Selected Systems from Al-Fe-V to Al-Ni-Zr (Landolt-Börnstein: Numerical Data and Functional Relationships in Science and Technology - New Series Physical Chemistry)

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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®


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.


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


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®


Introduction

XVII

General References [C.A.] [Curr.Cont.] [E] [G] [H] [L-B]

[Mas] [Mas2] [P] [S] [V-C] [V-C2]


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


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®


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]


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)


Al–Fe–V [1995Ant]

[1997And]

[1997Nis1]

[1997Nis2]

[1998Kat]

[1998Weh] [2000Ino]

[2000Kat]

[2000Sah]

[2000Zar]

[2001Fen]

[2001Han] [2001Kan]

[2001Lue]

[2001Mak]

[2001Nis1]

[2001Nis2]


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]


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


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


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


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


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].


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)


Al–Fe–Y [1976Mcn]

[1977Mur]

[1978Bus]

[1980Fel] [1981Fel]

[1982Erm]

[1983Erm1]

[1983Erm2]

[1985Gan]

[1986Plu]

[1986Sec] [1987Ric] [1988Che]

[1988Cun]

[1989Gsc]

[1992Gri]


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]


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]


Al–Fe–Y

17

Phase/ Temperature Range [°C]



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


oC8 Cmcm CrB

MSIT ®

Al–Fe–Y

18 Phase/ Temperature Range [°C]



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]


Al–Fe–Y Phase/ Temperature Range [°C]


* -1, Y(Fe1-xAlx)2 YFeAl

hP12 P63/mmc MgZn2

* -2, Y(FexAl1-x)12 YFe4Al8

tI26 I4/mmm ThMn12

YFe6Al6

* -3, YFe2Al10


oP52 Cmcm YbFe2Al10


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


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


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®


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®


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]


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]


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)


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


Comments/References

[2003Pis], solid solubility ranges from 54.5 to 62.5 at.% Al

MSIT ®

Al–Fe–Zn


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®



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


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


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


Al–Fe–Zn

33

Al


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


(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


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


(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


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


Al–Fe–Zn Fe Zn Al


35 0.00 99.45 0.55


L+ δ

L+η

+η

L+δ L+ζ+δ

L

L+ζ

Fe Zn Al

0.55 99.45 0.00


Zn Fe Zn Al

0.00 99.45 0.55


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


Γ1

Γ

80

0.00 80.00 20.00

ζ

Zn


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


Al–Fe–Zn

37

Fe Zn Al


0.00 99.45 0.55


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


0.00 60.00 40.00


(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.%


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.%


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


Al–Fe–Zn

39

Al

Data / Grid: at.%


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


(Al)

20

Fe4Al13

80

(Al)+(Zn)+Fe4Al13

40

60

60

40

80

20

(Zn)

Fe


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


Al–Fe–Zn


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


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


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®


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]


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


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]


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


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]


Al–Fe–Zr Zr Fe Al

Fig. 1: Al-Fe-Zr. Partial isothermal section at 1200°C

49 60.00 0.00 40.00


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


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


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


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®


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]


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)


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


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


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]


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


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®


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]


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)




(Li) < 180.6

cI2 Im3m W

a = 351.0


(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


MSIT ®

Al–H–Li




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]


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)


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®


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) = (dlnH/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

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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]


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)

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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)




(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]

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69




, 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


(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

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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


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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®


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) = dlnH / dx(Al) = +5.51 in Ti: JH(Al) = dlnH / 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].


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]


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


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


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


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


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


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


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%)


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®


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®


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]


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., Melńik, 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]


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


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


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]


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


[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


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


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


Al–Hf–Ni Hf Ni Al

Fig. 3: Al-Hf-Ni. Liquidus surface of the Ni corner

91 0.00 70.00 30.00


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


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


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


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


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®


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)


Al–Li–Mg [1956Lev] [1956Row]

[1965Fri]

[1968Pau]

[1971Fri1]

[1971Fri2]

[1972Sam]

[1973Dri1]

[1973Dri2]

[1973Tho] [1976Pad]

[1977Dri]

[1977Sab] [1978Sab]

[1979Gei]

[1979Vos]

[1980Sch1]


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]


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


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


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


-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


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


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



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


Al–Li–Mg

107

Al Fig. 10: Al-Li-Mg. Isothermal section at 400°C


(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


Mg


(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


20

Li3Al2+η+(Li)

(Mg)

(Li) 20

40

60

80

Mg

MSIT ®

Al–Li–Mg

108



(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


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


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®


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 E2É3´, 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®


