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 C Non-Ferrous Metal Systems Part 3 Selected Soldering and Brazing Systems
Editors G. Effenberg and S. Ilyenko Authors Materials Science International Team, MSIT®
ISSN 1615-2018 (Physical Chemistry) ISBN 978-3-540-25777-6
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/11C3: 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 2007 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: 10916025
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Editors:
Günter Effenberg Svitlana Ilyenko
MSI, Materials Science International Services GmbH Postfach 800749, D-70507, Stuttgart, Germany http://www.matport.com
Authors:
Materials Science International Team, MSIT®
The present series of books results from collaborative evaluation programs performed by MSI and authored by MSIT®. In this program data and knowledge are contributed by many individuals and accumulated over almost twenty years, now. The content of this volume is a subset of the ongoing MSIT® Evaluation Programs. Authors of this volume are: Nataliya Bochvar, Moscow, Russia
Evgenia Lysova, Moscow, Russia
Anatoliy Bondar, Kyiv, Ukraine
Pierre Perrot, Lille, France
Tatyana Dobatkina, Moscow, Russia
Alan Prince †, Harpenden, U.K.
Olga Fabrichnaya, Stuttgart, Germany
Qingsheng Ran, Stuttgart, Germany
Gautam Ghosh, Evanston, USA
Peter Rogl, Wien, Austria
Joachim Gröbner, Clausthal-Zellerfeld, Germany
Lazar Rokhlin, Moscow, Russia
Katharina Haußmann, Stuttgart, Germany
Myriam Sacerdote-Peronnet, Lyon, France
Volodymyr Ivanchenko, Kyiv, Ukraine
Stephan Schittny, Hanau, Germany
Yan Jialin, Guangxi, China
Rainer Schmid-Fetzer, Clausthal-Zellerfeld, Germany
Jozefien De Keyzer, Leuven, Belgium
Elena L. Semenova, Kyiv, Ukraine
Ulrich E. Klotz, Dübendorf, Switzerland
Oleh Shcherban, L'viv, Ukraine
Kostyantyn Korniyenko, Kyiv, Ukraine
Nuri Solak, Stuttgart, Germany
Tatiana Kosorukova, Kyiv, Ukraine
Vasyl Tomashik, Kyiv, Ukraine
Ortrud Kubashewski, Aachen, Germany
Ludmila Tretyachenko, Kyiv, Ukraine
K.C. Hari Kumar, Chennai, India
Mikhail Turchanin, Kramatorsk, Ukraine
Viktor Kuznetsov, Moscow, Russia
Tamara Velikanova, Kyiv, Ukraine
Nathalie Lebrun, Lille, France
Sigrid Wagner, Stuttgart, Germany
Yurii Liberov, Moscow, Russia
Andy Watson, Leeds, U.K.
Chunlei Liu, Dübendorf, Switzerland
Viktor Witusiewicz, Aachen, Germany
Jörg F. Löffler, Dübendorf, Switzerland
Matvei Zinkevich, Stuttgart, Germany
Hans Leo Lukas, Stuttgart, Germany
Institutions The content of this volume is produced by Materials Science International Services GmbH and the international team of materials scientists, MSIT®. Contributions to this volume have been made from the following institutions: Access, Aachen, Germany
Laboratoire des Multimatériaux et Interfaces, Université Claude Bernard Lyon I
The Baikov Institute of Metallurgy, Academy of Sciences, Moscow, Russia
Max-Planck Institut für Metallforschung, Institut für Werkstoffwissenschaft, Pulvermetallurgisches Laboratorium, Stuttgart, Germany
Donbass State Mechanical Engineering Academy, Kramatorsk, Ukraine
Moscow State University, Department of General Chemistry, Moscow, Russia
EMPA, Materials Science and Technology Laboratory of Joining and Interface Technology, Dübendorf, Switzerland
National University of L'viv, Cathedra of Inorganic Chemistry, L'viv, Ukraine
I.M. Frantsevich Institute for Problems of Materials Science, National Academy of Sciences, Kyiv, Ukraine
Northwestern University, Department of Materials Science and Engineering, Evanston, USA
Heraeus GmbH, Hanau, Germany
RWTH Aachen, Germany
Indian Institute of Technology Madras, Department of Metallurgical Engineering, Chennai, India
Technische Universität Clausthal, Metallurgisches Zentrum, Clausthal-Zellerfeld, Germany
Institute of Materials Science, Guangxi University, Nanning, China
Université de Lille I, Laboratoire de Métallurgie Physique, Villeneuve d’ASCQ, Cedex, France
Institute for Semiconductor Physics, National Academy of Sciences, Kyiv, Ukraine
Universität Wien, Institut für Physikalische Chemie, Wien, Austria
Katholieke Universiteit Leuven, Department Metaalkunde en Toegepaste Materiaalkunde, Heverlee, Belgium
University of Leeds, Department of Materials, School of Process, Environmental and Materials Engineering, Leeds, UK
G.V. Kurdyumov Institute for Metal Physics, National Academy of Sciences, Kyiv, Ukraine
Preface The growing need to develop Pb free solder alloys less damaging for the environment is causing increased research into candidate materials systems, such as those evaluated in this volume. However the lead-tin solder alloys are still commonly used and dominate the industrial market by far. A candidate lead free solder alloy must fulfill many requirements: wettability of the substrate, ability to form a strong chemical bond with the substrate, suitable melting and solidification behavior, good mechanical properties, corrosion resistance, electrical conductivity and it must not be harmful to health and environment. We included in this volume ternary systems from which one may arrive at alloys that satisfy the above requirements and are therefore promising candidates for the development of such soldering and brazing materials. But also other important solder or braze systems which are presently in use are included in this volume. The sub-series Ternary Alloy Systems of the Landolt-Börnstein New Series provides reliable and comprehensive descriptions of the materials constitution, based on critical intellectual evaluations of all data available at the time, and it critically weights the different findings, also with respect to their compatibility with today’s edge binary phase diagrams. Selected are ternary systems of importance to industrial alloy development and systems which gained scientific interest in the recent years otherwise. In a ternary materials system, however, one may find alloys for various applications, depending on the chosen composition. Reliable phase diagrams provide scientists and engineers with basic information of eminent importance for fundamental research and for the development and optimization of materials. So collections of such diagrams are extremely useful, if the data on which they are based have been subjected to critical evaluation, like in these volumes. Critical evaluation means: where contradictory information is published data and conclusions are being analyzed, broken down to the firm facts and reinterpreted 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 has been working together for 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, X-ray and microprobe analyses. In each published case the applicability of the chosen method has to be validated, the way of actually performing the experiment or computer modeling has to be validated 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 A-
B-X systems. Therefore combining systematically binary and ternary evaluations can lead to a level of increased 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 LandoltBörnstein 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 evaluation works 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 provides the reader with the data for each ternary system in a concise and consistent manner, as applied in the MSIT® Workplace: Phase Diagrams Online. The standard format and special features of the Landolt-Börnstein compendium are explained in the Introduction to the volume. In spite of the skill and labor that have been put into this volume, it will not be faultless. All criticisms and suggestions that can help us to improve our work are very welcome. Please contact us via
[email protected]. We hope that this volume will prove to be 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 we want to thank all those who contributed their comments and insight during the evaluation process. In particular we thank the reviewers – Andy Watson, Pierre Perrot, Rainer Schmid-Fetzer, Olga Fabrichnaya, Hari Kumar, Gabriele Cacciamani, Matvei Zinkevich, Artem Kozlov, Ludmila Tretyachenko, Joachim Gröbner, Hans Leo Lukas, Yong Du, Marina Bulanova. Special thanks go to Oleksandr Dovbenko who very competently helped the editors to supervise the review and helped the authors in discussing questionable matters. We all gratefully acknowledge the dedicated desk editing by Oleksandra Berezhnytska, Oleksandr Rogovtsov, and Vyacheslav Saltykov.
Günter Effenberg and Svitlana Ilyenko
Stuttgart, March 2006
Contents IV/11 Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data Subvolume C: Non-Ferrous Metal Systems Part 3: Selected Soldering and Brazing Systems
Introduction Data Covered ..................................................................................................................................XI General............................................................................................................................................XI Structure of a System Report ..........................................................................................................XI Introduction..........................................................................................................................XI Binary Systems ....................................................................................................................XI Solid Phases ....................................................................................................................... XII Quasibinary Systems......................................................................................................... XIII Invariant Equilibria ........................................................................................................... XIII Liquidus, Solidus, Solvus Surfaces................................................................................... XIII Isothermal Sections........................................................................................................... XIII Temperature – Composition Sections ............................................................................... XIII Thermodynamics .............................................................................................................. XIII Notes on Materials Properties and Applications............................................................... XIII Miscellaneous ................................................................................................................... XIII References.........................................................................................................................XVI General References .................................................................................................................... XVII
Ternary Systems Ag – Bi – Cu (Silver – Bismuth – Copper).......................................................................................1 Ag – Bi – Sn (Silver – Bismuth – Tin)..............................................................................................4 Ag – Cu – In (Silver – Copper – Indium) .......................................................................................18 Ag – Cu – Mn (Silver – Copper – Manganese) ..............................................................................26 Ag – Cu – Ni (Silver – Copper – Nickel) .......................................................................................30 Ag – Cu – P (Silver – Copper – Phosphorus) .................................................................................38 Ag – Cu – Sn (Silver – Copper – Tin) ............................................................................................47 Ag – Cu – Ti (Silver – Copper – Titanium) ....................................................................................63 Ag – Cu – Zn (Silver – Copper – Zinc) ..........................................................................................75 Ag – In – Sb (Silver – Indium – Antimony) ...................................................................................86 Ag – In – Sn (Silver – Indium – Tin)..............................................................................................96 Ag – Pb – Sn (Silver – Lead – Tin)...............................................................................................113 Ag – Sn – Zn (Silver – Tin – Zinc) ...............................................................................................121 Au – Bi – Sn (Gold – Bismuth – Tin) ...........................................................................................131 Au – Cu – Sn (Gold – Copper – Tin)............................................................................................138 Au – In – Sn (Gold – Indium – Tin) .............................................................................................149 B – Cr – Ni (Boron – Chromium – Nickel) ..................................................................................153 Bi – In – Sb (Bismuth – Indium – Antimony)...............................................................................168
Bi – In – Sn (Bismuth – Indium – Tin) .........................................................................................191 Bi – Sn – Zn (Bismuth – Tin – Zinc) ............................................................................................202 Co – Cu – Mn (Cobalt – Copper – Manganese)............................................................................212 Cr – Ni – P (Chromium – Nickel – Phosphorus) ..........................................................................222 Cr – Ni – Si (Chromium – Nickel – Silicon).................................................................................229 Cu – In – Sn (Copper – Indium – Tin)..........................................................................................249 Cu – Mn – Ni (Copper – Manganese – Nickel) ............................................................................274 Cu – Mn – Sn (Copper – Manganese – Tin) .................................................................................286 Cu – Ni – Sn (Copper – Nickel– Tin) ...........................................................................................303 Cu – Ni – Zn (Copper – Nickel – Zinc) ........................................................................................338 Cu – P – Sn (Copper – Phosphorus – Tin)....................................................................................355 Cu – Pb – Sn (Copper – Lead – Tin).............................................................................................368 Cu – Pd – Sn (Copper – Palladium – Tin) ....................................................................................390 Cu – Si – Zn (Copper – Silicon – Zinc) ........................................................................................393 Cu – Sn – Ti (Copper – Tin – Titanium).......................................................................................409 Cu – Sn – Zn (Copper – Tin – Zinc) .............................................................................................422 Cu – Ti – Zr (Copper – Titanium – Zirconium)............................................................................436 In – Ti – Zn (Indium – Tin – Zinc) ...............................................................................................465 Mn– Ti – Zr (Manganese – Titanium – Zirconium)......................................................................475 Ni – Pd – Si (Nickel – Palladium – Silicon) .................................................................................486
Free WEB Access to update information and more. Content updates of the Landolt-Börnstein sub-series IV/11 plus supplementary information are available from MSI, including: • • • •
Links to Literature (up-to-date bibliographic data base) Diagrams as Published (not MSIT®-evaluated diagrams) Research Results (published and proprietary data) Ternary Evaluations: These are LB IV/11 contents and their updates (if any) as interactive live diagrams & documents.
This service is free of charge for Landolt-Börnstein subscribers and applies for material systems included in the sub-series IV/11. As eligible Springer customer, please contact MSI for access at
[email protected] . Contents and supplementary information to the Landolt-Boernstein sub-series IV/11 are made by MSI, Materials Science International Services, GmbH, Stuttgart and its global team MSIT®, as part of their ongoing Phase Diagram Evaluation Programs. For details on “MSIT® Workplace, Phase Diagrams Online” see: http://www.matport.com .
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: Introduction The opening text reviews briefly the status of knowledge published on the system and outlines the experimental methods that have been applied. Furthermore, attention may be drawn to questions which are still open or to cases where conclusions from the evaluation work modified the published phase diagram. Binary Systems Where binary systems are accepted from standard compilations reference is made to these compilations. In other cases the accepted binary phase diagrams are reproduced for the convenience of the reader. The selection of the binary systems used as a basis for the evaluation of the ternary system was at the discretion of the assessor.
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Introduction
Heading Introduction Binary Systems Solid Phases Quasibinary 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 Quasibinary 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.
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Quasibinary Systems Quasibinary (pseudobinary) sections describe equilibria and can be read in the same way as binary diagrams. The notation used in quasibinary systems is the same as that of vertical sections, which are reported under “Temperature – Composition Sections”. Invariant Equilibria The invariant equilibria of a system are listed in the table “Invariant Equilibria” and, where possible, are described by a constitutional “Reaction Scheme” (Fig. 2). The sequential numbering of invariant equilibria increases with decreasing temperature, one numbering for all binaries together and one for the ternary system. Equilibria notations are used to indicate the reactions by which phases will be – decomposed (e- and E-type reactions) – formed (p- and P-type reactions) – transformed (U-type reactions) 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-quasibinary T-x sections (or vertical sections, isopleths, polythermal sections) show the phase fields where generally the tie lines are not in the same plane as the section. The notation employed for the latter (see Fig. 5) is the same as that used for binary and pseudobinary phase diagrams. 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.
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MSIT®
Ag-Tl
Tl-Bi
144 e9 (Tl)(h) Tl3Bi+(Tl)(r)
192 e8 l Tl3Bi+Tl2Bi3
202 e7 l (Bi)+Tl2Bi3
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)
303 e1 l (Tl)(h)+Tl3Bi
Fig 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
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XV
C
Data / Grid: at.% Axes: at.%
δ
p1
700
20
80
500°C isotherm, temperature is usually 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/11C3
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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.
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General References [C.A.] [Curr.Cont.] [E] [G] [H] [L-B]
[Mas] [Mas2] [P] [S] [V-C] [V-C2]
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Chemical Abstracts - pathways to published research in the world's journal and patent literature - http://www.cas.org/ Current Contents - bibliographic multidisciplinary current awareness Web resource http://www.isinet.com/products/cap/ccc/ Elliott, R.P., Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York (1965) Gmelin Handbook of Inorganic Chemistry, 8th ed., Springer-Verlag, Berlin Hansen, M. and Anderko, K., Constitution of Binary Alloys, McGraw-Hill, New York (1958) Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology (New Series). Group 3 (Crystal and Solid State Physics), Vol. 6, Eckerlin, P., Kandler, H. and Stegherr, A., Structure Data of Elements and Intermetallic Phases (1971); Vol. 7, Pies, W. and Weiss, A., Crystal Structure of Inorganic Compounds, Part c, Key Elements: N, P, As, Sb, Bi, C (1979); Group 4: Macroscopic and Technical Properties of Matter, Vol. 5, Predel, B., Phase Equilibria, Crystallographic and Thermodynamic Data of Binary Alloys, Subvol. a: Ac-Au ... Au-Zr (1991); Springer-Verlag, Berlin. Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, ASM, Metals Park, Ohio (1986) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Pearson, W.B., A Handbook of Lattice Spacings and Structures of Metals and Alloys, Pergamon Press, New York, Vol. 1 (1958), Vol. 2 (1967) Shunk, F.A., Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York (1969) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, ASM, Metals Park, Ohio (1985) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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IV/11 Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data Subvolume C: Non-Ferrous Metal Systems Part 3: Selected Soldering and Brazing Systems Ag - …
Au - …
B-…
Bi - …
Co - …
Ag - Bi - Cu Ag - Bi - Sn Ag - Cu - In Ag - Cu - Mn Ag - Cu - Ni Ag - Cu - P Ag - Cu - Sn Ag - Cu - Ti Ag - Cu - Zn Ag - In - Sb Ag - In - Sn Ag - Pb - Sn Ag - Sn - Zn
Au - Bi - Sn Au - Cu - Sn Au - In - Sn
B - Cr - Ni
Bi - In - Sb Bi - In - Sn Bi - Sn - Zn
Co - Cu - Mn
Cr - …
Cu - …
In - …
Mn - …
Ni - …
Cr - Ni - P Cr - Ni - Si
Cu - In - Sn Cu - Mn - Ni Cu - Mn - Sn Cu - Ni - Sn Cu - Ni - Zn Cu - P - Sn Cu - Pb - Sn Cu - Pd - Sn Cu - Si - Zn Cu - Sn - Ti Cu - Sn - Zn Cu - Ti - Zr
In - Ti - Zn
Mn - Ti - Zr
Ni - Pd - Si
Ag–Bi–Cu
1
Silver – Bismuth – Copper Joachim Gröbner Introduction The first data on the Ag-Bi-Cu system were reported by [1951Ava], who considered microstructure of alloys near the binary Ag-Cu eutectic depending on the cooling rate. [1976Zan] reported results of rapid quench-cooling experiments in the Ag rich part of the system. A eutectic reaction was found to exist at 240 5°C, but the phases in equilibrium were not specified. This work was reviewed by [1980Hen] and [1988Hen]. [1988Liu, 1989Liu] investigated the entire ternary system in 12 vertical sections by thermoanalysis (DTA) in an atmosphere of dry nitrogen. They found that the ternary Ag-Bi-Cu system is a simple eutectic one with the invariant reaction at 258°C. [1998Shu] has calculated the liquidus surface using a model of ideal solution of self-associates and compared the result with that measured by [1989Liu]. [2004Doi] has presented the primary crystallization surfaces and liquidus isotherms calculated using Calphad method and the thermodynamic parameters for the three binary systems. A good agreement was found between the results of calculation and the experimental data by [1989Liu]. Binary Systems The binary systems Ag-Cu and Bi-Cu are accepted from the MSIT Evaluation Program [2002Rom] and [2002Cac], respectively. The Ag-Bi phase diagram is accepted from [Mas2]. Solid Phases No ternary phases were found in this system. The crystallographic data of the phases in the Ag-Cu-Bi system and their ranges of stability are listed in Table 1. Invariant Equilibria One invariant eutectic reaction: L (Ag) + (Cu) + (Bi) at 258°C exists in the ternary system [1989Liu]. The composition of the liquid phase is given in Table 2. Liquidus Surface The liquidus surface after [1989Liu] is given in Fig. 1. Notes on Materials Properties and Applications The Ag-Bi-Cu system is a part of the Ag-Bi-Cu-Sn system which has been enumerated as a strong potential candidate alloy for a lead-free solder [2004Doi]. References [1951Ava] [1976Zan]
[1980Hen] [1988Hen]
Landolt-Börnstein New Series IV/11C3
Avakyan, S.V., Lashko, N.F., “Crystallization of Binary Eutectics in Ternary Systems” (in Russian), Zh. Fiz. Khim., 25, 1085-1091 (1951) (Experimental, 2) Zanicchi, G., Ferro, R., Marazza, R., Contardi, V., “Some Metastable Ag-Rich Alloys in the Silver-Bismuth-Copper System”, J. Less-Common Met., 50, 151-154 (1976) (Experimental, 6) Henig, E.T., “The Ag-Bi-Cu (Silver-Bismuth-Copper) System”, Bull. Alloy Phase Diagrams, 1(2), 62 (1980) (Abstract, 1) Henig, E.T., “Silver - Bismuth - Copper”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg G. (Ed.), MSI, Materials International Services GmbH, Stuttgart, Document ID: 10.12132.1.20, (1988) (Crys. Structure, Phase Diagram, Assessment, 1)
MSIT®
Ag–Bi–Cu
2 [1988Liu]
[1989Liu] [1994Sub] [1998Shu]
[2002Cac]
[2002Rom]
[2004Doi]
Liu, S., Sun, W., Wenqing, “Liquidus of the Silver-Copper-Bismuth System” (in Chinese), Jinshu Xuebao, 24(5), B376-B378 (1988) (Experimental, Phase Diagram, Phase Relations, #, 1) Liu, S., Sun, W., “Liquidus of Ag-Cu-Bi System”, Acta Metall. Sinica, 2B(2), 151-152 (1989) (Experimental, Phase Diagram, #, 1) Subramanian, P.R., Perepezko, J.H., “Ag-Cu (Silver - Copper)”, J. Phase Equilib., 14, 62-75 (1994) (Crys. Structure, Phase Diagram, Review, Thermodyn., 81) Shunyaev, K.Y., Tkachev, N.C., Vatolin, N.A., “Liquidus Surface and Association in Eutectic Ternary Alloys”, Thermochim. Acta, 314, 299-306 (1998) (Calculation, Phase Relations, Thermodyn., #, 21) Cacciamini, G., Cornish, L., Saltykov, P., “Bi-Cu (Bismuth - Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg G. (Ed.), MSI, Materials International Services GmbH, Stuttgart, Document ID: 20.19587.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, 13) Rompaey, T. van, Rogl, P., “Ag-Cu (Silver - Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg G. (Ed.), MSI, Materials International Services GmbH, Stuttgart, Document ID: 20.14511.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, 29) Doi, K., Othani, H., Hasebe, M., “Thermodynamic Study of the Phase Equilibria in the Sn-Ag-Bi-Cu Quaternary System”, Mater. Trans., 45(2), 380-383 (2004) (Calculation, Phase Diagram, Thermodyn., 9)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
(Cu) < 1084.62
cF4 Fm3m Cu
(Bi) < 271.442
pure Ag at 25°C [Mas2] dissolves 14 at.% Cu [Mas2, 2002Rom] dissolves 5 at.% Ag [Mas2, 2002Rom], 0.0193 at.% Bi [Mas2, 2002Cac] at 25°C [Mas2] melting point [1994Sub, 2002Rom]
a = 361.46
hR6 R3m As
a = 363.7
pure Cu at 19°C [Mas2, 2002Rom]
a = 454.613 c = 1186.152
at 31°C [V-C2, 2002Cac]
Table 2: Invariant Equilibria Reaction L (Ag) + (Cu) + (Bi)
MSIT®
T [°C]
258
Type
E
Phase
L
Composition (at.%) Ag
Cu
Bi
5.0
0.5
94.5
Landolt-Börnstein New Series IV/11C3
Ag–Bi–Cu
3
Bi Fig. 1: Ag-Bi-Cu. Liquidus surface projection
Data / Grid: at.% Axes: at.%
e3 400°C
(Bi) e2 E 20
600 80
650 700 40
60
500 750 60
40
800 850
(Cu)
80
20
900 950 (Ag)
Ag
Landolt-Börnstein New Series IV/11C3
1000 20
40 e1
60
80
Cu
MSIT®
4
Ag–Bi–Sn
Silver – Bismuth – Tin Andy Watson Introduction The Ag-Bi-Sn system has been the subject of intense scrutiny in recent years owing to its possible application as a lead free solder material. However, the vast majority of the interest has been with respect to physical and mechanical properties rather than phase equilibria. Numerous reviews have appeared in the literature [1997Hua, 1997Lee, 1999Ohn, 2001Kat, 2001Sug, 2003Ohn] covering mechanical and physical property studies but all refer ultimately to the same small body of work conducted on the phase equilibria. Studies of the phase equilibria in the ternary system began with a thermodynamic calculation by [1994Kat] that used thermodynamic descriptions of the associated binary systems to extrapolate into the three component system. Predicted liquidus temperatures agreed well with experimental observations using DTA studies of two alloys (Ag3.32-Bi1.14-Sn94.55 and Ag2.91-Bi42.84-Sn54.25). However, the first extensive study of the chemistry of the system was conducted by [1994Has]. Enthalpies of formation of the ternary liquid was measured by high temperature calorimetry at 783 and 878 K along sections Ag:Bi 1:3 and 1:1, Ag:Sn 1:4 and 1:1 and Bi:Sn 1:3, 1:1 and 3:1. High purity materials were used (Ag 99.998, Bi 99.999, Sn 99.999). The results were interpolated in order to construct isoenthalpic curves using models given in [1965Too, 1984Hoc]. By noting deviations in the enthalpy curves, it was possible to present an approximation of the liquidus surface. To complement the work on thermodynamic properties, [1998Has] made an extensive study of the phase equilibria in the ternary system using X-ray powder diffraction, DSC and DTA. Samples of Ag(4N), Bi(5N) and Sn(5N) were sealed in evacuated silica tubes before melting and cooling to an annealing temperature between 120 and 250°C, where they were held for 15 d. A variety of heating and cooling rates were used in the thermal studies (0.1-2°C#min–1), the lowest of which was used for the determination of the invariant temperatures. Three isoplethal sections were characterized (10 at.% Bi, 40 at.% Ag and 70 at.% Bi) and three ternary invariant reactions were found. The locations of two of these were within approximately 2°C the extrapolations of [1994Kat], the third however, was around 23°C higher. More recently, a complete thermodynamic assessment of the ternary system has been presented by [2001Oht] (which was used subsequently in a calculation of the quaternary Ag-Bi-Sn-Zn by [2004Oht]). Firstly, they performed DSC studies along isoplethal sections with 10, 20, 30, 40, 50, 60, 70 and 80 at.% Bi. Sn(99.999%), Ag(99.99%) and Bi(99.99%) were homogenized at 1100°C in sealed quartz tubes and quenched into iced brine before the DSC studies. Peaks were recorded on heating and cooling at 3°C#min–1. Isothermal sections at 200 and 400°C were produced from EDX studies of alloys which had been sealed in argon filled silica tubes and equilibrated for 14 and 7 d, respectively. Using the experimental data of [1994Has, 1998Has] and the results of their own DSC and EDX studies, they produced a thermodynamic description of the system that is in good agreement with all of the available experimental data. Interaction parameters were introduced for the liquid and the phases, and Bi was allowed to substitute for Sn in the model for the Ag3Sn compound. This latter aspect of the modeling was necessary to make the phase boundaries involving this phase in the ternary system fit the experimental data, despite the fact that little solubility of Bi in this phase has been found experimentally. A reaction scheme is also presented showing three invariant equilibria: 2 transition reactions and a ternary eutectic. However, the reaction given at 263.3°C is higher than the binary eutectic in the Ag-Bi system. In order for the ternary reaction to be a transition then there must be a maximum in the monovariant line from the binary Ag-Bi eutectic. This is suggested in [1998Has]. The assessed temperature of this reaction given by [2001Oht] is 263.6°C. [2001Lee] used thermodynamic data to predict viscosity and surface tension of the Ag-Bi-Sn liquid. However, they did not use the data of [2001Oht] but just the descriptions for the binary systems. Nevertheless, they found reasonable agreement with their experimental data.
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Bi–Sn
5
Binary Systems The Ag-Sn binary system is accepted from the thermodynamic assessment of [1999Oht] and is presented in Fig. 1 with an amendment to the Ag3Sn phase showing some homogeneity which was ignored in the modeling. The Ag-Bi and Bi-Sn systems are taken from the thermodynamic assessments of [1993Kar] and [1994Oht], respectively. Solid Phases Solid phases in the system are listed in Table 1. There are no ternary compounds and little experimental evidence of extension of binary phases into the ternary. However, [2001Oht] found that allowing substitution of Bi for Sn in Ag3Sn in the modeling of the system resulted in a better fit of the calculated phase boundaries to the experimental data. Invariant Equilibria Three invariant equilibria exist in this system. Details are given in Table 2 and the reaction scheme is given in Fig. 2. The reaction scheme is based on [2001Oht]. Liquidus, Solidus and Solvus Surfaces The liquidus surface is shown in Fig. 3. The primary crystallization fields are given by [2001Oht] whereas the liquidus isothermal lines are experimentally determined from calorimetric (605 and 510°C) and differential thermal analysis (477, 427 ant 377°C) measurements of [1994Has]. The liquidus isothermal lines in the Sn rich corner of the system (< 60 at.% Bi, < 7 at.% Ag), taken from [2001Kat] are given in Fig. 4. The enthalpies of mixing of the liquid phase as measured by [1994Has] are exothermic in the Ag rich corner, and endothermic in the main part of the diagram suggesting, along with the shape of the liquidus curves in the Ag-Bi and Bi-Sn systems, the presence of a metastable miscibility gap. This is shown in Fig. 5, taken from the assessment of [2001Oht]. Isothermal Sections Isothermal sections for temperatures of 400 and 200°C are given in Figs. 6 and 7, taken from [2001Oht]. Slight alterations have been made to the Ag3Sn phase to show a degree of homogeneity in agreement with Fig. 1. The extension of the Ag3Sn phase into the ternary system is a product of the optimization of the system without any supporting experimental evidence, but it was necessary for it to be included as discussed above. Temperature – Composition Sections Calculated temperature-composition sections taken from [2001Oht] are given in Figs. 8-13, and are based on experimental observations. Most of the phase boundaries in the calculations agree well with the experimental data which were used in the optimization. However, no experimental data were available for the Ag rich regions at low temperatures, and hence this part of the diagrams should be treated with a degree of caution. Thermodynamics [1994Has] measured by direct reaction calorimetry the enthalpies of formation of liquid (Ag,Bi) and (Ag,Bi,Sn) alloys at 510 and 605°C along different sections characterized by Ag/Bi = 1 (40 alloys) and 1/3 (38 alloys), Ag/Sn = 1 (35 alloys) and 1/4 (30 alloys), Bi/Sn = 1/3 (30 alloys), 1 (40 alloys) and 3 (29 alloys). Figure 14 shows the experimental isoenthalpic lines drawn at 605°C from [1994Has]. A general expression of the excess Gibbs energy of mixture in the liquid state has been derived by [2001Oht, 2004Oht] as a function of the temperature as part of the assessment of the ternary system. The calculated curves are in excellent agreement with the experimental data of [1994Has]: Landolt-Börnstein New Series IV/11C3
MSIT®
6
Ag–Bi–Sn
mixGxs = xAgxSnLAg,Sn + xBixSnLBi,Sn + xAgxBiLAg,Bi + xAgxBixSnLAg,Bi,Sn LAg,Sn = – 4902.5 + 4.30532T + (– 16 474 + 3.12507T)(xAg – xSn) – 7298.6 (xAg – xSn)2 LBi,Sn = 487 + 0.966T – (32 + 0.235T)(xBi – xSn) LAg,Bi = 4589.8 + 23.73047T – 3.93814TlnT – (5716.6 + 0.91452T)(xAg – xBi) + (– 2630.2 + 0.88522T)(xAg – xBi)2 LAg,Bi,Sn = (1700 + 76.2T)xAg + (11 000 + 4T)xBi + (20 000 – 38.95T)xSn Notes on Materials Properties and Applications Owing to increasing concerns over the use of lead, alloys of the Ag-Bi-Sn system have become of interest as possible replacements for the traditional Pb-Sn solder materials [1994Kat, 1997Hua, 1997Lee]. Consequently, a good deal of research has been conducted on the properties of alloys in this system focusing on Sn rich alloys and of Bi-Sn eutectic alloys with small additions of Ag [2002Kat]. With increasing Bi content, the melting temperature of Ag-Bi-Sn alloys decreases promoting wetting behavior [2002Cho]. When used in conjunction with a copper substrate, this can also be attributed to a reduction in the formation of Cu-Sn compounds owing to the presence of Ag. A small addition of Ag was found to improve significantly the ductility of Bi-Sn eutectic solder [1997McC]. The strength of a solder joint is an important issue; and this increases with increasing strain rate [2004Sho] and increasing Bi content owing to solid solution strengthening [2002Shi], with a corresponding loss in ductility. The bonding strength of Bi-Sn joints near room temperature is higher than that of the Sn-3.5Ag solder, but the ductility is lower [2001Kik]. Creep rate of the materials also increases with increasing Bi content [2002Yu], whereas [1998Kar1, 1998Kar2, 1999Kar, 2001Kar] pointed out than addition of Bi drastically degrades the thermal fatigue resistance of Sn-3.5 at.% Ag solder. The choice of substrate material also affects the strength of the solder joint. Cu increases the strength of the joint whereas Ni/Cu finishes are found to decrease it [2004Hwa]. Rare earth additions can improve the wettability of the liquid solder and increase the strength of the joint by refinement of the (Sn) phase [2002Xia1, 2004Wu]. Although Ag-Bi-Sn solders are superior to other candidates with respect to the melting properties and wettability [2001Bra], they are subject to the fillet-lifting (or lift-off) phenomenon [2001Tak]. Fillet-lifting was shown to occur in alloys (3 to 20 at.% Bi) with insufficient latent heat release near the final solidification temperature for the alleviation of the temperature gradient in the solder joint. Miscellaneous Solder powders can be prepared by mechanical alloying [2003Lai]. Grinding media is very important, ceramic balls providing a stronger impact force. As Ag-Bi-Sn solders have higher melting points (202-225°C) than the most popular eutectic Pb-Sn solder (183°C), [2003Shi] uses indium additions to lower the melting point of the solder alloy and to keep the good wetting performances and good mechanical properties which characterizes Ag-Bi-Sn solders [2002Xia2]. Applying ultrasonic waves to the soldering influences the microstructure [2001Kik] and increases the strength of the bonding interface in the solder joints. The surface tension together with the density was measured for the Ag-Bi-Sn liquid alloys (0 to 80 at.% Bi and 0 to 12 at.% Sn) in the temperature range 225 - 900°C [2001Mos] and for the Ag-Sn eutectic between 230 and 950°C [2004Gas]. A linear dependency with temperature was observed. The addition of Bi to liquid Sn and to Ag-Sn eutectic alloy (3.8 at.% Ag) reduces markedly the surface tension. The wettability of Ag-Bi-Sn solders on Cu was measured by Energy Dispersive X-ray Spectrometry [2002Tan]. The wettability and mechanical properties of solders based on Bi-Sn systems are lower than that of solders based on conventional Pb-Sn systems [1998Her, 1999Via2]. However, it may be enhanced by addition of 1 to
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Bi–Sn
7
4 mass% Ag to the solder. An investigation of the melting properties of the 91.84Sn-3.33Ag-4.83Bi solder [1999Via1] suggested the appearance of a metastable, short-range order as a result of thermal aging at low temperature. The kinetics growth of Cu-Sn layers at the interface solder/Cu was also investigated. References [1965Too] [1984Hoc]
[1993Kar]
[1994Has]
[1994Kat]
[1994Oht]
[1997Hua] [1997Lee] [1997McC]
[1998Has]
[1998Her]
[1998Kar1]
[1998Kar2]
[1999Kar]
[1999Ohn]
[1999Oht]
Landolt-Börnstein New Series IV/11C3
Toop, G.W., “Predicting Ternary Activities Using Binary Data”, Trans. Metall. Soc. AIME, 233(5), 850 (1965) (Calculation, 13) Hoch, M., Arpshofen, I., “Aggregation Model for the Calculation of Thermodynamic State Functions of Liquid Alloys” (in German), Z. Metallkd., 75(1), 23-29 (1984) (Calculation, 13) Karayaka, I., Thompson, W.T., “The Ag-Bi (Silver-Bismuth) System”, J. Phase Equlib., 14(4), 525-530 (1993) (Phase Diagram, Phase Relations, Thermodyn., Crys. Structure, Assessment, Review, 24) Hassam, S., Gambino, M., Bros, J.P., “Enthalpy of Formation of Liquid Ag-Bi and Ag-Bi-Sn Alloys”, Z. Metallkd., 85, 460-471 (1994) (Experimental, Phase Relations, Thermodyn., *, 27) Kattner, U.R., Boettinger, W.J., “On the Sn-Bi-Ag Ternary Phase Diagram”, J. Electron. Mater., 23(7), 603-610 (1994) (Phase Diagram, Phase Relations,Thermodyn., Calculation, 53) Ohtani, H., Ishida, K., “A Thermodynamic Study of the Phase Equilibria in the Bi-Sn-Sb System”, J. Electron. Mater., 23(8), 747-755 (1994) (Phase Relations, Thermodyn., Experimental, Assessment, 36) Hua, F., Glazer, J., “Lead-Free Solders For Electronic Assembly”, Des. Reliab. Solders Solder Interconnect., Proc. Symp., 65-73 (1997) (Experimental, 63) Lee, N.C., “Getting Ready for Lead-Free Solders”, Soldering Surf. Mount Technol., 9(2), 65-69 (1997) (Mechan. Prop., Phase Relations, Review, 0) McCormack, M., Chen, H.S., Kammlott, G.W., Jim, S., “Significantly Improved Mechanical Properties of Bi-Sn Solder Alloys by Ag-Doping”, J. Electron. Mater., 26(8), 954-958 (1997) (Experimental, Mechan. Prop., 12) Hassan, S., Dichi, E., Legendre, B., “Experimental Equilibrium Phase Diagram of the Ag-Bi-Sn System”, J. Alloys Compd., 268, 199-206 (1998) (Experimental, Phase Diagram, Phase Relations, *, 13) Hermandez, C.L., Vianco, P.T., Rejent, J.A., “Effect of Interface Microstructure on the Mechanical Properties of Pb-Free Hydrid Microcircuit Solder Joints”, IPC/SMTA Electr. Assem. Expo, S19(2), 1-8 (1998) (Experimental, Interface Phenomena, Morphology, Mechan. Prop., 9) Kariya, Y., Otsuka, M., “Effect of Bismuth on the Isothermal Fatigue Properties of Sn - 3.5 mass% Ag Solder Alloy”, J. Electron. Mater., 27(7), 866-870 (1998) (Experimental, Mechan. Prop., 16) Kariya, Y., Otsuka, M., “Mechanical Fatigue Characteristics of Sn-3.5Ag-X (X = Bi, Cu, Zn and In) Solder Alloys”, J. Electron. Mater., 27(11), 1229-1235 (1998) (Experimental, Mechan. Prop., 14) Kariya, Y., Hirata, Y., Otsuka, M., “Effect of Thermal Cycles on the Mechanical Strength of Quad Flat Pack Leads/Sn3.5Ag-X (X = Bi and Cu) Solder Joints”, J. Electron. Mater., 28(11), 1263-1269 (1999) (Experimental, Mechan. Prop., 15) Ohnuma, I., Liu, X.J., Ohtani, H., Ishida, K., “Thermodynamic Database for Phase Diagrams in Micro-Soldering Alloys”, J. Electron. Mater., 28(11), 1164-1171 (1999) (Phase Diagram, Phase Relations, Thermodyn., Calculation, 16) Ohtani, H., Miyashita, M., Ishida, K., “Thermodynamic Study of the Sn-Ag-Zn System”, J. Jpn. Inst. Met., 63 685-694 (1999) (Phase Relations, Phase Diagram, Thermodyn., Assessment, 68)
MSIT®
8 [1999Via1]
[1999Via2]
[2001Bra]
[2001Kar]
[2001Kat]
[2001Kik]
[2001Lee]
[2001Mos]
[2001Oht]
[2001Sug] [2001Tak]
[2002Cho]
[2002Kat] [2002Shi]
[2002Tan]
[2002Xia1]
[2002Xia2]
MSIT®
Ag–Bi–Sn Vianco, P.T., Rejent, J.A., “Properties of Ternary Sn-Ag-Bi Solder Alloys: Part I - Thermal Properties and Microstructural Analysis”, J. Electron. Mater., 28(10), 1127-1137 (1999) (Experimental, Thermodyn., 17) Vianco, P.T., Rejent, J.A., “Properties of Ternary Sn-Ag-Bi Solder Alloys: Part: II -Wettability and Mechanical Properties Analyses”, J. Electron. Mater., 28(10), 1138-1143 (1999) (Experimental, Mechan. Prop., 6) Bradley, E., III, Hranisavljevic, J., “Characterization of the Melting and Wetting of Sn-Ag-X Solders”, Trans. Electron. Packag. Manuf., 24(4), 255-260 (2001) (Experimental, Phase Relations, 16) Kariya, Y., Morihata, T., Hazawa, E., Otsuka, M., “Assessment of Low-Cycle Fatigue Life of Sn-3.5 mass% Ag-X (X = Bi or Cu) Alloy by Strain Range Partitioning Approach”, J. Electron. Mater., 30(9), 1184-1189 (2001) (Assessment, Mechan. Prop., 8) Kattner, U., Handwerker, C.A., “Calculation of Phase Equilibria in Candidate Solder Alloys”, Z. Metallkd., 92(7), 740-746 (2001) (Phase Relations, Thermodyn., Calculation, Review 26) Kikuchi, S., Nishimura, M., Suetsugu, K., Ikari, T., Matsushige, K., “Strength of Bonding Interface in Lead-Free Sn Alloy Solders”, Mater. Sci. Eng. A, A319-321, 475-479 (2001) (Experimental, Mechan. Prop., 4) Lee, J.H., Lee, D.N., “Use of Thermodynamic Data to Calculate Surface Tension and Viscosity of Sn-Based Soldering Alloy Systems”, J. Electron. Mater., 30(9), 1112-1119 (2001) (Experimental, Calculation, Phys. Prop., Thermodyn., 26) Moser, Z., Gasior, W., Pstrus, J., “Surface Tension Measurements of the Bi-Sn and Sn-Bi-Ag Alloys”, J. Electron. Mater., 30(9), 1104-1111 (2001) (Experimental, Mechan. Prop., Interface Phenomena, 33) Ohtani, H., Satoh, I., Miyashita, M., Ishida, K., “Thermodynamic Analysis of the Sn-Ag-Bi Ternary Phase Diagram”, Mater. Trans., JIM, 42(5), 722-731 (2001) (Phase Diagram, Phase Relations, Thermodyn., Assessment, #, 15) Suganuma, K., “Advances in Lead-Free Electronics Soldering”, Curr. Opin. Solid State Mater. Sci., 5, 55-64 (2001) (Experimental, Mechan. Prop., 75) Takao, H., Hasegawa, H., “Influence of Alloy Composition on Fillet-Lifting Phenomenon in Sn-Ag-Bi”, J. Electron. Mater., 30(9), 1060-1067 (2001) (Calculation, Phase Relations, 9) Choi, J.-W., Cha, H.-S., Oh, T.-S., “Mechanical Properties and Shear Strength of Sn-3.5Ag-Bi Solder Alloys”, Mater. Trans., JIM, 43(8), 1864-1867 (2002) (Experimental, Mechan. Prop., 15) Kattner, U.R., “Phase Diagrams for Lead-Free Solder Alloys”, JOM, 54(12), 45-51 (2002) (Phase Relations, Review, #, 47) Shimokawa, H., Soga, T., Serizawa, K., “Mechanical Properties of Tin-Silver-Bismuth Lead-Free Solder”, Mater. Trans., JIM, 43(8), 1808-1815 (2002) (Experimental, Mechan. Prop., 10) Tanabe, T., Harada, S., “Examination of Compound Formation at Interface of Tin-Bismuth-Silver Solder and Copper Substrate by Using Electron Probe Micro Analysis”, X-Ray Spectrom., 31, 3-6 (2002) (Electronic Structure, Phase Relations, Interface Phenomena, 17) Xia, Z., Chen, Z., Shi, Y., Mu, N., Sun, N., “Effect of Rare Earth Element Additions on the Microstrucutre and Mechanical Properties of Tin-Silver-Bismuth Solder”, J. Electron. Mater., 31(6), 564-567 (2002) (Experimental, Mechan. Prop., Morphology, Interface Phenomena, 9) Xia, Z., Shi, Y., Chen, Z., “Evaluation on the Characteristics of Tin-Silver-Bismuth Solder”, J. Mater. Eng. Perform., 11(1), 107-111 (2002) (Experimental, Review, Phys. Prop., Mechan. Prop., 10)
Landolt-Börnstein New Series IV/11C3
Ag–Bi–Sn [2002Yu]
[2003Lai]
[2003Ohn]
[2003Shi]
[2004Gas]
[2004Hwa]
[2004Oht]
[2004Sho]
[2004Wu]
9
Yu, J., Shin, S.W., “Rupture Time Analyses of the Sn-3.5Ag Solder Alloys Containing Cu or Bi”, Acta Mater., 50, 4315-4324 (2002) (Experimental, Mechan. Prop., Interface Phenomena, 23) Lai, H.L., Duh, J.G., “Lead-Free Sn-Ag and Sn-Ag-Bi Solder Powders Preparation by Mechanical Alloying”, J. Electron. Mater., 32(4), 215-220 (2003) (Experimental, Phys. Prop., Phase Relations, 46) Ohnuma, I., Miyashita, M., Liu, X.J., Ohtani, H., Ishida, K., “Phase Equilibria and Thermodynamic Properties of Sn-Ag Based Pb-Free Solder Alloys”, IEEE Trans., Electron. Pack. Manuf., 26(1), 84-89 (2003) (Phase Relations, Thermodyn., Assessment, Review, #, 21) Shiue, R.-K., Tsay, L.-W., Lin, Ch.-L., Ou, J.-L., “A Study of Sn-Bi-Ag-(In) Lead-Free Solders”, J. Mater. Sci., 38(6), 1269-1279 (2003) (Experimental, Interface Phenomena, Morphology, Phase Relations, 38) Gasior, W., Moser, Z., Pstrus, J., Bukat, K., Kisiel, R., Sitek, J., “(Sn-Ag)eut + Cu Soldering Materials, Part I: Wettability Studies”, J. Phase Equilib. Diffus., 25(2), 115-121 (2004) (Crys. Structure, Experimental, Interface Phenomena, Morphology, Phase Relations, Phys. Prop., 17) Hwang, C.-W., Suganuma, K., “Joint Reliability and High Temperature Stability of Sn-Ag-Bi Lead-Free Solder with Cu and Sn-Pb/Ni/Cu Substrates”, Mater. Sci. Eng. A, A373, 187-194 (2004) (Experimental, Mechan.. Prop., 15) Ohtani, H., Ono, S., Doi, K., Hasebe, M., “Thermodynamic Study of Phase Equilibria in the Sn-Ag-Bi-Zn Quaternary System”, Mater. Trans., 45(3), 614-624 (2004) (Thermodyn., Assessment, *, #, 25) Shohji, I., Yoshida, T., Takahashi, T., Hioki, S., “Tensile Properties of Sn-Ag Based Lead-Free Solders and Strain Rate Sensitivity”, Mater. Sci. Eng. A, A366, 50-55 (2004) (Experimental, Mechan. Prop., 17) Wu, C.M.L., Yu, D.Q., Law, C.M.T., Wang, L., “Properties of Lead-Free Solder Alloys with Rare Earth Element Additions”, Mater. Sci. Eng. R, 44, 1-44 (2004) (Mechan. Prop., Interface Phenomena, Review, 100)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
at 25°C [Mas2] Dissolves up to 2.76 at.% Bi [1993Kar], ~11.3 at.% Sn at 722.5°C [1999Oht]
(Bi) < 271.442
hR6 R3m As
a = 454.613 c = 1186.152
at 31°C [V-C2] Dissolves up to 3.8 at.% Sn at 140.2°C [1994Oht]
(Sn) 231.9681 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
at 25°C [Mas2] Dissolves ~16 at.% Bi at 140.2°C [1994Oht]
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
[Mas2]
Landolt-Börnstein New Series IV/11C3
MSIT®
Ag–Bi–Sn
10 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, (Ag-Sn) < 722.5
hP2 P63/mmc Mg
a = 296.58 c = 478.24
[1999Oht], [V-C2]
Ag3Sn < 476.5
oP8 Pmmn Cu3Ti
a = 478.02 b = 596.8 c = 518.43
[1999Oht], [V-C2] 23.5-25 at.% Sn [Mas2] Dissolves ~19 mass% Bi at 200°C (predicted) [2001Oht]
Table 2: Invariant Equilibria T [°C]
Reaction
Type
Phase
Composition (at.%) Ag
Bi
Sn
L (Ag) + (Bi)
264?
