% Ru, the hexagonal H-phase is located in a small range around Ru8.5Ni16.0Al75.5, and the decagonal D4 phase forms in the vicinity of the H phase around Ru11Ni16Al73 [2003Mi2]. The orthorhombic O, (Ru,Ni)4Al13 phase (O, Co4Al13 type) is observed in a small compositional range around
Landolt-Börnstein New Series IV/11A3
MSIT ®
390
Al–Ni–Ru
Ru16.0 Ni8.0Al76.0 while the C phase (C, Rh2Al5 type) forms in a compositional range of about 9.0-12.0 at.% Ni and 72.0-73.0 at.% Al [2003Mi2]. The diffraction pattern of the D4 phase is characterized in [2000Sun1, 2000Sun2, 2003Man]. The experiments made by [2004Mi1] point to a compositional dependence of the D range on temperature from the melting to at least 700°C. The D1 phase was characterized by [2000Sun1, 2003Man] and an I phase by [2001Sun]. These phases were not observed in annealed samples [2000Sun1, 2001Sun, 2003Mi1, 2003Mi2] and are considered to be metastable. The crystallographic data of the ternary phases and their stability against temperature are listed in Table 1. It was argued in [2004Mi1] that the stable ternary D4 phase is an extension of a metastable Al-Ru D4 phase. Pseudobinary Sections The RuAl-NiAl part of the phase diagram is suggested to be a pseudobinary section, between congruent RuAl and NiAl, at least at temperatures >1500°C. Invariant Equilibria A reaction scheme of the Al-Ni-Ru system has been presented in [1993Tre]. However, it had to be revised for the high-Al range because recently [2003Mi1, 2003Mi2] reported the formation of five ternary phases by incongruent reactions at the temperatures given in Table 1. The type of the reactions was not established. The presently accepted reaction scheme for the low-Al part is presented in Fig. 1. Liquidus Surface A tentative liquidus projection of the Al-Ni-Ru system was proposed in the previous evaluation published in [1993Tre]. Considering the ternary phases which now are supposed to form from the liquid, the high Al part of the liquidus surface can not be accepted anymore. For the Al-poor part of the system [1997Hor1] published an alternative version of a liquidus projection. However, considering the solubility data of Ni in (Ru), Ru in (Ni), and Al in (Ni) and (Ru), the indicated location of the eutectic is improbable because the liquid phase which takes part in the reaction LNi3Al+(Ni)+(Ru) does not lie inside the respective tie triangle. For the same reason, the liquidus projection for the Al-rich part of the diagram constructed in [2000Hoh] by using data obtained from the as-cast samples can not be accepted. Measured phase equilibria at the subsolidus temperatures are needed to support a decision. The Al-poor region of the liquidus projection still applies as described by [1993Tre], see Fig. 2. Isothermal Sections In the section of the Al-rich part of the Al-Ni-Ru system at 1600°C given by [1997Hor2] some of the samples investigated were already liquid. The partial isothermal section at 1250°C is presented in Fig. 3 according to [1993Tre], and the 1100°C isothermal section in Fig. 4 according to [2003Mi2]. The 1000°C isothermal section (Fig. 5) is combined from the data in [2003Mi2] for high Al compositions and from [1993Tre] for low-Al compositions considering the continuous b range. The isothermal sections at 900°C (Fig. 6), 800°C (Fig. 7) and 700°C (Fig. 8) are based on [2003Mi2], however the compositional limits of the investigated ranges are shifted to higher Al-concentrations in order to reach at lower temperatures equilibrium. Notes on Materials Properties and Applications Potential applications for (Ru,Ni)Al at high temperatures may arise from its high strength, reasonable high temperature toughness and good oxidation resistance as mentioned by [1997Wol1, 1997Wol2]. Synthesized by mechanical alloying, the (Ru,Ni)Al alloys with grain sizes of 20-40 nm show a high stability at elevated temperature [2001Liu].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ru
391
Miscellaneous The difficult mixing of Al and Ru in high-Ru binary and ternary alloys with small fractions of Ni requires repeated melting and prolonged annealing in order to obtain homogeneous materials. References [1963Obr]
[1965Eds] [1966Eds] [1968Eds] [1979Var] [1980Tsu]
[1985Cha1]
[1985Cha2]
[1985Sok]
[1986Cha] [1990Wan]
[1993Fle] [1993Tre]
[1996Bon1] [1996Bon2] [1997Hor1]
[1997Hor2]
[1997Poh]
Landolt-Börnstein New Series IV/11A3
Obrowski, W., “On the Alloys of Ruthenium with Boron, Berylium and Aluminium” (in German), Metallwissenschaft und Technik (Berlin), 17(2), 108-112 (1963) (Equi. Diagram, Crys. Structure, Experimental, 10) Edshammar, L-E., “The Crystal Structure of Ru4Al13”, Acta Chem. Scand., 19, 2124-2130 (1965) (Crys. Structure, Experimental, 5) Edshammar, L-E., “An X-Ray Investigation of Ruthenium-Aluminium Alloys”, Acta Chem. Scand., 20, 427-431 (1966) (Crys. Structure, Experimental, 3) Edshammar, L-E., “The Crystal Structure of RuAl6”, Acta Chem. Scand., 22, 2374-2400 (1968) (Crys. Structure, Experimental, 8) Varich, N.I., Petrunina, A.N., Russ. Metall., 90-91 (1979) (Crys. Structure, Experimental, 3) Tsurikov, V.F., Sokolovskaya, G.M., Kazakova, E.F., “Interaction of Nickel and Aluminium with Ruthenium” (in Russian), Vestn. Mosk. Univ., Khim., 21(5), 512-514 (1980) (Equi. Diagram, Experimental, 6) Chakravorty, S., West, D.R.F., “Phase Equilibria Between NiAl and RuAl in the Ni-Al-Ru System”, Scr. Metall., 19(11), 1355-1360 (1985) (Equi. Diagram, Crys. Structure, Experimental, 10) Chakravorty, S., Hashim, H., West, D.R.F., “The Ni3Al-Ni3Cr-Ni3Ru Section of the Ni-Cr-Al-Ru System”, J. Less-Common Met., 20, 2313-2322 (1985) (Equi. Diagram, Crys. Structure, Experimental, 31) Sokolovskaya, E.M., Tsurikov, V.F., Orybenkov, S.B., Makanov, U.M., “Phase Diagrams in Some Systems Containing Aluminum” (in Russian), Stable and Metastable Phase Equilibria in Metallic Systems, 86(6:72), 79-83 (1985) (Equi. Diagram, 11) Chakravorty, S., West, D.R.F., “The Constitution of the Ni-Al-Ru System”, J. Mater. Sci., 21(8), 2721-2730 (1986) (Equi. Diagram, Crys. Structure, Experimental, #, *, 23) Wang, Z.M., Gao, Y.Q., Kuo, K.H., “Quasicrystals of Rapidly Solidified Alloys of Al-Pt Group Metals – II. Quasicrystals in Rapidly Solidified of Al-Ru and Al-Os Alloys”. J. Less-Common Met., 163 (1990) (Experimental, Crys. Structure) Fleischer, R.L., “Boron and off-Stoichiometry Effects on the Strength and Quality of AlRu”, Metall. Trans. A, 24A, 227-230, (1993) (Experimental) Tretyachenko, L., Sheftel, E., Ibe, G., Grieb, B., Rogl, P., “Aluminum-Nickel-Ruthenium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16434.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 9) Bonniface, T.D., Cornish, L.A., “An Investigation of the High Aluminium end of the Al-Ru Phase Diagram”, J. Alloys Compd., 233, 241-245 (1996) (Experimental, 11) Bonniface, T.D., Cornish, L.A., “An Investigation of the Al-Ru Phase Diagram above 25 at.% Al”, J. Alloys Compd., 234, 275-279 (1996) (Equi. Diagram, Experimental, 15) Horner, L.J., Cornish, L.A., Witcomb, M.J., “A Study of the Al-Ni-Ru Ternary System Below 50 at.% Aluminium”, J. Alloys Compd., 256, 213-220 (1997) (Equi. Diagram, Experimental, 14) Horner, L.J., Cornish, L.A., Witcomb, M.J, “Constitution of the Al-Ni-Ru Ternary System Above 50 at.% Aluminium”, J. Alloys Compd., 256, 221-227 (1997) (Equi. Diagram, Experimental, 14) Pohla, C., Ryder, P.L., “Crystalline and Quasicrystalline Phases in Rapidly Solidified Al-Ni Alloys”, Acta Mater., 45, 2155-2166 (1997) (Crys. Structure, Experimental, 48)
MSIT ®
392 [1997Wol1] [1997Wol2]
[1998Hor]
[2000Hoh]
[2000Sun1]
[2000Sun2]
[2001Liu]
[2001Sun]
[2002Hir]
[2002Sun]
[2003Man]
[2003Mi1] [2003Mi2]
[2003Sal]
[2004Mi1]
[2004Mi2]
MSIT®
Al–Ni–Ru Wolff, I.M., “Towards a Better Understanding of Ruthenium Aluminide”, JOM, (1), 34-39 (1997) (Review, 58) Wolff, I.M., Sauthoff, G., Cornish, L.A., Steyn, H. de V., Coetzee, R., “Structure-Property-Application Relationships in Ruthenium Aluminide RuAl”, Structural Intermetallics, 1997, The Minerals, Metals & Materials Society, 815-823 (1997) (Crys. Structure, Electr. Prop., Experimental, Mechan. Prop., 41) Horner, I.J., Hall, N., Cornish, L.A., Witcomb, M.J., Cortie, M.B., Boniface, T.D., “An Investigation of the B2 Phase Between AlRu and AlNi in the Al-Ni-Ru Ternary System”, J. Alloys Compd., 264, 173-179 (1998) (Equi. Diagram, Experimental, 23) Hohls, J., Cornish, L.A., Ellis, P., Witcomb, M.J., “Solidification Phases and Liquidus Surface of the Al-Ni-Ru System Above 50 at.% Aluminium”, J. Alloys Compd., 308, 205-215 (2000) (Crys. Structure, Equi. Diagram, Experimental, 22) Sun, W., Hiraga, K., “Formation and Structures of Decagonal Quasi-Crystals in the Al-Ni-Ru System”, Mater. Sci. Eng. A, 294-296, 147-151 (2000) (Crys. Structure, Experimental, 12) Sun, W., Hiraga, K., “A New Highly Ordered Al-Ni-Ru Decagonal Quasicrystal with 1.6 nm Periodicity”, Philos. Mag. Lett., 80(3), 157-164 (2000) (Crys. Structure, Experimental, 29) Liu, K.W., Muecklich, F., Pitschke, W., Birringer, R., Wetzig, K., “Formation of Nanocrystalline B2-Structured (Ru,Ni)Al in the Ternary Ru-Al-Ni System by Mechanical Alloying and its Thermal Stability”, Mater. Sci. Eng. A, 313, 187-197 (2001) (Crys. Structure, Experimental, 30) Sun, W., Hiraga, K., “Structural Study of a Superlattice Al-Ni-Ru Decagonal Quasicrystal Using High-Resolution Electron Microscopy and a High-Angle Annual Dark-Field Technique”, Philos. Mag. Lett., 81(3), 187-195 (2001) (Crys. Structure, Experimental, 17) Hiraga, K., “The Structure of Quasicrystals Studied by Atomic-Scale Observations of Transmission Electron Microscopy”, Adv. Imag. Electr. Phys., 122, 1-86 (2002) (Review, Crys. Structure, 99) Sun, W., Hiraga, K., “A New Icosahedral Quasicrystal Coexisting with Decagonal Quasicrystals in the Al-Ni-Ru System”, J. Alloys Compd., 347, 110-114 (2002) (Crys. Structure, Experimental, 19) Mandal, P., Hashimoto, T., Suzuki, K, Hosono, K., Kamimura, Y., Edagawa, K., “Formation of Decagonal and Approximant Phases in the Al-Ni-Ru System”. Philos. Mag. Lett., 85, 315-323 (2003) (Experimental, Crys. Structure, 24) Mi, S., Grushko, B., Dong, C., Urban, K., “Ternary Al-Ni-Ru Phases”, J. Alloys Compd., 351, L1-L5 (2003) (Equi. Diagram, Crys. Structure, Experimental, 10) Mi, S., Grushko, B., Dong, C., Urban, K., “Isothermal Sections of the Al-Rich Part of the Al-Ni-Ru Phase Diagram”, J. Alloys Compd., 359, 193-197 (2003) (Equi. Diagram, Experimental, 21) 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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 25) Mi, S., Grushko, B., Dong, C., Urban, K., “Phase Equilibrium in the Vicinity of the Al-Ni-Ru Decagonal Phase”, J. Non-Cryst. Solids., 334-335, 214-217 (2004) (Equi. Diagram, Experimental, 15) Mi, S., Balanetskyy, S., Grushko, B., “A Study of the Al-Rich Part of the Al-Ru Alloy System”, Intermetallics, 11(7), 643-649 (2004) (Equi. Diagram, Crys. Structure, Experimental, 18)
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ru
393
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 (Ru) < 2334 (Ni) < 1455
RuAl6 < 734 Ru4Al13 < 1420
Ru2Al5 1340 - 1492 RuAl2 < 1805 Ru2Al3 < 1675 , (Ru1-xNix)yAl1-y
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu hP2 P63/mmc Mg cF4 Fm3m Cu
oC28 Cmcm Al6Mn mC102 C2/m Fe4Al13 oC* Cmcm Fe5Al2 oF24 Fddd TiS2 tI10 I4/mmm Os2Al3 cP2 Pm3m CsCl
NiAl < 1651
Landolt-Börnstein New Series IV/11A3
a = 404.96
dissolves 0.01 at.% Ni at 639.9°C [2003Sal]
a = 270.53 c = 428.20
at 25°C [V-C] dissolves 4 at.% Al at 1920°C [Mas2] dissolves 50 at.% Ni at ~1500°C [L-B] [Mas2] dissolves 20.2 at.% Al at 1385°C [2003Sal] dissolves 34.5 at.% Ru at ~1500°C [L-B] [1968Eds] [2004Mi2] dissolves < 0.5 at.% Ni [2003Mi2] [1965Eds], dissolves 7 at.% Ni at 1000°C [2003Mi2]
a = 352.40
a = 748.8 b = 655.6 c = 896.1 a = 1586.2 b = 818.8 c = 1273.6 = 107.88° a = 780 b = 660 c = 420 a = 801.2 b = 471.7 c = 878.5 a = 307.9 c = 1433
a = 293.9 a = 303 a = 293 a = 299.16
RuAl < at least 2100
Ni2Al
Lattice Parameters Comments/References [pm]
hP3 P3m1 CdI2
a = 299.16 a = 287 a = 288.72 0.02 a = 287.98 0.02 a = 288.64 a = 407 c = 499
[2004Mi2]
[1966Eds] [2004Mi2] [1966Eds] [2004Mi2] 0 x 1 [1985Sok] in Ru29 Ni24 Al47 annealed at 1600°C [1998Hor] [V-C] in 56 at.% Ru [1986Cha] [2001Liu] 42 to 69.2 at.% Ni [2003Sal] [1993Fle] at 63 at.% Ni at 50 at.% Ni at 54 at.% Ni [2001Liu] Metastable [2003Sal]
MSIT ®
Al–Ni–Ru
394 Phase/ Temperature Range [°C] Ni2Al < 1372
Pearson Symbol/ Space Group/ Prototype cP4 Pm3m AuCu3
Ni5Al3 < 723
oC16 Cmmm Pt5Ga3
Ni3Al4 < 702
cI112 Ia3d Ni3Ga4 hP5 P3m1 Ni2Al3 oP16 Pnma NiAl3 tP4 P4/mmm AuCu
Ni2Al3 < 1138 NiAl3 < 856 NixAl1-x
m**
Ni2Al9
mP22 P21/a Co2Al9
D1
P105mc or P105/mmc
D4 , (NiRu)
MSIT®
t** -
Lattice Parameters Comments/References [pm] a = 358.9 a = 356.32 a = 357.92 a = 356.77 a = 753 b = 661 c = 376 a = 1140.8
at 63 at.% Ni disordered ordered 73 to 76 at.% Ni [2003Sal] 63 to 68 at.% Ni at 63 at.% Ni [2003Sal]
[2003Sal]
a = 402.8 c = 489.1
36.8 to 40.5 at.% Ni [2003Sal]
a = 661.3 b = 736.7 c = 481.1
[2003Sal]
a = 379.5 c = 325.6 a = 379.5 c = 325.6 a = 375.1 c = 330.7 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 = 418 b = 271 c = 1448 = 93.4° a = 868.5 b = 623.2 c = 618.5 = 96.50º a = 373.3 c = 407.3
0.60 < x < 0.68 Martensite, metastable [2003Sal] at 62.5 at.% Ni at 63.5 at.% Ni at 66.0 at.% Ni at 64 at.% Ni at 65 at.% Ni [2003Sal] [2003Sal]
Metastable [2003Sal]
Decagonal in 24-30 at.% Ni [2003Sal]
a=? c 1600
Decagonal [2003Sal], contained some Si
a = 451.1 c = 362.0
Metastable [1979Var]
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ru Phase/ Temperature Range [°C] * (Ru,Ni)2Al9 < 783
Pearson Symbol/ Space Group/ Prototype mP22 P21/a Co2Al9
* O, Ru4Al13 < at least 1000
oP102 Pmn21 O-Co4Al13
*C, (Ru,Ni)2Al5 1233 to at least 1100 *H < 930
Pm3 or P23 C-Rh2Al5 h**
* D4 < 1057
395
Lattice Parameters Comments/References [pm] a = 863.6 b = 633.3 c = 627.3 = 95.12º a = 1496.0 b = 825.3 c = 1266.8 a = 767.4
in Ru5,6Ni12Al82 [2003Mi1, 2003Mi2]
Around Ru16.0Ni8Al76 [2003Mi1, 2003Mi2] High-temperature phase 9.0-12.0 at.% Ni and 72.0-73.0 at.% Al [2003Mi1, 2003Mi2] [2003Mi1, 2003Mi2]
a = 1213.2 c = 2702.0
* D1
Decagonal [2002Sun, 2003Mi1, 2003Mi2] [2003Man] Decagonal Metastable (?) [2000Sun1]
*I
Icosahedral Metastable (?) [2002Sun]
Ni-Ru
a = 248 c = 1670
Al-Ni
Al-Ni-Ru
Al-Ru
L+(Ru)+β
1550 p2 l + (Ru) (Ni) 1372 p4 l + (Ni) Ni3Al 1369 e2 l Ni3Al + β
1920 e1 l β + (Ru)
L+(Ru)+(Ni)
T3<1369 L + Ni3Al β + (Ni)
U2
(Ni) + β + Ni3Al L+β+(Ni) 1250
L β + (Ru) + (Ni)
E1
β + (Ru) + (Ni) Fig. 1: Al-Ni-Ru. Reaction scheme of the Al-poor part [1993Tre]
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ru
396
Al
Data / Grid: at.%
Fig. 2: Al-Ni-Ru. Liquidus projection of the Al-poor part
Axes: at.%
20
80
40
60
60
40
β
e1
Ni 3Al U2
80
e2 p4 20
E1
(Ni)
(Ru) 20
Ru
40
60
p2
Al
80
Ni
Data / Grid: at.%
Fig. 3: Al-Ni-Ru. Partial isothermal section at 1250°C
Axes: at.%
20
80
40
60
β 60
40
(Ru)+β
Ni3Al+(Ni)+β Ni3Al
β +(Ni)
(Ni)+(Ru)+β
80
20
(Ni) (Ru)+(Ni)
(Ru)
Ru
MSIT®
20
40
60
80
Ni
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ru
397
Al
Data / Grid: at.%
Fig. 4: Al-Ni-Ru. Partial isothermal section at 1100°C
Axes: at.%
10
90
20
80
M
L
30
RuAl2
70
C
Ni2Al3
40
60
β0 Ru Ni Al
10
45.00 0.00 55.00
20
30
40
Al
Ru Ni Al
0.00 45.00 55.00
Data / Grid: at.%
Fig. 5: Al-Ni-Ru. Partial isothermal section at 1000°C
Axes: at.%
L
20
M
80
D O Ni2Al3
40
60
β 60
40
β +Ni3Al (N i)+ Ni 3 Al +β
(Ru)+β 80
(Ru)+(Ni)+β
Ni3Al 20
(Ni) (Ru)+(Ni)
Ru
Landolt-Börnstein New Series IV/11A3
(Ru)
20
40
60
80
Ni
MSIT ®
Al–Ni–Ru
398
Al
Data / Grid: at.%
Fig. 6: Al-Ni-Ru. Partial isothermal section at 900ºC
Axes: at.%
10
90
L
20
80
M
O
D
H
30
70
Ni2Al3 Ru Ni Al
10
40.00 0.00 60.00
20
30
Al Fig. 7: Al-Ni-Ru. Partial isothermal section at 800ºC
Ru Ni Al
0.00 40.00 60.00
Data / Grid: at.% Axes: at.%
L
10
90
20
80
M
O
H
NiAl3
D 30
70
Ni2Al3 Ru Ni Al
MSIT®
40.00 0.00 60.00
10
20
30
Ru Ni Al
0.00 40.00 60.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ru
399
Al Fig. 8: Al-Ni-Ru. Partial isothermal section at 700ºC
Data / Grid: at.% Axes: at.%
L
10
90
RuAl6 m 20
80
NiAl3
M
Ru Ni Al
Landolt-Börnstein New Series IV/11A3
30.00 0.00 70.00
10
20
Ru Ni Al
0.00 30.00 70.00
MSIT ®
400
Al–Ni–Si
Aluminium – Nickel – Silicon Olga Fabrichnaya, Georg Beuers, Christian Bätzner and Hans Leo Lukas Literature Data The Al-rich corner was studied several times using thermal and microscopic analyses [1926His, 1930Ota, 1934Fus, 1939Wei, 1942Phi]. The phase relations at Ni contents up to 33.3 at.% have been recently studied by [2002Ric, 2003Ric1]. A ternary eutectic exists between (Al), NiAl3 and Si. The values given for temperature and concentration of the eutectic melt are between 560 and 568°C, 3.0 and 5.2 mass% (1.4 and 2.5 at.%) Ni, 11.0 and 11.8 mass% (10.8 and 11.7 at.%) Si. According to the measurements of [2003Ric1] the temperature of ternary eutectic is 565°C and the composition of the liquid is 2 at.% Ni and 11 at.% Si. Isopleths are reported for 2 [1942Phi], 6 and 14 mass% Si [1930Ota] and for 2 [1939Wei, 1942Phi], 2.5 [1930Ota], 3 [1990Kuz], 4 [1930Ota, 1939Wei], 5 [1959Phi], 7.5 and 12.5 [1930Ota] mass% Ni. The isopleths agree well though only [1939Wei] gives a Si solubility in (Al) in agreement with the binary Al-Si system. [1930Ota] ignores that totally and [1942Phi] gives a much lower value. Recently [2002Ric, 2004Ric] experimentally obtained isoplethal sections for 10, 20, 30, 33.3, 40, 45, 50, 55, 60 and 66.7 at.% Ni. [1934Fus] gave the Al-rich liquidus surface indicating two more invariant reactions. However the ternary phase Ni3(Al1-xSix)7 was not taken into account by [1934Fus]. Recently new data on the liquidus surface were reported by [2003Ric1] at compositions up to 33.3 at.% Ni and between 33.3 and 66.7 at.% Ni by combination of differential thermal analysis (DTA), powder X-ray diffraction (XRD), metallography and electron probe microanalysis (EPMA). The Ni-rich part with more than 50 at.% Ni was investigated by [1959Gua1]. Alloys were melted from carbonyl-Ni (99.9%), Al of 99.99% and Si of 99.98% purity, annealed at 1100 and 900°C and examined by metallography and X-ray diffraction. Solid solubility of Al in , Ni2Si, was studied by [1993Bos] and it was shown that , Ni2Si, could dissolve up to 21 at.% Al. This result has been confirmed by [2002Ric, 2004Ric]. [2004Ric] has reported lattice parameters for the solid solution of Al in ,Ni2Si, as function of composition. NiAl is reported to dissolve about 15 at.% Si [1959Gua1]. The Si solubility of more than 10 at.% Si in NiAl is confirmed by [1977Lit], by [2002Ric] (15 % of Si) and by [2004Ric] (20 % of Si). A partial isothermal section at 750°C with less than 50 at.% Ni content was given in [1969Pan]. The Si solubility in the phase Ni2Al3 was determined by [1969Pan, 1981Ger, 2003Ric1]. According to [1969Pan] approximately 17 % Al may be substituted by Si at 750°C, according to [1981Ger] it is 25 % at 600°C. According to recent measurements of [2003Ric1] 19.2 % Al can be substituted by Si at 550°C that corresponds to 11.5 at.% solubility of Si in the Ni2Al3 phase. [2004Ric] reported solid solubility of Si in the Ni2Al3 phase to be 18 at.% at 800 and 1000°C. The solubility of Si in NiAl3 was reported to be about 0.6 mass% Si by [1951Pra] and 0.7 at.% Si by [2003Ric1]. In NiSi2 33 % Si may be substituted by Al [1969Pan, 1981Ger]. According to [2003Ric1] maximum solubility of Al in NiSi2 at 550°C is 25.7 at.% that means that 38.5 % Si can be substituted by Al. The large ternary solubilities in NiAl, Ni2Al3 and NiSi 2 are compatible with the lattice parameter data of [1962Wit], although these data do not give exact ranges of homogeneity. Lattice parameters for NiSi2-xAlx in the whole homogeneity range up to 25.7 at.% Al have been measured by [2003Ric1]. Binary Systems The Al-Ni and Al-Si binaries are accepted from [2003Sal, 2003Luk]. The phase diagram for Ni-Si systems is accepted from [1999Du], but homogeneity ranges for 2 and 3 phase and phase relations involving J and J´ phases being adopted from [1987Nas]. Solid Phases The solid phases are given in Table 1. [1962Wit] mentioned the possibility that NiAl and NiSi2 may have a common range of homogeneity, regarding the CaF2 structure to be an ordered modification of the CsCl MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
401
structure with 50 % vacancies on the Ni sublattice. [1981Ger], however, gave clearly separated fields for these two phases. These phases have been considered as different phases by [2002Ric, 2003Ric1]. It has been shown by [2003Ric1] that Ni2Si could dissolve up to 25 at.% Al that corresponds to x = 0.77 for chemical formula NiSi2-xAlx. According to [2002Ric] NiAl could dissolve Si. Solid solutions containing ~15 at.% Si has been synthesized by [2002Ric] and lattice parameters for these solid solutions has been measured. [2004Ric] reported lattice parameters of ternary solid solutions of Si in NiAl at 45, 50 and 55 at.% Ni as function of composition in the range between 5 and 20 at.% Al. Some controversy exists regarding the mutual solid solubilities of the isostructural binary phases Ni3Al and Ni3Si. [1959Gua1] reports that Ni3Al at 1100°C may replace 2/3 of Al by Si. [1959Gua2], however, in comparing solubilities of different 3rd elements in Ni3Al claimed Si substitution of 50 % of the Al at 1150°C. [1981Ger, 1981Zar] on the other hand, reported a 600°C isothermal section showing really no Si solubility for Ni3Al. According to [1983Och, 1984Och1, 1984Och2] a continuous solid solution Ni3Al1-xSix with a linear decrease of the lattice parameter was reported for alloys annealed at 1000°C and quenched. The solubility of Al in Ni3Si2 and NiSi was found by [2004Ric] to be very small: 1.0 and 1.5 at.%, respectively. A ternary phase Ni2AlSi (-1) was first reported by [1956Sch, 1957Ess] and confirmed by [1962Wit, 1969Pan, 1981Ger] to have the FeSi structure type. In [1959Gua1] a phase close to this composition was also mentioned. Lattice parameters for Ni2AlSi phase with different Al and Si contents have been recently measured by [2002Ric]. Another ternary phase Ni3(Al1-xSix)7 (-2) (x 0.17) of the Ir3Ge7 type was first reported by [1962Wit] and confirmed by [1969Pan, 1981Ger, 2003Ric1]. The EPMA results of [2003Ric1] show that -2 phase exists in a small composition range from 9 to 11.4 at.% Si. The lattice parameters of -2 for compositions of 9 and 11.4 at.% Si are given in [2003Ric1]. A phase ' which is a superstructure of , Ni2-xSi, was reported by [1994Bos] and a formula Ni8-xAlySi4-y was designated to this phase. The stability of ´ phase has been confirmed by [2002Ric] and crystal structure has been carefully studied. The formula Ni13xAlySi9-y and name -3 has been designated to this phase by [2002Ric]. At 1000°C the extension of the homogeneity range of -3 was found to be much larger than at 800°C [2004Ric]. Based on experimental results of [2004Ric] there is no evidence for two separate phase fields for , Ni2-xSi, and -3. Since the structure of , Ni2-xSi, is not completely clear and structure determination of -3 from quenched samples is only possible in a small part of the homogeneity range, a detailed high temperature XRD study would be necessary to clarify if one single phase forms or closely related superstructures. A phase of approximate composition Ni4AlSi (Ni66Al17Si17) was first mentioned by [1959Gua1] and also reported by [1981Ger]. The X-ray pattern of this phase was complex and no structural analysis was made. Later it has been shown by [1993Bos, 2002Ric] that this phase is a part of , Ni2-xSi1-yAly, solid solution. Richter [2002Ric] has found a new ternary phase (-4) stable at temperature 550°C, but not at 800°C. The composition of this phase is Ni61 Al4Si35. The observed reflections could be indexed with an orthorhombic unit cell [2002Ric]. The space group for this phase is reported by [2004Ric]. Invariant Equilibria The invariant eutectic near the Al corner is well established. A partial reaction scheme, based on [1934Fus] has been recently changed by [2003Ric1] taking into account the ternary phase Ni3(Al1-xSix)7. The partial reaction scheme at Ni content up to 33.3 at.% based on [2003Ric1] data is presented in Fig. 1a. The partial reaction scheme for solid state reactions involving -3 and -4 phases is presented in Fig. 1b. The temperatures and compositions of phases taking part in invariant equilibria involving liquid phase are presented in Table 2.
Landolt-Börnstein New Series IV/11A3
MSIT ®
402
Al–Ni–Si
Liquidus Surface The part of the liquidus surface for compositions up to 66.7 at.% Ni is given in Fig. 2 based on works of [2003Ric1, 2004Ric] but made compatible with the accepted binary systems. The invariant equilibria containing the Ni3(Al1-xSix)7 ternary phase has been experimentally studied by [2003Ric1]. The invariant reactions involving the -1 and -3 phases have been studied by [2004Ric]. A systematic investigation was carried out to determine the solvus in Ni-Al-X ternary systems, with X being transition metal or subgroup B-elements, using the differential thermal analysis (DTA) in [1991Mis]. Solvus isotherms were presented for X = Si, Ga and Ge. In these systems a continuous solid solution was formed between Ni3Al and Ni3Si. However, in this work the solvus is not reproduced, because there was inconsistency between figure captions and figures. Isothermal Sections An isothermal section of the Ni-rich part (>50 at.% Ni) at 1100°C is given by [1959Gua1]. The (Ni) solvus is also given for 900°C. However, the isothermal sections based on these data and presented by [1993Beu] at 900 and 1100°C seem to be inconsistent with new findings of [1993Bos, 1994Bos, 2002Ric] that Ni4AlSi is a part of solid solution , Ni2-xSi1-yAly, and that there is a field of stability of -3 phase. The Al-rich part of the 600°C isothermal section presented by [1993Beu] is based on [1939Wei, 1941Han, 1942Phi, 1959Phi], the Al-poor part is based on [1981Ger] with the solubility of Si in Ni3Al changed according to [1959Gua1]. It should be noted that, isothermal section at 600°C presented by [1993Beu] is also inconsistent with data of [1993Bos, 2002Ric] concerning the Ni4AlSi phase and the existence of the new -4 phase. The isothermal section at 550°C combined from data [1981Ger, 1993Bos, 2002Ric, 2003Ric1, 2004Ric] is presented in Fig. 3a. According to the accepted Al-Ni binary diagram the Ni3Al4 phase is stable up to 710°C. This phase was not found in the ternary system by [1981Ger]. Tie lines between Ni3Al4, Ni1+xAl1-ySiy and Ni2(Al1-xSix)3 are shown tentatively in Fig. 3a. The phase relations at Ni contents up to 33 at.% at 550°C [2003Ric1] are the same as at 600°C [1981Ger]. The only difference is the appearance of a narrow stability field of the liquid phase in the Al-Si binary at 600°C. The phase relations at higher Ni content are assumed to be the same at 550 and 600°C because there is no change in phase stability in this temperature range. This part of phase diagram is accepted from [1993Beu] with corrections made according to data of [1993Bos, 2002Ric, 2004Ric]. Some modifications have been also made to comply the ternary phase diagram with the accepted binaries. The partition of Si between (Ni) and Ni3Al at 1000-1300°C and between Ni3Al and NiAl at 900-1300°C was investigated using diffusion couples by [1994Jia]. Partition coefficients
K SiNi3 Al /( Al ) = xSiNi3 Al / xSi( Ni ) and
K SiNi3 Al / NiAl = xSiNi3 Al / xSiNiAl were determined. It was shown that for the equilibrium between Ni3Al and (Ni) phases partition coefficient is slightly more than one and decreases with increasing temperature. For the equilibrium between Ni3Al and NiAl the partition coefficient is more than one at 900-1100°C and less than one at 1300°C. Isothermal sections at 800 and 1000°C from the experimental study of [2004Ric] are presented in Figs. 3b and 3c. They are based on XRD and EPMA data. The results of [2003Ric1] obtained at 800°C and Ni content between 0 and 33.3 at.% were taken into account by [2004Ric]. Besides the liquid phase which is present in the Al-rich corner of the phase diagram as well as in area adjacent to binary compound NiSi, the section at 1000°C is dominated by extended solid solution phase fields. As it is mentioned above, the experimental results by [2004Ric] could not distinguish between the phase fields of -3 and , Ni2Si.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
403
Temperature – Composition Sections Isoplethal sections at 10, 20, 30 and 33.3 at.% Ni from [2003Ric1] and at 40, 45, 50, 55, 60 and 66.7 at.% Ni from [2004Ric] are presented in Figs. 4 a-j, slightly modified for consistency with the accepted binary diagrams. Thermodynamics [1984Mar] measured the enthalpy of melting of the ternary eutectic LNiAl3+(Al)+(Si), to be 12.22 kJ#(mol-1 of atoms). The partial enthalpy of Ni at infinite dilution in Al-Si melts was measured by [1985Eml] ranging from -139.1 kJ#mol-1 in pure Al to -140.3 kJ#mol-1 in Al+45 at.% Si at 1547°C. [2000Wit] determined partial and integral enthalpies of mixing of liquid Al-Ni-Si alloys by high-temperature isoperibolic calorimetry for three sections with constant concentration ratio of Ni and Si at 1302°C. The results of [2000Wit] are shown in Fig. 5 (partial enthalpies of mixing) and Fig. 6 (integral enthalpies of mixing). The integral enthalpy of mixing of liquid Al-Ni-Si alloys exhibits a highly negative and strongly asymmetric dependence on composition with a minimum near Al0.26Ni0.56Al0.18, which gives evidence of short-range ordering. Using a regular associate model entropy and Gibbs energy of mixing for liquid Al-Ni-Si alloys have been calculated at 1302°C by [2000Wit]. The contribution of the ternary excess term is essential and the regular associate model description of enthalpy of mixing of liquid corresponds to the experimental data only if a ternary associate with the stoichiometry Ni2AlSi is assumed. The chemical potential of Al in Al-Ni-Si melt was derived from EMF measurements at 900°C and compositions with different ratio xNi/xSi = 0.066, 0.215 and 1.02. These data are presented in Fig. 7. It shows that the chemical potential of Al increases at high Al content (xAl > 0.75) and in contrast, decreases when xNi/xSi increases at low Al content. The derived activity of Al shows negative deviation from ideality. Addition of Ni to Al-Si alloys increases the deviation from ideality. The heat capacity of Ni3(Al1-xSix) alloys for x = 0, 0.05, 0.08 and 0.15 from 1.4 to 25 K obtained using semiadiabatic heat pulse method is presented in Fig. 8. Calculations of the ternary system have been performed by [1985Kau], however, without taking into account the ternary phases. Notes on Materials Properties and Applications Mechanical properties of Ni3(Al,Si) xSi = 0.025 single crystal with stress axes parallel to crystallographic orientation near [001] were investigated by both compressive creep and compression tests at temperature of 900°C by [1991Miu]. Magnetic properties of Ni3(Al,Si) at x = 0-0.1 were measured at temperatures 1.8-400 K by [1993Ful]. It was shown that when Si is substituted for Al, the Curie temperature decreases and goes to 0 K at a critical concentration of about 10 % Si. The electrical resistivity of NiSi2-xAlx phase was measured at 4.2-300 K at xAl = 0.15, 0.26 and 0.3 by [2003Ric2]. The studied solid solution is a promising materials for silicon epitaxy as it shows perfect lattice match to Si at composition xAl = 0.26. The conditions for precipitation of fine ductile (Ni) particles in the Ni3Al matrix were established by [1998Mer]. This could improve mechanical properties of Ni3Al alloy. Miscellaneous The Al-Ni2Si reactions were studied in lateral diffusion couples containing Al islands on Ni-Si multiple layers by [1990Liu]. The samples were first in situ annealed in transmission electron microscope at temperatures of 370°C to form Ni2Si phase in the multiple-layer area. Then they were in situ annealed at temperatures in the range of 498-545°C. During the second-step anneal a sequential formation of NiAl3, Ni2Al3 and Ni 3Si2 was observed. The lateral growth of NiAl3 and Ni2Al3 is a result of Al diffusion in Al-Ni silicide reaction, the lateral growth of Ni3Si2 is caused by the diffusion of Si atoms dissociated from the silicides. Diffusion of Si in the Ni3Al phase has been studied from 900 to 1325°C using the diffusion couple (Ni-24.2 Al (at.%), Ni-22.3Al-3.14Si (at.%)) by [1994Min]. The diffusion profiles in the annealed diffusion couple Landolt-Börnstein New Series IV/11A3
MSIT ®
404
Al–Ni–Si
were measured by electron probe microanalysis. The diffusion coefficient of Si was derived from the diffusion profiles and activation energies were calculated. The effect of alloying elements on the morphological stability of the interface between Ni3Al and NiAl phases was investigated using ternary diffusion couples annealed at temperatures in the range of 900-1300°C by [2001Kai]. Planar stable interfaces were found in couples with Si. The structure and thermal stability of rapidly solidified Al-Ni-Si alloys have been investigated using X-ray diffraction and thermal analysis measurements by [1986Dun]. Series of alloys Ni14Al86-xSix showed a region of stoichiometry that yields icosahedral symmetry and a region that yields an amorphous phase. References [1926His] [1930Ota] [1934Fus] [1939Wei] [1941Han] [1942Phi]
[1951Pra]
[1956Sch]
[1957Ess] [1959Gua1] [1959Gua2] [1959Phi] [1962Wit]
[1969Pan]
[1977Lit]
[1978Bha]
[1979Ell]
MSIT®