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]


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)


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)




(Li) < 180.6

cI2 Im3m W

a = 351.0


(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


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]


MSIT ®

Al–Li–Si




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]


Al–Li–Si

117




* -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


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


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


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®


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


[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


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


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)


Al

(Al)

Axes: at.%

20

80

40

60

Li12Si760

40

τ1

Li7Si3

τ3

Li13Si4

τ2

80

20

Li22Si5

L L

Li


20

40

LiAl

60

L

80

(Al)

Al

MSIT ®

Al–Li–Si

122

Si

Data / Grid: at.%

(Si)


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


Al–Li–Si

123

Si

Data / Grid: at.%

(Si)


Axes: at.%

20

80

40

60

60

40

τ1

τ3

80

20

L

Li


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®


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.


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)


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)




(Li) < 180.6

cI2 Im3m W

a = 351.0


(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


´, 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]


MSIT ®

Al–Li–Zn




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®


Al–Li–Zn T [°C]

Reaction

Type

129 Phase


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.%


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


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


(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


20

40

60

80

Zn

MSIT ®

Al–Li–Zn

132

Al Fig. 5: Al-Li-Zn. The isothermal section at 193°C


(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.%


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


experimental 20

calculated

80

40

60

τ2

-1.5 -1.0

9

60

40

τ1 -3.5

τ3

80

Li


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.

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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]


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)

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136 [1986Sak]

[1987Flo] [1987Tos] [1987Vec]

[1988Eun]

[1989Che]

[1989Gay] [1989Sau] [1991Des]

[1994Hos]

[1995Pav]

[1998Nob]

[2002Zat]

[2003Gro]

[2003Sch]

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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)


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


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



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]

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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

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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]


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


(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

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Al–Li–Zr

140

Al Fig. 3: Al-Li-Zr. Isothermal section at 197°C


(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


570.2 569.8 569.4

c

569.0 818 817

a

816 815 814 813 0

10

20

30

Li, at.%

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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


1728.2 1728.0

c

1727.8 1727.6 401.4 401.2

a 401.0 400.8 0

5

10

15

Li, at.%


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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

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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.


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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)




(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]


MSIT ®

Al–Mg–Mn




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


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®


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


(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


(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


(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


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®


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 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,


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®


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


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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®


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

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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].

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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

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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]


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)

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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]


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)


Al–Mo–Ti [1997Che]

[1997Nak]

[1997Sau] [1997Sin1]

[1997Sin2]

[1997Smi] [1997Zha] [1998Has]

[1998Joh]

[1998Kim]

[1998Tak]

[1998Woo]

[1999Hao]

[1999Flo] [1999Yan]

[2000Alo]

[2000Izu]

[2000Kai]

[2000Oka]


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


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]


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


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


, 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


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)


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]


Al–Mo–Ti Phase/ Temperature Range [°C] , TiAl < 1463

Pearson Symbol/ Space Group/ Prototype tP4 P4/mmm AuCu


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


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]


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


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

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



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


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


40

60

L L+α

α

50

L+β 50

α +β L+β

β


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


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


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


Al–Mo–Ti

313

Ti 40.00 Mo 0.00 Al 60.00



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


10

20

30

40

Ti 40.00 Mo 50.00 Al 10.00

MSIT ®

Al–Mo–Ti

314

Al

Data / Grid: at.%

L


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


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


Al–Mo–Ti

315

Al

Data / Grid: at.%


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


Mo


γ

60

40

α2

80

20

β2 α

Ti


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


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].


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)


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)




(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


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Al–N–Si

320

N Fig. 1: Al-N-Si. Isothermal section at 2400°C


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


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


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

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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.

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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

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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].

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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].


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References [1954Thy]

[1963Jei] [1964Now]

[1976Ivc1]

[1976Ivc2]

[1976Kuz]

[1977Ivc]

[1981Bee]

[1982Wit] [1983Psh]

[1984Bey] [1984Jon] [1984Sch] [1985Sch]

[1986Jeh]

[1986Kau]

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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)


Al–N–Ti [1986Kno] [1986Mue] [1987Hak]

[1987Ina] [1987Kno] [1988Hol] [1988Jeh] [1988Pen]

[1988Ran]

[1989Hol] [1990McI]

[1991Adi]

[1991And]

[1991Col] [1991Hul]

[1991Ike]

[1991Kaw]

[1992Koy]


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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)

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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)


Al–N–Ti [1999Ho]

[1999Mag] [2000Bar1]

[2000Bar2]

[2000Fin]

[2000Gam] [2000Par]

[2000Pro1] [2000Pro2]

[2000Raw]

[2001Hug]

[2001Per]

[2001Spe]

[2002Tan]

[2003Fer]

[2003Sch]


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]



(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]


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331




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


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]

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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


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


(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®


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®


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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´Ó´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

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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].