e1 (max)
-
-
-
-
L + (Ag) + (Bi)
263.6
U1
L
2.0
97.7
0.3
L + Ag3Sn + (Bi)
262.5
U2
L
2.6
97.3
0.1
L (Bi) + Ag3Sn + (Sn)
139.2
E
L
0.5
53.9
45.6
1000
Fig. 1: Ag-Bi-Sn. The Ag-Sn binary system
L
Temperature, °C
750
722.5
500
476.5
ζ 250
220.5 (β Sn)
0
Sn
(Ag)
Ag3Sn
20
40
60
80
Ag
Ag, at.%
MSIT®
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Ag-Bi
Ag-Bi-Sn
A-B-C
Ag-Sn
Bi-Sn
722.5 p1 l + (Ag) ζAg 476.5 p2 l + ζ Ag3Sn
264? e1(max) L (Ag) + (Bi) 263.6
L + (Ag) ζ + (Bi)
U1
(Ag)+ζ+(Bi) 262.3 e2 l (Ag) + (Bi)
262.5
L + ζ Ag 3Sn+(Bi)
U2
Ag–Bi–Sn
ζ+Ag3Sn+(Bi)
220.5 e3 l Ag3Sn + (βSn) 140.2 e4 l (Bi) + (βSn)
139.2 L (Bi) + Ag3Sn+(βSn)
E
(βSn)+Ag3Sn+(Bi)
11
MSIT®
Fig. 2: Ag-Bi-Sn. Reaction scheme
Ag–Bi–Sn
12
Sn Fig. 3: Ag-Bi-Sn. Liquidus surface projection
Data / Grid: at.% Axes: at.%
e3 (β Sn) 20
80
Ag3Sn
40
60
427
377°C
p2
e4
E
60
40
510
80
605
p1
ζ
(Bi) 20
477 (Ag) 20
Ag
Fig. 4: Ag-Bi-Sn. Liquidus lines near the Sn corner
40
60
U2 U1
e2
80
7
Bi
300 290
280
6
270 Ag3Sn
260
5
250
Ag, mass%
240 4
230 220 210
3
200
(βSn) 2
220
1
210
200 190
230
Sn
10
20
30
Bi, mass%
MSIT®
180
40
170 160 50
150 E e4
(Bi) Bi 60.00 Sn 40.00
Landolt-Börnstein New Series IV/11C3
Ag–Bi–Sn
13
Sn
Data / Grid: at.% Axes: at.%
Fig. 5: Ag-Bi-Sn. Calculated metastable miscibility gap in the liquid at 200 < T < 400°C
L1
20
80
200
40
60
250 300
60
40
400 80
20
L2 20
Ag
40
60
80
Sn
Bi
Data / Grid: at.% Axes: at.%
Fig. 6: Ag-Bi-Sn. Isothermal section at 400°C 20
L
80
40
60
60
40
Ag3Sn 80
20
ζ
L+Ag3Sn+ζ (Ag)+L+ζ
(Ag)
Ag
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Bi
MSIT®
Ag–Bi–Sn
14
Sn Fig. 7: Ag-Bi-Sn. Isothermal section at 200°C
Data / Grid: at.% Axes: at.%
(β Sn)
20
80
(β Sn)+L+Ag3Sn
L 40
60
60
40
Ag3Sn
(Bi)+L+Ag3Sn
80
20
ζ (Ag)
(Bi) 20
Ag
40
60
80
Bi
Fig. 8: Ag-Bi-Sn. Isopleth at 10 at.% Bi 750
L
Temperature, °C
L+(Ag) L+ζ 500
L+Ag3Sn+ζ 263.6
L+Ag3Sn
Ag3Sn
250
(β Sn)+Ag3Sn+L
L+Ag3Sn+(Bi)
139.2 (Bi)+Ag3Sn+ζ
(Bi)+(Ag) 0
Ag 90.00 Bi 10.00 (Bi)+(Ag)+ζ Sn 0.00
MSIT®
20
40
Sn, at.%
(β Sn)+Ag3Sn+(Bi) 60
80
Ag 0.00 Bi 10.00 Sn 90.00
Landolt-Börnstein New Series IV/11C3
Ag–Bi–Sn
Fig. 9: Ag-Bi-Sn. Isopleth at 20 mass% Bi, plotted in at.%
15
750
L
Temperature, °C
L+(Ag) L+ζ
500
L+Ag3Sn+ζ 263.6
L+Ag3Sn
Ag3Sn+(Bi)
250
L+Ag3Sn+(Bi)
(β Sn)+Ag3Sn+L 139.2
Ag3Sn+(Bi)+ζ
(Bi)+(Ag) 0
Ag 88.57 Bi 11.43 (Ag)+(Bi)+ζ Sn 0.00
20
(β Sn)+Ag3Sn+(Bi) 40
60
80
Sn, at.%
Ag 0.00 Bi 12.44 Sn 87.56
500
Fig. 10: Ag-Bi-Sn. Isopleth at 10 mass% Ag, plotted in at.%
L L+ζ
Temperature, °C
L+(Ag)
Ag3Sn+L 250
(β Sn)+Ag3Sn+L L+Ag3Sn+(Bi) 139.2 (β Sn)+Ag3Sn
Ag3Sn+(Bi) (β Sn)+Ag3Sn+(Bi)
(Bi)+(Ag) 0
Ag 17.71 Bi 82.29 Sn 0.00
Landolt-Börnstein New Series IV/11C3
20
40
Sn, at.%
60
80
Ag 10.90 0.00 Bi Sn 89.10
MSIT®
Ag–Bi–Sn
16
Temperature, °C
Fig. 11: Ag-Bi-Sn. Isopleth at 30 mass% Ag, plotted in at.%
L 500
L+ζ
Ag3Sn+L
L+(Ag)
250
(β Sn)+Ag3Sn+L L+Ag3Sn+(Bi) 139.2 Ag3Sn+(Bi)
(β Sn)+Ag3Sn (β Sn)+Ag3Sn+(Bi)
(Bi)+(Ag) 0
Ag 45.36 Bi 54.64 Sn 0.00
20
40
60
Sn, at.%
Ag 32.05 0.00 Bi Sn 67.95
Fig. 12: Ag-Bi-Sn. Isopleth at 70 at.% Bi L
Temperature, °C
500
L+(Ag)
L+ζ L+Ag3Sn
250
Ag3Sn+(Bi) L+Ag3Sn+(Bi) 139.2
0
Ag 30.00 Bi 70.00 Sn 0.00
MSIT®
(β Sn)+Ag3Sn+(Bi)
Ag3Sn+(Bi)+ζ
(Bi)+(Ag)
(Bi)+(Ag)+ζ
20
10
Ag, at.%
Ag 0.00 Bi 70.00 Sn 30.00
Landolt-Börnstein New Series IV/11C3
Ag–Bi–Sn
17
Fig. 13: Ag-Bi-Sn. Isopleth at 40 at.% Ag
L 500
Temperature, °C
L+ζ
L+(Ag) L+Ag3Sn
250
(β Sn)+Ag3Sn+L
(Bi)+ζ L+Ag3Sn+(Bi) 139.2 Ag3Sn+(Bi) (β Sn)+Ag3Sn+(Bi)
(Bi)+(Ag) 0
Ag 40.00 Bi 60.00 Sn 0.00
20
(β Sn)+Ag3Sn
Ag 40.00 0.00 Bi Sn 60.00
40
Sn, at.%
Sn
Data / Grid: at.% Axes: at.%
Fig. 14: Ag-Bi-Sn. Experimental isoenthalpic lines at 605°C. Full line: liquidus line at 605°C
20
80
500°C 0 40
60
-500 1000 -1500 60
40
-2500 1500 80
20
2000
Ag
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Bi
MSIT®
18
Ag–Cu–In
Silver – Copper – Indium Ortrud Kubaschewski, Alan Prince†, updated by Joachim Gröbner Introduction [1951Geb] and [1952Geb] investigated phase relationships in the indium poor region of this system by thermal analysis, metallography and electrical conductivity measurements. They established six four-phase equilibrium reactions and represented their findings by constructing three isothermal sections and four isopleths. [1955Val], [1967Mcd] and [1979Dri] reviewed the work of [1951Geb] and [1952Geb]. Later work by [1988Woy] is in substantial disagreement with [1951Geb] and [1952Geb]. The 1 phase occurring at high temperatures in the Cu-In system, forms a solid solution series with the low-temperature 2 phase of the Ag-In system. [1951Geb] and [1952Geb] assumed the bcc high temperature phases of Cu-In(1) and Ag-In(2) both to be stabilized in the ternary and to decompose eutectically at 490 and 475°C, respectively. [1988Woy] found both phases to be already decomposed at 505°C. Since [1951Geb] and [1952Geb] identified the phases by metallography only and not by X-ray diffraction, a partial explanation of the disagreement with [1988Woy] may be a misinterpretation of the phase as ternary solutions of the phases. Nevertheless only part of the micrographs of annealed specimen published by [1952Geb] can be interpreted by the isothermal sections given by [1988Woy]. Binary Systems The binary Ag-Cu system was taken from the recent, published in the MSIT Evaluation Program by [2002Rom]. Ag-In and Cu-In are accepted from [Mas2]. Solid Phases The solid phases of Ag-In and Cu-In are listed in Table 1. No ternary phases have been reported. Invariant Equilibria Almost the entire reaction scheme can be derived from the liquidus surface given by [1988Woy] (Fig. 1) but without temperatures for the four-phase equilibria. This disagrees with the partial reaction scheme given by [1952Geb] in as far the two phases react and disappear before the L + 1 + 2 three-phase field reacts with another three-phase field. Since temperatures of Gebhardt's [1952Geb] scheme are tentatively assigned to the invariant equilibria, which probably explain the thermal analysis effects of Gebhardt's isopleths, it is not possible to reinterpret all measured points of these isopleths to be consistent with the findings of [1988Woy]. Invariant solid-state reactions were not studied by [1988Woy]. The superstructure phases are ignored in the reaction scheme. Liquidus Surface In Fig. 2 the liquidus surface according to [1988Woy] is shown. In the In poor part some isotherms from [1951Geb] are inserted. For comparison the partial liquidus surface published by [1951Geb] is shown in Fig. 3. Isothermal Sections In Fig. 4 and Fig. 5 the 676 and 505°C isothermal sections as published by [1988Woy] are shown. Figure 5 shows the overwhelming influence of the phase on the ternary equilibria. The 676°C isothermal section (Fig. 4) does not correspond with the isopleths at 20 and 25 mass% In (Fig. 7 and Fig. 8). In particular [1952Geb] indicates that the liquid phase occurs at 676°C on the 20 mass% In section. [1988Woy] shows the liquidus isotherm located at higher In contents. It should be noted that the liquid apexes of the four tie triangles in Fig. 4 do not correspond with points on the monovariant curves in Fig. 2. The L + + field in MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–In
19
the 505°C isotherm is changed because the original disagrees with the liquidus surface. The L + 2 + 2 field in Fig. 4 fits more closely Gebhardt's liquidus surface (Fig. 3) than to Fig. 2. Gebhardt's 500°C isothermal section is shown in Fig. 6 for comparison. Temperature – Composition Sections Figures 7 and 8 show the isopleths at 20 and 25 mass% In, respectively, according to [1952Geb]. Notes on Materials Properties and Applications The effect of quenching on the stress-strain characteristics of Ag alloy with 2 mass% Cu and 0.5 mass% In was investigated in the temperature range from 500 to 650°C by [2004Nad]. References [1951Geb] [1952Geb] [1955Val] [1967Mcd] [1972Jai]
[1979Dri]
[1988Woy] [1994Sub]
[2002Rom]
[2004Nad]
Gebhardt, E., Dreher, M., “On the Constitution of the Copper-Silver-Indium System” (in German), Z. Metallkd., 42, 230-238 (1951) (Experimental, Phase Relations, 6) Gebhardt, E., Dreher, M., “On the Constitution of the Copper-Silver-Indium System”, Z. Metallkd., 43, 357-363 (1952) (Experimental, Phase Relations, 3) Valentiner, S., “Indium Alloys” (in German), Z. Metallkd., 46, 442-449 (1955) (Review, 86) McDonald, A.S., Price, B.R., Sistare, G.H., “Ternary and Higher-Order Alloys of Silver”, Silver-Economics, Metallurgy and Use, 272-303 (1967) (Phase Diagram, Review, 38) Jain, K.C., Ellner, M., Schubert, K., “On Phases in the Vicinity of the Composition Cu64In36” (in German), Z. Metallkd., 63, 456-461 (1972) (Experimental, Crys. Structure, Phase Relations, 6) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Cu-In-Ag” in “Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 128-129 (1979) (Phase Diagram, Review, 3) Woychik, C.G., Massalski, T.B., “Phase Stability Relationships and Glass Formation in the System Cu-Ag-In”, Metall. Trans., A19, 13-21 (1988) (Experimental, Phase Relations, 13) Subramanian, P.R., “Cu (Copper)”, in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 1-3 (1994) (Crys. Structure, Review, 16) Van Rompaey, T., Rogl, P., “Ag-Cu (Silver - Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14511.1.20, (2002) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 29) Nada, R.H., “Precipitation Kinetics in Ag-2wt.% Cu and Ag-2wt.% Cu-0.5wt.% In Alloys During Transformation”, Physica B, 349(1-4), 166-173 (2004) (Experimental, Kinetics, Mechan. Prop., Morphology, Phase Relations, 22)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
2, (Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
at 25°C [Mas2]
1, (Cu) < 1084.62
cF4 Fm3m Cu
a = 361.46
at 25°C [Mas2] melting point [1994Sub]
Landolt-Börnstein New Series IV/11C3
MSIT®
Ag–Cu–In
20 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(In) < 156.634
tI2 I4/mmm In
a = 325.3 c = 494.7
at 25°C [Mas2]
2, Ag3In(h2) 695 - 660
cI2 Im3m W
a = 336.82
[V-C2]
, Ag3In(r) < 670
hP2 P63/mmc Mg
a = 295.63 c = 478.57
[Mas2]
’, Ag3In < 187
cP4 Pm3m AuCu3
a = 414.4 0.4
[Mas2]
, (Cu,Ag)9In4
cP52 P43m Cu9Al4
a = 992.2
[V-C2]
a = 925.03
[V-C2]
2, Ag9In4(h) 314 1, Cu9In4(h) 684 - 614 ’, AgIn2 < 166
tI12 I4/mcm Al2Cu
a = 688.1 c = 562
[V-C2]
1, Cu4In(h) 710 - 574
cI2 Im3m W
a = 304.6
[Mas2, V-C]
, Cu7In3(r) < 631
aP40 P1 Cu7In3
a = 1007.1 b = 912.6 c = 672.4 = 90.22° = 82.84° = 106.81°
[V-C2]
, Cu2In(h3) 667 - 440
hP6 P63/mmc Ni2In
a = 412.0 c = 526.3
[V-C2]
Cu7In4(h2) 480 - 350
oP55 P*
a = 2137.5 b = 740.5 c = 521.8
[1972Jai] superstructure of the Ni2In type
Cu7In4(h1) 450 - 298
oP88 P*
a = 3419.4 b = 739.5 c = 526.2
[1972Jai] superstructure of the Ni2In type
Cu7In4(r) < 390
-
-
[1972Jai]
Cu15In8 < 355
-
-
[1972Jai]
Q, Cu11In9 < 310
mC20 C2/m CuAl
a = 1281.4 b = 435.43 c = 735.3 = 54.49°
[V-C2]
MSIT®
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Ag-Cu 779.1 e1 l α1 + α2
Ag-In
Ag-Cu-In
Cu-In 710 p1 l + α1 β1
695 p2 l + α2 β2
678 e2 l β1 + γ
670 p4 α2 + β2 ζ 660 e3 β2 ζ + l
β2 L + α2 + ζ
L + α2 γ + ζ
U3
β1 α1 + γ + δ
L+γ+ζ <505
614 e5 γ δ + η
β1+α1+γ
L+α2+γ
α2+γ+ζ
205 e7 ζγ+l
U2
α1+α2+γ
L+γη+ζ L+η+ζ
ζL+γ+η
470?
E2
574 e6 β1 α1 + δ
α1+γ+δ
U4
γ+η+ζ
616 e4 γ β1 + δ
α1 + γ α2 + δ α1+α2+δ
U5
γ+α2+δ
Ag–Cu–In
575?
U1
L+α1+γ
L + α1 α2 + γ
600? L+α2+ζ
L + β1 α1 + γ
E1
667 p3 l + γ η
310 p5 l + ηϕ
E3
L+γ+η L+ηγ+ϕ 166 p6 γ + l AgIn2 144 e9 l (In) + AgIn2
γ+η+ϕ
L+γ+ϕ L + γ ϕ + AgIn2 L+ϕ+AgIn2 L (In) + ϕ + AgIn2
U7
γ + ϕ + AgIn2 E4
U6
153 e8 l (In) + ϕ
Fig. 1: Ag-Cu-In. Reaction scheme
21
MSIT®
(In)+ϕ+AgIn2
Ag–Cu–In
22
In
E4 e8 (In) e
Fig. 2: Ag-Cu-In. Liquidus surface projection [1988Woy]
ϕ
20
U7 p5 U 6 E3
Data / Grid: at.% Axes: at.%
9
AgIn2 p6
γ e7 80
40
60
η ζ 60
40
U4
p3 e2 80
p1
γ U1
β1
e3
U3 E1
β2
p2
U2
20
α2
750
α1
850
850
950°C 20
Cu
e1 60
40
80
In
Ag
Data / Grid: at.% Axes: at.%
Fig. 3: Ag-Cu-In. Indium poor liquidus surface projection [1951Geb] 20
80
40
60
60
40
β1 p1
U
80
850
p2 20
750 950
U
p3
β2
650
(Ag)
(Cu)
850
1050°C
Cu
MSIT®
20
40
60 e
1
80
Ag
Landolt-Börnstein New Series IV/11C3
Ag–Cu–In
23
In
Data / Grid: at.% Axes: at.%
Fig. 4: Ag-Cu-In. Isothermal section at 676°C 20
80
L 40
60
60
40
L+γ
γ
L+β 1
γ +β 1 80
β1
L+β 2
β2
L+α1
β 2+α2
20
L+α2
α1+β 1
α2
α1
α1+α2 20
Cu
40
60
80
In
Ag
Data / Grid: at.% Axes: at.%
Fig. 5: Ag-Cu-In. Isothermal section at 505°C 20
80
L 40
60
L+η
η
L+ζ
60
γ +η
δ+η
L+γ
δ
40
γ +ζ γ
ζ
ζ +α 2
80
α 1+δ
α1+δ+γ
α1
Cu
Landolt-Börnstein New Series IV/11C3
20
α 1+α 2+γ
α2
α1+α2 20
40
60
80
Ag
MSIT®
Ag–Cu–In
24
In
Data / Grid: at.% Axes: at.%
Fig. 6: Ag-Cu-In. Isothermal section at 500°C 20
80
40
60
60
40
β2
(Cu)+β 1+δ
ζ
β1
80
(Ag)+β 2+ζ
20
(Cu)+(Ag)+β 2
(Cu)+β 1+β 2
(Ag)
(Cu) 20
Cu
Fig. 7: Ag-Cu-In. Isopleth at 20 mass% In, plotted in at.%
40
60
80
Ag
1000
900
L
Temperature, °C
800
700
600
500
L+(Cu) L+(Ag) L+(Cu)+β 1
L+(Ag)+β 2
L+(Cu)+β 2 (Cu)+β 1
L+(Ag)+(Cu)
(Cu)+β 2
(Cu)+(Ag)+β 2
(Cu)+β 1+β 2
(Ag)+β 2 (Ag)
(Cu)+β2+δ 400
(Cu)+δ
300
Cu 87.90 Ag 0.00 In 12.10
MSIT®
(Cu)+(Ag)+δ
(Cu)+β 1+δ
10
(Ag)+δ 20
30
40
Ag, at.%
50
60
70
0.10 Ag 80.90 In 19.00
80 Cu
Landolt-Börnstein New Series IV/11C3
Ag–Cu–In
Fig. 8: Ag-Cu-In. Isopleth at 25 mass% In, plotted in at.%
25
900
800
L
Temperature, °C
L+α
L+α '
700
L+α +β 600
L+α '+β '
L+α +α '
β '+γ'
L+α +β '
α +β
L+β '
α +β +δ
α +β +β '
500
α +β '
α '+γ'
δ +α '+α
Landolt-Börnstein New Series IV/11C3
10
α '+δ
α +β '+δ
400
Cu 84.40 Ag 0.00 In 15.60
α '+β '+γ' β'
α +α '+β '
α +δ +β '
α +δ
300
γ'
20
30
α '+δ +γ' 40
Ag, at.%
50
60
70
Cu 0.00 Ag 76.10 In 23.90
MSIT®
26
Ag–Cu–Mn
Silver – Copper – Manganese Ortrud Kubaschewski, updated by Lazar Rokhlin, Nataliya Bochvar, Evgeniya Lysova Introduction Based on the binary systems related to the ternary Ag-Cu-Mn system, [1909Jae] supposed the liquidus surface in this system with a minimum at about centre of the concentration triangle. This feature of the phase diagram was not still confirmed completely by the experiments. The only one experimental investigation was made by [1931Kei] who studied the ternary Ag-Cu-Mn system using mainly metallography supported by thermal analysis. [1931Kei] established extension of the liquid miscibility gap from the binary Ag-Mn system into the ternary system. Its boundaries were determined by chemical analysis. Besides, [1931Kei] assumed existence of a three-phase monovariant equilibrium in the Ag-Cu-Mn system which changed from the peritectic equilibrium in the Ag-Mn system to eutectic one in the Ag-Cu system. However, in the work [1931Kei] the reactions between the liquid phase and different allotropic forms of Mn in the Ag-Mn system were not taken into consideration. These reactions were shown later in the review of the Ag-Cu-Mn system [1977Cha]. [2002Jac] measured the melting temperature of the only 70Cu-25Mn-5Ag (mass%) alloy by differential thermal analysis. This temperature turned out between 800 and 830°C in agreement with [1931Kei]. [1949Jae, 1979Dri, 1980Bra] presented the reviews of the ternary Ag-Cu-Mn system basing on [1931Kei] without critical analysis. Binary Systems The binary Ag-Cu, Ag-Mn and Cu-Mn systems are accepted from [2002Rom], [2002Maz] and [2005Tur], respectively. Solid Phases No ternary compounds have been observed in the Ag-Cu-Mn system. The crystal structure and lattice parameters of the unary and binary phases of the system are presented in Table 1. Liquidus Surface On the liquidus surface the boundary of the miscibility gap L1+ L2 was determined only experimentally. It is shown in Fig. 1 after the publication [1977Cha], with some corrections of the compositions and temperatures of the invariant reactions taking into account the accepted binary Ag-Cu, Ag-Mn and Cu-Mn systems. The binary Ag-Mn liquid miscibility gap extends into the ternary phase diagram to a maximum at about 32 mass% Cu (~35 at.% Cu) at an unspecified temperature. Other features of liquidus surface were not established. Notes on Materials Properties and Applications Commercial alloy “Hi-Temp 095” with the nominal composition of 9.5Ag-52.5Cu-38Mn is used as high strength filler metal for joining carbides, steels and heat resistant alloys. Its liquidus temperature is 926°C, its solidus temperature is 879°C. Ag-Cu-Mn alloys are the part of the commercial alloy “WB850” (65Ag-28Cu-5Mn-2Ni) used for electronic packaging. [1942And] investigated a series of the Ag-Cu-Mn alloys aiming to substitute nickel in the present 5 cent coin. The resistivity, color, density and resistance to corrosion were determined for Cu rich alloys with Mn and Ag from 4 to 50 mass%. The diagrams characterizing dependence of the properties on the alloy compositions were presented.
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Mn
27
[1931Kei] measured the hardness of Ag based alloys with 2, 3, 4 mass% Cu and 15, 10, 5 mass% Mn after homogenization at 750°C, quenching and ageing at 280°C for 2, 5, 10 and 20 h. The maximum of hardness was reached on 91Ag-4Cu-5Mn (mass%) alloy after ageing at 280°C for 5 h. The increase of Mn content led to decrease of hardness. [2002Jac] investigated the properties of the 70Cu-25Mn-5Ag (mass%) alloy. The values of hardness 130 Hv and tensile strengths of 150 to 230 MPa were obtained. Miscellaneous [1969Gue] considered the ternary Ag-Cu-Mn system introduced by [1931Kei] and gave some comments on it. According to [1969Gue] the Mn corner of the system is scientifically interesting owing to the polymorphism of Mn and to its very different dissolving power for Ag and Cu atoms. Formation between Cu and Mn an ultimate series of the solid solutions in a definite temperature range might display special properties in presence of Ag in the alloys. Such properties could be brightness, color of copper, patina and chemical behavior. References [1909Jae] [1931Kei] [1942And] [1949Jae] [1969Gue]
[1977Cha]
[1979Dri]
[1980Bra] [2002Jac]
[2002Maz]
[2002Rom]
Landolt-Börnstein New Series IV/11C3
Jänecke, E., “Ternary Alloys of Cu, Ag, Au; Cr, Mn; Fe, Co, Ni; Pd, Pt Metals” (in German), Z. Phys. Chem., 67, 668-688 (1909) (Phase Diagram, Theory, 23) Keinert, M., “The System Silver - Copper - Manganese” (in German), Z. Physik. Chem., A156, 291-303 (1931) (Experimental, Phase Diagram, Phase Relations, Mechan. Prop., 10) Anderson, T., “Possible Substitutes for Nickel in the Present Five-Cent Coin”, US Dep. Int.Bur. Mines, 1-6 (1942) (Experimental, Mechan. Prop., Electr. Prop.) Jänecke, E., “Cu-Ni-Ag and Ag-Cu-Mn” (in German), Kurzgefasstes Handbuch aller Legierungen, Winter, Heidelberg, 406 (1949) (Phase Diagram, Review) Guertler, W., Guertler, M., Anastasiadias, E., “Copper-Silver-Manganese”, in “A Compendium of Constitutional Ternary Diagrams of Metallic Systems”, WADC Technical Report 58-615, Project No 7351, Wright-Patterson Air Force Base, Ohio, 126-127 (1969) (Theory, 1) Chang, Y.A., Goldberg, D., Neumann, J.P., “Phase Diagrams and Thermodynamic Properties of Ternary Copper - Silver Systems”, J. Phys. Chem. Ref. Data, 6, 621-673 (1977) (Review, Phase Diagram, #, 96) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Copper-Manganese-Silver” in “Binary and Multicomponent Copper-Base Systems” (in Russian), Nauka, Moscow, 211 (1979) (Phase Diagram, Phase Relations, Review, 1) Brandes, E.A., Flint, R.E., “Manganese Phase Diagram”, Manganese Centre, Paris, France, 72 (1980) (Phase Diagram, Review, 2) Jacobson, D.M., Sangha, S.P.S., Gales, A., Schmid, E.E., “The Development of New Silver-Free Brazing Alloys for Steel Tubular Assembly”, Welding J., (8), 149-155 (2002) (Experimental, Mechan. Prop., 7) Mazzone, D., “Ag-Mn (Silver - Manganese)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.24021.1.20, (2002) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 3) Van Rompaey, T., Rogl, P., “Ag-Cu (Silver - Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14511.1.20, (2002) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 29)
MSIT®
Ag–Cu–Mn
28 [2005Tur]
Turchanin, M., Agraval, P., Groeber, J., Matusch, D., Turkevich, V., “Cu-Mn (Copper - Manganese)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14136.1.20, (2005) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 25)
Table 1: Crystallographic Data of Solid Phases Phase Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
pure Ag at 25°C [Mas2] dissolves up to 14.1 at.% Cu at 779.1°C [2002Rom] and 43.6 at.% Mn at 989°C [2002Maz]
( Mn) 1246 - 1138
cI2 Im3m W
a = 308.0
pure Mn at >1138°C [Mas2] dissolves up to 13 at.% Cu at 1097°C [2005Tur] and 1.2 at.% Ag at 1207°C [2002Maz]
( Mn,Cu)
cF4 Fm3m Cu
( Mn) 1138 - 1087
a = 386.0
pure Mn [Mas2] dissolves up to 100 at.% Cu [2005Tur] and 0.7 at.% Ag at 1128°C [2002Maz]
(Cu) < 1084.62
a = 361.46
pure Cu at 25°C [Mas2] dissolves up to 5 at.% Ag at 779.1°C [2002Rom] and up to 100 at.% Mn [2005Tur]
(Mn) 1087 - 707
cP20 P4132 Mn
a = 631.52
pure Mn [Mas2] dissolves up to 2 at.% Cu at 706°C [2005Tur] and 0.4 at.% Ag at 1069°C [2002Maz]
(Mn) < 707
cI58 I43m Mn
a = 891.26
pure Mn at 25°C [Mas2] dissolves up to 2 at.% Cu at 706°C [2005Tur]
3 (Cu-Mn) 700
c*
-
[2005Tur]
2, MnCu3 450
c**
-
[2005Tur]
3 (Cu-Mn) 410
c**
-
[2005Tur]
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Mn
29
Cu
Data / Grid: at.% Axes: at.%
Fig. 1: Ag-Cu-Mn. Liquidus surface projection
1000 20
80
900°C 863°C min 40
60
60
(γ Mn,Cu)
e4,779.1 40
p1,1097 L1+L2
80
(γ Mn) (β Mn) (δMn)
(δ Mn)
Mn
Landolt-Börnstein New Series IV/11C3
20
e'1,1207
20
40 e'' e ,1128 1 2
60 p2,989 e3,1069
(Ag)
80
Ag
MSIT®
30
Ag–Cu–Ni
Silver – Copper – Nickel Ortrud Kubaschewski, updated by Myriam Sacerdote-Peronnet Introduction The system was studied initially by [1913Ces] and [1933Gue]. The latter determined the liquid miscibility gap by chemical analysis and their findings agree, in principle, with those of [1913Ces]. The results obtained by [1913Ces] and [1933Gue] are presented in the review published by [1967McD], which deals with silver alloys. [1977Cha] adjusted the liquidus surface formerly based on [1933Gue], in order to obtain consistency with the currently accepted binary data. By applying improved equipment and techniques, [1977Sie] redetermined the phase equilibria; their results show a somewhat lower ternary solubility limit for Ni below 1000°C. [1993Che] and [1996Luo] applied the diffusion couple method to the study of the interfacial reactions between liquid Ag-Cu alloys and pure nickel substrates. The results obtained at 860, 795 and 700°C confirm that (Ag) has almost no solubility of Ni. Binary Systems The Ag-Cu, Ag-Ni and Cu-Ni systems are taken from the recent evaluations made within the MSIT Binary Evaluation Program, by [2002Rom], [2002Iva] and [2002Leb], respectively. Solid Phases No intermediate phases have been reported in the binary systems or in the ternary system. The unary phases are listed in Table 1. Invariant Equilibria The reaction sequence is illustrated in Fig. 1. The three-phase equilibrium L1 + L2 + (Ni,Cu) disappears at a critical point of ~1250°C. Liquidus Surface According to the chemical analysis data of [1933Gue], the liquid miscibility gap extends from the binary Ag-Ni system to approximately 48 at.% Cu (42 mass% Cu) at 1250°C; see Fig. 2. The locations of some of the isotherms have been modified at the binary edges to ensure agreement with the accepted binary systems. Isothermal Sections Figures 3 to 7 present the results obtained by [1933Gue] based on chemical, thermal and metallographic analyses, for temperatures 1440, 1400, 1300, 1250 and 900°C. Some modifications have been made to ensure agreement with the accepted binary diagrams. The solubility of Ni in the liquid phase at 900°C is not accepted in this evaluation, following the more recent findings of [1996Luo] for a temperature of 860°C. Based on the review of [1977Cha] and the calculation of [1991Mur], the isothermal section for 793°C was constructed by [1993Che]. More recently, [1996Luo] confirmed the isothermal section given in [1993Che] and also determined phase equilibria at 860 and 700°C, by both experimental investigation and calculation. It was shown that (Ag) has some solubility of Cu but almost no solubility of Ni and the liquid phase has almost no Ni solubility [1993Che, 1996Luo, 1977Sie]. The isothermal sections established at 860, 795 in [1996Luo] are given in Figs. 8 and 9. The isothermal section for 700°C shown in Fig. 10 is a combination of data from [1933Gue] and [1996Luo].