Hisatsure, C., Suiyókuai Shi, 5, 52 (1926) (Experimental, Equi. Diagram) Otani, B., “Silumin and its Structure” (in Japanese), Kinzuku no Kenkyu, 7, 666-686 (1930) (Equi. Diagram, Experimental, 10) Fuss, V., “Metallography of Aluminium and its Alloys” (in German), Springer Verlag, Berlin, 143-145 (1934) (Equi. Diagram, Review, 1) Weisse, E., “The Al Corner of the Ternary Al-Ni-Si System” (in German), Aluminium Archiv, 26, 5-25 (1939) (Experimental, Equi. Diagram, 16) Hanemann, H., Schrader, A., “On the Ternary Systems of Al” (in German), Z. Metallkd., 33, 20-21 (1941) (Experimental, Equi. Diagram, 3) Phillips, H.W.L., “The Constitution of the Aluminium-Rich Alloys of the Aluminium-Nickel-Iron and Aluminium-Nickel-Silicon Systems”, J. Inst. Met., 68, 27-46 (1942) (Experimental, Equi. Diagram, 15) Pratt, J.N., Raynor, G.V., “The Intermetallic Compounds in the Alloys of Aluminium and Silicon with Chromium, Manganese, Iron, Cobalt and Nickel”, J. Inst. Met., 79, 211-232 (1951) (Experimental, Equi. Diagram, 32) Schubert, K., Burkhardt, W., Esslinger, P., Günzel, E., Meissner, H.G., Schütt, W., Wegst, J., Wilkens, M., “Some Structural Results on Metallic Phases” (in German), Naturwissenschaften., 43, 248-249 (1956) (Crys. Structure, 17) Esslinger, P., Schubert, K., “On the Systematics of the Structure Family NiAs” (in German), Z. Metallkd., 48, 126-136 (1957) (Experimental, Review, Crys. Structure, 19) Guard, R.W., Smith, E.A., “Constitution of Nickel-Base Ternary Alloys. III: Ni-Al-Si System”, J. Inst. Met., 88, 369-374 (1959) (Experimental, Equi. Diagram, #, 5) Guard, R.W., Westbrook, J.H., “Alloying Behavior of Ni3Al ( ' Phase)”, Trans. Met. Soc. AIME, 215, 807-814 (1959) (Experimental, Equi. Diagram, 27) Phillips, H.W.L., “Annotated Equilibrium Diagram of Some Aluminium Alloys Systems”, Inst. Metall, London, 84-86 (1959) (Equi. Diagram, Review, #, 6) Wittmann, A., Burger, K.O., Nowotny, H., “Investigations in the Ternary System, Ni-Al-Si as Well as of Mono- and Disilicides of Some Transition Metals” (in German), Monatsh. Chem., 93, 674-680 (1962) (Experimental, Crys. Structure, 20) Panday, P.K., Schubert, K., “Structure Investigations in Some Mixtures T-B3-B4 (T = Mn, Fe, Co, Ir, Ni, Pd; B3 = Al, Ga, Tl; B4 = Si, Ge)” (in German), J. Less-Common Met., 18, 175-202 (1969) (Experimental, Crys. Structure, 32) Litvinov, V.S., Lesnikova, Ye.G., “ Phase Stability in Ni-Al-Si Alloys”, Phys. Met. Metallogr., 44, 150-153, translated from Fiz. Met. Metalloved., 44, 1297-1299 (1977) (Experimental, 7) Bhan, S., Kudielka, H., “Ordered bcc Phases at High Temperature in Alloys of Transition Metals and B-Subgroup Elements”, Z. Metallkd., 66, 333-336 (1978) (Experimental, Crys. Structure, 18) Ellner, M., Heinrich, S., Bhargava, M.K., Schubert, K., “Structure Study of the Ni-Si System” (in German), J. Less-Common Met., 66, 163-173 (1979) (Experimental, Equi. Diagram, Crys. Structure, 22) Landolt-Börnstein New Series IV/11A3
Al–Ni–Si [1981Ger] [1981Zar]
[1983Och] [1984Mar]
[1984Och1]
[1984Och2] [1985Eml]
[1985Kau]
[1986Dun]
[1987Nas] [1987Hil]
[1990Kuz]
[1990Liu] [1991Mis]
[1991Miu]
[1991Ver]
[1993Beu]
[1993Bos]
Landolt-Börnstein New Series IV/11A3
405
German, N.V., “Ternary Systems Ni-Si-Al and Co-Si-Al” (in Russian), Vestn. Lvov. Univ. Ser. Khim., 23, 61-64 (1981) (Experimental, Equi. Diagram, 6) Zarechnyuk, O.S., German, N.V., Yanson, T.I., Rychal, R.M., Muravyeava, A.A., “Some Phase Diagrams of Aluminium with Transition Metals, Rare Earth Metals and Silicon” (in Russian), Fazovye Ravnovesiya v Metallicheskych Splavach, Nauka, Moscow, 69-73 (1981) (Crys. Structure, Equi. Diagram, Experimental, 5) Ochiau, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Addition”, Bull. P.M.E. (T.I.T.), 52, 1-16 (1983) (Experimental, Equi. Diagram, 7) Martynova, N.M., Rodionova, E.K., Tishura, T.A., Cherneeva, L.I., “Enthalpy of Melting of Metallic Eutectics”, Russ. J. Phys. Chem., 58, 616-617 (1984), translated from Zh. Fiz. Khim., 58, 1009-1010 (1984) (Thermodyn., Experimental, 6) Ochiau, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions”, Bull. P.M.E. (T.I.T.), 53, 15-28 (1984) (Crys. Structure, Experimental, 66) Ochiau, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32, 289-298 (1984) (Experimental, Theory, Thermodyn., 90) Emlin, B.I., Gizenko, N.V., “Investigation of Melts of Aluminium and Silicon With 3d-Metals and the Improvement of the Process of Production of Cast Alloys” (in Russian), Fiz. Khim. Issled. Malootkhod. Prots. Electrotkh., Nauka, Moscow, USSR, 186-194 (1985) (Experimental, Thermodyn., 10) Kaufman, L., “Application of Computer Methods for Calculation of Multicomponent Phase Diagrams of High Temperature Structure Ceramics”, AFOSR-TR-84-0972, 7-11 (1984) (Theory, 0) Dunlap, R.A., Dini, K., “Amorphization of Rapidly Quenched Quasicrystalline Al-Transition Metal Alloys by the Addition of Si”, J. Mater. Res., 1(3), 415-419 (1986) (Crys. Structure, Experimental, 19) Nash, P., Nash, A., “The Ni-Si (Nickel-Silicon) System”, Bull. Alloy Phase Diagrams, 8, 6-14 (1987) (Review, Equi. Diagram, 59) Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremer, F.T., Nickel, H., “Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch., 42a, 1327-1332 (1987) (Equi. Diagram, Experimental, 17) Kuznetsov, G.M., Kalulova, L.M., Mamzurin, O.B., “Phase Equilibria in the Al-Cu-Ni, Al-Cu-Si, Al-Ni-Si and Al-Cu-Ni-Si System Alloys”, Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., 2, 94-100 (1990) (Equi. Diagram, Experimental, Thermodyn., 7) Liu, J.C., Mayer, J.W., “Aluminum and Ni-Silicide Lateral Reactions”, J. Mater. Res., 5(2), 334-340 (1990) (Experimental, Equi. Diagram, Phys. Prop., 19) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130(1991) (Assessment, Experimental, Equi. Diagram, 5) Miura, S., Hayashi, T., Takekawa, M., Mishima, Y., Suzuki,T., “The Compression Creep Behavior of Ni 3Al-X Single Crystals?”, High-Temp.Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc, 213, 623-628 (1991) (Experimental, Phys. Prop., 9) Verhoeven, J.D., Lee, J.H., Laabs, F.C., Jones, L.L., “The Phase Equilibria of Ni3Al Evaluated by Directional Solidification and Diffusion Couple Experiments”, J. Phase Equilib., 12, 15-22 (1991) (Experimental, Equi. Diagram, #, 10) Beuers, G., Bätzner, C., Lukas. H.L., “Aluminium-Nickel-Silicon”, in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.10256.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 28) Bosselet, F., Viala, J.C., Colin, C., Mentzen, B.F., Bouix, J., “Solid State Solubility of Aluminum in the -Ni2Si Nickel Silicate”, J. Mat. Sci. Eng., A167, 147-154 (1993) (Crys. Structure, Equi. Diagram, Experimental, 17) MSIT ®
406 [1993Ful]
[1994Bos]
[1994Jia]
[1994Min]
[1998Mer]
[1999Du]
[2000Wit]
[2001Kai]
[2002Ric] [2003Luk]
[2003Ric1] [2003Ric2]
[2003Sal]
[2004Ric]
MSIT®
Al–Ni–Si Fuller, C.J., Lin, C.L., Mihalisin, T., “Thermodynamic and Magnetic Properties of (Ni1-xMx)3Al with M=Cu and Pd and Ni3(Al1-xSix)”, J. Appl.Phys., 73(10), 5338-5340 (1993) (Crys. Structure, Experimental, Phys. Prop., 13) Bosselet, F., Viala, J.C., Mentzen, B.F., Bouix, J., Colin, C., “'-Ni8-xSi4-yAly: A New Ternary Phase Deriving from -Ni2Si in the Al-Ni-Si System”, J. Mat. Sci. Lett., 13, 358-360 (1994) (Crys. Structure, Experimental, 11) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1), ´(L12) and (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473-485 (1994) (Crys. Structure, Experimental, Equi. Diagram, 25) Minamino, Y., Yamane, T., Saji, S., Hirao, K., Jung, S.B., Kohira, T., “Diffusion of Cu, Fe and Si in L1(2)-Type Intermetallic Compound Ni3Al” (in Japanese), J. Jpn. Inst. Met., 58(4),397-403 (1994) (Crys. Structure, Experimental, Kinetics, 28) Merabtine, R., Devaud-Rzepwski, J., Bertrandt, C. Dallas, J.-P., Trichet M.-F., Cornet, M., “Ductile Phase Precipitation in the L12 Ternary Intermetallic Alloy Ni3(AlSi)”, J. Alloys Compd., 278, 75-77 (1998) (Crys. Structure, Experimental, 11) Du, Y., Schuster, J.C., “Experimental Investigations and Thermodynamic Description of the Ni-Si and C-Ni-Si Systems”, Met. Trans. A, 88A, 2409-2418 (1999) (Equi. Diagram, Experimental, Theory, 44) Witusiewicz, V.T., Arpshofen, I., Seifert, H.J., Sommer, F., Aldinger, F., “Thermodynamics of Liquid and Undercooled Liquid Al-Ni-Si Alloys”, J. Alloys Comp., 305, 151-171 (2000) (Thermodyn., Experimental, 39) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of ´/ Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312,168-175 (2001) (Experimental, Thermodyn., 21) Richter, K.W., “Crystal Structure and Phase Relations of Ni13xAlySi9-y”, J. Alloys Comp., 338, 43-50 (Crys. Structure, Equi. Diagram, Experimental, 16) Lukas, H.L., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 29) Richter, K.W., Isper, H., “The Al-Ni-Si Phase Diagram Between 0 and 33.3 at.% Ni”, Intermetallics, 11, 101-109 (2003) (Crys. Structure, Equi. Diagram, Experimental, 10) Richter, K.W., Hiebl, K., “NiSi1.74Al0.26 and NiSi1.83Ga0.17: Two Materials with Perfect Lattice Match to Si”, Appl. Phys. Lett., 23(3), 497-499 (2003) (Crys. Structure, Electr. Prop., Experimental, 13) 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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164) Richter, K.W., Chandrasekaran, K., Ipser, H., “The Al-Ni-Si Phase Diagram. Part II: Phase Equilibria between 33.3 and 66.7 at.% Ni”, Intermetallics, 12(5), 545-554 (2004) (Crys. Structure, Experimental, Equi. Diagram, 24)
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
407
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Ni) < 1455 (Al) < 660.45 (Si) < 1414 Ni3Al1-xSix
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cF4 Fm3m Cu cF8 Fm3m C-diamond cP4 Pm3m Cu3Au
Ni3Al < 1372 1, Ni3Si < 1035 Ni5Al3 700 Ni1+xAl1-ySiy
oC16 Cmmm Pt5Ga3 cP2 Pm3m CsCl
NiAl < 1638
Ni2(Al1-xSix)3
Ni2Al3 < 1133
Landolt-Börnstein New Series IV/11A3
hP5 P3m1 Ni2Al3
Lattice Parameters Comments/References [pm] a = 352.40
at 25°C [Mas2]
a = 404.96
at 25°C [Mas2]
a = 543.06
at 25°C [Mas2]
a = 356.55 a = 356.9 a = 350 a = 351 a = 354 a = 744 b = 668 c = 372
0 x 1.0 [1984Och1, 1984Och2] 24.5 to 26 at.% Al at 700°C [1987Hil] 23.8 to 26.3 at.% Al at 1200°C [1991Ver] at x = 0.0 [V-C] at x = 0 [1993Bos] at x = 1.0 [1987Nas] at x = 1.0 [1984Och1] at x = 0.5 [1959Gua1] 32 to 36 at.% Al [Mas, V-C]
-0.35 x 0.55 [Mas] 0 y 0.5 [1962Wit] 30.8 to 58 at.% Al [Mas] a = 281.6 at x = 0; y = 0.5 [1962Wit] a = 288.64 at x = 0; y = 0 [V-C] a = 286.21 at x = 0.2020; y = 0.3303 [2002Ric] a = 286.32 at x = 0.2020; y = 0.3193 [2002Ric] a = 287.07 at x = 0.1739; y = 0.1913 [1993Bos] a = 286.89 at x = 0.2173; y = 0.1729 [1993Bos] a = 286.85 at x = 0.3419; y = 0.1686 [1993Bos] a = 285.7 at x = 0.276; y = 0.1479 [1993Bos] a = 285.91 to 282.8 at x = -0.1818; y = 0.091-0.3636 [2004Ric] a = 287.85 to 284.8 at x = 0; y = 0.1-0.4 [2004Ric] a = 286.96 at x = 0.2222; y = 0.1111 [2004Ric] 0 x 0.25 [1962Wit, 1981Ger] at x = 0.25 [1962Wit] a = 400.0 c = 479.1 a = 403.63 59.5 to 63.2 at.% Al [Mas] c = 490.04 at x = 0 [V-C] at x = 0 [2002Ric] a = 403.65 c = 490.03 at x = 0.19167 [2002Ric] a = 401.51 c = 482.31
MSIT ®
Al–Ni–Si
408 Phase/ Temperature Range [°C] Ni3Al4
NiAl3 < 854 3, Ni3Si(h2) 1200 - 1125 2, Ni25 Si9(h1) 1265 - 975
Pearson Symbol/ Space Group/ Prototype cI112 Ia3d Ni3Ga4 oP16 Pnma NiAl3 cP2 Pm3m CsCl hR34 hP34
, Ni31Si12 < 1242
, Ni 2-xSi(h) 1306 - 825
MSIT®
hP43 P321
hP6 P63/mmc Ni2Si
Lattice Parameters Comments/References [pm] a = 1140.8 0.1
[2003Sal]
a = 661.14 b = 736.62 c = 481.12 a = 280.08
[V-C, Mas] max. solubility of Si = 0.6 % [1951Pra]
a = 669.8 c = 2885.5 a = 669.8 c = 961.8 a = 667.1 c = 1228.8 a = 667.9 c = 1222.9 a = 383.6 to 380.2 c = 494.8 to 486.3
90 % of quenched sample [1979Ell] stacking variant, 10 % present in quenched sample [1979Ell] [V-C]
at 1153°C [1978Bha, V-C]
[1993Bos] 0.37 x 0.68 [1979Ell] 33.4 to 41 at.% Si [Mas2] parameters of splat cooled samples [1979Ell]
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si Phase/ Temperature Range [°C]
, Ni 2-xAlySi1-y(r)
Ni2Si < 1255
Landolt-Börnstein New Series IV/11A3
Pearson Symbol/ Space Group/ Prototype oP12 Pbnm Co2Si
409
Lattice Parameters Comments/References [pm] a = 502.2 b = 374.1 c = 708.8 a = 493.2 b = 374.9 c = 716.9 a = 499.5 b = 373.6 c = 707 a = 499.24 b = 374.9 c = 708.51 a = 498.24 b = 374.73 c = 711.04 a = 497.3 b = 374.9 c = 709 a = 497.75 b = 375.16 c = 712.09 a = 497.1 b = 375.61 c = 713.78 a = 496.6 b = 375.89 c = 715.05 a = 495.87 b = 376.92 c = 721.1 a = 492 b = 378.9 c = 732 a = 498 b = 375 c = 711.8 a = 496 b = 376 c = 717.5 a = 495 b = 376.8 c = 722.2 a = 495.8 b = 377.2 c = 723 a = 495.8 b = 378 c = 725.8
at x = 0; y = 0
at x = 0, y = 0.39 [V-C]
at x = 0.0606; y = 0.02939 [2002Ric]
at x = 0.1671; y = 0.1275 [1993Bos]
at x = 0.1751; y = 0.2345 [1993Bos]
at x = 0.2752; y = 0.1260 [1993Bos]
at x = 0.2826; y = 0.337 [1993Bos]
at x = 0.2452; y = 0.3636 [1993Bos]
at x = 0.2376; y = 0.3867 [1993Bos]
at x = 0.1751; y = 0.4689 [1993Bos]
at x = 0.0674; y = 0.6129 [1993Bos]
at x = 0; y = 0.05 [2004Ric]
at x = 0; y = 0.1 [2004Ric]
at x = 0; y = 0.15 [2004Ric]
at x = 0; y = 0.17 [2004Ric]
at x = 0; y = 0.2 [2004Ric]
MSIT ®
Al–Ni–Si
410 Phase/ Temperature Range [°C] J´, Ni3Si2(h) 845 - 800 J, Ni3Si2(r) < 830
NiSi < 992
NiSi2(h) 993 - 981 NiAlxSi2-x
NiSi2(r) < 981 * -1, Ni2AlSi
* -2, Ni3(Al1-xSix)7
MSIT®
Pearson Symbol/ Lattice Parameters Comments/References [pm] Space Group/ Prototype [Mas] oC80 Cmc21 Ni3Si2
oP8 Pnma MnP
cF12 Fm3m CaF2
cP8 P213 FeSi
cI40 Im3m Ir3Ge7
a = 1222.9 b = 1080.5 c = 692.4 a = 1225 b = 1082 c = 693 a = 518 b = 334 c = 562 a = 510.3 b = 333.3 c = 562.8 -
[V-C]
[1993Bos]
[V-C]
xAl = 0.015, xSi = 0.485
[Mas]
a = 551 a = 540.6 a = 541.5 a = 542.2 a = 542.5 a = 542.5 a = 543.0 a = 543.2 a = 543.8 a = 544.9 a = 546 a = 546.8 a = 547.9 a = 548.2 a = 540.6
0 x 0.77 x = 0.5 [1962Wit] x = 0 [V-C] [2003Ric1] x = 0.07[2003Ric1] x = 0.12 [2003Ric1] x = 0.15 [2003Ric1] x = 0.17 [2003Ric1] x = 0.23 [2003Ric1] x = 0.3 [2003Ric1] x = 0.36 [2003Ric1] x = 0.5 [2003Ric1] x = 0.53 [2003Ric1] x = 0.6 [2003Ric1] x = 0.72 [2003Ric1] x = 0.75 [2003Ric1] x = 0 [V-C]
a = 455.9 a = 453.1 to 455.3 a = 453.7 a = 452.99 a = 455.16 a = 829.1 a = 829.1 a = 831.59 a = 830.53
[1956Sch, 1957Ess] [1962Wit] [1981Ger] xAl = 0.165, xSi = 0.32 [2003Ric1] xAl = 0.26, xSi = 0.235 [2003Ric1] x 0.17 [1962Wit] [1981Ger] x = 0.1286 [2003Ric1] x = 0.1629 [2003Ric1]
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si Phase/ Temperature Range [°C] * -3, Ni13xAlySi9-y
* -4, Ni61Al4Si35
Pearson Symbol/ Space Group/ Prototype hP66 P3121 Ga3Ge6Ni13 (designated before as GaGe2Ni4)
oC104 Cmcm Ni16 AlSi9
411
Lattice Parameters Comments/References [pm] a = 766.3 c = 1467 a = 765.3 c = 1466.5 a = 770.2 c = 1472 a = 770.4 c = 1474 a = 770.2 c = 1474 a = 771.2 c = 1473.2 a = 1213.7 b = 1126.5 c = 853.3
x = -0.5714; y = 1.0714 [2002Ric] x = -0.4998; y = 0.9 [2003Ric1] x = 0.5; y = 1.9125 [2003Ric1] x = 0.78481; y = 1.93671 [2003Ric1] x = 1.0769; y = 1.615385 [2003Ric1] x = 0.5; y = 2.25 [1994Bos] [2003Ric1, 2004Ric]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
L + Ni1+xAl1-ySiy Ni2(Al1-xSix)3
1155
e1(max)
L (Si) + NiAlxSi2-x
1085
e2(max)
L Ni1+xAl1-ySiy + NiAlxSi2-x
1080
e3(max)
L Ni1+xAl1-ySiy Ni2(Al1-xSix)3 L (Si) NiAlxSi2-x L Ni1+xAl1-ySiy NiAlxSi2-x L Ni1+xAl1-ySiy Ni2(Al1-xSix)3 NiAlxSi2-x L -3 Ni1+xAl1-ySiy -1 L Ni1+xAl1-ySiy -1 NiAlxSi2-x L NiAlxSi2-x NiSi -1
L + Ni1+xAl1-ySiy Ni2(Al1-xSix)3 + 1071 NiAlxSi2-x
U1
L + -3 + Ni1+xAl1-ySiy -1
998
P1
L + Ni1+xAl1-ySiy -1 + NiAlxSi2-x
969
U2
L + NiAlxSi2-x NiSi + -1
928
U3
Landolt-Börnstein New Series IV/11A3
Composition at.% Al Ni 60 29 49 44 52 40 17 30 0 0 20 33 25.5 37.5 34 34 21 21 33 34 45 34 44 40 22 34 13 49 8 59 29 50 25 50 13 47 35 44 25 50 19 34 6 51 17 34 1 50 23 50
Si 11 7 8 53 100 47 37 21 45 33 21 16 44 38 33 21 25 40 21 25 47 43 49 49 27
MSIT ®
Al–Ni–Si
412 Reaction L NiSi + -3 + -1
T [°C]
Type
925
E1
L + NiAlxSi2-x Ni2(Al1-xSix)3 + (Si) 839
U4
L + NiAl3 + Ni2(Al1-xSix)3 -2
778
P2
L + Ni2(Al1-xSi) 3 -2 + (Si)
775
U5
L + -2 NiAl3 + (Si)
659
U6
L (Al) + (Si) + NiAl3
565
E2
MSIT®
Phase L NiSi -3 -1 L NiSi2-xAlx Ni2(Al1-xSix)3 (Si) L NiAl3 Ni2(Al1-xSix)3 -2 L Ni2(Al1-xSix)3 -2 (Si) L NiAl3 (Si) L (Al) (Si) NiAl3
Composition at.% Al Ni 6 52 1 50 3.5 57.5 20 50.5 56 16 29 33 45 40 0 0 68 12 75 25 50 40 60 30 66 12 50 40 60 30 0 0 76 8 59 30 74 25 0 0 87 2 0 100 0 0 74 25
Si 42 49 39 29.5 28 38 15 100 20 0 10 10 22 10 10 100 16 11 1 100 11 0 100 1
Landolt-Börnstein New Series IV/11A3
Landolt-Börnstein New Series IV/11A3
862
p2
565
659
775
NiAl3+(Al)+(Si)
L NiAl3 + (Al) + (Si)
τ2+NiAl3+(Si)
L + τ2 ΝiAl3 + (Si)
P1
E2
U6
Ni2(Al1xSix)3+τ2+(Si)
L+Ni2(Al1-xSix)3τ2+(Si)
U5
1085 e2max L (Si)+NiAlxSi2-x
L+NiAlxSi2-xNi2Al3+(Si) U4
Al-Ni-Si
NiSi2-xAlx+Ni2Al3+(Si)
839
L+NiAl3+Ni2(Al1-xSix)3τ2
NiAl3+Ni2(Al1-xSix)3
778
Fig. 1a: Al-Ni-Si. Reaction scheme
640 e6 l NiAl3 + (Al)
l + Ni2Al3 NiAl3
Al-Ni
577 e7 L (Al) + (Si)
Al-Si
970 p1 L + (Si) NiSi2
Ni-Si
Al–Ni–Si 413
MSIT ®
Al–Ni–Si
414 Al-Ni-Si τ3(θ)+Ni3Al1-xSix+Ni1+xAl1-ySiy
Ni-Si
τ3(θ)+Ni3Al1-xSix+δ
930 τ3(θ)+Ni3Al1-xSixδ+Ni1+xAl1-ySiy
U
845 p NiSi + θ Ni3Si2(ε/ε')
Ni1+xAl1-ySiy+Ni3Al1-xSix+δ 786 τ3(θ) + δ + Ni3Si2(ε/ε') τ4
P
δ+τ4+Ni3Si2(ε/ε')
τ3(θ)+τ1+Ni1+xAl1-ySiy τ3(θ) + Ni1+xAl1-ySiy δ + τ1 Ni1+xAl1-ySiy+τ1+δ
825 e θ δ + Ni3Si2(ε/ε')
U
τ3(θ) + δτ4 + τ1
U
Ni2(Al1-xSix)3+τ2+(Si) τ3(θ) + τ4 τ1 + Ni3Si2(ε/ε')
U
τ4+τ1+Ni3Si2(ε/ε') τ3(θ)+τ1+NiSi τ3(θ) τ1 + Ni3Si2(ε/ε') + NiSi
770
E
τ1+NiSi+Ni3Si2(ε/ε')
Fig. 1b: Al-Ni-Si. Proposed ternary reaction scheme for solid state reactions according to [2004Ric]. No difference is assumed for ε and ε ' in the Ni-Si binary and for θ and τ 3 in the Al-Ni-Si ternary systems. Temperatures of p and e reactions in the Ni-Si binary system are corrected according to [1987Nas]
Si
Fig. 2: Al-Ni-Si. Partial liquidus surface projection including fields of primary crystallization
Data / Grid: at.% Axes: at.%
1400°C
20
1300°C
80
(Si)
1200°C p1 e4
40
NiSi e5
60
U3
60
E1 τ 1
θ /τ 3
°C 1100
e2max U2
°C 1000
NiAlxSi2-x e3max
40
P1 U1
Ni 2 Al 3
900°C U4
80
800°C U5 τ
2
NiAl
P2
U6
e1max
20
600°C e7 E2 (Al)
Ni MSIT®
20
40
60
p3 80 p2
NiAl3 e6
Al
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
415
Si Fig. 3a: Al-Ni-Si. Isothermal section at 550°C
Data / Grid: at.%
(Si)
Axes: at.%
20
80
NiAlxSi2-x 40
60
NiSi
ε
60
δ Ni3Si
40
τ4
γ
τ1
80
20
τ2 (Ni)
Ni3Al1-xSix 20
Ni
(Al) 60 80 Ni5Al3 40 Ni Al Ni1+xAl1-ySiy 3 4 Ni2(Al1-xSix)3 NiAl3
Si Fig. 3b: Al-Ni-Si. Isothermal section at 800°C
Al
Data / Grid: at.%
(Si)
Axes: at.%
20
80
NiAlxSi2-x 40
60
NiSi
ε
60
40
τ3(θ )
δ
τ1
80
20
L
Ni1-xAl1-ySiy
Ni
Landolt-Börnstein New Series IV/11A3
20
40
(Al) 60
NiAl3 80
Ni2(Al1-xSix)3
Al
MSIT ®
Al–Ni–Si
416
Si Fig. 3c: Al-Ni-Si. Isothermal section at 1000°C
Data / Grid: at.%
(Si)
Axes: at.%
20
80
40
60
NiAlxSi2-x L 60
40
δ τ 4 (θ ) L
80
Ni2-xAlySi1-y
Ni1+xAl1-ySiy
20
Ni
20
40
60
Ni2(Al1-xSix)3
80
Al
Fig. 4a: Al-Ni-Si. Vertical section at 10 at.% Ni 1250
L+(Si)
L
e2max
Temperature, °C
(Si)+NiAlxSi2-x 1000
970°C
L+(Si)+NiAlxSi2-x
L+τ2 750
L+τ2+(Si)
τ2+(Si)+Ni2(Al1-xSix)3
U6
L+NiAl3
U4
P2 L+Ni2(Ai1-xSix)3+(Si)
NiAl3+τ2+(Si)
L+NiAl3+(Si) E2
Ni Al Si
MSIT®
500
10.00 90.00 0.00
(Al)+NiAl3+(Si) 20
40
Si, at.%
60
80
Ni Al Si
10.00 0.00 90.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
Fig. 4b: Al-Ni-Si. Vertical section at 20 at.% Ni
1250
417
L L+(Si) e2max
Temperature, °C
L+NiAlxSi2-x 1000
970°C L+(Si)+NiAlxSi2-x
L+Ni2(Al1-xSix)3 L+NiAl3+Ni2(Al1-xSix)3
U4 L+(Si)+Ni2(Al1-xSix)3
P2 750
L+τ2+(Si) L+τ2+NiAl3
L+NiAl3
(Si)+NiAlxSi2-x
U5
U6
Ni2(Al1-xSix)3+ (Si)+NiAlxSi2-x
L+(Si)+NiAl3
NiAl3+τ2+(Si)
NiAl3+(Al)+(Si)
Ni Al Si
E2
500 20
20.00 80.00 0.00
40
Ni Al Si
60
Ni2(Al1-xSix)3 +τ2+(Si)
Si, at.%
Fig. 4c: Al-Ni-Si. Vertical section at 30 at.% Ni
20.00 0.00 80.00
L 1250
L+NiAl L+NiAlxSi2-x e2max
Temperature, °C
L+Ni2(Al1-xSix)3
L+(Si)
1000
L+Ni2(Al1-xSix)3+NiAlxSi2-x
970°C
L+Ni2(Al1-xSix)3+(Si)
U4
P2
U5
750
L+(Si)+NiAlxSi2-x (Si)+NiAlxSi2-x
Ni2(Al1-x)3+τ2+NiAl3 Ni2(Al1-xSix)3+ (Si)+NiAlxSi2-x
NiAl3+Ni2(Al1-xSix)3 Ni2(Al1-xSix)3+τ2+(Si)
τ2 Ni Al Si
Landolt-Börnstein New Series IV/11A3
500
30.00 70.00 0.00
20
40
Si, at.%
60
Ni Al Si
30.00 0.00 70.00
MSIT ®
Al–Ni–Si
418
1200
Fig. 4d: Al-Ni-Si. Vertical section at 33 at.% Ni
1100
Temperature, °C
L+(Si) L+NiAlxSi2-x L+(Si)+NiAlxSi2-x
1000
970°C NiAlxSi2-x 900
L+Ni2(Al1-xSix)3+NiAlxSi2-x U4 (Si)+Ni2(Al1-xSix)3+NiAlxSi2-x
Ni Al Si
800
33.00 32.00 35.00
40
50
Ni Al Si
60
Si, at.%
Fig. 4e: Al-Ni-Si. Vertical section at 40 at.% Ni
33.00 0.00 67.00
L 1500
L+Ni1+xAl1-ySiy+NiSi2
Temperature, °C
L+Ni1+xAl1-ySiy 1250
NiSi2+τ1+Ni1+xAl1-ySiy
L+Ni2(Al1-xSix)3+Ni1+xAl1-ySiy
L+(Si)
(Ni2(Al1-xSix)3) 1000
500
40.00 60.00 0.00
20
40
Si, at.%
NiSi2+τ 1+NiSi
Ni2(Al1-xSix)3+NiSi2
L+NiSi+NiSi2
NiSi2 +τ 1
Ni 1+x Al1-y Siy+NiSi 2
750
Ni Al Si
L+NiSi2+ (Si)
L+NiSi2 NiSi2+Ni2(Al1-xSix)3+ Ni1+xAl1-ySiy
MSIT®
L+NiSi2+τ1
NiSi+NiSi2
Ni Al Si
40.00 0.00 60.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
419
Fig. 4f: Al-Ni-Si. Vertical section at 45 at.% Ni
L 1500
Temperature, °C
Ni1+xAl1-ySiy
L+Ni1+xAl1-ySiy+NiSi2
1250
L+NiSi2+τ1
L+(Ni1+xAl1-ySiy)
L+Ni1+xAl1-ySiy+τ1 1000
L+NiSi2
L+NiSi L+NiSi+ NiSi2
Ni1+xAl1-ySiy+NiSi2 NiSi2+τ1
750
NiSi+NiSi2 NiSi2+τ1+NiSi
NiSi2+τ1+Ni1+xAl1-ySiy
Ni Al Si
500 20
45.00 55.00 0.00
Ni Al Si
40
Si, at.%
Fig. 4g: Al-Ni-Si. Vertical section at 50 at.% Ni
45.00 0.00 55.00
L 1500
Temperature, °C
Ni1+xAl1-ySiy 1250
L+τ3(θ )
L+(Ni1+xAl1-ySiy) L+Ni1+xAl1-ySiy+τ3(θ )
L+NiSi2 L+τ1
1000
P1
L+NiSi
L+τ1+τ3(θ )
Ni1+xAl1-ySiy+τ1
L+τ1+NiSi2
U3 L+NiSi+NiSi2
750
τ1
Ni Al Si
Landolt-Börnstein New Series IV/11A3
NiSi
NiSi+τ1
500
50.00 50.00 0.00
20
40
Si, at.%
Ni Al Si
50.00 0.00 50.00
MSIT ®
Al–Ni–Si
420
L
Ni2-xAlySi1-y(δ )+Ni1+xAl1-ySiy+τ3(θ )
Fig. 4h: Al-Ni-Si. Vertical section at 55 at.% Ni
Ni2-xAlySi1-y(δ )+Ni1+xAl1-ySiy+τ1 1500
τ1+τ4+Ni2-xAlySi1-y(δ )
Temperature, °C
L+Ni1+xAl1-ySiy
Ni1+xAl1-ySiy+τ3(θ )+τ1
L+Ni1+xAl1-ySiy+τ3(θ )
1250
τ3(θ )+NiSi L+τ3(θ )+NiSi Ni1+xAl1-ySiy+τ3(θ )
L+τ3(θ )+τ1
P1
1000
τ3(θ )+τ1 Ni1+xAl1-ySiy
L+τ3(θ )
E1
τ1+τ3(θ )+NiSi
750
Ni3Si2(ε)+NiSi+ τ 3( θ )
Ni1+xAl1-ySiy+ Ni2-xAlySi1-y(δ )
Ni Al Si
Ni3Si2(ε)+NiSi+τ1
500 10
55.00 45.00 0.00
Fig. 4i: Al-Ni-Si. Vertical section at 60 at.% Ni
20
30
τ1+τ4+Ni3Si2(ε)
Ni Al Si
40
Si, at.%
55.00 0.00 45.00
Ni1+xAl1-ySiy+τ3(θ )+Ni3Al
L
Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ )+τ3(θ ) 1500
L+Ni1+xAl1-ySiy+τ3(θ )
L+Ni1+xAl1-ySiy
Ni2-x AlySi1-y(δ )+Ni1+xAl1-ySiy+τ1
Temperature, °C
Ni2-xAlySi1-y(δ )+τ1+τ4 1250
Ni3Si2(ε)+τ1+τ4 L+τ3(θ )
1000
750
Ni1+xAl1-ySiy
Ni1+xAl1-ySiy+τ3(θ )
τ 3(θ )
Ni1+xAl1-ySiy+Ni3Al Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ )+Ni3Al
Ni3Si2(ε)
Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ )
Ni Al Si
MSIT®
500
60.00 40.00 0.00
10
20
Si, at.%
30
Ni Al Si
60.00 0.00 40.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
Fig. 4j: Al-Ni-Si. Vertical section at 66.7 at.% Ni
421
L L+Ni1+xAl1-ySiy+τ3(θ )
Ni1+xAl1-ySiy 1500
L+τ3(θ )
Temperature, °C
L+(Ni1+xAl1-ySiy) 1250
?
Ni1+xAl1-ySiy+τ3(θ )
τ 3(θ ) Ni3Al+
1000
Ni1+xAl1-ySiy+τ3(θ ) Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ )+τ3(θ )
Ni3Al+Ni1+xAl1-ySiy 750
Ni3Al+Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ ) Ni2-xAlySi1-y(δ )
Ni Al Si
-1 ∆H , kJ·mol-1 ∆HNi Ni, kJ·mol
Fig. 5a: Al-Ni-Si. Partial enthalpy of mixing of nickel of ternary liquid and undercooled liquid Al-Ni-Si alloys at 1302 3°C. Standard states: Al(l), Ni(l) and Si(l) [2000Wit]
500 10
66.70 33.30 0.00
20
30
Si, at.%
Ni Al Si
66.70 0.00 33.30
00
Al-Ni Al-Ni Si0.2 Al-Ni Al-Ni0.8 0.8Si 0.2 Al-Ni0.5 Si0.5 Al-Ni 0.5Si 0.5 Al-Ni Al-Ni0.2 Si0.8 0.2Si 0.8 -40 -40
-80 -80
-120 -120
-160 -160
00
Ni1-y Ni Siyy 1-ySi
Landolt-Börnstein New Series IV/11A3
20 20
40 40
Al,at.% Al, at.%
60 60
80 80
100 100
Al Al
MSIT ®
Al–Ni–Si
422
Fig. 5b: Al-Ni-Si. Partial enthalpy of mixing of silicon of ternary liquid and undercooled liquid Al-Ni-Si alloys at 1302 3°C. Standard states: Al(l), Ni(l) and Si(l) [2000Wit]
0
∆HSi, kJ·mol-1
-50
-100
-150
Al-Si Al-Ni0.8Si0.2 Al-Ni0.5Si0.5 Al-Ni0.2Si0.8
-200
-250 0
20
Ni1-ySiy
Fig. 5c: Al-Ni-Si. Partial enthalpy of mixing of aluminum of ternary liquid and undercooled liquid Al-Ni-Si alloys at 1302 3°C. Standard states: Al(l), Ni(l) and Si(l) [2000Wit]
40
60
80
100
Al
Al, at.%
0
∆HAl, kJ·mol-1
-40
-80
-120
Al-Ni Al-Ni0.8Si0.2 Al-Ni0.5Si0.5 Al-Ni0.2Si0.8 Al-Si -160 0
Ni1-ySiy
MSIT®
20
40
Al, at.%
60
80
100
Al
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
423
0
Fig. 6a: Al-Ni-Si. Integral enthalpy of mixing of liquid and undercooled liquid Al-Ni-Si alloys at 1302 3°C. Standard states: Al(l), Ni(l) and Si(l) [2000Wit]
-10
-20
∆H, kJ·mol-1
-30
-40
-50
Al-Ni Al-Ni0.8Si0.2 Al-Ni0.5Si0.5 Al-Ni0.2Si0.8 Al-Si
-60
-70 0
40
20
Ni1-ySiy
80
60
100
Al
Al, at.%
Si Fig. 6b: Al-Ni-Si. Isolines for integral enthalpy of mixing based on experimental data of [2000Wit]
Data / Grid: at.% Axes: at.%
-10 20
80
-20 40
60
-30
-40
60
40
-50 80
-30 -20 -10
Ni
Landolt-Börnstein New Series IV/11A3
-50 -40
20
20
-69 -60
40
60
80
Al
MSIT ®
Al–Ni–Si
424
0
-2
µAl, kJ·mol-1
Fig. 7: Al-Ni-Si. Partial molar free enthalpy of Al in Al-Ni-Si melts at 900°C with respect to mole fraction of Al and p=0.066 (1), 0.215 (2) and 1.020 (3)
-4
1 -6
2 3 -8 0.6
0.8
1.0
XAl
55
Fig. 8: Al-Ni-Si. C/T vs T2 for Ni3(Al1-xSi) with x=0.05, 0.08 and 0.15
Ni3(Al1-xSix) 50
C/T, mJ·mol-1·K-2
45
x=0.08
40
x=0.15 35
x=0.05 30
25
0
25
50
75
100
T2, K2
MSIT®
Landolt-Börnstein New Series IV/11A3
400
Al–Ni–Si
Aluminium – Nickel – Silicon Olga Fabrichnaya, Georg Beuers, Christian Bätzner and Hans Leo Lukas Literature Data The Al-rich corner was studied several times using thermal and microscopic analyses [1926His, 1930Ota, 1934Fus, 1939Wei, 1942Phi]. The phase relations at Ni contents up to 33.3 at.% have been recently studied by [2002Ric, 2003Ric1]. A ternary eutectic exists between (Al), NiAl3 and Si. The values given for temperature and concentration of the eutectic melt are between 560 and 568°C, 3.0 and 5.2 mass% (1.4 and 2.5 at.%) Ni, 11.0 and 11.8 mass% (10.8 and 11.7 at.%) Si. According to the measurements of [2003Ric1] the temperature of ternary eutectic is 565°C and the composition of the liquid is 2 at.% Ni and 11 at.% Si. Isopleths are reported for 2 [1942Phi], 6 and 14 mass% Si [1930Ota] and for 2 [1939Wei, 1942Phi], 2.5 [1930Ota], 3 [1990Kuz], 4 [1930Ota, 1939Wei], 5 [1959Phi], 7.5 and 12.5 [1930Ota] mass% Ni. The isopleths agree well though only [1939Wei] gives a Si solubility in (Al) in agreement with the binary Al-Si system. [1930Ota] ignores that totally and [1942Phi] gives a much lower value. Recently [2002Ric, 2004Ric] experimentally obtained isoplethal sections for 10, 20, 30, 33.3, 40, 45, 50, 55, 60 and 66.7 at.% Ni. [1934Fus] gave the Al-rich liquidus surface indicating two more invariant reactions. However the ternary phase Ni3(Al1-xSix)7 was not taken into account by [1934Fus]. Recently new data on the liquidus surface were reported by [2003Ric1] at compositions up to 33.3 at.% Ni and between 33.3 and 66.7 at.% Ni by combination of differential thermal analysis (DTA), powder X-ray diffraction (XRD), metallography and electron probe microanalysis (EPMA). The Ni-rich part with more than 50 at.% Ni was investigated by [1959Gua1]. Alloys were melted from carbonyl-Ni (99.9%), Al of 99.99% and Si of 99.98% purity, annealed at 1100 and 900°C and examined by metallography and X-ray diffraction. Solid solubility of Al in , Ni2Si, was studied by [1993Bos] and it was shown that , Ni2Si, could dissolve up to 21 at.% Al. This result has been confirmed by [2002Ric, 2004Ric]. [2004Ric] has reported lattice parameters for the solid solution of Al in ,Ni2Si, as function of composition. NiAl is reported to dissolve about 15 at.% Si [1959Gua1]. The Si solubility of more than 10 at.% Si in NiAl is confirmed by [1977Lit], by [2002Ric] (15 % of Si) and by [2004Ric] (20 % of Si). A partial isothermal section at 750°C with less than 50 at.% Ni content was given in [1969Pan]. The Si solubility in the phase Ni2Al3 was determined by [1969Pan, 1981Ger, 2003Ric1]. According to [1969Pan] approximately 17 % Al may be substituted by Si at 750°C, according to [1981Ger] it is 25 % at 600°C. According to recent measurements of [2003Ric1] 19.2 % Al can be substituted by Si at 550°C that corresponds to 11.5 at.% solubility of Si in the Ni2Al3 phase. [2004Ric] reported solid solubility of Si in the Ni2Al3 phase to be 18 at.% at 800 and 1000°C. The solubility of Si in NiAl3 was reported to be about 0.6 mass% Si by [1951Pra] and 0.7 at.% Si by [2003Ric1]. In NiSi2 33 % Si may be substituted by Al [1969Pan, 1981Ger]. According to [2003Ric1] maximum solubility of Al in NiSi2 at 550°C is 25.7 at.% that means that 38.5 % Si can be substituted by Al. The large ternary solubilities in NiAl, Ni2Al3 and NiSi 2 are compatible with the lattice parameter data of [1962Wit], although these data do not give exact ranges of homogeneity. Lattice parameters for NiSi2-xAlx in the whole homogeneity range up to 25.7 at.% Al have been measured by [2003Ric1]. Binary Systems The Al-Ni and Al-Si binaries are accepted from [2003Sal, 2003Luk]. The phase diagram for Ni-Si systems is accepted from [1999Du], but homogeneity ranges for 2 and 3 phase and phase relations involving J and J´ phases being adopted from [1987Nas]. Solid Phases The solid phases are given in Table 1. [1962Wit] mentioned the possibility that NiAl and NiSi2 may have a common range of homogeneity, regarding the CaF2 structure to be an ordered modification of the CsCl MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
401
structure with 50 % vacancies on the Ni sublattice. [1981Ger], however, gave clearly separated fields for these two phases. These phases have been considered as different phases by [2002Ric, 2003Ric1]. It has been shown by [2003Ric1] that Ni2Si could dissolve up to 25 at.% Al that corresponds to x = 0.77 for chemical formula NiSi2-xAlx. According to [2002Ric] NiAl could dissolve Si. Solid solutions containing ~15 at.% Si has been synthesized by [2002Ric] and lattice parameters for these solid solutions has been measured. [2004Ric] reported lattice parameters of ternary solid solutions of Si in NiAl at 45, 50 and 55 at.% Ni as function of composition in the range between 5 and 20 at.% Al. Some controversy exists regarding the mutual solid solubilities of the isostructural binary phases Ni3Al and Ni3Si. [1959Gua1] reports that Ni3Al at 1100°C may replace 2/3 of Al by Si. [1959Gua2], however, in comparing solubilities of different 3rd elements in Ni3Al claimed Si substitution of 50 % of the Al at 1150°C. [1981Ger, 1981Zar] on the other hand, reported a 600°C isothermal section showing really no Si solubility for Ni3Al. According to [1983Och, 1984Och1, 1984Och2] a continuous solid solution Ni3Al1-xSix with a linear decrease of the lattice parameter was reported for alloys annealed at 1000°C and quenched. The solubility of Al in Ni3Si2 and NiSi was found by [2004Ric] to be very small: 1.0 and 1.5 at.%, respectively. A ternary phase Ni2AlSi (-1) was first reported by [1956Sch, 1957Ess] and confirmed by [1962Wit, 1969Pan, 1981Ger] to have the FeSi structure type. In [1959Gua1] a phase close to this composition was also mentioned. Lattice parameters for Ni2AlSi phase with different Al and Si contents have been recently measured by [2002Ric]. Another ternary phase Ni3(Al1-xSix)7 (-2) (x 0.17) of the Ir3Ge7 type was first reported by [1962Wit] and confirmed by [1969Pan, 1981Ger, 2003Ric1]. The EPMA results of [2003Ric1] show that -2 phase exists in a small composition range from 9 to 11.4 at.% Si. The lattice parameters of -2 for compositions of 9 and 11.4 at.% Si are given in [2003Ric1]. A phase ' which is a superstructure of , Ni2-xSi, was reported by [1994Bos] and a formula Ni8-xAlySi4-y was designated to this phase. The stability of ´ phase has been confirmed by [2002Ric] and crystal structure has been carefully studied. The formula Ni13xAlySi9-y and name -3 has been designated to this phase by [2002Ric]. At 1000°C the extension of the homogeneity range of -3 was found to be much larger than at 800°C [2004Ric]. Based on experimental results of [2004Ric] there is no evidence for two separate phase fields for , Ni2-xSi, and -3. Since the structure of , Ni2-xSi, is not completely clear and structure determination of -3 from quenched samples is only possible in a small part of the homogeneity range, a detailed high temperature XRD study would be necessary to clarify if one single phase forms or closely related superstructures. A phase of approximate composition Ni4AlSi (Ni66Al17Si17) was first mentioned by [1959Gua1] and also reported by [1981Ger]. The X-ray pattern of this phase was complex and no structural analysis was made. Later it has been shown by [1993Bos, 2002Ric] that this phase is a part of , Ni2-xSi1-yAly, solid solution. Richter [2002Ric] has found a new ternary phase (-4) stable at temperature 550°C, but not at 800°C. The composition of this phase is Ni61 Al4Si35. The observed reflections could be indexed with an orthorhombic unit cell [2002Ric]. The space group for this phase is reported by [2004Ric]. Invariant Equilibria The invariant eutectic near the Al corner is well established. A partial reaction scheme, based on [1934Fus] has been recently changed by [2003Ric1] taking into account the ternary phase Ni3(Al1-xSix)7. The partial reaction scheme at Ni content up to 33.3 at.% based on [2003Ric1] data is presented in Fig. 1a. The partial reaction scheme for solid state reactions involving -3 and -4 phases is presented in Fig. 1b. The temperatures and compositions of phases taking part in invariant equilibria involving liquid phase are presented in Table 2.