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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


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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®


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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

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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®


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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

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[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)


Al–Nb–Ti [1978Zak]

[1980Jor]

[1981Bla] [1981Ish] [1981Kin] [1982Mii]

[1982Str]

[1983Tro]

[1984Arg] [1984Zak]

[1986Ibr]

[1986Kon]

[1987Ban]

[1987Mur] [1988Ban] [1988Has]

[1988Str]

[1988Sub] [1989Ben]


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)


Al–Nb–Ti [1990Sau] [1990Sch]

[1990Sho]

[1990Wey]

[1991Akk]

[1991Ban]

[1991Ben]

[1991Cha] [1991Chr] [1991Li] [1991Moh] [1991Pav]

[1991Row1]

[1991Row2]

[1991Row3]

[1991Smi]

[1991Spa]

[1991Zak]


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)


Al–Nb–Ti [1992Qua] [1992Roz] [1992Smi]

[1992Sur] [1992Tre]

[1992Voz]

[1992Yan] [1992Zak]

[1992Zha]

[1993Ebr]

[1993Ers]

[1993Jac] [1993Lon] [1993Mur] [1993Nak1]

[1993Nak2] [1993Oka1] [1993Oka2] [1993Rub]

[1993Sob]


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)


Al–Nb–Ti [1995Zdz]

[1996Che]

[1996Ebr] [1996Jac]

[1996Men]

[1996Sad]

[1996Sik]

[1996Tre]

[1996Vas]

[1996Woo]

[1997Bul]

[1997Jew]

[1997Nak]

[1997Wan]

[1998Che]

[1998Din]

[1998Hel]


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)


Al–Nb–Ti [2000Kai]

[2000Leo1]

[2000Leo2]

[2000Mur] [2000Oka] [2000Sad]

[2000Wel] [2000Yan]

[2001Bra]

[2001Cao]

[2001Col]

[2001Kan]

[2001Li1]

[2001Li2] [2001Mis]

[2001Sad]

[2001Ser]

[2001Sun]


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]


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


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]


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


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



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]


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


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®


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]


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 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®


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]


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)


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


hP5 P3m1 Ni2Al3


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]


Al–Ni–Si Phase/ Temperature Range [°C]

, Ni 2-xAlySi1-y(r)

Ni2Si < 1255


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]


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]


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


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



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


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


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


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


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


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


Al–Ni–Si

419

Fig. 4f: Al-Ni-Si. Vertical section at 45 at.% Ni

L 1500

Temperature, °C

Ni1+xAl1-ySiy


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


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


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

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


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

?


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


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


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]


-10 20

80

-20 40

60

-30

-40

60

40

-50 80

-30 -20 -10

Ni


-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®


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®


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.


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®


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]


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)


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


hP5 P3m1 Ni2Al3


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]


Al–Ni–Si Phase/ Temperature Range [°C]

, Ni 2-xAlySi1-y(r)

Ni2Si < 1255


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]


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]


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


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



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


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


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


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


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


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


Temperature, °C

L+Ni1+xAl1-ySiy 1250


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


Al–Ni–Si

419

Fig. 4f: Al-Ni-Si. Vertical section at 45 at.% Ni

L 1500

Temperature, °C

Ni1+xAl1-ySiy


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


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


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


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

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


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


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

?


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


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


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]


-10 20

80

-20 40

60

-30

-40

60

40

-50 80

-30 -20 -10

Ni


-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®


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


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®


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.


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., Dýakonova, 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)


Al–Ni–Ta [1979Nas] [1984Och]

[1984Wil]

[1985Mis]

[1986Hua] [1986Wil1]

[1986Wil2]

[1987Hil]

[1987Kha] [1988Bre]

[1990Kum]

[1991Bon]

[1991Eno]

[1991Hay]

[1991Kau] [1991Mas]

[1991Mis]

[1991Nas]


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




(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


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




, 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®


Al–Ni–Ta

433




* -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.


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


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®


Al–Ni–Ta

435

830

Fig. 1: Al-Ni-Ta. Lattice constants a (triangle) and c (circle) of the -1 phase

820


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

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