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Ni
31
Notes on Materials Properties and Application Knowledge of the phase equilibria in the Ag-Cu-Ni system is important because Ni is widely used as a barrier layer material for the Ag-Cu-Sn alloys, which are the most prominent candidates for Pb-free solders. Ag-Sn-Cu/Ni and Ag-Cu/Ni contacts are frequently used in microelectronic products [1993Che, 2004Che]. The Ag-Cu-Ni system is important for the ceramic packaging industry: Ag-Cu alloys are the most commonly used brazes in ceramic packaging, and nickel pads are usually plated on ceramic substrates in order to improve the brazability between the ceramic and the metals [1996Luo]. Ag-Cu/Ni/Ag-Cu clad tapes for YBCO (YBa2Cu3) superconducting tape were prepared by [2001Yos]. A high Jc value of about 105 A/cm2 was obtained on Ag-0.1%Cu/Ag-10%Ni clad tape. Miscellaneous A number of Ag-Cu-Ni alloys exhibit age-hardening phenomena. For this reason, [1934Pfi1] studied the Cu rich Ag-Cu-Ni alloys and [1934Pfi2] the precipitation in ternary alloys. [1981Duk] and [1981Pav] studied the dendritic segregation of silver-copper alloys with the addition of nickel. Knowledge of phase equilibria of the Ag-Cu-Ni ternary system is also important for the understanding of the interfacial reactions in Ag-Cu/Ni contacts, which are frequently encountered in microelectronic products. [1993Che] studied the interfacial reaction between liquid Ag-Cu eutectic alloys and pure nickel substrates at 793°C. [1996Luo] determined the phases formed at the interfaces of Ni and Ag-Cu alloys at three different temperatures: 860, 795 and 700°C. The diffusion paths proposed by the author are in agreement with the isothermal sections. [1997Zhu] reports experimental results obtained at 850 and 800°C, in terms of kinetics and microstructures of the products formed at the interface of Cu/Ni/Ag-Cu eutectic melts. Using Auger Electron Spectroscopy, [2000Weh] determined the depth distributions of each element in either Ni/Ag-Cu or Ag/Ni-Cu samples. [2004Che] used the isothermal section at 250°C taken from [1977Cha] and [1996Luo] to obtain a representation of the phase equilibria in the Ag-Cu-Ni-Sn quaternary system with the aim of better understanding the reactions occurring at the Ag-Cu-Sn/Ni interfaces. Isoplethal sections at 60, 70, 80 and 90 at.% Sn of the 250°C isothermal section of the Ag-Cu-Ni-Sn quaternary system are represented. References [1913Ces] [1933Gue] [1934Pfi1] [1934Pfi2]
[1967McD] [1977Cha]
[1977Sie] [1981Duk]
[1981Pav]
Landolt-Börnstein New Series IV/11C3
De Cesaris, P., “The Ternary Alloys Ni-Cu-Ag”, Gazz. Chim. Ital., 43, 368-379 (1913) (Experimental, Phase Diagram, 12) Guertler, W., Bergmann, A., “Studies of the Ternary Silver-Copper-Nickel System”, Z. Metallkd., 25, 53-57 (1933) (Experimental, Phase Diagram, Phase Relations, 11) Pfister, H., Wiest, P., “Nickel-Silver-Copper Alloys”, Metallurgist, 9, 154-156 (1934) (Kinetics, Mechan. Prop., 4) Pfister, H., Wiest, P., “Influence of Ni Addition on the Solubility and the Precipitation Course in Ag-Cu System” (in German), Metallwirtschaft, 13, 317-320 (1934) (Experimental, 13) McDonald, A.S., Price, B.R., Sistare, G.H., “Ternary and Higher Order Alloys of Silver”, Silver-Economics, Metallurg., 272-303 (1967) (Phase Diagram, 25) Chang, Y.A., Goldberg, D., Neumann, J.P., “Phase Diagrams and Thermodynamic Properties of Ternary Copper-Silver Systems”, J. Phys. Chem. Ref. Data, 6, 621-673 (1977) (Phase Diagram, Review, 2) Siewert, T.A., Heine, W.R., “Recent Look at Ag-Cu-Ni System”, Metall Trans., 8A, 515-518 (1977) (Experimental, Phase Relations, #, 2) Dukiet-Zawadska, B., Pawlowski, A., Ciah, R., Tasior-Grabianowska, K., Wolczynski, W., “Dendritic Segregation of Ag-Cu Alloys with the Addition of Ti, Al, Mg and Ni”, Arch. Hutn., 26(3), 429-448 (1981) (Experimental, Morphology, 20) Pavlovski, A., “Quantitative analysis of the Dendritic Segregation in Ternary Alloys on the Basis of the Silver-Copper System” (in Russian), Fiz. Met. Metalloved., 52(4), 760-766 (1981) (Experimental, Morphology, 8) MSIT®
Ag–Cu–Ni
32 [1991Mur] [1993Che] [1994Sub]
[1996Luo]
[1997Zhu]
[2000Weh]
[2001Yos]
[2002Leb]
[2002Iva]
[2002Rom]
[2004Che]
Murray, J.L., Alcoa Internal Commun., 1991-04-30 (1991), as quoted in [1993Che] Chen, S.W, “Interfacial Reactions in the Liquid (Ag, Cu)/Ni Diffusion Couples”, Mater. Chem. Phys., 33, 271-276 (1993) (Experimental, Phase Diagram, 11) Subramanian, P.R., “Cu (Copper)”, in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 1-3 (1994) (Crys. Structure, Review, 16) Luo, H.T., Chen, S.W., “Phase Equilibria of the Ternary Ag-Cu-Ni System and the Interfacial Reactions in the Ag-Cu/Ni Couples”, J. Mater. Res., 31, 5059-5067 (1996) (Calculation, Experimental, Phase Diagram, Thermodyn., 29) Zhuang, W.D., Eagar, T.W., “Diffusional Breakdown of Nickel Protective Coatings on Copper Substrate in Silver-Copper Eutectic Melts”, Metall. Mater. Trans. A, 28, 969-977 (1997) (Experimental, Kinetics, Transport Phenomena, 25) Wehr, A., Rylski, A., “The Effect of Microstructure on Interphase Interaction in Layered Polycristalline Films of Silver-Nickel and Silver-Copper-Nickel”, J. Mater. Sci., 35, 777-781 (2000) (Electronic Structure, Experimental, Transport Phenomena, 8) Yoshino, H., Yamazaki, M., Thanh, T.D., Kudo, Y., Kubota, H., “Preparation of Ag-Cu/Ni/Ag-Cu Clad Tapes for YBCO Superconducting Tape and its Textured Properties”, Physica C, 357-360, 923-930 (2001) (Experimental, Mechan. Prop., Supercond., 6) Lebrun, N., “Cu-Ni (Copper-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14832.1.20, (2002) (Phase Diagram, Assessment, Crys. Structure, 51) Ivanchenko, V., “Ag-Ni (Silver-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.29294.1.20, (2002) (Phase Diagram, Assessment, Crys. Structure, 7) Van Rompaey, T., Rogl, P., “Ag-Cu (Silver-Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14511.1.20, (2002) (Phase Diagram, Crys. Structure, Assessment, 28) Chen, S.W., Chang, C.A., “Phase eaquilibria of the Sn-Ag-Cu-Ni Quaternary System at the Sn-rich Corner”, J. Electron. Mater., 33(10), 1071-1079 (2004) (Experimental, Phase Diagram, 25)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
(Ni1–xCux)
cF4 Fm3m Cu
(Ni) < 1455 (Cu) < 1084.62
MSIT®
pure Ag at 25°C [Mas2]
0x1 a = 352.40
pure Ni at 25°C [Mas2]
a = 361.46
pure Cu at 25°C [Mas2] melting point [1994Sub]
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Ni
Ag-Cu
Ag-Cu-Ni
33
A-B-C
Ag-Ni
1435 e1 l' l'' + (Ni) L1+L2+(Ni,Cu) ca. 1250
960 e2 l (Ag) + (Ni) L+(Ag)+(Ni,Cu)
779.1 e3 l (Ag) + (Cu) Fig. 1: Ag-Cu-Ni. Reaction scheme
Ag
Data / Grid: at.%
e2
Fig. 2: Ag-Cu-Ni. Liquidus surface projection
Axes: at.%
e''1
20
80
e3
40
60
0 90
L1+L2 60
40
00 10
80
20
00 11
Landolt-Börnstein New Series IV/11C3
20
1350 40
60
(Ni,Cu)
00 12
Ni
1400
°C 50 12
e'1
13 00
80
Cu
MSIT®
Ag–Cu–Ni
34
Ag
Data / Grid: at.% Axes: at.%
Fig. 3: Ag-Cu-Ni. Isothermal section at 1440°C 20
80
L1
40
60
60
40
80
20
L
L2
Ni
20
(Ni,Cu)
40
60
80
Ag
Cu
Data / Grid: at.% Axes: at.%
Fig. 4: Ag-Cu-Ni. Isothermal section at 1400°C 20
80
L1
40
60
60
40
L1+L2+(Ni,Cu) 80
20
L L2
Ni
MSIT®
(Ni,Cu)
20
40
60
80
Cu
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Ni
35
Ag
Data / Grid: at.% Axes: at.%
Fig. 5: Ag-Cu-Ni. Isothermal section at 1300°C 20
80
L1 40
60
60
40
80
20
L1+L2+ (Ni,Cu)
L L2
20
Ni
40
60
80
Cu
(Ni,Cu)
Ag
Data / Grid: at.% Axes: at.%
Fig. 6: Ag-Cu-Ni. Isothermal section at 1250°C 20
80
40
60
60
40
80
20
L
Ni
Landolt-Börnstein New Series IV/11C3
(Ni,Cu)
20
40
60
80
Cu
MSIT®
Ag–Cu–Ni
36
Ag Fig. 7: Ag-Cu-Ni. Isothermal section at 900°C
Data / Grid: at.% Axes: at.%
(Ag) (Ni,Cu)+(Ag)+L
20
80
L 40
60
60
40
80
20
20
Ni
40
60
80
Cu
(Ni,Cu)
Ag
Data / Grid: at.% Axes: at.%
Fig. 8: Ag-Cu-Ni. Isothermal section at 860°C
(Ag) (Ni,Cu)+(Ag)+L 20
80
L 40
60
60
40
(Ni,Cu)+(Ag)
80
20
(Ni,Cu)+L
(Ni,Cu)
Ni
MSIT®
20
40
60
80
Cu
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Ni
37
Ag
Data / Grid: at.% Axes: at.%
Fig. 9: Ag-Cu-Ni. Isothermal section at 795°C
(Ag)
20
80
(Ni,Cu)+(Ag)+L 40
60
L
60
40
(Ni,Cu)+(Ag)
L+(Ni,Cu)
80
20
(Ni,Cu) 20
Ni
40
60
80
Ag
Data / Grid: at.% Axes: at.%
Fig. 10: Ag-Cu-Ni. Isothermal section at 700°C
(Ag)
20
80
40
60
80
Ni
40
(Ni,Cu)+ (Ag)
60
Landolt-Börnstein New Series IV/11C3
Cu
20
20
40
60
80
(Ni,Cu)
Cu
MSIT®
38
Ag–Cu–P
Silver – Copper – Phosphorus Ortrud Kubashewski, updated by Elena L. Semenova Introduction Ag-Cu-P alloys have been used for joining copper and brasses. These alloys have found favour because of their low melting temperature and the fact that they can be used to braze copper without use of a flux. Compositions used for fillers range from binary Cu-P alloys to those containing up to 15 mass% Ag. Both Ag and P additions decrease the melting range of an alloy, but P has a greater influence. In order to operate within the melting temperature range reliable information is necessary on the Ag-Cu-P phase diagram (first of all on the liquidus and solidus). The first reports on the Ag-Cu-P system [1932Mos, 1956Wei, 1963Bal] were summarized by [1977Cha] followed by the MSIT evaluation by [1988Kub]. [1932Mos] determined that three constituents of the Ag-Cu-Cu3P system have either little or no mutual solubility in the solid. The Ag-Cu3P section was shown to be quasibinary. Other vertical sections as well as the liquidus surface of Ag-Cu-Cu3P system were constructed. Two vertical sections passing through or in the vicinity of the alloys used in soldering (BCuP-3, BCuP-4) were constructed by [1956Wei]. [1963Bal] conducted a study to ascertain the melting ranges of the ternary alloys in the Cu rich region. Flow characteristics were deduced from the data. High purity materials, such as oxygen free high conductivity copper and 99.99% silver, were used by [1987Tak1, 1987Tak2] for investigation of the Ag-Cu-Cu3P system to obtain the liquidus surface and several vertical sections. One of the vertical sections constructed by [1956Wei] has been corrected here. Crystal structure studies by [1965Olo] of ternary alloys with compositions (Cu0.67Ag0.33)P2, (Cu0.50Ag0.50)P2 and (Cu0.33Ag0.67)P2 that had been prepared by sintering mixtures of the binary phosphides at 600°C for 7 days indicated complete mutual solid solubility between CuP2 and AgP2. [1991Zen] studied alloys in the Cu rich part of the system with P contents up to 38 at.%. Alloys were prepared from the components having purities of 99.99, 99.9 and 99.3% for Ag, Cu and P respectively. Some solubility of the components was found. The isothermal section at 500°C was constructed. A number of reviews of the Ag-Cu-Cu3P system have been published: [1933Fro, 1949Jae, 1967Mcd, 1973Sis, 1979Dri] citing [1932Mos]. Brief details of the experimental works on the Ag-Cu-P system are listed in Table 1. Binary Systems Assessments of the Ag-Cu system by [2002Rom], of the Ag-P system by [2006Per] and of the Cu-P system by [2002Per] are accepted. The Ag-Cu and Cu-P systems as well as the Ag-Cu3P quasibinary system bounding the Ag-Cu-Cu3P region [1932Mos] are of the eutectic type. Solid Phases Crystallographic data of solid phases are given in Table 2. Quasibinary Systems The quasibinary character of the Ag-Cu3P section (Fig. 1) was established by [1932Mos]. Invariant Equilibria The data relating to the invariant equilibria in the Ag-Cu-Cu3P system are given in Table 3. The eutectic temperature of 650°C is accepted from [1991Zen]. The corresponding reaction scheme is shown in Fig. 2.
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–P
39
Liquidus Surface The liquidus surface of the Ag-Cu-Cu3P system is shown in Fig. 3, taken from [1932Mos]. Small variations in the phosphorus content affect the liquidus temperature to a great extent in the high phosphorus region. Isothermal Sections The isothermal section of the Ag-Cu-Cu3P system at 500°C is given in Fig. 4, taken from [1991Zen]. Temperature – Composition Sections Two sections of the Ag-Cu-Cu3P system, through 6 mass% P and 2 mass% Ag, respectively, are shown in Figs. 5 and 6. However, [1991Zen] ignored the mutual solubility of the components that was previously indicated by them in the isothermal section at 500°C. Notes on Materials Properties and Applications Due to the possible application of Ag-Cu-P alloys as soldering materials, the investigation of their electro-physical, electro-chemical and mechanical properties has been carried out in parallel with the phase diagram studies. The effect of phosphorus additions on physical and chemical properties of the Ag-Cu-P alloys was investigated by [1932Fro, 1945Coo, 1956Wei, 1963Bal, 1981Ple]. According to the diagram of shear strength for the Ag-Cu-P alloys, the best characteristics were observed for compositions with 15Ag-Cu-(30-40)Cu3P (mass%). The diagram of ductility suggested that the most conductive compositions were those rich in Ag (~85 mass% Ag) and those with 40-60 mass% Ag and 7-15 mass% Cu3P [1956Wei]. The14.5Ag-Cu-7P alloy was chosen as the most ductile by [1981Ple]. It was noted that being less strong and ductile 2-6 mass% Ag alloys could be applied for somewhat lower brazing temperatures. The alloys of the Ag-Cu-P system richest in Ag were investigated by [1932Fro]. Alloys containing 83.5 mass% Ag and between 0.3-1 mass% P, quenched from 520 and 600°C showed an appreciable combination of the tensile strength and elongation. Maximum elongation was seen for alloys with 50 mass% Ag and 1 mass% P, whereas tensile strength was found to increase with increasing P content. Minimum Brinell hardness was found for alloys containing 83.5 mass% Ag and quenched from 520, 600 and 700°C when the P content was 0.3 mass%. For the alloys with 50 mass% Ag the P content giving minimum hardness was about 1 mass% P. The deep extension of these alloys did not depend much on the phosphorous content (0-3 mass%). Corrosion resistance of the alloys with 80 mass% Ag and 0-1.0 mass% P was increased in acetic acid and decreased in nitric acid with increasing phosphorus content. [1945Coo] studied the effect of small additions of P (< 0.2 mass%) on physical and mechanical properties of the 1 mass% Ag-Cu alloy in the wrought and deposited states. Phosphorous was shown not to have an appreciable effect upon the melting range of the ternary alloys as well as upon density of the materials in the wrought condition. The increase in the P content from 0.025 to 0.19 mass% in the deposited state and in the welds resulted in a decrease in the density. The tensile strength of the wrought alloys was observed to increase with phosphorus addition, while it did not notably affect the ductility. The strength of the alloys in the deposited state was reduced when the phosphorous content exceeded about 0.1 mass% and the alloys with higher P content were exposed to hot shortness. On decreasing the Ag content (from 14.5 to 2 mass%) with simultaneous increasing of P (from 4.5 to 6 mass%) the ductility of the ternary alloys fell, the yield stress and breaking point decreased while electrical resistance increased [1981Ple]. There was observed no correlation between flow characteristics with changing Ag content (0-15 mass%), whereas a higher P content (5-7.3 mass%) resulted in higher fluidity. The largest spread was at 10 Ag and 7.3 P (mass%) [1963Bal].
Landolt-Börnstein New Series IV/11C3
MSIT®
40
Ag–Cu–P
From the work of [1987Tak1, 1987Tak2] it follows that the spread area of the Ag-Cu-7P filler metals on copper plate does not depend on Ag content (10 or 15 mass%) at a brazing temperature 700°C. It reaches the same size at 670°C for an alloy with 15% Ag. The best result at 700°C was obtained for a 15Ag-Cu-8.5P alloy. The capillarity rise did not deteriorate with phosphorus in the presence of Ag, but without Ag all characteristics of wettability deteriorated [1981Ple]. Miscellaneous The eutectic alloy, having the lowest temperature of melting, was found to be useless for practical purposes owing to its extreme brittleness [1956Wei]. References [1932Fro]
[1932Mos] [1933Fro] [1945Coo]
[1949Jae] [1956Wei] [1963Bal] [1965Olo]
[1967Mcd]
[1973Sis]
[1977Cha]
[1979Dri]
[1981Ple]
[1987Tak1]
MSIT®
Fröhlich, K.W., “About Behavior of Phosphorus with Cooper and Silver” (in German), Mitt. Forsch. Inst. Edelmetalle, 6, 69-74 (1932) (Experimental, Phase Diagram, Phase Relations, 1) Moser, H., Fröhlich, K.W., Raub, E., “Ternary Silver-Copper-Phosphous System” (in German), Z. Anorg. Chem., 208, 225-237 (1932) (Phase Diagram, Phase Relations, 6) Fröhlich, K.W., “Phosphorus in Silver-Copper Alloys - II.” (in German), Mitt. Forsch. Inst. Edelmetalle, 7, 91-96 (1933) (Experimental, Mechan. Prop., 3) Cook, H., Davis, E., “Effect of Phosphorus on the Properties and Welding Characteristics of Arsenical and Non- Arsenical Copper and on Copper-Silver Alloy Filler Rod”, Trans. Inst. Neld., 8(1), 19-29 (1945) (Experimental, Mechan. Prop., 3) Jänecke, E., “Ag-Cu-P”, Kurzgefasstes Handbuch Aller Legierungen, Winter, Heidelberg, 461 (1949) (Phase Diagram, Review, 1) Weigert, K.M., “Physical Properties of Commercial Silver-Copper-Phosphorus Brazing Alloys”, Welding J., 35, 672 (1956) (Phase Diagram, Phase Relations, Experimental, 6) Ballentine, R.E., “Selecting Phosphorus Brazing Filter Metals”, Machinery Lloyd (Overseas Edn.), 69, 138 (1963) (Phase Diagram, Experimental, 0) Olofsson, O., “The Crystal Structure of CuP2 and AgP2 with Some Phase- Analytical Data of the Cu-P and Ag-P Systems”, Acta Chem. Scand., 19, 229-241 (1965) (Crys. Structure, Phase Relations, Experimental, 15) McDonald, A.S., Price, B.R., Sistare, G.H., “Ternary and Higher-Order Alloys of Silver”, Silver-Economics, Metallurgy and Use, Butts, A. (Ed.), Princeton, NJ: Van Nostrand, 272-303 (1967) (Phase Diagram, Review, 1) Sistare, G.H., “Ag-Cu-Cu3P (Silver-Copper-Copper Phosphide)”, Metals Handbook, ASM International, Metals Park, Ohio, Vol. 8, 379 (1973) (Phase Diagram, Phase Relations, Assessment, 2) Chang, Y.A., Goldberg, D., Neumann, J.P., “Phase Diagrams and Thermodynamic Properties of Ternary Cu-Ag Systems”, J. Phys. Chem. Ref. Data, 6, 621-673 (1977) (Phase Diagram, Review, 3) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Cu-Ag-P” (in Russian), in “Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 205-206 (1979) (Phase Diagram, Phase Relations, Review, 1) Pleter, O., “Technologial Possibilities for the Decreasing Silver Content in Soldering Alloys”, Metalurgia (Bukarest), 33(7), 379-390 (1981) (Experimental, Mechan. Prop., Phys. Prop., 8) Takemoto, T., Okamoto, I., Matsumura, J., “Phase Diagrams of Copper-Silver-Phosphorus and Copper-Tin- Phosphorus Ternary Brazing Filler Metals-Copper Phosphorus Brazing Filler Metals with Low Melting Temperature (II)”, Trans. JWRI, 16(2), 301-307 (1987) (Phase Diagram, Experimental, 8)
Landolt-Börnstein New Series IV/11C3
Ag–Cu–P [1987Tak2]
[1988Kub]
[1991Zen]
[2002Per]
[2002Rom]
[2006Per]
41
Takemoto, T., Okamoto, I., Matsumura, J., “Phase Diagrams of Copper- Silver-Phosphorus and Copper - Tin - Phosphorus Ternary Filler Metals Copper Phosphorus Brazing Filler Metals with Low Melting Temperatures (Report II)” (in Japanese), Yosetsu Gakkai Ronbunshu, 5(1), 81-86 (1987) (Phase Diagram, Experimental, Phase Relations, Crys Structure, 13) Kubashewski, O., “Silver-Copper-Phosphorus”, MSIT Ternary Evaluation Program in MSIT Workplace, Effenberg G. (Ed.), MSI, Material Science International Services GmBH, Stuttgart, Document ID: 10.16994.1.20 (1988) (Crys Structure, Phase Diagram, Phase Relations, Assessment, 5) Zeng, L., Zhuang, Y., Li, D., Lin, W., “A 500°C Isothermal Section of the Ag-Cu-P System”, Acta Metall. Sin. (China), 27(2), B140-B142 (1991) (Phase Diagram, Phase Relations, Experimental, 4) Perrot, P., Batista, S., Xing, X., “Cu-P (Copper - Phosphorus)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.16300.1.20, (2002) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 7) van Rompaey, T., Rogl, P., “Ag-Cu (Silver - Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14511.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, Phase Relations, 29) Perrot, P., “Ag-P (Silver - Phosphorus)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2006) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 5)
Table 1: Investigation of the Ag-Cu-P Phase Relations, Structures and Thermodynamics Reference
Experimental Technique
[1932Mos]
Thermal analysis, metallography 600 - 1100°C, in Ag-Cu-Cu3P region
[1956Wei]
Thermal analysis, metallography 600 - 1100°C, in Ag-Cu-Cu3P region
[1963Bal]
Determination of liquidus surface, 700 - 1100°C, in copper rich area; up to about establishing of melting ranges 10 mass% (4.7 at.%) P and 20 mass% (30.7 at.%) Ag
[1965Olo]
X-ray diffraction
[1987Tak1, 1987Tak2]
DTA, optical microscopy, EPMA, 650 - 900°C, Cu - 15 mass% (23 at.%)Ag X-ray analyses 8 mass% (3.4 at.%) P
[1991Zen]
X-ray diffraction, DTA, metallography
Landolt-Börnstein New Series IV/11C3
Temperature/Composition/Phase Range Studied
AgP2-CuP2 section
500-1083°C, in Ag-Cu-Cu3P region
MSIT®
Ag–Cu–P
42 Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comment/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
at 25°C, pure Ag
a = 403.95
at 12.1 at.% Cu
a = 402.1
at 12.6 at.% Cu, at 7 GPa, at 890°C, dissolves 14 at.% Cu [2002Rom]; liquid Ag dissolves 9.7 at.% P at 880°C [2006Per]; at 500°C dissolves 1.02 at.% Cu [1991Zen]
(Cu) < 1084.62
cF4 Fm3m Cu
a = 361.46
pure Cu at 19°C at 500°C dissolves 2.6 at.% Ag and 0.54 at.% P [1991Zen]
a = 363.7
at 3.7 at.% Ag
a = 366.3
at 4 at.% Ag, 7 GPa, 890°C [2002Rom]
(P) (I)
cP1 Pm3m Po
a = 237.7 a = 229
high pressure phase, above ~10 GPa [2006Per]
(P) (II)
hR6 R3m As
a = 337.7 c = 880.6
high pressure phase 5 to 11.1 GPa [2006Per]
(P) (red)
c*66
a = 1131
subl. at 417C, 1 bar [2006Per]
(P) (white) < 44.14
c* ? P (white)
a = 718
at 25°C, common form of elemental P, less stable than P (red) at this temperature [2006Per]
(P) (black)
oC8 Cmca P (black)
a = 331.6 b = 1047.8 c = 437.63
at 25°C [2006Per]
Ag3P11 458
m* Cm
a = 1299.9 b = 755.5 c = 661.2 = 118.84
[2006Per]
Cu3P < 1022
hP8 P3m1 Cu3P
a = 409.2 c = 718.6
at 560°C, congruent melting at 1 bar, 25 to 31 at.% P, at ~833°C [2002Per]
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–P Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Cu3P
hP24 P63cm Cu3P
(AgxCu1–x)P2
43
Lattice Parameters Comment/References [pm] low temperature phase [2002Per]
a = 695.93 c = 714.3
0 x 100 [1965Olo]
mP12 P21/c CuP2
AgP2 554
a = 621.8 b = 505.6 c = 780.4 = 113.48
[2006Per]
CuP2 < 891
a = 580.04 b = 480.63 c = 752.63 = 112.70
congruent melting at 15.2 bar [2002Per]
Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition, at.% Ag
Cu
P
L (Ag) + Cu3P
796
e1(max)
-
-
-
-
L (Ag) + (Cu) + Cu3P
650
E
L
10.6
74.6
14.8
Landolt-Börnstein New Series IV/11C3
MSIT®
Ag–Cu–P
44
1100
Fig. 1: Ag-Cu-P. The quasibinary system Ag - Cu3P
1022°C 1000
L
Temperature, °C
961.78°C
900
L+(Ag)
L+(Cu3P)
800
700
(Cu3P)+(Ag) 600
20
Ag
40
60
Cu, at.%
Ag-Cu
Ag-Cu-P
A-B-C
796 e1(max) l (Ag) + Cu3P
Ag 0.00 Cu 75.00 P 25.00
Cu-P
714 e3 l (Cu) + Cu3P
779.1 e2 l (Ag) + (Cu)
650
L (Ag) + (Cu) + Cu3P
E
(Ag) + (Cu) + Cu3P
Fig. 2: Ag-Cu-P. Partial reaction scheme in the Ag - Cu - Cu3P region
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–P
45
P
Data / Grid: at.% Axes: at.%
Fig. 3: Ag-Cu-P. Partial liquidus surface projection 20
80
40
60
60
40
80
e1(max)
700
800 (Ag)
900 20
Ag
Cu 3P
900 800 E 800
(Cu)
e2 40
60
900
20
e3
°C 00 10
80
P
Cu
Data / Grid: at.% Axes: at.%
Fig. 4: Ag-Cu-P. Partial isothermal section at 500°C 20
80
40
60
60
40
Cu3P
80
20
(Ag)+(Cu)+Cu3P
Ag
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Cu
MSIT®
Ag–Cu–P
46
Fig. 5: Ag-Cu-P. Isopleth through 6 mass% P, plotted in at.%
Temperature, °C
L
750
L+(Cu) L+(Ag) L+(Cu)+(Ag) L+(Cu)+Cu3P L+(Ag)+Cu3P (Cu)+(Ag)+Cu3P
500
Ag 34.34 Cu 51.30 P 14.35
Fig. 6: Ag-Cu-P. Isopleth through 2 mass% Ag, plotted in at.%
60
70
Ag 0.00 Cu 88.42 P 11.58
80
Cu, at.%
1000
Temperature, °C
L
L+(Cu)
L+(Cu)+(Ag)
L+Cu3P
750
L+(Ag)+Cu3P L+(Cu)+Cu3P (Cu)+(Ag)+Cu3P
Ag 1.19 Cu 98.81 0.00 P
MSIT®
10
20
P, at.%
Ag 1.03 Cu 73.75 P 25.22
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Sn
47
Silver – Copper – Tin K.C. Hari Kumar, Ortrud Kubaschewski Introduction The phase diagram of the Ag-Cu-Sn system is fundamental to the development of technology of lead-free solder alloys. Recent legislations in many countries to eliminate the use of lead containing solder alloys have created renewed interest in this system, especially in the Sn rich region. Apart from this, Ag-Cu-Sn also forms the basis of several Cu rich dental amalgams. Constitutional data of the Ag-Cu-Sn system were reviewed by several authors [1967Mcd, 1977Cha, 1977Dar, 1979Dri, 1982Fed, 1988Kub, 2002Lew]. Present report updates an earlier evaluation by [1988Kub] made within the MSIT Evaluation Program. First comprehensive experimental work on the phase equilibria was reported by [1959Geb], employing thermal, metallographic and X-ray analyses. That investigation mainly concentrated in the Ag-Cu region and reported a liquidus projection, isothermal sections at 500 and 600°C, and vertical sections at 5, 10, 15, 20 and 25 mass% Sn. The Sn rich region was investigated by [1959Wag] (thermal analysis) and by [1954Cro] (metallography), results of these studies were incorporated in [1959Geb]. Using X-ray and metallographic techniques [1956Gla] claimed that the intermetallic systems Ag3Sn()-Cu3Sn(J1) and Ag5Sn(J2)-Cu3Sn(J1) show only limited solid solubility. Phase relation in the system at 37°C was reported by [1972Jen, 1973Wir, 1980Kra] independently, that was used to explain characteristics of certain Cu rich dental amalgams. [1994Mil] was the first to report that Sn rich invariant reaction, observed in alloys close to typical solder compositions currently in use, is of eutectic type instead of transition type as stated in the work of [1959Geb]. This was later confirmed by several investigators: [1997And], [2000Loo] (DSC), [2000Moo, 2004Moo] (DTA), and [2002Lew] (SEM, EPMA, LOM). Using DTA [1997He] determined liquidus and solidus of seven Cu rich alloys, typically used as brazing filler. [1997He] also reported that all alloys in the composition range studied (15-25 mass% Ag, 15-24 mass% Sn) contained the phase (Cu41Sn11), that decreased ductility. Using SEM-EDS and optical microscopy [2000Ohn] established phase relation in nine samples in the temperature range 300-500°C having compositions 20-48 mass% Ag and 15-60 mass% Cu. Isopleths of the ternary Ag-Cu-Sn system were published by two teams, [1959Geb] and [1982Fed]. The vertical sections at constant Sn content, namely 5, 10, 15, 20 and 25 mass% are based on thermal, microscopic and X-ray studies by [1959Geb]. Three isopleths were constructed by [1982Fed]: one parallel to the Ag-Sn system containing 10 mass% Cu, one cross section from Ag to 80Sn-17.8Cu and one from Cu to 50Ag-50Sn (mass%). Apart from experimental investigations, several authors have reported calculated phase equilibria in the system [1997Lee, 2000Moo, 2000Ohn, 2003Kim, 2003Ohn], employing the CALPHAD method, however, only [2000Moo] reported ternary parameters used for the calculation. The phase diagrams presented in the present evaluation are calculated using the Calphad approach (see section “Thermodynamics”). [2004Lue] determined isoenthalpy curves for liquid alloys at 500, 700, 900°C using a Calvet type calorimeter. Physical properties of typical solder compositions were also studied by several authors: notably density [2004Gas], viscosity [2001Lee, 2002Wan, 2003Ohn], surface tension [2001Lee, 2002Mos, 2003Ohn, 2004Gas], and wetting properties [2004Gas]. Electrical properties were studied by [2003Coo, 2004Kis]. Coarsening kinetics of intermetallics of the solder alloys were determined by [2004All]. Binary Systems The binary system Ag-Cu [2002Rom] is from the MSIT Binary Evaluation Program. Ag-Sn is from [1999Oht] and Cu-Sn is from [2000Ohn]. The version of the Cu-Sn phase diagram used here differs from that of [Mas2] with respect to the phase. It has been confirmed by the careful investigations by [2000Ohn, 2004Liu] that there is no eutectoid reaction (Cu) + in the Cu-Sn system, instead the phase boundary between and is of higher-order type involving a two-stage chemical ordering reaction (A2) ’ (B2) (D03). In the evaluation of the ternary system, the order-disorder equilibria are neglected since the ternary extensions of them are undetermined. Landolt-Börnstein New Series IV/11C3
MSIT®
48
Ag–Cu–Sn
Solid Phases No ternary compounds have been reported. The solid phases of the Ag-Cu-Sn ternary systems are listed in Table 1. Invariant Equilibria The invariant equilibria associated with the liquid phase are given in Table 2. All invariant reaction temperatures and compositions are calculated using the Calphad approach. They are in reasonable agreement with experimental data. Figure 1 depicts the reaction scheme for the system relevant to reactions involving the liquid phase. Liquidus Surface The calculated liquidus surface is shown in Fig. 2. It should be noted that the calculated phase boundaries along the Ag-Cu side deviate slightly from the accepted version of the Ag-Cu binary system [2002Rom]. Wherever necessary the phase boundaries are adjusted such that it matches with those given in [2002Rom]. The invariant reaction close to the Sn corner occurring at 217.7°C is of eutectic type (E1). The corresponding eutectic composition is 3.24Ag-0.57Cu-Sn (mass%). These are very close to the very precisely experimentally determined eutectic temperature (217.2°C) and composition (3.58Ag-0.96Cu-Sn) [2004Moo]. Figure 3 is the enlarged Sn rich corner of the liquidus projection in detail. Isothermal Sections Figures 4 to 7 show isothermal sections at 600, 500, 400, and 37°C, respectively, calculated using the Calphad method. Here also the phase boundaries along the Ag-Cu side have been adjusted to match with the accepted binary system. It is assumed that except and J2, none of the intermediate phases extends into the ternary region. The calculated diagrams compare well with available experimental data [1959Geb, 1980Kra, 2000Ohn]. Temperature – Composition Sections Isopleths at 3.66 mass% Ag, 10 mass% Sn, and 25 mass% Sn are shown in Figs. 8, 9 and 10, respectively. These diagrams reproduce fairly well the experimental findings of [1959Geb] and [1982Fed]. Thermodynamics The phase diagrams presented in the present evaluation are calculated using the Calphad approach. The binary Gibbs energy model parameters are from [1986Hay] (Ag-Cu), [1999Oht] (Ag-Sn), and [2001Liu] (Cu-Sn). The ternary parameters are obtained in the present work and they are listed in Table 3. Calorimetrically measured [2004Lue] isoenthalpy curves for liquid alloys at 500 and 700°C are shown in Figs. 11 and 12, respectively. The data indicate temperature dependence for the enthalpy of mixing. Notes on Materials Properties and Applications Ag-Cu-Sn alloys form the basis of lead-free solders and dental amalgams. Knowledge of the phase relationships in the ternary and quaternary (Ag-Cu-Sn-Hg) systems is essential for the understanding and development of new dental amalgams. Ag3Sn() is known as “dental alloy”. If mixed with an equal amount of Hg, the undesirable phase Sn7.8Hg is formed during the hardening process. The addition of Cu to the basic Ag-Sn alloys suppresses the formation of this phase [1980Kra]. Renewed interest in Ag-Cu-Sn system is particularly due to several Sn rich lead-free solder compositions close to the eutectic E1. Physical properties such as surface tension and viscosity of the liquid phase are important quantities for quantifying wetting properties of solders. Several such studies are mentioned in the section “Introduction”. In Figs. 13 and 14, calculated surface tension contours for the liquid are shown at
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Sn
49
800 and 1000°C, respectively [2000Ohn, 2003Ohn]. Figures 15 and 16 show isoviscosity contours for the liquid at 800 and 1000°C [2000Ohn, 2003Ohn], respectively. References [1954Cro] [1956Gla]
[1959Geb] [1959Wag] [1967Mcd] [1972Jen] [1973Wir]
[1977Cha]
[1977Dar] [1979Dri]
[1980Kra]
[1982Fed]
[1986Hay]
[1987Kar]
[1988Kub]
[1994Mil] [1997And] [1997He]
Landolt-Börnstein New Series IV/11C3
Crowell, W.S., “The Metallography of Dental Amalgam Alloys”, J. Dent. Res., 33, 592-595 (1954) as quoted in [1959Geb] Galdyshevsky, E.I., Cherkashin, E.E., “Solid Solutions Based on Metallic Compounds” (in Russian), Zh. Neorg. Khim., 1, 1394-1401 (1956) (Crys. Structure, Phase Relations, Experimental, 4) Gebhardt, E., Petzow, G., "The Constitution of the System Ag-Cu-Sn” (in German), Z. Metallkd., 50, 597-605 (1959), (Experimental, Phase Diagram, #, *, 15) Wagner, E., Degussa Pforzheim, (personal communication to [1959Geb]) McDonald, A.S., Price, B.R., Sistare, G.H., “Ternary and Higher-Order Alloys of Silver”, Silver-Economics, Metallurgy and Use, 272-303 (1967) (Review, Phase Relations, 38) Jensen, S.J., “Cu-Sn Phases in Dental Silver Amalgam Alloy”, Scand. J. Dent. Res., 80, 158-161 (1972) (Experimental, Phase Diagram, #, 9) Wirjasumarto, H., “Physical Metallurgy of Ag-Sn Dental Amalgam”, Thesis, Univ. Kentucky, USA, (1973) (Experimental, Crys. Structure, Phase Relations, Phase Diagram, 50) Chang, Y.A., Goldberg, D., Neumann, J.P., “Phase Diagrams and Thermodynamic Properties of Ternary Cu-Ag Systems”, J. Phys. Chem. Ref. Data, 6, 621-673 (1977) (Review, Phase Diagram, #, 6) Darvell, B.W., “Some Studies on Dental Amalgam. Pt. 3. The Constitution of Amalgam and Alloy”, Surf. Technol., 5, 487-499 (1977) (Review, Phase Relations, #, 6) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Cu-Sn-Ag” (in Russian), in “Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 189-190 (1979) (Assessment, Phase Diagram, Phase Relations, 1) Kraft, W., Petzow, G., Aldinger, F., “Constitution of Copper-Rich Dental Amalgams” (in German), Z. Metallkd., 71, 699-703 (1980) (Experimental, Phase Relations, Phase Diagram, #, 24) Fedorov, V.N., Osintsev, O.E., Yushkina, E.T., “Phase Diagrams of Metallic Systems”, Ageev, N.V., Petrova, L.A. (Eds.), Vol. 26, Acad. Sci. USSR, Moscow, 149-150 (1964-1982) (Review, Phase Diagram, Phase Relations, #) Hayes, F.H., Lukas, H.L., Effenberg, G., Petzow, G., “Thermodynamic Optimisation of the Cu-Ag-Pb System”, Z. Metallkd., 77, 749-754 (1986) (Theory, Phase Diagram, Thermodyn., 35) Karakaya, I., Thompson, W.T., “The Ag-Sn (Silver-Tin) System”, Bull. Alloy Phase Diagrams, 8, 340-347 (1987) (Review, Phase Diagram, Phase Relations, Thermodyn., Crys. Structure, #, 51) Kubaschewski, O., “Silver - Copper - Tin”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16022.1.20, (1988) (Crys. Structure, Phase Diagram, Assessment, 11) Miller, C.M., Anderson, I.E., Smith, J.F., J. Electron. Mater., 23, 595-601 (1994) (Experimental) as quoted in [2000Loo] Anderson, I.E., Terpstra, R.L., Unal, O., “Reliability of Solders and Solder Joints”, Symposium, TMS Annual Meeting, Orlando, FL, (1997) as quoted in [2000Loo] He, Z., Ding, L., “Investigation on Ag-Cu-Sn Brazing Filler Metals”, Mater. Chem. Phys., 49, 1-6 (1997) (Experimental, 5)
MSIT®
50 [1997Lee]
[1999Oht]
[2000Loo]
[2000Moo]
[2000Ohn]
[2001Lee]
[2001Liu]
[2002Lew]
[2002Mos]
[2002Rom]
[2002Wan] [2003Coo]
[2003Kim]
[2003Ohn]
[2004All]
MSIT®
Ag–Cu–Sn Lee, B.-J., Hwang, N.M., Lee, H.M., “Prediction of Interface Reaction Products Between Cu and Various Solder Alloys by Thermodynamic Calculation”, Acta Mater., 45, 1867-1874 (1997) (Theory, Phase Diagram, Thermodyn., #, 26) Ohtani, H., Miyashita, M., Ishida, K., “Thermodynamic Study of Phase Equilibria in the Sn-Ag-Zn System” (in Japanese), J. Jpn. Inst. Met., 63, 685-694 (1999) (Theory, Phase Diagram, Thermodyn., Phase Relations, 57) Loomans, M.E., Fine, M.E., “Tin-Silver-Copper Eutectic Temperature and Composition”, Metall. Mater. Trans., 31A, 1155-1162 (2000) (Experimental, Phase Relations, Phase Diagram, 6) Moon, K.-W., Boettinger, W.J., Kattner, U.R., Biancaniello, F.S., Handwerker, C.A., “Experimental and Thermodynamic Assessment of Sn-Ag-Cu Solder Alloys”, J. Electron. Mater., 29, 1122-1136 (2000) (Experimental, Theory, Thermodyn., Phase Diagram, Phase Relations, #, 24) Ohnuma, I., Miyashita, M., Anzai, K., Liu, X.J., Ohtani, H., Kainuma, R., Ishida, K., “Phase Equilibria and the Related Properties of Sn-Ag-Cu Based Pb-Free Solder Alloys”, J. Electron. Mater., 29, 1137-1144 (2000) (Experimental, Theory, Phase Relations, Thermodyn., #, *, 22) Lee, J.H., Lee, D.N., “Use of Thermodynamic Data to Calculate Surface Tension and Viscosity of Sn-Based Soldering Alloy Systems”, J. Electron. Mater., 30, 1112-1119 (2001) (Theory, Phase Relations, #, 26) Liu, X.J., Lu, H.S., Ohnuma, I., Kainuma, R., Ishida, K, Itabashi, S., Kameda, K., Yamaguchi, K., “Experimental Determination and Thermodynamic Calculation of the Phase Equilibria in the Cu-In-Sn System”, J. Electron. Mater., 30, 1093-1099 (2001) (Experimental, Theory, Phase Diagram, Phase Relations, Thermodyn., 19) Lewis, D., Allen, S., Notis, M., Scotch, A., “Determination of the Eutectic Structure in the Ag-Cu-Sn System”, J. Electron. Mater., 31, 161-167 (2002) (Experimental, Phase Diagram, Phase Relations, Review, 13) Moser, Z., Gasior, W., Pstrus, J., Ksiezarek, S., “Surface-Tension Measurements of the Eutectic Alloy (Ag-Sn 96.2 at.%) with Cu Additions”, J. Electron. Mater., 31, 1225-1229 (2002) (Experimental, Theory, 22) Van Rompaey, T., Rogl, P., “Ag-Cu (Silver - Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14511.1.20, (2002) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 29) Wang, X., Bao, H., Li, W., “Estimation of Viscosity of Ternary-Metallic Melts”, Metall. Trans. A, 33, 3201-3204 (2002) (Theory, 18) Cook, B.A., Anderson, I.E., Harringa, J.L., Kang, S.K., “Isothermal Aging of Near-Eutectic Sn-Ag-Cu Solder Alloys and its Effect on Electrical Resistivity”, J. Electron. Mater., 32, 1384-1390 (2003) (Experimental, Electr. Prop., 11) Kim, K.S., Huh, S.H., Suganuma, K., “Effects of Intermetallic Compounds on Properties of Sn-Ag-Cu Lead-free Soldered Joints”, J. Alloys Compd., 352, 226-236 (2003) (Experimental, Theory, Phase Relations, 23) Ohnuma, I., Miyashita, M., Liu, X.J., Ohtani, H., Ishida, K., “Phase Equilibria and Thermodynamic Properties of Sn-Ag Based Pb-Free Solder Alloys”, IEEE Trans., Electron. Pack. Manuf., 26, 84-89 (2003) (Theory, Phase Diagram, Phase Relations, Thermodyn., #, *, 21) Allen, S.L., Notis, M.R., Chromik, R.R., Vinci, R.P., Lewis, D.J., Schaefer, R., “Microstructural Evolution in Lead-free Solder Alloys: Part II. Directionally Solidified Sn-Ag-Cu, Sn-Cu and Sn-Ag”, J. Mater. Res., 19, 1425-1431 (2004) (Experimental, Morphology, 24)
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Ag–Cu–Sn [2004Gas]
[2004Kis]
[2004Liu]
[2004Lue]
[2004Moo]
51