Landolt-Börnstein New Series IV/11A3
MSIT ®
402
Al–Ni–Si
Liquidus Surface The part of the liquidus surface for compositions up to 66.7 at.% Ni is given in Fig. 2 based on works of [2003Ric1, 2004Ric] but made compatible with the accepted binary systems. The invariant equilibria containing the Ni3(Al1-xSix)7 ternary phase has been experimentally studied by [2003Ric1]. The invariant reactions involving the -1 and -3 phases have been studied by [2004Ric]. A systematic investigation was carried out to determine the solvus in Ni-Al-X ternary systems, with X being transition metal or subgroup B-elements, using the differential thermal analysis (DTA) in [1991Mis]. Solvus isotherms were presented for X = Si, Ga and Ge. In these systems a continuous solid solution was formed between Ni3Al and Ni3Si. However, in this work the solvus is not reproduced, because there was inconsistency between figure captions and figures. Isothermal Sections An isothermal section of the Ni-rich part (>50 at.% Ni) at 1100°C is given by [1959Gua1]. The (Ni) solvus is also given for 900°C. However, the isothermal sections based on these data and presented by [1993Beu] at 900 and 1100°C seem to be inconsistent with new findings of [1993Bos, 1994Bos, 2002Ric] that Ni4AlSi is a part of solid solution , Ni2-xSi1-yAly, and that there is a field of stability of -3 phase. The Al-rich part of the 600°C isothermal section presented by [1993Beu] is based on [1939Wei, 1941Han, 1942Phi, 1959Phi], the Al-poor part is based on [1981Ger] with the solubility of Si in Ni3Al changed according to [1959Gua1]. It should be noted that, isothermal section at 600°C presented by [1993Beu] is also inconsistent with data of [1993Bos, 2002Ric] concerning the Ni4AlSi phase and the existence of the new -4 phase. The isothermal section at 550°C combined from data [1981Ger, 1993Bos, 2002Ric, 2003Ric1, 2004Ric] is presented in Fig. 3a. According to the accepted Al-Ni binary diagram the Ni3Al4 phase is stable up to 710°C. This phase was not found in the ternary system by [1981Ger]. Tie lines between Ni3Al4, Ni1+xAl1-ySiy and Ni2(Al1-xSix)3 are shown tentatively in Fig. 3a. The phase relations at Ni contents up to 33 at.% at 550°C [2003Ric1] are the same as at 600°C [1981Ger]. The only difference is the appearance of a narrow stability field of the liquid phase in the Al-Si binary at 600°C. The phase relations at higher Ni content are assumed to be the same at 550 and 600°C because there is no change in phase stability in this temperature range. This part of phase diagram is accepted from [1993Beu] with corrections made according to data of [1993Bos, 2002Ric, 2004Ric]. Some modifications have been also made to comply the ternary phase diagram with the accepted binaries. The partition of Si between (Ni) and Ni3Al at 1000-1300°C and between Ni3Al and NiAl at 900-1300°C was investigated using diffusion couples by [1994Jia]. Partition coefficients
K SiNi3 Al /( Al ) = xSiNi3 Al / xSi( Ni ) and
K SiNi3 Al / NiAl = xSiNi3 Al / xSiNiAl were determined. It was shown that for the equilibrium between Ni3Al and (Ni) phases partition coefficient is slightly more than one and decreases with increasing temperature. For the equilibrium between Ni3Al and NiAl the partition coefficient is more than one at 900-1100°C and less than one at 1300°C. Isothermal sections at 800 and 1000°C from the experimental study of [2004Ric] are presented in Figs. 3b and 3c. They are based on XRD and EPMA data. The results of [2003Ric1] obtained at 800°C and Ni content between 0 and 33.3 at.% were taken into account by [2004Ric]. Besides the liquid phase which is present in the Al-rich corner of the phase diagram as well as in area adjacent to binary compound NiSi, the section at 1000°C is dominated by extended solid solution phase fields. As it is mentioned above, the experimental results by [2004Ric] could not distinguish between the phase fields of -3 and , Ni2Si.
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
403
Temperature – Composition Sections Isoplethal sections at 10, 20, 30 and 33.3 at.% Ni from [2003Ric1] and at 40, 45, 50, 55, 60 and 66.7 at.% Ni from [2004Ric] are presented in Figs. 4 a-j, slightly modified for consistency with the accepted binary diagrams. Thermodynamics [1984Mar] measured the enthalpy of melting of the ternary eutectic LNiAl3+(Al)+(Si), to be 12.22 kJ#(mol-1 of atoms). The partial enthalpy of Ni at infinite dilution in Al-Si melts was measured by [1985Eml] ranging from -139.1 kJ#mol-1 in pure Al to -140.3 kJ#mol-1 in Al+45 at.% Si at 1547°C. [2000Wit] determined partial and integral enthalpies of mixing of liquid Al-Ni-Si alloys by high-temperature isoperibolic calorimetry for three sections with constant concentration ratio of Ni and Si at 1302°C. The results of [2000Wit] are shown in Fig. 5 (partial enthalpies of mixing) and Fig. 6 (integral enthalpies of mixing). The integral enthalpy of mixing of liquid Al-Ni-Si alloys exhibits a highly negative and strongly asymmetric dependence on composition with a minimum near Al0.26Ni0.56Al0.18, which gives evidence of short-range ordering. Using a regular associate model entropy and Gibbs energy of mixing for liquid Al-Ni-Si alloys have been calculated at 1302°C by [2000Wit]. The contribution of the ternary excess term is essential and the regular associate model description of enthalpy of mixing of liquid corresponds to the experimental data only if a ternary associate with the stoichiometry Ni2AlSi is assumed. The chemical potential of Al in Al-Ni-Si melt was derived from EMF measurements at 900°C and compositions with different ratio xNi/xSi = 0.066, 0.215 and 1.02. These data are presented in Fig. 7. It shows that the chemical potential of Al increases at high Al content (xAl > 0.75) and in contrast, decreases when xNi/xSi increases at low Al content. The derived activity of Al shows negative deviation from ideality. Addition of Ni to Al-Si alloys increases the deviation from ideality. The heat capacity of Ni3(Al1-xSix) alloys for x = 0, 0.05, 0.08 and 0.15 from 1.4 to 25 K obtained using semiadiabatic heat pulse method is presented in Fig. 8. Calculations of the ternary system have been performed by [1985Kau], however, without taking into account the ternary phases. Notes on Materials Properties and Applications Mechanical properties of Ni3(Al,Si) xSi = 0.025 single crystal with stress axes parallel to crystallographic orientation near [001] were investigated by both compressive creep and compression tests at temperature of 900°C by [1991Miu]. Magnetic properties of Ni3(Al,Si) at x = 0-0.1 were measured at temperatures 1.8-400 K by [1993Ful]. It was shown that when Si is substituted for Al, the Curie temperature decreases and goes to 0 K at a critical concentration of about 10 % Si. The electrical resistivity of NiSi2-xAlx phase was measured at 4.2-300 K at xAl = 0.15, 0.26 and 0.3 by [2003Ric2]. The studied solid solution is a promising materials for silicon epitaxy as it shows perfect lattice match to Si at composition xAl = 0.26. The conditions for precipitation of fine ductile (Ni) particles in the Ni3Al matrix were established by [1998Mer]. This could improve mechanical properties of Ni3Al alloy. Miscellaneous The Al-Ni2Si reactions were studied in lateral diffusion couples containing Al islands on Ni-Si multiple layers by [1990Liu]. The samples were first in situ annealed in transmission electron microscope at temperatures of 370°C to form Ni2Si phase in the multiple-layer area. Then they were in situ annealed at temperatures in the range of 498-545°C. During the second-step anneal a sequential formation of NiAl3, Ni2Al3 and Ni 3Si2 was observed. The lateral growth of NiAl3 and Ni2Al3 is a result of Al diffusion in Al-Ni silicide reaction, the lateral growth of Ni3Si2 is caused by the diffusion of Si atoms dissociated from the silicides. Diffusion of Si in the Ni3Al phase has been studied from 900 to 1325°C using the diffusion couple (Ni-24.2 Al (at.%), Ni-22.3Al-3.14Si (at.%)) by [1994Min]. The diffusion profiles in the annealed diffusion couple Landolt-Börnstein New Series IV/11A3
MSIT ®
404
Al–Ni–Si
were measured by electron probe microanalysis. The diffusion coefficient of Si was derived from the diffusion profiles and activation energies were calculated. The effect of alloying elements on the morphological stability of the interface between Ni3Al and NiAl phases was investigated using ternary diffusion couples annealed at temperatures in the range of 900-1300°C by [2001Kai]. Planar stable interfaces were found in couples with Si. The structure and thermal stability of rapidly solidified Al-Ni-Si alloys have been investigated using X-ray diffraction and thermal analysis measurements by [1986Dun]. Series of alloys Ni14Al86-xSix showed a region of stoichiometry that yields icosahedral symmetry and a region that yields an amorphous phase. References [1926His] [1930Ota] [1934Fus] [1939Wei] [1941Han] [1942Phi]
[1951Pra]
[1956Sch]
[1957Ess] [1959Gua1] [1959Gua2] [1959Phi] [1962Wit]
[1969Pan]
[1977Lit]
[1978Bha]
[1979Ell]
MSIT®
Hisatsure, C., Suiyókuai Shi, 5, 52 (1926) (Experimental, Equi. Diagram) Otani, B., “Silumin and its Structure” (in Japanese), Kinzuku no Kenkyu, 7, 666-686 (1930) (Equi. Diagram, Experimental, 10) Fuss, V., “Metallography of Aluminium and its Alloys” (in German), Springer Verlag, Berlin, 143-145 (1934) (Equi. Diagram, Review, 1) Weisse, E., “The Al Corner of the Ternary Al-Ni-Si System” (in German), Aluminium Archiv, 26, 5-25 (1939) (Experimental, Equi. Diagram, 16) Hanemann, H., Schrader, A., “On the Ternary Systems of Al” (in German), Z. Metallkd., 33, 20-21 (1941) (Experimental, Equi. Diagram, 3) Phillips, H.W.L., “The Constitution of the Aluminium-Rich Alloys of the Aluminium-Nickel-Iron and Aluminium-Nickel-Silicon Systems”, J. Inst. Met., 68, 27-46 (1942) (Experimental, Equi. Diagram, 15) Pratt, J.N., Raynor, G.V., “The Intermetallic Compounds in the Alloys of Aluminium and Silicon with Chromium, Manganese, Iron, Cobalt and Nickel”, J. Inst. Met., 79, 211-232 (1951) (Experimental, Equi. Diagram, 32) Schubert, K., Burkhardt, W., Esslinger, P., Günzel, E., Meissner, H.G., Schütt, W., Wegst, J., Wilkens, M., “Some Structural Results on Metallic Phases” (in German), Naturwissenschaften., 43, 248-249 (1956) (Crys. Structure, 17) Esslinger, P., Schubert, K., “On the Systematics of the Structure Family NiAs” (in German), Z. Metallkd., 48, 126-136 (1957) (Experimental, Review, Crys. Structure, 19) Guard, R.W., Smith, E.A., “Constitution of Nickel-Base Ternary Alloys. III: Ni-Al-Si System”, J. Inst. Met., 88, 369-374 (1959) (Experimental, Equi. Diagram, #, 5) Guard, R.W., Westbrook, J.H., “Alloying Behavior of Ni3Al ( ' Phase)”, Trans. Met. Soc. AIME, 215, 807-814 (1959) (Experimental, Equi. Diagram, 27) Phillips, H.W.L., “Annotated Equilibrium Diagram of Some Aluminium Alloys Systems”, Inst. Metall, London, 84-86 (1959) (Equi. Diagram, Review, #, 6) Wittmann, A., Burger, K.O., Nowotny, H., “Investigations in the Ternary System, Ni-Al-Si as Well as of Mono- and Disilicides of Some Transition Metals” (in German), Monatsh. Chem., 93, 674-680 (1962) (Experimental, Crys. Structure, 20) Panday, P.K., Schubert, K., “Structure Investigations in Some Mixtures T-B3-B4 (T = Mn, Fe, Co, Ir, Ni, Pd; B3 = Al, Ga, Tl; B4 = Si, Ge)” (in German), J. Less-Common Met., 18, 175-202 (1969) (Experimental, Crys. Structure, 32) Litvinov, V.S., Lesnikova, Ye.G., “ Phase Stability in Ni-Al-Si Alloys”, Phys. Met. Metallogr., 44, 150-153, translated from Fiz. Met. Metalloved., 44, 1297-1299 (1977) (Experimental, 7) Bhan, S., Kudielka, H., “Ordered bcc Phases at High Temperature in Alloys of Transition Metals and B-Subgroup Elements”, Z. Metallkd., 66, 333-336 (1978) (Experimental, Crys. Structure, 18) Ellner, M., Heinrich, S., Bhargava, M.K., Schubert, K., “Structure Study of the Ni-Si System” (in German), J. Less-Common Met., 66, 163-173 (1979) (Experimental, Equi. Diagram, Crys. Structure, 22) Landolt-Börnstein New Series IV/11A3
Al–Ni–Si [1981Ger] [1981Zar]
[1983Och] [1984Mar]
[1984Och1]
[1984Och2] [1985Eml]
[1985Kau]
[1986Dun]
[1987Nas] [1987Hil]
[1990Kuz]
[1990Liu] [1991Mis]
[1991Miu]
[1991Ver]
[1993Beu]
[1993Bos]
Landolt-Börnstein New Series IV/11A3
405
German, N.V., “Ternary Systems Ni-Si-Al and Co-Si-Al” (in Russian), Vestn. Lvov. Univ. Ser. Khim., 23, 61-64 (1981) (Experimental, Equi. Diagram, 6) Zarechnyuk, O.S., German, N.V., Yanson, T.I., Rychal, R.M., Muravyeava, A.A., “Some Phase Diagrams of Aluminium with Transition Metals, Rare Earth Metals and Silicon” (in Russian), Fazovye Ravnovesiya v Metallicheskych Splavach, Nauka, Moscow, 69-73 (1981) (Crys. Structure, Equi. Diagram, Experimental, 5) Ochiau, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al with Ternary Addition”, Bull. P.M.E. (T.I.T.), 52, 1-16 (1983) (Experimental, Equi. Diagram, 7) Martynova, N.M., Rodionova, E.K., Tishura, T.A., Cherneeva, L.I., “Enthalpy of Melting of Metallic Eutectics”, Russ. J. Phys. Chem., 58, 616-617 (1984), translated from Zh. Fiz. Khim., 58, 1009-1010 (1984) (Thermodyn., Experimental, 6) Ochiau, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions”, Bull. P.M.E. (T.I.T.), 53, 15-28 (1984) (Crys. Structure, Experimental, 66) Ochiau, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32, 289-298 (1984) (Experimental, Theory, Thermodyn., 90) Emlin, B.I., Gizenko, N.V., “Investigation of Melts of Aluminium and Silicon With 3d-Metals and the Improvement of the Process of Production of Cast Alloys” (in Russian), Fiz. Khim. Issled. Malootkhod. Prots. Electrotkh., Nauka, Moscow, USSR, 186-194 (1985) (Experimental, Thermodyn., 10) Kaufman, L., “Application of Computer Methods for Calculation of Multicomponent Phase Diagrams of High Temperature Structure Ceramics”, AFOSR-TR-84-0972, 7-11 (1984) (Theory, 0) Dunlap, R.A., Dini, K., “Amorphization of Rapidly Quenched Quasicrystalline Al-Transition Metal Alloys by the Addition of Si”, J. Mater. Res., 1(3), 415-419 (1986) (Crys. Structure, Experimental, 19) Nash, P., Nash, A., “The Ni-Si (Nickel-Silicon) System”, Bull. Alloy Phase Diagrams, 8, 6-14 (1987) (Review, Equi. Diagram, 59) Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremer, F.T., Nickel, H., “Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch., 42a, 1327-1332 (1987) (Equi. Diagram, Experimental, 17) Kuznetsov, G.M., Kalulova, L.M., Mamzurin, O.B., “Phase Equilibria in the Al-Cu-Ni, Al-Cu-Si, Al-Ni-Si and Al-Cu-Ni-Si System Alloys”, Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., 2, 94-100 (1990) (Equi. Diagram, Experimental, Thermodyn., 7) Liu, J.C., Mayer, J.W., “Aluminum and Ni-Silicide Lateral Reactions”, J. Mater. Res., 5(2), 334-340 (1990) (Experimental, Equi. Diagram, Phys. Prop., 19) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130(1991) (Assessment, Experimental, Equi. Diagram, 5) Miura, S., Hayashi, T., Takekawa, M., Mishima, Y., Suzuki,T., “The Compression Creep Behavior of Ni 3Al-X Single Crystals?”, High-Temp.Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc, 213, 623-628 (1991) (Experimental, Phys. Prop., 9) Verhoeven, J.D., Lee, J.H., Laabs, F.C., Jones, L.L., “The Phase Equilibria of Ni3Al Evaluated by Directional Solidification and Diffusion Couple Experiments”, J. Phase Equilib., 12, 15-22 (1991) (Experimental, Equi. Diagram, #, 10) Beuers, G., Bätzner, C., Lukas. H.L., “Aluminium-Nickel-Silicon”, in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.10256.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 28) Bosselet, F., Viala, J.C., Colin, C., Mentzen, B.F., Bouix, J., “Solid State Solubility of Aluminum in the -Ni2Si Nickel Silicate”, J. Mat. Sci. Eng., A167, 147-154 (1993) (Crys. Structure, Equi. Diagram, Experimental, 17) MSIT ®
406 [1993Ful]
[1994Bos]
[1994Jia]
[1994Min]
[1998Mer]
[1999Du]
[2000Wit]
[2001Kai]
[2002Ric] [2003Luk]
[2003Ric1] [2003Ric2]
[2003Sal]
[2004Ric]
MSIT®
Al–Ni–Si Fuller, C.J., Lin, C.L., Mihalisin, T., “Thermodynamic and Magnetic Properties of (Ni1-xMx)3Al with M=Cu and Pd and Ni3(Al1-xSix)”, J. Appl.Phys., 73(10), 5338-5340 (1993) (Crys. Structure, Experimental, Phys. Prop., 13) Bosselet, F., Viala, J.C., Mentzen, B.F., Bouix, J., Colin, C., “'-Ni8-xSi4-yAly: A New Ternary Phase Deriving from -Ni2Si in the Al-Ni-Si System”, J. Mat. Sci. Lett., 13, 358-360 (1994) (Crys. Structure, Experimental, 11) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1), ´(L12) and (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473-485 (1994) (Crys. Structure, Experimental, Equi. Diagram, 25) Minamino, Y., Yamane, T., Saji, S., Hirao, K., Jung, S.B., Kohira, T., “Diffusion of Cu, Fe and Si in L1(2)-Type Intermetallic Compound Ni3Al” (in Japanese), J. Jpn. Inst. Met., 58(4),397-403 (1994) (Crys. Structure, Experimental, Kinetics, 28) Merabtine, R., Devaud-Rzepwski, J., Bertrandt, C. Dallas, J.-P., Trichet M.-F., Cornet, M., “Ductile Phase Precipitation in the L12 Ternary Intermetallic Alloy Ni3(AlSi)”, J. Alloys Compd., 278, 75-77 (1998) (Crys. Structure, Experimental, 11) Du, Y., Schuster, J.C., “Experimental Investigations and Thermodynamic Description of the Ni-Si and C-Ni-Si Systems”, Met. Trans. A, 88A, 2409-2418 (1999) (Equi. Diagram, Experimental, Theory, 44) Witusiewicz, V.T., Arpshofen, I., Seifert, H.J., Sommer, F., Aldinger, F., “Thermodynamics of Liquid and Undercooled Liquid Al-Ni-Si Alloys”, J. Alloys Comp., 305, 151-171 (2000) (Thermodyn., Experimental, 39) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of ´/ Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312,168-175 (2001) (Experimental, Thermodyn., 21) Richter, K.W., “Crystal Structure and Phase Relations of Ni13xAlySi9-y”, J. Alloys Comp., 338, 43-50 (Crys. Structure, Equi. Diagram, Experimental, 16) Lukas, H.L., “Al-Si (Aluminium-Silicon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 29) Richter, K.W., Isper, H., “The Al-Ni-Si Phase Diagram Between 0 and 33.3 at.% Ni”, Intermetallics, 11, 101-109 (2003) (Crys. Structure, Equi. Diagram, Experimental, 10) Richter, K.W., Hiebl, K., “NiSi1.74Al0.26 and NiSi1.83Ga0.17: Two Materials with Perfect Lattice Match to Si”, Appl. Phys. Lett., 23(3), 497-499 (2003) (Crys. Structure, Electr. Prop., Experimental, 13) 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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164) Richter, K.W., Chandrasekaran, K., Ipser, H., “The Al-Ni-Si Phase Diagram. Part II: Phase Equilibria between 33.3 and 66.7 at.% Ni”, Intermetallics, 12(5), 545-554 (2004) (Crys. Structure, Experimental, Equi. Diagram, 24)
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
407
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Ni) < 1455 (Al) < 660.45 (Si) < 1414 Ni3Al1-xSix
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cF4 Fm3m Cu cF8 Fm3m C-diamond cP4 Pm3m Cu3Au
Ni3Al < 1372 1, Ni3Si < 1035 Ni5Al3 700 Ni1+xAl1-ySiy
oC16 Cmmm Pt5Ga3 cP2 Pm3m CsCl
NiAl < 1638
Ni2(Al1-xSix)3
Ni2Al3 < 1133
Landolt-Börnstein New Series IV/11A3
hP5 P3m1 Ni2Al3
Lattice Parameters Comments/References [pm] a = 352.40
at 25°C [Mas2]
a = 404.96
at 25°C [Mas2]
a = 543.06
at 25°C [Mas2]
a = 356.55 a = 356.9 a = 350 a = 351 a = 354 a = 744 b = 668 c = 372
0 x 1.0 [1984Och1, 1984Och2] 24.5 to 26 at.% Al at 700°C [1987Hil] 23.8 to 26.3 at.% Al at 1200°C [1991Ver] at x = 0.0 [V-C] at x = 0 [1993Bos] at x = 1.0 [1987Nas] at x = 1.0 [1984Och1] at x = 0.5 [1959Gua1] 32 to 36 at.% Al [Mas, V-C]
-0.35 x 0.55 [Mas] 0 y 0.5 [1962Wit] 30.8 to 58 at.% Al [Mas] a = 281.6 at x = 0; y = 0.5 [1962Wit] a = 288.64 at x = 0; y = 0 [V-C] a = 286.21 at x = 0.2020; y = 0.3303 [2002Ric] a = 286.32 at x = 0.2020; y = 0.3193 [2002Ric] a = 287.07 at x = 0.1739; y = 0.1913 [1993Bos] a = 286.89 at x = 0.2173; y = 0.1729 [1993Bos] a = 286.85 at x = 0.3419; y = 0.1686 [1993Bos] a = 285.7 at x = 0.276; y = 0.1479 [1993Bos] a = 285.91 to 282.8 at x = -0.1818; y = 0.091-0.3636 [2004Ric] a = 287.85 to 284.8 at x = 0; y = 0.1-0.4 [2004Ric] a = 286.96 at x = 0.2222; y = 0.1111 [2004Ric] 0 x 0.25 [1962Wit, 1981Ger] at x = 0.25 [1962Wit] a = 400.0 c = 479.1 a = 403.63 59.5 to 63.2 at.% Al [Mas] c = 490.04 at x = 0 [V-C] at x = 0 [2002Ric] a = 403.65 c = 490.03 at x = 0.19167 [2002Ric] a = 401.51 c = 482.31
MSIT ®
Al–Ni–Si
408 Phase/ Temperature Range [°C] Ni3Al4
NiAl3 < 854 3, Ni3Si(h2) 1200 - 1125 2, Ni25 Si9(h1) 1265 - 975
Pearson Symbol/ Space Group/ Prototype cI112 Ia3d Ni3Ga4 oP16 Pnma NiAl3 cP2 Pm3m CsCl hR34 hP34
, Ni31Si12 < 1242
, Ni 2-xSi(h) 1306 - 825
MSIT®
hP43 P321
hP6 P63/mmc Ni2Si
Lattice Parameters Comments/References [pm] a = 1140.8 0.1
[2003Sal]
a = 661.14 b = 736.62 c = 481.12 a = 280.08
[V-C, Mas] max. solubility of Si = 0.6 % [1951Pra]
a = 669.8 c = 2885.5 a = 669.8 c = 961.8 a = 667.1 c = 1228.8 a = 667.9 c = 1222.9 a = 383.6 to 380.2 c = 494.8 to 486.3
90 % of quenched sample [1979Ell] stacking variant, 10 % present in quenched sample [1979Ell] [V-C]
at 1153°C [1978Bha, V-C]
[1993Bos] 0.37 x 0.68 [1979Ell] 33.4 to 41 at.% Si [Mas2] parameters of splat cooled samples [1979Ell]
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si Phase/ Temperature Range [°C]
, Ni 2-xAlySi1-y(r)
Ni2Si < 1255
Landolt-Börnstein New Series IV/11A3
Pearson Symbol/ Space Group/ Prototype oP12 Pbnm Co2Si
409
Lattice Parameters Comments/References [pm] a = 502.2 b = 374.1 c = 708.8 a = 493.2 b = 374.9 c = 716.9 a = 499.5 b = 373.6 c = 707 a = 499.24 b = 374.9 c = 708.51 a = 498.24 b = 374.73 c = 711.04 a = 497.3 b = 374.9 c = 709 a = 497.75 b = 375.16 c = 712.09 a = 497.1 b = 375.61 c = 713.78 a = 496.6 b = 375.89 c = 715.05 a = 495.87 b = 376.92 c = 721.1 a = 492 b = 378.9 c = 732 a = 498 b = 375 c = 711.8 a = 496 b = 376 c = 717.5 a = 495 b = 376.8 c = 722.2 a = 495.8 b = 377.2 c = 723 a = 495.8 b = 378 c = 725.8
at x = 0; y = 0
at x = 0, y = 0.39 [V-C]
at x = 0.0606; y = 0.02939 [2002Ric]
at x = 0.1671; y = 0.1275 [1993Bos]
at x = 0.1751; y = 0.2345 [1993Bos]
at x = 0.2752; y = 0.1260 [1993Bos]
at x = 0.2826; y = 0.337 [1993Bos]
at x = 0.2452; y = 0.3636 [1993Bos]
at x = 0.2376; y = 0.3867 [1993Bos]
at x = 0.1751; y = 0.4689 [1993Bos]
at x = 0.0674; y = 0.6129 [1993Bos]
at x = 0; y = 0.05 [2004Ric]
at x = 0; y = 0.1 [2004Ric]
at x = 0; y = 0.15 [2004Ric]
at x = 0; y = 0.17 [2004Ric]
at x = 0; y = 0.2 [2004Ric]
MSIT ®
Al–Ni–Si
410 Phase/ Temperature Range [°C] J´, Ni3Si2(h) 845 - 800 J, Ni3Si2(r) < 830
NiSi < 992
NiSi2(h) 993 - 981 NiAlxSi2-x
NiSi2(r) < 981 * -1, Ni2AlSi
* -2, Ni3(Al1-xSix)7
MSIT®
Pearson Symbol/ Lattice Parameters Comments/References [pm] Space Group/ Prototype [Mas] oC80 Cmc21 Ni3Si2
oP8 Pnma MnP
cF12 Fm3m CaF2
cP8 P213 FeSi
cI40 Im3m Ir3Ge7
a = 1222.9 b = 1080.5 c = 692.4 a = 1225 b = 1082 c = 693 a = 518 b = 334 c = 562 a = 510.3 b = 333.3 c = 562.8 -
[V-C]
[1993Bos]
[V-C]
xAl = 0.015, xSi = 0.485
[Mas]
a = 551 a = 540.6 a = 541.5 a = 542.2 a = 542.5 a = 542.5 a = 543.0 a = 543.2 a = 543.8 a = 544.9 a = 546 a = 546.8 a = 547.9 a = 548.2 a = 540.6
0 x 0.77 x = 0.5 [1962Wit] x = 0 [V-C] [2003Ric1] x = 0.07[2003Ric1] x = 0.12 [2003Ric1] x = 0.15 [2003Ric1] x = 0.17 [2003Ric1] x = 0.23 [2003Ric1] x = 0.3 [2003Ric1] x = 0.36 [2003Ric1] x = 0.5 [2003Ric1] x = 0.53 [2003Ric1] x = 0.6 [2003Ric1] x = 0.72 [2003Ric1] x = 0.75 [2003Ric1] x = 0 [V-C]
a = 455.9 a = 453.1 to 455.3 a = 453.7 a = 452.99 a = 455.16 a = 829.1 a = 829.1 a = 831.59 a = 830.53
[1956Sch, 1957Ess] [1962Wit] [1981Ger] xAl = 0.165, xSi = 0.32 [2003Ric1] xAl = 0.26, xSi = 0.235 [2003Ric1] x 0.17 [1962Wit] [1981Ger] x = 0.1286 [2003Ric1] x = 0.1629 [2003Ric1]
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si Phase/ Temperature Range [°C] * -3, Ni13xAlySi9-y
* -4, Ni61Al4Si35
Pearson Symbol/ Space Group/ Prototype hP66 P3121 Ga3Ge6Ni13 (designated before as GaGe2Ni4)
oC104 Cmcm Ni16 AlSi9
411
Lattice Parameters Comments/References [pm] a = 766.3 c = 1467 a = 765.3 c = 1466.5 a = 770.2 c = 1472 a = 770.4 c = 1474 a = 770.2 c = 1474 a = 771.2 c = 1473.2 a = 1213.7 b = 1126.5 c = 853.3
x = -0.5714; y = 1.0714 [2002Ric] x = -0.4998; y = 0.9 [2003Ric1] x = 0.5; y = 1.9125 [2003Ric1] x = 0.78481; y = 1.93671 [2003Ric1] x = 1.0769; y = 1.615385 [2003Ric1] x = 0.5; y = 2.25 [1994Bos] [2003Ric1, 2004Ric]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
L + Ni1+xAl1-ySiy Ni2(Al1-xSix)3
1155
e1(max)
L (Si) + NiAlxSi2-x
1085
e2(max)
L Ni1+xAl1-ySiy + NiAlxSi2-x
1080
e3(max)
L Ni1+xAl1-ySiy Ni2(Al1-xSix)3 L (Si) NiAlxSi2-x L Ni1+xAl1-ySiy NiAlxSi2-x L Ni1+xAl1-ySiy Ni2(Al1-xSix)3 NiAlxSi2-x L -3 Ni1+xAl1-ySiy -1 L Ni1+xAl1-ySiy -1 NiAlxSi2-x L NiAlxSi2-x NiSi -1
L + Ni1+xAl1-ySiy Ni2(Al1-xSix)3 + 1071 NiAlxSi2-x
U1
L + -3 + Ni1+xAl1-ySiy -1
998
P1
L + Ni1+xAl1-ySiy -1 + NiAlxSi2-x
969
U2
L + NiAlxSi2-x NiSi + -1
928
U3
Landolt-Börnstein New Series IV/11A3
Composition at.% Al Ni 60 29 49 44 52 40 17 30 0 0 20 33 25.5 37.5 34 34 21 21 33 34 45 34 44 40 22 34 13 49 8 59 29 50 25 50 13 47 35 44 25 50 19 34 6 51 17 34 1 50 23 50
Si 11 7 8 53 100 47 37 21 45 33 21 16 44 38 33 21 25 40 21 25 47 43 49 49 27
MSIT ®
Al–Ni–Si
412 Reaction L NiSi + -3 + -1
T [°C]
Type
925
E1
L + NiAlxSi2-x Ni2(Al1-xSix)3 + (Si) 839
U4
L + NiAl3 + Ni2(Al1-xSix)3 -2
778
P2
L + Ni2(Al1-xSi) 3 -2 + (Si)
775
U5
L + -2 NiAl3 + (Si)
659
U6
L (Al) + (Si) + NiAl3
565
E2
MSIT®
Phase L NiSi -3 -1 L NiSi2-xAlx Ni2(Al1-xSix)3 (Si) L NiAl3 Ni2(Al1-xSix)3 -2 L Ni2(Al1-xSix)3 -2 (Si) L NiAl3 (Si) L (Al) (Si) NiAl3
Composition at.% Al Ni 6 52 1 50 3.5 57.5 20 50.5 56 16 29 33 45 40 0 0 68 12 75 25 50 40 60 30 66 12 50 40 60 30 0 0 76 8 59 30 74 25 0 0 87 2 0 100 0 0 74 25
Si 42 49 39 29.5 28 38 15 100 20 0 10 10 22 10 10 100 16 11 1 100 11 0 100 1
Landolt-Börnstein New Series IV/11A3
Landolt-Börnstein New Series IV/11A3
862
p2
565
659
775
NiAl3+(Al)+(Si)
L NiAl3 + (Al) + (Si)
τ2+NiAl3+(Si)
L + τ2 ΝiAl3 + (Si)
P1
E2
U6
Ni2(Al1xSix)3+τ2+(Si)
L+Ni2(Al1-xSix)3τ2+(Si)
U5
1085 e2max L (Si)+NiAlxSi2-x
L+NiAlxSi2-xNi2Al3+(Si) U4
Al-Ni-Si
NiSi2-xAlx+Ni2Al3+(Si)
839
L+NiAl3+Ni2(Al1-xSix)3τ2
NiAl3+Ni2(Al1-xSix)3
778
Fig. 1a: Al-Ni-Si. Reaction scheme
640 e6 l NiAl3 + (Al)
l + Ni2Al3 NiAl3
Al-Ni
577 e7 L (Al) + (Si)
Al-Si
970 p1 L + (Si) NiSi2
Ni-Si
Al–Ni–Si 413
MSIT ®
Al–Ni–Si
414 Al-Ni-Si τ3(θ)+Ni3Al1-xSix+Ni1+xAl1-ySiy
Ni-Si
τ3(θ)+Ni3Al1-xSix+δ
930 τ3(θ)+Ni3Al1-xSixδ+Ni1+xAl1-ySiy
U
845 p NiSi + θ Ni3Si2(ε/ε')
Ni1+xAl1-ySiy+Ni3Al1-xSix+δ 786 τ3(θ) + δ + Ni3Si2(ε/ε') τ4
P
δ+τ4+Ni3Si2(ε/ε')
τ3(θ)+τ1+Ni1+xAl1-ySiy τ3(θ) + Ni1+xAl1-ySiy δ + τ1 Ni1+xAl1-ySiy+τ1+δ
825 e θ δ + Ni3Si2(ε/ε')
U
τ3(θ) + δτ4 + τ1
U
Ni2(Al1-xSix)3+τ2+(Si) τ3(θ) + τ4 τ1 + Ni3Si2(ε/ε')
U
τ4+τ1+Ni3Si2(ε/ε') τ3(θ)+τ1+NiSi τ3(θ) τ1 + Ni3Si2(ε/ε') + NiSi
770
E
τ1+NiSi+Ni3Si2(ε/ε')
Fig. 1b: Al-Ni-Si. Proposed ternary reaction scheme for solid state reactions according to [2004Ric]. No difference is assumed for ε and ε ' in the Ni-Si binary and for θ and τ 3 in the Al-Ni-Si ternary systems. Temperatures of p and e reactions in the Ni-Si binary system are corrected according to [1987Nas]
Si
Fig. 2: Al-Ni-Si. Partial liquidus surface projection including fields of primary crystallization
Data / Grid: at.% Axes: at.%
1400°C
20
1300°C
80
(Si)
1200°C p1 e4
40
NiSi e5
60
U3
60
E1 τ 1
θ /τ 3
°C 1100
e2max U2
°C 1000
NiAlxSi2-x e3max
40
P1 U1
Ni 2 Al 3
900°C U4
80
800°C U5 τ
2
NiAl
P2
U6
e1max
20
600°C e7 E2 (Al)
Ni MSIT®
20
40
60
p3 80 p2
NiAl3 e6
Al
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
415
Si Fig. 3a: Al-Ni-Si. Isothermal section at 550°C
Data / Grid: at.%
(Si)
Axes: at.%
20
80
NiAlxSi2-x 40
60
NiSi
ε
60
δ Ni3Si
40
τ4
γ
τ1
80
20
τ2 (Ni)
Ni3Al1-xSix 20
Ni
(Al) 60 80 Ni5Al3 40 Ni Al Ni1+xAl1-ySiy 3 4 Ni2(Al1-xSix)3 NiAl3
Si Fig. 3b: Al-Ni-Si. Isothermal section at 800°C
Al
Data / Grid: at.%
(Si)
Axes: at.%
20
80
NiAlxSi2-x 40
60
NiSi
ε
60
40
τ3(θ )
δ
τ1
80
20
L
Ni1-xAl1-ySiy
Ni
Landolt-Börnstein New Series IV/11A3
20
40
(Al) 60
NiAl3 80
Ni2(Al1-xSix)3
Al
MSIT ®
Al–Ni–Si
416
Si Fig. 3c: Al-Ni-Si. Isothermal section at 1000°C
Data / Grid: at.%
(Si)
Axes: at.%
20
80
40
60
NiAlxSi2-x L 60
40
δ τ 4 (θ ) L
80
Ni2-xAlySi1-y
Ni1+xAl1-ySiy
20
Ni
20
40
60
Ni2(Al1-xSix)3
80
Al
Fig. 4a: Al-Ni-Si. Vertical section at 10 at.% Ni 1250
L+(Si)
L
e2max
Temperature, °C
(Si)+NiAlxSi2-x 1000
970°C
L+(Si)+NiAlxSi2-x
L+τ2 750
L+τ2+(Si)
τ2+(Si)+Ni2(Al1-xSix)3
U6
L+NiAl3
U4
P2 L+Ni2(Ai1-xSix)3+(Si)
NiAl3+τ2+(Si)
L+NiAl3+(Si) E2
Ni Al Si
MSIT®
500
10.00 90.00 0.00
(Al)+NiAl3+(Si) 20
40
Si, at.%
60
80
Ni Al Si
10.00 0.00 90.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
Fig. 4b: Al-Ni-Si. Vertical section at 20 at.% Ni
1250
417
L L+(Si) e2max
Temperature, °C
L+NiAlxSi2-x 1000
970°C L+(Si)+NiAlxSi2-x
L+Ni2(Al1-xSix)3 L+NiAl3+Ni2(Al1-xSix)3
U4 L+(Si)+Ni2(Al1-xSix)3
P2 750
L+τ2+(Si) L+τ2+NiAl3
L+NiAl3
(Si)+NiAlxSi2-x
U5
U6
Ni2(Al1-xSix)3+ (Si)+NiAlxSi2-x
L+(Si)+NiAl3
NiAl3+τ2+(Si)
NiAl3+(Al)+(Si)
Ni Al Si
E2
500 20
20.00 80.00 0.00
40
Ni Al Si
60
Ni2(Al1-xSix)3 +τ2+(Si)
Si, at.%
Fig. 4c: Al-Ni-Si. Vertical section at 30 at.% Ni
20.00 0.00 80.00
L 1250
L+NiAl L+NiAlxSi2-x e2max
Temperature, °C
L+Ni2(Al1-xSix)3
L+(Si)
1000
L+Ni2(Al1-xSix)3+NiAlxSi2-x
970°C
L+Ni2(Al1-xSix)3+(Si)
U4
P2
U5
750
L+(Si)+NiAlxSi2-x (Si)+NiAlxSi2-x
Ni2(Al1-x)3+τ2+NiAl3 Ni2(Al1-xSix)3+ (Si)+NiAlxSi2-x
NiAl3+Ni2(Al1-xSix)3 Ni2(Al1-xSix)3+τ2+(Si)
τ2 Ni Al Si
Landolt-Börnstein New Series IV/11A3
500
30.00 70.00 0.00
20
40
Si, at.%
60
Ni Al Si
30.00 0.00 70.00
MSIT ®
Al–Ni–Si
418
1200
Fig. 4d: Al-Ni-Si. Vertical section at 33 at.% Ni
1100
Temperature, °C
L+(Si) L+NiAlxSi2-x L+(Si)+NiAlxSi2-x
1000
970°C NiAlxSi2-x 900
L+Ni2(Al1-xSix)3+NiAlxSi2-x U4 (Si)+Ni2(Al1-xSix)3+NiAlxSi2-x
Ni Al Si
800
33.00 32.00 35.00
40
50
Ni Al Si
60
Si, at.%
Fig. 4e: Al-Ni-Si. Vertical section at 40 at.% Ni
33.00 0.00 67.00
L 1500
L+Ni1+xAl1-ySiy+NiSi2
Temperature, °C
L+Ni1+xAl1-ySiy 1250
NiSi2+τ1+Ni1+xAl1-ySiy
L+Ni2(Al1-xSix)3+Ni1+xAl1-ySiy
L+(Si)
(Ni2(Al1-xSix)3) 1000
500
40.00 60.00 0.00
20
40
Si, at.%
NiSi2+τ 1+NiSi
Ni2(Al1-xSix)3+NiSi2
L+NiSi+NiSi2
NiSi2 +τ 1
Ni 1+x Al1-y Siy+NiSi 2
750
Ni Al Si
L+NiSi2+ (Si)
L+NiSi2 NiSi2+Ni2(Al1-xSix)3+ Ni1+xAl1-ySiy
MSIT®
L+NiSi2+τ1
NiSi+NiSi2
Ni Al Si
40.00 0.00 60.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
419
Fig. 4f: Al-Ni-Si. Vertical section at 45 at.% Ni
L 1500
Temperature, °C
Ni1+xAl1-ySiy
L+Ni1+xAl1-ySiy+NiSi2
1250
L+NiSi2+τ1
L+(Ni1+xAl1-ySiy)
L+Ni1+xAl1-ySiy+τ1 1000
L+NiSi2
L+NiSi L+NiSi+ NiSi2
Ni1+xAl1-ySiy+NiSi2 NiSi2+τ1
750
NiSi+NiSi2 NiSi2+τ1+NiSi
NiSi2+τ1+Ni1+xAl1-ySiy
Ni Al Si
500 20
45.00 55.00 0.00
Ni Al Si
40
Si, at.%
Fig. 4g: Al-Ni-Si. Vertical section at 50 at.% Ni
45.00 0.00 55.00
L 1500
Temperature, °C
Ni1+xAl1-ySiy 1250
L+τ3(θ )
L+(Ni1+xAl1-ySiy) L+Ni1+xAl1-ySiy+τ3(θ )
L+NiSi2 L+τ1
1000
P1
L+NiSi
L+τ1+τ3(θ )
Ni1+xAl1-ySiy+τ1
L+τ1+NiSi2
U3 L+NiSi+NiSi2
750
τ1
Ni Al Si
Landolt-Börnstein New Series IV/11A3
NiSi
NiSi+τ1
500
50.00 50.00 0.00
20
40
Si, at.%
Ni Al Si
50.00 0.00 50.00
MSIT ®
Al–Ni–Si
420
L
Ni2-xAlySi1-y(δ )+Ni1+xAl1-ySiy+τ3(θ )
Fig. 4h: Al-Ni-Si. Vertical section at 55 at.% Ni
Ni2-xAlySi1-y(δ )+Ni1+xAl1-ySiy+τ1 1500
τ1+τ4+Ni2-xAlySi1-y(δ )
Temperature, °C
L+Ni1+xAl1-ySiy
Ni1+xAl1-ySiy+τ3(θ )+τ1
L+Ni1+xAl1-ySiy+τ3(θ )
1250
τ3(θ )+NiSi L+τ3(θ )+NiSi Ni1+xAl1-ySiy+τ3(θ )
L+τ3(θ )+τ1
P1
1000
τ3(θ )+τ1 Ni1+xAl1-ySiy
L+τ3(θ )
E1
τ1+τ3(θ )+NiSi
750
Ni3Si2(ε)+NiSi+ τ 3( θ )
Ni1+xAl1-ySiy+ Ni2-xAlySi1-y(δ )
Ni Al Si
Ni3Si2(ε)+NiSi+τ1
500 10
55.00 45.00 0.00
Fig. 4i: Al-Ni-Si. Vertical section at 60 at.% Ni
20
30
τ1+τ4+Ni3Si2(ε)
Ni Al Si
40
Si, at.%
55.00 0.00 45.00
Ni1+xAl1-ySiy+τ3(θ )+Ni3Al
L
Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ )+τ3(θ ) 1500
L+Ni1+xAl1-ySiy+τ3(θ )
L+Ni1+xAl1-ySiy
Ni2-x AlySi1-y(δ )+Ni1+xAl1-ySiy+τ1
Temperature, °C
Ni2-xAlySi1-y(δ )+τ1+τ4 1250
Ni3Si2(ε)+τ1+τ4 L+τ3(θ )
1000
750
Ni1+xAl1-ySiy
Ni1+xAl1-ySiy+τ3(θ )
τ 3(θ )
Ni1+xAl1-ySiy+Ni3Al Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ )+Ni3Al
Ni3Si2(ε)
Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ )
Ni Al Si
MSIT®
500
60.00 40.00 0.00
10
20
Si, at.%
30
Ni Al Si
60.00 0.00 40.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
Fig. 4j: Al-Ni-Si. Vertical section at 66.7 at.% Ni
421
L L+Ni1+xAl1-ySiy+τ3(θ )
Ni1+xAl1-ySiy 1500
L+τ3(θ )
Temperature, °C
L+(Ni1+xAl1-ySiy) 1250
?