Gasior, W., Moser, Z., Pstrus, J., Bukat, K., Kisiel, R., Sitek, J., “(Sn-Ag)eut + Cu Soldering Materials, Part I: Wettability Studies”, J. Phase Equilib. Diffus., 25, 115-121 (2004) (Experimental, 17) Kisiel, R., Gasior, W., Moser, Z., Bukat, K., Sitek, J., “(Sn-Ag)eut + Cu Soldering Materials, Part II: Electrical and Mechanical Studies”, J. Phase Equilib. Diffus., 25, 122-124 (2004) (Experimental, 5) Liu, X.J., Wang, C.P., Ohnuma, I., Kainuma, R., Ishida, K., “Experimental Investigation and Thermodynamic Calculation of the Phase Equilibria in the Cu-Sn and Cu-Sn-Mn Systems”, Metall. Mater. Trans., 35A, 1641-1654 (2004) (Theory, Experimental, Phase Diagram, Phase Relations, Thermodyn., 55) Luef, C., Flandorfer, H., Ipser, H., “Lead-Free Solder Materials: Experimental Enthalpies of Mixing in the Ag-Cu-Sn and Cu-Ni-Sn Ternary Systems”, Z. Metallkd., 95, 151-163 (2004) (Experimental, Theory, Thermodyn., 11) Moon, K.-W., Boettinger, W.J., “Accurately Determining Eutectic Compositions: The Sn-Ag-Cu Ternary Eutectic”, J. Metals, 56, 22-27 (2004) (Experimental, Phase Relations, Phase Diagram, #, 18)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57 a = 414.7
pure, at 25°C [Mas2] 14 at.% Sn [V-C]
(Cu) < 1084.62
cF4 Fm3m Cu
a = 361.46 a = 370.46
pure, at 25°C [Mas2] 9.1 at.% Sn, [Mas2]
(Sn) 231.9681 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
pure, at 25°C [Mas2]
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
pure, [Mas2]
, Cu13Sn3 796 - 586
cI2 Im3m W
a = 297.81
13.4 at.% Sn [Mas2] 12.6 - 17 at.% Sn [2004Liu]
´ 775 - 574
cP2 Pm3m CsCl
, Cu4Sn 722 - 515
cF16 Fm3m BiF3
a = 606.05
710°C, 16.6 at.% Sn [Mas2] 13.7 - 27.5 at.% Sn [2004Liu]
, Cu41Sn11 569 - 350
cF416 F43m Cu41Sn11
a = 1798
20.5 at.% Sn [Mas2]
Landolt-Börnstein New Series IV/11C3
13.7 - 22.8 at.% Sn [2004Liu]
MSIT®
Ag–Cu–Sn
52 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, Cu10Sn3 635 - 540
hP26 P63 Cu10Sn3
a = 733.0 c = 786.4
23.1 at.% Sn [Mas2]
, Cu6Sn5 (HT) 415 - 187
hP4 P63/mmc NiAs
a = 419.0 c = 508.6
45.45 at.% Sn [Mas2]
´, Cu6Sn5 (LT) < 187
hR22
a = 2087.0 c = 2508.1
45.45 at.% Sn [Mas2] Hexagonal superlattice based on NiAs
J1, Cu3Sn < 666
oC80 Cmcm Cu3Sn
a = 552.9 b = 4775.0 c = 432.3
25 at.% Sn [Mas2]
J2, Ag5Sn < 722
hP2 P63/mmc Mg
a = 294.36 c = 478.45
11.8-22.85 at.% Sn [1987Kar]
, Ag3Sn < 477
oP8 Pmmn Cu3Ti
a = 596.82 b = 478.02
23.7-25 at.% Sn [1987Kar]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Ag
Cu
Sn
L + (Cu) (Ag) +
586
U1
L (Cu) (Ag)
37.20 1.10 90.00 2.84
47.06 91.57 3.61 82.83
15.74 7.33 6.39 14.33
L + (Ag) + J1
556
U2
L (Ag) J1
38.42 2.15 88.45 0.00
41.58 79.56 2.33 75.00
20.00 18.29 9.22 25.00
L + (Ag) J1 + J2
544
U3
L (Ag) J1 J2
42.10 87.06 0.00 86.02
34.60 1.68 75.00 0.79
23.30 11.26 25.00 13.19
L + J2 J1 +
436
U4
L J2 J1
40.88 77.98 0.00 75.00
16.51 0.10 75.00 0.00
42.61 21.92 25.00 25.00
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Sn Reaction
T [°C]
Type
Phase
53 Composition (at.%) Ag
Cu
Sn
L + J1 +
375
U5
L J1 0
26.68 0.00 75.00 0.00
8.75 75.00 0.00 54.53
64.57 25.00 25.00 45.47
L + + (Sn)
217.7
E1
L (Sn)
3.54 75.00 0.00 0.07
1.06 0.00 54.50 0.01
95.40 25.00 45.50 99.92
L + J1
661
e2 (max)
L J1
0.27 4.54 0.00
74.13 61.28 75.00
25.60 34.18 25.00
Table 3: Thermodynamic Description of the Ag-Cu-Sn System Ag-Cu: [1986Hay] Ag-Sn: [1999Oht] Cu-Sn: [2001Liu] Ternary parameters for the liquid phase (in J#mol–1) L(LIQUID,AG,CU,SN;0) = –63565 L(LIQUID,AG,CU,SN;1) = –99000 L(LIQUID,AG,CU,SN;2) = –22418
Landolt-Börnstein New Series IV/11C3
MSIT®
54
MSIT®
Ag-Cu
Ag-Sn
Ag-Cu-Sn
Cu-Sn 818 p1 l + (Cu) β
779 e1 L (Cu) + (Ag) 772 p2 L + (Ag) ε2
661 e2(max) β L + ε1 L + (Cu) (Ag) + β
586
L + (Ag) + β
U3 (Ag) + ε1 + ε2
L + ε1 + ε2 436
L + ε2 ε1+ θ
ε1 + ε2 + θ
U4
415 p4 L + ε1 η
L + ε1 + θ L + ε1 θ + η
375
221 e5 L (Sn) + θ
Ag–Cu–Sn
L + (Ag) ε1+ ε2
544
U2
(Ag) + β + ε1
L + (Ag) + ε1
477 p3 L + ε2 θ1
ε1 + θ + η
217.7
U5
L+θ+η
L (Sn) + θ + η
Landolt-Börnstein New Series IV/11C3
(Sn) + θ + η
Fig. 1: Ag-Cu-Sn. Part of the reaction scheme involving the liquid phase
U1
(Cu) + (Ag) + β
L + β (Ag) + ε1
556
640 e3 β L + ε1
E1
227 e4 L (Sn) + η
Ag–Cu–Sn
55
Sn (β Sn) e4 E1
Fig. 2: Ag-Cu-Sn. Liquidus surface projection
Data / Grid: at.% Axes: at.%
e5
η
p4
20
80
θ U5
40
60
400 p3 U4
60 0
e3
60 0
60
ε2
ε1
0 50
40
500 e2
p2
700
U3 U2
80
β
20
U1
700
p1
800
90 0
Ag
20
) (Cu
800
(Ag) 40 e
00 10
900
60
1
80
Sn
Cu
Data / Grid: at.% Axes: at.%
Fig. 3: Ag-Cu-Sn. Sn rich corner of the projection of the liquidus surface
(β Sn) e4
e3
E1 250 300
10
350
η
90
400
p5
θ ε1
Ag Cu Sn Landolt-Börnstein New Series IV/11C3
20.00 0.00 80.00
10
Ag Cu Sn
0.00 20.00 80.00
MSIT®
Ag–Cu–Sn
56
Sn
Data / Grid: at.% Axes: at.%
Fig. 4: Ag-Cu-Sn. Isothermal section at 600°C 20
80
40
60
L
60
40
L+ε1
ε2
80
ε1
L+ε2
20
ζ
β
L+(Ag) L+(Ag) L+(Ag)+(Cu)
(Ag)
(Cu)
(Ag)+(Cu) 20
Ag
40
60
80
Sn
Cu
Data / Grid: at.% Axes: at.%
Fig. 5: Ag-Cu-Sn. Isothermal section at 500°C 20
80
L
40
60
L+ε1
60
40
L+ε2
ε2
ε1
80
δ
20
(Ag)+(Cu)+δ (Ag)
Ag
MSIT®
(Cu)
(Ag)+(Cu) 20
40
60
80
Cu
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Sn
57
Sn
Data / Grid: at.% Axes: at.%
Fig. 6: Ag-Cu-Sn. Isothermal section at 400°C 20
80
L 40
60
L+ε1
η
60
L+ε1+θ
ε1+ε2+θ
θ ε2
40
ε1
80
δ
20
ε 1+ε 2 δ +(Ag)+(Cu)
(Ag)
ε1+(Ag) 20
Ag
(Cu)
(Ag)+(Cu) 40
60
80
Sn
Cu
Data / Grid: at.% Axes: at.%
Fig. 7: Ag-Cu-Sn. Isothermal section at 37°C 20
80
(β Sn)+θ +η'
40
60
η'
60
40
ε1+η'+θ
θ 80
ε2
ε1
ε1+ε2+θ
20
ε 1+ε 2
(Cu)+(Ag)+ε1
ε1+(Ag) (Ag)
Ag
Landolt-Börnstein New Series IV/11C3
(Cu) 20
40
60
80
Cu
MSIT®
Ag–Cu–Sn
58
Temperature, °C
Fig. 8: Ag-Cu-Sn. Vertical section at 3.66 mass% Ag, plotted in at.%
L+ε1
L
400
L+η
300
L+θ L+η+θ (β Sn)+θ +η
200
(β Sn)+θ +η' 100
Ag 4.01 Cu 0.00 Sn 95.99
Fig. 9: Ag-Cu-Sn. Vertical section at 10 mass% Sn, plotted in at.%
Ag 3.64 Cu 20.00 Sn 76.36
10
Cu, at.%
1000
L 750
Temperature, °C
L+(Cu) (Ag)+L
(Cu) (Ag)+(Cu)+L
(Ag)
β +(Ag)
(Ag)+(Cu)+β
(Ag)+(Cu)
500
(Ag)+δ
250
ε1+(Ag)+(Cu)
(Ag)+ε1
(Ag)+ε1+ε2
Ag 90.83 Cu 0.00 Sn 9.17
MSIT®
20
40
Cu, at.%
60
80
Ag 0.00 Cu 94.39 Sn 5.61
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Sn
Fig. 10: Ag-Cu-Sn. Vertical section at 25 mass% Sn, plotted in at.%
59
1000
L
Temperature, °C
750
L+β
L+ε2 500
L+(Ag) (Ag)+δ +β
L+ε1+ε2
(Ag)+β
(Ag)+δ +(Cu)
ε1+ε2 (Ag)+δ +ε1
250
(Ag)+ε1
ε1+ε2+θ
Ag 76.75 Cu 0.00 Sn 23.25
(Ag)+(Cu)+ε1 20
40
60
80
Cu, at.%
Sn
Ag 0.00 Cu 84.86 Sn 15.14
Data / Grid: at.% Axes: at.%
Fig. 11: Ag-Cu-Sn. Isoenthalpy curves for Ag-Cu-Sn liquid alloys at 500°C 0
20
80
-1 40
60
-2
60
40
-6 -7 -8 80
Ag
Landolt-Börnstein New Series IV/11C3
0
-9 20
+1
20 +2
+3
40
+4
60
80
Cu
MSIT®
Ag–Cu–Sn
60
Sn
Data / Grid: at.% Axes: at.%
Fig. 12: Ag-Cu-Sn. Isoenthalpy curves for Ag-Cu-Sn liquid alloys at 700°C 0
20
40
80
-1
60
-2 60
40
-4 -3 80
Ag
0
20
+1
20
+2
40 +3
60
+4
80
Sn
Cu
Data / Grid: at.% Axes: at.%
Fig. 13: Ag-Cu-Sn. Calculated surface tension (in mN/m) contours for liquid Ag-Cu-Sn alloys at 800°C
20
80
550 40
60
600 (mN/m) 60
40
650 750 80
850 900
20
950 0 105
0 100
Ag
MSIT®
20
40
60
80
Cu
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Sn
61
Sn
Data / Grid: at.% Axes: at.%
Fig. 14: Ag-Cu-Sn. Calculated surface tension (in mN/m) contours for liquid Ag-Cu-Sn alloys at 1000°C
500 20
80
40
60
600
60
40
700(mN/m)
80
800
20
900 1000 1100 20
Ag
40
60
80
Sn
Cu
Data / Grid: at.% Axes: at.%
Fig. 15: Ag-Cu-Sn. Calculated isoviscosity (in mPa#s) contours for liquid Ag-Cu-Sn alloys at 800°C
20
1.0
80
40
60
1.5 (mPa.s) 60
40
2.0 80
2.5
20
3.0
Ag
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Cu
MSIT®
Ag–Cu–Sn
62
Sn
Data / Grid: at.% Axes: at.%
Fig. 16: Ag-Cu-Sn. Calculated isoviscosity (in mPa#s) contours for liquid Ag-Cu-Sn alloys at 1000°C
20
80
1.0 40
60
1.25(mPa.s)
60
40
2.0 80
20
1.5 2.5 4.0
Ag
MSIT®
20
40
60
80
Cu
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Ti
63
Silver – Copper – Titanium Ortrud Kubaschewski, updated by Jozefien De Keyzer, Rainer Schmid-Fetzer, Oleh Shcherban, Vasyl Tomashik, Yan Jialin, Ludmila Tretyachenko Introduction The first evaluation within the ongoing MSIT Evaluation Programs was made by [1988Kub], which is updated by the present work. The ternary Ag-Cu-Ti system has been investigated by [1969Ere1, 1969Ere2, 1970Ere1, 1970Ere2] using thermal and metallographic analyses supported by microhardness and X-ray measurements. This is the only systematic experimental study of equilibria in this system up to now. The system was later reassessed by [1977Cha] and reviewed in [1979Dri]. The system is characterized by the existence of a wide miscibility gap in the liquid state, formation of which is in agreement with criteria expressed in [1991But] and thermodynamic calculations of [1995Pau1, 1995Pau2, 1994Gub]. Thermodynamics of Ti oxidation was studied in Ag-Cu-Ti melts. Influence of Ti addition on segregation in Ag rich alloys [1981Duk, 1981Pav] and ageing of Cu rich alloys [1980Bzo] were studied in the ternary Ag-Cu-Ti system. Phase equilibria in the solid state are characterized by the formation of a continuous solid solution Ti2(Ag,Cu). The equilibria in the ternary Ag-Cu-Ti system with participation of Ti2Cu3 and TiCu2 phases are doubtful due to the discrepancies with regard to the formation and transformation of these phases in the binary Cu-Ti system [2002Ans]. Binary Systems All three binary systems have been reevaluated recently [2002Rom, 2002Ans, 2002Li] and are accepted here. Respective changes were applied to the here evaluated ternary Ag-Cu-Ti system to fit the phase diagrams of boundary binary systems. Only in the Cu-Ti system uncertainty in the concentration range 55-70 at.% Cu exists. The phase diagram of the binary Cu-Ti system was studied originally by Eremenko et al. and refined during the investigation of ternary Ag-Cu-Ti system [1970Ere2]. According to the accepted phase diagram by Eremenko et al. and [2002Ans] the Ti2Cu phase is a congruently melting compound contrary to the peritectic formation assessed in [Mas2]. Such formation of Ti2Cu leads to the existence of a eutectic in the Ti rich region of the Cu-Ti system. Main discrepancies in the Cu-Ti system refer to the Ti2Cu3 phase which according to the accepted phase diagram forms through a peritectoid reaction Ti3Cu4 + TiCu2 Ti2Cu3 at 875°C and is stable at room temperature, opposite to a peritectic formation Ti3Cu4 + L Ti2Cu3 at 890°C and doubtful eutectoid decomposition at 800°C determined by [1970Ere2]. Such formation of the Ti2Cu3 phase affects equilibria of the high-temperature TiCu2 compound which according to the accepted phase diagram forms at higher temperature through a peritectic reaction Ti3Cu4 + L TiCu2 at 890°C. Solid Phases No ternary compounds form in the Ag-Cu-Ti system. A regular solid solution exists between isotypic Ti2Ag and Ti2Cu [1969Ere1, 1970Ere1, 1970Ere2]. The dependence of cell parameters in this solid solution obeys Vegad’s law [1970Ere1], see Table 1. Crystallographic data of the phases of the Ag-Ti and Cu-Ti systems are listed in Table 1 (the notation of phases adopted originally by Eremenko et al. added). Invariant Equilibria Invariant equilibria reported by [1970Ere1, 1970Ere2] were modified to fit the accepted phase transformations in the binary systems and are presented in Table 2 and Figs. 1a and 1b. Parts of the reaction scheme (Figs. 1a, 1b) have not been experimentally observed and therefore the scheme should be considered as tentative.
Landolt-Börnstein New Series IV/11C3
MSIT®
64
Ag–Cu–Ti
The composition of liquid was determined for four invariant transformations (Table 2) and critical points of liquid immiscibility in the ternary system K1 (64Ag-2Cu-34Ti (at.%)) and K2 (30Ag-61Cu-9Ti (at.%)). Due to a wide immiscibility region in the ternary system most transformations are located close to the boundary of the binary Cu-Ti system in a narrow 5 at.% Ag region (Figs. 2a and 2b) and should be considered approximate. The decomposition of Ti2Cu3 phase in ternary system assumed by Eremenko et al. to be of the eutectoid type and to take place at 803°C was omitted. The invariant transformation U6 (Ti3Cu4 + TiCu2 L + Ti2Cu3) was added to the reaction scheme proposed by [1970Ere2] at 875-851°C (Figs. 1a, 1b, 2a, 2b) in order to be consistent with the accepted binary Cu-Ti system. Liquidus Surface The liquidus surface based on thermal, metallographic and X-ray analyses undertaken by [1970Ere1] is presented in Figs. 2a and 2b. It was adjusted to be consistent with the binaries. The liquid immiscibility was determined to exist at temperatures above 850°C in the ranges of compositions 5-94.5 at.% Ag, 2-68 at.% Cu and 1.5-65 at.% Ti [1969Ere2]. In the region of existence of the congruent melting compound TiCu, the liquidus surface is raised with the estimated temperature of the maximum critical points L´, L´´ to be 970°C. Isothermal Sections Five isothermal sections at 1300, 1005, 960, 900 and 700°C are studied in the ternary Ag-Cu-Ti system [1969Ere1, 1970Ere2] using X-ray and metallographic analysis. The isothermal section at 1300°C is characterized by a liquid miscibility gap (Fig. 3). The isothermal section at 700°C (Fig. 4) presents the ternary Ag-Cu-Ti system in the solid state. All solid phases except (Ti) are in equilibrium with (Ag) at this temperature. Isotypic Ti2Cu and Ti2Ag form a series of solid solutions, but the cross section at 66.7 at.% Ti is not quasibinary due to equilibrium of TiAg and (Ti) at temperatures above 940°C. Crystallographic data of phases in equilibria at 700°C were determined [1970Ere1]. The solubility of a third component in binary phases was determined approximately to be 5 at.% Ag in TiCu, 2 at.% Cu in TiAg, and less than 2 at.% in others [1969Ere1]. The equilibrium of Ti2Cu3 phase and (Ag) omitted by [1969Ere1] added in the assessment of [1977Cha] is accepted here as the only probable. The phase boundaries, homogeneity ranges, etc. in the isothermal sections of Ag-Cu-Ti system presented in Figs. 3 and 4 were corrected to fit the accepted binary systems. Temperature – Composition Sections Three polythermal sections of ternary Ag-Cu-Ti system at 5 at.% Ag, 60 at.% Ag and TiAg-Cu, based on thermal analysis, have been presented by [1970Ere2]. Due to serious discrepancies with the accepted binary systems they are not presented here. Thermodynamics [1990Pak] determined the activity of Ti in Ag-Cu-Ti melts at 1000°C by measuring the oxygen potential, pO2, in equilibrium with Ti in the melts and titanium oxide. This titanium oxide phase was determined by X-ray diffraction and was identified as Ti2O. In order to determine the activity the free energy of formation of this phase had to be estimated, which results in an uncertainty of the activities of a factor 1.8. The results (Table 3) indicate a positive deviation from the ideal solution behavior. Table 4 shows the effect of Ag content on Ti activity in Ag-Cu melts at 1000°C. The activities of Ti in Ag-28Cu melts at 1000°C also have been measured by an oxygen sensor [2002Ron] and are shown in Table 5. Contrary to the result of [1990Pak], a negative deviation from the ideal solution behavior was found. The difference might be due to the assumption of [1990Pak] that the reaction layer is Ti2O, while [2002Ron] defines the equilibrium titanium oxide phase as TiO, which is according to the phase diagram. [2002Ron] also measured the effect of silver (Table 6) and copper (Table 7) content on the activity. Ag increases the Ti activity coefficient significantly. [1994Gub] calculated two isothermal sections at 1300°C and 800°C. The calculation at 1300°C shows a three phase equilibrium between (Ti) and two liquids which is not found in the experimental diagram. The MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Ti
65
calculation at 800°C is incomplete and lacks several equilibria. [1995Pau1] calculated two isothermal sections at 950 and 1000°C based on the measurements of [1970Ere2] and own optical and SEM/EDX/EPMA observation of diffusion couples of Ti and Ag-Cu (AgCu25%, AgCu40%, AgCu50% and AgCu75% - compositions given in at.%). The calculation is more complete than the one from [1994Gub], but some three-phase equilibria are wrong or lacking at both temperatures. This could be caused by the use of wrong or inaccurate temperatures for the binary reactions. In all calculations, the miscibility gap is smaller than the experimental one, which is probably due to the use of a regular solution model. Notes on Materials Properties and Applications Ag-Cu-Ti alloys are used as an interactive brazing material for joining ceramics [1990Pak, 1995Pau2, 1994Gub, 1998Nak]. Miscellaneous [1981Duk] found that the addition of Ti to Ag-Cu alloys resulted in a non-equilibrium ternary eutectic and the formation of the intermetallic compound TiCu4. [1981Pav] investigated the mechanisms of precipitation in Ag-Cu alloys with small additions of titanium. A vast amount of research was done on the use of AgCuTi as a brazing material. Several authors, [1992Kur], [1993Sue], [1993Kat], [1995Hon] [1997Hao], [1998Pau], [2001Jan], [2003Shi], [2003Bai], [2004Shi] investigated the reaction mechanism and the interfacial morphology. Others did research on the wetting and spreading behavior [1999Lop], [1999Ich], [2001Jan], [2001Sci], [2001Abe], [2001Pal], [2002Iwa], [2003Nov], [2004Muo], and the oxidation behavior [1999Lee]. References [1969Ere1]
[1969Ere2]
[1970Ere1]
[1970Ere2]
[1977Cha]
[1979Dri]
[1980Bzo] [1981Duk]
[1981Pav]
Landolt-Börnstein New Series IV/11C3
Eremenko, V.N., Buyanov, Yu.I., Panchenko, N.M., “Phase Equilibria in the Ti-Cu-Ag System at 700°C” (in Russian), Izv. Akad. Nauk SSSR, Met., 3, 188-192 (1969) (Phase Diagram, Experimental, #, 13) Eremenko, V.N., Buyanov, V.I., Panckenko, V.N., “Phase Separation in the Molten State in the Ti-Cu-Ag System” (in Russian), Izv. Akad. Nauk SSSR, Met., 5, 200-202 (1969) (Phase Diagram, Experimental, #, 6) Eremenko, V.N., Buyanov, Yu.I., Panchenko, N.M., “The Liquidus Surface of the System Ti-Cu-Ag” (in Russian), Poroshk. Metall., 4, 44-48 (1970) (Phase Diagram, Crys. Structure, Experimental, #, 3) Eremenko, V.N., Buyanov, Yu.I., Panchenko, N.M., “Polythermal and Isothermal Cross Sections of the Titanium-Copper-Silver System” (in Russian), Poroshk. Metall., 5, 73-78 (1970), translated in Sov. Powder Metall. Met. Ceram., 10, 410-414 (1970) (Phase Diagram, Experimental, #, 6) Chang, Y.A., Goldberg, D., Neumann, J.P., “Phase Diagrams and Thermodynamic Properties of Ternary Copper-Silver Systems”, J. Phys. Chem. Ref. Data, 6(3), 621-673 (1977) (Phase. Diagram, Review, #, 96) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Copper-Silver-Titanium” (in Russian), in “Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 203-205 (1979) (Phase Diagram, Review, 5) Bzowski, S., “Structure Changes in the CuAg2Ti3 Alloy during Aging” (in Polish), Arch. Hutn., 25(4), 559-606 (1980) (Experimental, 20) Dukiet-Zawadska, B., Pawlowski, A., Ciach, R., Tasior-Grabianowska, K., Wolczynski, W., “Dendritic Segregation of Ag-Cu Alloys with the Addition of Ti, Al, Mg and Ni” (in Polish), Arch. Hutn., 26(3), 429-448 (1981) (Experimental, 20) Pavlovski, A., “Quantitative Analysis of the Dendritic Segregation in Ternary Alloys on the Basis of the Silver-Copper System” (in Russian), Fiz. Met. Metalloved., 52(4), 760-766 (1981) (Experimental, 8) MSIT®
66 [1988Kub]
[1990Pak] [1991But] [1992Kur]
[1993Kat]
[1993Sue]
[1994Gub]
[1995Hon]
[1995Pau1]
[1995Pau2]
[1997Hao]
[1998Nak]
[1998Pau]
[1999Ich]
[1999Lee] [1999Lop]
[2001Abe]
MSIT®
Ag–Cu–Ti Kubaschewki, O., “Silver - Copper - Titanium”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.19244.1.20, (1988) (Crys. Structure, Phase Diagram, Assessment, 11) Pak, J.J., Santella, M.L., Fruehan, R.J., “Thermodynamics of Ti in Ag-Cu Alloys”, Metall. Trans. B, 21B(2), 349-355 (1990) (Thermodyn., Experimental, 16) Butt, M.T.Z., Bodsworth, C., “Liquid Immiscibility in Ternary Metallic Systems”, Mater. Sci. Tech., 7(9), 795-802 (1991) (Phase Relations, Review, 39) Kurihara, Y., Takahashi, Sh., Ogihara, S., Kurosu, T., “Bonding Mechanism Between Aluminum Nitride Substrate and Ag-Cu-Ti Solder”, IEEE Trans., Comp. Hyb. Man. Technol., 15(3), 361-368 (1992) (Experimental, Interface Phenomena, Morphology, 22) Kato, S., Yano, T., Iseki, T., “Interfacial Structures Between Ag-Cu-Ti Alloy and Sintered SiC with Various Additives”, J. Ceram. Soc. Jpn., 101(3), 325-330 (1993) (Experimental, Interface Phenomena, Morphology, 24) Suenaga, S., Koyama, M., Arai, Sh., Nakahashi, M., “Solid-State Reactions of the Ag-Cu-Ti Thin Film-Al2O3 Substrate System”, J. Mater. Res., 8(8), 1805-1811 (1993) (Crys. Structure, Experimental, 13) Gubbels, G.H.M., Heikinheimo, L.S.K., Klomp, J.T., “A Comparison Between Titanium-Aluminium Diffusion Bonding and Titanium Active Brazing”, Z. Metallkd., 85(12), 828-832 (1994) (Thermodyn., Phase Relations, Experimental, 13) Hongqi, H., Yonglan, W., Zhihao, J., Xiaotian, W., “Interfacial Reaction of Alumina with Ag-Cu-Ti Alloy”, J. Mater. Sci., 30, 1233-1239 (1995) (Experimental, Interface Phenomena, Morphology, Thermodyn., 15) Paulatso, M., van Loo, F.J.J., Kivilahti, J.K., “Thermodynamic and Experimental Study of Ti-Ag-Cu Alloys”, J. Alloys Compd., 220, 136-141 (1995) (Thermodyn., Phase Relations, Experimental, 15) Paulasto, M., Kivilahti, J.K., “Formation of Interfacial Microstructure in Brazing of Si3N4 wiht Ti-activated Ag-Cu Filler Alloys”, Scr. Metall. Mater., 32(8), 1209-1214 (1995) (Thermodyn., Phase Relations, Experimental, 15) Hao, H., Wang, Y., Jin, Zh., Wang, X., “Interfacial Morphologies Between Alumina and Silver-Copper-Titanium Alloy”, J. Mater. Sci., 32, 5011-5015 (1997) (Experimental, Morphology, 12) Nakamura, M., Mabuchi, M., Saito, N., Yamada, Y., Nakanishi, M., Shimojima, K., Shigematsu, I., “Joining of a Si-Ti-C-O Fiber-Bonded Ceramic and an Fe-Cr-Ni Stainless Steel with a Ag-Cu-Ti Brazing Alloy”, J. Ceram. Soc. Jpn., 106(9), 927-930 (1998) (Experimental, Mechan. Prop., Morphology, 8) Paulasto, M., Kivilahti, J., “Metallurgical Reactions Controlling the brazing of Al2O3 with Ag-Cu-Ti Filler Alloys”, J. Mater. Res., 13(2), 343-352 (1998) (Calculation, Morphology, Phase Diagram, Thermodyn., 29) Ichimori, T., Iwamoto, Ch., Tanaka, S., “Nanoscopic Analysis of a Ag-Cu-Ti/Sapphire Brazed Interface”, Mater. Sci. Forum, 294-296, 337-340 (1999) (Experimental, Morphology, 9) Lee, D.B., Woo, J.H., Park, S.W., “Oxidation Behavior of Ag-Cu-Ti Brazing Alloys”, Mater. Sci. Eng. A, 268, 202-207 (1999) (Crys. Structure, Experimental, Thermodyn., 25) Lopez-Cuevas, J., Jones, H., Atkinson, H.V., “The Effect of Surface Preoxidation of Sintered Silicon Carbide on its Wettability by Silver-Copper Based Brazing Alloys in Vacuo”, Mater. Sci. Eng. A, 266, 161-166 (1999) (Experimental, Interface Phenomena, 17) Abed, A., Jalham, I.S., Hendry, A., “Wetting and Reaction Between ´-Sialon, Stainless Steel and Cu-Ag Brazing Alloys Containing Ti”, J. Eur. Ceram. Soc., 21, 283-290 (2001) (Experimental, Interface Phenomena, Morphology, 12)
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Ti [2001Jan]
[2001Pal]
[2001Sci]
[2002Ans]
[2002Iwa] [2002Li]
[2002Rom]
[2002Ron] [2003Bai]
[2003Nov]
[2003Shi]
[2004Muo]
[2004Shi]
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67
Janickovic, D., Sebo, P., Duhaj, P., Svec, P., “The Rapidly Quenched Ag-Cu-Ti Ribbons for Active Joining of Ceramics”, Mater. Sci. Eng. A, 304-306, 569-573 (2001) (Experimental, Morphology, 9) Palavra, A., Fernandes, A.J.S., Serra, C., Costa, F.M., Rocha, L.A., Silva, R.F., “Wettability Studies of Reactive Brazing Alloys on CVD Diamond Plates”, Diam. Relat. Mater., 10, 775-780 (2001) (Experimental, Morphology, 18) Sciti, D., Bellosi, A., Esposito, L., “Bonding of Zirconia to Super Alloy with the Active Brazing Technique”, J. Eur. Ceram. Soc., 21, 45-52 (2001) (Experimental, Interface Phenomena, Morphology, 29) Ansara, I., Ivanchenko, V., “Cu-Ti (Copper-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services GmbH, Stuttgart; Document ID: 20.11457.1.20, (2002) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 26) Iwamoto, Ch., Tanaka, Sh.-I., “Atomic Morphology and Chemical Reaction of the Reactive Wetting Front”, Acta Mater., 50, 749-755 (2002) (Experimental, Morphology, 21) Li, C., Lebrun, N., Dobatkina, T., Kusnetsov, V., “Ag - Ti (Silver - Titanium)”, MSIT Binary Evaluation Program,in MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services GmbH, Stuttgart; Document ID: 20.26006.1.20, (2002) (Phase Diagram, Crys. Structure, Assessment, 5) Van Rompaey, T., Rogl, P., “Ag-Cu (Silver - Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, Document ID: 20.14511.1.20, (2002) (Phase Diagram, Crys. Structure, Assessment, 28) Rongti, L., Wei, P., Jian, C., Jie, L., “Thermodynamic Properties of Ti in Ag-Cu-Ti Alloys”, Mater. Sci. Eng. A, 335, 21-25 (2002) (Experimental, Thermodyn., 11) Bai, S.Q., Chen, L.D., Yamamura, A., “Bonding of Copper to Aluminum Nitride Substrate Using Active Alloy Interlayer”, Mater. Sci. Forum, 423-425, 301-304 (2003) (Experimental, Morphology, 9) Novakovic, R., Ricci, E., Muolo, M.L., Giuranno, D., Passerone, A., “On the Application of Modelling to Study the Surface and Interfacial Phenomena in Liquid Alloy-Ceramic Substrate Systems”, Intermetallics, 11, 1301-1311 (2003) (Experimental, Interface Phenomena, Phase Relations, Theory, Thermodyn., 73) Shiue, R.K., Wu, S.K., Chen, S.Y., “Infrared Brazing of TiAl Intermetallic Using BAg-8 Braze Alloy”, Acta Mater., 51(7), 1991-2004 (2003) (Crys. Structure, Experimental, Mechan. Prop., Morphology, Phase Diagram, 30) Muolo, M.L., Ferrera, E., Morbelli, L., Passerone, A., “Wetting, Spreading and Joining in the Alumina-Zirconi-Inconel 738 System”, Scr. Mater., 50, 325-330 (2004) (Experimental, Kinetics, Morphology, 31) Shiue, R.K., Wu, S.K., Chan, C.H., “The Interfacial Reactions of Infrared Brazing Cu and Ti with two Silver-Based Braze Alloys”, J. Alloys Compd., 372(1-2), 148-157 (2004) (Experimental, Morphology, Phase Relations, 16)
MSIT®
Ag–Cu–Ti
68 Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
pure Ag at 25°C [Mas2] dissolves 5 at.% Ti at 959°C dissolves 14 at.% Cu at 780°C
(Cu) < 1084.87
cF4 Fm3m Cu
a = 361.46
pure Cu at 25°C [Mas2] dissolves 8 at.% Ti at 885°C dissolves 0.8 at.% Ti at 450°C dissolves 5 at.% Ag at 780°C
(Ti) 1670 - 790
cI2 Im3m W
a = 330.65
pure Ti(h) at 25°C [Mas2] dissolves 13.5 at.% Cu at 1005°C dissolves 15.5 at.% Ag at 1020°C
(Ti) < 882
hP2 P63/mmc Mg
a = 295.06 c = 468.35
pure Ti(r) at 25°C [Mas2] dissolves 1.6 at.% Cu at 790°C dissolves 4.7 at.% Ag at 855°C
, TiAg < 1020
tP4 P4/nmm TiCu
a = 290.3 c = 574
48 to 50 at.% Ag [Mas2, V-C2]
, TiCu < 982
tP4 P4/nmm TiCu
a = 310.8 to 311.8 48 to 52 at.% Cu [Mas2, V-C2] c = 588.7 to 592.1
J, Ti3Cu4 < 925
tI14 I4/mmm Ti3Cu4
a = 313.0 c = 1994
[Mas2, V-C2]
, Ti2Cu3 < 875
tP10 P4/nmm Ti2Cu3
a = 313 c = 1395
[Mas2, V-C2]
, TiCu2 890 - 870
oC12 Amm2 VAu2
a = 436.3 b = 797.7 c = 447.8
[Mas2, V-C2]
TiCu4 885 - ~400
oP20 Pnma ZrAu4
a = 452.5 b = 434.1 c = 1295.3
~78 to ~80.9 at.% Cu [Mas2, V-C2]
TiCu4 500
tI10 I4/m MoNi4
Ti2Ag1–xCux
tI6 I4/mmm MoSi2
~78 to ~80.9 at.% Cu [Mas2]
a = 295.5 to 294.3 x = 0 to 1 [1970Ere2] c = 1185 to 1077 linear dependence
Ti2Ag < 940
a = 295.2 c = 1185
[Mas2, V-C2]
Ti2Cu < 1012
a = 295.3 c = 1073.4
[Mas2, V-C2]
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Ti
69
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%)* Ag
Cu
Ti
L1 + (Ti) L2 + Ti2(Ag,Cu)
982
U1
L1 L2
13 88
30 10
57 2
L + (Ti) Ti2(Ag,Cu) + TiAg
960
U2
-
-
-
-
L3 L4 + Ti2(Ag,Cu) + TiCu
954
E1
L3 L4
10 84
38 14
52 2
L + TiAg Ti2(Ag,Cu) + (Ag)
929
U3
-
-
-
-
L + Ti2(Ag,Cu) TiCu + (Ag)
908
U4
-
-
-
-
L5 + TiCu L6 + Ti3Cu4
900
U5
L5 L6
6 66
61 32
33 2
Ti3Cu4 + TiCu2 L + Ti2Cu3
875-851
U6
-
-
-
-
L + TiCu Ti3Cu4 + (Ag)
860
U7
-
-
-
-
TiCu2 L + Ti2Cu3 + TiCu4
851
E2
L
5
72
23
L + Ti3Cu4 Ti2Cu3 + (Ag)
843
U8
-
-
-
-
L + Ti2Cu3 TiCu4 + (Ag)
808
U9
-
-
-
-
L + TiCu4 (Cu) + (Ag)
783
U10
-
-
-
-
* - the composition of L estimated from a figure of the liquidus surface [1970Ere1], agrees with those given in mass% in [1977Cha]
Table 3: Activity of Diluted Ti in Eutectic 72Ag-28Cu (mass%) Melts at 1000°C [1990Pak] xTi (final mol fraction Ti in the melt)
aTi # 10
log Ti
0.03087
1.13
0.563
0.03087
1.30
0.624
0.0190
0.753
0.598
0.00863
0.403
0.669
0.00432
0.264
0.786
0.0105
0.468
0.649
0.00414
0.213
0.712
0.0198
0.827
0.621
Landolt-Börnstein New Series IV/11C3
MSIT®
Ag–Cu–Ti
70
Table 4: Activity of Diluted Ti in Eutectic 72Ag-28Cu (mass%) Melts at 1000°C [1990Pak] Alloys (composition in mass%)
xTi aTi # 10 (final mol fraction Ti in the melt)
log Ti
Ag-1Ti
0.0016
1.27
1.90
Ag-10Cu-1Ti
0.0067
1.34
1.30
Ag-28Cu-1Ti
0.0105
0.468
0.649
Table 5: Activity of Diluted Ti in Eutectic Ag-Cu (mass%) Melts at 1000°C [2002Ron] Ti (mass%)
Cu (mass%)
xTi (final fraction Ti in the melt) # 10–3
aTi
ln Ti
0.4
28.74
0.577
4.00 # 10–5
–2.67
0.6
28.4
1.374
4.82 # 10–4
–2.02
–4
0.8
29.15
2.728
5.09 # 10
–1.68
1.0
29.95
4.64
1.01 # 10–3
–1.53
–3
–0.95
1.5
28.54
3.47 # 10
9.002
Table 6: Effect of Silver Content on the Activity Coefficient of Ti in Eutectic Ag-Cu Melts Containing 1 mass% Titanium at 1000°C [2002Ron] xTi
xAg
aTi # 10–4
Ti
0.00142
0.0094
1.51
0.1065
0.00139
0.0198
3.50
0.2516
0.00135
0.0385
6.14
0.4550
0.001295
0.0405
13.1
1.018
Table 7: Effect of Copper Content on the Activity Coefficient of Ti in Eutectic Ag-Cu Melts Containing 1 mass% Titanium at 1000°C [2002Ron] xTi
xCu
aTi # 10–4
Ti
0.00126
0.00474
2.692
0.2136
0.00122
0.0219
0.952
0.0781
0.00117
0.0326
0.658
0.0563
0.00115
0.0355
0.379
0.0330
MSIT®
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Cu-Ti
Ag-Cu-Ti
Ag-Ti 1020 p1 l (βTi) + TiAg
K1 1100
1005 e1 l (βTi) + Ti2Cu
Lx+Ly+(βTi) 982
L1+(βTi)L2+Ti2(Ag,Cu)
U1
Lx+Ly+Ti2(Ag,Cu)
970 max L' + L'' + TiCu
Lx+(βTi)+Ti2(Ag,Cu)
960 e2 l Ti2Cu + TiCu
960 L+(βTi)Ti2(Ag,Cu)+TiAg
Ag-Cu
U2
960 e3 l (βTi) + TiAg
(βTi)+Ti2(Ag,Cu)+TiAg E1 940 p2 (βTi) + TiAgTi2Ag 929 L+TiAgTi2(Ag,Cu)+(Ag) 925 p3 l + TiCu Ti3Cu4
U3
Ag–Cu–Ti
954 L3L4 + Ti2(Ag,Cu) + TiCu
L+Ti2(Ag,Cu)+TiAg
L+Ti2Cu+TiCu
Ti2(Ag,Cu)+TiAg+(Ag) L+Ti2(Ag,Cu)+(Ag) 908 L+Ti2(Ag,Cu)TiCu+(Ag) L+TiCu+(Ag)
900
Ti2(Ag,Cu)+TiCu+(Ag)
L5 + TiCu L6 + Ti3Cu4
U5
71
MSIT®
Fig. 1a: Ag-Cu-Ti. Reaction scheme
U4
72
MSIT®
Cu-Ti
Ag-Cu-Ti
890 p4 l + Ti3Cu4 TiCu2
Ag-Ti
Ag-Cu
Lx + Ly + Ti3Cu4 L5 + TiCu + Ti3Cu4
885 p5 l + (Cu) βTiCu4 875 p6 Ti3Cu4 + TiCu2Ti2Cu3
875-851 Ti3Cu4+TiCu2L+Ti2Cu3 U6
875 e4 l TiCu2 + βTiCu4
L+Ti2Cu3+Ti3Cu4
870 e5 TiCu2Ti2Cu3 + TiCu4
860 L + TiCu Ti3Cu4 + (Ag)
851
TiCu2 L+Ti2Cu3+TiCu4
L+Ti2Cu3+TiCu4
843
E2
500 p7 βTiCu4 + (Cu)αTiCu4 Landolt-Börnstein New Series IV/11C3
400 e9 βTiCu4αTiCu4+Ti2Cu3 Fig. 1b: Ag-Cu-Ti. Reaction scheme (continued)
783
855 e6 (βTi)(αTi) +Ti2Ag
U8
L+Ti2Cu3 +(Ag)
L + Ti2Cu3 TiCu4 +(Ag)
Ti2Cu3+TiCu4+(Ag) 790 e7 (βTi)(αTi) + Ti2Cu
850
(βTi)(αTi)+Ti2Ag
L + Ti3Cu4 Ti2Cu3 +(Ag)
Ti3Cu4 +Ti2Cu3 +(Ag) 808
K2
L+Ti3Cu4+(Ag)
Ag–Cu–Ti
TiCu+Ti3Cu4+(Ag)
L5+TiCu2+Ti2Cu3
U7
U9
L+TiCu4+(Ag)
L + Ti2Cu3 TiCu4 +(Ag)
TiCu4 +(Cu)+(Ag)
U10
780 e8 l (Ag) + (Cu)
Ag–Cu–Ti
73
Cu
Data / Grid: at.% Axes: at.%
Fig. 2a: Ag-Cu-Ti. Liquidus surface projection
(Cu) p5,885 20
β TiCu4
e4,875 p4,890 p3,925
U6,875-851
40
60
L5 K2,850
Ti3Cu4
982 e2,960 TiCu 60
900
L´
Ti2Cu
U10,783 e8,780 40
L3
U9,808 U8,843 U7,860
970
1012 e1,1050
80
E2,851 TiCu2
954
L1
L6
982 80
20
L´´ L4 (β Ti) 20
Ti
L2
K1,1100 40
60
Ti Ag Cu
Fig. 2b: Ag-Cu-Ti. Liquidus surface; enlarged view of Fig. 2a in the Ag rich corner
E2,851 U6,875-851 K2,850
80
0.00 50.00 50.00 p5,885
U4,908 U3,929 U2,960 (Ag)
p1,1020 e3,960
Ag
Data / Grid: at.% Axes: at.%
U10,783 e8,780
10
40
U9,808
900
U8,843 L6
20
U7,860 30
970 30
954
20
L´´ 982 U4,908 L4
40
U3,929 10
L2
(Ag)
K1,1100 Ti Ag Cu Landolt-Börnstein New Series IV/11C3
50.00 50.00 0.00
60
70
U2,960
80
90
p1 e3
Ag
MSIT®
Ag–Cu–Ti
74
Cu
Data / Grid: at.% Axes: at.%
Fig. 3: Ag-Cu-Ti. Isothermal section at 1300°C 20
80
40
60
L1+L2
60
40
80
20
L (β Ti)
β Ti+L 20
Ti
40
60
80
Cu Fig. 4: Ag-Cu-Ti. Isothermal section at 700°C
Ag
Data / Grid: at.% Axes: at.%
(Cu)
20
80
TiCu4
Ti2Cu340 Ti3Cu4
60
TiCu 60
40
Ti2Cu
80
20
Ti2(Ag,Cu) (Ag) (β Ti)
Ti
MSIT®
20
Ti2Ag 40
TiAg
60
80
Ag
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Zn
75
Silver – Copper – Zinc Stephan Schittny, updated by Hans Leo Lukas Introduction After early investigations by Ueno [1929Uen], Keinert [1932Kei] and Weigert [1954Wei], Gebhardt et al. [1962Geb] reexamined the phase relations in the ternary system thoroughly and published a set of very consistent diagrams on the constitution of the system Ag-Cu-Zn. Investigations were carried out by means of differential thermal analysis (DTA) and metallographic and dilatometric techniques. In order to prepare 130 different alloys 99.97% Ag, 99.99% Zn and electrolytic copper were melted together under a suitable slag or in an atmosphere of Ar inert gas. Heat treatments were carried out in sealed evacuated glass tubes at 600, 500 and 350°C with subsequent water quenching. Few alloys were annealed also at 550 or 400°C. From the results the following diagrams were drawn: a liquidus projection, isotherms at 600, 500 and 350°C, isopleths at 20, 40 and 60 mass% Ag (12.84-13.16 at.% Ag, 28.20-28.78 at.% Ag, or 46.92-47.62 at.% Ag, respectively), 40 and 60 mass% Cu (53.08-40.68 at.% Cu or 71.80-60.68 at.% Cu, respectively) and 20 mass% Zn (29.20-19.55 at.% Zn) and a projection of the three- and four-phase equilibria containing liquid onto the concentration triangle. Several reviews of the Ag-Cu-Zn phase diagram were published [1949Jae, 1967McD, 1969Gue, 1973Sis, 1977Cha, 1978Cha, 1979Dri], except two early ones [1949Jae, 1969Gue] mainly based on [1962Geb]. The order-disorder transformation ´ and the transformation of quenched metastable to were studied in detail by several authors [1967Nak, 1967Yon, 1971Mur1, 1971Mur2, 1980Mat, 1982Gra, 1983Mur, 1984Mur, 1985Nak, 1987Mat]. Moeller [1943Moe] measured lattice parameters of the hexagonal close packed J solid solution phase at different compositions and confirmed the continuous range of solid solutions between the Ag-Zn and Cu-Zn J phases. [1997Roe] measured diffusion paths between (Ag) and (Cu) solid solutions, one of them containing up to 25 at.% Zn, the other one without Zn. They determined the compositions of coexisting (Ag) and (Cu) phases after long time (700-850 h) annealing at 550, 650, 670 and 700°C. The Zn vapor pressure of Ag-Cu-Zn alloys was measured by [1965Arg] at 727°C and by [1990Sir] at 1000°C. By means of a high temperature isoperibolic calorimeter [2002Wit] measured partial enthalpies of mixing of the three components in liquid Ag-Cu-Zn alloys at 795 and 840°C and calculated integral enthalpies of mixing from the results. Binary Systems The binary phase diagrams are accepted: Ag-Cu from the MSIT evaluation [2002Rom], Ag-Zn from [Mas2] and Cu-Zn from [1994Mio]. Additional information on metastable Ag-Cu alloys was reviewed by Giessen [1980Gie]. Thermodynamic datasets were assessed for all three binary systems: Ag-Cu by [1986Hay, 1988Jon], Ag-Zn by [1998Gom, 1999Oht], Cu-Zn by [1993Kow]. [1988Jon] seems to contain typing errors for the unary Ag and Cu terms of the liquid description. If, however, the reported excess terms are used and the unary terms of all phases are replaced by the newer ones of [1991Din], the resulting calculated phase diagram reproduces very well the accepted one. [1999Oht] contains a typing error: instead of the excess terms of the phase those of the phase are repeated, thus the description of is incomplete. Solid Phases The solid phases are listed in Table 1. Continuous solid solutions designated , and J connect the corresponding isostructural binary Ag-Zn and Cu-Zn phases. The phase undergoes an order-disorder transformation. In the vicinity of the binary Ag-Zn system this transformation is in a metastable range slightly below the temperature of formation of the stable phase. This range can be easily accessed by
Landolt-Börnstein New Series IV/11C3
MSIT®
76
Ag–Cu–Zn
quenching. [1967Yon, 1971Mur1, 1971Mur2, 1973Mur, 1980Mat, 1982Gra, 1983Mur, 1984Mur, 1985Nak, 1987Mat] studied this transformation in detail. [1983Mur, 1984Mur] found additionally spinodal phase separation in the phase, leading to a modulated microstructure coarsening with annealing time. The mutual solubilities of the (Ag) and (Cu) phases increase somewhat with increasing Zn content [1962Geb, 1997Roe]. The phase dissolves about 1.7 at.% Cu [1980Mat]. The Ag solubility in the phase was estimated by [1962Geb] as 25 mass% (16.7 at.%). Ternary phases were not found in this system. Invariant Equilibria Two invariant four-phase reactions were identified by [1962Geb]. Their temperatures and phase compositions are shown in Table 2. The complete reaction scheme is adopted from [1962Geb] and shown in Fig. 1. The temperatures of the two binary invariant reactions p7 and p8 are changed to the values assessed by [1994Mio], 600 to 598°C and 424 to 425°C, respectively. Liquidus Surface The liquidus surface after Gebhardt et al. [1962Geb], transformed to at.%, is shown in Fig. 2. Isothermal Sections Three isothermal sections at 600, 500 and 350°C are shown in Figs. 3a, 3b, 3c [1962Geb], converted to at.%. The Zn corner of the 600°C isotherm drawn by [1962Geb] corresponds to a slightly higher temperature, it disagrees with the 600°C isotherm of the liquidus surface and the temperature of 598°C for the binary L+ +J three-phase equilibrium. Therefore Fig. 3a was changed slightly compared with [1962Geb]. On the 350°C isotherm, the and ´ (ordered) regions were labeled incorrectly as ´ and , respectively, in the original paper [1962Geb]. Temperature – Composition Sections Two vertical sections at 20 mass% Ag (12.84-13.16 at.% Ag) and 20 mass% Zn (29.20-19.55 at.% Zn) are shown in Figs. 4a, 4b, converted to at.%. Figure 5 shows the section at constant Zn content of 45 at.%. The order-disorder temperature of the /´ transformation and the spinodal of the phase, determined by [1984Mur] are included in Fig. 5. Some of the lines are measured only partially and drawn incompletely in the diagram. Thermodynamics The vapor pressure of Zn in the (Ag) and (Cu) phases was measured by [1965Arg] at 727°C. Replacing Ag by Cu in (Ag) decreases the activity coefficient of Zn, but replacing Cu by Ag in (Cu) within the accuracy of the measurements does not change the activity coefficient of Zn. Two tie lines between (Ag) and (Cu) were constructed from the results: (Ag +10.7 at.% Cu +9.8 at.% Zn) to (Cu +5.6 at.% Ag +13.5 at.% Zn) and (Ag +10.7 at.% Cu +14.6 at.% Zn) to (Cu +6.2 at.% Ag +18.6 at.% Zn) (digitized from a graph). Witusiewicz et al. [2002Wit] determined partial enthalpies of all three elements in liquid at 795 or 840°C along the two lines with Ag:Cu ratios 2:3 and 3:2. An assessment of a thermodynamic dataset was referenced by [1997Roe] as “to be published”. However, no publication could be found until now. The calculated isothermal section at 670°C, shown by 1997Roe], deviates by several at.% from the experimental results of [1962Geb]. Notes on Materials Properties and Applications Alloys near the eutectic composition E1 are proposed as Cd free brazing alloys to connect TiAl based alloys with Cr-steel [2002Hui]. Ag with 3% Zn, 1.75% Cu and 1.25% Sn (mass%) after internal oxidation is proposed to replace the toxic standard Ag-CdO material for electric contacts [1993Dev].
MSIT®
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77
References [1929Uen] [1932Kei] [1943Moe] [1949Jae] [1954Wei] [1962Geb] [1965Arg] [1967McD] [1967Nak]
[1967Yon]
[1969Gue] [1971Mur1]
[1971Mur2]
[1973Mur]
[1973Sis] [1977Cha]
[1978Cha]
[1979Dri]
[1980Mat]
[1980Gie]
Landolt-Börnstein New Series IV/11C3
Ueno, S., “The System Ag-Cu-Zn”, Mem. Coll. Eng. Kyoto Imp. Univ., A12, 347-373 (1929) (Phase Relations, Experimental, 3) Keinert, M., “The System Silver-Copper-Zinc”, Z. Phys. Chem., 160, 15-33 (1932) (Phase Relations, Phase Diagram, Experimental, 12) Moeller, K., “X-ray and Microscopic Investigation in the Four-Component Zn-Mn-Cu-Ag System”, Z. Metallkd., 35, 27-28 (1943) (Crys. Structure, Experimental, 4) Jaenecke, E., “Ag-Cu-Zn” (in German), Kurzgefasstes Handbuch aller Legierungen, 563-565 (1949) (Phase, Relations, Phase Diagram) Weigert, K.M., “Constitution and Properties of Ag-Cu-Zn Brazing Alloys”, JOM - J. Min. Met. Mat. Sci., 200, 233-237 (1954) (Phase Relations, Experimental, 6) Gebhardt, E., Petzow, G., Krauss, W., “On the Constitution of the Cu-Ag-Zn System”, Z. Metallkd., 53, 372-379 (1962) (Phase Relations, Phase Diagram, Experimental, #, 8) Argent, B.B., Lee, K.T., “Thermodynamic Properties of Solid Solutions”, Trans. Faraday Soc., 61, 826-833 (1965) (Thermodyn., Experimental, 15) McDonald, A.S., Price, B.R., Sistare, G.H., “Ternary and Higher-Order Alloys of Silver”, Paper from Silver-Economics, Metallurgy and Use, 272-303 (1967) (Phase Diagram, 25) Nakanishi, N., Takehora, H., Murakami, Y., Senda, Y., Sugiyama, H., Kachi, S., “Low Temperature Martensitic Transformation in Ag-Cu-Zn Alloys”, Jpn. J. Appl. Phys., 67, 1341 (1967) (Phase Relations, Experimental, 5) Yono, M., Asano, H., Nakanishi, N., Kachi, S., “-Phase Stability in Hume-Rothery Type Cu-Ag-Zn Ternary Alloys”, Trans. JIM, 8, 277-278 (1967) (Phase Diagram, Phase Relations, Experimental, *, 10) Guertler, M., Anastasiadias, E., “Copper-Zinc-Silver”, A Comp. of Const. Ternary Diagr. Met. Systems, Isr. Pro. Sci. Tr., Jerusalem, 608 (1969) (Phase Diagram, 4) Murakami, Y., Nakanishi, N., Kachi, S., “Superlattice Formation in the Ternary Phase Alloys. Pt. I: Au-Cu-Zn and Ag-Cu-Zn Alloys”, Acta Metall., 19, 93-96 (1971) (Phase Diagram, Phase Relations, Experimental, *, 10) Murakami, Y., Kachi, S., Nakanishi, N. Takehara, H., “Superlattice Formation in the Ternary Phase Alloys, Pt. II: Application of the Statistical Thermodynamics”, Acta Metall., 19, 97-105 (1971) (Crys. Structure, Experimental, 21) Murakami, Y., Nakanishi, N., Takehara, H., Kachi, S., “Electric and Magnetic Properties of Au-Cu-Zn and Ag-Cu-Zn Phase Alloys”, J. Phys. Soc. Jpn., 34, 557 (1973) (Electr. Prop., Experimental, Magn. Prop., 6) Sistare, G.H., “Ag-Cu-Zn (Silver-Copper-Zinc)”, Metals Handbook, 8, 380-381 (1973) (Phase Diagram, Review, 1) Chang, Y.A., Goldberg, D., Neumann, J.P., “Phase Diagrams and Thermodynamic Properties of Ternary Copper Silver Systems”, J. Phys. Chem. Ref. Data, 6(2), 621-673 (Ag-Cu-Zn 669-673) (1977) (Phase Diagram, Phase Relations, Review, 6) Chang, Y.A., Neumann, J.P., Choudary, U.V., “Evaluations of Phase Diagrams and Thermodynamic Properties of Ternary Copper Alloy Systems”, NBS Special Publication 496, Application of Phase Diagrams in Metallurgy and Ceramics, 1, Carter, G.C. (Ed.), 237-246 (1978) (Phase Diagram, Phase Relations, Review, 6) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Cu-Ag-Zn” (in Russian) in “Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 206-208 (1979) (Phase Diagram, 6) Matsuo, Y., “The Effect of Additional Elements on the ´ Transformation in Equiatomic AgZn Alloy”, Trans. Jpn. Inst. Met., 21(3), 174-178 (1980) (Phase Relations, Phase Diagram, Experimental, 7) Giessen, W.C., “Metastable Ag-Cu System” in “The Ag-Cu System, Provisional”, Bull. Alloy Phase Diagrams, 1, 43-44 (1980) (Phase Diagram, Phase Relations, Review, 5) MSIT®
78 [1982Gra]
[1983Mur]
[1984Mur]
[1985Nak]
[1986Hay]
[1987Mat]
[1988Jon]
[1990Sir]
[1991Din] [1993Dev] [1993Kow] [1994Mio]
[1997Roe]
[1998Gom] [1999Oht]
[2002Hui]
[2002Rom]
[2002Wit]
MSIT®
Ag–Cu–Zn Granovsky, M., Resta Levi, M., Arias, D., “Phase Transformations in Ag48Zn50Au2 and Ag48Zn50Cu2 Alloys”, Metallography, 15(3), 213-224 (1982) (Phase Relations, Experimental, 17) Murakami, Y., Nakanishi, N., “Modulated Structures and the Early Stage of Prepitation in Phase Ag-Cu-Zn Alloys” (in Japanese), J. Jpn. Inst. Met., 47(6), 470-475 (1983) (Phase Relations, Experimental, 21) Murakami, Y., Kachi, S., Nakanishi, N., “Ordering and Phase Separation in the Ternary Ag-Cu-Zn -Phase Alloys”, Acta Met., 32(5), 629-636 (1984) (Experimental, Phase Relations, Phase Relations, 30) Nakashini, N., Murakami, Y., “The Effect of Silver Addition on the Bainitic Transformation of a CuZn Phase Alloy” (in Japanese), J. Jpn Inst. Met., 49(5), 332-336 (1985) (Crys. Structure, Experimental, Phase Relations, 12) Hayes, F.H., Lukas, H.L., Effenberg, G., Petzow, G., “A Thermodynamic Optimisation of the Cu-Ag-Pb System”, Z. Metallkd., 77, 749-754 (1986) (Calculation, Thermodyn., Phase Relations, 36) Matsuo, Y., Torii, Y., “Dynamical Process of the ´- Transformation in Silver-Zinc Alloys” (in Japanese), J. Jpn. Inst. Met., 51(1), 31-36 (1987) (Electr. Prop., Experimental, Kinetics, Phase Relations, 10) Jönsson, B., &gren, J., “Thermodynamic and Kinetic Aspects of Crystallisation of Supercooled Ag-Cu Liquids”, J. Less-Common Met., 145, 153-166 (1988) (Calculation, Thermodyn., Phase Relations, 30) Siromakha, A.K., Shorikov, Yu.S., “Thermodynamic Analysis of the Vacuum Distillation of Zinc from Copper-Silver-Zinc Alloys” (in Russian), Metally, (3), 36-39 (1990) (Calculation, Phase Relations, Thermodyn., 6) Dinsdale, A.T., “SGTE Data for Pure Elements”, Calphad, 15(4), 317-425 (1991) (Review, Thermodyn.) Dev, S.C., Basak, O., Mohanty, O.N., “Defelopment of Cadmium-Free Silver Metal-Oxide Contact Materials”, J. Mater. Sci., 28, 6440-6544 (1993) (Experimental, Morphology, 11) Kowalski, M., Spencer, P.J., “Thermodynamic Reevaluation of the Cu-Zn System”, J. Phase Equilib., 14(4), 432-438 (1993) (Assessment, Thermodyn., Phase Relations, 36) Miodownik, A.P., “Cu-Zn (Copper-Zinc)”, Phase Diagrams of Binary Copper Alloys, Subramanian, P.R., Chakrabarti, D.J. Laughlin, D.E. (Eds.), ASM International, Metals Park, OH 487-496 (1994) (Assessment, Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., 98) Roenkae, K.J., van Loo, F.J.J., Kivilathi, J.K., “Thermodynamic and Kinetic Study of Diffusion Paths in the System Cu-Ag-Zn”, Z. Metallkd., 88(1), 9-13 (1997) (Experimental, Phase Relations, Thermodyn., 10) G?mez-Acebo, T., “Thermodynamic Assessment of the Ag-Zn System”, Calphad, 22(2), 203-220 (1998) (Assessment, Thermodyn., Phase Relations, 32) Ohtani, H., Miyashita, M., Ishida, K., “Thermodynamic Study of Phase Equilibria in the Sn-Ag-Zn System” (in Japanese), J. Jpn Inst. Met., 63(6), 685-694 (1999) (Assessment, Thermodyn., Phase Relations, 68) Huijie, L., Jicai, F., “Vacuum Brazing TiAl-Based Alloy to 40Cr Steel Using Ag-Cu-Zn Filler Metal”, J. Mater. Sci. Lett., 21, 9-10 (2002) (Experimental, Mechan. Prop., Morphology, 9) van Rompaey, T., Rogl., P., “Ag - Cu (Silver - Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14511.1.20, (2002) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 29) Witusiewicz, V.T., Hecht, U., Rex, S., Sommer, F., “Partial and Integral Enthalpies of Mixing of Liquid Ag-Al-Cu and Ag-Cu-Zn Alloys”, J. Alloys Compd., 337, 189-201 (2002) (Experimental, Thermodyn., 30) Landolt-Börnstein New Series IV/11C3
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79
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.62
pure element, 25°C [Mas2]
(Cu) < 1084.87
cF4 Fm3m Cu
a = 361.47 a = 369.74
pure element, 25°C [Mas2] Cu65.44Zn33.56 [P]
(Zn) < 419.58
hP2 P63/mmc Mg
a = 266.49 c = 494.68
pure element, 25°C [Mas2]
, (Ag,Cu)Zn (h)
cI2 Im3m W
AgZn 710 - 258
a = 311.0
AgZn [V-C2]
CuZn 903 - 454
a = 299.67
CuZn, 871°C [P]
´, (Ag,Cu)Zn (l) < 468
cP2 Pm3m CsCl
a = 315.58 a = 295.39
Ag50.1Zn49.9 quenched [P] Cu52.34Zn47.66 [P]
, AgZn (l) < 274
hP9 P3 AgZn(l)
a = 763.55 c = 282.00
Ag50Zn50 [P]
, CuZn3 700 - 560
hP3 P6 CuZn3
a = 427.5
[V-C2]
Landolt-Börnstein New Series IV/11C3
MSIT®
Ag–Cu–Zn
80 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, (Ag,Cu)5Zn8 Ag5Zn8 < 661
cI52 I43m Cu5Zn8
Cu5Zn8 < 834 J, (Ag,Cu)Zn4
hP2 P63/mmc Mg
J (Ag-Zn) < 631
CuZn4 < 574
Lattice Parameters Comments/References [pm]
a = 934.07
[V-C2]
a = 886.9
[V-C2]
a = 276.0 c = 431.6
Ag4Cu15Zn81 [1943Moe]
a = 277.7 c = 434.5
Ag10Cu10Zn80 [1943Moe]
a = 279.8 c = 439.4
Ag16Cu5Zn79 [1943Moe]
a = 282.6 c = 448.49
Ag33Zn67 [P]
a = 281.16 c = 439.96
Ag12Zn88 [P]
a = 273.83 c = 429.37
Cu21Zn79 [P]
a = 276.53 c = 429.78
Cu13Zn87 [P]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Ag
Cu
Zn
L (Cu) + (Ag) +
665
E1
L (Cu) (Ag)
43.2 5.5 64.4 37.4
26.2 66.0 7.3 29.2
30.6 28.5 28.3 33.4
L+ J+
630
U1
L J
17.5 16.8 17.5 16.8
10.9 13.6 12.0 12.5
71.6 69.6 70.5 70.7
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Landolt-Börnstein New Series IV/11C3
Ag-Cu
Ag-Zn
Ag-Cu-Zn
Cu-Zn 902 p1 l + (Cu) β 834 p2 l+βγ
779 e1 l (Cu) + (Ag) 710 p3 l + (Ag) β
665
L (Cu) + (Ag) + β
E1 700 p4 l+γδ
(Cu) + (Ag) + β 661 p5 l+βγ
630
L+γε+δ
Ag–Cu–Zn
631 p6 l+γε U1 598 p7 l+δε
431 p8 l + ε (Zn)
560 e2 δγ+ε 425 p9 l + ε (Zn)
274 p10 β+γζ 258 e3 β (Ag) + ξ
81
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Fig. 1: Ag-Cu-Zn. Reaction scheme
Ag–Cu–Zn
82
Zn (Zn) p9
Fig. 2: Ag-Cu-Zn. Liquidus surface projection
p7 650 p4
20
Data / Grid: at.% Axes: at.%
p8 500°C
550
600
δ
80
ε
U1 800 750
40
p6
γ
650 p5
60
p2
700 850
β
60
40
p1
p3
E1
800
80
0 75
850
900
950
0 70
750
20
0 80 (Ag)
(Cu)
0 85
1050 1000 20
Cu
60e1
40
0 90
80
Zn
Ag
Data / Grid: at.% Axes: at.%
Fig. 3a: Ag-Cu-Zn. Isothermal section at 600°C
L
L+δ
δ
20
δ+ε γ +ε
γ +δ 40
80
L+ε
ε
γ
60
β+γ β
60
40
β +(Cu) (Ag)+(Cu)+β
(Cu)
β +(Ag) (Ag)
80
20
(Cu)+(Ag)
Cu
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40
60
80
Ag
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Zn
83
Zn
Data / Grid: at.% Axes: at.%
Fig. 3b: Ag-Cu-Zn. Isothermal section at 500°C
L L+ε 20
80
ε γ +ε γ
40
60
β +γ β 60
(Cu)+β
40
(Ag)+β
(Ag)+(Cu)+β
(Cu)
(Ag)
80
20
(Cu)+(Ag)
20
Cu
40
60
80
Zn Fig. 3c: Ag-Cu-Zn. Isothermal section at 350°C
Ag
Data / Grid: at.% Axes: at.%
(Zn)
ε+(Zn) ε
20
80
γ +ε γ
40
β ´+γ
60
(Cu)
β´
(Cu)+β ´
60
β +γ β 40
(Ag)+β
(Ag)+(Cu)+β ´
(Ag)
80
20
(Cu)+(Ag)
Cu
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Ag
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Ag–Cu–Zn
84
Fig. 4a: Ag-Cu-Zn. Vertical section at 20 mass% Ag, plotted in at.%
900
L+(Cu)+β L+(Cu)
L 800
L+β +γ
L+β
Temperature, °C
L+(Cu)+(Ag) 700
L+γ+δ L+δ δ L+δ +ε
600
L+γ
γ+δ
β
γ+δ +ε L+ε
(Cu)+(Ag)+β
δ +ε
γ
500
ε
(Cu)+β
γ+ε
β +γ
(Cu)+(Ag)
400
(Cu)+(Ag)+β '
β ´+γ β ´ 300
Cu 0.00 Ag 13.20 Zn 86.80
Temperature, °C
Fig. 4b: Ag-Cu-Zn. Vertical section at 20 mass% Zn, plotted in at.%
20
(Cu)+β ´
40
60
Cu 87.20 Ag 12.80 Zn 0.00
80
Cu, at.%
900
800
L L+(Cu)
700
L+(Ag)
(Cu) (Ag)
L+(Cu)+(Ag) 600
(Cu)+(Ag) 500
400
300
Cu 80.50 Ag 0.00 Zn 19.50
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20
40
Ag, at.%
60
Cu 0.00 Ag 70.80 Zn 29.20
Landolt-Börnstein New Series IV/11C3
Ag–Cu–Zn
Temperature, °C
Fig. 5: Ag-Cu-Zn. Vertical section at 45 at.% Zn with spinodal of the phase and critical temperature of the /´-ordering. Some lines are not completely determined and not drawn outside the investigated ranges
85
900
L 800
700
β1 600
β 1+β 2
β2
500
400
β ´ (ordered)
β ´+(Cu) 300
200
Ag 0.00 Cu 55.00 Zn 45.00
Landolt-Börnstein New Series IV/11C3
β ´+(Ag) ζ 10
20
30
Ag, at.%
40
50
Ag 55.00 Cu 0.00 Zn 45.00
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86
Ag–In–Sb
Silver – Indium – Antimony Nuri Solak Introduction The experimental base on this system was laid by [1965Gub, 1969Kuz, 1971Kuz] and later assessed by [1988Sch]. In the early study of [1965Gub] a sample of composition Ag3InSb was fused directly from the elements, studied by thermal analysis and metallography and found not to be a single phase material. Using the same technology [1969Kuz] confirmed this in studies of 14 samples along the vertical section InSb-Ag up to 71.9 at.% Ag (70 mass% Ag). However, the primary crystallization is interpreted by [1965Gub] to be InSb whereas for the same sample composition an unidentified ternary phase was reported to crystallize first by [1969Kuz]. Further, the temperatures of primary and eutectic crystallization disagree [1965Gub, 1969Kuz]. The later published study by [1971Kuz] is not helpful in this respect as the published isopleth based on DTA, X-ray analysis and metallography is virtually identical to that of [1969Kuz]. [2004Liu] reinvestigated the system using differential scanning calorimetry (DSC), metallography and electron probe microanalysis (EPMA). They also modeled the system using the CALPHAD approach. In contradiction with [1971Kuz] they reported that there is no ternary compound in the system. The phase extends in the ternary from the Ag-In to the Ag-Sb side and corresponds to one of the ternary compounds mentioned by [1971Kuz]. According to [2004Liu] the second ternary compound mentioned by [1971Kuz] is the (Sb) solid solution with very small amounts of Ag and In. Binary Systems None of the phase diagrams published on Ag-In reflects fully the present knowledge. The Ag-In binary system published in 1990 by [Mas2] shows an unlikely asymmetric /+ boundary. Further, [2001Mos] investigated phase boundaries in the Ag-In system using diffusion couples, DSC and metallography. They determined for the phase a narrower homogeneity range than reported earlier. However, the ' and the phases were not taken into account when they constructed the phase diagram, although DSC peaks belonging to these phases were determined. Therefore, combining data from [Mas2, 2001Mos and 1970Cam] results in a new binary Ag-In phase diagram, given in Fig. 1. The In-Sb binary system is accepted from [Mas2] and Ag-Sb from [1993Oka] which actually is the work of [1992Fes]. Solid Phases [2004Liu] determined that no ternary compound exists in this system. The phase extents continuously from the edge binary Ag-In system to the Ag-Sb side. Very small Ag solubility in InSb was determined by [1971Kuz] and confirmed by [2004Liu]. The solid phases are listed in Table 1. Invariant Equilibria [2004Liu] omitted in their computer calculations the phase in the Ag-In binary system and arrived at two ternary eutectic (E) and three ternary transition (U) reactions. The phase has been re-introduced here on the basis of the accepted binary systems. This generates a new four-phase equilibrium of unknown type and temperature which is connected with p1, p2 and e1 binary reactions. It has been assumed and shown in the reaction scheme given by Fig. 2. The invariant reactions and the changing composition of the liquid phase associated with the phase reactions are given in Table 2. Liquidus Surface The liquidus surface of the system, omitting the phase, was calculated by [2004Liu] based on experimental work of [1971Kuz] along the InSb-Ag vertical section and on their own DSC work along the
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Ag–In–Sb
87
AgIn-Sb vertical section. In Fig. 3 the phase is added the liquidus surface has been slightly modified in order to meet the accepted binary phase boundaries. Isothermal Sections Sections at 350 and 400°C given in Fig. 4 and Fig. 5, are based on those calculated by [2004Liu]. They determined two tie lines in the InSb + phase field and one tie line in (Sb) + having almost the same composition at both temperatures. No three-phase samples were reported. Although calculated sections are close to the experimental results of [1971Kuz] and [2004Liu], slight corrections such as the homogeneity range of the J phase have been done in agreement with the accepted binaries. The three-phase fields are drawn by dashed lines, since tie lines are not sufficient to describe boundaries. Temperature – Composition Sections Isopleths along InSb-Ag and AgIn-Sb are given in Fig. 6 and Fig. 7, respectively. The vertical section InSb-Ag up to 71.9 at.% Ag (70 mass% Ag) was experimentally determined by [1971Kuz] and later calculated by [2004Liu] up to pure Ag. The Ag rich part of the diagram is described here by dashed lines, since there were no experimental data available on homogeneity ranges of (Ag) and phase. [2004Liu] calculated also the AgIn-Sb section using their DSC results which were conducted on 4 samples. The temperatures determined for the invariant lines are consistent with the work of [1971Kuz]. In the Sb rich part of the diagram a calculated (Sb) + two phase field exists up to 3.5 mass% Sb at 424°C which is not consistent with the calculated isothermal sections and EPMA analysis results. On the other hand the system was calculated down to 100°C but the results are not consistent with experimental works. Both isopleths are redrawn according to experimental results but the part calculated below 300°C was not taken into account. Thermodynamics [2001Ita] measured the activity of indium in the Ag-In-Sb system using a galvanic cell with zirconia solid electrolyte in temperature ranges between 697 and 1007°C. Isoactivity curves at 827 and 927°C are given in Fig. 8. They are derived from combining the activity data of the In-Sb and Ag-In binaries. The shape of the curves exhibits a slight bend to the Ag-Sb binary side. [1987Gat] determined enthalpy of mixing of the ternary liquid in a heat flow calorimeter. The measured ternary mixing isoenthalpy curves are given in Fig. 9. Miscellaneous Surface tension and viscosity of liquid phase calculated on the basis of the thermodynamic modeling were given in the study of [2004Liu]. References [1965Gub]
[1969Kuz]
[1970Cam]
[1971Kuz]
Landolt-Börnstein New Series IV/11C3
Gubskaya, G.F., Ping-Nan, W., Luzhnaya, N.P., Kudryavtsev, D.L., “Reaction in Ternary Systems Ag-BIII-CV”, Izv. Akad. Nauk SSSR, Neorg. Mater., 1, 188-192 (1965) (Experimental, Phase Relations, 4) Kuznetsov, G.M., Bobrov, A.P., “Reaction of Silver, Copper and Cadmium with Indiumantimonide”, Izv. Vyssh. Ucheb. Zaved. Tsvet. Met., 12, 126-128 (1969) (Phase Relations, Experimental, 8) Campbell A.N. Wagemann,R., Ferguson,R.B. “The Silver-Indium System: Thermal Analysis, Photomicro-graphy, Electron Microprobe, and X-Ray Powder Diffraction Results”, Canad. J. Chem., 48, 1703-1715 (1970) (Phase Diagram, Phase Relations, Experimental, *, 10) Kuznetsov, G.M., Bobrov, A.P., “Reaction of InSb with Metals”, Izv. Akad. Nauk SSSR, Neorg. Mater., 7, 766-768 (1971) (Phase Relations, Experimental, 6)
MSIT®
Ag–In–Sb
88 [1987Gat]
[1988Sch]
[1992Fes]
[1993Oka] [2001Ita]
[2001Mos]
[2004Liu]
Gather, B., Schroter, P., Blachnik, R., “The Enthalpies of Mixing in the Liquid State of the Ternary Systems Ag-In-Sn, Ag-Sn-Sb, Ag-In-Sb and In-Pb-Sb” (in German), Z. Metallkd., 78(4), 280-285 (1987) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 20) Schmid-Fetzer, R., “The Silver-Indium-Antimony”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.14857.1.20, (1988) (Crys. Structure, Phase Diagram, Assessment, 3) Feschotte, P., Monachon, F., Durussel, Ph., “The Binary System Sb-Ag: a Revision of the Ag3Sb Phase Boundaries”, J. Alloys Compd., 186, L17-L18 (1992) (Phase Diagram, Phase Relations, 4) Okamoto, H., “Ag-Sb Phase Diagram”, J. Phase Equilib., 14, 531-532, (1993) (Phase Diagram, 2) Itabashi, S., Yamaguchi, K., Kameda, K., “Activity Measurement of Indium in Liquid In-Sb-Ag and In-Sb-Te Alloys by EMF Method with Zirconia Solid Electrolyte”, (in Japanese), J. Japan Inst. Metals, 65, 888-892 (2001) (Thermodyn., Experimental, 7) Moser, Z., Gasior W., Pstrus J., Zakulski W., Ohnuma I., Liu X.J., Inohana Y., Ishida K., “Studies of the Ag-In Phase Diagram and Surface Tension Measurements”, J. Electron. Mater., 30, 1120-1128, (2001) (Phase Diagram, Phase Relations, Calculation, Experimental, #, 43) Liu, X.J., Yamaki, T., Ohnuma, I., Kainuma, R., Ishida, K., "Thermodynamic Calculations of Phase Equilibria, Surface Tension and Viscosity in the In-Ag-X (X=Bi, Sb) System", Mater. Trans., JIM, 45, 637-645, (2004) (Phase Diagram, Phase Relations, Calculation, Experimental, #, 30)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
at 25°C [Mas2]
(In) < 156.634
tI2 I4/mmm In
a = 325.3 c = 494.7
at 25°C [Mas2]
(Sb) < 630.755
hR6 R3m As
a = 430.84 c = 1127.4
at 25°C [V-C2]
InSb < 525.7
cF8 F43m ZnS (sphalerite)
a = 647.84
[V-C2] up to 0.16 mass% Ag solubility at 350-400°C [2004Liu]
J, Ag3Sb < 562
oP8 Pmmn Cu3Ti
a = 599 b = 485 c = 524
[V-C2]
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Landolt-Börnstein New Series IV/11C3
Ag–In–Sb Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
hP2 P63/mmc Mg
89
Lattice Parameters Comments/References [pm] Solid solution from Ag-Sb to Ag-In
Ag7Sb
a = 296.16 0.3 c = 479.90 0.3
Pure Ag7Sb [V-C2]
Ag3In
a = 295.63 c = 478.57
Pure Ag3In [V-C2]
694 - 660
cI2 Im3m W
a = 336.82
Ag-In binary 25-30 at.% In [Mas2] High temperature phase
' < 187
cP4 Pm3m AuCu3
a = 414.4 0.4
Ag-In binary 25 at.% In [Mas2], [V-C2]
, Ag9In4 < 301
cP52 P43m Al4Cu9
a = 992.2 0.4
[V-C2]
AgIn2 < 167
tI12 I4/mcm Al2Cu
a = 688.1 0.4 c = 562.0 0.4
[V-C2], [Mas2]
Table 2: Invariant Equilibria Reaction
Type
T [°C]
Phase
Composition (at.%) Ag
In
Sb
L, , (Ag)
? (min)
500 - 600
-
-
-
-
L + InSb ?
e? (max)
417 - 450
-
-
-
-
L (Sb) + + InSb
E1
417
L
45.54
21.14
33.32
L InSb + AgIn2 + (In)
E2
144
L
2.34
97.33
0.33
L + J + (Sb)
U1
455
L
57.42
5.04
37.54
+ InSb L + Ag2In
U2
197
L
6.05
93.01
0.94
L + Ag2In InSb + AgIn2
U3
167
L
4.57
94.91
0.52
L, , , (Ag)
?
?
-
-
-
-
Landolt-Börnstein New Series IV/11C3
MSIT®
Ag–In–Sb
90
Fig. 1: Ag-In-Sb. Ag-In binary system
1000
961.93°C
750
694
Temperature, °C
L
β
671 660
500
(Ag)
ζ
ζ +(Ag) 250
197
187
156.634°C
167
α'
144
γ
(In)
AgIn2 0
Ag
20
40
60
80
In
In, at.%
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Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Ag-In
Ag-Sb 702.5 l + (Ag)
Ag-In-Sb
In-Sb
p1 694 l + (Ag)
p2
660
e1 l ?
562 l+ 488 l
?