Ni1+xAl1-ySiy+τ3(θ )
τ 3(θ ) Ni3Al+
1000
Ni1+xAl1-ySiy+τ3(θ ) Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ )+τ3(θ )
Ni3Al+Ni1+xAl1-ySiy 750
Ni3Al+Ni1+xAl1-ySiy+Ni2-xAlySi1-y(δ ) Ni2-xAlySi1-y(δ )
Ni Al Si
-1 ∆H , kJ·mol-1 ∆HNi Ni, kJ·mol
Fig. 5a: Al-Ni-Si. Partial enthalpy of mixing of nickel of ternary liquid and undercooled liquid Al-Ni-Si alloys at 1302 3°C. Standard states: Al(l), Ni(l) and Si(l) [2000Wit]
500 10
66.70 33.30 0.00
20
30
Si, at.%
Ni Al Si
66.70 0.00 33.30
00
Al-Ni Al-Ni Si0.2 Al-Ni Al-Ni0.8 0.8Si 0.2 Al-Ni0.5 Si0.5 Al-Ni 0.5Si 0.5 Al-Ni Al-Ni0.2 Si0.8 0.2Si 0.8 -40 -40
-80 -80
-120 -120
-160 -160
00
Ni1-y Ni Siyy 1-ySi
Landolt-Börnstein New Series IV/11A3
20 20
40 40
Al,at.% Al, at.%
60 60
80 80
100 100
Al Al
MSIT ®
Al–Ni–Si
422
Fig. 5b: Al-Ni-Si. Partial enthalpy of mixing of silicon of ternary liquid and undercooled liquid Al-Ni-Si alloys at 1302 3°C. Standard states: Al(l), Ni(l) and Si(l) [2000Wit]
0
∆HSi, kJ·mol-1
-50
-100
-150
Al-Si Al-Ni0.8Si0.2 Al-Ni0.5Si0.5 Al-Ni0.2Si0.8
-200
-250 0
20
Ni1-ySiy
Fig. 5c: Al-Ni-Si. Partial enthalpy of mixing of aluminum of ternary liquid and undercooled liquid Al-Ni-Si alloys at 1302 3°C. Standard states: Al(l), Ni(l) and Si(l) [2000Wit]
40
60
80
100
Al
Al, at.%
0
∆HAl, kJ·mol-1
-40
-80
-120
Al-Ni Al-Ni0.8Si0.2 Al-Ni0.5Si0.5 Al-Ni0.2Si0.8 Al-Si -160 0
Ni1-ySiy
MSIT®
20
40
Al, at.%
60
80
100
Al
Landolt-Börnstein New Series IV/11A3
Al–Ni–Si
423
0
Fig. 6a: Al-Ni-Si. Integral enthalpy of mixing of liquid and undercooled liquid Al-Ni-Si alloys at 1302 3°C. Standard states: Al(l), Ni(l) and Si(l) [2000Wit]
-10
-20
∆H, kJ·mol-1
-30
-40
-50
Al-Ni Al-Ni0.8Si0.2 Al-Ni0.5Si0.5 Al-Ni0.2Si0.8 Al-Si
-60
-70 0
40
20
Ni1-ySiy
80
60
100
Al
Al, at.%
Si Fig. 6b: Al-Ni-Si. Isolines for integral enthalpy of mixing based on experimental data of [2000Wit]
Data / Grid: at.% Axes: at.%
-10 20
80
-20 40
60
-30
-40
60
40
-50 80
-30 -20 -10
Ni
Landolt-Börnstein New Series IV/11A3
-50 -40
20
20
-69 -60
40
60
80
Al
MSIT ®
Al–Ni–Si
424
0
-2
µAl, kJ·mol-1
Fig. 7: Al-Ni-Si. Partial molar free enthalpy of Al in Al-Ni-Si melts at 900°C with respect to mole fraction of Al and p=0.066 (1), 0.215 (2) and 1.020 (3)
-4
1 -6
2 3 -8 0.6
0.8
1.0
XAl
55
Fig. 8: Al-Ni-Si. C/T vs T2 for Ni3(Al1-xSi) with x=0.05, 0.08 and 0.15
Ni3(Al1-xSix) 50
C/T, mJ·mol-1·K-2
45
x=0.08
40
x=0.15 35
x=0.05 30
25
0
25
50
75
100
T2, K2
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
425
Aluminium – Nickel – Tantalum Viktor Kuznetsov Literature Data Phase equilibria and intermetallic phase formation has been reviewed by [1990Kum]. However, this was followed by a thorough assessment of the data published up to 1991 by [1993Zak]. They presented graphically the Ni3Al-TaNi3 section, an assessed Scheil reaction scheme, liquidus and solidus projections, the solvus of the ´ (Ni3Al) based phase and two partial isothermal sections for 1000 and 1250°C. The existence of six ternary phases, TaNiAl, TaNi2Al, Ta0.5Ni3Al0.5, Ta5Ni2Al3, Ta~55 Ni~10 Al~35 and TaNiAl2 was accepted. However, some earlier work was not mentioned in [1993Zak]. [1965Ram] had indicated that in addition to the TaNiAl and TaNi2Al phases, which had been established initially by [1964Mar], a phase with a structure “closely resembling” that of NiTi2 (in Table 1 of [1965Ram] denoted as NiTi2) was present in alloys of gross compositions Ta25Ni38Al37 and Ta25Ni25Al50. The phase was found in both the as cast state and after annealing for 7 days at 900°C, but with an amount significantly less after the heat treatment. Moreover, it was found in the as cast sample with a composition of Ta25Ni50Al25, but later transformed almost entirely to TaNi2Al after annealing. Unfortunately, no compositional data for the phase was given. Later, [1974Ali] performed a DTA study of 5 alloys in the Ni3Al-TaNi3 section in the course of a study of the Ni3Al-Ni3Ta-Ni 3Nb pseudoternary system. In more recent years the phase equilibria in this system have been investigated in much detail. [1994Joh] studied five arc-melted alloys with compositions close to NiAl+15 at.% Ta (on the eutectic line) in the as cast and directionally solidified state by using scanning electron microscopy with EDS to measure phase composition. From the results, a fragment of the liquidus projection (for NiAl-Ni2TaAl-NiTaAl composition region) was constructed suggesting a peritectic formation for the Ni2AlTa ternary phase. [2001Miu] used DTA to determine liquidus and solidus temperatures of alloys made by arc-melting Al, Ni and Ta of purities 99.99, 99.95 and 99.9 mass%, respectively, followed by a homogenization treatment of 1000°C for 24 h. [1991Mis] determined the solvus line of the phase at temperatures between 827 and 1327°C using DTA. Energy-dispersive X-ray spectroscopy was used to confirm the phase constitution of the alloys. [1994Jia] studied the partition of Ta between and ´, as well as between the ´ and phases using a diffusion couple technique. The results are presented in tabular form with phase composition and partition coefficients for 1300, 1200, 1100, 1000 and 800°C and also rendered graphically as partial sections for some selected temperatures. [1996Pal] re-investigated two partial isothermal sections for Ta contents of < 50 at.% for 1000 and 1250°C in order to confirm the work of [1993Zak]. 32 compositions were prepared from components of purities of 99.95 mass% Ni, 99.99 mass% Al and 99.97 mass% Ta using levitation melting. Heat treatment at 1000°C was performed in Ar filled silica ampoules for 168 h for alloys in the NiAl+TaNiAl composition region and for 500 h for alloys of all other compositions. Water quenching followed the heat treatment. At 1250°C, the heat treatment was carried out in a box made from Ta sheets; each specimen was wrapped into Ta foil, and the box was filled with Ti-filings. The heat treatment was carried out in an Ar atmosphere for between 100 and 20 h with subsequent cooling under flowing gas. Samples were examined by metallography, X-ray diffraction and electron microprobe. The results show significant differences from the assessed data of [1993Zak]. [1999Sun] studied the partition of Al and Ta between the liquid and fcc phases in samples quenched from the two phase liquid + fcc state. The compositions of the phases were measured by EPMA. Equilibrium conditions were confirmed by the measure of homogeneity of the solid phase. In addition, they performed a simultaneous regression analysis of their own data, the published data of [1993Zak] and data for the Ni-Cr-Al-Ta quaternary. Good agreement (within approx. 1%, i.e. 7 to 10 K) between the different sets was found. Data for liquid compositions and partition coefficients for Al and Ta were tabulated. Very little work has been carried out on the thermodynamic properties in this system. [1999Roc] measured the low-temperature (3.2 to 10.3 K) heat capacity of the TaNi2Al phase and calculated its electron structure by the LMTO technique. Combining the results of both, the electron-phonon interaction constant was
Landolt-Börnstein New Series IV/11A3
MSIT ®
426
Al–Ni–Ta
derived. Some phase boundaries have been calculated using CALPHAD and ab initio techniques. [1991Kau] performed an approximate CALPHAD calculation of the phase equilibria. However, ternary phases were not taken into account, although dissolution of Al in the TaNi binary compound was allowed in the calculation. [1991Eno] calculated the equilibria between the and ´ phases at 1000°C, using the cluster variation method based on empirical Lennard-Jones type interatomic pair potentials. Good agreement with experimental data was obtained. A number of investigations of mechanical properties have been made. [1991Sas] noticed the precipitation of Ta enriched phase whilst studying the mechanical properties of (NiAl)0.95 Ta0.05. [1996Mac] measured lattice spacing and mechanical properties of the TaNiAl ternary phase. Mechanical properties were also studied by [1991Bon], [1991Hay], [1991Mas], [1991Sas]. Binary Systems The Ni-Ta system is taken from [Mas2], [1991Nas]. For the Al-Ni binary, the latest version [2003Sal] evaluated within the MSIT Binary Evaluation Program is accepted; it does not differ significantly from that of [1987Hil, 1988Bre], which was used by [1993Zak]. The Al-Ta system is taken from [2003Cor], who accepted results of the thermodynamic assessment of the system performed by [1996Du]. Solid Phases [1993Zak] accepted the existence of six ternary phases, TaNiAl, TaNi2Al, Ta0.5Ni3Al0.5, Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2. The TaNiAl phase has a wide solubility range for Al (11 to 50 at.%), but restricted for Ta (32.5 to 37.5 at.%). [1996Pal] noted, that in comparing calculated and observed intensities of X-ray diffractions lines, the suggestion is that Al substitutes for Ni on two different crystallographic sites which exist in the MgZn2 structure to a similar extent. The lattice constants of that phase seem to depend on cooling rate; the reason for this is unclear, but because no peak broadening was observed, it is not likely to be due to stacking faults introduced by thermal stresses on cooling [1996Pal]. The true composition of the TaNi2Al phase was found to be off-stoichiometric: 51 to 55% Al and 22.5 to 25% Ta at 1000°C; 52 to 58% Al and 17.5 to 24% Ta at 1250°C [1996Pal]. [1996Pal] did not find any trace of the Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2 phases as presented by [1993Zak], nor the NiTi2 phase reported by [1965Ram]. The existence of the first three was explicitly rejected; the latter was not considered anyhow by [1996Pal], but no such phase was detected in the composition range studied by [1965Ram]. As [1993Zak] noted weak support for the existence of Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2, these phases are considered in present review to be non-existent. The ternary phase proposed by [1965Ram] seems to be metastable; also, phases with that structure are often stabilized by impurities such as C, N or O. Crystallographic information for the solid phases, including the probably metastable ternary phase, is summarized in Table 1. Detailed data for the concentration dependence of the lattice spacing of TaNiAl [1996Pal] are given on Fig. 1. For 20 < xNi < 50 (xNi in at.%) that dependence is essentially linear: a (pm) = 487.6+0.413(50-xNi), c (pm) = 791.5+0.476(50-xNi), though marked deviations from that can be seen for less Ni [1996Pal]. For lattice spacing of the phase, linearity holds for all compositions studied: a (pm) = 487.5+0.386(50-xNi), c (pm) = 2653+2.90(50-xNi) (also for 20 < xNi < 50 at.%). Pseudobinary Systems No pseudobinary sections have been found in the system, though some authors suggested such behavior for the Ni3Al-TaNi3 section, see section “Temperature – Composition Sections”. Invariant Equilibria Data for the invariant equilibria and Scheil reaction scheme (Fig. 2) were assessed by [1993Zak], and are accepted here with some alterations. Table 2 is based on [1993Zak], but with a corrected error in the temperature of the U4 reaction, noted by [1996Pal]. The reaction U5, presented by [1993Zak] has been omitted as it was shown to be unlikely by [1996Pal]. The eutectic e1(min) L+TaNiAl, is added from MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
427
[1994Joh]. It is not possible to include the ternary peritectic reaction +TaNiAl+LTaNi2Al and surrounding univariant eutectic L+TaNiAl and peritectics +LTaNi2Al and TaNiAl+LTaNi2Al suggested by [1994Joh] in either Table 2 or the reaction scheme as neither temperatures nor phase compositions were determined. (See however discussion of liquidus below). Liquidus, Solidus and Solvus Surfaces The liquidus data from [2001Miu] are in good agreement with [1993Zak] for both edge systems, but differ markedly for intermediate compositions. The data of [2001Miu] are preferred as they result from detailed work and seem to be more reliable. On the other hand, [2001Miu] presents mono- and invariant equilibria lines taken from [1986Wil] which is the main source for [1993Zak]. The liquidus surface is presented here in Fig. 3. It is a composite of the liquidus taken from [2001Miu] and the liquidus surfaces of ´, -3 and phases taken from [1993Zak]. The partial liquidus projection from [1994Joh] is added tentatively, although its connection with other parts of liquidus surface remains rather unclear. Figure 4 provides isotherms of the solidus from [2001Miu]. Figures 5 and 6 present the data of [2001Miu] showing the dependence of the liquidus and solidus temperatures on Al variation at parametric Ta content, and on Ta variation at parametric Al content, respectively. These data give more detailed representation than is possible in Figs. 3 and 4. Figure 7 presents the isotherms of the /( + ´) solvus surface as determined by [1991Mis]. Isothermal Sections Isothermal sections at 1273 and 1000°C are presented in Figs. 8 and 9, respectively, generally accepted from [1996Pal]. The results differ significantly from those of the earlier assessment of [1993Zak]. On the other hand, the data given in the original work disagree with the accepted Al-Ta binary system (and even with the binary accepted by the authors [1996Pal] themselves). To maintain consistency, it was necessary to change the region adjacent to Al-Ta system, which in any event is based on just two alloys. In particular, the homogeneity range of TaAl3 phase is removed, and that of Ta2Al3 is split into stoichiometric Ta5Al7 and Ta39Al69 phases at 1250°C (Fig. 8) and into Ta5Al7 and Ta2Al3 at 1000°C (Fig. 9). These changes were suggested by [1996Du] who analyzed the results of [1996Pal] during their assessment of the Al-Ta binary and is accepted here. Also, the position of the phase corners of +TaNiAl+TaAl3 and L+TaNiAl+TaAl3 tie triangles had to be shifted somewhat to make them compatible with the accepted version of the Al-Ni binary. The data of [1994Jia] for - ´ and ´- equilibria, presented in tabular form, are reproduced in Tables 3 and 4. Temperature – Composition Sections [1993Zak] suggested the section Ni3Al-Ni3Ta to be “partly pseudobinary” and mentioned some experiments on directional growth of a “pseudobinary eutectic” [1972Hub, 1974Mol]; the reported composition of the latter is indeed in very good agreement with the composition of the e2 reaction of [1993Zak]. The DTA study of [1974Ali] is also in agreement, though the authors themselves interpreted their results as indication of a simple pseudobinary section with a single eutectic. As indicated by [1993Zak], the Ni3Al-Ni3Ta section cannot be pseudobinary due to the incongruent formation of Ni3Al. Moreover, in the presently accepted version of the Al-Ni binary, the Ni3Al phase becomes off-stoichiometric starting from approx. 1347°C up to the melting point [1987Hil, 1988Bre]. Also, the phase boundaries of ternary TaNi2Al phase as determined for 1000°C by [1996Pal] are not crossed by the Ni3Al-Ni3Ta join. No account of these phenomena was taken by [1993Zak]. On the other hand, the assessed liquidus-solidus region of that section is indeed independently confirmed by the results of [1974Ali] and by directional solidification experiments, reported by [1993Zak]. So, this fragmentary section is reproduced from [1993Zak] with minor corrections and given as Fig. 10, though the true phase relations should be much more complicated both in a region closer to the Ni3Al side and at lower temperatures.
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ta
428 Thermodynamics
No thermodynamic studies have been carried out except for low-temperature (3.2 to 10.3 K) measurements of the heat capacity of the TaNi2Al phase performed by [1999Dar]. Their results, when treated in the standard way (Cp(T) = elT + CD(/T)), give el = 10.01 0.14 mJ#mol#K-2, D = 299 1.9 K. This equation is valid only below approximately 7 K. Notes on Materials Properties and Applications The influence of Ta additions on mechanical properties of NiAl was studied in [1991Mas, 1991Sas]. Such properties of Ta alloyed single crystals of ´ Ni 3Al can be found in [1991Bon]; creep behavior of that phase was studied by [1991Hay]. Some mechanical properties of the Laves phase TaNiAl were measured by [1996Mac]. Miscellaneous [2001Ter] suggested the usage of thermal conductivity measurements for determination of site preferences in the ´ Ni3Al phase. The results are in broad agreement with the phase diagram determinations, which suggest that Ta substitutes for Al in Ni3Al. [2001Kai] investigated the morphological stability of the interface between ´(L1 2) and (B2) phases in diffusion couples. In addition, the results of an unpublished calculation of thermodynamic properties are cited and used in the discussion of the results. References [1964Mar]
[1965Gie] [1965Ram]
[1968Hun]
[1972Hub]
[1972Min]
[1974Ali]
[1974Mol]
[1974Var]
MSIT®
Markiv, V.Ya., Voroshilov, Yu.V., Kripyakevich, P.I., Cherkashin, E.E., “New Compounds of the MnCu 2Al and MgZn2 Types Containing Aluminium and Gallium”, Sov. Phys.-Crystallogr., 9, 619-620 (1964), translated from Kristallografiya, 9, 737-738 (1964) (Crys. Structure, 4) Giessen, B.C., Grant, N.J., “New Intermediate Phases in Transition Metal Systems. II”, Acta Crystallogr., 18, 99 (1965) (Crys. Structure, 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) (Equi. Diagram, Experimental, 14) Hunt, C.R., Raman, A., “Alloy Chemistry of )(U)-Related Phases. I. Extens Ion of - and Occurrence of ´-Phases in the Ternary Systems Nb(Ta)-X-Al (X = Fe, Co, Ni, Cu, Cr, Mo)”, Z. Metallkd., 59(9), 701-707 (1968) (Crys. Structure, Equi. Diagram, 14) Hubert, J.-C., Kurz, W., Lux, B., “Growth by Directed Solidification of the Ni3Al-Ni3Ta Quasibinary Eutectic” (in French), J. Cryst. Growth, 13-14, 757-764 (1972) (Equi. Diagram, 15) Mints, R.S., D´yakonova, N.P., Umansky, Ya.S., Bondarenko, Yu.A., Bondarenko, T.A., “Interaction of the Phase Ni 3Al with Ni3Ta”, Sov. Physics Doklady, 17(9), 904-906 (1973) translated from Dokl. Akad. Nauk SSSR, 206(1), 87-88 (1972) (Crys. Structure, Experimental, 5) Alikhanov, V.A., Pyatnitskii, V.N., Sokolovskaya, E.M., “Phase Diagram of the System Ni3Al-Ni3Nb-Ni3Ta” (in Russian), Vestn. Mosk. Univ., Ser. 2:Khim., 15, 698-701 (1974) (Equi. Diagram, Experimental, 5) Mollard, F., Lux, B., Hubert, J.C., “Directionally Solidified Composites Based on the Ternary Eutectic Ni-Ni3Al-Ni3Ta ( / ´ - )”, Z. Metallkd., 65, 461-468 (1974) (Equi. Diagram, Experimental, 6) Varli, K.V., D’yakonova, N.P., Umansky, Ya.S., Bondarenko, Yu.A., Putman, A.M., “Crystal Structure of the Ternary Phase of the Ni-Ta-Al System”, Vses. Konf. Kristallokhim. Intermet., 2nd, Tezisy Dokl., Lvov Gos. Univ.: Lvov, USSR, 49 (1974) (Crys. Structure, 0)
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta [1979Nas] [1984Och]
[1984Wil]
[1985Mis]
[1986Hua] [1986Wil1]
[1986Wil2]
[1987Hil]
[1987Kha] [1988Bre]
[1990Kum]
[1991Bon]
[1991Eno]
[1991Hay]
[1991Kau] [1991Mas]
[1991Mis]
[1991Nas]
Landolt-Börnstein New Series IV/11A3
429
Nash, P., West, D.T.F., “Phase Equilibria in the Ni-Ta-Al System”, Met. Sci., 13(12), 670-676 (1979) (Equi. Diagram, Crys. Structure, Experimental, 22) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni ( ), Ni3Al ( ´) and Ni3Ga ( ´) Solid Solutions”, Bull. P. M. E.,(T. I. T.), 53, 15-28 (1984) (Crys. Structure, Experimental, Rewiew, 56) Willemin, P., Dugue, O., Durand-Charre, M., Davidson, J., “High-Temperature Phase Equilibria in the Ni-Al-Ta System”, Superall. 1984 Champ., MS/AIME, Conf: Pa. USA, 637-647 (1984) (Equi. Diagram, Crys. Structure, Experimental, 13) Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni( ), Ni3Al( ´) and Ni3Ga( ´) Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33(6), 1161-1169 (1985) (Crys. Structure, Review, 64) 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) Willemin, P., Dugue, O., Durand-Charre, M. J., Davidson, H., “Experimental Determination of Nickel-Rich Corner of Ni-Al-Ta Phase Diagram”, Mater. Sci. Technol., 2(4), 344-348 (1986) (Equi. Diagram, 13) Willemin, P., Durand-Charre,, M., Ansara, I., “Liquid-Solid Equilibria in the System Ni3Al-Ni3Ta and Ni3Al-Ni3Ti”, High Temp. Alloys Cas Turbines Other Appl., Pt.2, Comm. Euro. Communicates, Rep. EUR 10567, 955-964 (1986) (Equi. Diagram, Thermodyn., 8) Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremer, F.T., Nickel, H., “Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch., 42a, 1327-1332 (1987) (Equi. Diagram, Experimental, 17) Khadkikar, P.S., Vedula, K., “An Investigation of the Ni5Al3 Phase”, J. Mater. Res., 2(2), 163-167 (1987) (Crys. Structure, Experimental, 7) Bremer, F.J., Beyss, M., Karthaus, E., Hellwig, A., Schober, T., Welter, J.-M., Wenzl, H., “Experimental Analysis of the Ni-Al Phase Diagram”, J. Cryst. Growth, 87, 185-192 (1988) (Equi. Diagram, Experimental, 16) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35, 293-327 (1990) (Crys. Structure, Equi. Diagram, Review, 158) Bonneville, J., Martin, J.L., “The Strain Rate Sensitivity of Ni3(Al,Ta) Single Crystals”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc., 213, 629-634 (1991) (Mechan. Prop., Experimental, 17) Enomoto, M., Harada, H., Yamazaki, M., “Calculation of ´/ Equilibrium Phase Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15, 143-158 (1991) (Equi. Diagram, Calculation, 34) Hayashi, T., Shinoda, T., Mishima, Y., Suzuki, T., “Effect of Off-Stoichiometry on the Creep Behavior of Binary And Ternary Ni3Al”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc., 213, 617-622 (1991) (Mechan. Prop., Experimental, 7) Kaufman, L., “Calculation of the Multicomponent Tantalum Based Phase Diagrams”, Calphad, 15, 261-282 (1991) (Equi. Diagram, Calculation, 15) Maslenkov, S.B., Filin, S.A., Abramov, V.O., “Effect of Structural State and Alloying of Transition Metals on the Degree of Hardening of Ternary Solid Solutions Based on Nickel Monoaluminide”, Russ. Metall. (Engl. Transl.), (1), 115-118 (1991) (Mechan. Prop., Experimental, 10) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Equi. Diagram, Experimental, 5) Nash, A., Nash, P., “The Ni-Ta (Nickel-Tantalum) System”, in “Phase Diagrams of Binary Nickel Alloys, Monograph Series on Alloy Phase Diagrams”,Vol. 6, ASM-Intl., Materials Park, Ohio, 320-325 (1991) (Equi. Diagram, Crys. Structure, Review, 38)
MSIT ®
430 [1991Sas]
[1991Zha]
[1993Kha]
[1993Zak]
[1994Jia]
[1994Joh] [1996Du]
[1996Mac] [1996Pal] [1996Pau]
[1999Roc]
[1999Sun] [2001Kai]
[2001Miu]
[2001Ter]
[2003Cor]
[2003Sal]
MSIT®
Al–Ni–Ta Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on the Solidified Structure of NiAl”, Intermetal. Comp. - Struct. Mechan. Prop., Proc. Conf., 877-881 (1991) (Equi. Diagram, Mechan. Prop., Abstract, 10) Zhao, J.T., Celato, L., Parthe, E., “Structure Refinement of Monoclinic 12-Layer TaNi3 with -NbPt3 Type. New Crystallographic Descriptions of this Type and of the Nb3Rh 5 Type Based on Smaller Unit Cells”, Acta Crystallogr., Sect. C: Crys. Struct. Commun., C47, 479-483 (1991) (Crys. Structure, Experimental, 11) Khadkikar, P.S., Locci, I.E., Vedula, K., Michal, G.M., “Transformation to Ni5Al3 in a 63.0 at.% Ni-Al Alloy”, Metall. Trans. A, 24A, 83-94 (1993) (Equi. Diagram, Crys. Structure, Experimental, 28) Zakharov, A., “Aluminium - Nickel - Tantalum”, in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.14883.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 28) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1), `(L12) and (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, A25, 473-485 (1994) (Equi. Diagram, Experimental, 25) Johnson, D.R., Oliver, B.F., “Ternary Peritectic Solidification in the NiAl-Ni2AlTa-NiAlTa System”, Mater. Lett., 20, 129-133 (1994) (Equi. Diagram, Experimental, 11) Du, Y., Schmid-Fetzer, R., “Thermodynamic Modelling of the Al-Ta System”, J. Phase Equilib., 17, 311-324 (1996) (Equi. Diagram, Crys. Structure, Thermodyn., Assessment, Calculation, 55) Machon, L., Sauthoff, G., “Deformation Behavior of Al-Containing C14 Laves Phase Alloys”, Intermetallics, 4, 469-481 (1996) (Crys. Structure, Experimental, 41) Palm, M., Sanders, W., Sauthoff, G., “Phase Equilibria in the Ni-Al-Ta System”, Z. Metallkd., 87, 390-398 (1996) (Equi. Diagram, Crys. Structure, Experimental, 27) 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) 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), B269, 154-162 (1999) (Thermodyn., Phys. Prop., Experimental, Calculation, 20) Sung, P.K., Poirier, D.R., “Liquid-Solid Partition Ratios in Nickel-Base Alloys”, Metall. Mater. Trans. A, A30, 2173-2181 (1999) (Equi. Diagram, Experimental, 41) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of ´/ Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, A312, 168-175 (2001) (Kinetics, Thermodyn., Experimental, 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) (Equi. Diagram, Experimental, 9) Terada, Y., Ohkubo, K., Mohri, T., Suzuki, T., “Site Preference Determination in Intermetallic Compounds by Thermal Conductivity Measurements”, J. Mater. Res., 16, 2314-2320 (2001) (Thermal Conduct., Crys. Structure, Experimental, Calculation, 63) Cornish, L., Dolotko, O., Rogl, P., “Al-Ta (Aluminium - Tantalum)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram, Assessment, 3) 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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164) Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
431
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Al)
hP2 P63/mmc Mg
a = 269.3 c = 439.8
at 25°C, 20.5 GPa [Mas2]
(Al) < 660.452
cF4 Fm3m Cu
a = 404.96
at 25°C [Mas2]
(Ta) < 3020
cI2 Im3m W
a = 330.30
at 25°C [Mas2]
, (Ni) < 455
cF4 Fm3m Cu
a = 352.40
at 25°C [Mas2]
a = 357.8
at 8.% Ta, 1150°C, linear da/dx [1984Och2, 1985Mis] at 14.% Al, linear da/dx [1984Och2]
TaxNi1-x Ni1-xAlx TaxNi1-x-yAly
a = 355.0
a = 355.3 to 357.5 at 7.5 - 10 at.% Ta, 75 - 80 at.% Ni, 1000°C, quenched, sample contained ´ and -3 [1979Nas] a = 359.3 at 10 at.% Ta, 80 at.% Ni, 1250°C, quenched, sample contained -3 [1979Nas] ´, Ni3Al < 372
cP4 Pm3m AuCu3
[1986Hua] at 63 at.% Ni [1993Kha]
a = 359.0 to 362.4 at 3 - 12 at.% Ta, 58.6 - 80 at.% Ni, 1000 - 1200°C, multiphase samples quenched, linear da/dx [1972Min, 1979Nas, 1984Och2, 1985Mis]
Ta1-xNi3Alx
Ni5Al3 723
oC16 Cmmm Pt5Ga3
, NiAl < 1638 Tax(Ni1-yAly)1-x
cP2 a) Pm3m CsCl
TaNi8 < 307 (Ta1-xAlx)Ni8
tI36 NbNi8
Landolt-Börnstein New Series IV/11A3
a = 356.77 a = 358.9
a = 753 b = 661 c = 376 a = 286.0 a = 287 a = 288.72 0.02 a = 287.98 0.02 a = 286.6 to 296.8
a = 760.5 c = 358.5 a = 767 c = 348
32 to 36 at.% Al at 63 at.% Ni [1993Kha] 42 to 69.2 at.% Ni [Mas2] [1987Kha] at 63 at.% Ni [1993Kha] at 50 at.% Ni [1996Pau] at 54 at.% Ni [1996Pau] 3.0 - 20 at.% Ta, 50 - 70 at.% Ni. Quenched from 1250 - 1000°C. Samples were multiphase. [1979Nas] at 11.1 at.% Ta [1991Nas] at 11.8 at.%, 83.3 at.% Ni, from EMPA, 1250°C, quenched, [1979Nas]
MSIT ®
Al–Ni–Ta
432 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, TaNi 3 < 1547
mP48 b) P21/m TaPt3
a = 452.3 b = 512.6 c = 2544 = 90°
TaNi2 < 1404
tI6 I4/mmm MoSi2
, TaNi < 1570
hR13 R3m W6Fe7
Ta(Ni,Al)
a = 315.4 c = 790.5
at 22.5 to 28.5 at.% Ta [1991Nas] [1991Zha] single crystal
32.5 to 35 at.% Ta [1991Nas] at 33.3 at.% Ta [1991Nas]
50 to 54 at.% Ta [1991Nas] a = 492.1 at 50 at.% Ta [1991Nas] c = 2690.5 a = 491.9 to 497.8 50 - 55 at.% Ta, 35 - 23 at.% Ni c = 2714 to 2735 [1968Hun] a = 496.1 c = 2504
9 at.% Ta, 58.8 at.% Ni 1250°C, quenched, alloy with , , -2 [1979Nas]
a = 428.3 c = 2649
20 at.% Ta, 50 at.% Ni 1250°C, quenched, alloy with -11 [1979Nas]
a = 986.4 c = 521.5
at ~20 to 40 at.% Al [Mas2, V-C]
), Ta2Al < 2061
tP30 P42/mnm )CrFe
TaAl < 1446
mP*
[1996Du]
Ta5Al7 < 1345
hP*
[1996Du]
Ta2Al3 < 1226
cF*
[1996Du]
Ta39Al69 1548 - 1183
cF432 F43m
[1996Du]
TaAl3 < 1608
tI8 I4/mmm TiAl3
a = 383.7 c = 855.0
[V-C]
* -1, TaNiAl
hP12 P63/mmc MgZn2
a = 496.9 c = 798.5 a = 501.5 c = 817.1
[V-C] alloy 20 at.% Ta, 50 at.% Ni 1000°C, quenched, alloy with and -2 [1979Nas]
* -2, TaNi2Al
cF16 Fm3m BiF3
a = 594.9 a = 580 to 594
[V-C] 9 - 20 at.% Ta, 50 - 58.8 at.% Ni, 1000 - 1250°C, quenched. Multiphase samples [1979Nas]
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
433
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
* -3, Ta0.5Ni3Al0.5 < 1393
hP16 P63/mmc TiNi3
a = 510.5 to 513.7 Al-rich [1965Gie, 1972Min, 1974Var, c = 831.9 to 836.6 1979Nas, V-C, 1984Wil, 1986Wil2]
* -4
cF96 Fd3m NiTi2
a = 1150
[1965Ram] most probably metastable
a)
When quenching from 1250°C NiAl transformed to a body centered tetragonal martensite with a = 261.0 and c = 337.6 pm [1979Nas].
b)
As a result of heavy cold work TaNi3 transforms to “TaNi3-cw” with the TiAl3 type, tI8, a = 362.7, c = 745.5 pm [1991Nas]. This form was obtained as phase by [1979Nas] when quenching Al-Ni-Ta alloys from 1250°C; the lattice parameters in these alloys varied from a = 357.1 to 364.8 pm and c = 741.9 to 748.7 pm. “TaNi3” with the TiCu3 type, oP8, a = 512.2, b = 452.2 and c = 423.5 pm was listed by [1991Nas] as a metastable phase due to surface contamination TaNi3Ox. This phase was observed as phase in Al-Ni-Ta alloys when quenched from 1000°C [1979Nas]; the lattice parameters in these alloys varied from a = 509.4 to 512 pm, b = 437 to 452.7 and c = 423 to 424.7 pm.
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Ta
Ni
Al
L + -1
~1550
e1 (min)
L -1
15.5 1.0 30.0
42.25 48.5 36.0
42.25 50.5 34.0
L -3 +
1387
e2 (max)
L -3
16 15 22
75 75 75
9 10 3
L ´ + -3
1372
e3 (max)
L ´ -3
11 10 13.5
75 75 75
14 15 11.5
L + ´ + -3
~1365
U1
L ´ -3
13 11.5 6 13.5
72.5 73.5 71 74
14.5 15 23 12.5
L + ´ + -3
~1360
U2
L ´ -3
11.5 10 11 13
78.5 78.5 84 76
10 11.5 5 11
L + -3 +
~1360
U3
L -3
14 14.5 7 22
71.5 71.5 70 74
14.5 14 23 4
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ta
434 T [°C]
Reaction
Type
L + -3 +
~1355
+ TaNi8 + -3
1330-1250 U4
Phase
E
Composition (at.%) Ta
Ni
Al
L -3
14.5 12.5 15 20
78.5 83.5 77.5 77
7 4 7.5 3
-
-
-
-
Table 3: Equilibrium Compositions of and ´ Phases and Ta Partition Coefficient [1994Jia] Temperature [°C]
(at.%)
´ (at.%)
Ta
Al
Ta
Al
Partition coefficient kTa / ´
0.09
19.9
0.25
22.9
2.78
0.21
19.4
0.56
22.3
2.67
0.86
18.4
1.52
20.8
1.77
1200
0.57
18.9
1.20
22.4
2.11
1100
0.23
15.9
0.68
19.9
2.96
0.40
15.4
1.23
19.2
3.08
0.45
14.7
1.05
20.2
2.33
0.61
13.0
2.01
20.7
3.30
0.64
10.4
1.79
14.4
2.80
1300
1000 800
Table 4: Equilibrium Compositions of ' and Phases and Ta Partition Coefficient [1994Jia] Temperature [°C]
´ (at.%)
(at.%)
Ta
Al
Ta
Al
Partition coefficient kTa ´/
1300
0.9
23.0
0.39
34.0
2.31
1100
0.78
23.3
0.34
34.0
2.29
1000
1.44
27.2
0.09
40.0
-
2.64
25.9
0.26
39.3
10.2
0.74
27.2
0.39
34.3
1.90
900
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
435
830
Fig. 1: Al-Ni-Ta. Lattice constants a (triangle) and c (circle) of the -1 phase
820
Lattice parameter, pm
810
800
790 505
500
495
490
485 40
50
30
20
10
Ni, at.%
Al-Ni
Al-Ni-Ta
Ni-Ta
ca.1550 e1min L β + τ1 1387 e2max L τ3 + δ
1372 e3max L γ´ + τ3
1372 p1 l + γ γ´ 1369 e4 l γ´ + β
ca.1365 L + γ´ β + τ3
L+β+τ3
γ´+β+τ3 ca.1360
L + γ´ γ + τ3
γ+γ´+τ3
U1
τ3+β+δ
L+γ+τ3 ca.1355
L + τ3 β + δ
ca.1360
U2
L γ + τ3 + δ
γ+TaNi8+τ3
1360 e5 lγ+δ
E
γ+τ3+δ <1330 γ + δ TaNi8 + τ3
L+β+δ
U3
U4
ca.1330 p2 γ + δ TaNi8
δ+τ3+TaNi8
Fig. 2: Al-Ni-Ta. Reaction scheme Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ta
436
Ta Ni Al
Fig. 3: Al-Ni-Ta. Partial liquidus surface
0.00 40.00 60.00
10
20
Data / Grid: at.% Axes: at.%
50
e1min
40
1500
50
U1
e3max
τ3 E
δ Ta Ni Al
60.00 40.00 0.00
50
60
70
Ta Ni Al
Fig. 4: Al-Ni-Ta. Partial solidus surface
20
γ'
xx maa ee 22m
15 00
β
80
U2
e5
0.00 70.00 30.00
10
11445500
13 140 75 0 U3 145 0
e4 p1
11440000
7755 1133
40
30
1450 1400
11442255
30
γ
90
Ni
Data / Grid: at.% Axes: at.%
10
1425° C
Maximum solid solubility
14 00 °C
20
20
γ
Ta Ni Al
MSIT®
30.00 70.00 0.00
80
90
1450°C
10
Ni
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
437
1500
Fig. 5: Al-Ni-Ta. Dependence of the phase liquidus (solid lines) and solidus (dash lines) temperatures vs Al content at constant Ta contents
1450
1 − Al-Ni system 2 − 2 at.% Ta 3 − 4 at.% Ta 4 − 6 at.% Ta 5 − 8 at.% Ta 6 − 10 at.% Ta
2
3
1
Temperature, °C
4 5 1400
6
Maximum solid solubility
1350
0
5
10
15
20
Al, at.%
Fig. 6: Al-Ni-Ta. Dependence of the phase liquidus (solid lines) and solidus (dash lines) temperatures vs Ta content at constant Al contents
1500
1 − Ni-Ta system 2 − 3 at.% Al 3 − 7 at.% Al 4 − 9 at.% Al 5 − 11 at.% Al 6 − 13 at.% Al 7 − 15 at.% Al 8 − 17 at.% Al
2
Temperature, °C
1450
1400
8
7
5 6
1 4
3
1350
0
5
10
15
20
Ta, at.%
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ta
438
Ta Ni Al
Fig. 7: Al-Ni-Ta. The ´/( + ´) solvus isotherms
0.00 75.00 25.00
Data / Grid: at.% Axes: at.%
20
10
132 7°C 122 7°C
γ +γ ´
11 10 27°C 2 7° 9 82 27°C C 7° C
10
20
γ
Ta Ni Al
80
25.00 75.00 0.00
90
Ni
Al
Data / Grid: at.%
Fig. 8: Al-Ni-Ta. Isothermal section at 1250°C
Axes: at.%
20
80
TaAl3
L Ta39Al69 40
60
Ta5Al7 TaAl
β
60
40
τ1 τ2
80
20
τ3
γ'
µ
Ta
MSIT®
20
40
γ 60
TaNi2
δ
80
TaNi8
Ni
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
439
Al Fig. 9: Al-Ni-Ta. Isothermal section at 1000°C
Data / Grid: at.% Axes: at.%
L 20
80
TaAl3 Ni2Al3
Ta2Al3 40 Ta5Al7
60
TaAl
β
60
40
τ1 τ2 γ'
80
20
τ3 γ
µ 20
Ta
40
60
TaNi2
δ
80
Ni
TaNi8
1600
Fig. 10: Al-Ni-Ta. Experimentally determined partial vertical section Ni3Al - TaNi3
1550°C 1500
Temperature, °C
L 1400
δ
τ3
γ' 1300
τ3+δ 1200
γ'+τ3 1100
1000
0.00 Ta Ni 75.00 Al 25.00
Landolt-Börnstein New Series IV/11A3
10
20
Ta, at.%
Ta 25.00 Ni 75.00 0.00 Al
MSIT ®
Al–Ni–Ta
425
Aluminium – Nickel – Tantalum Viktor Kuznetsov Literature Data Phase equilibria and intermetallic phase formation has been reviewed by [1990Kum]. However, this was followed by a thorough assessment of the data published up to 1991 by [1993Zak]. They presented graphically the Ni3Al-TaNi3 section, an assessed Scheil reaction scheme, liquidus and solidus projections, the solvus of the ´ (Ni3Al) based phase and two partial isothermal sections for 1000 and 1250°C. The existence of six ternary phases, TaNiAl, TaNi2Al, Ta0.5Ni3Al0.5, Ta5Ni2Al3, Ta~55 Ni~10 Al~35 and TaNiAl2 was accepted. However, some earlier work was not mentioned in [1993Zak]. [1965Ram] had indicated that in addition to the TaNiAl and TaNi2Al phases, which had been established initially by [1964Mar], a phase with a structure “closely resembling” that of NiTi2 (in Table 1 of [1965Ram] denoted as NiTi2) was present in alloys of gross compositions Ta25Ni38Al37 and Ta25Ni25Al50. The phase was found in both the as cast state and after annealing for 7 days at 900°C, but with an amount significantly less after the heat treatment. Moreover, it was found in the as cast sample with a composition of Ta25Ni50Al25, but later transformed almost entirely to TaNi2Al after annealing. Unfortunately, no compositional data for the phase was given. Later, [1974Ali] performed a DTA study of 5 alloys in the Ni3Al-TaNi3 section in the course of a study of the Ni3Al-Ni3Ta-Ni 3Nb pseudoternary system. In more recent years the phase equilibria in this system have been investigated in much detail. [1994Joh] studied five arc-melted alloys with compositions close to NiAl+15 at.% Ta (on the eutectic line) in the as cast and directionally solidified state by using scanning electron microscopy with EDS to measure phase composition. From the results, a fragment of the liquidus projection (for NiAl-Ni2TaAl-NiTaAl composition region) was constructed suggesting a peritectic formation for the Ni2AlTa ternary phase. [2001Miu] used DTA to determine liquidus and solidus temperatures of alloys made by arc-melting Al, Ni and Ta of purities 99.99, 99.95 and 99.9 mass%, respectively, followed by a homogenization treatment of 1000°C for 24 h. [1991Mis] determined the solvus line of the phase at temperatures between 827 and 1327°C using DTA. Energy-dispersive X-ray spectroscopy was used to confirm the phase constitution of the alloys. [1994Jia] studied the partition of Ta between and ´, as well as between the ´ and phases using a diffusion couple technique. The results are presented in tabular form with phase composition and partition coefficients for 1300, 1200, 1100, 1000 and 800°C and also rendered graphically as partial sections for some selected temperatures. [1996Pal] re-investigated two partial isothermal sections for Ta contents of < 50 at.% for 1000 and 1250°C in order to confirm the work of [1993Zak]. 32 compositions were prepared from components of purities of 99.95 mass% Ni, 99.99 mass% Al and 99.97 mass% Ta using levitation melting. Heat treatment at 1000°C was performed in Ar filled silica ampoules for 168 h for alloys in the NiAl+TaNiAl composition region and for 500 h for alloys of all other compositions. Water quenching followed the heat treatment. At 1250°C, the heat treatment was carried out in a box made from Ta sheets; each specimen was wrapped into Ta foil, and the box was filled with Ti-filings. The heat treatment was carried out in an Ar atmosphere for between 100 and 20 h with subsequent cooling under flowing gas. Samples were examined by metallography, X-ray diffraction and electron microprobe. The results show significant differences from the assessed data of [1993Zak]. [1999Sun] studied the partition of Al and Ta between the liquid and fcc phases in samples quenched from the two phase liquid + fcc state. The compositions of the phases were measured by EPMA. Equilibrium conditions were confirmed by the measure of homogeneity of the solid phase. In addition, they performed a simultaneous regression analysis of their own data, the published data of [1993Zak] and data for the Ni-Cr-Al-Ta quaternary. Good agreement (within approx. 1%, i.e. 7 to 10 K) between the different sets was found. Data for liquid compositions and partition coefficients for Al and Ta were tabulated. Very little work has been carried out on the thermodynamic properties in this system. [1999Roc] measured the low-temperature (3.2 to 10.3 K) heat capacity of the TaNi2Al phase and calculated its electron structure by the LMTO technique. Combining the results of both, the electron-phonon interaction constant was
Landolt-Börnstein New Series IV/11A3
MSIT ®
426
Al–Ni–Ta
derived. Some phase boundaries have been calculated using CALPHAD and ab initio techniques. [1991Kau] performed an approximate CALPHAD calculation of the phase equilibria. However, ternary phases were not taken into account, although dissolution of Al in the TaNi binary compound was allowed in the calculation. [1991Eno] calculated the equilibria between the and ´ phases at 1000°C, using the cluster variation method based on empirical Lennard-Jones type interatomic pair potentials. Good agreement with experimental data was obtained. A number of investigations of mechanical properties have been made. [1991Sas] noticed the precipitation of Ta enriched phase whilst studying the mechanical properties of (NiAl)0.95 Ta0.05. [1996Mac] measured lattice spacing and mechanical properties of the TaNiAl ternary phase. Mechanical properties were also studied by [1991Bon], [1991Hay], [1991Mas], [1991Sas]. Binary Systems The Ni-Ta system is taken from [Mas2], [1991Nas]. For the Al-Ni binary, the latest version [2003Sal] evaluated within the MSIT Binary Evaluation Program is accepted; it does not differ significantly from that of [1987Hil, 1988Bre], which was used by [1993Zak]. The Al-Ta system is taken from [2003Cor], who accepted results of the thermodynamic assessment of the system performed by [1996Du]. Solid Phases [1993Zak] accepted the existence of six ternary phases, TaNiAl, TaNi2Al, Ta0.5Ni3Al0.5, Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2. The TaNiAl phase has a wide solubility range for Al (11 to 50 at.%), but restricted for Ta (32.5 to 37.5 at.%). [1996Pal] noted, that in comparing calculated and observed intensities of X-ray diffractions lines, the suggestion is that Al substitutes for Ni on two different crystallographic sites which exist in the MgZn2 structure to a similar extent. The lattice constants of that phase seem to depend on cooling rate; the reason for this is unclear, but because no peak broadening was observed, it is not likely to be due to stacking faults introduced by thermal stresses on cooling [1996Pal]. The true composition of the TaNi2Al phase was found to be off-stoichiometric: 51 to 55% Al and 22.5 to 25% Ta at 1000°C; 52 to 58% Al and 17.5 to 24% Ta at 1250°C [1996Pal]. [1996Pal] did not find any trace of the Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2 phases as presented by [1993Zak], nor the NiTi2 phase reported by [1965Ram]. The existence of the first three was explicitly rejected; the latter was not considered anyhow by [1996Pal], but no such phase was detected in the composition range studied by [1965Ram]. As [1993Zak] noted weak support for the existence of Ta5Ni2Al3, Ta~55Ni~10Al~35 and TaNiAl2, these phases are considered in present review to be non-existent. The ternary phase proposed by [1965Ram] seems to be metastable; also, phases with that structure are often stabilized by impurities such as C, N or O. Crystallographic information for the solid phases, including the probably metastable ternary phase, is summarized in Table 1. Detailed data for the concentration dependence of the lattice spacing of TaNiAl [1996Pal] are given on Fig. 1. For 20 < xNi < 50 (xNi in at.%) that dependence is essentially linear: a (pm) = 487.6+0.413(50-xNi), c (pm) = 791.5+0.476(50-xNi), though marked deviations from that can be seen for less Ni [1996Pal]. For lattice spacing of the phase, linearity holds for all compositions studied: a (pm) = 487.5+0.386(50-xNi), c (pm) = 2653+2.90(50-xNi) (also for 20 < xNi < 50 at.%). Pseudobinary Systems No pseudobinary sections have been found in the system, though some authors suggested such behavior for the Ni3Al-TaNi3 section, see section “Temperature – Composition Sections”. Invariant Equilibria Data for the invariant equilibria and Scheil reaction scheme (Fig. 2) were assessed by [1993Zak], and are accepted here with some alterations. Table 2 is based on [1993Zak], but with a corrected error in the temperature of the U4 reaction, noted by [1996Pal]. The reaction U5, presented by [1993Zak] has been omitted as it was shown to be unlikely by [1996Pal]. The eutectic e1(min) L+TaNiAl, is added from MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
427
[1994Joh]. It is not possible to include the ternary peritectic reaction +TaNiAl+LTaNi2Al and surrounding univariant eutectic L+TaNiAl and peritectics +LTaNi2Al and TaNiAl+LTaNi2Al suggested by [1994Joh] in either Table 2 or the reaction scheme as neither temperatures nor phase compositions were determined. (See however discussion of liquidus below). Liquidus, Solidus and Solvus Surfaces The liquidus data from [2001Miu] are in good agreement with [1993Zak] for both edge systems, but differ markedly for intermediate compositions. The data of [2001Miu] are preferred as they result from detailed work and seem to be more reliable. On the other hand, [2001Miu] presents mono- and invariant equilibria lines taken from [1986Wil] which is the main source for [1993Zak]. The liquidus surface is presented here in Fig. 3. It is a composite of the liquidus taken from [2001Miu] and the liquidus surfaces of ´, -3 and phases taken from [1993Zak]. The partial liquidus projection from [1994Joh] is added tentatively, although its connection with other parts of liquidus surface remains rather unclear. Figure 4 provides isotherms of the solidus from [2001Miu]. Figures 5 and 6 present the data of [2001Miu] showing the dependence of the liquidus and solidus temperatures on Al variation at parametric Ta content, and on Ta variation at parametric Al content, respectively. These data give more detailed representation than is possible in Figs. 3 and 4. Figure 7 presents the isotherms of the /( + ´) solvus surface as determined by [1991Mis]. Isothermal Sections Isothermal sections at 1273 and 1000°C are presented in Figs. 8 and 9, respectively, generally accepted from [1996Pal]. The results differ significantly from those of the earlier assessment of [1993Zak]. On the other hand, the data given in the original work disagree with the accepted Al-Ta binary system (and even with the binary accepted by the authors [1996Pal] themselves). To maintain consistency, it was necessary to change the region adjacent to Al-Ta system, which in any event is based on just two alloys. In particular, the homogeneity range of TaAl3 phase is removed, and that of Ta2Al3 is split into stoichiometric Ta5Al7 and Ta39Al69 phases at 1250°C (Fig. 8) and into Ta5Al7 and Ta2Al3 at 1000°C (Fig. 9). These changes were suggested by [1996Du] who analyzed the results of [1996Pal] during their assessment of the Al-Ta binary and is accepted here. Also, the position of the phase corners of +TaNiAl+TaAl3 and L+TaNiAl+TaAl3 tie triangles had to be shifted somewhat to make them compatible with the accepted version of the Al-Ni binary. The data of [1994Jia] for - ´ and ´- equilibria, presented in tabular form, are reproduced in Tables 3 and 4. Temperature – Composition Sections [1993Zak] suggested the section Ni3Al-Ni3Ta to be “partly pseudobinary” and mentioned some experiments on directional growth of a “pseudobinary eutectic” [1972Hub, 1974Mol]; the reported composition of the latter is indeed in very good agreement with the composition of the e2 reaction of [1993Zak]. The DTA study of [1974Ali] is also in agreement, though the authors themselves interpreted their results as indication of a simple pseudobinary section with a single eutectic. As indicated by [1993Zak], the Ni3Al-Ni3Ta section cannot be pseudobinary due to the incongruent formation of Ni3Al. Moreover, in the presently accepted version of the Al-Ni binary, the Ni3Al phase becomes off-stoichiometric starting from approx. 1347°C up to the melting point [1987Hil, 1988Bre]. Also, the phase boundaries of ternary TaNi2Al phase as determined for 1000°C by [1996Pal] are not crossed by the Ni3Al-Ni3Ta join. No account of these phenomena was taken by [1993Zak]. On the other hand, the assessed liquidus-solidus region of that section is indeed independently confirmed by the results of [1974Ali] and by directional solidification experiments, reported by [1993Zak]. So, this fragmentary section is reproduced from [1993Zak] with minor corrections and given as Fig. 10, though the true phase relations should be much more complicated both in a region closer to the Ni3Al side and at lower temperatures.
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ta
428 Thermodynamics
No thermodynamic studies have been carried out except for low-temperature (3.2 to 10.3 K) measurements of the heat capacity of the TaNi2Al phase performed by [1999Dar]. Their results, when treated in the standard way (Cp(T) = elT + CD(/T)), give el = 10.01 0.14 mJ#mol#K-2, D = 299 1.9 K. This equation is valid only below approximately 7 K. Notes on Materials Properties and Applications The influence of Ta additions on mechanical properties of NiAl was studied in [1991Mas, 1991Sas]. Such properties of Ta alloyed single crystals of ´ Ni 3Al can be found in [1991Bon]; creep behavior of that phase was studied by [1991Hay]. Some mechanical properties of the Laves phase TaNiAl were measured by [1996Mac]. Miscellaneous [2001Ter] suggested the usage of thermal conductivity measurements for determination of site preferences in the ´ Ni3Al phase. The results are in broad agreement with the phase diagram determinations, which suggest that Ta substitutes for Al in Ni3Al. [2001Kai] investigated the morphological stability of the interface between ´(L1 2) and (B2) phases in diffusion couples. In addition, the results of an unpublished calculation of thermodynamic properties are cited and used in the discussion of the results. References [1964Mar]
[1965Gie] [1965Ram]
[1968Hun]
[1972Hub]
[1972Min]
[1974Ali]
[1974Mol]
[1974Var]
MSIT®
Markiv, V.Ya., Voroshilov, Yu.V., Kripyakevich, P.I., Cherkashin, E.E., “New Compounds of the MnCu 2Al and MgZn2 Types Containing Aluminium and Gallium”, Sov. Phys.-Crystallogr., 9, 619-620 (1964), translated from Kristallografiya, 9, 737-738 (1964) (Crys. Structure, 4) Giessen, B.C., Grant, N.J., “New Intermediate Phases in Transition Metal Systems. II”, Acta Crystallogr., 18, 99 (1965) (Crys. Structure, 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) (Equi. Diagram, Experimental, 14) Hunt, C.R., Raman, A., “Alloy Chemistry of )(U)-Related Phases. I. Extens Ion of - and Occurrence of ´-Phases in the Ternary Systems Nb(Ta)-X-Al (X = Fe, Co, Ni, Cu, Cr, Mo)”, Z. Metallkd., 59(9), 701-707 (1968) (Crys. Structure, Equi. Diagram, 14) Hubert, J.-C., Kurz, W., Lux, B., “Growth by Directed Solidification of the Ni3Al-Ni3Ta Quasibinary Eutectic” (in French), J. Cryst. Growth, 13-14, 757-764 (1972) (Equi. Diagram, 15) Mints, R.S., D´yakonova, N.P., Umansky, Ya.S., Bondarenko, Yu.A., Bondarenko, T.A., “Interaction of the Phase Ni 3Al with Ni3Ta”, Sov. Physics Doklady, 17(9), 904-906 (1973) translated from Dokl. Akad. Nauk SSSR, 206(1), 87-88 (1972) (Crys. Structure, Experimental, 5) Alikhanov, V.A., Pyatnitskii, V.N., Sokolovskaya, E.M., “Phase Diagram of the System Ni3Al-Ni3Nb-Ni3Ta” (in Russian), Vestn. Mosk. Univ., Ser. 2:Khim., 15, 698-701 (1974) (Equi. Diagram, Experimental, 5) Mollard, F., Lux, B., Hubert, J.C., “Directionally Solidified Composites Based on the Ternary Eutectic Ni-Ni3Al-Ni3Ta ( / ´ - )”, Z. Metallkd., 65, 461-468 (1974) (Equi. Diagram, Experimental, 6) Varli, K.V., D’yakonova, N.P., Umansky, Ya.S., Bondarenko, Yu.A., Putman, A.M., “Crystal Structure of the Ternary Phase of the Ni-Ta-Al System”, Vses. Konf. Kristallokhim. Intermet., 2nd, Tezisy Dokl., Lvov Gos. Univ.: Lvov, USSR, 49 (1974) (Crys. Structure, 0)
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta [1979Nas] [1984Och]
[1984Wil]
[1985Mis]
[1986Hua] [1986Wil1]
[1986Wil2]
[1987Hil]
[1987Kha] [1988Bre]
[1990Kum]
[1991Bon]
[1991Eno]
[1991Hay]
[1991Kau] [1991Mas]
[1991Mis]
[1991Nas]
Landolt-Börnstein New Series IV/11A3
429
Nash, P., West, D.T.F., “Phase Equilibria in the Ni-Ta-Al System”, Met. Sci., 13(12), 670-676 (1979) (Equi. Diagram, Crys. Structure, Experimental, 22) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni ( ), Ni3Al ( ´) and Ni3Ga ( ´) Solid Solutions”, Bull. P. M. E.,(T. I. T.), 53, 15-28 (1984) (Crys. Structure, Experimental, Rewiew, 56) Willemin, P., Dugue, O., Durand-Charre, M., Davidson, J., “High-Temperature Phase Equilibria in the Ni-Al-Ta System”, Superall. 1984 Champ., MS/AIME, Conf: Pa. USA, 637-647 (1984) (Equi. Diagram, Crys. Structure, Experimental, 13) Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni( ), Ni3Al( ´) and Ni3Ga( ´) Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33(6), 1161-1169 (1985) (Crys. Structure, Review, 64) 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) Willemin, P., Dugue, O., Durand-Charre, M. J., Davidson, H., “Experimental Determination of Nickel-Rich Corner of Ni-Al-Ta Phase Diagram”, Mater. Sci. Technol., 2(4), 344-348 (1986) (Equi. Diagram, 13) Willemin, P., Durand-Charre,, M., Ansara, I., “Liquid-Solid Equilibria in the System Ni3Al-Ni3Ta and Ni3Al-Ni3Ti”, High Temp. Alloys Cas Turbines Other Appl., Pt.2, Comm. Euro. Communicates, Rep. EUR 10567, 955-964 (1986) (Equi. Diagram, Thermodyn., 8) Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremer, F.T., Nickel, H., “Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch., 42a, 1327-1332 (1987) (Equi. Diagram, Experimental, 17) Khadkikar, P.S., Vedula, K., “An Investigation of the Ni5Al3 Phase”, J. Mater. Res., 2(2), 163-167 (1987) (Crys. Structure, Experimental, 7) Bremer, F.J., Beyss, M., Karthaus, E., Hellwig, A., Schober, T., Welter, J.-M., Wenzl, H., “Experimental Analysis of the Ni-Al Phase Diagram”, J. Cryst. Growth, 87, 185-192 (1988) (Equi. Diagram, Experimental, 16) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35, 293-327 (1990) (Crys. Structure, Equi. Diagram, Review, 158) Bonneville, J., Martin, J.L., “The Strain Rate Sensitivity of Ni3(Al,Ta) Single Crystals”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc., 213, 629-634 (1991) (Mechan. Prop., Experimental, 17) Enomoto, M., Harada, H., Yamazaki, M., “Calculation of ´/ Equilibrium Phase Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15, 143-158 (1991) (Equi. Diagram, Calculation, 34) Hayashi, T., Shinoda, T., Mishima, Y., Suzuki, T., “Effect of Off-Stoichiometry on the Creep Behavior of Binary And Ternary Ni3Al”, High-Temp. Ordered Intermetallic Alloys IV, Mater. Res. Soc. Symp. Proc., 213, 617-622 (1991) (Mechan. Prop., Experimental, 7) Kaufman, L., “Calculation of the Multicomponent Tantalum Based Phase Diagrams”, Calphad, 15, 261-282 (1991) (Equi. Diagram, Calculation, 15) Maslenkov, S.B., Filin, S.A., Abramov, V.O., “Effect of Structural State and Alloying of Transition Metals on the Degree of Hardening of Ternary Solid Solutions Based on Nickel Monoaluminide”, Russ. Metall. (Engl. Transl.), (1), 115-118 (1991) (Mechan. Prop., Experimental, 10) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Equi. Diagram, Experimental, 5) Nash, A., Nash, P., “The Ni-Ta (Nickel-Tantalum) System”, in “Phase Diagrams of Binary Nickel Alloys, Monograph Series on Alloy Phase Diagrams”,Vol. 6, ASM-Intl., Materials Park, Ohio, 320-325 (1991) (Equi. Diagram, Crys. Structure, Review, 38)
MSIT ®
430 [1991Sas]
[1991Zha]
[1993Kha]
[1993Zak]
[1994Jia]
[1994Joh] [1996Du]
[1996Mac] [1996Pal] [1996Pau]
[1999Roc]
[1999Sun] [2001Kai]
[2001Miu]
[2001Ter]
[2003Cor]
[2003Sal]
MSIT®
Al–Ni–Ta Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on the Solidified Structure of NiAl”, Intermetal. Comp. - Struct. Mechan. Prop., Proc. Conf., 877-881 (1991) (Equi. Diagram, Mechan. Prop., Abstract, 10) Zhao, J.T., Celato, L., Parthe, E., “Structure Refinement of Monoclinic 12-Layer TaNi3 with -NbPt3 Type. New Crystallographic Descriptions of this Type and of the Nb3Rh 5 Type Based on Smaller Unit Cells”, Acta Crystallogr., Sect. C: Crys. Struct. Commun., C47, 479-483 (1991) (Crys. Structure, Experimental, 11) Khadkikar, P.S., Locci, I.E., Vedula, K., Michal, G.M., “Transformation to Ni5Al3 in a 63.0 at.% Ni-Al Alloy”, Metall. Trans. A, 24A, 83-94 (1993) (Equi. Diagram, Crys. Structure, Experimental, 28) Zakharov, A., “Aluminium - Nickel - Tantalum”, in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.14883.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 28) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1), `(L12) and (B2) Phases in Ni-Al Base Systems”, Metall. Mater. Trans. A, A25, 473-485 (1994) (Equi. Diagram, Experimental, 25) Johnson, D.R., Oliver, B.F., “Ternary Peritectic Solidification in the NiAl-Ni2AlTa-NiAlTa System”, Mater. Lett., 20, 129-133 (1994) (Equi. Diagram, Experimental, 11) Du, Y., Schmid-Fetzer, R., “Thermodynamic Modelling of the Al-Ta System”, J. Phase Equilib., 17, 311-324 (1996) (Equi. Diagram, Crys. Structure, Thermodyn., Assessment, Calculation, 55) Machon, L., Sauthoff, G., “Deformation Behavior of Al-Containing C14 Laves Phase Alloys”, Intermetallics, 4, 469-481 (1996) (Crys. Structure, Experimental, 41) Palm, M., Sanders, W., Sauthoff, G., “Phase Equilibria in the Ni-Al-Ta System”, Z. Metallkd., 87, 390-398 (1996) (Equi. Diagram, Crys. Structure, Experimental, 27) 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) 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), B269, 154-162 (1999) (Thermodyn., Phys. Prop., Experimental, Calculation, 20) Sung, P.K., Poirier, D.R., “Liquid-Solid Partition Ratios in Nickel-Base Alloys”, Metall. Mater. Trans. A, A30, 2173-2181 (1999) (Equi. Diagram, Experimental, 41) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of ´/ Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, A312, 168-175 (2001) (Kinetics, Thermodyn., Experimental, 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) (Equi. Diagram, Experimental, 9) Terada, Y., Ohkubo, K., Mohri, T., Suzuki, T., “Site Preference Determination in Intermetallic Compounds by Thermal Conductivity Measurements”, J. Mater. Res., 16, 2314-2320 (2001) (Thermal Conduct., Crys. Structure, Experimental, Calculation, 63) Cornish, L., Dolotko, O., Rogl, P., “Al-Ta (Aluminium - Tantalum)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram, Assessment, 3) 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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164) Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
431
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Al)
hP2 P63/mmc Mg
a = 269.3 c = 439.8
at 25°C, 20.5 GPa [Mas2]
(Al) < 660.452
cF4 Fm3m Cu
a = 404.96
at 25°C [Mas2]
(Ta) < 3020
cI2 Im3m W
a = 330.30
at 25°C [Mas2]
, (Ni) < 455
cF4 Fm3m Cu
a = 352.40
at 25°C [Mas2]
a = 357.8
at 8.% Ta, 1150°C, linear da/dx [1984Och2, 1985Mis] at 14.% Al, linear da/dx [1984Och2]
TaxNi1-x Ni1-xAlx TaxNi1-x-yAly
a = 355.0
a = 355.3 to 357.5 at 7.5 - 10 at.% Ta, 75 - 80 at.% Ni, 1000°C, quenched, sample contained ´ and -3 [1979Nas] a = 359.3 at 10 at.% Ta, 80 at.% Ni, 1250°C, quenched, sample contained -3 [1979Nas] ´, Ni3Al < 372
cP4 Pm3m AuCu3
[1986Hua] at 63 at.% Ni [1993Kha]
a = 359.0 to 362.4 at 3 - 12 at.% Ta, 58.6 - 80 at.% Ni, 1000 - 1200°C, multiphase samples quenched, linear da/dx [1972Min, 1979Nas, 1984Och2, 1985Mis]
Ta1-xNi3Alx
Ni5Al3 723
oC16 Cmmm Pt5Ga3
, NiAl < 1638 Tax(Ni1-yAly)1-x
cP2 a) Pm3m CsCl
TaNi8 < 307 (Ta1-xAlx)Ni8
tI36 NbNi8
Landolt-Börnstein New Series IV/11A3
a = 356.77 a = 358.9
a = 753 b = 661 c = 376 a = 286.0 a = 287 a = 288.72 0.02 a = 287.98 0.02 a = 286.6 to 296.8
a = 760.5 c = 358.5 a = 767 c = 348
32 to 36 at.% Al at 63 at.% Ni [1993Kha] 42 to 69.2 at.% Ni [Mas2] [1987Kha] at 63 at.% Ni [1993Kha] at 50 at.% Ni [1996Pau] at 54 at.% Ni [1996Pau] 3.0 - 20 at.% Ta, 50 - 70 at.% Ni. Quenched from 1250 - 1000°C. Samples were multiphase. [1979Nas] at 11.1 at.% Ta [1991Nas] at 11.8 at.%, 83.3 at.% Ni, from EMPA, 1250°C, quenched, [1979Nas]
MSIT ®
Al–Ni–Ta
432 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, TaNi 3 < 1547
mP48 b) P21/m TaPt3
a = 452.3 b = 512.6 c = 2544 = 90°
TaNi2 < 1404
tI6 I4/mmm MoSi2
, TaNi < 1570
hR13 R3m W6Fe7
Ta(Ni,Al)
a = 315.4 c = 790.5
at 22.5 to 28.5 at.% Ta [1991Nas] [1991Zha] single crystal
32.5 to 35 at.% Ta [1991Nas] at 33.3 at.% Ta [1991Nas]
50 to 54 at.% Ta [1991Nas] a = 492.1 at 50 at.% Ta [1991Nas] c = 2690.5 a = 491.9 to 497.8 50 - 55 at.% Ta, 35 - 23 at.% Ni c = 2714 to 2735 [1968Hun] a = 496.1 c = 2504
9 at.% Ta, 58.8 at.% Ni 1250°C, quenched, alloy with , , -2 [1979Nas]
a = 428.3 c = 2649
20 at.% Ta, 50 at.% Ni 1250°C, quenched, alloy with -11 [1979Nas]
a = 986.4 c = 521.5
at ~20 to 40 at.% Al [Mas2, V-C]
), Ta2Al < 2061
tP30 P42/mnm )CrFe
TaAl < 1446
mP*
[1996Du]
Ta5Al7 < 1345
hP*
[1996Du]
Ta2Al3 < 1226
cF*
[1996Du]
Ta39Al69 1548 - 1183
cF432 F43m
[1996Du]
TaAl3 < 1608
tI8 I4/mmm TiAl3
a = 383.7 c = 855.0
[V-C]
* -1, TaNiAl
hP12 P63/mmc MgZn2
a = 496.9 c = 798.5 a = 501.5 c = 817.1
[V-C] alloy 20 at.% Ta, 50 at.% Ni 1000°C, quenched, alloy with and -2 [1979Nas]
* -2, TaNi2Al
cF16 Fm3m BiF3
a = 594.9 a = 580 to 594
[V-C] 9 - 20 at.% Ta, 50 - 58.8 at.% Ni, 1000 - 1250°C, quenched. Multiphase samples [1979Nas]
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
433
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
* -3, Ta0.5Ni3Al0.5 < 1393
hP16 P63/mmc TiNi3
a = 510.5 to 513.7 Al-rich [1965Gie, 1972Min, 1974Var, c = 831.9 to 836.6 1979Nas, V-C, 1984Wil, 1986Wil2]
* -4
cF96 Fd3m NiTi2
a = 1150
[1965Ram] most probably metastable
a)
When quenching from 1250°C NiAl transformed to a body centered tetragonal martensite with a = 261.0 and c = 337.6 pm [1979Nas].
b)
As a result of heavy cold work TaNi3 transforms to “TaNi3-cw” with the TiAl3 type, tI8, a = 362.7, c = 745.5 pm [1991Nas]. This form was obtained as phase by [1979Nas] when quenching Al-Ni-Ta alloys from 1250°C; the lattice parameters in these alloys varied from a = 357.1 to 364.8 pm and c = 741.9 to 748.7 pm. “TaNi3” with the TiCu3 type, oP8, a = 512.2, b = 452.2 and c = 423.5 pm was listed by [1991Nas] as a metastable phase due to surface contamination TaNi3Ox. This phase was observed as phase in Al-Ni-Ta alloys when quenched from 1000°C [1979Nas]; the lattice parameters in these alloys varied from a = 509.4 to 512 pm, b = 437 to 452.7 and c = 423 to 424.7 pm.
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Ta
Ni
Al
L + -1
~1550
e1 (min)
L -1
15.5 1.0 30.0
42.25 48.5 36.0
42.25 50.5 34.0
L -3 +
1387
e2 (max)
L -3
16 15 22
75 75 75
9 10 3
L ´ + -3
1372
e3 (max)
L ´ -3
11 10 13.5
75 75 75
14 15 11.5
L + ´ + -3
~1365
U1
L ´ -3
13 11.5 6 13.5
72.5 73.5 71 74
14.5 15 23 12.5
L + ´ + -3
~1360
U2
L ´ -3
11.5 10 11 13
78.5 78.5 84 76
10 11.5 5 11
L + -3 +
~1360
U3
L -3
14 14.5 7 22
71.5 71.5 70 74
14.5 14 23 4
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ta
434 T [°C]
Reaction
Type
L + -3 +
~1355
+ TaNi8 + -3
1330-1250 U4
Phase
E
Composition (at.%) Ta
Ni
Al
L -3
14.5 12.5 15 20
78.5 83.5 77.5 77
7 4 7.5 3
-
-
-
-
Table 3: Equilibrium Compositions of and ´ Phases and Ta Partition Coefficient [1994Jia] Temperature [°C]
(at.%)
´ (at.%)
Ta
Al
Ta
Al
Partition coefficient kTa / ´
0.09
19.9
0.25
22.9
2.78
0.21
19.4
0.56
22.3
2.67
0.86
18.4
1.52
20.8
1.77
1200
0.57
18.9
1.20
22.4
2.11
1100
0.23
15.9
0.68
19.9
2.96
0.40
15.4
1.23
19.2
3.08
0.45
14.7
1.05
20.2
2.33
0.61
13.0
2.01
20.7
3.30
0.64
10.4
1.79
14.4
2.80
1300
1000 800
Table 4: Equilibrium Compositions of ' and Phases and Ta Partition Coefficient [1994Jia] Temperature [°C]
´ (at.%)
(at.%)
Ta
Al
Ta
Al
Partition coefficient kTa ´/
1300
0.9
23.0
0.39
34.0
2.31
1100
0.78
23.3
0.34
34.0
2.29
1000
1.44
27.2
0.09
40.0
-
2.64
25.9
0.26
39.3
10.2
0.74
27.2
0.39
34.3
1.90
900
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
435
830
Fig. 1: Al-Ni-Ta. Lattice constants a (triangle) and c (circle) of the -1 phase
820
Lattice parameter, pm
810
800
790 505
500
495
490
485 40
50
30
20
10
Ni, at.%
Al-Ni
Al-Ni-Ta
Ni-Ta
ca.1550 e1min L β + τ1 1387 e2max L τ3 + δ
1372 e3max L γ´ + τ3
1372 p1 l + γ γ´ 1369 e4 l γ´ + β
ca.1365 L + γ´ β + τ3
L+β+τ3
γ´+β+τ3 ca.1360
L + γ´ γ + τ3
γ+γ´+τ3
U1
τ3+β+δ
L+γ+τ3 ca.1355
L + τ3 β + δ
ca.1360
U2
L γ + τ3 + δ
γ+TaNi8+τ3
1360 e5 lγ+δ
E
γ+τ3+δ <1330 γ + δ TaNi8 + τ3
L+β+δ
U3
U4
ca.1330 p2 γ + δ TaNi8
δ+τ3+TaNi8
Fig. 2: Al-Ni-Ta. Reaction scheme Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ta
436
Ta Ni Al
Fig. 3: Al-Ni-Ta. Partial liquidus surface
0.00 40.00 60.00
10
20
Data / Grid: at.% Axes: at.%
50
e1min
40
1500
50
U1
e3max
τ3 E
δ Ta Ni Al
60.00 40.00 0.00
50
60
70
Ta Ni Al
Fig. 4: Al-Ni-Ta. Partial solidus surface
20
γ'
xx maa ee 22m
15 00
β
80
U2
e5
0.00 70.00 30.00
10
11445500
13 140 75 0 U3 145 0
e4 p1
11440000
7755 1133
40
30
1450 1400
11442255
30
γ
90
Ni
Data / Grid: at.% Axes: at.%
10
1425° C
Maximum solid solubility
14 00 °C
20
20
γ
Ta Ni Al
MSIT®
30.00 70.00 0.00
80
90
1450°C
10
Ni
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
437
1500
Fig. 5: Al-Ni-Ta. Dependence of the phase liquidus (solid lines) and solidus (dash lines) temperatures vs Al content at constant Ta contents
1450
1 − Al-Ni system 2 − 2 at.% Ta 3 − 4 at.% Ta 4 − 6 at.% Ta 5 − 8 at.% Ta 6 − 10 at.% Ta
2
3
1
Temperature, °C
4 5 1400
6
Maximum solid solubility
1350
0
5
10
15
20
Al, at.%
Fig. 6: Al-Ni-Ta. Dependence of the phase liquidus (solid lines) and solidus (dash lines) temperatures vs Ta content at constant Al contents
1500
1 − Ni-Ta system 2 − 3 at.% Al 3 − 7 at.% Al 4 − 9 at.% Al 5 − 11 at.% Al 6 − 13 at.% Al 7 − 15 at.% Al 8 − 17 at.% Al
2
Temperature, °C
1450
1400
8
7
5 6
1 4
3
1350
0
5
10
15
20
Ta, at.%
Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Ta
438
Ta Ni Al
Fig. 7: Al-Ni-Ta. The ´/( + ´) solvus isotherms
0.00 75.00 25.00
Data / Grid: at.% Axes: at.%
20
10
132 7°C 122 7°C
γ +γ ´
11 10 27°C 2 7° 9 82 27°C C 7° C
10
20
γ
Ta Ni Al
80
25.00 75.00 0.00
90
Ni
Al
Data / Grid: at.%
Fig. 8: Al-Ni-Ta. Isothermal section at 1250°C
Axes: at.%
20
80
TaAl3
L Ta39Al69 40
60
Ta5Al7 TaAl
β
60
40
τ1 τ2
80
20
τ3
γ'
µ
Ta
MSIT®
20
40
γ 60
TaNi2
δ
80
TaNi8
Ni
Landolt-Börnstein New Series IV/11A3
Al–Ni–Ta
439
Al Fig. 9: Al-Ni-Ta. Isothermal section at 1000°C
Data / Grid: at.% Axes: at.%
L 20
80
TaAl3 Ni2Al3
Ta2Al3 40 Ta5Al7
60
TaAl
β
60
40
τ1 τ2 γ'
80
20
τ3 γ
µ 20
Ta
40
60
TaNi2
δ
80
Ni
TaNi8
1600
Fig. 10: Al-Ni-Ta. Experimentally determined partial vertical section Ni3Al - TaNi3
1550°C 1500
Temperature, °C
L 1400
δ
τ3
γ' 1300
τ3+δ 1200
γ'+τ3 1100
1000
0.00 Ta Ni 75.00 Al 25.00
Landolt-Börnstein New Series IV/11A3
10
20
Ta, at.%
Ta 25.00 Ni 75.00 0.00 Al
MSIT ®
440
Al–Ni–W
Aluminium – Nickel – Tungsten Konstyantyn Korniyenko, Vasyl Kublii, Olga Fabrichnaya, Natalia Bochvar Literature Data Experimental investigations of the phase equilibria in the Al-Ni-W system are limited to the Ni-rich range of compositions [1958Bud, 1978Gul, 1983Nas, 1986Nov, 1986Udo, 1991Udo, 1994Jia]. Solid solubility of tungsten in the Ni3Al ( ´) phase is presented in [1966Arb, 1983Och, 1984Och1, 1984Och2, 1985Mis, 1985Nas] and the , (Ni) solvus surface is described in [1989Hon1, 1989Hon2] and [1991Mis]. [1991Sas] studied the alloying effect of tungsten, on the solidification of the NiAl () phase. [1958Bud] investigated about 60 ternary alloys, that were prepared by high-frequency melting using corundum crucibles under a protective layer of basic slag. The starting components were A-000 aluminium, nickel (99.9 mass%) and tungsten (99.98 mass%). After subsequent stepwise annealing in high vacuum at 1200°C for 24 h, at 1000°C for an additional 100 h and at 800°C for another 100 h, respectively, each batch of alloys was partly quenched and cooled to room temperature. All losses were less than 0.2 to 0.5 mass% due to the fact that the Al was introduced by means of premelted Al-Ni master alloys. Three isothermal sections of the Al-Ni-W system in the Ni rich range at 1200, 1000 and 800°C and part of the NiAl-W pseudobinary section are plotted. Since then the Ni-NiAl-W partial system has been reinvestigated frequently and partial isothermal sections have been established by various research groups: at 1370°C [1986Udo], 1250°C [1983Nas], 1200°C [1986Udo, 1987Pri], 1150°C [1991Udo], 900°C [1991Udo] and for the temperature range from the end of alloy solidification up to the beginning of solid state reactions, i.e. from ~1350 to 1000°C [1986Nov]. [1978Gul] recorded the liquidus and solidus temperatures in the Ni-rich corner by means of calibrated thermocouples. [1986Nov] presented also the liquidus surface projection and reaction scheme for the partial Ni-NiAl-W system. In most cases the alloys were prepared by arc melting of the elements with 99.9 mass% minimal purity on a water-cooled copper hearth under an argon atmosphere using a nonconsumable tungsten electrode [1966Arb, 1983Nas, 1983Och, 1984Och1, 1984Och2, 1985Mis, 1985Nas, 1986Nov, 1986Udo, 1987Pri, 1989Hon1, 1989Hon2, 1991Mis, 1991Sas, 1991Udo], by induction melting under an argon atmosphere [1994Jia] as well as by high-frequency melting [1986Nov, 1991Udo]. Due to rather sluggish reaction kinetics, it was difficult to obtain equilibrium and homogenization treatments up to 500 h at various temperatures were carried out. After heat treatment the samples usually were quenched to room temperature. Methods of experimental investigation of the alloys were: X-ray diffraction [1966Arb, 1983Nas, 1983Och, 1984Och1, 1984Och2, 1985Mis, 1986Udo, 1986Nov, 1987Pri], metallography [1983Nas, 1986Nov, 1986Udo, 1987Pri, 1991Sas, 1991Udo, 1994Jia], electron microprobe analysis (EMPA) [1983Nas, 1986Nov, 1987Pri, 1989Hon1, 1989Hon2, 1991Mis, 1991Udo, 1994Jia], differential thermal analysis (DTA) [1986Nov, 1989Hon1, 1989Hon2, 1991Mis] and [1991Udo] as well as microhardness measurements [1987Pri]. The critical review of literature data on the phase equilibria in the system was carried out by [1993Ale] within the MSIT Evaluation Program and is continued and updated by the present evaluation. Binary Systems The descriptions of the Al-Ni, Al-W and Ni-W systems are accepted from [2003Sal], [2003Sch] and [Mas2], respectively. Solid Phases No ternary phases have been found. Crystallographic data on the known unary and binary phases are listed in Table 1. Based on earlier investigations of [1966Arb], the extent of the ´ phase field on alloying with tungsten as well as the mode of atom substitution has been studied by [1983Och, 1984Och1, 1984Och2] and [1985Mis] who arrived at a maximum solubility of less than 5 at.% W at 1000°C with W replacing Al
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–W
441
in WxNi3Al1-x [1983Och, 1984Och1, 1984Och2, 1985Mis], whereas in general W was found to substitute for both lattice sites in Wx(W yNi1-y)3Al1-x [1966Arb]. Data on solubility of W in the ´ phase are contradictory and need more accurate definition. So, [1973Mul] and [2001Sav] prepared single crystals WxNi3Al1-x with 3 at.% W replacing Al. A slightly larger solubility of 6 at.% W in ´ at 1000°C was obtained by [1983Nas] on the basis of EMPA data, and the solubility of W in the ´ phase was found to decrease from ~6 to ~4.5 at.% as the temperature increases from 1000 to 1250°C [1983Nas, 1985Nas]. But [1987Pri] obtained the solubility of W in the ´ phase ~2 at.% at 1200°C, while [1991Udo] reported ~4 at.% W at this temperature (Table 1). Similarly the solubility of W in the phase was shown to be ~0.2 at.% at 1250°C [1983Nas, 1985Nas], whereas [1958Bud] reported a value of 10 at.% at 1600°C, and 6 at.% at 1500°C. The value of W solubility of ~2 at.% at 1250°C, obtained by [1993Ale] from extrapolation of [1958Bud] data, is ten times higher than the result of [1983Nas]. As a whole, it ought to agree with judgement of [1987Pri] that interpretation of the obtained experimental data is very difficult because the ´ and phases crystal structures are the superstructures of fcc and bcc lattices, respectively. In the ternary system in the ranges of their coexistence with W or Ni it is very difficult to fix the superstructure reflexes at small relative amounts of intermetallic phases. The and ´ phases in the ranges of their existence are very similar both by chemical compositions, lattice parameters, and by microhardness. Pseudobinary Systems The partial pseudobinary section NiAl-W given by [1958Bud] has been corrected by [1993Ale] to account for the results of [1983Nas] revealing a much lower solid solubility of W in NiAl (see Section “Solid Phases”); the composition of the eutectic point is now at 1.4 at.% W. The pseudobinary section NiAl-W in the range of compositions 0 to 10 at.% W is presented in Fig. 1 according to [1991Ale], with small changes in the melting temperature of the phase as 1651°C [2003Sal], whereas [1993Ale] accepted 1640°C according to [Mas2] data. Invariant Equilibria Based on a theoretical analysis of the phase reactions in the Ni-NiAl-W ternary as well as from the experimental investigation of six selected alloys, [1986Udo] reported two invariant equilibria: L++ ´ at 1380 10°C and L+ + ´ at 1340 15°C. The same reaction scheme has been constructed by [1986Nov] based on the experimental liquidus surface, thereby confirming the phase triangulation of [1983Nas] and superseding earlier results by [1968Bud]. The calculated liquidus surface and the calculated isothermal section at 1323°C [1974Kau, 1975Kau], however, suggest two invariant equilibria: L+ + ´ and L++ at 1367°C which are different from [1986Udo]. The reason for this inconsistency could be the differences in binary systems used by [1975Kau] and accepted by [1986Udo]. The other reason could be that [1975Kau] used data of [1958Bud], which contradict recent experimental results. [1993Ale] proposed a new reaction scheme (Fig. 2) based on more recent experimental data by [1991Udo] with some adjustments of data of [1986Nov, 1986Udo, 1987Pri] to the Al-Ni phase diagram of [1987Hil], which is essentially the same as assessed by [2003Sal]. The main feature of this reaction scheme is the change of the character of invariant reactions due to the changes in character of l+ ´ reaction in the Al-Ni system. According to [1993Ale] L++ ´ is a peritectic reaction and L++ ´ is a transition reaction. The tentative compositions of the invariant equilibria that were derived are based on the boundary binary phase diagrams, on the experimental isothermal section at 1250°C and on extrapolations of the solubility of W in the ´ phase from 1000°C (6 at.%) through 1250°C (4.5 at.%) to 1400°C (3.5 at.%), see Table 2. Liquidus, Solidus and Solvus Surfaces The liquidus surface shown in Fig. 3 contains four fields of primary crystallization: , , and ' and is primarily based on the findings of [1986Nov] and on the data of [1978Gul] for the Ni-rich corner. With respect to the adopted boundary systems the paths of the tie-triangle liquidus vertex and univariant curves shown in Fig. 3 are adjusted to the reaction scheme (Fig. 2) and to the compositions of the invariant equilibria (Table 2). The content of W in the liquid for the reaction L+ is taken from the data of Landolt-Börnstein New Series IV/11A3
MSIT ®
442
Al–Ni–W
[1993Ale]. One of the main features of the liquidus surface is the relatively small area of primary crystallization of the ' phase. The (Ni) solvus was determined by [1991Mis] using DTA as major experimental technique. Chemical analyses using energy-dispersive X-ray spectroscopy proved high accuracy of DTA results and provided information on the phase relations such as the three-phase triangle neighboring the + ´ two-phase field. The results on the determination of the solvus in Al-Ni-W system are presented in the form of solvus isotherms in Fig. 4. The solvus is shown by solid lines connected to broken lines at clearly defined inflections, which indicates the appearance of the three-phase equilibrium + + ´. Isothermal Sections [1958Bud] presented the isothermal section of the partial Ni-NiAl-W system at 1200, 1000 and 800°C. But calculations of isothermal sections carried out by [1975Kau] for 1423, 1323, 1123 and 923°C, showed essentially different phase equilibria, compared with those by [1958Bud]. [1975Kau] showed the existence of the ´+ two-phase region, which was confirmed later by experimental results of [1983Nas, 1986Nov, 1986Udo, 1987Pri, 1989Hon1, 1989Hon2, 1991Eno, 1991Udo]. Besides that, [1958Bud] and [1986Nov] showed a high solubility of tungsten in the phase (up to 10 at.% at 1600°C), but [1983Nas] established, that it is not more than 0.2 at.% at 1250°C. These different values can not be reconlied but the data [1983Nas] can be preferred because a direct method (EMPA) was used to determine the phase composition, while [1958Bud] used indirect methods of DTA and X-ray diffraction. Results of [1994Jia] obtained by diffusion couples method, also indicate a small solubility of tungsten in the phase, not higher than 0.53 at.% in the temperature range 1300 to 900°C. The isothermal section at 1250°C on the basis of [1983Nas] data is presented in Fig. 5. The experimental phase compositions given in [1986Udo] and [1991Udo] are in good agreement with this isothermal section. But according to the [1987Pri] data for 1200°C, the three-phase + + ´ field is much narrower along the Al-content. Therefore further investigations of the isothermal sections are necessary. The position of the phase boundary in Fig. 5 is slightly corrected according to the accepted Ni-W binary system. The composition of the vertex of the + + ´ three-phase field has been taken from [1989Hon1, 1989Hon2, 1991Mis] for 1227°C. The position of the boundary of the ´ phase at the Ni side calculated by [1991Eno] is slightly shifted in the direction of increasing Ni contents, compared with [1983Nas]. [1991Eno] employed the cluster variation method (CVM) which utilizes the tetrahedron approximation and the phenomenological Lennard-Jones pair interaction potential. Thermodynamics Information on thermodynamic properties of the Al-Ni-W alloys is not quite complete. The thermodynamic activities of Al in the ternary system for the Al-Ni0.9162W0.0838 section with aluminium content from 0 to 9 at.% have been determined by [1968Mal] using the emf method. The measurements were conducted at temperatures of 772 and 907°C. The obtained values for the excess integral Gibbs energies and for the activity coefficients at 772 and 907°C are presented in Table 3. Notes on Materials Properties and Applications Al-Ni alloys with additions of a refractory metal, in particular tungsten, are interesting materials for the production of in situ composites of eutectic superalloys. In spite of very complicated alloy compositions, commercial superalloys generally consist mainly of two phases, namely, and ´. The phase has been used for surface coating of the superalloys because of its high resistance against oxidation [1987Woo]. [1991Sas], studying the alloying effect of tungsten on the solidification of the phase, classified it as the eutectic-phase containing compounds. The grain sizes appeared to correlate with the melting temperatures of the compounds, similar to NiAl based ternary phases, formed by other elements. [1995Juj] investigated the tensile properties of the ´ phase reinforced with continuous tungsten fibers. Model composites were fabricated by isothermal forging of sandwiched tungsten fibers between boron-doped ´ plates at temperatures from 1100 to 1200°C. It was found that the use of cold rolled ´ plates for hot forging enables better consolidation and a lower forging temperature than the use of recrystallized ´ plates. Tensile test of MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–W
443
the / ´ composites at ambient temperature to 1050°C reveals that the composites are stronger than monolithic ´ alloys at the test temperatures. [2001Kai] studied the effect of tungsten on the morphological stability of the interface between the ´ and phases using the Al-Ni-W ternary diffusion couples annealed at temperatures ranging from 900 to 1300°C. Nonplanar interfaces with the Widmanstaetten-like structure were formed in the couples. Measurements of electrical resistivity were used by [2001Sav] to study the kinetics of ordering in the ´ phase. Investigations were conducted contactless levitating the sample by a rotary magnetic field. The ordering and disordering processes are described by C-shaped time-temperature transition diagrams. The activation energy values of the ordering and disordering are estimated using the Arrhenius equation. References [1958Bud]
[1966Arb]
[1968Mal]
[1973Mul] [1974Kau]
[1975Kau] [1978Gul]
[1983Nas] [1983Och] [1984Och1] [1984Och2]
[1985Mis]
[1985Nas]
[1986Nov]
[1986Udo]
Landolt-Börnstein New Series IV/11A3
Budberg, P.B., “Study of Alloys of the Ternary System Nickel-Aluminum-Tungsten”, Russ. J. Inorg. Chem., USSR, 3, translated from Zh. Neorg. Khim., 3(3), 694-698 (1958) (Equi. Diagram, Experimental, #, 8) Arbuzov, M.P., Kachkovskaya, E.T., Khayenko, B.V., “Structural X-Ray Diffraction Study of the Compound Ni3Al Alloyed with Ti, Cr and W”, Russ. Met. Phys. Met. Sci., 21(6), 46-49 (1966), translated from Fiz. Met. Metalloved., 21(6), 854-857 (1966) (Crys. Structure, Experimental, 15) Malkin, V.I., Pokidyshev, V.V., “The Effect of Alloying Elements on the Thermodynamic Properties of Ni-Al Alloys” (in Russian), Sb. Tekhn. Trud. Nauchno-Issled. Inst. Chern. Met., 59, 94-99 (1968) (Equi. Diagram, Experimental, Thermodyn., 4) Mulford, R.M., Pope, D.P., “The Yield Stress of Ni3(Al,W)”, Acta Metall., 21, 1375-1380 (1973) (Experimental, 24) as quoted by [1993Ale] Kaufman, L., Nesor, H., “Computer Calculated Phase Diagrams for the Ni-W-Al, Ni-Al-Hf, Ni-Cr-Hf and Co(Cr,Ni)-Ta-C Systems”, Report No. NASA CR-134608, 55 (1974) (Equi. Diagram, Theory, #, 28) as quoted by [1993Ale] Kaufman, L., Nesor, H., “Calculation of the Ni-Al-W, Ni-Al-Hf and Ni-Cr-Hf Systems”, Can. Metall. Q., 14, 221-232 (1975) (Equi. Diagram, Theory, #, 22) Gulyaev, B.B., Grigorash, E.F., Efimova, M.N., “Investigation of Solidification Ranges of Nickel Alloys” (in Russian), Metalloved. Term. Obrab. Metallov., 11, 34-37 (1978) (Equi. Diagram, Experimental, 8) 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) (Equi. Diagram, Experimental, 20) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al With Ternary Additions”, Bull. P.M.E., 52, 1-17 (1983) (Equi. Diagram, Experimental, 7) as quoted by [1993Ale] Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32(2), 289-298 (1984) (Experimental, 90) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data on Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions”, Bull. P.M.E. (T.I.T.), 53, 15-28 (1984) (Crys. Structure, Experimental, 66) Mishima, Y., Ochiai, S., Suzuki, “Lattice Parameters of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33(6), 1161-1169 (1985) (Crys. Structure, Experimental, 64) Nash, P., “Nickel-Base Intermetallics for High Temperature Alloy Design”, High-Temp. Ordered Intermet. Alloys, Boston, Mat. Res. Soc. Conf., 423-427 (1985) (Equi. Diagram, Review, 15) Novikova, M.B., Budberg, P.B., “Phase State of Cast Alloys of Ni-NiAl-W System”, Russ. Metall., (4), 407-111 (1986), translated from Izv. Akad. Nauk SSSR, Met., (4), 104-108 (1986) (Equi. Diagram, Experimental, #, 6) Udovskii, A.L., Alekseeva, Z.M., Lukovkin, A.I., “Phase Equilibrium Diagram of the Nickel-Aluminum-Tungsten System in the Range 1200-2000°C for the Concentration MSIT ®