L, , , Ag
< 600 ? (min) L, (Ag),
p3
492.5 e2 l InSb + Sb
e3 + Sb
455
L+
+ Sb
U1
Ag–In–Sb
417-450 e ? (max) L InSb +
+ + Sb 417
L
E1
InSb + Sb InSb+ Sb
197
p4 Ag2In + l
167 l + Ag2In
197
+ InSb
U2
+InSb+Ag2In
p5 AgIn2
167
L+ Ag2In
InSb + AgIn2
Ag2In+InSb+AgIn2 144 e5 l (In) + AgIn2
L + Ag2In
144
L
InSb + In
AgIn2
U3 153.9 e4 l InSb + In
E2
InSb+ In AgIn2
91
MSIT®
Fig. 2: Ag-In-Sb. Reaction scheme
Ag–In–Sb
92
Sb
Data / Grid: at.% Axes: at.%
Fig. 3: Ag-In-Sb. Liquidus surface
20
80
(Sb) e2 40
60
500 e3
ε
60
U1
40
450
InSb E1
p3
450
500
ζ
80
450
p1
20
400
700
600 300°C
(Ag)
U3 E2
? 20
Ag
p2
β e1 40
60
80
Sb
5
e4
In
Data / Grid: at.%
(Sb)
Fig. 4: Ag-In-Sb. Isothermal section at 350°C
U2 p4 p e 5
Axes: at.%
20
80
40
60
InSb InSb+ζ+(Sb)
60
40
(Sb)+ζ+ε
ε 80
20
L+InSb+ζ
ζ
L
(Ag)
Ag
MSIT®
20
40
60
80
In
Landolt-Börnstein New Series IV/11C3
Ag–In–Sb
93
Sb Fig. 5: Ag-In-Sb. Isothermal section at 400°C
Data / Grid: at.% Axes: at.%
(Sb)
20
80
40
60
InSb (Sb)+ζ+ε
InSb+ζ +(Sb)
60
40
ε 80
20
b+ζ InS L+
ζ
L (Ag) 20
Ag
Fig. 6: Ag-In-Sb. Temperature composition cut along Ag - InSb
40
60
80
In
1000
900
L
Temperature, °C
800
700
(Ag) 600
500
InSb+L
444°C
ζ
424
400
L+InSb+ζ
InSb+(Sb)+ζ 300
Ag
10
20
30
In, at.%
Landolt-Börnstein New Series IV/11C3
40
Ag 0.00 In 49.96 Sb 50.04
MSIT®
Ag–In–Sb
94
Fig. 7: Ag-In-Sb. Temperature composition cut along AgIn-Sb
700
Temperature, °C
600
L
500
L+(Sb) L+ζ
L+InSb+(Sb) 424
400
InSb+ζ
InSb+ζ +(Sb)
L+InSb+ζ 300
Ag 49.96 In 50.04 Sb 0.00
40
30
20
10
Sb
Ag, at.%
Sb
Data / Grid: at.% Axes: at.%
Fig. 8: Ag-In-Sb. Isoactivity curves of indium at 827 and 927°C 20
80
aIn
827°C
aIn
927°C
0.1
40
60
0.2 0.3
60
40
0.4 0.5 0.6
80
20
0.7 0.8
0.9
Ag
MSIT®
20
40
60
80
In
Landolt-Börnstein New Series IV/11C3
Ag–In–Sb
95
Sb
Data / Grid: at.% Axes: at.%
Fig. 9: Ag-In-Sb. Ternary mixing isoenthalpy curves at 980°C 20
1000 40
80
500
60
0
-500 60
40
-2000
-1000 -1500 80
20
-3500 -4000
-3000
-2500 -2000
-4500
Ag
Landolt-Börnstein New Series IV/11C3
20
-1500 40
60
80
-1000 -500
In
MSIT®
96
Ag–In–Sn
Silver – Indium – Tin Vasyl Tomashik, Rainer Schmid-Fetzer, Ludmila Tretyachenko, Jozefien De Keyzer Introduction The phase equilibria in the Ag-In-Sn system were determined experimentally by [2002Liu] using differential scanning calorimetry (DSC), energy dispersive X-ray spectroscopy, XRD and metallographic techniques and have been summarized in [2003Ohn]. Isothermal sections between 180-600°C were determined together with a number of vertical sections. A thermodynamic assessment of this system was carried out using experimental thermodynamic property and phase equilibria data using the CALPHAD method [2000Ohn, 2002Liu, 2003Ohn]. It was shown that the present dataset can be applied for the prediction not only of phase equilibria but also other properties, such as non equilibrium solidification, surface tension and viscosity in solder alloys. Specimens used for the experimental study were equilibrated at 180, 250, 400 and 600°C before quenching into iced water [2002Liu]. The Ag-In-Sn system has been studied both experimentally and by thermodynamic modeling by [1998Kor]. Six ternary alloys were melted, annealed at 250°C for 100 h, quenched in water and analyzed by optical microscopy (OM) and SEM/EPMA. The interfacial reactions between Ag and binary In-Sn alloys were investigated by diffusion couple experiments which were carried out at 250°C for either 1000 min or 100 h. DSC experiments were carried out to determine the melting temperature of some of the alloys. Calculated isothermal sections at 250 and 113°C were given. [2001Sug] used DSC to investigate the melting/solidification behavior of the alloys in the Sn rich corner of the ternary system. It was shown that strong segregation takes place during cooling of some In-containing solder alloys as the In component is rejected from the solid with excess In remaining in the liquid. At high enough In concentrations (>16 mass%) this segregation results in partial melting of the solder at temperatures above 113°C. Interfacial reactions between Ag substrates and liquid In-49Sn (here and further compositions of alloys and phases are given in mass%, unless otherwise indicated) [2000Hua, 2001Chu] and 8Ag-In20-Sn [2002Chi] solders were studied in the temperature ranges from 200 to 350°C and from 225 to 325°C, respectively, by means of OM, SEM and EPMA. The formation of Ag2In and AgIn2 compounds was observed by [2000Hua, 2001Chu]. It was shown that the interfacial reaction products were a scallop-shaped Ag2In phase enveloped in an outer layer of AgIn2 [2001Chu]: both were identified as single crystals using polarized light under an optical microscope. During the aging of the In-49Sn/Ag soldered specimens at various temperatures ranging from 25 to 100°C, the Ag2In and AgIn2 intermetallic compounds compete for a dominant role [2001Chu]. At temperatures below 75°C, the AgIn2 expands in proportion to the composition of the Ag2In phase and the dominant reaction is: Ag2In + 3In 2AgIn2. At aging temperatures above 75°C, another type of solid/solid reaction becomes dominant: AgIn2 + 3Ag 2Ag2In. According to the data of [2002Chi], the intermetallic phase Ag2+x(In,Sn) appeared at the interface. Its chemical composition changed from 68.6Ag-22.8In-8.6Sn at 225°C to 69.6Ag-18.9In-11.5Sn at 325°C. The composition of the matrix of the as-cast solder alloy deviated from Sn-20In-2.8Ag to Sn-22.6In-1.2Ag due to the precipitation of Ag2In particles. DSC indicated a eutectic temperature of 187.9°C for this alloy, but according to the data of [1997Lee], the liquidus and solidus temperatures were reported to be 175 and 187°C respectively. Using DSC, XRD and SEM, [2002Yeh] studied an alloy of the same composition which had been homogenized at 95°C for 100 h. The solidus and liquidus temperatures were determined to be 170.8 and 195.5°C, respectively. [1994Woo] used XRD to establish that the solid solubility in the Ag-In-Sn system extends to alloys containing at least 10 at.% In and 2 at.% Ag. However, In decreases the solidus slightly and widens the melting range of the equivalent Ag-Sn alloy. There are two potential candidates (Sn-3.5Ag-5In and Sn-2.8Ag-20In) for Pb-free solders in this system [2002Liu]. It was found that the Sn-2.8Ag-20In alloy is composed of the (~ 90%), Ag2In and AgIn2 phases in the solid state, the Sn-3.5Ag-5In alloy is composed of the , (Sn) and phases in the temperature
MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–In–Sn
97
range from 150-200°C and of the (Sn) and the Ag2In phase below 150°. It should be noted that the formation of the liquid phase caused by segregation under non equilibrium conditions increases the freezing range. In particular, for the Sn-3.5Ag-5In alloy, a very narrow freezing range exists under equilibrium conditions, while a small fraction of liquid phase exists down to 115°C under non equilibrium conditions. Samples of about 8 at.% (In + Sn) in Ag, covering the range of 1-8 at.% In, were found to be single phase by metallography [1980Iga, 1988Sch] Measurement of thermodynamic properties was carried out by [1956Kle, 1987Gat, 2002Liu, 2003Zha]. Binary Systems The binary systems are accepted from recent thermodynamic calculations supported by additional experimental work as follows: Ag-In [2001Mos], Ag-Sn [1999Oht] and In-Sn [1996Lee]. The three phase diagrams are reproduced in [2002Liu] since they form the basis of that work. Solid Phases All crystallographic data are listed in Table 1. Ternary phases were not found in this system over the investigated temperature range [1998Kor]; however, the (hcp) phase forms a continuous series of solid solutions from the Ag-Sn side to the Ag-In side [2002Liu]. The mutual solubilities of In in the Ag3Sn and that of Sn in the Ag2In phase are very small. Invariant Equilibria The ternary invariant equilibria are given in Table 2 and Fig. 1 [2000Ohn, 2002Liu]. All solid state reactions are taken from the vertical sections presented in [2002Liu], supposedly of U type. According to the work of [1998Kor], a ternary eutectic melts in this system at 114°C (this temperature was reported also in the reviews of [2002Kat, 2003Oul]). The calculation gave two different possibilities for this eutectic reaction: L Ag2In + + or L + + , in which would transform to Ag2In by a solid state reaction at some temperature lower than the eutectic temperature. This latter option is not accepted here. The composition Sn-50In-2.5Ag (mass%) (Ag0.027In0.507Sn0.466) was placed close to the ternary eutectic [2003Oul], which is not far from the accepted value in Table 2. A number of typographical errors in the table given by [2002Liu] were corrected. For example, the temperature of 227°C given for U1 in [2002Liu] is impossible for this transition type reaction since it is higher than 221°C of the related Ag-Sn eutectic, e2. The correction was made by recalculating all the ternary and binary phase equilibria using the same thermodynamic parameters as [2002Liu], resulting in 217°C for U1. From this new calculation, the compositions of all other phases involved in the invariant reactions are given in Table 2; they are not given in [2002Liu]. Also confusion over the phases Ag2In and AgIn2 is corrected. It is noted that in this, otherwise very extensive and useful work on the ternary system, some other typographical errors are misleading concerning the references of the binary systems used: for Ag-Sn [1999Oht] (not 1995Ohtani) and for In-Sn [1996Lee] (not 1994Lee) is correct. Liquidus Surface Based on the optimized thermodynamic parameters, the calculated liquidus projection of the Ag-In-Sn system is shown in Fig. 2 [2000Ohn, 2002Liu, 2003Ohn], in which there are five transition type reactions and one eutectic reaction near the In-Sn edge. This diagram has been checked by recalculation as detailed in the previous section. Isothermal Sections The calculated isothermal sections are in good agreement with experimental data [2002Liu, 1998Kor]. Four examples of the calculated and optimized isothermal sections at 400, 250, 180 and 113°C are given in Figs. 3-6 [2002Liu]. The isothermal sections at 250°C and 113°C are not given in the paper by [2002Liu]
Landolt-Börnstein New Series IV/11C3
MSIT®
98
Ag–In–Sn
but are calculated here using the parameters of [2002Liu]. This was done in order to enable a quantitative comparison of the two isothermal sections given at these temperatures by [1998Kor]. The isothermal section at 113°C, Fig. 6, slightly below the ternary eutectic is in qualitative agreement with the diagram given by [1998Kor], and these authors confirm that 113°C is very close to the eutectic temperature. There is, however, some quantitative difference in the phase boundaries between these two different thermodynamic calculations of [2002Liu] and [1998Kor]. At 250°C, Fig. 4, the agreement between these authors is even better. Temperature – Composition Sections Very satisfactory agreement between the calculated vertical sections and the experimental data was obtained by [2002Liu]. Some of the temperature-composition sections are shown in Figs. 7a-7g [2000Ohn, 2002Liu]. The effect of alloying indium on the phase equilibria of the Sn-3.66Ag eutectic alloy is shown in Fig. 8 [2003Ohn]. The influence of In on the eutectic temperature in the Ag-Sn system (221°C) was calculated by [2002Liu] and is shown in Fig. 9. It is seen that the addition of In decreases the eutectic temperature, and the eutectic temperature reaches 200°C when the In content is 15 mass%. The effects of Ag on the liquidus and solidus of the Sn-20In alloy (Fig. 10a) and on the liquidus temperature of the Sn-51.7In eutectic alloy (Fig. 10b) have been studied by [2000Ohn]. It should be noted that many phase fields have wrong designations and some phase boundaries are missing in the vertical sections in [2002Liu]. This was corrected in the diagrams presented in the present evaluation. Thermodynamics The enthalpy change associated with the formation of the ternary Ag-In-Sn alloys from the binary Ag-Sn and In-Sn alloys was studied by quasibinary experiments in the Sn rich range [1956Kle] and the enthalpies of mixing in the liquid state were determined in a heat flow calorimeter at 980°C (Fig. 11) [1987Gat]. The enthalpies of mixing of the alloys in the sections Ag0.9Sn0.1-In0.2Sn0.8 and Ag0.9Sn0.1-In0.4Sn0.6 were determined at 893°C. The enthalpies of mixing were calculated from the data for the binary systems and found to be in good agreement with the experimental results. Based on the phase diagrams, the measured activities, the mass action law and the coexistence theory of metallic melts involving compound formation, models for the mass action concentrations for the Ag-In-Sn metallic melts have been formulated by [2003Zha]. The calculated results agree well with experimental values, showing that the models formulated can embody the structural characteristics of the corresponding melts. The calculated activities of Ag, Sn and In at 1150°C are shown in Fig. 12 [2002Liu, 2003Zha]. Notes on Materials Properties and Applications Ag-In-Sn alloys are among the most promising candidates for lead-free solders due to their mechanical strength and creep resistance [1997Hua, 1997Lee, 1998Kar, 1998Kor, 2002Kat, 2002Liu, 2002Yeh]. Ultimate tensile strength (UTS) and elongation of the Sn-3.5Ag-5In composition were reported in the review of [1997Hua]. The soldering temperature was given to be 190°C for the alloy Sn-2.8Ag-20In [1997Hua]. UTS and elongation of this alloy at a strain rate of 10–3 s–1 at room temperature were given as 59.3 MPa and 50.2%, respectively [2002Yeh]. UTS decreased but its elongation increased with increasing testing temperature. The activation energy for creep at a constant strain rate of 10–4 s–1 was estimated to be 71.4 kJ#mol–1 at a low testing temperature (< 50°C) and 51.0 kJ#mol–1 in the temperature range from 50 to 100°C [2002Yeh]. It was shown by [1998Kar] that In additions of up to 2% slightly decreased the fatigue life of Sn-3.5Ag alloy due to a loss in ductility. The tensile strength of Sn-3.5Ag-In increases but its ductility decreases with increasing In content. Microstructural observation suggests that In atoms are dissolved in Sn. Thus, the increase in the tensile strength through In addition is probably caused by a solid solution hardening mechanism. Furthermore, this alloy also contains coarse plate like precipitates of Ag3Sn. The loss of ductility in Sn-3.5Ag-In alloy should also be ascribed to the irregularly shaped second phases of Ag3Sn. MSIT®
Landolt-Börnstein New Series IV/11C3
Ag–In–Sn
99
The Sn-3.5Ag-In alloy has a relatively high fatigue resistance and the addition of In to less than 2% practically does not affect the fatigue life of the binary alloy [1998Kar]. The fatigue life of this alloy remains superior to that of a Sn-37Pb eutectic alloy, even after the addition of 5% In. Rapidly solidified powder of the composition Ag-6.0Sn-3.0In has commercial importance as an electrical contact material [1992Ver]. The melting point of this alloy is 835.6°C. These alloy powders can be prepared by gas atomization. Important physical properties of such powders (electrical conductivity, hardness, density and microstructure) were determined by [1992Ver]. The corrosion behavior of three Ag-In-Sn alloys in an aqueous NaCl (3%) solution has been studied using a standard three electrode cell [2003Oul]. Good corrosion resistance of the alloys (xIn/xSn = 0.12 and xIn/xSn = 0.2 and xAg = 0.03 and 0.1) was confirmed and it was shown that an increase in Ag addition induces an increase in the corrosion resistance while the In content has little effect. Miscellaneous The AgIn2 and Ag2In intermetallic compounds dominating at different aging temperatures could affect material properties, which should be taken into consideration in the application of In-49Sn/Ag solder joints [2001Chu]. The growth of the Ag2+x(In,Sn) intermetallic compound has been shown to be a diffusion-controlled process and the activation energy of its growth rate was calculated to be 41.6 kJ#mol–1 [2002Chi]. The oxidation behavior of the Ag-In-Sn alloys containing 8 at.% (In + Sn) was studied between 600 and 860°C in air and at 700°C in oxygen [1980Iga]. Silver increases both the surface tension and viscosity of the Sn-20In alloy [2000Ohn]. Density and surface tension of the (Sn-3.8Ag)eut + 10In liquid alloys have been measured using the maximum bubble-pressure and dilatometric methods in the temperature range from 160-917 and 242-930°C, respectively [2002Liu]. It was shown that the addition of In slightly increases the density in comparison with the binary-eutectic Ag-Sn at 250°C, with the opposite tendency at 960°C. Structure, electrical resistivity, wettability, melting point, internal friction and elastic moduli of the 2.8Ag-25In-72.2Sn (I) and 2.5Ag-20In-77.5Sn (II) rapidly solidified Pb-free solder alloys have been investigated by [2004Elb]. According to XRD data, these alloys contain a tetragonal Sn phase, tetragonal In3Sn, orthorhombic Ag3Sn, cubic Ag9In4 and a hexagonal InSn4 phase. Youngs modulus, bulk modulus and shear modulus of the I and II alloys are 25.739 and 21.446 GPa, 35.541 and 28.571 GPa, 9.33 and 7.799 GPa, respectively. These alloys have lower melting points (173.11 and 162.01°C for the I and II alloy, respectively), electrical resistivity and higher elastic modulus than commercial Pb-Sn alloys. In addition, these alloys have good wetting behavior [2004Elb]. References [1956Kle] [1980Iga]
[1987Gat]
[1988Sch]
[1992Ver]
Landolt-Börnstein New Series IV/11C3
Kleppa, O.J., “A Calorimetric Investigation of Some Binary and Ternary Liquid Alloys Rich in Tin”, J. Phys. Chem., 60, 842 (1956) (Thermodyn., Experimental, 15) Igarashi, T., Shibata, M., Kodama Y., “The Precipitation Behavior of Oxide During Internal Oxidation of Ag-In-Sn Alloys” (in Japanese), Nippon Kinzoku Gakkai Shi, 44(4), 378-386 (1980) (Crys. Structure, Experimental, 30) Gather, B., Schroeter, P., Blachnik, R., “The Enthalpies of Mixing in the Liquid State on the Ternary Systems Ag-In-Sn, Ag-Sn-Sb, Ag-In-Sb and In-Pb-Sb”, Z. Metallkd., 78, 280-285 (1987) (Thermodyn., Experimental, 20) Schmid-Fetzer, R., “Silver - Indium - Tin”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.14858.1.20, (1988) (Crys. Structure, Phase Diagram, Assessment, 1) Verma, A., Anantharaman, T.R., “Internal Oxidation of Rapidly Solidified Silver-Tin-Indium Alloy Powders”, J. Mater. Sci., 27, 5623-5628 (1992) (Experimental, Electr. Prop., Mechan. Prop., Morphology, Phys. Prop., 13)
MSIT®
100 [1994Woo] [1996Lee]
[1997Hua] [1997Lee] [1998Kar]
[1998Kor]
[1999Oht]
[2000Hua]
[2000Ohn]
[2001Mos]
[2001Chu] [2001Sug] [2002Chi]
[2002Kat] [2002Liu]
[2002Yeh] [2003Ohn]
[2003Oul]
MSIT®
Ag–In–Sn Wood, E.P., Nimmo, K.L., “In Search of New Lead-Free Electronic Solders”, J. Electron. Mater., 23(8), 709-713 (1994) (Phase Diagram, Experimental, 9) Lee, B.J., Oh, C.S., Shim, J.H., “Thermodynamic Assessments of the Sn-In and Sn-Bi Binary Systems”, J. Electron. Mater. 25(6), 983-991 (1996) (Calculation, Phase Diagram, Phase Relations, Thermodyn., #, 50) Hua, F., Glazer, J., “Lead-Free Solders for Electronic Assembly”, Des. Reliab. Solders Solder Interconnect., Proc. Symp., 65-73 (1997) (Review, 63) Lee, N.-C., “Getting Ready for Lead-Free Solders” (in German), Soldering Surf. Mount. Technol., 9(2), 65-69 (1997) (Mechan. Prop., Review, 0) Kariya, Y., Otsuka, M., “Mechanical Fatigue Characteristics of Sn-3.5Ag-X (X = Bi, Cu, Zn and In) Solder Alloys”, J. Electron. Mater., 27(11), 1229-1235 (1998) (Experimental, Mechan. Prop., 14) Korhonen, T.-M., Kivilahti, J.K., “Thermodynamics of the Sn-In-Ag Solder System”, J. Electron. Mater., 27(3), 149-158 (1998) (Calculation, Phase Diagram, Thermodyn., #, *, 34) Ohtani, H., Miyashita, M., Ishida, K., “Thermodynamic Study of Phase Equilibria in the Sn-Ag-Zn System”, J. Japan. Inst. Metals, 63(6), 685-694 (1999) (Calculation, Phase Diagram, Phase Relations, Thermodyn., Experimental, #, 68) Huang, Y.T., Chuang, T.H., “Interfacial Reactions between Liquid In-49Sn Solders and Ag Substrates”, Z. Metallkd., 91(12), 1002-1005 (2000) (Kinetics, Phase Relations, Morphology, Experimental, 12) Ohnuma, I., Cui, Y., Liu, X.J., Inohana, Y., Ishihara, S., Ohtani, H., Kainuma, R., Ishida, K., “Phase Equilibria of Sn-In Based Micro-Soldering Alloys”, J. Electron. Mater., 29(10), 1113-1121 (2000) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 27) Moser, Z., Gasior, W., Pstrus, J., Zakulski, W., Ohnuma, I., Liu, X.J., Inohana, Y., Ishida, K., “Studies of the Ag-In Phase Diagram and Surface Tension Measurements”, J. Electron. Mater. 30(9), 1120-1128 (2001) (Calculation, Phase Diagram, Phase Relations, Thermodyn., Experimental, #, 27) Chuang. T.H., Huang, Y.T., Tsao, L.C., “AgIn2/Ag2In Transformation in an In-49Sn/Ag Soldered Joint under Thermal Aging”, J. Electron. Mater., 30(8), 945-950 (2001) Suganuma, K., “Advances in Lead-Free Electronics Soldering”, Curr. Opin. Solid State Mater. Sci., 5, 55-64 (2001) (Review, Mechan. Prop., 75) Chiang, M.I., Chuang, T.H., “Interfacial Reaction Between Liquid Sn-20In-2.8Ag Solder and Ag Substrate”, Z. Metallkd., 93(12), 1194-1198 (2002) (Kinetics, Phase Relations, Morphology, Experimental, 15) Kattner, U.R., “Phase Diagrams for Lead-Free Solder Alloys”, JOM, 12, 45-51 (2002) (Phase Diagram, Review, 47) Liu, X.J., Inohana, Y., Takaku, Y., Ohnuma, I., Kainuma, R., Ishida, K., Moser, Z., Gasior, W., Pstrus, J., “Experimental Determination and Thermodynamic Calculation of the Phase Equilidria and Surface Tension in the Sn-Ag-In System”, J. Electron. Mater., 31(11), 1139-1151 (2002) (Experimental, Assessment, Interface Phenomena, Phase Diagram, Phase Relations, Phys. Prop., Thermodyn., #, *, 28) Yeh, M.S., “Evaluation of the Mechanical Properties of a Ternary Sn-20In-2.8Ag Solder”, J. Electron. Mater., 31(9), 953-956 (2002) (Mechan. Prop., Experimental, 9) Ohnuma, I., Miyashita, M., Liu, X.J., Ohtani, H., Ishida, K., “Phase Equilibria and Thermodynamic Properties of Sn-Ag Based Pb-Free Solder Alloys”, IEEE Trans., Electron. Pack. Manuf., 26(1), 84-89 (2003) (Experimental, Assessment, Phase Diagram, Thermodyn., 21) Oulfajrite, H., Sabbar, A., Boulghallat, M., Jouaiti, A., Lbibb, R., Zrineh, A., “Electrochemical Behavior of a New Solder Material (Sn-In-Ag)”, Mater. Let., 57, 4368-4371 (2003) (Electrochemistry, Experimental, 13)
Landolt-Börnstein New Series IV/11C3
Ag–In–Sn [2003Zha] [2004Elb]
101
Zhang, J., “Calculating Models of Mass Action Concentration for Metallic Melts Ag-In-Sn”, Calphad, 27(1), 9-17 (2003) (Calculation, Thermodyn., 13) El-Bediwi, A., El-Bahay, M.M., “Influence of Silver on Structural, Electrical, Mechanical and Soldering Properties of Tin-Indium Based Alloys”, Rad. Eff. Defects Solids, 159, 133-140 (2004) (Crys. Structure, Electr. Prop., Experimental, 18)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
at 25°C [Mas2], dissolves up to 21 at.% In and up to 11.5 at.% Sn
(In) < 156.634
tI2 I4/mmm In
a = 325.3 c = 494.7
at 25°C [Mas2], dissolves up to 10 at.% Sn
( Sn)
tI2 ? Sn
a = 370 c = 337
at 25°C, 9.0 GPa [Mas2]
(Sn) 231.9681 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
at 25°C [Mas2]
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
[Mas2]
, (Ag-In) 695 - 660
cP2 Pm3m CsCl
25-30 at.% In [Mas2] Not considered in [2001Mos]
´, (Ag-In) < 187
cP* Pm3m ?
25 at.% In [Mas2] Not considered in [2001Mos]
,
hP2 P63/mmc Mg
[1998Kor, Mas2, V-C2]
Ag7In3 < 675
a = 296.1 0.2 c = 477.8 0.4
21-36 at.% In [2001Mos] [V-C2]
Ag4Sn < 722
a = 296.58 c = 478.24
10-23 at.% Sn [1999Oht] [V-C2]
a = 992.2 0.4
32 at.% In [2001Mos] 31-33.5 at.% In [Mas2] [V-C2]
Ag2In < 300
cP52 P43m Cu9Al4
AgIn2 < 166
tI12 I4/mcm CuAl2
Landolt-Börnstein New Series IV/11C3
a = 688.1 0.4 c = 562.0 0.4
67 at.% In [2001Mos] 66.7 at.% In [Mas2] [V-C2]
MSIT®
Ag–In–Sn
102 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
J, Ag3Sn < 477
oP8 Pmmn Cu3Ti
Lattice Parameters Comments/References [pm]
a = 596.8 0.9 b = 478.02 0.4 c = 518.43 0.9
, In3Sn < 142
tI2 I4/mmm In
, InSn4 or In2Sn9 < 223
hP1 P6/mmm BiIn
25 at.% Sn [1999Oht] 23.7-25 at.% Sn [Mas2] [V-C2]
12.5-44.5 at.% Sn [1996Lee] [V-C2]
a = 346.3 c = 440.4 a = 320.8 c = 299.7 a = 321.77 0.01 c = 299.88 0.01
76-97.5 at.% Sn [1996Lee] for In2Sn9 at –10°C [V-C2] for InSn4 at 49°C [V-C2]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Ag
In
Sn
L + Ag3Sn (Sn) +
217
U1
L Ag3Sn (Sn)
4.22 75 0.07 73.31
1.95 0.54 7.96
93.83 25 99.39 18.73
L + (Sn) +
209
U2
L (Sn)
3.64 0.06 70.98
6.94 2.01 3.74 19.35
89.42 97.93 96.26 9.67
L + + Ag2In
180
U3
L Ag2In
2.69 67.8 68
26.10 28.56 12.76 32
71.21 3.64 87.24
+ (Sn) + Ag2In
~ 139
U4 (?)
-
-
-
-
L + (In) + AgIn2
135
U5
L (In) AgIn2
1.45 33
84.55 88.26 87.93 67
14.0 11.74 12.61 -
L + AgIn2 + Ag2In
119
U6
L AgIn2 Ag2In
1.1 33 68
61.81 67 64.79 32
37.09 35.21
L + + Ag2In
114
E
L Ag2In
0.94 68
52.97 56.09 23.35 32
46.09 43.91 76.65 -
Ag2In + + AgIn2
~ 104
U7 (?)
-
-
-
-
+ Ag2In AgIn2 + (Sn)
~ 23.6
U8 (?)
-
-
-
-
MSIT®
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Ag-In
Ag-In-Sn
A-B-C
In-Sn
Ag-Sn 722.4 p1 l + (Ag) ζ
675.4 p2 (Ag) + l ζ
476.6 p3 l+ζε 223.5 p4 l + (βSn) γ 217
L + ε (βSn) + ζ
ε+(βSn)+ζ
194.8 e2 ζ Ag2In + l
209
L + (βSn) γ + ζ
L + ζ γ + Ag2In
180
U1 U2
U3
139 135 119
L + (In) β + AgIn2
L + Ag2In β + AgIn2
L β + γ + Ag2In
117.9 e4 lβ+γ
E
U7 23.6 Ag2In + γ (βSn) + AgIn2 (βSn)+γ+AgIn2
U8
(βSn)+Ag2In+AgIn2
103
MSIT®
Fig. 1: Ag-In-Sn. Reaction scheme
(βSn)+ζ+Ag2In
(In)+β+AgIn2
Ag2In + β γ + AgIn2
(In)+β+AgIn2
U5
U4
U6
114 104
γ + ζ (βSn) + Ag2In
141.6 p6 l + (In) β
Ag–In–Sn
167.8 p5 l + Ag2In AgIn2 144.8 e3 lAgIn2+(In)
220.5 e1 l (βSn) + ε
Ag–In–Sn
104
Sn e1
Fig. 2: Ag-In-Sn. Calculated liquidus surface projection
Data / Grid: at.% Axes: at.%
(β Sn) p4
U1
U2 20
80
Ag3Sn
γ U3
300
40
60
400 e4 E Ag2In
p3 60
ζ
40
U6
500 p1 80
600
β
700
20
U5
800
p6
900°C (Ag)
(In) 20
Ag
p2
40
60
80
e2 p5
Sn
AgIn2
e3
In
Data / Grid: at.% Axes: at.%
Fig. 3: Ag-In-Sn. Calculated isothermal section at 400°C 20
80
40
60
60
40
L Ag3Sn 80
20
ζ (Ag)
Ag
MSIT®
20
40
60
80
In
Landolt-Börnstein New Series IV/11C3
Ag–In–Sn
105
Sn
Data / Grid: at.% Axes: at.%
Fig. 4: Ag-In-Sn. Calculated isothermal section at 250°C L+Ag3Sn
20
80
40
ζ+L+Ag3Sn
60
L
ζ+L
60
40
Ag3Sn 80
20
ζ (Ag) 20
Ag
Ag2In
40
60
Sn Fig. 5: Ag-In-Sn. Calculated isothermal section at 180°C
80
In
Data / Grid: at.%
(β Sn) (β Sn)+γ +ζ
Axes: at.%
γ
20
80
L+ζ+γ 40
60
(β Sn)+Ag3Sn+ζ 60
40
Ag3Sn 80
20
ζ L
(Ag)
Ag
Landolt-Börnstein New Series IV/11C3
20
40
Ag2InAg In+ζ+γ 2
60
80
In
MSIT®
Ag–In–Sn
106
Sn
Data / Grid: at.%
(β Sn)
Fig. 6: Ag-In-Sn. Calculated isothermal section at 113 °C
Axes: at.%
(β Sn)+Ag2In+γ
γ 20
Ag 2 In+ γ
80
40
(βS n)+ ζ+A gS 3 n
Ag2In+γ +β 40
Sn)
60
60
β
ζ+(β
Ag3Sn
20
Ag I 2 n+
80
ζ
Ag2In+AgIn2+β
(Ag) 20
Ag
Temperature, °C
Fig. 7a: Ag-In-Sn. The calculated vertical section at 10 mass% Ag, plotted in at.%
Ag2In
40
60
AgIn2+(In)
(In)
80
AgIn2
In
400
L
300
L+ζ
L+Ag3Sn+ζ L+(β Sn)+ζ 209
200
L+AgIn2
180
L+Ag2In L+Ag2In+γ
119 135 100
(In)+AgIn2
0
Ag 10.58 In 89.42 Sn 0.00
MSIT®
AgIn2+β +(In)
AgIn2+β
114 Ag2In+AgIn2+β AgIn2+β
20
Ag2In+β +γ
Ag2In+AgIn2+γ AgIn2+γ+β (β Sn)+AgIn2+γ 40
Sn, at.%
ζ +γ (β Sn)+ Ag3Sn+ζ
~139
Ag2In+γ
(β Sn)+ Ag2In+ζ
~104
(β Sn)+ Ag2In+γ
~23.6°C
60
217
80
(β Sn)+Ag2In+AgIn2
Ag 10.90 0.00 In Sn 89.10
Landolt-Börnstein New Series IV/11C3
Ag–In–Sn
Temperature, °C
Fig.7b: Ag-In-Sn. The calculated vertical section at 20 mass% Ag, plotted in at.%
107
400
L L+ζ 300
L+Ag3Sn+ζ (β Sn)+ζ L+ζ +γ
200
L+AgIn2
180 L+Ag2In+AgIn2
135 100
L+Ag2In L+Ag2In+γ 114°C
119 β +AgIn2 ~104
~139
(β Sn)+ Ag2In+ζ (β Sn)+ Ag2In+γ
Ag2In+AgIn2+(β Sn)
0
Ag 21.02 In 78.98 Sn 0.00
20
(β Sn)+ Ag3Sn+ζ
Ag2In+γ
Ag2In+AgIn2+γ
~23.6
Temperature, °C
ζ +γ
Ag2In+AgIn2+β
AgIn2+γ+β
40
Ag 21.58 0.00 In Sn 78.42
60
AgIn2+(β Sn)+γ
Sn, at.%
500
L
400
L+ζ L+ζ +Ag3Sn 300
(β Sn)+ζ +Ag3Sn 200
L+Ag2In
L+Ag2In+AgIn2 135 100
AgIn2+γ+(In) 0
Ag 31.33 In 68.67 Sn 0.00
Landolt-Börnstein New Series IV/11C3
217
Ag2In+γ+β
(In)+AgIn2
Fig. 7c: Ag-In-Sn. The calculated vertical section at 30 mass% Ag, plotted in at.%
209
119°C 104
114 Ag2In+β +AgIn2
AgIn2+γ+β
ζ +γ
180 L+Ag2In+γ
139
209
217
(Sn)+ζ
Ag2In+γ+β
Ag2In+γ+AgIn2 (β Sn)+Ag2In+γ ~23.6 (β Sn)+Ag2In+AgIn2 20
40
Sn, at.%
60
(β Sn)+Ag2In+ζ
Ag 32.05 0.00 In Sn 67.95
MSIT®
Ag–In–Sn
108
Fig. 7d: Ag-In-Sn. The calculated vertical section at 40 mass% Ag, plotted in at.%
500
L
400
Temperature, °C
L+ζ 300
200
L+Ag2In L+Ag2In+AgIn2 104 Ag2In+AgIn2+β 23.6 Ag2In+AgIn2+(β Sn)
Ag 41.51 In 58.49 Sn 0.00
Temperature, °C
Ag2In+γ+β Ag2In+γ+(β Sn) 20
40
Sn, at.%
750
217
(Sn)+ζ 139°C
Ag2In+γ 114
Ag2In+AgIn2+γ 0
209
ζ +γ
180
119
Ag 42.32 0.00 In Sn 57.68
Ag2In+ζ +(β Sn)
(Ag)+L L
500
ζ L+ζ
250
(Ag)+ζ
ζ +γ (β Sn)+ζ +γ (β Sn)+ζ
0
Ag 80.98 In 19.02 Sn 0.00
MSIT®
(Sn)+ζ +Ag3Sn
L+ζ +γ
100
Fig. 7e: Ag-In-Sn. The calculated vertical section at 20 mass% In, plotted in at.%
L+ζ +Ag3Sn
γ+Ag2In+L γ+ζ +L γ+Ag2In+L
139
γ+Ag2In+AgIn2
ζ +(β Sn)+Ag2In 20
180°C Ag2In+γ
~23.6 40
60
Ag2In+γ+(β Sn)
Sn, at.%
Ag 0.00 In 20.54 Sn 79.46
Landolt-Börnstein New Series IV/11C3
Ag–In–Sn
Fig. 7f: Ag-In-Sn. The calculated vertical section at 40 mass% In, plotted in at.%
109
700
600
L
Temperature, °C
500
400
L+ζ 300
200
Ag2In+AgIn2+L
Ag2In+L+γ
Ag2In+L
119 100
114°C
0
Ag 61.49 In 38.51 Sn 0.00
Fig. 7g: Ag-In-Sn. The calculated vertical section at 80 mass% Sn, plotted in at.%
20
AgIn2+β +γ
Ag 0.00 In 40.80 Sn 59.20
40
Sn, at.%
400
L L+Ag3Sn 300
Temperature, °C
104
Ag2In+γ+β Ag2In+AgIn2+γ
Ag2In+AgIn2+β
L+Ag3Sn+ζ L+ζ 217
200
209°C
γ+L 180 L+Ag2In+γ
γ+L+ζ
ζ +γ 139
Ag2In+γ
ζ +(β Sn)
100
γ+(β Sn)+Ag2In Ag 0.00 In 20.54 Sn 79.46
Landolt-Börnstein New Series IV/11C3
10
Ag, at.%
ζ +(β Sn)+ +Ag3Sn 20
Ag 21.58 0.00 In Sn 78.42
MSIT®
Ag–In–Sn
110
(β Sn)+L+Ag3Sn
L
217 L+ζ
209
200
Temperature, °C
Fig. 8: Ag-In-Sn. Effect of In on the phase equilibria of the Sn-3.66Ag (mass%) eutectic alloys
ζ +γ
(β Sn)+ζ
180°C
γ+Ag2In+ζ (β Sn)+ζ +Ag3Sn
139
γ+Ag2In 100
γ+Ag2In+AgIn2
(β Sn)+Ag2In+γ (β Sn)+Ag2In
Ag 4.01 0.00 In Sn 95.99
3.99 In 20.46 Sn 75.55
10
20 Ag
In, at.%
Ag 10.00 Sn 90.10
Ag, mass% 2
Fig. 9: Ag-In-Sn. Effect of In on the eutectic temperature in Ag-Sn alloys
4
6
8
300
Ag-Sn binary system
Temperature, °C
2 mass% In 5 mass% In 10 mass% In 15 mass% In 221°C
200
Sn
2
4
6
Ag, at.%
MSIT®
8
10
Ag 10.90 Sn 98.10
Landolt-Börnstein New Series IV/11C3
Ag–In–Sn
Temperature, °C
Fig. 10a:Ag-In-Sn. Effects of Ag on the liquidus and solidus of Sn-20In (mass%) alloy
111
300
L
Ag(Liquidus) 200
Ag(Solidus)
100
Ag 0.00 In 20.54 Sn 79.46
Temperature, °C
Fig. 10b:Ag-In-Sn. Effects of Ag on the liquidus temperature of Sn-51.7In (mass%) eutectic alloy
2
4
Ag 5.44 In 20.44 Sn 74.13
4
Ag 5.38 In 52.27 Sn 42.35
Ag, at.%
300
L
200
100
Ag 0.00 In 52.53 Sn 47.47
Landolt-Börnstein New Series IV/11C3
2
Ag, at.%
MSIT®
Ag–In–Sn
112
Sn
Data / Grid: at.% Axes: at.%
Fig. 11: Ag-In-Sn. Enthalpies of mixing in the liquid state at 980°C 20
80
0
40
60
-500 -1000 60
-1500 -2000
40
-2500 -3000 80
20
-3500 -4000 -4500 20
Ag
40
60
80
Sn Fig. 12: Ag-In-Sn. Calculated activities of Ag, In and Sn at 1150°C
In
Data / Grid: at.% Axes: at.%
activity of Sn
0.9
activity of Ag 20
activity of In
80
0.8 0.1 0.7
40
0.6 0.2
60
60
0.5 0.4
0.3
40
0.3 0.4 80
0.2 20
0.5 0.1
0.6 0.7
0.1
0.8 0.9
Ag
MSIT®
20
0.2
0.3 0.4 40
0.6
0.5 60
0.7
0.8 80
0.9
In
Landolt-Börnstein New Series IV/11C3
Ag–Pb–Sn
113
Silver – Lead – Tin Qingsheng Ran, Hans Leo Lukas, updated by Volodymyr Ivanchenko Introduction Based on thermal analysis and metallography of 102 alloys, [1911Par] reported seven vertical sections, the liquidus surface and two invariant equilibria. The results were later discussed by the same author again [1913Par]. Because the binary systems adopted are too old, the data are not very reliable. [1946Ear] investigated the liquidus surface of the Pb corner and the ternary eutectic by thermal analysis, micrographic observation, segregation study of slowly cooled alloys and small-scale melting experiments of a total of 59 alloys. [1933Gar] reported liquidus and solidus curves along the section Pb-Ag3Sn with up to 7 mass% Ag3Sn and put into evidence the formation of a solid solution between Pb and Ag3Sn; however, the points scatter markedly. [1944Rus] determined the solidification ranges of three Ag-Pb-Sn alloys by thermal analysis. A review of the ternary system was presented by [1988Ran]. [1987Tar] calculated the eutectic temperature depression in the Sn-Pb system coarsened by addition of silver as –3.3°C/mass% Ag. This value is in excellent agreement with their own experimental values, which ranged from –3.3 to –2.6°C, and with value of –3.7°C/mass% Ag derived from [1946Ear]. [1996Fri] studied the changes occurring in a eutectic Ag-Sn solder due to addition of lead. It was shown, that already at very small lead contents (< 0.5 at.%) a ternary eutectic reaction takes place at 178°C. [2001Kat] calculated isopleth (Sn-3.5 mass% Ag)-(Sn-37 mass% Pb). [2003Zen] presented calculated liquidus surface projection as well as vertical section of the Ag-Sn-Pb phase diagram from the binary Ag-Sn eutectic composition to pure Pb. By means of emf, using the cell Sn(l), SnO2(s)/calcia stabilized zirconia/Sn(alloy), SnO2(s), [1973Jag] studied the partial free energy of mixing of Sn in the molten Ag-Pb-Sn system at 1073 K, along the concentration lines with xAg/xPb ratios of 3/7.1 and 7/3. On each line nine alloys were examined. Using the results of [1973Jag, 1978Cho] and [1983Lee, 1984Lee] derived thermodynamic descriptions of the liquid phase. [1992Lee] optimized the available thermodynamic values and calculated activities of Ag and Pb in the Ag-Sn-Pb system at 800°C. The morphology of eutectic in Sn-36Pb-2Ag (mass%) alloy was discussed by [2002Nog]. [1979Zah] studied 6 alloys that were positioned near the corners and in the sides of the following composition triangle: 20Ag20Sn60Pb - 20Ag60Sn20Pb - 60Ag20Sn20Pb (at.%). The density, surface tension and their temperature coefficients were examined in the temperature interval from liquidus to 800°C. [1992Sca] investigated the effects of accelerated fatigue tests on crack formation in 95.5Pb-2Sn-2.5Ag (mass%) solder joints. [1998Vei] presented the results of the experimental determination of the viscoelastic tensile and shear properties of the eutectic 62Sn-36Pb-2Ag (mass%) solder alloy in the temperature range between –20 and 100°C. The aims of [2001Bra, 2002Ama, 2004Isl] are concerned with the development of lead-free solders, the 62Sn-36Pb-2Ag (mass%) solder was used for comparative experiments. The characteristics of melting and wetting were studied by [2001Bra]. Strain rate as a function of steady state stress for 62Sn-36Pb-2Ag solder was presented by [2002Ama]. [2004Isl] presented the interfacial reactions of 62Sn-36Pb-2Ag solder with Ni. The shear test of 62Sn-36Pb-2Ag balls bonding placed on copper pad and cooled at 10 K#min–1 to 200 K#min–1 has been performed by [2002Nak]. [1998Sch] presented experimental data for Cu-Sn intermetallic growth between copper and 62Sn-36Pb-2Ag solder. The morphology of eutectic in Sn-36Pb-2Ag alloy was discussed by [2002Nog]. Binary Systems The accepted Ag-Pb phase diagram is taken from [2002Luk]. The Ag-Sn binary system is taken from [1987Kar]. The Pb-Sn phase diagram is taken from [1988Kar].