444
[1987Hil]
[1987Pri]
[1987Woo] [1989Hon1]
[1989Hon2]
[1991Eno]
[1991Mis]
[1991Sas]
[1991Udo]
[1993Ale]
[1994Jia]
[1995Juj] [2001Kai]
[2001Sav]
[2003Sal]
[2003Sch]
MSIT®
Al–Ni–W Region Ni-Ni0.5Al0.5W”, Sov. Phys., Dokl., 288(4), 496-499 (1986), translated from Dokl. Akad. Nauk SSSR, 288(4), 935-939 (1986) (Equi. Diagram, Experimental, #, 13) Hilpert, K., Kobertz, D., Venugopol, V., Miller, M., Gerads, H., Bremer, F.J., Nickel, H., “Phase Diagram Studies of the Al-Ni System”, Z. Naturforsch., 42A, 1327-1332 (1987) (Equi. Diagram, Experimental, #, 17) Prima, S.B., “The Isothermal Section of the W-Ni-Al Phase Diagram in the Range of W-Ni-NiAl at 1200°C” (in Russian) in “Stabilnye i Metastabil'nye Fasy v Materialakh”, Stable and Metastable Phases in Materials, Kiev, IPM, 97-105 (1987) (Equi. Diagram, Experimental, #, 9) Wood, J.E., Goldman, E., in “Superalloys II”, Sims, C.T., Stoloff, N.S., Hagel, W.C. (Eds.), New York, John Willey & Sons, 359-384 (1987) (Experimental) as quoted by [1994Jia] Hong, Y.M., Nakajima, H., Mishima, Y., Suzuki, T., “The Solvus Surface in Ni-Al-X (X: Cr, Mo and W) Ternary Systems”, I.S.I.J. International, 29(1), 78-84 (1989) (Equi. Diagram, Experimental, 25) Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of ´ Solvus in Ni-Al-X Ternary Systems”, Mat. Res. Soc. Symp. Proc., 133, 429-440 (1989) (Equi. Diagram, Experimental, 35) as quoted by [1993Ale] Enomoto, M., Harada, H., Yamazaki, M., “Calculation of ´/ Equilibrium Phase Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15(2), 143-158 (1991) (Assessment, Calculation, Equi. Diagram, 34) Mishima, Y., Hong, Y., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Equi. Diagram, Experimental, 5) Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on the Solidified Structure of NiAl”, Intermetal. Comp. - Struct. Mechan. Prop., Proc. Conf., 877-881 (1991) (Abstract, Equi. Diagram, Experimental, Mechan. Prop., 10) Udovskii, A.L., Oldakovskii, I.V., Moldavskii, V.G., “Theoretical and Experimental Investigations of Phase Equilibria in the Al-Ni-W System in the Range 900 to 1500°C” (in Russian), Izv. Akad. Nauk. SSSR Met., 4, 112-123 (1991) (Equi. Diagram, Experimental) Alekseeva, Z.M., “Al-Ni-W (Aluminium - Nickel - Tungsten)”, 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, Equi. Diagram, Assessment, 23) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1), ´ (L12) and (B2) Phases In Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473-485 (1994) (Crys. Structure, Equi. Diagram, Experimental, 25) Jujur, I.N., Hanada, S., “Tensile Properties of W/Ni 3Al Composites at Elevated Temperatures”, Mater. Sci. Eng. A, 192/193, 848-855 (1995) (Equi. Diagram, Review, 20) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of ´/ Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168-175 (2001) (Equi. Diagram, Experimental, Thermodyn., 21) 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, 18) 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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164) Schuster, J., “Al-W (Aluminium - Tungsten)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; submitted for publication, (2003) (Crys. Structure, Equi. Diagram, Assessment, 22) Landolt-Börnstein New Series IV/11A3
Al–Ni–W
445
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 , WxNi1-x-yAly < 1455 WxNi1-x
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype cF4 Fm3m Cu a = 404.96 cF4 Fm3m Cu
a = 352.40 a = 356.35 a = 358.8 a = 357.4 a = 357.3 a = 357.0 , (W) < 3422
cI2 Im3m W a = 316.52
´, Ni3Al < 1372
cP4 Pm3m AuCu3
a = 356.6 a = 356.77 a = 356.32 a = 357.92 a = 357.30 a = 357.0 a = 357.8 a = 357.3 a = 357.6 a = 357.6 a = 357.0 a = 358.87
Ni5Al3 < 723
Landolt-Börnstein New Series IV/11A3
oC16 Cmmm Pt5Ga3
a = 744 b = 668 c = 372
Comments/References
(Al) dissolves up to 0.1 at.% of Ni at 639.9°C and up to 0.024 at.% W at 650°C pure Al, T = 25°C [Mas2] 0 x 0.11 0 y 0.2 at 1250°C [1983Nas] 0 y 0.203 at 1200°C [1987Pri] 0 x 0.16 y = 0 at 1002°C [Mas2] 0 x 0.175 at 1495°C [Mas2] 0 y 0.202 x = 0 at 1385°C [Mas2] pure Ni, at 25°C [1984Och2, Mas2] x = 0.09 [1985Mis], linear da/dx x = 0.155 [1984Och2], linear da/dx scaled from diagram x = 0.05 y = 0.1783, annealed at 1200°C, together with phase [1987Pri] x = 0.053 y = 0.2018, annealed at 1200°C, together with phase [1987Pri] x = 0.0431 y = 0.2027, annealed at 1200°C, together with phase [1987Pri] dissolves up to 2.6 at.% Al [2003Sch]; up to 0.05 Ni at T = 1187 °C and 0.6 at.% Ni at 1927°C [Mas2] pure W, at 25°C [Mas2] ´ contains 73 to 76 at.% Ni [Mas2] and up to 4-6 at.% W [1983Nas, 1984Och1] [2003Sal] [2003Sal] Disordered [2003Sal] Ordered [2003Sal] at 75 at.%Ni [1966Arb] at 75 at.% Ni [1984Och2] 3W-75Ni (at.%) [1984Och2], linear da/dx [1985Mis] 1.9W-73.38Ni (at.%), annealed at 1200C, together with phase [1987Pri] 0.94W-74.12Ni (at.%), annealed at 1200°C, together with and phases [1987Pri] 2W-74.12Ni (at.%), annealed at 1200°C, together with and phases [1987Pri] 1.4 at.% W [1966Arb] 3 at.% W, 75 at.% Ni, annealed at 1227°C (6 h) [2001Sav] 63 to 68 at.% Ni [2003Sal, Mas2] 63 at.% Ni [2003Sal]
MSIT ®
Al–Ni–W
446 Phase/ Temperature Range [°C] , NiAl < 1638
Pearson Symbol/ Lattice Parameters [pm] Space Group/ Prototype cP2 Pm3m CsCl a = 287.04 a = 287.26 a = 286.0 a = 287.0 a = 288.72 0.02 a = 287.98 0.02 a = 286.6 a = 286.4
´, Ni3Al4 < 702
, Ni 2Al3 < 1138 J, NiAl3 < 856 Ni2Al9
NixAl1-x
cI112 Ia3d Ni3Ga4 hP5 P3m1 Ni2Al3 oP16 Pnma NiAl3 mP22 P2 1/c Ni2Al9 tP4 P4/mmm AuCu
m**
MSIT®
a = 1140.8 0.1
a = 402.8 c = 489.1 a = 661.3 0.1 b = 736.7 0.1 c = 481.1 0.1 a = 868.5 0.6 b = 623.2 0.4 c = 618.5 0.4 = 96.50 0.5° 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 = 90° = 93.4° = 90°
Comments/References
, NiAl contains 42 to 69.2 at.% Ni [Mas2] 0.2-1 at.% W [1983Nas, 1985Nas, 1993Ale] 57.7 at.% Ni [L-B] 46.6 at.% Ni [L-B] [2003Sal] 63 at.% Ni [2003Sal] 50 at.% Ni [2003Sal] 54 at.% Ni [2003Sal]; 0.18W-65.51Ni (at.%), annealed at 1200°C, together with ´ phase [1987Pri] 0.02W-63.12Ni (at.%), annealed at 1200°C, together with and ´ phase [1987Pri] [2003Sal, V-C]
36.8 to 40.5 at.% Ni [Mas2] [2003Sal, V-C] [2003Sal]
Metastable [2003Sal]
Martensite, metastable, 0.60 x 0.68 [2003Sal] 62.5 at.% Ni [2003Sal] 63.5 at.% Ni [2003Sal] 66.0 at.% Ni [2003Sal] 64 at.% Ni [2003Sal] 65 at.% Ni [2003Sal] [2003Sal] [2003Sal]
Landolt-Börnstein New Series IV/11A3
Al–Ni–W Phase/ Temperature Range [°C] Ni2Al
D1 D4 WAl12 < 697 WAl6
WAl5 < 870 WAl4 < 1326
Pearson Symbol/ Space Group/ Prototype hP3 P3m1 CdI 2 P105mc or P105/mmc cI26 Im3 WAl12 mC56 C2/c MoAl6 hP12 P6 3 WAl5 mC30 Cm WAl4
Lattice Parameters Comments/References [pm]
a = 407 b = 499 a = 373.3 c = 407.3 a = 758.03
a = 514.4 0.3 b = 1298.9 0.5 c = 1348.4 0.4 = 94.03 0.39° a = 490.20 c = 885.70 a = 527.2 b = 1777.1 c = 521.8 = 100.2°
Q(Al-W)
c**
a = 787.0 0.4 c = 2380 3 a = 714.5 c = 787.4 a = 692 8
cI*
a = 766.4
h** t** Q´(Al-W)
(Al-W) 1300 < T < 1344 ´(Al-W) 1317 < T 1420 ´´(Al-W) 1335 < T 1650 W50 Al50
t**
WNi 1060
o** MoNi
W2Ni
tI96
WNi4
tI10 I4/m MoNi4
Landolt-Börnstein New Series IV/11A3
447
Metastable [2003Sal] [2003Sal] Decagonal, contained some Si [2003Sal] [2003Sch]
In the Al-B-W alloys rich in aluminium [2003Sch]
[2003Sch]
[2003Sch]
More rich in Al than WAl12, T < 580°C [2003Sch] T = 650°C [2003Sch] [2003Sch] Metastable (?), T = 200°C; transforms into WAl12 [2003Sch] [2003Sch] at 24 at.% W [2003Sch] at 30 at.% W [2003Sch] at 33 at.% W [2003Sch]
a = 613 c = 418 a = 776 b = 1248 c = 710 a = 1040 c = 1090 a = 573 1 c = 355.3 0.1
Body-centered (?), metastable (?), from TEM data [2003Sch] at 50 at.% W [V-C]
at 66 at.% W [V-C] at 20 at.% W [V-C]
MSIT ®
Al–Ni–W
448 Table 2: Invariant Equilibria T [°C]
Reaction
Type
Phase
Composition (at.%) Al
Ni
W
L+
~1400
U
L ´
25.0 1.5 30.8 23.0
73.0 1.5 69.0 73.0
2.0 97.0 0.2 4.0
L + + ´
~1405
P
L ´
23.5 1.5 10.0 23.5
73.5 1.5 80.0 73.0
3.0 97.0 10.0 3.5
Table 3: Integral Excess Gibbs Energies of Al-Ni-W Alloys and Activity Coefficients of Aluminium at 772 and 907°C [1968Mal] fGes [J#mol-1]
xAl 0,01 0.03 0.05 0.07 0.09
lg Al
772°C
907°C
772°C
907°C
- 1710 - 3680 - 5580 - 7240 - 8860
-1550 -3430 -5180 -6810 -8440
-5.17 -4.79 -4.47 -4.22 -3.99
-4.34 -4.00 -3.76 -3.62 -3.52
1800
Fig. 1: Al-Ni-W. Pseudobinary section NiAl-W in the range of compositions 0 to 10 at.% W
Temperature, °C
1700
L
α +L
1600°C
1600
~1.4
MSIT®
1500
10.00 45.00 45.00
8
6
4
W, at.%
~1
β
α +β
W Ni Al
~1651°C
2
W Ni Al
0.00 50.00 50.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–W
Ni-W
449
Al-Ni-W
Al-Ni
1600 e1(max) Lβ+α
1495 e2 lγ+α
L + α + γ γ´
1405
P L + α γ´
α + γ + γ´
L + α β + γ´
1400
U
α + β + γ´
1372 p l + γ γ´ 1369 e3 l β + γ´
Fig. 2: Al-Ni-W. Partial reaction scheme
W Ni Al
Fig. 3: Al-Ni-W. Liquidus surface of the Ni-rich region
0.00 50.00 50.00
Data / Grid: at.% Axes: at.%
e1,1600
10
1637 1600
40
β
20
30
e3,1369 U P
γ'
p,1372
30
20
α
1440
40
10
γ 1486
W Ni Al Landolt-Börnstein New Series IV/11A3
50.00 50.00 0.00
60
70
e2,1495 80
90
Ni
MSIT ®
Al–Ni–W
450
W Ni Al
Fig. 4: Al-Ni-W. The (Ni) solvus surface
0.00 70.00 30.00
Data / Grid: at.% Axes: at.%
10
20
γ +γ ´ °C 27 12 °C 27 11 °C 27 10 C 7° 92 °C 7 2 8
20
10
γ
W Ni Al
80
30.00 70.00 0.00
90
Ni
Al
Data / Grid: at.%
Fig. 5: Al-Ni-W. Partial isothermal section at 1250°C
Axes: at.%
20
80
40
60
β
60
40
α +β γ'
α+β +γ '
80
20
α+γ +γ ' α
W
MSIT®
γ
α+γ 20
40
60
80
Ni
Landolt-Börnstein New Series IV/11A3
440
Al–Ni–W
Aluminium – Nickel – Tungsten Konstyantyn Korniyenko, Vasyl Kublii, Olga Fabrichnaya, Natalia Bochvar Literature Data Experimental investigations of the phase equilibria in the Al-Ni-W system are limited to the Ni-rich range of compositions [1958Bud, 1978Gul, 1983Nas, 1986Nov, 1986Udo, 1991Udo, 1994Jia]. Solid solubility of tungsten in the Ni3Al ( ´) phase is presented in [1966Arb, 1983Och, 1984Och1, 1984Och2, 1985Mis, 1985Nas] and the , (Ni) solvus surface is described in [1989Hon1, 1989Hon2] and [1991Mis]. [1991Sas] studied the alloying effect of tungsten, on the solidification of the NiAl () phase. [1958Bud] investigated about 60 ternary alloys, that were prepared by high-frequency melting using corundum crucibles under a protective layer of basic slag. The starting components were A-000 aluminium, nickel (99.9 mass%) and tungsten (99.98 mass%). After subsequent stepwise annealing in high vacuum at 1200°C for 24 h, at 1000°C for an additional 100 h and at 800°C for another 100 h, respectively, each batch of alloys was partly quenched and cooled to room temperature. All losses were less than 0.2 to 0.5 mass% due to the fact that the Al was introduced by means of premelted Al-Ni master alloys. Three isothermal sections of the Al-Ni-W system in the Ni rich range at 1200, 1000 and 800°C and part of the NiAl-W pseudobinary section are plotted. Since then the Ni-NiAl-W partial system has been reinvestigated frequently and partial isothermal sections have been established by various research groups: at 1370°C [1986Udo], 1250°C [1983Nas], 1200°C [1986Udo, 1987Pri], 1150°C [1991Udo], 900°C [1991Udo] and for the temperature range from the end of alloy solidification up to the beginning of solid state reactions, i.e. from ~1350 to 1000°C [1986Nov]. [1978Gul] recorded the liquidus and solidus temperatures in the Ni-rich corner by means of calibrated thermocouples. [1986Nov] presented also the liquidus surface projection and reaction scheme for the partial Ni-NiAl-W system. In most cases the alloys were prepared by arc melting of the elements with 99.9 mass% minimal purity on a water-cooled copper hearth under an argon atmosphere using a nonconsumable tungsten electrode [1966Arb, 1983Nas, 1983Och, 1984Och1, 1984Och2, 1985Mis, 1985Nas, 1986Nov, 1986Udo, 1987Pri, 1989Hon1, 1989Hon2, 1991Mis, 1991Sas, 1991Udo], by induction melting under an argon atmosphere [1994Jia] as well as by high-frequency melting [1986Nov, 1991Udo]. Due to rather sluggish reaction kinetics, it was difficult to obtain equilibrium and homogenization treatments up to 500 h at various temperatures were carried out. After heat treatment the samples usually were quenched to room temperature. Methods of experimental investigation of the alloys were: X-ray diffraction [1966Arb, 1983Nas, 1983Och, 1984Och1, 1984Och2, 1985Mis, 1986Udo, 1986Nov, 1987Pri], metallography [1983Nas, 1986Nov, 1986Udo, 1987Pri, 1991Sas, 1991Udo, 1994Jia], electron microprobe analysis (EMPA) [1983Nas, 1986Nov, 1987Pri, 1989Hon1, 1989Hon2, 1991Mis, 1991Udo, 1994Jia], differential thermal analysis (DTA) [1986Nov, 1989Hon1, 1989Hon2, 1991Mis] and [1991Udo] as well as microhardness measurements [1987Pri]. The critical review of literature data on the phase equilibria in the system was carried out by [1993Ale] within the MSIT Evaluation Program and is continued and updated by the present evaluation. Binary Systems The descriptions of the Al-Ni, Al-W and Ni-W systems are accepted from [2003Sal], [2003Sch] and [Mas2], respectively. Solid Phases No ternary phases have been found. Crystallographic data on the known unary and binary phases are listed in Table 1. Based on earlier investigations of [1966Arb], the extent of the ´ phase field on alloying with tungsten as well as the mode of atom substitution has been studied by [1983Och, 1984Och1, 1984Och2] and [1985Mis] who arrived at a maximum solubility of less than 5 at.% W at 1000°C with W replacing Al
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–W
441
in WxNi3Al1-x [1983Och, 1984Och1, 1984Och2, 1985Mis], whereas in general W was found to substitute for both lattice sites in Wx(W yNi1-y)3Al1-x [1966Arb]. Data on solubility of W in the ´ phase are contradictory and need more accurate definition. So, [1973Mul] and [2001Sav] prepared single crystals WxNi3Al1-x with 3 at.% W replacing Al. A slightly larger solubility of 6 at.% W in ´ at 1000°C was obtained by [1983Nas] on the basis of EMPA data, and the solubility of W in the ´ phase was found to decrease from ~6 to ~4.5 at.% as the temperature increases from 1000 to 1250°C [1983Nas, 1985Nas]. But [1987Pri] obtained the solubility of W in the ´ phase ~2 at.% at 1200°C, while [1991Udo] reported ~4 at.% W at this temperature (Table 1). Similarly the solubility of W in the phase was shown to be ~0.2 at.% at 1250°C [1983Nas, 1985Nas], whereas [1958Bud] reported a value of 10 at.% at 1600°C, and 6 at.% at 1500°C. The value of W solubility of ~2 at.% at 1250°C, obtained by [1993Ale] from extrapolation of [1958Bud] data, is ten times higher than the result of [1983Nas]. As a whole, it ought to agree with judgement of [1987Pri] that interpretation of the obtained experimental data is very difficult because the ´ and phases crystal structures are the superstructures of fcc and bcc lattices, respectively. In the ternary system in the ranges of their coexistence with W or Ni it is very difficult to fix the superstructure reflexes at small relative amounts of intermetallic phases. The and ´ phases in the ranges of their existence are very similar both by chemical compositions, lattice parameters, and by microhardness. Pseudobinary Systems The partial pseudobinary section NiAl-W given by [1958Bud] has been corrected by [1993Ale] to account for the results of [1983Nas] revealing a much lower solid solubility of W in NiAl (see Section “Solid Phases”); the composition of the eutectic point is now at 1.4 at.% W. The pseudobinary section NiAl-W in the range of compositions 0 to 10 at.% W is presented in Fig. 1 according to [1991Ale], with small changes in the melting temperature of the phase as 1651°C [2003Sal], whereas [1993Ale] accepted 1640°C according to [Mas2] data. Invariant Equilibria Based on a theoretical analysis of the phase reactions in the Ni-NiAl-W ternary as well as from the experimental investigation of six selected alloys, [1986Udo] reported two invariant equilibria: L++ ´ at 1380 10°C and L+ + ´ at 1340 15°C. The same reaction scheme has been constructed by [1986Nov] based on the experimental liquidus surface, thereby confirming the phase triangulation of [1983Nas] and superseding earlier results by [1968Bud]. The calculated liquidus surface and the calculated isothermal section at 1323°C [1974Kau, 1975Kau], however, suggest two invariant equilibria: L+ + ´ and L++ at 1367°C which are different from [1986Udo]. The reason for this inconsistency could be the differences in binary systems used by [1975Kau] and accepted by [1986Udo]. The other reason could be that [1975Kau] used data of [1958Bud], which contradict recent experimental results. [1993Ale] proposed a new reaction scheme (Fig. 2) based on more recent experimental data by [1991Udo] with some adjustments of data of [1986Nov, 1986Udo, 1987Pri] to the Al-Ni phase diagram of [1987Hil], which is essentially the same as assessed by [2003Sal]. The main feature of this reaction scheme is the change of the character of invariant reactions due to the changes in character of l+ ´ reaction in the Al-Ni system. According to [1993Ale] L++ ´ is a peritectic reaction and L++ ´ is a transition reaction. The tentative compositions of the invariant equilibria that were derived are based on the boundary binary phase diagrams, on the experimental isothermal section at 1250°C and on extrapolations of the solubility of W in the ´ phase from 1000°C (6 at.%) through 1250°C (4.5 at.%) to 1400°C (3.5 at.%), see Table 2. Liquidus, Solidus and Solvus Surfaces The liquidus surface shown in Fig. 3 contains four fields of primary crystallization: , , and ' and is primarily based on the findings of [1986Nov] and on the data of [1978Gul] for the Ni-rich corner. With respect to the adopted boundary systems the paths of the tie-triangle liquidus vertex and univariant curves shown in Fig. 3 are adjusted to the reaction scheme (Fig. 2) and to the compositions of the invariant equilibria (Table 2). The content of W in the liquid for the reaction L+ is taken from the data of Landolt-Börnstein New Series IV/11A3
MSIT ®
442
Al–Ni–W
[1993Ale]. One of the main features of the liquidus surface is the relatively small area of primary crystallization of the ' phase. The (Ni) solvus was determined by [1991Mis] using DTA as major experimental technique. Chemical analyses using energy-dispersive X-ray spectroscopy proved high accuracy of DTA results and provided information on the phase relations such as the three-phase triangle neighboring the + ´ two-phase field. The results on the determination of the solvus in Al-Ni-W system are presented in the form of solvus isotherms in Fig. 4. The solvus is shown by solid lines connected to broken lines at clearly defined inflections, which indicates the appearance of the three-phase equilibrium + + ´. Isothermal Sections [1958Bud] presented the isothermal section of the partial Ni-NiAl-W system at 1200, 1000 and 800°C. But calculations of isothermal sections carried out by [1975Kau] for 1423, 1323, 1123 and 923°C, showed essentially different phase equilibria, compared with those by [1958Bud]. [1975Kau] showed the existence of the ´+ two-phase region, which was confirmed later by experimental results of [1983Nas, 1986Nov, 1986Udo, 1987Pri, 1989Hon1, 1989Hon2, 1991Eno, 1991Udo]. Besides that, [1958Bud] and [1986Nov] showed a high solubility of tungsten in the phase (up to 10 at.% at 1600°C), but [1983Nas] established, that it is not more than 0.2 at.% at 1250°C. These different values can not be reconlied but the data [1983Nas] can be preferred because a direct method (EMPA) was used to determine the phase composition, while [1958Bud] used indirect methods of DTA and X-ray diffraction. Results of [1994Jia] obtained by diffusion couples method, also indicate a small solubility of tungsten in the phase, not higher than 0.53 at.% in the temperature range 1300 to 900°C. The isothermal section at 1250°C on the basis of [1983Nas] data is presented in Fig. 5. The experimental phase compositions given in [1986Udo] and [1991Udo] are in good agreement with this isothermal section. But according to the [1987Pri] data for 1200°C, the three-phase + + ´ field is much narrower along the Al-content. Therefore further investigations of the isothermal sections are necessary. The position of the phase boundary in Fig. 5 is slightly corrected according to the accepted Ni-W binary system. The composition of the vertex of the + + ´ three-phase field has been taken from [1989Hon1, 1989Hon2, 1991Mis] for 1227°C. The position of the boundary of the ´ phase at the Ni side calculated by [1991Eno] is slightly shifted in the direction of increasing Ni contents, compared with [1983Nas]. [1991Eno] employed the cluster variation method (CVM) which utilizes the tetrahedron approximation and the phenomenological Lennard-Jones pair interaction potential. Thermodynamics Information on thermodynamic properties of the Al-Ni-W alloys is not quite complete. The thermodynamic activities of Al in the ternary system for the Al-Ni0.9162W0.0838 section with aluminium content from 0 to 9 at.% have been determined by [1968Mal] using the emf method. The measurements were conducted at temperatures of 772 and 907°C. The obtained values for the excess integral Gibbs energies and for the activity coefficients at 772 and 907°C are presented in Table 3. Notes on Materials Properties and Applications Al-Ni alloys with additions of a refractory metal, in particular tungsten, are interesting materials for the production of in situ composites of eutectic superalloys. In spite of very complicated alloy compositions, commercial superalloys generally consist mainly of two phases, namely, and ´. The phase has been used for surface coating of the superalloys because of its high resistance against oxidation [1987Woo]. [1991Sas], studying the alloying effect of tungsten on the solidification of the phase, classified it as the eutectic-phase containing compounds. The grain sizes appeared to correlate with the melting temperatures of the compounds, similar to NiAl based ternary phases, formed by other elements. [1995Juj] investigated the tensile properties of the ´ phase reinforced with continuous tungsten fibers. Model composites were fabricated by isothermal forging of sandwiched tungsten fibers between boron-doped ´ plates at temperatures from 1100 to 1200°C. It was found that the use of cold rolled ´ plates for hot forging enables better consolidation and a lower forging temperature than the use of recrystallized ´ plates. Tensile test of MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–W
443
the / ´ composites at ambient temperature to 1050°C reveals that the composites are stronger than monolithic ´ alloys at the test temperatures. [2001Kai] studied the effect of tungsten on the morphological stability of the interface between the ´ and phases using the Al-Ni-W ternary diffusion couples annealed at temperatures ranging from 900 to 1300°C. Nonplanar interfaces with the Widmanstaetten-like structure were formed in the couples. Measurements of electrical resistivity were used by [2001Sav] to study the kinetics of ordering in the ´ phase. Investigations were conducted contactless levitating the sample by a rotary magnetic field. The ordering and disordering processes are described by C-shaped time-temperature transition diagrams. The activation energy values of the ordering and disordering are estimated using the Arrhenius equation. References [1958Bud]
[1966Arb]
[1968Mal]
[1973Mul] [1974Kau]
[1975Kau] [1978Gul]
[1983Nas] [1983Och] [1984Och1] [1984Och2]
[1985Mis]
[1985Nas]
[1986Nov]
[1986Udo]
Landolt-Börnstein New Series IV/11A3
Budberg, P.B., “Study of Alloys of the Ternary System Nickel-Aluminum-Tungsten”, Russ. J. Inorg. Chem., USSR, 3, translated from Zh. Neorg. Khim., 3(3), 694-698 (1958) (Equi. Diagram, Experimental, #, 8) Arbuzov, M.P., Kachkovskaya, E.T., Khayenko, B.V., “Structural X-Ray Diffraction Study of the Compound Ni3Al Alloyed with Ti, Cr and W”, Russ. Met. Phys. Met. Sci., 21(6), 46-49 (1966), translated from Fiz. Met. Metalloved., 21(6), 854-857 (1966) (Crys. Structure, Experimental, 15) Malkin, V.I., Pokidyshev, V.V., “The Effect of Alloying Elements on the Thermodynamic Properties of Ni-Al Alloys” (in Russian), Sb. Tekhn. Trud. Nauchno-Issled. Inst. Chern. Met., 59, 94-99 (1968) (Equi. Diagram, Experimental, Thermodyn., 4) Mulford, R.M., Pope, D.P., “The Yield Stress of Ni3(Al,W)”, Acta Metall., 21, 1375-1380 (1973) (Experimental, 24) as quoted by [1993Ale] Kaufman, L., Nesor, H., “Computer Calculated Phase Diagrams for the Ni-W-Al, Ni-Al-Hf, Ni-Cr-Hf and Co(Cr,Ni)-Ta-C Systems”, Report No. NASA CR-134608, 55 (1974) (Equi. Diagram, Theory, #, 28) as quoted by [1993Ale] Kaufman, L., Nesor, H., “Calculation of the Ni-Al-W, Ni-Al-Hf and Ni-Cr-Hf Systems”, Can. Metall. Q., 14, 221-232 (1975) (Equi. Diagram, Theory, #, 22) Gulyaev, B.B., Grigorash, E.F., Efimova, M.N., “Investigation of Solidification Ranges of Nickel Alloys” (in Russian), Metalloved. Term. Obrab. Metallov., 11, 34-37 (1978) (Equi. Diagram, Experimental, 8) 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) (Equi. Diagram, Experimental, 20) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data in Ni3Al With Ternary Additions”, Bull. P.M.E., 52, 1-17 (1983) (Equi. Diagram, Experimental, 7) as quoted by [1993Ale] Ochiai, S., Oya, Y., Suzuki, T., “Alloying Behaviour of Ni3Al, Ni3Ga, Ni3Si and Ni3Ge”, Acta Metall., 32(2), 289-298 (1984) (Experimental, 90) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data on Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions”, Bull. P.M.E. (T.I.T.), 53, 15-28 (1984) (Crys. Structure, Experimental, 66) Mishima, Y., Ochiai, S., Suzuki, “Lattice Parameters of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions with Additions of Transition and B-Subgroup Elements”, Acta Metall., 33(6), 1161-1169 (1985) (Crys. Structure, Experimental, 64) Nash, P., “Nickel-Base Intermetallics for High Temperature Alloy Design”, High-Temp. Ordered Intermet. Alloys, Boston, Mat. Res. Soc. Conf., 423-427 (1985) (Equi. Diagram, Review, 15) Novikova, M.B., Budberg, P.B., “Phase State of Cast Alloys of Ni-NiAl-W System”, Russ. Metall., (4), 407-111 (1986), translated from Izv. Akad. Nauk SSSR, Met., (4), 104-108 (1986) (Equi. Diagram, Experimental, #, 6) Udovskii, A.L., Alekseeva, Z.M., Lukovkin, A.I., “Phase Equilibrium Diagram of the Nickel-Aluminum-Tungsten System in the Range 1200-2000°C for the Concentration MSIT ®
444
[1987Hil]
[1987Pri]
[1987Woo] [1989Hon1]
[1989Hon2]
[1991Eno]
[1991Mis]
[1991Sas]
[1991Udo]
[1993Ale]
[1994Jia]
[1995Juj] [2001Kai]
[2001Sav]
[2003Sal]
[2003Sch]
MSIT®
Al–Ni–W Region Ni-Ni0.5Al0.5W”, Sov. Phys., Dokl., 288(4), 496-499 (1986), translated from Dokl. Akad. Nauk SSSR, 288(4), 935-939 (1986) (Equi. Diagram, Experimental, #, 13) Hilpert, K., Kobertz, D., Venugopol, V., Miller, M., Gerads, H., Bremer, F.J., Nickel, H., “Phase Diagram Studies of the Al-Ni System”, Z. Naturforsch., 42A, 1327-1332 (1987) (Equi. Diagram, Experimental, #, 17) Prima, S.B., “The Isothermal Section of the W-Ni-Al Phase Diagram in the Range of W-Ni-NiAl at 1200°C” (in Russian) in “Stabilnye i Metastabil'nye Fasy v Materialakh”, Stable and Metastable Phases in Materials, Kiev, IPM, 97-105 (1987) (Equi. Diagram, Experimental, #, 9) Wood, J.E., Goldman, E., in “Superalloys II”, Sims, C.T., Stoloff, N.S., Hagel, W.C. (Eds.), New York, John Willey & Sons, 359-384 (1987) (Experimental) as quoted by [1994Jia] Hong, Y.M., Nakajima, H., Mishima, Y., Suzuki, T., “The Solvus Surface in Ni-Al-X (X: Cr, Mo and W) Ternary Systems”, I.S.I.J. International, 29(1), 78-84 (1989) (Equi. Diagram, Experimental, 25) Hong, Y.M., Mishima, Y., Suzuki, T., “Accurate Determination of ´ Solvus in Ni-Al-X Ternary Systems”, Mat. Res. Soc. Symp. Proc., 133, 429-440 (1989) (Equi. Diagram, Experimental, 35) as quoted by [1993Ale] Enomoto, M., Harada, H., Yamazaki, M., “Calculation of ´/ Equilibrium Phase Compositions in Nickel-Base Superalloys by Cluster Variation Method”, Calphad, 15(2), 143-158 (1991) (Assessment, Calculation, Equi. Diagram, 34) Mishima, Y., Hong, Y., Suzuki, T., “Determination of the Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng. A, 146, 123-130 (1991) (Equi. Diagram, Experimental, 5) Sasaki, K., Morinaga, M., Yukawa, N., “Alloying Effect on the Solidified Structure of NiAl”, Intermetal. Comp. - Struct. Mechan. Prop., Proc. Conf., 877-881 (1991) (Abstract, Equi. Diagram, Experimental, Mechan. Prop., 10) Udovskii, A.L., Oldakovskii, I.V., Moldavskii, V.G., “Theoretical and Experimental Investigations of Phase Equilibria in the Al-Ni-W System in the Range 900 to 1500°C” (in Russian), Izv. Akad. Nauk. SSSR Met., 4, 112-123 (1991) (Equi. Diagram, Experimental) Alekseeva, Z.M., “Al-Ni-W (Aluminium - Nickel - Tungsten)”, 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, Equi. Diagram, Assessment, 23) Jia, C.C., Ishida, K., Nishizawa, T., “Partition of Alloying Elements Between (A1), ´ (L12) and (B2) Phases In Ni-Al Base Systems”, Metall. Mater. Trans. A, 25, 473-485 (1994) (Crys. Structure, Equi. Diagram, Experimental, 25) Jujur, I.N., Hanada, S., “Tensile Properties of W/Ni 3Al Composites at Elevated Temperatures”, Mater. Sci. Eng. A, 192/193, 848-855 (1995) (Equi. Diagram, Review, 20) Kainuma, R., Ichinose, M., Ohnuma, I., Ishida, K., “Formation of ´/ Interface Morphologies in Ni-Al-X Ternary Diffusion Couples”, Mater. Sci. Eng. A, 312, 168-175 (2001) (Equi. Diagram, Experimental, Thermodyn., 21) 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, 18) 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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164) Schuster, J., “Al-W (Aluminium - Tungsten)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; submitted for publication, (2003) (Crys. Structure, Equi. Diagram, Assessment, 22) Landolt-Börnstein New Series IV/11A3
Al–Ni–W
445
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) < 660.452 , WxNi1-x-yAly < 1455 WxNi1-x
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype cF4 Fm3m Cu a = 404.96 cF4 Fm3m Cu
a = 352.40 a = 356.35 a = 358.8 a = 357.4 a = 357.3 a = 357.0 , (W) < 3422
cI2 Im3m W a = 316.52
´, Ni3Al < 1372
cP4 Pm3m AuCu3
a = 356.6 a = 356.77 a = 356.32 a = 357.92 a = 357.30 a = 357.0 a = 357.8 a = 357.3 a = 357.6 a = 357.6 a = 357.0 a = 358.87
Ni5Al3 < 723
Landolt-Börnstein New Series IV/11A3
oC16 Cmmm Pt5Ga3
a = 744 b = 668 c = 372
Comments/References
(Al) dissolves up to 0.1 at.% of Ni at 639.9°C and up to 0.024 at.% W at 650°C pure Al, T = 25°C [Mas2] 0 x 0.11 0 y 0.2 at 1250°C [1983Nas] 0 y 0.203 at 1200°C [1987Pri] 0 x 0.16 y = 0 at 1002°C [Mas2] 0 x 0.175 at 1495°C [Mas2] 0 y 0.202 x = 0 at 1385°C [Mas2] pure Ni, at 25°C [1984Och2, Mas2] x = 0.09 [1985Mis], linear da/dx x = 0.155 [1984Och2], linear da/dx scaled from diagram x = 0.05 y = 0.1783, annealed at 1200°C, together with phase [1987Pri] x = 0.053 y = 0.2018, annealed at 1200°C, together with phase [1987Pri] x = 0.0431 y = 0.2027, annealed at 1200°C, together with phase [1987Pri] dissolves up to 2.6 at.% Al [2003Sch]; up to 0.05 Ni at T = 1187 °C and 0.6 at.% Ni at 1927°C [Mas2] pure W, at 25°C [Mas2] ´ contains 73 to 76 at.% Ni [Mas2] and up to 4-6 at.% W [1983Nas, 1984Och1] [2003Sal] [2003Sal] Disordered [2003Sal] Ordered [2003Sal] at 75 at.%Ni [1966Arb] at 75 at.% Ni [1984Och2] 3W-75Ni (at.%) [1984Och2], linear da/dx [1985Mis] 1.9W-73.38Ni (at.%), annealed at 1200C, together with phase [1987Pri] 0.94W-74.12Ni (at.%), annealed at 1200°C, together with and phases [1987Pri] 2W-74.12Ni (at.%), annealed at 1200°C, together with and phases [1987Pri] 1.4 at.% W [1966Arb] 3 at.% W, 75 at.% Ni, annealed at 1227°C (6 h) [2001Sav] 63 to 68 at.% Ni [2003Sal, Mas2] 63 at.% Ni [2003Sal]
MSIT ®
Al–Ni–W
446 Phase/ Temperature Range [°C] , NiAl < 1638
Pearson Symbol/ Lattice Parameters [pm] Space Group/ Prototype cP2 Pm3m CsCl a = 287.04 a = 287.26 a = 286.0 a = 287.0 a = 288.72 0.02 a = 287.98 0.02 a = 286.6 a = 286.4
´, Ni3Al4 < 702
, Ni 2Al3 < 1138 J, NiAl3 < 856 Ni2Al9
NixAl1-x
cI112 Ia3d Ni3Ga4 hP5 P3m1 Ni2Al3 oP16 Pnma NiAl3 mP22 P2 1/c Ni2Al9 tP4 P4/mmm AuCu
m**
MSIT®
a = 1140.8 0.1
a = 402.8 c = 489.1 a = 661.3 0.1 b = 736.7 0.1 c = 481.1 0.1 a = 868.5 0.6 b = 623.2 0.4 c = 618.5 0.4 = 96.50 0.5° 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 = 90° = 93.4° = 90°
Comments/References
, NiAl contains 42 to 69.2 at.% Ni [Mas2] 0.2-1 at.% W [1983Nas, 1985Nas, 1993Ale] 57.7 at.% Ni [L-B] 46.6 at.% Ni [L-B] [2003Sal] 63 at.% Ni [2003Sal] 50 at.% Ni [2003Sal] 54 at.% Ni [2003Sal]; 0.18W-65.51Ni (at.%), annealed at 1200°C, together with ´ phase [1987Pri] 0.02W-63.12Ni (at.%), annealed at 1200°C, together with and ´ phase [1987Pri] [2003Sal, V-C]
36.8 to 40.5 at.% Ni [Mas2] [2003Sal, V-C] [2003Sal]
Metastable [2003Sal]
Martensite, metastable, 0.60 x 0.68 [2003Sal] 62.5 at.% Ni [2003Sal] 63.5 at.% Ni [2003Sal] 66.0 at.% Ni [2003Sal] 64 at.% Ni [2003Sal] 65 at.% Ni [2003Sal] [2003Sal] [2003Sal]
Landolt-Börnstein New Series IV/11A3
Al–Ni–W Phase/ Temperature Range [°C] Ni2Al
D1 D4 WAl12 < 697 WAl6
WAl5 < 870 WAl4 < 1326
Pearson Symbol/ Space Group/ Prototype hP3 P3m1 CdI 2 P105mc or P105/mmc cI26 Im3 WAl12 mC56 C2/c MoAl6 hP12 P6 3 WAl5 mC30 Cm WAl4
Lattice Parameters Comments/References [pm]
a = 407 b = 499 a = 373.3 c = 407.3 a = 758.03
a = 514.4 0.3 b = 1298.9 0.5 c = 1348.4 0.4 = 94.03 0.39° a = 490.20 c = 885.70 a = 527.2 b = 1777.1 c = 521.8 = 100.2°
Q(Al-W)
c**
a = 787.0 0.4 c = 2380 3 a = 714.5 c = 787.4 a = 692 8
cI*
a = 766.4
h** t** Q´(Al-W)
(Al-W) 1300 < T < 1344 ´(Al-W) 1317 < T 1420 ´´(Al-W) 1335 < T 1650 W50 Al50
t**
WNi 1060
o** MoNi
W2Ni
tI96
WNi4
tI10 I4/m MoNi4
Landolt-Börnstein New Series IV/11A3
447
Metastable [2003Sal] [2003Sal] Decagonal, contained some Si [2003Sal] [2003Sch]
In the Al-B-W alloys rich in aluminium [2003Sch]
[2003Sch]
[2003Sch]
More rich in Al than WAl12, T < 580°C [2003Sch] T = 650°C [2003Sch] [2003Sch] Metastable (?), T = 200°C; transforms into WAl12 [2003Sch] [2003Sch] at 24 at.% W [2003Sch] at 30 at.% W [2003Sch] at 33 at.% W [2003Sch]
a = 613 c = 418 a = 776 b = 1248 c = 710 a = 1040 c = 1090 a = 573 1 c = 355.3 0.1
Body-centered (?), metastable (?), from TEM data [2003Sch] at 50 at.% W [V-C]
at 66 at.% W [V-C] at 20 at.% W [V-C]
MSIT ®
Al–Ni–W
448 Table 2: Invariant Equilibria T [°C]
Reaction
Type
Phase
Composition (at.%) Al
Ni
W
L+
~1400
U
L ´
25.0 1.5 30.8 23.0
73.0 1.5 69.0 73.0
2.0 97.0 0.2 4.0
L + + ´
~1405
P
L ´
23.5 1.5 10.0 23.5
73.5 1.5 80.0 73.0
3.0 97.0 10.0 3.5
Table 3: Integral Excess Gibbs Energies of Al-Ni-W Alloys and Activity Coefficients of Aluminium at 772 and 907°C [1968Mal] fGes [J#mol-1]
xAl 0,01 0.03 0.05 0.07 0.09
lg Al
772°C
907°C
772°C
907°C
- 1710 - 3680 - 5580 - 7240 - 8860
-1550 -3430 -5180 -6810 -8440
-5.17 -4.79 -4.47 -4.22 -3.99
-4.34 -4.00 -3.76 -3.62 -3.52
1800
Fig. 1: Al-Ni-W. Pseudobinary section NiAl-W in the range of compositions 0 to 10 at.% W
Temperature, °C
1700
L
α +L
1600°C
1600
~1.4
MSIT®
1500
10.00 45.00 45.00
8
6
4
W, at.%
~1
β
α +β
W Ni Al
~1651°C
2
W Ni Al
0.00 50.00 50.00
Landolt-Börnstein New Series IV/11A3
Al–Ni–W
Ni-W
449
Al-Ni-W
Al-Ni
1600 e1(max) Lβ+α
1495 e2 lγ+α
L + α + γ γ´
1405
P L + α γ´
α + γ + γ´
L + α β + γ´
1400
U
α + β + γ´
1372 p l + γ γ´ 1369 e3 l β + γ´
Fig. 2: Al-Ni-W. Partial reaction scheme
W Ni Al
Fig. 3: Al-Ni-W. Liquidus surface of the Ni-rich region
0.00 50.00 50.00
Data / Grid: at.% Axes: at.%
e1,1600
10
1637 1600
40
β
20
30
e3,1369 U P
γ'
p,1372
30
20
α
1440
40
10
γ 1486
W Ni Al Landolt-Börnstein New Series IV/11A3
50.00 50.00 0.00
60
70
e2,1495 80
90
Ni
MSIT ®
Al–Ni–W
450
W Ni Al
Fig. 4: Al-Ni-W. The (Ni) solvus surface
0.00 70.00 30.00
Data / Grid: at.% Axes: at.%
10
20
γ +γ ´ °C 27 12 °C 27 11 °C 27 10 C 7° 92 °C 7 2 8
20
10
γ
W Ni Al
80
30.00 70.00 0.00
90
Ni
Al
Data / Grid: at.%
Fig. 5: Al-Ni-W. Partial isothermal section at 1250°C
Axes: at.%
20
80
40
60
β
60
40
α +β γ'
α+β +γ '
80
20
α+γ +γ ' α
W
MSIT®
γ
α+γ 20
40
60
80
Ni
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr
451
Aluminium – Nickel – Zirconium Gautam Ghosh Literature Data [1966Mar1] was the first to report the isothermal section of the Al corner at 800°C. They prepared 99 ternary alloys, containing up to 25 at.% Zr and 75 at.% Ni, in an arc furnace under Ar atmosphere using elemental metals of the following purity: AV-000 grade Al (99.997 mass%), very pure grade Ni (99.99 mass%) and iodide Zr (99.96 mass%). The alloys were annealed at 800°C for 700 h in evacuated quartz tubes together with Zr foil followed by quenching in cold water. Later, [1969Bur] investigated the constitutional equilibria of the Zr corner. The authors prepared 150 ternary alloys using metals of the above mentioned purity and the alloys were annealed at 800°C for 2100 h followed by quenching in cold water. In both cases phase analysis was performed by microstructural and X-ray diffraction techniques. The constitutional equilibria of the Ni corner at 1100 and 1000°C were determined by [1983Jay1, 1983Jay2], using metallography, electron microprobe analysis, and X-ray diffraction techniques. These results were reviewed by [1991Nas] and [1993Gho]. Only a brief review of phase equilibria was presented by [1990Kum]. Recently, Miura et al. [1999Miu, 2001Miu] investigated the solid-liquid phase equilibria of Ni-rich ternary alloys using DTA, XRD and SEM-WDS analysis. [2001Miu] prepared three ternary alloys using 99.99 mass% Al, 99.95 mass% Ni, and 99.6 mass% Zr. [1991Mis] determined the solvus boundary of (Ni) using DTA and SEM-EDX analysis. Other recent investigations of the ternary system involve measurement of heat of formation of liquid phase by calorimetry [1999Wit, 1999Zho], prediction of glass-forming ranges [2002Shi, 2003Shi], and synthesis of amorphous alloys by liquid quenching/mechanical alloying, and study their crystallization behavior [1990Bha, 1990Ino, 1995Gaf, 1997Kuh, 1998Sri, 1998Tur, 1999Hel, 2000Ill, 2001Cho, 2003Ele, 2003Yan]. Binary Systems The Al-Ni binary phase diagram is accepted from [2003Sal] and the Al-Zr binary phase diagram is accepted from [2003Sch]. Miura et al. [1999Miu, 2001Miu] determined the liquidus of Ni-rich alloys containing up to 13 at.% Al. Unlike Hilpert et al. [1987Hil], Miura et al. [2001Miu] observed a maximum (1466°C) in the liquidus at about 2 at.% Al. Except for [2001Miu], this feature has not been considered in the thermodynamic modeling of the Al-Ni system [2003Sal]. The Ni-Zr binary phase diagram is accepted from [1984Nas] in which the assessed temperature for invariant reaction L (Ni) + ZrNi5 is 1170°C. However, recent experiment [2001Miu] shows that it occurs at 1196°C. Solid Phases The solid solubility of Ni in (Al) decreases from 0.11 at.% at 639.9°C to 0.01 at.% at 500°C. By rapid solidification processing enhancement of the solid solubility is observed to as much as 7.7 at.% Ni [Mas2]. (Ni) dissolves both Al and Zr. The limit of solubility is given up to 21.3 at.% at the peritectic temperature 1372°C by [2003Sal]. Solution of Al in (Ni) causes a linear increase in the lattice parameter from 352.32 pm for pure Ni to 353.88 pm at 8 at.% Al [1985Mis]. Also, the limit of solubility of Zr 1.6 at.%, and this is associated with a linear increase in the lattice parameter. The rate of increase in the lattice parameter, da/dc, is reported to be 1.0 pm/at.% Zr [1984Och, 1985Mis]. Figure 1 shows the solubility isotherms of (Ni) [1991Mis]. (Zr) can dissolve up to 25.5 at.% Al [Mas] and 2.92 at.% Ni [1984Nas], and the corresponding values for (Zr) are 9.5 at.% Al [Mas] and 0.2 at.% Ni [1984Nas]. None of the binary intermediate phases is reported to dissolve more than 1 at.% of the third element [1966Mar1, 1969Bur] at 800°C. Contrary to [1969Bur], [1971Bla] found that ZrNi5 can dissolve up to
Landolt-Börnstein New Series IV/11A3
MSIT ®
452
Al–Ni–Zr
about 16.7 at.% Al at 800°C by replacing Ni and causing a linear increase in the lattice parameter to 674.8 pm at 16.7 at.% Zr; however, it is not known whether the stability of ZrNi4(Ni,Al) is enhanced by the presence of oxygen. [1959Gua] reported that Ni3Al dissolves very small amount of Zr at 1150°C though no specific value was given. On the other hand [1969Tho] found that, in directionally solidified Ni3Al-Zr2Ni7 eutectic, Ni3Al dissolves about 2.7 at.% Zr as measured by an electron probe microanalyzer. In a review paper [1985Nas] reported that Ni3Al can dissolve about 5 at.% Zr and Zr2Ni7 can dissolve about 11 at.% Al, but the source of such information was not mentioned. Careful X-ray diffraction measurements [1983Och, 1984Och, 1985Mis] showed, however, that Ni3Al can dissolve about 1 at.% Zr at 1000°C. Solution of Zr in Ni3Al causes a linear increase in the lattice parameter, and the rate of increase in the lattice parameter, da/dc, is 0.79 pm/at.% Zr [1984Och]. So far eight ternary phases have been reported. The ternary phase ZrNiAl (-1) was first reported by [1964Mar] and subsequently confirmed by [1966Mar2, 1967Kri, 1968Dwi, 1974Fer]. According to [1968Dwi], the structure of the ZrNiAl phase can be better described by introducing a slight variation in stacking sequence and by doubling the c-parameter. Originally, [1964Sch, 1965Ram] reported that ZrNiAl has the Cu2Mg type structure with lattice parameter a = 735 and 734 pm. Since subsequent investigations confirmed that ZrNiAl has the Fe2P type structure, the ZrNiAl phase as designated by [1964Sch] and [1965Ram] is certainly the 2 phase (ZrNixAl2-x, 0.2 x 0.5) which has the Cu2Mg type structure and similar lattice parameter as confirmed by [1966Mar2] and [1969Bur]. The lattice parameter of the 2 phase increases with increasing Al content [1966Mar2]. The existence of ZrNi2Al (-2) has been confirmed several times [1962Hei, 1964Mar, 1964Sch, 1965Ram]. [1962Hei] reported its structure to be of the CsCl type with a = 302.0 pm, but subsequent investigations confirmed the structure to be of the MnCu2Al type with a = 609 to 612.3 pm [1964Mar, 1964Sch, 1965Ram, 1967Hof]. The ternary phase Zr2NiAl5 (-5) was reported to be present in the as-cast alloy [1965Ram], but was not reported by [1969Bur] in the 800°C isothermal section. Also, [1966Mar1] did not observe the -5 phase in the annealed alloys (900°C for 700 h). Nevertheless, minor impurity levels can significantly influence the stability of the AuCu3 type (or L12) phase [1990Kum]. Originally, ZrNi2Al5 (-6) was designated as Zr3Ni6Al16 [1969Bur]. The crystal structure of Zr5Ni4Al (-7) is not known [1969Bur]. The details of the crystal structures and lattice parameters of stable solid phases are listed in Table 1. Pseudobinary Systems Based on the DTA results and microstructure observations, [2001Miu] proposed that (Ni,Al) and ZrNi5 form a pseudobinary eutectic (e2(max)). Even though the details are not known, the eutectic temperature must be greater than 1196°C [2001Miu]. Ni3Al and Zr2Ni7 form a pseudobinary eutectic (e3(max)) at 1193°C [1969Tho]. This was also confirmed by [1978Hao] and [1983Jay1]. The composition of the eutectic point was claimed by [1969Tho] to be at 10.9Zr-73.4Ni-15.7Al (at.%), however this can not be correct because this point would appear outside of the Ni3Al-Zr2Ni7 section. Solid solubilities of Zr in Ni3Al were also measured by electron microprobe and X-ray analyses after annealing the samples at 1100 and 1000°C [1983Jay1, 1991Nas], and they were found to be 3.8 and 3.1 at.%, respectively. It is believed that Zr resides primarily on the Al-sublattice of Ni3Al [2001Ter]. The solid solubility of Zr in Ni3Al, as measured by electron microprobe on dendrites adjacent to the eutectic, was 2.7 at.% [1969Tho]. Invariant Equilibria Figure 2 shows the tentative reaction scheme for the solidification of Ni-rich alloys. Two saddle points, e2(max) and e3(max), are due to [2001Miu] and [1969Tho], respectively. Since e2(max) feeds the binary eutectic L(Ni)+ZrNi5 at 1196°C [2001Miu], it must occur above 1196°C. [2001Miu] observed both (Ni)+ZrNi5 and Ni3Al+ZrNi5 microstructures in as-cast alloys. Based on these observations, they proposed the existence of (Ni)+Ni3Al+ZrNi5 phase field which is the product of the ternary eutectic reaction E1 estimated to be occurring around 1186°C. It is important to note that the three-phase field, (Ni)+Ni3Al+ZrNi5, has also been observed at 800°C [1966Mar1, 1969Bur]. The invariant reaction U2 gives rise to a three-phase field Ni5Al+ZrNi5+Zr2Ni7 as proposed by [1969Tho] but it was not considered by MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr
453
[1983Jay1] and [1991Nas]. Other three ternary invariant reactions account for the observed microstructures, but their temperatures of occurrence are not known. Liquidus Surface [1983Jay1] and [1991Nas] presented a speculative liquidus surface in the Ni-corner, which involves two ternary eutectic reactions, two transition reactions and a pseudobinary eutectic reaction. They proposed a ternary eutectic reaction L(Ni)+ZrNi5+Ni7Zr2 resulting in a three-phase field (Ni)+ZrNi5+Zr2Ni7 on isothermal sections [1991Nas]. However, this conflicts with the observation of Ni3Al+ZrNi 5 microstructure by [2001Miu]. Therefore, the existence of ternary eutectic reaction L(Ni)+ZrNi5+Zr2Ni7 can be ruled out. Based on the aforementioned results, the probable liquidus surface of the Ni-corner is shown in Fig. 3. [1999Miu] determined the effect of Zr addition on the liquidus temperature of Ni-rich alloys. They found that in alloys containing up to 13 at.% Al addition of Zr increases the liquidus temperature, while in alloys containing more than 13 at.% addition of Zr decreases the liquidus temperature. Isothermal Sections Figures 4 and 5 show the partial isothermal sections in the Ni corner at 1100 and 1000°C [1983Jay1, 1991Nas], respectively. Some of the phase boundaries involving (Ni), Ni3Al, Zr2Ni7 and ZrNi 5 are not known exactly because of limited experimental work and also due to inconsistency between the X-ray data and electron microprobe analysis. A number of amendments in Figs. 4 and 5 have been made to comply with the accepted binary phase diagrams. [1969Tho] observed the Ni3Al-Zr2Ni7 eutectic microstructure in the directionally solidified Ni-15Al-10.9Zr (at.%) alloy. Based on this observation, [1969Tho] proposed a tentative phase diagram of the Ni corner characterized by the presence of a Ni3Al+ZrNi5+Zr2Ni7 phase field and a Ni3Al+Zr2Ni7 phase field, which are consistent with the reaction scheme shown in Fig. 2. [1978Hao] also observed Ni3Al-Zr2Ni7 as-cast eutectic microstructure in the above ternary alloy. Even though some phase boundaries are shown dashed in Figs. 4 and 5, it is important to note that the above ternary alloy falls in the Ni3Al+Zr2Ni7 phase field. There are at least two reasons to doubt the results of [1971Bla] where it is claimed that ZrNi5 and ZrNi4Al form a continuous solid solution. First, [1971Bla] noted a poor agreement between calculated and observed X-ray intensities of AlNi4Zr which was attributed to preferred orientation due to cleavage parallel to (311) plane, and the presence of oxygen in some positions in the structure. Secondly, [1983Jay1] reported that ZrNi4Al is not a single-phase alloy. Figure 6 shows the isothermal section at 800°C [1966Mar1, 1969Bur]. It should be noted that the solid solubility ranges of Ni3Al, Zr2Ni7 and -2 are shown to be drastically reduced between 1000 and 800°C but to be virtually unchanged between 1100 and 1000°C, which may be considered unlikely. The phase relations of the Ni corner at 800°C as reported by [1969Bur] differ from that at 1000°C given by [1983Jay1]. Also, the observation of equilibrium between ZrAl3 and -3 phases at 780°C is inconsistent with the 800°C isothermal section, but it would be consistent if one assumes a transition reaction 2+-6ZrAl3+-3 occurring just below 800°C [1991Nas]. Similarly, the presence of transition reaction at higher temperatures given in Fig. 2 is not necessarily in conflict with Fig. 6. The absence of phase fields -2+NiAl+Zr2Ni7, Ni3Al+ZrNi5+Zr2Ni7 and Ni 3Al+NiAl+Zr2Ni7 in the 800°C isothermal section can be due to various solid-state reactions that may take place between 800 and 1000°C. Thermodynamics [1999Wit] determined the heat of formation of Al-Ni-Zr liquid alloys at 1292 5°C by high-temperature calorimetry. Also, [1999Wit] used an empirical relationship for excess entropy, and derived the Gibbs energy of mixing of liquid alloys. The heat capacity of undercooled Zr60Ni25Al15 liquid alloy was measured by differential scanning calorimeter [1994Zap] and by an adiabatic calorimeter [1999Zho]. Both heat of formation of liquid alloys and the heat capacity of undercooled liquid have been analyzed in terms of an association model [1999Zho, 2000Kru]. A maximum in the capacity near the liquidus temperature was
Landolt-Börnstein New Series IV/11A3
MSIT ®
454
Al–Ni–Zr
attributed to temperature dependent chemical short-range ordering [1999Zho]. Hoch [1995Hoc] proposed an alternate model to describe the specific heat of undercooled liquid alloys. The heat of formation of ternary amorphous alloys has been determined by solution calorimetry [1998Tur] and direct reaction calorimetry [1999Hel]. These results also indicate the evidence of strong chemical short-range ordering in amorphous alloys. [1999Dar] measured the low-temperature (3.2 to 10.3 K) specific heat of ZrNi2Al (-2) using an adiabatic calorimeter, and analyzed the specific heat data in terms of electronic, Debye lattice and Einstein models. The analysis of experimental data yields the Debye temperature D = 5°C. They also calculated the electronic structure by tight-binding linearized muffin-tin orbital (TB-LMTO) method. Their results underscore the importance of electron-phonon coupling on the phase stability. Notes on Materials Properties and Applications [1998Sri] studied the microstructure and hardness of rapidly solidified Zr1Ni10Al89 alloy which was subsequently aged at 150, 250, 350 and 450°C. The as-solidified ribbons, containing nanoscale precipitates of NiAl3, exhibit hardness up to 4.5GPa. Aging treatment results in the precipitation of metastable cubic ZrAl3, and the hardness decreases. They also observed that during aging the metastable ZrAl3, precipitates in (Al) matrix and inside NiAl3 phase. Miscellaneous Existence of metastable phases have been reported in mechanically alloyed specimens, and also during crystallization of amorphous alloys. Mechanical alloying of elemental powders produced an amorphous phase having composition Al-12.5Ni-25Zr (at.%) (i.e. Zr2NiAl 5 or 2), which crystallizes into ZrAl3 (L12) and Zr6Ni8Al15 (-3) upon heating to 780°C [1991Des] which is not the equilibrium state according to Fig. 6. [1995Gaf] reported the formation of an fcc phase with lattice parameter of 460 pm in mechanically alloyed Zrx(NiAl)1-x, 0.05 x 0.5. [2003Ele] observed the formation of two cubic phases during mechanical alloying of Zr60Ni25Al15. The phase that forms first has lattice parameter of 1228.2 pm, and upon further milling it transforms to another cubic phase with lattice parameter of 454.49 pm. They also found that during crystallization both these phases give same end products, viz., Zr5Ni4Al (-7) and Zr6NiAl2 (-4); however, the crystallization temperatures are different. The structure, low-temperature specific heat [1987Yam] and crystallization behavior [1990Bha, 1997Kuh, 2001Cho] of some Al-Ni-Zr metallic glasses have also been reported. Mechanical alloying of elemental powders produced an amorphous phase having composition Al-12.5Ni-25Zr (at.%) (i.e. Zr2NiAl5 or 2), which crystallizes into ZrAl3 and Zr6Ni8Al15 (-3) upon heating to 780°C [1991Des], which is not the equilibrium state according to Fig. 6. [1990Ino] reported the formation of amorphous alloys in the composition range of 3 to 67 at.% Ni and 0 to 37 at.% Al by melt spinning. In these alloys, the difference between glass transition temperature (Tg) and crystallization (Tg) can be as large as 77°C. Also, the reduced glass transition temperature (Tg/Tm) can be as high as 0.64. Shindo et al. [2002Shi, 2003Shi] have employed a quasi-chemical approach to predict the critical composition range for bulk metallic glasses. [1995Gaf] noted that mechanical alloying of (NiAl)1-xZr x, 0.05 x 0.5, does not yield fully amorphous phase. They obtained up to 50% amorphous phase when x = 0.5. On the other hand, [1997Kuh] claimed to obtain fully amorphous phase by mechanical alloying of Zr55Ni25Al20 and by rapid solidification of Zr52Ni26Al22. Also, [1997Kuh] found that the observed phases in fully crystallized specimens do not correspond to the expected equilibrium phases. [2001Cho] studied crystallization of Zr55+xNi25Al20-x, 0 x 10, amorphous alloys prepared by rapid solidification. Based on the change in crystallization temperature, they concluded that local environment of atomic pairs is important for the stability of amorphous and supercooled liquid, and the retardation of crystallization process. These amorphous alloys crystallize to Zr3Al2, (Zr) and ZrNi phases. Once again, these do not correspond to the expected equilibrium phases in Fig. 6. A study of liquid-quenched Zr60Ni25Al15 glassy alloy suggests the presence of high density of quenched-in nuclei [2003Yan].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr
455
References [1959Gua] [1962Hei]
[1964Mar]
[1964Sch] [1965Ram]
[1966Gan] [1966Mar1]
[1966Mar2]
[1967Hof] [1967Kri]
[1968Dwi]
[1969Bur] [1969Tho]
[1970Kri]
[1971Bla] [1974Fer]
[1978Hao] [1982Mar]
Landolt-Börnstein New Series IV/11A3
Guard, R.W., Westbrook, J.H., “Alloying Behavior of Ni3Al ( ´ Phase)”, Trans. AIME, 215, 807-813 (1959) (Crys. Structure, Experimental, 27) Heine, W., Zwicker, U., “Phases of the B2 Type (CsCl) in Ternary Systems Containing Cu and Ni” (in German), Naturwissenschaften, 49, 391 (1962) (Crys. Structure, Experimental, 1) Markiv, V.Ya., Voroshilov, Yu.V., Kripyakevich, P.I., Cherkashin, E.E., “New Compounds of MnCu 2Al and MgZn 2 Types Containing Al and Ga”, Sov. Phys.- Crystallogr., 9, 619-620 (1965), translated from Kristallografiya, 9, 737-738 (1964) (Crys. Structure, Experimental, 4) Schubert, K., Raman, A., Rossteutscher, W., “Some Structure Data on Metallic Phases” (in German), Naturwissenschaften, 51, 506 (1964) (Crys. Structure, Experimental, 0) 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 Alloy Systems” (in German), Z. Metallkd., 56, 99-104 (1965) (Crys. Structure, Experimental, 14) Ganglberger, E., Nowotny, H., Benesovsky, F., “On Some New G-Phases” (in German), Monatsh. Chem., 97, 219-220 (1966) (Crys. Structure, Experimental, 3) Markiv, V.Ya., Matushevskaya, N.F., Rozum, N.S., Kuzma, Yu.B., “Investigation of Al-Rich Zr-Ni-Al Alloys” (in Ukrainian), Izv. Akad. Nauk SSSR, Neorg. Mater., 2, 1581-1585 (1966) (Equi. Diagram, Experimental, #, *, 21) Markiv, V.Ya., Kripyakevich, P.I., “Compounds of the Type R(X'X") 2 in Systems with R= Ti, Zr, Hf; X'= Fe, Co, Ni, Cu; and X" = Al or Ga and Their Crystal Structures”, Sov. Phys.- Crystallogr., 11, 733-738 (1967), translated from Kristallografiya, 11, 859-865 (1966) (Crys. Structure, Experimental, 15) Hofer, G., Stadelmaier, H.H., “Co, Ni and Cu Phases of the Ternary MnCu2Al Type” (in German), Monatsh. Chem., 98, 408-411 (1967) (Crys. Structure, Experimental, 9) Kripyakevich, P.I., Markiv, V.Ya., Melnik, Ya.V., “Crystal Structure of Zr-Ni-Al, Zr-Cu-Ga and Analogous Compounds” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A8), 750-753 (1967) (Crys. Structure, Experimental, 9) Dwight, A.E., Mueller, M.H., Conner, R.A., Downey, J.W., Knott, H., “Ternary Compounds with the Fe2P-Type Structure”, Trans. TMS-AIME, 242, 2075-2080 (1968) (Crys. Structure, Experimental, 14) Burnashova, V.V., Markiv, V.Ya., “A Study of the Zr-Ni-Al System” (in Ukrainian), Vest. Lvov. Univ. (Khim.), (11), 34-37 (1969) (Equi. Diagram, Experimental, #, *, 8) Thompson, E.R., Lemkey, F.D., “Structure and Properties of Ni3Al( ') Eutectic Alloys Produced by Unidirectional Solidification”, Trans. ASM, 62, 140-154 (1969) (Equi. Diagram, Experimental, #, *, 35) Kripyakevich, P.I., Burnashova, V.V., Markiv, V.Ya, “Crystal Structures of the Compounds Zr 6FeAl2, Zr6CoAl2, Zr6NiAl2” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A9), 828-831 (1970) (Crys. Structure, Experimental, 10) Blazina, Z., Ban, Z., “X-Ray Studies in the System ZrNi5-xAlx”, Croat. Chem. Acta., 43, 59-63 (1971) (Crys. Structure, Experimental, 4) Ferro, R., Marazza, R., Rambaldi, G., “Equiatomic Ternary Phases in the Alloys of the Rare Earths with Indium and Nickel or Palladium”, Z. Metallkd., 65, 37-39 (1974) (Crys. Structure, Experimental, 2) Haour, G., Mollard, F., Lux, B., Wright, I.G., “New Eutectics Based on Fe, Co and Ni”, Z. Metallkd., 69, 149-154 (1978) (Equi. Diagram, Experimental, 14) Markiv, V.Ya., Kripyakevich, P.I., Belyavina, N.M., “Crystal Structure of the Compound ZrNi2Al5” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, A3, 76-78 (1982) (Crys. Structure, Experimental, 6)
MSIT ®
456 [1983Jay1] [1983Jay2]
[1983Och] [1984Nas] [1984Och]
[1985Mis]
[1985Nas]
[1987Hil]
[1987Yam]
[1990Bha]
[1990Ino]
[1990Kum]
[1991Des]
[1991Mis] [1991Nas] [1993Gho]
[1994Zap] [1995Gaf]
[1995Hoc]
MSIT®
Al–Ni–Zr Jayanth, C.S., “Phase Equilibria in the Ni-Al-Zr and Ni-Al-V Systems”, M.S. Thesis, Illinois Institute of Technology, (1983) (Equi. Diagram, Experimental, #, *, quoted in [1991Nas] Jayanth, C.S., Nash, P., “Phase Equilibria in the Ni-Rich Region of the Al-Ni-Zr System”, Bennett, L.H., Massalski, T.B., Giessen, B.C, (Eds.), Materials Research Society, Pittsburg, PA, 395-398 (1983) (Experimental, 10) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data of Ni3Al with Ternary Additions”, Bull. P. M. E., (52), 1-17 (1983) (Equi. Diagram, Experimental, Review, 39) Nash, P., Jayanth, C.S., “The Ni-Zr (Nickel-Zirconium) System”, Bull. Alloy Phase Diagrams, 5, 144-148 (1984) (Equi. Diagram, Review, #, *, 38) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions”, Bull. P. M. E., (53), 15-28 (1984) (Crys. Structure, Experimental, 66) Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions with Additions of Taransition and B-Subgroup Elements”, Acta Metall., 33 1161-1169 (1985) (Crys. Structure, Experimental, 64) Nash, P., “Ni-Base Intermetallics for High-Temperature Alloy Design” in “High-Temperature Ordered Intermetallic Alloys”, Koch, C.C., Liu, C.T., Stoloff, N.S., (Eds.), Materials Res. Soc., Pittsburg, PA, 423-427 (1985) (Equi. Diagram, Review, 15) Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremmer, F.J., Nickel, H., “Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch. A, 42A, 1327-1332 (1987) (Equi. Diagram, Experimental, #, 17) Yamada, Y., Iton, Y., Mizutani, U., Shibagaki, N., Tanaka, K., “Low-Temperature Specific Heat and Soft X-Ray Spectroscopic Studies of Ni33Zr67-Based Metallic Glasses Containing H, B, Al and Si”, J. Phys. F: Met. Phys., 17, 2303-2311 (1987) (Experimental, 12) Bhatnagar, A.K., Rhie, K.W., Naugle, D.G., Wolfender, A., Zhang, B.H., Callaway, T.O., Bruton, W.D., Hu, C.-R., “The Effect of Simple Metal (Al, Ga) Addition on the Crystallization and Density of Amorphous Zr-Ni Alloys”, J. Phys.: Condens. Matter., 2, 2625-2636 (1990) (Experimental, 24) Inoue, A., Zhang, T., Masumoto, T., “Zr-Al-Ni Amorphous Alloys with High Glass Transition Temperature and Significant Supercooled Liquid Region”, Mater. Trans., JIM, 31, 177-183 (1990) (Experimental, 12) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35, 293-327 (1990) (Equi. Diagram, Review, #, 158) Desch, P., Schwarz, R.B., Nash, P., “Formation of Metastable L1 2 Phases in Al3Zr and Al-12% X-25% Zr (X = Li, Cr, Fe, Ni, Cu)”, J. Less-Common Met., 168, 69-80 (1991) (Experimental, 25) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination g Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng., A146 123-130 (1991) (Equi. Diagram, Experimental, #, *, 5) Nash, P., Pan, Y.Y., “The Al-Ni-Zr System (Aluminum-Nickel- Zirconium)”, J. Alloy Phase Equilibria, 12, 105-113 (1991) (Equi. Diagram, Review, #, *, 49) Ghosh, G., “Aluminium-Nickel-Zirconium”, in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.13050.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 35) Zappel, B., Sommer, F., “Structural Enthalpy Relaxation in the Glass Transition Range”, Mater. Sci. Eng. A, A179-180, 283-287 (1994) (Thermodyn., 10) Gaffet, E., “Structural Investigation of Mechanicall Alloyed (NiAl)1-x(M) x (M=Fe,Zr) Nanocrystalline and Amorphous Phases”, Nano-Structured Mater., 5(4), 393-409 (1995) (Crys. Structure, Experimental, 58) Hoch, M., “The Heat Capacity C p of Undercooled Liquid Metals and Alloys”, Z. Metallkd., 86(8), 557-560 (1995) (Theory, Thermodyn., 14) Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr [1997Jou]
[1997Kuh]
[1998Sri]
[1998Tur]
[1999Dar]
[1999Hel]
[1999Miu]
[1999Wit] [1999Zav]
[1999Zho]
[2000Ill]
[2000Kru] [2001Cho]
[2001Miu]
[2001Ter]
[2002Shi] [2003Ele]
Landolt-Börnstein New Series IV/11A3
457
Joubner, J.-M., Cerny, R., Yvon, K., Latroche, M., Persheron-Guegan, A., “Zirconium-Nickel, Zr7Ni10: Spase Group Revision for the Stoichiometric Phase”, Acta Crystallogr., Sect. C; Cryst. Struct. Commun., C53(11), 1536-1538 (1997) (Crys. Structure, Experimental, 12) Kuhnast, F.A., Held, O., Ragnier, F., Illekova, E., “Calorimetric and Structural Analyses of Mechanically Alloyed and Rapidly Quenched Zn-Ni-Al Alloys”, Mater. Sci. Eng. A, A226-228, 463-467 (1997) (Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 10) Srinivasan, D., Chattopadhyay, K., “Formation and Coarsening of a Nanodispersed Microstructure in Melt Spun Al-Ni-Zr Alloy”, Mater. Sci. Eng. A, A255, 107-116 (1998) (Equi. Diagram, Experimental, 18) Turchanin, A.A., Tomilin, I.A., “Experimental Investigations of the Enthalpies of Formation of Zr-Based Metallic Amorphous Binary and Ternary Alloys”, Ber. Bunsen-Ges. Phys. Chem., 102(9), 1252-1258 (1998) (Experimental, Thermodyn., 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, Theory, Thermodyn., 20) Held, O., Braganti, J.P., Kuhnast, F.A., “Calorimetric and Structural Analysis of the New Phase Al33Ni16Zr 51 Produced by Direct Synhesis and Mechanical Alloying”, J. Alloys Compd., 290, 197-202 (1999) (Experimental, Thermodyn., 15) Miura, S., Hong, Y.-M., Suzuki, T., Mishima, Y., “Liquidus and Solidus Temperatures of Ni-Solid Solution in Ni-Al-X (X: Ti, Zr, and Hf) Ternary Systems”, J. Phase Equilib., 20(3), 193-198 (1999) (Equi. Diagram, Experimental, #, *,11) Witusiewicz, V.T., Sommer, F., “Thermodynamics of liquid Al-Ni-Zr and Al-Cu-Ni-Zr Alloys”, J. Alloys Compd., 289, 152-167 (1999) (Experimental, Thermodyn., 11) Zavaliy, I. Yu., Pecharsky, V.K., Miler, G.J., Akselrud, L.G., “Hydrogenation of Zr6MeX 2 Intermetallic Compounds (Me=Fe, Co, Ni; X=Al, Ga, Sn): Crystallographic and Theoretical Analysis”, J. Alloys Compd., 283, 106-116 (1999) (Crys. Structure, Experimental, 31) Zhou, S.H., Sommer, F., “Calometric Study of Liquid and Undercooled Liquid Al-Ni-Zr Alloys”, J. Non-Cryst. Solids, 250-252, 572-576 (1999) (Calculation, Experimental, Thermodyn., 14) Illekova, E., Jergel, M., Kuhnast, F.-A., “On Structural and Thermal Relaxation in Non-Crystalline Zr-Ni-Al Alloys”, Mater. Sci. Eng. A, A278, 27-35 (2000) (Crys. Structure, Experimental, 32) Krull, H.G., Singh, R.N., Sommer, F., “Generalized Association Model”, Z. Metallkd., 91(5), 356-365 (2000) (Review, Thermodyn., 46) Choi, H.W., Cho, J.H., Kim, J.E., Kim, Y.H., Yang, Y.S., “Calorimetric and Structural Properties of Amorphous Zr-Al-Ni Alloys”, Scr. Mater., 44(8-9), 2027-2030 (2001) (Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 13) Miura, S., Unno, H., Yamazaki, T., Takizawa, S., Mohri. T., “Reinvestigation of Ni-Solid Solution/Liquid Equilibria in Ni-Al Binary and Ni-Al-Zr Ternary Systems”, J. Phase Equilib., 22, 457-462 (2001) (Equi. Diagram, Experimental, #, *, 9) 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) Shindo, T., Waseda, Y., Inoue, A., “Prediction of Glass-Forming Ranges in Zr-Ni-Al Alloys”, Mater. Trans., 43, 2502-2508 (2002) (Thermodyn., Theory, 25) El-Eskandarany, M.S., Saida, J., Inoue, A., “Structural and Calorimetric Evolutions of Mechanically-Induced Solid-State Devitrificated Zr60Ni25Al15 Glassy Alloy Powder”, Acta Mater., 51, 1481-1492 (2003) (Crys. Structure, Experimental, 41)
MSIT ®
Al–Ni–Zr
458 [2003Sal]
[2003Sch]
[2003Shi]
[2003Yan]
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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164) Schuster, J.C., “Al-Zr (Aluminium-Zirconium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram, Assessment, 151) Shindo, T., Waseda, Y., Inoue, A., “Prediction of Critical Compositions for Bulk Glass Formation in La-Based, Cu-Based and Zr-Based Ternary Alloys”, Mater. Trans., 44, 351-352 (2003) (Thermodyn., Theory, 20) Yan, Z., Li, J., He, S., Wang, H., Zhou, Y., “Study of the Crystallization Kinetics of Zr 60Ni25Al15 Glassy Alloy by Differential Scanning Calorimetry”, Mater. Trans., 44, 709-712 (2003) (Experiemntal, 17)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) 660.452 (Ni) 1455 (Zr)(h) 1855 - 863 (Zr)(r) < 863 NiAl3 854 Ni2Al3 1133 NiAl 1638 Ni5Al3 700 Ni3Al 1372 ZrAl3 1580
MSIT®
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cF4 Fm3m Cu cI2 Im3m W hP2 P63/mmc Mg oP16 Pnma NiAl3 hP5 P3m1 Ni2Al3 cP2 Pm3m CsCl oC16 Cmmm Pt5Ga3 cP4 Pm3m AuCu3 cP4 Pm3m Cu3Au
Lattice Parameters Comments/References [pm] a = 404.88
pure Al at 24°C [V-C]
a = 352.32
pure Ni at 20°C [V-C]
a = 356.8
[V-C]
a = 323.2 c = 514.7
[V-C]
a = 661.3 b = 736.7 c = 481.1 a = 402.8 c = 489.1
[2003Sal]
a = 286.00 to 288.72
[2003Sal], solid solubility ranges from 28.7 to 57.9 at.% Al
a = 753.0 b = 661.0 c = 376.0 a = 356.77 to 358.90
[2003Sal], solid solubility ranges from 31.8 to 37.6 at.% Al
a = 399.93 c = 1728.3
[2003Sch]
[2003Sal] 58.7 to 63.9 at.% Al
[2003Sal], solid solubility ranges from 23.7 to 27.4 at.% Al
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr Phase/ Temperature Range [°C] ZrAl2 1660 Zr2Al3 1590 ZrAl 1275 25 Zr5Al4 (h) 1550 - 1000 Zr4Al3 < 1030 Zr3Al2 < 1480 Zr5Al3 (h) < 1400 Zr5Al3 (r)
Zr2Al < 1350 Zr3Al < 1019 ZrNi5 1300 Zr2Ni7 1440
ZrNi3 920 Zr8Ni21 1180
Landolt-Börnstein New Series IV/11A3
Pearson Symbol/ Space Group/ Prototype hP12 P63/mmc MgZn2 oF40 Fdd2 Zr2Al3 oC8 Cmcm CrB hP18 P6/mcm Ti5Ga4 hP7 P6/mmm Zr4Al3 tP20 P42/mnm Zr3Al2 tI32 I4/mcm W5Si3 hP16 P63/mcm Mn5Si3 hP6 P63/mmc Ni2In cP4 Pm3m AuCu3 cF24 F43m AuBe5 mC36 C2/m Zr2Ni7 hP8 P63/mmc Ni3Sn aP29 P1 Hf8Ni21
459
Lattice Parameters Comments/References [pm] a = 528.24 c = 874.82
[2003Sch]
a = 960.1 b = 1390.6 c = 557.4 a = 335.9 b = 1088.7 c = 427.4 a = 844.8 c = 580.5
[2003Sch]
a = 543.3 c = 539.0
[2003Sch]
a = 763.0 c = 699.8
[2003Sch]
a = 1104.4 c = 539.1
[2003Sch]
a = 817.4 c = 569.8
[2003Sch]
a = 489.39 c = 592.83
[2003Sch]
a = 437.2
[2003Sch]
a = 670.64 to 670.72
[1984Nas], 15.0 to 18.0 at.% Zr
a = 469.8 0.9 b = 823.5 1.3 c = 1219.3 1.6 = 95.83° a = 530.9 c = 430.3
[V-C]
a = 647.21 b = 806.45 c = 858.75 = 75.18° = 68.00° = 75.20°
[2003Sch]
[2003Sch]
[1984Nas], 24.5 to 26.0 at.% Zr
[1984Nas]
MSIT ®
Al–Ni–Zr
460 Phase/ Temperature Range [°C] Zr7Ni10 1160 Zr9Ni11 978 - 1170 ZrNi 1260 Zr2Ni 1120
* -1, ZrNiAl
* -2, ZrNi2Al
* -3, Zr6Ni8Al15
* -4, Zr6NiAl2
* -5, Zr2NiAl5
* -6, ZrNi2Al5
* -7, Zr5Ni4Al * 2, ZrNixAl2-x
MSIT®
Pearson Symbol/ Space Group/ Prototype oC68 Cmca Zr7Ni10 tI40 I4/m Zr9Pt11 oC8 Cmcm CrB tI12 I4/mcm CuAl2 hP9 P62m Fe2P cF16 Fm3m MnCu2Al cP2 Pm3m CsCl cF116 Fm3m Th6Mn23 hP9 P62m Zr6CoAl2 cP4 Pm3m AuCu3 tI16 I4/mmm ZrNi2Al5 cF24 Fd3m Cu2Mg
Lattice Parameters Comments/References [pm] a = 1238.1 1.0 b = 918.5 0.5 c = 922.1 1.1 a = 988.0 c = 661.0
[1997Jou], measured on single crystal with 799 refections [1984Nas]
a = 326.8 b = 990.3 c = 410.7 a = 647.7 to 648.3 c = 524.1 to 526.7 a = 691.57 c = 694.12 a = 692.1 c = 346.7 a = 611.47
[V-C]
a = 302.0
[1962Hei]
a = 1208.0
[1966Gan, 1966Mar1]
a = 792.0 c = 334.0 a = 792.8 c = 334.7 a = 406.0
[1969Bur, 1970Kri]
a = 402.3 c = 1444.0 a = 401.0 c = 1441 a = 734.3 to 746.4 a = 746.4 a = 734.3
[V-C]
[1968Dwi] [1974Fer] [1999Dar]
[1999Zav] [1964Sch, 1965Ram], observed in as-cast alloy [1982Mar] [1969Bur] [1969Bur] 0.2 x 0.5 [1966Mar1, 1966Mar2] at x = 0.2 [1966Mar1] at x = 0.5 [1966Mar1]
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr Zr Ni Al
Fig. 1: Al-Ni-Zr. Solubility isotherms of (Ni)
461 0.00 80.00 20.00
Data / Grid: at.% Axes: at.%
1127°C 1027 927 827
10
10
γ
Zr Ni Al
90
20.00 80.00 0.00
Al-Ni 1372 p1 l+(Ni)Ni3Al 1369 e1 lNi3Al+NiAl
Ni
Al-Ni-Zr
? L+NiAl+τ2 ?
A-B-C
Ni-Zr
? L+τ2+Zr2Ni7
L+τ2NiAl+Zr2Ni7 U1 τ2+NiAl+Zr2Ni7
L+NiAl+Zr2Ni7
>1196 e2 L(Ni)+ZrNi5
?
1300 p3 l+Zr2Ni7ZrNi5
1193 e3 LNi3Al+Zr2Ni7
L+Zr2Ni7Ni3Al+ZrNi5 U2
L+Ni3Al+ZrNi5
Ni3Al+ZrNi5+Zr2Ni7
1170 e4 l(Ni)+ZrNi5
1150 L(Ni)+Ni3Al+ZrNi5 E1
(Ni)+Ni3Al+ZrNi5 ?