Landolt-Börnstein New Series IV/11C3
MSIT®
114
Ag–Pb–Sn
Solid Phases Solid phases observed in this system are given in Table 1. Quasibinary Systems [1911Par] reported an invariant equilibrium L (Pb) + Ag3Sn with maximal solidus temperature of about 300°C. A ternary maximum at 310°C was reported by [1946Ear] too. These results are doubtful since the Ag-Sn system contains the Ag4Sn phase and two U type nonvariant equilibria must be realized in the Pb corner of the phase diagram. Invariant Equilibria A ternary eutectic is reported by [1946Ear]. The only known values are the temperature and equilibrium composition of the liquid, which are taken from [1946Ear]. They are presented in Table 2. Two U reactions, in which Ag4Sn is involved, are also qualitatively presented in Table 2. Liquidus Surface [2003Zen] calculated liquidus surface projection of the Ag-Pb-Sn phase diagram using CALPHAD technique. It is presented in Fig. 1 and the reaction scheme is shown in Fig. 2. The triangle drawn by dashed lines shows the calculated compositions of the three solid phases that are formed in the eutectic reaction L (Sn) + (Pb) + Ag3Sn. Temperature – Composition Sections The vertical sections 96.2Sn3.8Ag - 74.79Sn25.21Pb and 96.2Sn3.8Ag - Pb were calculated by [2001Kat] and [2003Zen], respectively. They are presented in Fig. 3 and Fig. 4. Thermodynamics The activities of Ag, Pb and Sn in the liquid alloys at 800°C are shown in Fig. 5. Activities of tin have been experimentally determined by [1973Jag], using the following cell: Sn(l),SnO2(s)/Calcia stabilized zirconia/Sn(alloy),SnO2(s) Activities of lead and silver have been then obtained by Gibbs-Duhem integration. These results have been used by [1983Lee, 1984Lee, 1992Lee] for modelling the liquid alloy. Experimental data at 800°C could be excellently represented by the Redlih-Kister Muggianu model including the ternary interaction parameters as follows: Gxs/J#mol–1 = – 2929xAgxSnxPb + xAgxSnxPb (47680xAg + 10137xSn + 6296xPb) Notes on Materials Properties and Applications Soldering is the most important method for joining mechanical components in electronics. The 62Sn-36Pb-2Ag (mass%) eutectic alloy is widely used as a solder and it is one of the major bonding materials in electronic devices. [1979Zah] showed that temperature dependences of density of Ag-Pb-Sn alloys are linear. The isobaric coefficients of thermal expansion are minimal ( 8#10–5°C–1 and 9#10–5°C–1) at Ag contents of ~ 55 and ~ 20 at.% and Pb contents of 20 and 55 at.%, respectively. The higher Ag content leads to the higher surface tension coefficient, and vice versa, the higher Pb content leads to the lower one. For example, the 20Sn-60Ag-20Pb (at.%) alloy coefficient of surface tension is equal to ) = 527 mJ#m–2 and for the 20Sn-20Ag-60Pb (at.%) alloy the coefficient of surface tension is equal to 415 mJ#m–2 at 750°C. The investigation of the effect of accelerated fatigue tests on crack formation in the 62Sn-36Pb-2Ag (mass%) solder joint shows that the crack occurs at a critical number of cycles when a Sn-depleted region is formed, yielding weaker inner layers with lower shear strength. The increasing of mechanical stress leads to increasing in tin depletion. This depletion is not homogeneously distributed along the surface. MSIT®
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The apparently spontaneous fractures take place when the crack can propagate along Sn depleted inner regions which are continuously extended from one end of alloy to another [1992Sca]. The experimental results of [1998Vei] show that the temperature dependences of the dynamic Young’s modulus and shear modulus of the eutectic 62Sn-36Pb-2Ag (mass%) solder alloy in the temperature range between –2 and 100°C can be approximated by linear functions: E(T)/GPa = – 0.055T + 41.90; G(T)/GPa) = – 0.022T + 15.16. The dynamic Young’s and shear moduli decrease by approximately 6% and the viscous damping coefficient increases by approximately 80% for a temperature rise from room temperature to 60°C. The eutectic 62Sn-36Pb-2Ag (mass%) solder alloy is used in a ball grid array technique, that is a bonding process which solder balls are arranged on Cu pads in a print board and heated in a reflow instrument after setting IC chips on them. This process makes semiconductor devices higher integrated, smaller, thinner and lighter than conventional ones. It must be noted, that there is tremendous interest presently in Pb-free solder assembly in the surface mount assembly industry in response to recent Japanese and European initiatives and proposed governmental restrictions regarding Pb usage and disposal. References [1911Par] [1913Par] [1933Gar] [1944Rus] [1946Ear] [1973Jag] [1978Cho]
[1979Zah]
[1983Lee]
[1984Lee] [1987Kar] [1987Tar]
[1988Kar] [1988Ran]
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Parravano, N., “The Ternary Silver - Tin - Lead System” (in German), Int. Z. Metallogr., 1, 89-108 (1911) (Phase Relations, Experimental, 7) Parravano, N., “The Ternary Silver - Tin - Lead System” (in German), Int. Z. Metallogr., 3, 15-22 (1913) (Phase Relations, Theory, 4) Garre, B., Vollmert, F., “Härtbare Bleilegierungen”, Z. Anorg. Allg. Chemie, 210, 77-80 (1933) (Phase Relations, Experimental, 1) Russell, J.B., Mack, J.O., “Substitute Solders of the 15-85 Tin - Lead Type”, Metals Technology, October (1944) (Experimental, 25) Earle, L.G., “Some Data Concerning Lead Alloys Containing Tin (up to 63%) and Silver (up to 3%)”, J. Inst. Metals, 72, 403-413 (1946) (Phase Relations, Experimental, #, 9) Jagannathan, K.P., Ghosh, A., “A Thermodynamic Study of the Molten Pb-Sn-Ag System”, Metall. Trans., 4, 1577-1583 (1973) (Thermodyn., Experimental, 23) Chou, K.-C., “Thermodynamics for Ternary and Multicomponent Systems”, Sci. Sinica, 11, 73-86 (1978), translated from Acta Metall. Sinica 12, 232-244 (1976) (Thermodyn., Theory, 15) Zaharova, T.V., Popel, S.I., Gavrilova, A.V., “Densities and Surface Tensions of Pb-Sn-Ag Melts” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., 1, 86-90 (1979) (Experimental, Interface Phenomena, Thermodyn., 5) Lee, D.N., Lee, S.K., “Calculation of the Partial Molar Properties of Ternary Alloys from Binary Data Using the Shortest Distance Composition Path”, Scr. Metall., 17(7), 861-866 (1983) (Calculation, Thermodyn., 17) Lee, H.M., Lee, S.K., Lee, D.N., “Optimization of Ternary Thermodynamic Data”, J. Korean Inst. Metals, 22, 398-407 (1984) (Thermodyn., Theory, 30) Karakaya, I., Thompson, W.T., “Ag-Sn (Silver-Tin)”, Bull. Alloy Phase Diagrams, 8(4) 340-347 (1987) (Phase Diagram, Phase Relations, Crys. Structure, Assessment, #, 51) Tarby, S.K., Notis, M.R., “Effect of Small Additions of Silver on the Eutectic Temperature in the Lead-Tin System”, Metall. Trans. B, 17B, 829-832 (1987) (Experimental, Thermodyn., 16) Karakaya, I., Thompson, W.T., “Pb-Sn (Lead-Tin)”, Bull. Alloy Phase Diagrams, 9(2), 200-202 (1988) (Phase Diagram, Crys. Structure, Assessment, #, 61) Ran, Q., Lukas, H.L., “Silver-Lead-Tin” MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH,
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[1992Lee] [1992Sca]
[1996Fri]
[1998Sch]
[1998Vei] [2001Bra]
[2001Kat]
[2002Ama]
[2002Luk]
[2002Nak]
[2002Nog]
[2003Zen]
[2004Isl]
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Ag–Pb–Sn Stuttgart; Document ID: 10.11471.1.20, (1988) (Crys. Structure, Phase Diagram, Assessment, 9) Lee, H.M., “Interaction Parameters of Lead-Base Ternary Systems”, Calphad, 16(1), 47-52 (1992) (Experimental, Thermodyn., 21) Scandurra, A., Licciardello, A., Torrisi, A., La Mantia, A., Puglisi, O., “Fatigue Failure in Pb-Sn-Ag Alloy During Plastic Deformation: A 3D-SIMS Imaging Study”, J. Mater. Res., 7(9), 2395-2402 (1992) (Experimental, Mechan. Prop., 19) Fritz, K., Schmitt-Thomas, Kh.G., “Lead Influence on the Eutectic SnAg3.5” (in German), Metall, 50(10), 633-635 (1996) (Experimental, Phase Diagram, Phase Relations, Morphology, #, 6) Schaefer, M., Fournelle, R.A., Liang, J., “Theory for Intermetallic Phase Growth Between Cu and Liquid Sn-Pb Solder Based on Grain Boundary Diffusion Control”, J. Electron. Mater., 27(11), 1167-1176 (1998) (Experimental, Morphology, Theory, Transport Phenomena, 26) Veidt, M., “Viscoelastic Tensile and Shear Properties of the 62 wt% Sn - 36 wt% Pb - 2 wt% Ag Solder Alloy”, J. Mater. Sci., 33, 1607-1610 (1998) (Experimental, Mechan. Prop., 18) Bradley, E., III, Hranisavljevic, J., “Characterization of the Melting and Wetting of Sn-Ag-X Solders”, Trans. Electr. Pack. Manuf., 24(4), 255-260 (2001) (Experimental, Phase Relations, 16) Kattner, U.R., Handwerker, C.A., “Calculation of Phase Equilibria in Candidate Solder Alloys”, Z. Metallkd., 92(7), 740-746 (2001) (Calculation, Phase Diagram, Phase Relations, #, 6) Amagai, M., Watanabe, M., Omiya, M., Kishimoto, K., Shibuya, T., “Mechanical Characterization of Sn-Ag-Based Lead-Free Solders”, Microelectron. Reliability, 42, 951-966 (2002) (Mechan. Prop., Morphology, Phys. Prop., 22) Lukas, H.L., “Ag-Pb (Silver-Lead), MSI Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.12122.1.20 (2005) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 6) Nakamori, T., Ikeda, M., Noguchi, K., Shimizu, I., Ohno, Y., “Influence of Shear Heigh on Shear Strength of Tin-Lead Solder Ball Bonding”, Mater. Trans., JIM, 43(8), 2130-2136 (2002) (Experimental, Mechan. Prop., 15) Nogichi, K., Ikeda, M., Shimizu, I., Ohno, Y., Nakamori, T., “Influence of Silver Addition and Reflow Cooling Rate on Microstructure and Strength of Tin-Lead System Solder Ball Bonding” (in Japanese), J. Jpn. Inst. Met., 66(8), 799-807 (2002) (Experimental, Mechan. Prop., Morphology, 11) Zeng, X., “Thermodynamic Analysis of Influence of Pb Contamination on Pb-Free Solder Joints Reliability”, J. Alloys Compd., 348, 184-188 (2003) (Calculation, Phase Relations, Thermodyn., #, 26) Islam, M.N., Chan, Y.C., Sharif, A., “Interfacial Reactions of Sn-Cu and Sn-Pb-Ag Solder with Au/Ni During Extended Time Reflow in Ball Grid Array Packages”, J. Mater. Res., 19(10), 2897-2904 (2004) (Experimental, Kinetics, 25)
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Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 962
cF4 Fm3m Cu
a = 408.57
pure Ag at 25°C [Mas2]
(Pb) < 327.5
cF4 Fm3m Cu
a = 495.02
pure Pb at 25°C [Mas2]
(Pb)
hP2 P63/mmc Mg
a = 326.5 c = 538.7
pure Pb at 25°C, > 10.3 GPa [Mas2]
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
pure Sn, [Mas2]
(Sn) 232 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
Pure Sn at 25°C [Mas2]
( Sn)
tI2 ? Sn
a = 370.0 c = 337.0
pure Sn at 25°C, > 9.0 GPa [Mas2]
Ag4Sn < 724
hP2 P63/mmc Mg
a = 294.36 c = 478.45
[1987Kar]
Ag3Sn < 480
oP8 Pmmn Cu3Ti
a = 596.82 b = 478.02 c = 518.43
[1987Kar]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Ag
Pb
Sn
L + (Ag) Ag4Sn + (Pb)
182
U1
L
4.5
94.5
1
L + Ag4Sn Ag3Sn + (Pb)
~ 180
U2
L
4
87
9
L Ag3Sn + (Pb) + (Sn)
176
E
L
1.75
24.45
73.80
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Sn
Data / Grid: at.% Axes: at.%
e2
Fig. 1: Ag-Pb-Sn. Calculated liquidus surface projection
(Sn)
20
80
E
e3
40
60
Ag3Sn p2 60
40
Ag4Sn
p1 80
20
(Ag)
20
Ag
Ag-Pb
U2 40
Ag-Sn
60
80
Ag-Pb-Sn A-B-C
(Pb) e1
U1
Pb
Sn-Pb
724 p1 L + (Ag) Ag4Sn
304 e1 L (Ag) + (Pb)
480 p2 L+Ag4SnAg3Sn 221 e2 L Ag3Sn + (Sn) 182
(Ag) +L Αg4Sn + Pb
U1
183 e3 l (Pb) + (Sn)
(Ag)+Ag4Sn+(Pb) ~180 Ag4Sn +L Αg3Sn + (Pb) U2 Ag4Sn+Ag3Sn+(Pb) 176
L Αg3Sn + (Pb) + (Sn)
E
Ag3Sn+(Pb)+(Sn) Fig. 2: Ag-Pb-Sn. Reaction scheme
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Ag–Pb–Sn
Fig. 3: Ag-Pb-Sn. Calculated vertical section from the binary Ag-Sn eutectic composition e2 to Sn25.21Sn74.79
119
L 200
Temperature, °C
L+(Sn)+Ag3Sn
L+(Sn) 176°C L+(Sn)+(Pb)
(Sn)+Ag3Sn (Sn)+(Pb)
(Sn)+Ag3Sn+(Pb)
100
Ag 3.80 Pb 0.00 Sn 96.20
10
20
Pb, at.%
Temperature, °C
Fig. 4: Ag-Pb-Sn. Part of the calculated vertical section from the binary Ag-Sn eutectic composition e2 to pure Pb
Ag 0.00 Pb 25.21 Sn 74.79
L
200
L+Ag3Sn
L+Ag3Sn+(Sn) 176°C
Ag3Sn+(Sn) Ag3Sn+(Sn)+(Pb)
100
Ag 3.80 Pb 0.00 Sn 96.20
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10
20
Pb, at.%
Ag 2.75 Pb 27.50 Sn 69.75
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Sn
Data / Grid: at.% Axes: at.%
Fig. 5: Ag-Pb-Sn. Isoactivity lines of Ag (solid lines), Pb (dotted lines), Sn (dashed lines) in liquid alloys at 800°C
0.9 20
0.1 0.8
80
0.7
0.2 40
0.6
60
0.3 0.5 60
0.4 40
0.4
0.5 0.6
0.3
0.7
80
20
0.8 0.2 0.9 0.2 0.1
Ag
MSIT®
0.9 0.4
0.3 20
0.5
0.6 40
0.7 60
0.1
0.8 80
Pb
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Silver – Tin – Zinc Lazar Rokhlin, Evgeniya Lysova and Matvei Zinkevich Introduction The ternary Ag-Sn-Zn system was investigated by [1994Kar, 1996Kar, 1997Hua, 1997Wnu, 1999Oht, 1999Ohn1, 1999Ohn2, 2001Vas, 2002Vas]. [1994Kar, 1996Kar] determined thermodynamic properties of the Ag-Sn-Zn alloys from electromotive force measurements. In the experiments, the liquidus temperatures of the alloys along the sections with Ag:Sn atomic ratios of 3:1, 1:1 and 1:3 were measured. Using vapor pressure measurements [1997Wnu], determined the activity of Zn in alloys, where the total content of Zn and Ag was less than 30 at.%. According to [1997Hua], differential scanning calorimetry suggested that the ternary eutectic was at the composition 3.8Ag-1.8Zn-Sn (at.%). Addition of 1.8 at.% Zn reduced the eutectic temperature of the Sn-3.8Ag (at.%) eutectic by 4°C (217°C) [1997Hua]. [1999Oht, 1999Ohn1, 1999Ohn2] studied the Ag-Sn-Zn phase diagram over the entire concentration range by calculations based on the CALPHAD approach, together with experimental results from differential scanning calorimetry and electron probe microanalysis. The liquidus surface, two isothermal sections at 420 and 190°C and five vertical sections were presented. [1999Oht, 1999Ohn1, 1999Ohn2] also determined the invariant equilibria with the liquid phase. The composition and eutectic temperature in the Sn rich alloys was at 4.03Ag-1.62 Zn-94.35Sn (at.%) and 216.4°C according to [1999Oht, 1999Ohn1, 1999Ohn2] and was in good agreement with [1997Hua]. The liquidus surface of [1999Oht, 1999Ohn1, 1999Ohn2] was unusual in having two separate fields of the primary phase, but was accepted by later investigators. [2001Vas] constructed a complete isothermal section of the Ag-Sn-Zn system at 380°C, using optical and scanning electron microscopy, X-ray analysis, microhardness measurements and differential scanning calorimetry. The isothermal section constructed by [2001Vas] was in general agreement with those of [1999Oht, 1999Ohn1, 1999Ohn2], although there were discrepancies with the + L and + L + fields. [2001Vas] established the non-zero solubility of Zn in the binary Ag-Sn phases, and believed this was a reason for the discrepancy. The Ag-Sn-Zn isothermal section [2001Vas] was confirmed by later work [2002Vas], where diffusion couples were used to study the phase equilibria at 380°C. [2003Ohn] reported a thermodynamic database for micro-soldering alloys, including those of the Ag-Sn-Zn system, which mentioned the successful application of thermodynamics for calculating phase diagrams [1999Oht, 1999Ohn1, 1999Ohn2]. An additional vertical section was constructed in accordance with [1999Oht, 1999Ohn1, 1999Ohn2]. Binary Systems The boundary binary systems Ag-Sn, Ag-Zn and Sn-Zn are taken from [Mas2]. Solid Phases No ternary phases were observed in the system. The terminal solid solutions and the binary phases of the system are described in Table 1. The solubilities in Table 1 are shown after [2001Vas] because this work is the most detailed, and is considered the most reliable. Invariant Equilibria Based on calculations [1999Oht, 1999Ohn1, 1999Ohn2], the invariant reactions associated with the full liquidus surface are shown in Table 2. The reaction scheme is shown in Fig. 1, and is constructed considering the four-phase invariant reactions suggested by [1999Oht, 1999Ohn1, 1999Ohn2], as well as three-phase invariant reactions connected with the four-phase invariant reactions shown on the liquidus surface, but not described by [1999Oht, 1999Ohn1, 1999Ohn2]. The peritectic nature of the three–phase Landolt-Börnstein New Series IV/11C3
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Ag–Sn–Zn
invariant reactions shown on the scheme was determined from the vertical sections of [1999Oht, 1999Ohn1, 1999Ohn2], although their temperatures remain unknown. However, some of the invariant reactions seem to be missing, although they can be inferred from existence of the three-phase invariant reactions, connected with transformation of the phase into in the solid state during cooling in the binary Ag-Zn system [Mas2]. The assumed binary system Sn-Zn [Mas2] had a eutectic transformation at a higher temperature (198.5°C) than that of the four-phase transition reaction with L, (Zn), (Sn) and J(Ag-Zn) at 193.7°C [1999Oht, 1999Ohn1, 1999Ohn2] (Table 1), which is assumed to be the ternary eutectic, E2. However, [1999Oht, 1999Ohn1, 1999Ohn2] assumed this reaction to be the transition one because the Sn-Zn phase diagram used then had the eutectic temperature at 181°C. Liquidus Surface Figures 2 and 3 show the liquidus surface of the Ag-Sn-Zn phase diagram constructed after [1999Oht, 1999Ohn1, 1999Ohn2] with some corrections according to the accepted binary phase diagrams [Mas2]. Taking into account the accepted Sn-Zn phase diagram, the corresponding composition of the liquid phase in the four-phase invariant reaction with the phases (Zn), (Sn) and J(Ag-Zn) is shown as eutectic E2. Isothermal Sections Figure 4 shows isothermal section at 380°C. It is constructed after [2001Vas, 2002Vas] considering results of these works to be the most reliable. Compared to [1999Oht, 1999Ohn1, 1999Ohn2] which gave the isothermal sections at 190°C and 420°C, [2001Vas, 2002Vas] used more detailed experiments and, in general, confirmed the results of [1999Oht, 1999Ohn1, 1999Ohn2], which were obtained mainly by calculations. Temperature – Composition Sections Figures 5 and 6 show two vertical sections of the Ag-Sn-Zn phase diagram. They are reproduced after [1999Oht, 1999Ohn1, 1999Ohn2] with minor corrections according to the accepted binary phase diagrams Ag-Sn and Sn-Zn [Mas2]. Thermodynamics [1994Kar, 1996Kar] determined the partial free energies of Zn in 30 liquid Ag-Sn-Zn alloys as a function of composition and temperature. These data were used by [1999Oht, 1999Ohn1, 1999Ohn2], together with their own measurements to obtain a self-consistent thermodynamic assessment. The derived thermodynamic parameters were incorporated in the database for micro-soldering alloys [2003Ohn] thus enabling many useful subsequent calculations. [1998Pen] compared experimental free energy and enthalpy of mixing values [1994Kar, 1996Kar] with those calculated using different models for the extrapolation of thermodynamic functions from binary subsystems, and found some large discrepancies. [1997Wnu] measured the activity of Zn in Ag-Sn-Zn alloys at 630°C, describing the activity coefficient of Zn as: ln fZn = 0.414 – 0.424( 0.13) # xZn – 1.063(0.20) # xAg; xZn + xAg < 0.30 Notes on Materials Properties and Applications The Sn rich alloys of the Ag-Sn-Zn system are considered as the basis of non-toxic lead-free solders for electronic assemblies [1994McC, 1997Hua, 1997Lee, 1999Oht, 1999Ohn1, 1999Ohn2, 2001Vas, 2002Che, 2002Tsa, 2003Cha1, 2003Cha2, 2003Ohn, 2003Son2]. [1997Hua, 1997Lee] gave a review on the lead-free solder alloys for electronics, together with their properties, including those of the Ag-Sn-Zn system. Ag based alloys containing Sn and Zn after internal oxidation are of interest as materials for electrical contacts [1993Dev]. [2002Tsa] studied thermal expansion coefficients and melting ranges of the Sn rich alloys containing 15.2 at.% Zn and 0-3.58 at.% Ag. The thermal expansion coefficient was determined to be in the range MSIT®
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18-22 10–6#K–1 when the Ag content was less than 2.05 at.%, and in the range 24-27#10–6#K–1 when the Ag content was greater than 2.05 at.%. The onset of melting (~198°C) did not change significantly with increasing Ag content. The melting range increased with increasing Ag content from 7.7°C for the binary alloy 15.2Zn-Sn (at.%), up to 27.4°C for the alloy 3.58Ag-15.2Zn-Sn (at.%). [2001Vas] reported microhardness measurements of the Ag-Sn-Zn alloys. [2003Son2] studied mechanical properties of the Sn rich Ag-Sn-Zn alloys containing 15.2 at.% Zn and 0.51 - 3.58 at.% Ag. Addition of Ag decreased the strength and increased the ductility of the alloys. Miscellaneous [1956San] studied effect of Sn on the evaporation rate of Zn in the Ag-Zn alloys at 650°C in vacuum, and the addition of 0.19 at.% Sn to alloy of 34.5 at.% Zn decreased evaporation rate of Zn. [2003Son1] investigated behavior of intermetallic compounds in the molten Ag-Sn-Zn alloys for solders during cooling and remelting. References [1956San]
[1987Kar]
[1993Dev] [1994Kar] [1994McC]
[1996Kar]
[1997Hua] [1997Lee] [1997Wnu]
[1998Pen]
[1999Ohn1]
[1999Ohn2] [1999Oht]
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Santalov, F.A., “The Influence of Additions of Low-Melting-Point Metals on the Structure of Specimens (of Silver- Zinc and Silver-Cadmium Alloys) and of the Rate of Evaporation of the Volatile Component of the Solid Solutions” (in Russian), Fiz. Met. Metalloved., 3(2), 247-253 (1956) (Experimental, 8) Karakaya, I., Tompson, W.P., “The Ag-Sn (Silver - Tin) System“, Bull. Alloy Phase Diagrams, 8(4), 340-347 (1987) (Assesment, Phase Diagram, Crys. Structure, Thermodyn., 51) Dev, S.C., Basak, O., Mohanty, O.N., “Defelopment of Cadmium-Free Silver Metal-Oxide Contact Materials”, J. Mater. Sci., 28, 6440-6544 (1993) (Experimental, Morphology, 11) Karlhuber, S., Kurt, L., Komarek, K.L., Mikula, A., “Thermodynamic Properties of Liquid Ag-Sn-Zn Alloys”, Z. Metalkd., 85, 307-311 (1994) (Experimental, Thermodyn., 33) McCormack, M., Kammlott, G.W., Chen, H.S., Jin, S., “New Lead-Free, Sn-Ag-Zn-Cu Solder Alloy with Improved Mechanical Properties”, Appl. Phys. Lett., 65(10), 1233-1235 (1994) (Experimental, Mechan. Prop., Morphology, 6) Karlhuber, S., Peng, M., Mikula, A., “Thermodynamic Properties of Ternary Liquid Zinc-Alloys”, J. Non-Cryst. Solids, 205-207, 421-424 (1996) (Experimental, Thermodyn., 11) Hua Fay, Glaser, J., “Lead-Free Solders for Electronic Assembly”, Des. Reliab. Solders Solder Interconnect., Proc. Symp., 65-73 (1997) (Experimental, 63) Lee, N.C., “Getting Ready for Lead-Free Solders”, Soldering Surf. Mount Technol., 9(2), 65-69 (1997) (Mechan. Prop., Review, 0) Wnuk, G., Pomianek, T., Romanowska, J., Rychlewski, M., “Influence of Ag on the Activity of Zn in {(1 - x1 - x2)Sn + x1Zn + x2Ag} (I) at T = 903K”, J. Chem. Thermodyn., 29, 931-939 (1997) (Kinetics, Thermodyn., 33) Peng, M., Qiao, Z., Mikula, A., “Comparison between Calculated and Measured Thermodynamic Data of Liquid (Ag, Au, Cu)-Sn-Zn Alloys”, Calphad, 22(4), 459-468 (1998) (Calculation, Thermodyn., 19) Ohnuma, I., Liu, X.J., Ohtani, H., Ishida, K., “Thermodynamic Database for Phase Diagrams in Micro-Soldering Alloys”, J. Electron. Mater., 28(11), 1164 -1171 (1999) (Calculation, Phase Relations, 22) Ohnuma, I., Liu, X.J., Ohtani, H., Ishida, K., “Phase Diagrams of Pb – Free Solder Alloys” (in Japanese), Materia Japan, 38, 923-926 (1999) (Calculation, Phase Diagram, 13) Ohtani, H., Miyashita, M., Ishida, K., “Thermodynamic Study of Phase Equilibria in the Sn-Ag-Zn System” (in Japanese), J. Jpn. I. Met., 63(6), 685-694 (1999) (Assessment, Calculation, Experimental, Phase Diagram, Thermodyn., *, 68)
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[2002Che]
[2002Tsa]
[2002Vas]
[2003Cha1]
[2003Cha2]
[2003Ohn]
[2003Son1] [2003Son2]
Vassilev, G.P., Dobrev, E.S., Evtimova, S.K., Tedenac, J.-C., “Studies of the Phase Equilibria in the Ag-Sn-Zn System”, J. Alloys Compd., 327, 285-291 (2001) (Experimental, Phase Diagram, Mechan. Prop.,*, 26) Cheng, S.Ch., Lin, K.L., “The Thermal Properties of Lead-Free Sn-8.55Zn-1Ag-XAl Solder Alloys and Their Wetting Interaction with Cu”, J. Electron. Mater., 31(9), 940-945 (2002) (Experimental, Interface Phenomena, Phase Relations, 27) Tsai, Y-L., Hwang, W-S., Wang, H-S., Wu, M-H., “Measurements of Thermal Expansion Coefficients and Melting Ranges for Sn-9Zn-xAg Lead-Free Solder Alloys”, Int. J. Cast Metal Res., 15, 181-186 (2002) (Experimental, 6) Vassilev, G.P., Evtimova, S.K., Tedenac, J.-C., Dobrev. E.S., “Experimental Study of the Ternary Ag-Sn-Zn System through Diffusion Couples”, J. Alloys Compd., 334, 182-186 (2002) (Experimental, Phase Diagram, Mechan. Prop., 7) Chang, T.-C., Hsu, Y.-T., Hon, M.-H., Wang, M.-C., “Enhancement of the Wettability and Solder Joint Reliability at the Sn-9Zn-0,5Ag Lead-Free Solder Alloy-Cu Interface by Ag Precoating”, J. Alloys Compd., 360, 217-224 (2003) (Experimental, Mechan. Prop., Morphology, 21) Chang, T.-C., Wang, M.-C., Hon, M.-H., “Effect of Aging on the Growth of Intermetallic Compounds at the Interface of Sn-9Zn-xAg/Cu Substrates”, J. Cryst. Growth, 250, 236-243 (2003) (Crys. Structure, Experimental, Interface Phenomena, Morphology, Phase Relations, 14) Ohnuma, I., Miyashita, M., Liu, X.J., Ohtani, H., Ishida, K., “Phase Equilibria and Thermodynamic Properties of Sn-Ag Based Pb-Free Solder Alloys”, IEEE T. Electron. PA. M., 26(1), 84-89 (2003) (Calculation, Phase Diagram, Thermodyn., 21) Song, J.M., Lin, K.L., “Behavior of Intermetallics in Liquid Sn-Zn-Ag Solder Alloys”, J. Mater. Res., 18(9), 2060-2067 (2003) (Experimental, 14) Song, J.M., Lan, G.F., Lui, T.S., Chen, L.H., “Microstructure and Tensile Properties of Sn-9Zn-xAg Lead – Free Solder Alloys”, Scr. Mater., 48, 1047-1051 (2003) (Experimental, Mechan. Prop., 10)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ag) < 961.93
cF4 Fm3m Cu
a = 408.57
hP2 P63/mmc Mg
a = 266.50 c = 494.70
tI4 I41/amd Sn
a = 583.18 c = 318.18
cF8 Fd3m C(diamond)
a = 648.92
(Zn) < 419.58
(Sn) 231.9681 - 13
(Sn) < 13
MSIT®
pure Ag at 25°C [Mas2] dissolves up to 32.1 at.% Zn and 11.5 at.% Sn [Mas2] pure Zn at 25°C [Mas2] dissolves up to 5.0 at.% Ag and 0.039 at.% Sn [Mas2] pure Sn at 25°C [Mas2] dissolves up to 0.09 at.% Ag and 0.6 at.% Zn [Mas2] [Mas2]
Landolt-Börnstein New Series IV/11C3
Ag–Sn–Zn
125
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, AgZn < 274
hP9 P3 , AgZn
a = 763.60 c = 281.79
cI2 Im3m W
a = 315.58
cI52 I43m Cu5Zn8
a = 934.07
, AgZn 710 - 258
, Ag5Zn8 < 661
J (Ag-Zn) < 631
37.0 to ~ 51.2 at.% Zn [Mas2, P] dissolves up to ~1 at.% Sn at 190°C [1999Oht] 36.7 to 58.6 at.% Zn [Mas2, P] dissolves up to ~13 at.% Sn at 380°C [2001Vas] 58.5 to 64.7 at.% Zn [Mas2, P] dissolves up to ~2.6 at.% Sn at 380°C [2001Vas]
hP2 P63/mmc Mg
~66.2 to 89 at.% Zn [Mas2, P] at 22.5 at.% Ag [P]
a = 282.27 c = 442.74
dissolves up to ~0.2 at.% Sn at 380°C [2001Vas] (Ag-Sn) < 724
J, Ag3Sn < 480
hP2 P63/mmc Mg
11.8 to 22.85 at.% Sn [Mas2, 1987Kar]
oP8 Pmmn Cu3Ti
a = 294.36 c = 478.45
at 16.0 at.% Sn [P]
a = 596.82 b = 478.02 c = 518.43
at 23.7 to 25 at.% Sn [Mas2, 1987Kar]
dissolves up to ~6 at.% Zn at 380°C [2001Vas]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Ag
Sn
Zn
L+
?
p5
L
12.69
52.62
34.69
L + (Ag-Zn)
?
p8
L
5.66
91.13
3.21
L + (Sn)
?
p10
L
1.95
94.98
3.07
L + + J(Ag-Zn)
560
U1
L
28.67
11.26
60.07
L + + J(Ag-Zn)
482
U2
L
11.83
49.28
38.89
L + (Ag-Sn) (Ag) +Ag3Sn
437.2
U3
L
51.11
38.21
10.68
Landolt-Börnstein New Series IV/11C3
MSIT®
Ag–Sn–Zn
126
Reaction
T [°C]
Type
Phase
Composition (at.%) Ag
Sn
Zn
L + (Ag) + (Ag-Zn)
266.3
U4
L
6.37
90.60
3.03
L + + (Ag-Zn)
252.5
U5
L
3.99
92.45
3.56
L + (Ag) (Ag-Zn) + Ag3Sn
240.8
U6
L
5.42
92.43
2.15
L + (Sn) + (Ag-Zn)
217.7
U7
L
2.82
94.66
2.52
L (Sn) + (Ag-Zn) + Ag3Sn
216.4
E1
L
4.03
94.35
1.62
L + (Sn) + J (Ag-Zn)
209.7
U8
L
0.42
91.87
7.71
L J(Ag-Zn) + (Sn) + (Zn)
193.7
E2
L
0.04
85.70
14.26
The composition of liquid in p5 is taken from Fig. 2, the compositions of liquid in p8, p10 and E2 are taken from Fig. 3 [1999Oht]
MSIT®
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Ag-Sn
Ag-Sn-Zn
Ag-Zn
724 p1 L + (Ag) ζ(Ag-Sn)
Sn-Zn
710 p2 L + (Ag) β 661 p3 L+βγ 560
L + γ β + ε(Ag-Zn)
γ + β + ε(Ag-Zn)
L + β + ε(Ag-Zn) 482 480 p6 L+ζ(Ag-Sn)Ag3Sn
L + β γ + ε(Ag-Zn)
L + γ + ε(Ag-Zn) 437.2
L + β(Ag) + ζ(Ag-Zn)
216.4
252.5
L(βSn)+ζ(Ag-Zn)+Ag3Sn
L + β γ + ζ(Ag-Zn)
L+γ+ζ(Ag-Zn)
U6
L+(βSn)+ζ(Ag-Zn)
274 p9 β + γ ζ(Ag-Zn)
U4
217.7
U5
258 e1 β (Ag) + ζ(Ag-Zn)
β+γ+ζ(Ag-Zn) ? p10 L + γ (βSn)
(Ag)+ζ(Ag-Zn)+Ag3Sn
L+ζ(Ag-Zn)+Ag3Sn 221 e2 L Ag3Sn + (βSn)
? p8 L + β ζ(Αg-Zn)
β+(Ag)+ζ(Ag-Zn)
L+(Ag)ζ(Ag-Zn)+Ag3Sn
431 p7 L + ε(Ag-Zn) (Zn)
U3
ζ(Ag-Sn)+(Ag)+Ag3Sn
L+(Ag)+ζ(Ag-Zn)
240.8
U2
p5 L+βγ
Ag–Sn–Zn
266.3
?
β + γ + ε(Ag-Zn)
L+ζ(Ag-Sn)(Ag)+Ag3Sn
L+(Ag)+Ag3Sn
631 p4 L + γ ε(Ag-Zn)
U1
L + γ (βSn) + ζ(Ag-Zn)
U7
γ+(βSn)+ζ(Ag-Zn)
E1
(βSn)+ζ(Ag-Zn)+Ag3Sn 209.7
L + γ (βSn) + ε(Ag-Zn)
γ+(βSn)+ε(Ag-Zn)
U8
198.5 e3 L (Zn) + (βSn)
L+(βSn)+ε(Ag-Zn) 193.7
Lε(Ag-Zn)+(βSn) +(Zn)
E2
ε(Ag-Zn)+(βSn)+(Zn)
127
MSIT®
Fig. 1: Ag-Sn-Zn. Reaction scheme
Ag–Sn–Zn
128
Sn Fig. 2: Ag-Sn-Zn. Liquidus surface projection
Data / Grid: at.% Axes: at.%
(β Sn) 250
e2 300
e3
20
80
350 400
Ag3Sn
γ
40
60
p6
450 U2
p5
ζ(Ag-Sn)
60
40
U3 500
p1
β
80
ε(Ag-Zn)
20
550 U1
600
(Ag)
γ
650°C p2 40
20
Ag i 2
i id
60 p
(Zn) p4
3
80
p7
Zn
f
Sn
Data / Grid: at.% Axes: at.%
Fig. 3: Ag-Sn-Zn. Liquidus surface of the Sn rich corner (β Sn)
e2
U7
E1
p10
ζ(Ag-Zn)
U6 10
Ag3Sn (Ag)
U5 p8
U4
U8
γ
90
β ε(Ag-Zn)
E2 e3 (Zn)
Ag Zn Sn
MSIT®
20.00 0.00 80.00
10
Ag Zn Sn
0.00 20.00 80.00
Landolt-Börnstein New Series IV/11C3
Ag–Sn–Zn
129
Sn
Data / Grid: at.% Axes: at.%
Fig. 4: Ag-Sn-Zn. Isothermal section at 380°C L 20
80
40
60
60
40
β +L
Ag3Sn
ε(Ag-Zn)+L
γ +L
80
20
ζ(Ag-Sn)
β
γ
(Ag) 20
Ag
Fig. 5: Ag-Sn-Zn. Vertical section at 40 mass% Sn, plotted in at.%
40
60
80
L 500
β +L
Temperature, °C
Zn
ζ (Ag-Sn)+L ε(Ag-Zn)+L
(Ag)+L
Ag3Sn+L
γ+L
250
(Zn)+L
ζ (Ag-Zn)+L
(Zn)+(β Sn)
Ag3Sn+ζ (Ag-Zn)+(β Sn)
ζ (Ag-Zn)+(β Sn) 0
Ag 62.27 Zn 0.00 Sn 37.73
Landolt-Börnstein New Series IV/11C3
(Zn)
ε(Ag-Zn)
20
γ+(β Sn) 40
Zn, at.%
ε(Ag-Zn)+(β Sn) 60
Ag 0.00 Zn 73.14 Sn 26.86
MSIT®
Ag–Sn–Zn
130
Temperature, °C
Fig. 6: Ag-Sn-Zn. Vertical section at 60 mass% Sn, plotted in at.%
600
L 500
400
Ag3Sn+L
(Ag)+L
β +L
ε(Ag-Zn)+L
γ+L
300
(Zn)+L 200
100
ζ (Ag-Zn)+L Ag3Sn+ζ (Ag-Zn)+(β Sn) (Zn)+(β Sn)
(β Sn)+γ
ζ (Ag-Zn)+(β Sn) 0
Ag 42.32 Zn 0.00 Sn 57.68
MSIT®
10
ε(Ag-Zn)+(β Sn) 20
30
Zn, at.%
40
50
Ag 0.00 Zn 54.76 Sn 45.24
Landolt-Börnstein New Series IV/11C3
Au–Bi–Sn
131
Gold – Bismuth – Tin Ortrud Kubaschewski Introduction The partial system AuSn-Bi-Sn has been investigated experimentally [1985Hum, 1987Max]. [1985Hum] studied the AuSn-Bi section and found it to be quasibinary, Fig. 1. [1987Max] determined the region of the miscibility gap at 386°C. The partial system is rather complex, including not only a miscibility gap but also three invariant reactions. The data reported by [1985Hum, 1987Max] are accepted. The liquidus surface and a reaction scheme constructed from their results are shown in Fig. 2 and Fig. 3. Binary Systems The Au-Sn system, evaluated by [Mas2], shows uncertainties concerning the equilibrium reactions for Au rich alloys. [1987Leg] have shown the existence of two peritectic reactions at 532 and 519°C, respectively. The solid state transformation requires further study. The Au-Bi system [Mas2] with a peritectic temperature of 377.5°C [1983Eva] is accepted. The Bi-Sn diagram by [Mas2] is accepted. Solid Phases No ternary phases have been found in the partial system. The phases of the partial AuSn-Bi-Sn system are listed in Table 1. Quasibinary Systems The AuSn-Bi section is quasibinary, Fig. 1 [1987Max]. Bi may be soluble in AuSn but the solubility has not been determined. A solubility of about 8 at.% Bi in AuSn is indicated by [1985Hum] at the eutectic temperature. A monotectic reaction occurs at 386°C: L' L'' + AuSn. Invariant Equilibria Three ternary invariant equilibria, Table 2, are reported for the partial ternary system. The reaction scheme is presented in Fig. 3. Liquidus Surface Figure 2 illustrates the projection of the liquidus surface [1986Max]. The extent of the liquid immiscibility entering from the AuSn-Bi section is not known and shown here as a dashed curve, disappearing at a critical tie line L AuSn below 386°C. The regions of primary crystallization of (Bi) and (Sn) are confined to narrow regions adjoining the Bi-Sn binary edge. Isothermal Sections Figures 4, 5, 6 and 7 display the isothermal sections at 240, 200, 160 and 139°C respectively, for the partial Sn-AuSn-Bi system. The last one corresponds to the temperature of the invariant reaction E which is shown by dash lines in Fig. 7. The isothermal sections have been slightly amended from [1986Max] to be consistent with the vertical sections. Temperature – Composition Sections The vertical section at Au50Bi50 - Sn [1985Hum, 1986Max] is shown in Fig. 8.
Landolt-Börnstein New Series IV/11C3
MSIT®
Au–Bi–Sn
132 References [1983Eva] [1985Hum] [1986Max] [1987Max] [1987Leg]
Evans, D.S., Prince, A., “The Au-Bi Phase Diagram”, CALPHAD XII Meeting, Liége (1983) Humpston, G., “The Constitution of Some Ternary Au-Based Solder Alloys”, Ph.D. Thesis, Brunel University, 67-68 and 95, Figs. 46-50 (1985) (Phase Diagram, Experimental, 2) Maxwell, C.A., M.Sc. Thesis, Univ. Manchester, Faculty of Technology (1986) Maxwell, C.A., Hayes, F.H., private communication Legendre, B., Hancheng, C.H., Hayes, F., Maxwell, C., Evans, D.S., Prince, A., “Contribution Towards Clarification of the Au-Sn Phase Diagram for Sn Contents Less than 25 at.%”, Mater. Sci. Technol., 3, 875-876 (1987) (Phase Diagram, Experimental, 6)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Au) < 1064.43
cF4 Fm3m Cu
a = 407.82
at 25°C [Mas2]
(Bi) < 271.442
hR6 R3m As
a = 454.613 c = 1186.152
at 31°C [V-C2]
(Sn) 231.97 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
at 25°C [Mas2]
AuSn < 419.3
hP4 P63/mmc NiAs
a = 432.18 c = 552.30
[V-C2]
AuSn2 < 309
oP24 Pbca AuSn2
a = 690.9 b = 703.7 c = 1178.9
[V-C2]
AuSn4 < 252
oC20 Aba2 PtSn4
a = 651.24 b = 651.62 c = 1170.65
at 25°C [V-C2]
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Au
Bi
Sn
L' L'' + AuSn
368
e1 (max)
L' L''
<35 <17.8
30 64.4
35 > 17.8>
L’ (Bi) + AuSn
262.5
e2 (max)
L
<2.65
94.7
2.65>
L + AuSn AuSn2 + (Bi)
224
U1
L
1.9
76.8
21.3
L + AuSn2 (Bi) + AuSn4
146
U2
L
2.7
44.7
52.6
L (Bi) + (Sn) + AuSn4
137
E
L
0.8
40.0
59.2
values in brackets are estimated from Fig. 1.