LNi3Al+NiAl+Zr2Ni7 E2
Ni3Al+NiAl+Zr2Ni7 Fig. 2: Al-Ni-Zr. A tentative reaction scheme for the solidification of Ni-rich alloys Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Zr
462
Al
Data / Grid: at.%
Fig. 3: Al-Ni-Zr. A tentative liquidus surface of the Ni-corner
Axes: at.%
20
80
40
60
U1
60
40
τ2 NiAl
e1 p1
E2 80
Zr8Ni21 Zr2Ni7
20
Zr
40
Fig. 4: Al-Ni-Zr. Partial isothermal section at 1100°C. The dashed lines represent interpolated phase boundaries
e3(max)
E1
U2
60
Zr Ni Al
20
Ni3Al
e2(max) (Ni)
80 p 3
Ni
e ZrNi5 4
0.00 40.00 60.00
Data / Grid: at.% Axes: at.%
10
50
NiAl
20
40
30
30
Ni3Al
τ2
40
20
50
10
(Ni)
Zr2Ni7 Zr Ni Al
MSIT®
60.00 40.00 0.00
50
60
70
80
ZrNi5
90
Ni
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr Zr Ni Al
Fig. 5: Al-Ni-Zr. Partial isothermal section at 1000°C. The dashed lines represent interpolated phase boundaries
463 0.00 40.00 60.00
Data / Grid: at.% Axes: at.%
10
50
NiAl 20
40
30
30
Ni3Al
τ2 40
20
50
10
(Ni) Zr2Ni7 Zr Ni Al
50
60.00 40.00 0.00
60
70
80
Al Fig. 6: Al-Ni-Zr. Isothermal section at 800°C
Axes: at.%
20
80
NiAl3
ZrAl2
τ6 Zr2Al3
Ni
Data / Grid: at.%
L
ZrAl3
90
ZrNi5
Ni2Al3
40
60
λ2
ZrAl Zr4Al3
τ3
Zr3Al2 60
NiAl 40
τ1
Zr2Al
τ2
Zr3Al
τ4
Ni3Al
80
20
τ7 (Ni)
Zr
Landolt-Börnstein New Series IV/11A3
(Zr)
20
Zr2Ni
40
60 Zr8Ni21
ZrNi Zr7Ni10
80
ZrNi3 Zr2Ni7
ZrNi5
Ni
MSIT ®
Al–Ni–Zr
451
Aluminium – Nickel – Zirconium Gautam Ghosh Literature Data [1966Mar1] was the first to report the isothermal section of the Al corner at 800°C. They prepared 99 ternary alloys, containing up to 25 at.% Zr and 75 at.% Ni, in an arc furnace under Ar atmosphere using elemental metals of the following purity: AV-000 grade Al (99.997 mass%), very pure grade Ni (99.99 mass%) and iodide Zr (99.96 mass%). The alloys were annealed at 800°C for 700 h in evacuated quartz tubes together with Zr foil followed by quenching in cold water. Later, [1969Bur] investigated the constitutional equilibria of the Zr corner. The authors prepared 150 ternary alloys using metals of the above mentioned purity and the alloys were annealed at 800°C for 2100 h followed by quenching in cold water. In both cases phase analysis was performed by microstructural and X-ray diffraction techniques. The constitutional equilibria of the Ni corner at 1100 and 1000°C were determined by [1983Jay1, 1983Jay2], using metallography, electron microprobe analysis, and X-ray diffraction techniques. These results were reviewed by [1991Nas] and [1993Gho]. Only a brief review of phase equilibria was presented by [1990Kum]. Recently, Miura et al. [1999Miu, 2001Miu] investigated the solid-liquid phase equilibria of Ni-rich ternary alloys using DTA, XRD and SEM-WDS analysis. [2001Miu] prepared three ternary alloys using 99.99 mass% Al, 99.95 mass% Ni, and 99.6 mass% Zr. [1991Mis] determined the solvus boundary of (Ni) using DTA and SEM-EDX analysis. Other recent investigations of the ternary system involve measurement of heat of formation of liquid phase by calorimetry [1999Wit, 1999Zho], prediction of glass-forming ranges [2002Shi, 2003Shi], and synthesis of amorphous alloys by liquid quenching/mechanical alloying, and study their crystallization behavior [1990Bha, 1990Ino, 1995Gaf, 1997Kuh, 1998Sri, 1998Tur, 1999Hel, 2000Ill, 2001Cho, 2003Ele, 2003Yan]. Binary Systems The Al-Ni binary phase diagram is accepted from [2003Sal] and the Al-Zr binary phase diagram is accepted from [2003Sch]. Miura et al. [1999Miu, 2001Miu] determined the liquidus of Ni-rich alloys containing up to 13 at.% Al. Unlike Hilpert et al. [1987Hil], Miura et al. [2001Miu] observed a maximum (1466°C) in the liquidus at about 2 at.% Al. Except for [2001Miu], this feature has not been considered in the thermodynamic modeling of the Al-Ni system [2003Sal]. The Ni-Zr binary phase diagram is accepted from [1984Nas] in which the assessed temperature for invariant reaction L (Ni) + ZrNi5 is 1170°C. However, recent experiment [2001Miu] shows that it occurs at 1196°C. Solid Phases The solid solubility of Ni in (Al) decreases from 0.11 at.% at 639.9°C to 0.01 at.% at 500°C. By rapid solidification processing enhancement of the solid solubility is observed to as much as 7.7 at.% Ni [Mas2]. (Ni) dissolves both Al and Zr. The limit of solubility is given up to 21.3 at.% at the peritectic temperature 1372°C by [2003Sal]. Solution of Al in (Ni) causes a linear increase in the lattice parameter from 352.32 pm for pure Ni to 353.88 pm at 8 at.% Al [1985Mis]. Also, the limit of solubility of Zr 1.6 at.%, and this is associated with a linear increase in the lattice parameter. The rate of increase in the lattice parameter, da/dc, is reported to be 1.0 pm/at.% Zr [1984Och, 1985Mis]. Figure 1 shows the solubility isotherms of (Ni) [1991Mis]. (Zr) can dissolve up to 25.5 at.% Al [Mas] and 2.92 at.% Ni [1984Nas], and the corresponding values for (Zr) are 9.5 at.% Al [Mas] and 0.2 at.% Ni [1984Nas]. None of the binary intermediate phases is reported to dissolve more than 1 at.% of the third element [1966Mar1, 1969Bur] at 800°C. Contrary to [1969Bur], [1971Bla] found that ZrNi5 can dissolve up to
Landolt-Börnstein New Series IV/11A3
MSIT ®
452
Al–Ni–Zr
about 16.7 at.% Al at 800°C by replacing Ni and causing a linear increase in the lattice parameter to 674.8 pm at 16.7 at.% Zr; however, it is not known whether the stability of ZrNi4(Ni,Al) is enhanced by the presence of oxygen. [1959Gua] reported that Ni3Al dissolves very small amount of Zr at 1150°C though no specific value was given. On the other hand [1969Tho] found that, in directionally solidified Ni3Al-Zr2Ni7 eutectic, Ni3Al dissolves about 2.7 at.% Zr as measured by an electron probe microanalyzer. In a review paper [1985Nas] reported that Ni3Al can dissolve about 5 at.% Zr and Zr2Ni7 can dissolve about 11 at.% Al, but the source of such information was not mentioned. Careful X-ray diffraction measurements [1983Och, 1984Och, 1985Mis] showed, however, that Ni3Al can dissolve about 1 at.% Zr at 1000°C. Solution of Zr in Ni3Al causes a linear increase in the lattice parameter, and the rate of increase in the lattice parameter, da/dc, is 0.79 pm/at.% Zr [1984Och]. So far eight ternary phases have been reported. The ternary phase ZrNiAl (-1) was first reported by [1964Mar] and subsequently confirmed by [1966Mar2, 1967Kri, 1968Dwi, 1974Fer]. According to [1968Dwi], the structure of the ZrNiAl phase can be better described by introducing a slight variation in stacking sequence and by doubling the c-parameter. Originally, [1964Sch, 1965Ram] reported that ZrNiAl has the Cu2Mg type structure with lattice parameter a = 735 and 734 pm. Since subsequent investigations confirmed that ZrNiAl has the Fe2P type structure, the ZrNiAl phase as designated by [1964Sch] and [1965Ram] is certainly the 2 phase (ZrNixAl2-x, 0.2 x 0.5) which has the Cu2Mg type structure and similar lattice parameter as confirmed by [1966Mar2] and [1969Bur]. The lattice parameter of the 2 phase increases with increasing Al content [1966Mar2]. The existence of ZrNi2Al (-2) has been confirmed several times [1962Hei, 1964Mar, 1964Sch, 1965Ram]. [1962Hei] reported its structure to be of the CsCl type with a = 302.0 pm, but subsequent investigations confirmed the structure to be of the MnCu2Al type with a = 609 to 612.3 pm [1964Mar, 1964Sch, 1965Ram, 1967Hof]. The ternary phase Zr2NiAl5 (-5) was reported to be present in the as-cast alloy [1965Ram], but was not reported by [1969Bur] in the 800°C isothermal section. Also, [1966Mar1] did not observe the -5 phase in the annealed alloys (900°C for 700 h). Nevertheless, minor impurity levels can significantly influence the stability of the AuCu3 type (or L12) phase [1990Kum]. Originally, ZrNi2Al5 (-6) was designated as Zr3Ni6Al16 [1969Bur]. The crystal structure of Zr5Ni4Al (-7) is not known [1969Bur]. The details of the crystal structures and lattice parameters of stable solid phases are listed in Table 1. Pseudobinary Systems Based on the DTA results and microstructure observations, [2001Miu] proposed that (Ni,Al) and ZrNi5 form a pseudobinary eutectic (e2(max)). Even though the details are not known, the eutectic temperature must be greater than 1196°C [2001Miu]. Ni3Al and Zr2Ni7 form a pseudobinary eutectic (e3(max)) at 1193°C [1969Tho]. This was also confirmed by [1978Hao] and [1983Jay1]. The composition of the eutectic point was claimed by [1969Tho] to be at 10.9Zr-73.4Ni-15.7Al (at.%), however this can not be correct because this point would appear outside of the Ni3Al-Zr2Ni7 section. Solid solubilities of Zr in Ni3Al were also measured by electron microprobe and X-ray analyses after annealing the samples at 1100 and 1000°C [1983Jay1, 1991Nas], and they were found to be 3.8 and 3.1 at.%, respectively. It is believed that Zr resides primarily on the Al-sublattice of Ni3Al [2001Ter]. The solid solubility of Zr in Ni3Al, as measured by electron microprobe on dendrites adjacent to the eutectic, was 2.7 at.% [1969Tho]. Invariant Equilibria Figure 2 shows the tentative reaction scheme for the solidification of Ni-rich alloys. Two saddle points, e2(max) and e3(max), are due to [2001Miu] and [1969Tho], respectively. Since e2(max) feeds the binary eutectic L(Ni)+ZrNi5 at 1196°C [2001Miu], it must occur above 1196°C. [2001Miu] observed both (Ni)+ZrNi5 and Ni3Al+ZrNi5 microstructures in as-cast alloys. Based on these observations, they proposed the existence of (Ni)+Ni3Al+ZrNi5 phase field which is the product of the ternary eutectic reaction E1 estimated to be occurring around 1186°C. It is important to note that the three-phase field, (Ni)+Ni3Al+ZrNi5, has also been observed at 800°C [1966Mar1, 1969Bur]. The invariant reaction U2 gives rise to a three-phase field Ni5Al+ZrNi5+Zr2Ni7 as proposed by [1969Tho] but it was not considered by MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr
453
[1983Jay1] and [1991Nas]. Other three ternary invariant reactions account for the observed microstructures, but their temperatures of occurrence are not known. Liquidus Surface [1983Jay1] and [1991Nas] presented a speculative liquidus surface in the Ni-corner, which involves two ternary eutectic reactions, two transition reactions and a pseudobinary eutectic reaction. They proposed a ternary eutectic reaction L(Ni)+ZrNi5+Ni7Zr2 resulting in a three-phase field (Ni)+ZrNi5+Zr2Ni7 on isothermal sections [1991Nas]. However, this conflicts with the observation of Ni3Al+ZrNi 5 microstructure by [2001Miu]. Therefore, the existence of ternary eutectic reaction L(Ni)+ZrNi5+Zr2Ni7 can be ruled out. Based on the aforementioned results, the probable liquidus surface of the Ni-corner is shown in Fig. 3. [1999Miu] determined the effect of Zr addition on the liquidus temperature of Ni-rich alloys. They found that in alloys containing up to 13 at.% Al addition of Zr increases the liquidus temperature, while in alloys containing more than 13 at.% addition of Zr decreases the liquidus temperature. Isothermal Sections Figures 4 and 5 show the partial isothermal sections in the Ni corner at 1100 and 1000°C [1983Jay1, 1991Nas], respectively. Some of the phase boundaries involving (Ni), Ni3Al, Zr2Ni7 and ZrNi 5 are not known exactly because of limited experimental work and also due to inconsistency between the X-ray data and electron microprobe analysis. A number of amendments in Figs. 4 and 5 have been made to comply with the accepted binary phase diagrams. [1969Tho] observed the Ni3Al-Zr2Ni7 eutectic microstructure in the directionally solidified Ni-15Al-10.9Zr (at.%) alloy. Based on this observation, [1969Tho] proposed a tentative phase diagram of the Ni corner characterized by the presence of a Ni3Al+ZrNi5+Zr2Ni7 phase field and a Ni3Al+Zr2Ni7 phase field, which are consistent with the reaction scheme shown in Fig. 2. [1978Hao] also observed Ni3Al-Zr2Ni7 as-cast eutectic microstructure in the above ternary alloy. Even though some phase boundaries are shown dashed in Figs. 4 and 5, it is important to note that the above ternary alloy falls in the Ni3Al+Zr2Ni7 phase field. There are at least two reasons to doubt the results of [1971Bla] where it is claimed that ZrNi5 and ZrNi4Al form a continuous solid solution. First, [1971Bla] noted a poor agreement between calculated and observed X-ray intensities of AlNi4Zr which was attributed to preferred orientation due to cleavage parallel to (311) plane, and the presence of oxygen in some positions in the structure. Secondly, [1983Jay1] reported that ZrNi4Al is not a single-phase alloy. Figure 6 shows the isothermal section at 800°C [1966Mar1, 1969Bur]. It should be noted that the solid solubility ranges of Ni3Al, Zr2Ni7 and -2 are shown to be drastically reduced between 1000 and 800°C but to be virtually unchanged between 1100 and 1000°C, which may be considered unlikely. The phase relations of the Ni corner at 800°C as reported by [1969Bur] differ from that at 1000°C given by [1983Jay1]. Also, the observation of equilibrium between ZrAl3 and -3 phases at 780°C is inconsistent with the 800°C isothermal section, but it would be consistent if one assumes a transition reaction 2+-6ZrAl3+-3 occurring just below 800°C [1991Nas]. Similarly, the presence of transition reaction at higher temperatures given in Fig. 2 is not necessarily in conflict with Fig. 6. The absence of phase fields -2+NiAl+Zr2Ni7, Ni3Al+ZrNi5+Zr2Ni7 and Ni 3Al+NiAl+Zr2Ni7 in the 800°C isothermal section can be due to various solid-state reactions that may take place between 800 and 1000°C. Thermodynamics [1999Wit] determined the heat of formation of Al-Ni-Zr liquid alloys at 1292 5°C by high-temperature calorimetry. Also, [1999Wit] used an empirical relationship for excess entropy, and derived the Gibbs energy of mixing of liquid alloys. The heat capacity of undercooled Zr60Ni25Al15 liquid alloy was measured by differential scanning calorimeter [1994Zap] and by an adiabatic calorimeter [1999Zho]. Both heat of formation of liquid alloys and the heat capacity of undercooled liquid have been analyzed in terms of an association model [1999Zho, 2000Kru]. A maximum in the capacity near the liquidus temperature was
Landolt-Börnstein New Series IV/11A3
MSIT ®
454
Al–Ni–Zr
attributed to temperature dependent chemical short-range ordering [1999Zho]. Hoch [1995Hoc] proposed an alternate model to describe the specific heat of undercooled liquid alloys. The heat of formation of ternary amorphous alloys has been determined by solution calorimetry [1998Tur] and direct reaction calorimetry [1999Hel]. These results also indicate the evidence of strong chemical short-range ordering in amorphous alloys. [1999Dar] measured the low-temperature (3.2 to 10.3 K) specific heat of ZrNi2Al (-2) using an adiabatic calorimeter, and analyzed the specific heat data in terms of electronic, Debye lattice and Einstein models. The analysis of experimental data yields the Debye temperature D = 5°C. They also calculated the electronic structure by tight-binding linearized muffin-tin orbital (TB-LMTO) method. Their results underscore the importance of electron-phonon coupling on the phase stability. Notes on Materials Properties and Applications [1998Sri] studied the microstructure and hardness of rapidly solidified Zr1Ni10Al89 alloy which was subsequently aged at 150, 250, 350 and 450°C. The as-solidified ribbons, containing nanoscale precipitates of NiAl3, exhibit hardness up to 4.5GPa. Aging treatment results in the precipitation of metastable cubic ZrAl3, and the hardness decreases. They also observed that during aging the metastable ZrAl3, precipitates in (Al) matrix and inside NiAl3 phase. Miscellaneous Existence of metastable phases have been reported in mechanically alloyed specimens, and also during crystallization of amorphous alloys. Mechanical alloying of elemental powders produced an amorphous phase having composition Al-12.5Ni-25Zr (at.%) (i.e. Zr2NiAl 5 or 2), which crystallizes into ZrAl3 (L12) and Zr6Ni8Al15 (-3) upon heating to 780°C [1991Des] which is not the equilibrium state according to Fig. 6. [1995Gaf] reported the formation of an fcc phase with lattice parameter of 460 pm in mechanically alloyed Zrx(NiAl)1-x, 0.05 x 0.5. [2003Ele] observed the formation of two cubic phases during mechanical alloying of Zr60Ni25Al15. The phase that forms first has lattice parameter of 1228.2 pm, and upon further milling it transforms to another cubic phase with lattice parameter of 454.49 pm. They also found that during crystallization both these phases give same end products, viz., Zr5Ni4Al (-7) and Zr6NiAl2 (-4); however, the crystallization temperatures are different. The structure, low-temperature specific heat [1987Yam] and crystallization behavior [1990Bha, 1997Kuh, 2001Cho] of some Al-Ni-Zr metallic glasses have also been reported. Mechanical alloying of elemental powders produced an amorphous phase having composition Al-12.5Ni-25Zr (at.%) (i.e. Zr2NiAl5 or 2), which crystallizes into ZrAl3 and Zr6Ni8Al15 (-3) upon heating to 780°C [1991Des], which is not the equilibrium state according to Fig. 6. [1990Ino] reported the formation of amorphous alloys in the composition range of 3 to 67 at.% Ni and 0 to 37 at.% Al by melt spinning. In these alloys, the difference between glass transition temperature (Tg) and crystallization (Tg) can be as large as 77°C. Also, the reduced glass transition temperature (Tg/Tm) can be as high as 0.64. Shindo et al. [2002Shi, 2003Shi] have employed a quasi-chemical approach to predict the critical composition range for bulk metallic glasses. [1995Gaf] noted that mechanical alloying of (NiAl)1-xZr x, 0.05 x 0.5, does not yield fully amorphous phase. They obtained up to 50% amorphous phase when x = 0.5. On the other hand, [1997Kuh] claimed to obtain fully amorphous phase by mechanical alloying of Zr55Ni25Al20 and by rapid solidification of Zr52Ni26Al22. Also, [1997Kuh] found that the observed phases in fully crystallized specimens do not correspond to the expected equilibrium phases. [2001Cho] studied crystallization of Zr55+xNi25Al20-x, 0 x 10, amorphous alloys prepared by rapid solidification. Based on the change in crystallization temperature, they concluded that local environment of atomic pairs is important for the stability of amorphous and supercooled liquid, and the retardation of crystallization process. These amorphous alloys crystallize to Zr3Al2, (Zr) and ZrNi phases. Once again, these do not correspond to the expected equilibrium phases in Fig. 6. A study of liquid-quenched Zr60Ni25Al15 glassy alloy suggests the presence of high density of quenched-in nuclei [2003Yan].
MSIT®
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr
455
References [1959Gua] [1962Hei]
[1964Mar]
[1964Sch] [1965Ram]
[1966Gan] [1966Mar1]
[1966Mar2]
[1967Hof] [1967Kri]
[1968Dwi]
[1969Bur] [1969Tho]
[1970Kri]
[1971Bla] [1974Fer]
[1978Hao] [1982Mar]
Landolt-Börnstein New Series IV/11A3
Guard, R.W., Westbrook, J.H., “Alloying Behavior of Ni3Al ( ´ Phase)”, Trans. AIME, 215, 807-813 (1959) (Crys. Structure, Experimental, 27) Heine, W., Zwicker, U., “Phases of the B2 Type (CsCl) in Ternary Systems Containing Cu and Ni” (in German), Naturwissenschaften, 49, 391 (1962) (Crys. Structure, Experimental, 1) Markiv, V.Ya., Voroshilov, Yu.V., Kripyakevich, P.I., Cherkashin, E.E., “New Compounds of MnCu 2Al and MgZn 2 Types Containing Al and Ga”, Sov. Phys.- Crystallogr., 9, 619-620 (1965), translated from Kristallografiya, 9, 737-738 (1964) (Crys. Structure, Experimental, 4) Schubert, K., Raman, A., Rossteutscher, W., “Some Structure Data on Metallic Phases” (in German), Naturwissenschaften, 51, 506 (1964) (Crys. Structure, Experimental, 0) 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 Alloy Systems” (in German), Z. Metallkd., 56, 99-104 (1965) (Crys. Structure, Experimental, 14) Ganglberger, E., Nowotny, H., Benesovsky, F., “On Some New G-Phases” (in German), Monatsh. Chem., 97, 219-220 (1966) (Crys. Structure, Experimental, 3) Markiv, V.Ya., Matushevskaya, N.F., Rozum, N.S., Kuzma, Yu.B., “Investigation of Al-Rich Zr-Ni-Al Alloys” (in Ukrainian), Izv. Akad. Nauk SSSR, Neorg. Mater., 2, 1581-1585 (1966) (Equi. Diagram, Experimental, #, *, 21) Markiv, V.Ya., Kripyakevich, P.I., “Compounds of the Type R(X'X") 2 in Systems with R= Ti, Zr, Hf; X'= Fe, Co, Ni, Cu; and X" = Al or Ga and Their Crystal Structures”, Sov. Phys.- Crystallogr., 11, 733-738 (1967), translated from Kristallografiya, 11, 859-865 (1966) (Crys. Structure, Experimental, 15) Hofer, G., Stadelmaier, H.H., “Co, Ni and Cu Phases of the Ternary MnCu2Al Type” (in German), Monatsh. Chem., 98, 408-411 (1967) (Crys. Structure, Experimental, 9) Kripyakevich, P.I., Markiv, V.Ya., Melnik, Ya.V., “Crystal Structure of Zr-Ni-Al, Zr-Cu-Ga and Analogous Compounds” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A8), 750-753 (1967) (Crys. Structure, Experimental, 9) Dwight, A.E., Mueller, M.H., Conner, R.A., Downey, J.W., Knott, H., “Ternary Compounds with the Fe2P-Type Structure”, Trans. TMS-AIME, 242, 2075-2080 (1968) (Crys. Structure, Experimental, 14) Burnashova, V.V., Markiv, V.Ya., “A Study of the Zr-Ni-Al System” (in Ukrainian), Vest. Lvov. Univ. (Khim.), (11), 34-37 (1969) (Equi. Diagram, Experimental, #, *, 8) Thompson, E.R., Lemkey, F.D., “Structure and Properties of Ni3Al( ') Eutectic Alloys Produced by Unidirectional Solidification”, Trans. ASM, 62, 140-154 (1969) (Equi. Diagram, Experimental, #, *, 35) Kripyakevich, P.I., Burnashova, V.V., Markiv, V.Ya, “Crystal Structures of the Compounds Zr 6FeAl2, Zr6CoAl2, Zr6NiAl2” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, (A9), 828-831 (1970) (Crys. Structure, Experimental, 10) Blazina, Z., Ban, Z., “X-Ray Studies in the System ZrNi5-xAlx”, Croat. Chem. Acta., 43, 59-63 (1971) (Crys. Structure, Experimental, 4) Ferro, R., Marazza, R., Rambaldi, G., “Equiatomic Ternary Phases in the Alloys of the Rare Earths with Indium and Nickel or Palladium”, Z. Metallkd., 65, 37-39 (1974) (Crys. Structure, Experimental, 2) Haour, G., Mollard, F., Lux, B., Wright, I.G., “New Eutectics Based on Fe, Co and Ni”, Z. Metallkd., 69, 149-154 (1978) (Equi. Diagram, Experimental, 14) Markiv, V.Ya., Kripyakevich, P.I., Belyavina, N.M., “Crystal Structure of the Compound ZrNi2Al5” (in Ukrainian), Dopov. Akad. Nauk Ukr. RSR, A3, 76-78 (1982) (Crys. Structure, Experimental, 6)
MSIT ®
456 [1983Jay1] [1983Jay2]
[1983Och] [1984Nas] [1984Och]
[1985Mis]
[1985Nas]
[1987Hil]
[1987Yam]
[1990Bha]
[1990Ino]
[1990Kum]
[1991Des]
[1991Mis] [1991Nas] [1993Gho]
[1994Zap] [1995Gaf]
[1995Hoc]
MSIT®
Al–Ni–Zr Jayanth, C.S., “Phase Equilibria in the Ni-Al-Zr and Ni-Al-V Systems”, M.S. Thesis, Illinois Institute of Technology, (1983) (Equi. Diagram, Experimental, #, *, quoted in [1991Nas] Jayanth, C.S., Nash, P., “Phase Equilibria in the Ni-Rich Region of the Al-Ni-Zr System”, Bennett, L.H., Massalski, T.B., Giessen, B.C, (Eds.), Materials Research Society, Pittsburg, PA, 395-398 (1983) (Experimental, 10) Ochiai, S., Oya, Y., Suzuki, T., “Solubility Data of Ni3Al with Ternary Additions”, Bull. P. M. E., (52), 1-17 (1983) (Equi. Diagram, Experimental, Review, 39) Nash, P., Jayanth, C.S., “The Ni-Zr (Nickel-Zirconium) System”, Bull. Alloy Phase Diagrams, 5, 144-148 (1984) (Equi. Diagram, Review, #, *, 38) Ochiai, S., Mishima, Y., Suzuki, T., “Lattice Parameter Data of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions”, Bull. P. M. E., (53), 15-28 (1984) (Crys. Structure, Experimental, 66) Mishima, Y., Ochiai, S., Suzuki, T., “Lattice Parameters of Ni( ), Ni3Al( ') and Ni3Ga( ') Solid Solutions with Additions of Taransition and B-Subgroup Elements”, Acta Metall., 33 1161-1169 (1985) (Crys. Structure, Experimental, 64) Nash, P., “Ni-Base Intermetallics for High-Temperature Alloy Design” in “High-Temperature Ordered Intermetallic Alloys”, Koch, C.C., Liu, C.T., Stoloff, N.S., (Eds.), Materials Res. Soc., Pittsburg, PA, 423-427 (1985) (Equi. Diagram, Review, 15) Hilpert, K., Kobertz, D., Venugopal, V., Miller, M., Gerads, H., Bremmer, F.J., Nickel, H., “Phase Diagram Studies on the Al-Ni System”, Z. Naturforsch. A, 42A, 1327-1332 (1987) (Equi. Diagram, Experimental, #, 17) Yamada, Y., Iton, Y., Mizutani, U., Shibagaki, N., Tanaka, K., “Low-Temperature Specific Heat and Soft X-Ray Spectroscopic Studies of Ni33Zr67-Based Metallic Glasses Containing H, B, Al and Si”, J. Phys. F: Met. Phys., 17, 2303-2311 (1987) (Experimental, 12) Bhatnagar, A.K., Rhie, K.W., Naugle, D.G., Wolfender, A., Zhang, B.H., Callaway, T.O., Bruton, W.D., Hu, C.-R., “The Effect of Simple Metal (Al, Ga) Addition on the Crystallization and Density of Amorphous Zr-Ni Alloys”, J. Phys.: Condens. Matter., 2, 2625-2636 (1990) (Experimental, 24) Inoue, A., Zhang, T., Masumoto, T., “Zr-Al-Ni Amorphous Alloys with High Glass Transition Temperature and Significant Supercooled Liquid Region”, Mater. Trans., JIM, 31, 177-183 (1990) (Experimental, 12) Kumar, K.S., “Ternary Intermetallics in Aluminium-Refractory Metal-X Systems (X = V, Cr, Mn, Fe, Co, Ni, Cu, Zn)”, Int. Mater. Rev., 35, 293-327 (1990) (Equi. Diagram, Review, #, 158) Desch, P., Schwarz, R.B., Nash, P., “Formation of Metastable L1 2 Phases in Al3Zr and Al-12% X-25% Zr (X = Li, Cr, Fe, Ni, Cu)”, J. Less-Common Met., 168, 69-80 (1991) (Experimental, 25) Mishima, Y., Hong, Y.M., Suzuki, T., “Determination g Solvus Surface in Ni-Al-X Ternary Systems”, Mater. Sci. Eng., A146 123-130 (1991) (Equi. Diagram, Experimental, #, *, 5) Nash, P., Pan, Y.Y., “The Al-Ni-Zr System (Aluminum-Nickel- Zirconium)”, J. Alloy Phase Equilibria, 12, 105-113 (1991) (Equi. Diagram, Review, #, *, 49) Ghosh, G., “Aluminium-Nickel-Zirconium”, in MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.13050.1.20, (1993) (Crys. Structure, Equi. Diagram, Assessment, 35) Zappel, B., Sommer, F., “Structural Enthalpy Relaxation in the Glass Transition Range”, Mater. Sci. Eng. A, A179-180, 283-287 (1994) (Thermodyn., 10) Gaffet, E., “Structural Investigation of Mechanicall Alloyed (NiAl)1-x(M) x (M=Fe,Zr) Nanocrystalline and Amorphous Phases”, Nano-Structured Mater., 5(4), 393-409 (1995) (Crys. Structure, Experimental, 58) Hoch, M., “The Heat Capacity C p of Undercooled Liquid Metals and Alloys”, Z. Metallkd., 86(8), 557-560 (1995) (Theory, Thermodyn., 14) Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr [1997Jou]
[1997Kuh]
[1998Sri]
[1998Tur]
[1999Dar]
[1999Hel]
[1999Miu]
[1999Wit] [1999Zav]
[1999Zho]
[2000Ill]
[2000Kru] [2001Cho]
[2001Miu]
[2001Ter]
[2002Shi] [2003Ele]
Landolt-Börnstein New Series IV/11A3
457
Joubner, J.-M., Cerny, R., Yvon, K., Latroche, M., Persheron-Guegan, A., “Zirconium-Nickel, Zr7Ni10: Spase Group Revision for the Stoichiometric Phase”, Acta Crystallogr., Sect. C; Cryst. Struct. Commun., C53(11), 1536-1538 (1997) (Crys. Structure, Experimental, 12) Kuhnast, F.A., Held, O., Ragnier, F., Illekova, E., “Calorimetric and Structural Analyses of Mechanically Alloyed and Rapidly Quenched Zn-Ni-Al Alloys”, Mater. Sci. Eng. A, A226-228, 463-467 (1997) (Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 10) Srinivasan, D., Chattopadhyay, K., “Formation and Coarsening of a Nanodispersed Microstructure in Melt Spun Al-Ni-Zr Alloy”, Mater. Sci. Eng. A, A255, 107-116 (1998) (Equi. Diagram, Experimental, 18) Turchanin, A.A., Tomilin, I.A., “Experimental Investigations of the Enthalpies of Formation of Zr-Based Metallic Amorphous Binary and Ternary Alloys”, Ber. Bunsen-Ges. Phys. Chem., 102(9), 1252-1258 (1998) (Experimental, Thermodyn., 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, Theory, Thermodyn., 20) Held, O., Braganti, J.P., Kuhnast, F.A., “Calorimetric and Structural Analysis of the New Phase Al33Ni16Zr 51 Produced by Direct Synhesis and Mechanical Alloying”, J. Alloys Compd., 290, 197-202 (1999) (Experimental, Thermodyn., 15) Miura, S., Hong, Y.-M., Suzuki, T., Mishima, Y., “Liquidus and Solidus Temperatures of Ni-Solid Solution in Ni-Al-X (X: Ti, Zr, and Hf) Ternary Systems”, J. Phase Equilib., 20(3), 193-198 (1999) (Equi. Diagram, Experimental, #, *,11) Witusiewicz, V.T., Sommer, F., “Thermodynamics of liquid Al-Ni-Zr and Al-Cu-Ni-Zr Alloys”, J. Alloys Compd., 289, 152-167 (1999) (Experimental, Thermodyn., 11) Zavaliy, I. Yu., Pecharsky, V.K., Miler, G.J., Akselrud, L.G., “Hydrogenation of Zr6MeX 2 Intermetallic Compounds (Me=Fe, Co, Ni; X=Al, Ga, Sn): Crystallographic and Theoretical Analysis”, J. Alloys Compd., 283, 106-116 (1999) (Crys. Structure, Experimental, 31) Zhou, S.H., Sommer, F., “Calometric Study of Liquid and Undercooled Liquid Al-Ni-Zr Alloys”, J. Non-Cryst. Solids, 250-252, 572-576 (1999) (Calculation, Experimental, Thermodyn., 14) Illekova, E., Jergel, M., Kuhnast, F.-A., “On Structural and Thermal Relaxation in Non-Crystalline Zr-Ni-Al Alloys”, Mater. Sci. Eng. A, A278, 27-35 (2000) (Crys. Structure, Experimental, 32) Krull, H.G., Singh, R.N., Sommer, F., “Generalized Association Model”, Z. Metallkd., 91(5), 356-365 (2000) (Review, Thermodyn., 46) Choi, H.W., Cho, J.H., Kim, J.E., Kim, Y.H., Yang, Y.S., “Calorimetric and Structural Properties of Amorphous Zr-Al-Ni Alloys”, Scr. Mater., 44(8-9), 2027-2030 (2001) (Crys. Structure, Equi. Diagram, Experimental, Thermodyn., 13) Miura, S., Unno, H., Yamazaki, T., Takizawa, S., Mohri. T., “Reinvestigation of Ni-Solid Solution/Liquid Equilibria in Ni-Al Binary and Ni-Al-Zr Ternary Systems”, J. Phase Equilib., 22, 457-462 (2001) (Equi. Diagram, Experimental, #, *, 9) 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) Shindo, T., Waseda, Y., Inoue, A., “Prediction of Glass-Forming Ranges in Zr-Ni-Al Alloys”, Mater. Trans., 43, 2502-2508 (2002) (Thermodyn., Theory, 25) El-Eskandarany, M.S., Saida, J., Inoue, A., “Structural and Calorimetric Evolutions of Mechanically-Induced Solid-State Devitrificated Zr60Ni25Al15 Glassy Alloy Powder”, Acta Mater., 51, 1481-1492 (2003) (Crys. Structure, Experimental, 41)
MSIT ®
Al–Ni–Zr
458 [2003Sal]
[2003Sch]
[2003Shi]
[2003Yan]
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; to be published, (2003) (Crys. Structure, Equi. Diagram, Assessment, 164) Schuster, J.C., “Al-Zr (Aluminium-Zirconium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; submitted for publication (2003) (Crys. Structure, Equi. Diagram, Assessment, 151) Shindo, T., Waseda, Y., Inoue, A., “Prediction of Critical Compositions for Bulk Glass Formation in La-Based, Cu-Based and Zr-Based Ternary Alloys”, Mater. Trans., 44, 351-352 (2003) (Thermodyn., Theory, 20) Yan, Z., Li, J., He, S., Wang, H., Zhou, Y., “Study of the Crystallization Kinetics of Zr 60Ni25Al15 Glassy Alloy by Differential Scanning Calorimetry”, Mater. Trans., 44, 709-712 (2003) (Experiemntal, 17)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C] (Al) 660.452 (Ni) 1455 (Zr)(h) 1855 - 863 (Zr)(r) < 863 NiAl3 854 Ni2Al3 1133 NiAl 1638 Ni5Al3 700 Ni3Al 1372 ZrAl3 1580
MSIT®
Pearson Symbol/ Space Group/ Prototype cF4 Fm3m Cu cF4 Fm3m Cu cI2 Im3m W hP2 P63/mmc Mg oP16 Pnma NiAl3 hP5 P3m1 Ni2Al3 cP2 Pm3m CsCl oC16 Cmmm Pt5Ga3 cP4 Pm3m AuCu3 cP4 Pm3m Cu3Au
Lattice Parameters Comments/References [pm] a = 404.88
pure Al at 24°C [V-C]
a = 352.32
pure Ni at 20°C [V-C]
a = 356.8
[V-C]
a = 323.2 c = 514.7
[V-C]
a = 661.3 b = 736.7 c = 481.1 a = 402.8 c = 489.1
[2003Sal]
a = 286.00 to 288.72
[2003Sal], solid solubility ranges from 28.7 to 57.9 at.% Al
a = 753.0 b = 661.0 c = 376.0 a = 356.77 to 358.90
[2003Sal], solid solubility ranges from 31.8 to 37.6 at.% Al
a = 399.93 c = 1728.3
[2003Sch]
[2003Sal] 58.7 to 63.9 at.% Al
[2003Sal], solid solubility ranges from 23.7 to 27.4 at.% Al
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr Phase/ Temperature Range [°C] ZrAl2 1660 Zr2Al3 1590 ZrAl 1275 25 Zr5Al4 (h) 1550 - 1000 Zr4Al3 < 1030 Zr3Al2 < 1480 Zr5Al3 (h) < 1400 Zr5Al3 (r)
Zr2Al < 1350 Zr3Al < 1019 ZrNi5 1300 Zr2Ni7 1440
ZrNi3 920 Zr8Ni21 1180
Landolt-Börnstein New Series IV/11A3
Pearson Symbol/ Space Group/ Prototype hP12 P63/mmc MgZn2 oF40 Fdd2 Zr2Al3 oC8 Cmcm CrB hP18 P6/mcm Ti5Ga4 hP7 P6/mmm Zr4Al3 tP20 P42/mnm Zr3Al2 tI32 I4/mcm W5Si3 hP16 P63/mcm Mn5Si3 hP6 P63/mmc Ni2In cP4 Pm3m AuCu3 cF24 F43m AuBe5 mC36 C2/m Zr2Ni7 hP8 P63/mmc Ni3Sn aP29 P1 Hf8Ni21
459
Lattice Parameters Comments/References [pm] a = 528.24 c = 874.82
[2003Sch]
a = 960.1 b = 1390.6 c = 557.4 a = 335.9 b = 1088.7 c = 427.4 a = 844.8 c = 580.5
[2003Sch]
a = 543.3 c = 539.0
[2003Sch]
a = 763.0 c = 699.8
[2003Sch]
a = 1104.4 c = 539.1
[2003Sch]
a = 817.4 c = 569.8
[2003Sch]
a = 489.39 c = 592.83
[2003Sch]
a = 437.2
[2003Sch]
a = 670.64 to 670.72
[1984Nas], 15.0 to 18.0 at.% Zr
a = 469.8 0.9 b = 823.5 1.3 c = 1219.3 1.6 = 95.83° a = 530.9 c = 430.3
[V-C]
a = 647.21 b = 806.45 c = 858.75 = 75.18° = 68.00° = 75.20°
[2003Sch]
[2003Sch]
[1984Nas], 24.5 to 26.0 at.% Zr
[1984Nas]
MSIT ®
Al–Ni–Zr
460 Phase/ Temperature Range [°C] Zr7Ni10 1160 Zr9Ni11 978 - 1170 ZrNi 1260 Zr2Ni 1120
* -1, ZrNiAl
* -2, ZrNi2Al
* -3, Zr6Ni8Al15
* -4, Zr6NiAl2
* -5, Zr2NiAl5
* -6, ZrNi2Al5
* -7, Zr5Ni4Al * 2, ZrNixAl2-x
MSIT®
Pearson Symbol/ Space Group/ Prototype oC68 Cmca Zr7Ni10 tI40 I4/m Zr9Pt11 oC8 Cmcm CrB tI12 I4/mcm CuAl2 hP9 P62m Fe2P cF16 Fm3m MnCu2Al cP2 Pm3m CsCl cF116 Fm3m Th6Mn23 hP9 P62m Zr6CoAl2 cP4 Pm3m AuCu3 tI16 I4/mmm ZrNi2Al5 cF24 Fd3m Cu2Mg
Lattice Parameters Comments/References [pm] a = 1238.1 1.0 b = 918.5 0.5 c = 922.1 1.1 a = 988.0 c = 661.0
[1997Jou], measured on single crystal with 799 refections [1984Nas]
a = 326.8 b = 990.3 c = 410.7 a = 647.7 to 648.3 c = 524.1 to 526.7 a = 691.57 c = 694.12 a = 692.1 c = 346.7 a = 611.47
[V-C]
a = 302.0
[1962Hei]
a = 1208.0
[1966Gan, 1966Mar1]
a = 792.0 c = 334.0 a = 792.8 c = 334.7 a = 406.0
[1969Bur, 1970Kri]
a = 402.3 c = 1444.0 a = 401.0 c = 1441 a = 734.3 to 746.4 a = 746.4 a = 734.3
[V-C]
[1968Dwi] [1974Fer] [1999Dar]
[1999Zav] [1964Sch, 1965Ram], observed in as-cast alloy [1982Mar] [1969Bur] [1969Bur] 0.2 x 0.5 [1966Mar1, 1966Mar2] at x = 0.2 [1966Mar1] at x = 0.5 [1966Mar1]
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr Zr Ni Al
Fig. 1: Al-Ni-Zr. Solubility isotherms of (Ni)
461 0.00 80.00 20.00
Data / Grid: at.% Axes: at.%
1127°C 1027 927 827
10
10
γ
Zr Ni Al
90
20.00 80.00 0.00
Al-Ni 1372 p1 l+(Ni)Ni3Al 1369 e1 lNi3Al+NiAl
Ni
Al-Ni-Zr
? L+NiAl+τ2 ?
A-B-C
Ni-Zr
? L+τ2+Zr2Ni7
L+τ2NiAl+Zr2Ni7 U1 τ2+NiAl+Zr2Ni7
L+NiAl+Zr2Ni7
>1196 e2 L(Ni)+ZrNi5
?
1300 p3 l+Zr2Ni7ZrNi5
1193 e3 LNi3Al+Zr2Ni7
L+Zr2Ni7Ni3Al+ZrNi5 U2
L+Ni3Al+ZrNi5
Ni3Al+ZrNi5+Zr2Ni7
1170 e4 l(Ni)+ZrNi5
1150 L(Ni)+Ni3Al+ZrNi5 E1
(Ni)+Ni3Al+ZrNi5 ?
LNi3Al+NiAl+Zr2Ni7 E2
Ni3Al+NiAl+Zr2Ni7 Fig. 2: Al-Ni-Zr. A tentative reaction scheme for the solidification of Ni-rich alloys Landolt-Börnstein New Series IV/11A3
MSIT ®
Al–Ni–Zr
462
Al
Data / Grid: at.%
Fig. 3: Al-Ni-Zr. A tentative liquidus surface of the Ni-corner
Axes: at.%
20
80
40
60
U1
60
40
τ2 NiAl
e1 p1
E2 80
Zr8Ni21 Zr2Ni7
20
Zr
40
Fig. 4: Al-Ni-Zr. Partial isothermal section at 1100°C. The dashed lines represent interpolated phase boundaries
e3(max)
E1
U2
60
Zr Ni Al
20
Ni3Al
e2(max) (Ni)
80 p 3
Ni
e ZrNi5 4
0.00 40.00 60.00
Data / Grid: at.% Axes: at.%
10
50
NiAl
20
40
30
30
Ni3Al
τ2
40
20
50
10
(Ni)
Zr2Ni7 Zr Ni Al
MSIT®
60.00 40.00 0.00
50
60
70
80
ZrNi5
90
Ni
Landolt-Börnstein New Series IV/11A3
Al–Ni–Zr Zr Ni Al
Fig. 5: Al-Ni-Zr. Partial isothermal section at 1000°C. The dashed lines represent interpolated phase boundaries
463 0.00 40.00 60.00
Data / Grid: at.% Axes: at.%
10
50
NiAl 20
40
30
30
Ni3Al
τ2 40
20
50
10
(Ni) Zr2Ni7 Zr Ni Al
50
60.00 40.00 0.00
60
70
80
Al Fig. 6: Al-Ni-Zr. Isothermal section at 800°C
Axes: at.%
20
80
NiAl3
ZrAl2
τ6 Zr2Al3
Ni
Data / Grid: at.%
L
ZrAl3
90
ZrNi5
Ni2Al3
40
60
λ2
ZrAl Zr4Al3
τ3
Zr3Al2 60
NiAl 40
τ1
Zr2Al
τ2
Zr3Al
τ4
Ni3Al
80
20
τ7 (Ni)
Zr
Landolt-Börnstein New Series IV/11A3
(Zr)
20
Zr2Ni
40
60 Zr8Ni21
ZrNi Zr7Ni10
80
ZrNi3 Zr2Ni7
ZrNi5
Ni
MSIT ®