MSIT®
Landolt-Börnstein New Series IV/11C3
Au–Bi–Sn
Fig. 1: Au-Bi-Sn. Quasibinary system Bi-AuSn
133
L'+AuSn
L
L'+L"
400
Temperature, °C
386°C
L"+AuSn 300
271.44 262.5
AuSn+(Bi)
(Bi) 10
Bi
AuSn
20
30
Au 50.00 Sn 50.00 0.00 Bi
40
Au, at.%
Bi Fig. 2: Au-Bi-Sn. Partial liquidus surface in the region AuSn-Bi-Sn
Data / Grid: at.% Axes: at.%
e2
20
80
U1
e"1 40
60
(Bi) U2 60
e4
AuSn2
200
e'1
E
240
280
40
AuSn
20
AuSn4
360°C
80
(Sn)
Au
Landolt-Börnstein New Series IV/11C3
20
40
60
p1
80
p2
e3
Sn
MSIT®
134
MSIT® Bi-Sn
AuSn-Bi-Sn
AuSn-Sn
386 e1 L' L'' + AuSn L' + L'' + AuSn < 386
309 p1 l + AuSn AuSn2
262.5 e2 LAuSn + (Bi)
252 p2 l + AuSn2 AuSn4 224
L + AuSn AuSn2 + (Bi) U1
146
L + AuSn2 AuSn4 + (Bi) U2
L +AuSn4 + (Bi)
AuSn2 +AuSn4 + (Bi)
139 e4 l(Bi) + (Sn) 137
L (Bi) +(Sn) + AuSn4
Landolt-Börnstein New Series IV/11C3
(Bi) + (Sn) +AuSn4 Fig. 3: Au-Bi-Sn. Reaction scheme
E
Au–Bi–Sn
L +AuSn2 + (Bi)
217 e3 l (Sn) + AuSn4
AuSn +AuSn2 + (Bi)
Au–Bi–Sn
135
Bi
Data / Grid: at.% Axes: at.%
Fig. 4: Au-Bi-Sn. Isothermal section at 240°C L+AuSn+(Bi) 20
80
L+ AuSn 40
60
60
40
L+ AuSn2
L
80
20
AuSn+ AuSn2+L L+AuSn2+AuSn4 20
Au
40
AuSn 60
80
AuSn4
AuSn2
Bi
L+AuSn4
Sn
Data / Grid: at.% Axes: at.%
Fig. 5: Au-Bi-Sn. Isothermal section at 200°C 20
L+(Bi)+AuSn2 80
L+(Bi)
40
60
60
40
L+AuSn2 L 80
20
AuSn+AuSn2+(Bi)
Au
Landolt-Börnstein New Series IV/11C3
20
40
L+AuSn4 L+(Sn)
(Sn) 80 AuSn4 (Sn)+AuSn AuSn 60 AuSn2 4Sn L+AuSn2+AuSn4 L+(Sn)+AuSn4
MSIT®
Au–Bi–Sn
136
Bi
Data / Grid: at.% Axes: at.%
Fig. 6: Au-Bi-Sn. Isothermal section at 160°C 20
80
L+(Bi)+AuSn2 L+(Bi)
40
60
L+AuSn2 60
L
40
L+AuSn4 L+(Sn)
80
20
L+(Sn)+AuSn4
AuSn+AuSn2+(Bi) (Sn) (Sn)+AuSn4 20
Au
40
AuSn 60AuSn2
80AuSn4 L+AuSn2+AuSn4
Bi
Sn
Data / Grid: at.% Axes: at.%
Fig. 7: Au-Bi-Sn. Isothermal section at 139°C 20
80
40
60
L 60
40
(Bi)+(Sn)+AuSn4 80
20
AuSn+ AuSn2+ (Bi)
Au
MSIT®
20
40
AuSn 60
AuSn2+ AuSn4+ (Bi) AuSn2
(Sn) (Sn)+AuSn4 80 AuSn4
Sn
Landolt-Börnstein New Series IV/11C3
Au–Bi–Sn
Fig. 8: Au-Bi-Sn. Part of the vertical section Au50Bi50-Sn
137
400
L'+L"+AuSn L L+AuSn L+AuSn+AuSn2 300
Temperature, °C
L+AuSn+(Bi) L+AuSn2
L+AuSn2+AuSn4 L+AuSn4
224°C
L+(Sn)
L+AuSn2+(Bi)
200
(Sn) 146
137
AuSn+AuSn2+(Bi)
L+(Bi)+AuSn4
L+(Sn)+AuSn4
100
AuSn2+AuSn4+(Bi)
Au 33.34 Sn 33.33 Bi 33.33
Landolt-Börnstein New Series IV/11C3
40
50
60
AuSn4+(Bi)+(Sn) 70
80
AuSn4+(Sn) 90
Sn
Sn, at.%
MSIT®
138
Au–Cu–Sn
Gold – Copper – Tin Nataliya Bochvar, Yurii Liberov Introduction Interdiffusion in a diffusion sample of the composition electrolytic Cu (99.95%, 5000 m)/Ag (15 m)/Au (13 m)/Sn-Pb solder (41.5% Pb, 100 m) resulted in the occurrence of three ternary compounds, AuCu5Sn5, Au2Cu4Sn5 and Au3Cu3Sn5. These were formed between 110 and 212°C [1971Cre]. However, the existence of the first two phases was not confirmed in subsequent studies. The Au poor part of the Au-Cu-Sn system was investigated by [1988Roe1, 1988Roe2] using metallography and differential scanning calorimetry. Starting materials were 99.999% pure Au, Cu and Sn. Alloys were induction melted under an argon atmosphere, annealed under vacuum in silica or pyrex ampoules for various time periods and water quenched. Diffusion couples were prepared from binary Au-Cu alloys which were hot-worked and annealed for 48 h at temperatures ~25 K below the solidus temperature. After polishing they were pressed against sheets of Sn and annealed at 170°C in evacuated capsules. 232 h was the annealing time for the AuCu/Sn and Au3Cu/Sn couples and 256 h for the AuCu3/Sn couple. The specimens were water quenched and examined by electron probe microanalysis. Solid state diffusion couples and ternary alloys were examined to establish an isothermal section at 170°C. Three ternary phases, 0 (~25 at.% Au, 55 at.% Cu, 20 at.% Sn), 3 (~33.3 at.% Au, 33.3 at.% Cu, 33.3 at.% Sn), and (approximately, Au3Cu3Sn5) were found and their crystal structures were determined by convergent beam electron diffraction. The 0 phase, denoted -1 in this assessment, extends along the Au-Cu axis. The 3 compound - -2 in this review - has an AuCuSn formula. The (Au3Cu3Sn5) compound, denoted -3 in the present review, was earlier reported by [1971Cre]. A liquidus surface and a reaction scheme for the Au poor part of the system (< 50 at.% Au) were proposed. [1990Kar] and [1992Kar1] performed density measurements and used X-ray diffraction, metallography, scanning electron microscopy and differential thermal analysis to study the phase equilibria of the Au-Cu-Sn system. They prepared specimens by fusing (for 1 min at 1100 to 1200°C under shaking) 99.99% Au, 99.9% Cu, and 99.98% pure Sn in evacuated and sealed silica-glass tubes. The molten specimens were quenched into water and then annealed at 360°C. An isothermal section at 360°C was established. It contains three ternary compounds A, B, and C labeled -4, -1 and -2 respectively, in the present review. B and C correspond to those phases detected by [1988Roe1, 1988Roe2] and A was found to contain 20 at.% Sn, 42 to 47% Au and 38 to 33% Cu. According to [1990Kar, 1992Kar1, 1992Kar2, 1992Kar3] -1 forms through a solid-state reaction, whereas [1988Roe2] revealed this phase in equilibrium with the liquid (see Figs. 2 and 3). -2 is formed through a solid state reaction at about 370°C. [1990Kar] and [1992Kar1] found the (AuSn) and 1(Cu6Sn5) phases to form a continuous series of solid solutions and the phases and (both Au-Sn) as well as J1 (Cu3Sn), to extend into the ternary system. [1992Kar1] examined the phase constitution in the section Au80–xCuxSn20. Two invariant four-phase equilibria at ~515 and ~320°C were revealed. Based on these data, [1992Kar1] derived tentative partial isothermal sections at ~450°C, ~320°C as well as just above and below 320°C. [1990Kar, 1992Kar2] and [1992Kar3] determined the crystal structure of the ternary phases -1, -2 and -4. Using X-ray powder diffraction, scanning electron microscopy/electron microprobe analysis, and other analytical techniques, [1994Pep] revealed the ternary Au8Cu8Sn4 and found it homogeneous in the range from Au7Cu9Sn4 to Au9Cu7Sn4. The alloys were prepared from pure metals by melting in corundum crucibles. The samples were molted for a while, and cooled with the rate of 2 K#min–1 to room temperature. The ingots were annealed at 175°C for 7 and 240 d. The above Au8Cu8Sn4 ternary phase exhibited the crystal structure same as the [1988Roe1, 1988Roe2] and [1990Kar, 1992Kar1] compounds, both denoted -1 in the present assessment. Based on the X-ray diffraction data, [2002Luc] constructed a partial vertical section AuCu-Sn up to 8 at.% Sn at the temperatures below 450°C. [2002Luc] prepared the alloys by melting the charge sealed in evacuated silica capsules in a high-frequency induction furnace. The ingots were forged at room
MSIT®
Landolt-Börnstein New Series IV/11C3
Au–Cu–Sn
139
temperature, annealed during 20000 min at different temperatures from 450 to 200°C, and quenched into ice brine from an annealing temperature. In the AuCu-Sn section [2002Luc] showed two low-temperature ternary phases. One of them, AuCu(III) (-5 in the present assessment) appeared orthorhombic with eight atoms per unit cell similar to the ternary phase of the Au-Cu-Ga system and isomorphous to a Au2CuZn phase of the Au-Cu-Zn system. The other low-temperature ternary AuCu(III)´ phase exhibited the cubic Mn type structure like the 0 [1988Roe1, 1988Roe2], [1990Kar, 1992Kar1], or Au8Cu8Sn4 [1994Pep], which we denote -1 here. [2002Huh] examined three alloys in the Sn corner of the Au-Cu-Sn system with 1, 2 and 3 mass% Au and constant 0.7 mass% Cu content, using differential scanning calorimetry, scanning electron microscopy, electron probe microanalysis and X-ray diffraction. [2002Huh] confirmed the solidification of ternary (Sn) + + ´ eutectic and a large solid solubilities of Au in the ´ phase and Cu in the phase. In a review on employing X-ray microprobe to study diffusion, [1989Gol] discussed the data reported by [1988Roe1, 1988Roe2]. Binary Systems The binary systems, Au-Cu and Cu-Sn, were accepted as given by [Mas2] and the Au-Sn binary phase diagram, as given by [1993Oka] (Fig. 1). Solid Phases Five ternary compounds were confirmed to exist in the Au-Cu-Sn system. [1988Roe2] reported the -1 and -3 compounds to precipitate from the liquid; however no data on the formation reactions is available. In the solidified alloys, the -1 compound was found to exist at 360°C [1990Kar, 1992Kar1], below 320°C [2002Luc], at 175°C [1994Pep], and 170°C [1988Roe1, 1988Roe2]; the -2 compound, at 170°C [1988Roe1, 1988Roe2] and 360°C [1990Kar, 1992Kar1]; -3, between 110 and 220°C [1971Cre, 1988Roe1, 1988Roe2]; -4, at 360°C [1990Kar, 1992Kar1], and -5, below 270°C [2002Luc]. The -1 compound has a homogeneity range that, at 360°C, extends from 11 to 38 at.% Au, from 69 to 42 at.% Cu, and from 18 to 20.5 at.% Sn. This homogeneity range becomes wider at 300°C, its lower limit shifting to 35 at.% Cu, and seems to reach even lower Cu contents below 300°C [1992Kar1]. The homogeneity range of the -2 phase extends from 33 to 37 at.% Au (33.7 to 29.7 at.% Cu) at 33.3 at.% Sn at 360°C [1990Kar, 1992Kar1]. The -3 phase was found homogeneous between 22.1 and 25.9 at.% Au, 22.5 and 26.8 at.% Cu, and 51.1 and 51.7 at.% Sn at 170°C [1988Roe1, 1988Roe2]. The -4 phase is homogeneous from 42 to 47 at.% Au and from 38 to 33 at.% Cu for 20 at.% Sn and 360°C. The minimum gold limit of the homogeneity range shifts to 15 at.% Au (65 at.% Cu) at 550°C [1990Kar]. Most of the binary phases extend into the ternary system. (AuSn) and 1 (Cu6Sn5) form a continuous series of solid solutions (Au1–xCux)Sn1–y at 360°C [1990Kar, 1992Kar1], denoted in the present assessment. At 170°C, the phase dissolves about 15 at.% Cu and the ’ phase, about 8 at.% Au [1988Roe1, 1988Roe2]. At 360°C, the phase dissolves up to 50 at.% Cu, the J1 phase (Cu3Sn), up to 17 at.% Au, and the phase, up to 20 at.% Cu. The structural data for the binary and ternary phases are given in Table 1. Invariant Equilibria [1988Roe2] proposed a solidification reaction scheme for the Au poor part of the Au-Cu-Sn system. It is presented slightly modified in Fig. 2. [1988Roe2] deduced nine transition (U1 to U9) and one eutectic (E2) reaction from the metallographic data obtained for the ternary alloys. Besides, the U9 and E2 reactions were also revealed by differential scanning calorimetry. Reactions U9, L + -3 + 1, at 279°C, and U8, L + J -3 + , above the U9 temperature [1988Roe2], are unlikely as implying the phase to be stable at higher temperatures in the ternary than in the binary system, which would only be possible with Cu soluble in . Therefore, we describe the U9 reaction as P and the U8 reaction, as E1. The solid-state invariant reactions are unknown and omitted in Fig. 2.
Landolt-Börnstein New Series IV/11C3
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140
Au–Cu–Sn
Two four-phase reactions in the Au rich part of the system were reported by [1992Kar1]. These are Au1–xCux + + L at ~515°C and L + -4 + ´ at ~320°C. The binary reactions p3, p4 and e2 indicated in Fig. 2 are connected with these four-phase equilibria. Liquidus Surface Based on microstructural examination and differential scanning calorimetry, a liquidus projection for the partial system below 35 at.% Au was mapped [1988Roe2] and is shown here (Fig. 3) corrected with respect to the P and E1 reactions. Isothermal Sections Figure 4 demonstrates the isothermal section at 360°C in the Au-Cu-Sn system [1992Kar1], with corrections based on the binary data quoted by [Mas2, 1993Oka]. The section features the existence of the liquid phase in two single-phase, four two-phase and two three-phase regions. The two-phase regions are marked with tie lines. The 1 phase and region which involves ordered AuCu3(I) have not been examined by experiment and are added tentatively by dashed lines. Figure 5 depicts the partial isothermal section at 170°C based on [1988Roe2]. Phase equilibria in the Sn corner up to 50 at.% Cu and 50 at.% Au as well as the homogeneity ranges of the ternary -1, -2 and -3 compounds were experimentally determined by [1988Roe2]. We added a -5 ternary phase in the section shown in Fig. 5, reported by [2002Luc]. The binary phases on the Au-Cu side are added by dashed lines according to [Mas2]. Also [1988Roe2] indicated diffusion paths which may be interpreted as tie lines showing equilibrium between the phases -1 and -2, -2 and , -1 and (Au), -1 and AuCu3(I), -1 and AuCu(I), -1 and Au3Cu, -1 and ´, -1 and J1. But there is no determination of the three phase equilibria with less than 50 at.% Sn. Temperature – Composition Sections Figure 6 demonstrates a partial vertical section AuCu-Sn [2002Luc] up to 8 at.% Sn in a temperature range from 450 to 200°C. The section reveals single-phase and multiphase regions including ordered AuCu(I), AuCu(II), -1, and -5 phases and disordered Au1–xCux solid solution. The -5 phase appears below 320°C for the Sn contents of and above 2 at.% in the alloys, and the -1 phase, below 270°C with the Sn content growing from 4 at.%. Thermodynamics [1969She] determined heats of solution of gold and copper in dilute Au-Cu-Sn alloys at 720°C, using a liquid metal solution calorimeter. Notes on Materials Properties and Applications [1974Alc] measured the chemical potentials of tin in Au1–xCux solid solutions with Cu + Au alloys using a gas-solid equilibration technique. [1984Kim] examined the microstructure, electrical resistivity, hardness and soldering contact strength of the Cu based alloys added with about 1 mass% Au and from 1 to 25 mass% Sn; [1984Kim] also examined the microstructure of thin films of these alloys obtained by e-beam evaporation onto aluminium oxide substrates. All samples were finally annealed for 30 min at 250°C. [1984Kim] pointed out that the structures of annealed thin-film and bulky alloys are similar to one another and, therefore, the thin-film properties can be predicted as measured with the bulky specimens and respectively corrected for the scale factor. [2002Huh] examined influence of small (to 1 mass%) additions of Au on the tensile properties of a Sn-0.7Cu (mass%) alloy. The best combination of strength and ductility was achieved with 0.3 mass% Au owing to solution hardening.
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141
[2004Kis] measured the electrical resistivity and solder tensile strength of solder eutectic Au-Sn alloy with 0.46 at.% Cu and shown the promise of this alloy for providing a replacement for Pb-Sn alloys in electronic assemblies. Miscellaneous [1993Zak] examined the effect of Kirkendall pore formation in the Au-Cu-Sn system on the adhesion strength of the inner lead bond contacts. The contact-diffusion layer was formed between contacting 80/20 eutectic Au-Sn alloy and Cu. The sandwich specimens were annealed at the temperatures between 125 and 200°C for up to 2000 h. Adhesion in the inner lead bond contacts was found to primarily depend on the mechanism of pore formation, growth, and coagulation. [1993Zak] reported on two new ternary phases, defined as ternary compounds. The first phase consisted of 40 to 50 at.% Cu, 40 to 50 at.% Au, and up to 20 at.% Sn and the second phase, of 15 to 20 at.% Cu, 65 to 75 at.% Au and 10 to 12 at.% Sn. Judging from the [1992Kar1, 1992Kar2, 1992Kar3] data, we can identify the ternary -1 compound with the first phase, whereas the second phase should be considered as a ternary solid solution based on the binary phase of the Au-Sn system. [1997Ven] used differential scanning calorimetry to investigate the solidification kinetics of a transient liquid phase in Au and Sn layers electroplated on Cu foil. [1997Ven] showed that this liquid phase had a gold rich eutectic composition for the used thicknesses of Sn (2 k) and Au (6 k) and a heating rate of 10 K#min–1. The solidification reaction takes place between this liquid and the excess solid gold. The solidification time is a strong function on the gold-to-tin thickness ratio and tends to infinity as the net composition approaches the phase solidus. At temperatures up to 295°C, the diffusion of copper into gold is not significant enough to affect the solidification kinetics. [2001Lee, 2003Dua] investigated the effect of the under bump metallurgy (UBM) structure on the redeposition rate of Au-containing intermetallic compounds at the solder/UBM interface during the aging treatments. Au/Ni/Cu [2001Lee] and Ni(P)/Au and Ni(V)/Cu [2003Dua] compositions were used as the UBM structures and Sn-37Pb (mass%), Sn-3.5Ag (mass%), Sn-3.5Ag- 0.7Cu (mass%), as the solder alloys. Copper dissolution from UBM into the solder, as well as the addition of Cu into the solder, were found to be effective in retarding the redeposition of the Au containing intermetallic compounds and in preventing the Au embrittlement. References [1969She]
[1971Cre] [1973Gan]
[1974Alc]
[1984Kim] [1988Roe1]
[1988Roe2]
Landolt-Börnstein New Series IV/11C3
Shen, S.S., Spencer, P.J., Pool, M.J., “Thermodynamic Analysis of Dilute Ternary Systems: III. The Au-Cu-Sn System”, Trans. AIME, 245, 1009-1013 (1969) (Thermodyn., Experimental, 6) Creydt, M., Fichter, R., “Diffusion in Galvanically Deposited Layers and Solders Between 23°C and 212°C” (in German), Metall, 25(10), 1124-1127 (1971) (Experimental, 10) Gangulee, A., Das, G.S., Bever, M.B., “An X-Ray Diffraction and Calorimetric Investigation of the Compound Cu6Sn5”, Metall. Trans., 4, 2063-2066 (1973) (Crys. Structure, Experimental, 21) Alcock, C.B., Jacob, K.T., “Solute-Solute and Solvent-Solute Interaction in -Solid Solutions of Cu + Sn and Cu + Au + Sn Alloys”, Acta Met., 22(5), 539-544 (1974) (Experimental, 22) Kim, J., “Characterization of Ternary Cu-Sn-Au Bulk Alloys and Thin Films”, J. Electron. Mater., 13(1), 191-209 (1984) (Experimental, 15) Roeder, J.F., Notis, M.R., Goldstein, J.I., “Compound Formation and Interfacial Instability in the Au-Cu-Sn System at Low Temperature”, Defect Diffus. Forum, 59, 271-278 (1988) (Phase Relations, Experimental, #, 18) Roeder, J.F., “Phase Equilibria and Compound Formation in the Au-Cu-Sn System at Low Temperature”, Ph. D. Thesis, Lehigh University (1988) (Phase Relations, Phase Diagram, Crys. Structure, Experimental, #, 88)
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142 [1989Gol]
[1990Kar] [1992Kar1] [1992Kar2] [1992Kar3]
[1993Oka] [1993Zak]
[1994Pep] [1997Ven]
[2001Lee]
[2002Huh]
[2002Luc]
[2003Dua]
[2004Kis]
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Au–Cu–Sn Goldstein, J.I., Notis, M.R., Romig, A.D., “The Use of X-Ray Microanalysis in the Study of Diffusion”, Diffus. Anal. Appl., Romig, A.D., Dayananda, M.A, (Eds.), The Minerals and Materials Society, 115-142 (1989) (Review, 54) Karlsen, O.B., Kjekshus, A., Roest, E., “The Ternary Phases in the System Au-Cu-Sn”, Acta Chem. Scand., 44, 197-198 (1990) (Phase Diagram, Phase Relations, Experimental, 7) Karlsen, O.B., Kjekshus, A., Roest, E., “The Ternary System Au-Cu-Sn”, Acta Chem. Scand., 46, 147-156 (1992) (Phase Diagram, Phase Relations, Experimental, #, *, 23) Karlsen, O.B., Kjekshus, A., Roemming, C., Roest, E., “The Crystal Structure of AuCuSn”, Acta Chem. Scand., 46, 442-445 (1992) (Crys. Structure, Experimental, 11) Karlsen, O.B., Kjekshus, A., Roemming, C., Roest, E., “The Crystal Structure of the Low-Temperature Au80-vCuvSn20 Phase”, Acta Chem. Scand., 46, 1076-1082 (1992) (Crys. Structure, Experimental, 22) Okamoto, H., “Au-Sn (Gold-Tin)”, J. Phase Equilib., 14(6), 765-766 (1993) (Phase Diagram, Phase Relations, Rewiev, #, 5) Zakel, E., “Au-Sn Bonding Metallurgy of TAB Contacts and Its Influence on the Kirkendall Effect in the Ternary Cu-Au-Sn”, IEEE Trans., Comp. Hyb. Man. Technol., 16(3), 323-332 (1993) (Experimental, 18) Peplinski, B., Zakel, E., “X-Ray Powder Diffraction Investigation of Ternary Au-Cu-Sn”, Mater. Sci. Forum, 166-169, 443-448 (1994) (Crys. Structure, Experimental, 7) Venkatraman, R., Wilcox, J.R., Cain, S.R., “Experimental Study of the Kinetics of Transsient Liquid Phase Solidification Reaction in Electroplated Gold-Tin Layers on Copper”, Metall. Mater. Trans. A, 28, 699-706 (1997) (Experimental, 22) Lee, J.-H., Park, J.-H., Shin, D.-H., Lee, Y.-H., Kim, Y.-S., “Kinetics of Au-Containing Ternary Intermetallic Redeposition at Solder/UBM Interface”, J. Electron. Mater., 30(9), 1138-1144 (2001) (Experimental, 21) Huh, S.-H., Kim, K.-S., Suganuma, K., “Effect of Au Addition on Microstructural and Mechanical Properties of Sn-Cu Eutectic Solder”, Mater. Trans., JIM, 43(2), 239-245 (2002) (Experimental, Phase Relations, Mechan. Prop., 20) Luciano, Ch.N.C., Udoh, K., Nakagawa, M., Matsuya, S., Ohta, M., “AuCu-Sn Pseudobinary Phase Diagram”, J. Alloys Compd., 337, 289-295 (2002) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, 10) Duan, N., Scheer, J., Bielen, J., van Kleef, M. “The influence of Sn-Cu-Ni(Au) and Sn-Au Intermetallic Compounds on the Solder Joint Reliability of Flip Chips on Low Temperature Co-Fired Ceramic Substrates”, Microelectron. Reliab., 43, 1317-1327 (2003) (Experimental, 15) Kisiel, R., Gasior, W., Mozer, Z., Pstrus, J., Bukat, K., Sitek, J., “(Sn-Ag)-eut + Cu Soldering Materials. Part II: Electrical and Mechanical Studies”, J. Phase Equilib. Diffus., 25(2), 122-124 (2004) (Experimental, Electr. Prop., Mechan. Prop., 5)
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143
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Au1–xCux
cF4 Fm3m Cu
Lattice Parameters Comments/References [pm] 0x1
(Au) < 1064.43
a = 407.82
at x = 0 and 25°C [Mas2]
(Cu) < 1084.87
a = 361.46
at x = 1 and 25°C [Mas2]
Sn(r) 231.97 - 13
tI4 I41/amd Sn(r)
a = 583.18 c = 318.18
at 25°C [Mas2]
Sn(l) < 13
cF8 Fm3m C(diamond)
a = 648.92
[Mas2]
, Au10Sn < 532
hP16 P63/mmc TiNi3
< 521
hP2 P63/mmc Mg
´, Au5Sn < 190
hR6 R3 Au5Sn
´, (Au1–xCux)Sn1–y
hP4 P63/mmc NiAs
8.2 to 9.1 at.% Sn [1993Oka], dissolves up to 20 at.% Cu at 360°C [1992Kar1] a = 290.2 c = 951.0
[V-C]
-
9.1 to 17.6 at.% Sn [1993Oka], dissolves up to 50 at.% Cu at 360°C [1992Kar1]
a = 509.2 c = 1433.3 -
, AuSn < 419.3 a = 432.18 c = 552.30 1, Cu6Sn5 415 - 186 a = 419.2 c = 503.7 J, AuSn2 < 309
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oP24 Pbca AuSn2
a = 690.9 b = 703.7 c = 1178.9
[1993Oka] [V-C] at 360°C, 0 x 1 and 0 y 0.17 [1990Kar, 1992Kar1]
at x = 0, y = 0 [1993Oka], dissolves up to 18 at.% Cu at 170°C [1988Roe2] [V-C2]
at x = 1, y = 0.17, [1993Oka], dissolves up to 15 at.% Au at 170°C [1988Roe2] quenched [Mas2, V-C2] [1993Oka, V-C2]
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Au–Cu–Sn
144 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, AuSn4 252 - ~60
oC20 Aba2 PtSn4
a = 650.2 b = 654.3 c = 1170.5
[1993Oka, V-C2]
1, Cu17Sn3 798 - 586
cI2 Im3m W
a = 302.61
13.1 to 16.5 at.% Sn [Mas2, V-C2]
1, Cu3Sn 755 - 520
cF16 Fm3m BiF3
a = 611.66
15.5 to 27.5 at.% Sn [Mas2] at 700°C [V-C2]
1, Cu41Sn11 590 - ~350
cF416 F43m Cu41Sn11
a = 1798.0
20 to 21 at.% Sn [Mas2, V-C2]
1, Cu10Sn3 638 - 582
hP26 P63 Cu10Sn3
a = 733.0 c = 786.4
20.3 to 22.5 at.% Sn [Mas2, V-C2]
J1, Cu3Sn < 676
oC80 Cmcm Cu3Sn a = 552.9 b = 477.5 c = 432.3
´, Cu6Sn5 < 189
h** superstructure of NiAs
a = 2087.0 c = 2508.1
24.5 to 25.9 at.% Sn [Mas2], dissolves up to 17 at.% Au at 360°C [1992Kar1] at 25 at.% Sn [V-C2]
45 at.% Sn [Mas2], dissolves up to 15 at.% Au at 170°C [1988Roe2] [1973Gan]
Au3Cu < 240
cP4 Pm3m AuCu3
a = 396.5
10 to 38.5 at.% Cu [Mas2, V-C2]
AuCu(II) < 410
oI40 Imma AuCu(II)
a = 367.6 b = 395.6 c = 397.2
38.5 to 63 at.% Cu [Mas2, V-C2]
AuCu(I) < 385
tP4 P4/mmm AuCu
a = 396.3 c = 367.1
42 to 57 at.% Cu [Mas2, V-C2]
AuCu3(II) 390 - 255(?)
tP28 P4mm PdCu3
-
66 to ? at.% Cu [Mas2]
AuCu3(I) < 390
cP4 Pm3m AuCu3
a = 374.8
67 to 81 at.% Cu [Mas2, V-C2]
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Landolt-Börnstein New Series IV/11C3
Au–Cu–Sn Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
* -1 , Au80–xCuxSn20
cP20 P4132 Mn
145
Lattice Parameters Comments/References [pm]
a = 676.2 0.2 a = 666.2 0.2 a = 673.61 0.01 a = 678.32 0.03 a = 681.29 0.01
42 x 69, at 360°C [1992Kar1,1992Kar3] at x = 45 [1992Kar3] at x = 60 [1992Kar3] at x = 47, annealed at 175°C [1994Pep] at x = 40 [1994Pep] at x = 30 [1994Pep]
* -2, AuCuSn
tI12 I4/mmm La2Sb
a = 407.8 0.1 c = 1293.8 0.1 a = 408.8 0.3 c = 1300.7 0.6
at 360°C for Au24.9Cu31.8Sn33.3 [1992Kar2] at 360°C for Au37Cu29.7Sn33.3 [1990Kar]
* -3, Au3Cu3Sn5 or AuCuSn2
-
-
[1971Cre, 1988Roe1, 1988Roe2]
* -4, Au80–xCuxSn20
cI52 I43m Cu5Zn8
* -5
31 x 37, at 360°C or 31 x 65, at 550°C [1990Kar]
oP8 Pbam Au2CuZn
Fig. 1: Au-Cu-Sn. Au-Sn phase diagram [1993Oka]
a = 918.4 0.3 a = 948.4 0.6
at x = 60, quenched from 550°C at x = 32, quenched from 360°C
a = 899.1 b = 454.2 c = 285.9
[2002Luc]
1064.43°C 1000
L
Temperature, °C
750
532
(Au) 500
521
Au
309
280
ζ
250
0
419.3
190
β
252
20
40
217 (β Sn) 13
δ
ζ´
231.9681°C
η
ε 60
(α Sn) 80
Sn
Sn, at.%
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Au–Cu–Sn
146
Cu-Sn
Au-Cu-Sn
Au-Sn
798 p1 l + (Cu) β1 755 p2 l + β1 γ1
L+β1Au1-xCux+γ1
?
640 e1 γ ε1 + l
U1
?
Au1-xCux+β1+γ1 L+Au1-xCux+γ1 ?
?
L+α+τ1
L+Au1-xCuxγ1+τ1 U2
?
?
532 p3 l + (Au) β
?
519 p4 l+βζ
L+γ1+τ3 L + γ 1 ε1 + τ 3
? 415 p5 l + ε1 η1
U3 L+γ1+τ3
L+ε1+τ3
L + γ1 τ 1 + τ 3
?
?
L + τ1 δ' +τ3
?
τ1+δ'+τ3
τ3+η1+L 279
L+δ'+τ3 L + δ' τ3 + ε
U6
τ3+δ'+ε
?
U7
280 e2 l δ' + ζ
ε+τ3+η1
L + τ3 + η η1
η+τ3+η1
P1
τ3+L+ε
τ3+η+L η+η1+L <251
227 e3 l η + Sn(r)
252 p7 l+εη L ε + τ3 + η τ3+η+ε
211
309 p6 l + δ' ε
?
U5
?
L + ε1 η1 + τ3
U4
γ3+τ1+τ3
L+τ1+τ3
L+δ1'+τ3
?
L+γ1+τ1
?
L η+η1+Sn(r)
E1 217 e4 l η + Sn(r)
E2
η+η1+Sn Fig. 2: Au-Cu-Sn. Reaction scheme
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147
Sn Fig. 3: Au-Cu-Sn. Liquidus surface for the composition range below 35 at.% Au
e3,227 p5,415
η
P1
80
η1
ε
p6,309
E1 U7
40
U6
U3
60
τ3
p2,755 p1,798
U5 40
γ1
80
60
δ
ε1 e1,640
Axes: at.%
(Sn) e4,217 p7,252
E2
20
Data / Grid: at.%
U4
τ1
U1 U2
20
β1
(Cu)
20
Cu
40
60
80
Sn
Au
Data / Grid: at.% Axes: at.%
Fig. 4: Au-Cu-Sn. Isothermal section at 360°C
L 20
80
40
60
δ
δ´
η1 60
40
τ2
L
δ 1 ε1 80
20
τ4
τ1
ζ β
AuCu3(I)
Cu
Landolt-Börnstein New Series IV/11C3
Au1-xCux 20
40
AuCu(II)
60
AuCu(I) AuCu(II)
80
Au
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Au–Cu–Sn
148
Sn
Data / Grid: at.% Axes: at.%
Fig. 5: Au-Cu-Sn. Partial isothermal section at 170°C η
20
80
ε 40
60
η'
δ
τ3
60
40
τ2 ε1
τ1
80
20
ζ´ ζ (Au)
Cu
(Cu)
Fig. 6: Au-Cu-Sn. Partial vertical section AuCu - Sn
20
40
AuCu3(I)
60
AuCu(I)
Au3Cu
80
AuCu(II)
Au
Au1-xCux
AuCu(I)+AuCu(II) 400
AuCu(I)+AuCu(II)+Au1-xCux
Temperature, °C
AuCu(II)+Au1-xCux
AuCu(I)+Au1-xCux 300
AuCu(I)+τ5+Au1-xCux
Au1-xCux+τ5
AuCu(I)+τ5
AuCu(I)
Au1-xCux+τ1+τ5 AuCu(I)+τ5+τ1
200
Au 50.00 Cu 50.00 Sn 0.00
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2
4
Sn, at.%
6
Au 46.00 Cu 46.00 Sn 8.00
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Au–In–Sn
149
Gold – Indium – Tin Katharina Haußmann Introduction [1958Sch] reported the ternary Au4In3Sn3 phase in a short communication. [1959Sch] investigated a section of the ternary system. They showed alloy equilibria after heat treatments at various temperatures for 0 to 340 d. Samples were melted in evacuated silica tubes. Solidified alloys as well as powder samples were vacuum heat-treated before they were examined by X-ray diffraction methods and metallography, respectively. One half of the 20 alloys were heat treated at 250°C, the other half was examined as cast or heat treated between 100 and 350°C. [1968Sch] investigated the crystallographic structure of an alloy with the composition 50Au-37In-13Sn (at.%) and [1959Bur] that of 75Au-22In-3Sn (at.%). The alloy investigated by [1959Bur] contained J' with some . The same alloy was quoted by [1959Sch] to be single-phase J' as a result of X-ray powder diffraction and 95% J' + 5% AuSn in the grain boundaries according to metallographic examination of a sample heat treated at 250°C for 48 d. The Au50In37Sn13alloy [1968Sch] is probably within the homogeneity range of AuIn. [1985Hum] worked on the section AuIn-AuSn and found a two-phase region below 325°C and the section not to be quasibinary. There is a considerable solubility of Sn in AuIn (10 at.% at 250°C) and extensive solubility of In in AuSn (17 at.% at 250°C). The ternary compound, Au4In3Sn3, melts congruently at 431°C. Binary Systems The binary systems Au-In, Au-Sn and In-Sn are accepted from [Mas2]. Solid Phases Crystal structure data of the ternary Au4In3Sn3 compound and the binary phases stable at 250°C and below are given in Table 1. Isothermal Sections Figure 1 shows the isothermal section at approximately 250°C. At this temperature the In-Sn binary alloys are all molten, the Au-In alloys have a L+AuIn2 region ranging up to 98.8 at.% In and Au-Sn alloys have an L + AuSn4 region from 80 to 89 at.% Sn, followed by a liquid region up to Sn. Therefore, the equilibria between solid phases on the Au-poor side of the AuIn2-AuSn4 joint, given by [1959Sch], present at 100°C section. This part of the section is omitted in Fig. 1. The Au9In phase forms a continuous series of solid solution (1) with the Au10Sn phase, over the temperature range of the Au10Sn stability. Au10Sn was not detected in the isothermal section. Therefore it is probable that it transforms to (Au) + at a temperature higher than 250°C. The phase is a continuous series of solid solution between Au17Sn3 and Au4In phases. Instead of one three-phase region J'+1/+AuSn as given by [1959Sch] the isothermal section presented in Fig. 1 shows two three-phase regions: J' +1 + AuSn and 1 + + AuSn. References [1958Sch]
[1959Bur]
Landolt-Börnstein New Series IV/11C3
Schubert, K., Breimer, H., Gohle, R., Lukas, H.L., Meissner, H.G. Stolz, E., “Some Structural Results on Metallic Phases III” (in German), Naturwiss., 45, 360-361 (1958) (Crys. Structure, Experimental, 2) Burkhardt, W., Schubert, K., “On Brass Like Phases with A3 Related Structure” (in German), Z. Metallkd., 50, 442-452 (1959) (Phase Diagram, Crys. Structure, Experimental, 40)
MSIT®
Au–In–Sn
150 [1959Sch]
[1968Sch]
[1985Hum]
Schubert, K., Breimer, H., Gohle, R., “On the Constitution of the Gold - Indium, Gold - Tin, Gold - Indium - Tin and Gold - Tin - Antimony Systems” (in German), Z. Metallkd., 50, 146-155 (1959) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, 30) Schubert, K., Bhan, S., Biswas, T.K., Frank, K., Pandau, P.K., “Some Structural Data of Metallic Phases” (in German), Naturwiss., 55, 542-543 (1968) (Crys. Structure, Experimental, 1) Humpston, G., “The Constitution of Some Ternary Au-Based Solder Alloys”, Ph.D. Thesis, Brunel University (1985) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, 80)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Au) < 1064.43
cF4 Fm3m Cu
a = 407.82
at 25°C [Mas2]
(Sn)(r) 231.97 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
at 25°C [Mas2]
(Sn)(l) < 13
cF8 Fd3m C (diamond)
a = 648.92
[Mas2]
(In) < 156.63
tI2 I4/mmm In
a = 325.3 c = 494.70
at 25°C [Mas2]
1
hP16 P63/mmc TiNi3
Au9In < 649.25 Au10Sn < 649.25
a = 410.60
12 to 14.3 at.% In [Mas2] at 20°C [Mas2, V-C2]
-
[Mas2]
1, Au7In2 < 295
hP26 P3 Cu10Sb3
a = 1054.5 c = 476.9
21.7 to 22.5 at.% In [Mas2, V-C2]
J', Au3In 339.5
oP8 Pmmm TiCu3
a = 586 b = 474 c = 516
24.5 to 25 at.% In [Mas2, V-C2]
a = 590 b = 476 c = 518
Au75In22Sn3 [1959Bur]
a = 1221.5 c = 850.9
29.8 to 30.6 at.% In [Mas2, V-C2]
', Au7In3 < 374.6
MSIT®
hP60 P3 Au7In3
Landolt-Börnstein New Series IV/11C3
Au–In–Sn
151
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
5, Au3In2 457.5 - 224.3
hP5 P3m1 Ni2Al3
a = 453.7 c = 565.9
35.3 to 39.5 at.% In [Mas2, V-C2]
AuIn < 509.6
aP*
a = 430 b = 1059 c = 356 = 90.54° = 90.0° = 90.17°
[Mas2, V-C2]
AuIn2 < 540.7
cF12 Fm3m CaF2
a = 650.2
[Mas2, V-C2]
’, Au5Sn < 195
hR6 R3 Au5Sn
a = 509.2 c = 1433.3
16.7 at.% Sn [Mas2, V-C2]
, AuSn < 419.3
hP4 P63/mmc NiAs
a = 432.18 c = 552.30
50 to 50.5 at.% Sn [Mas2, V-C2]
J, AuSn2 < 309
oP24 Pbca AuSn2
a = 690.9 b = 703.7 c = 1178.9
66.7 at.% Sn [Mas2, V-C2]
, AuSn4 252 - 50
oC20 Aba2 PtSn4
a = 651.24 b = 651.62 c = 1170.65
80 at.% Sn [Mas2, V-C2], lattice parameters at 25°C
, In3Sn < 143
tI2 I4/mmm In
a = 346.3 c = 440.4
12 to 44 at.% Sn [Mas2, V-C2], lattice parameter at 24°C
, InSn4 < 224
hP1 P6/mmm BiIn
a = 320.8 c = 299.7
72 to ? at.% Sn [Mas2, V-C2], lattice parameter at –10°C
hP2 P63/mmc Mg
a = 293.05 c = 477.01
10 to 18.5 at.% Sn [Mas2, V-C2]
a = 291.57 c = 479.30
13 to 23 at.% In [Mas2, V-C2]
Au17Sn3 Au4In * -1, Au4In3Sn3 < 431
hP10 P63/mmc Pt2Sn3
a = 449.9 c = 1304.7
[1958Sch], melting point [1985Hum]
* -2?, Au5In4Sn
oC8 Cmcm Tl I
a = 360 b = 1049 c = 428
Au50In37Sn13 [1968Sch], probably (Au,In) solid solution and not a ternary compound
Landolt-Börnstein New Series IV/11C3
MSIT®
Au–In–Sn
152
Sn
Data / Grid: at.%
Fig. 1: Au-In-Sn. Partial isothermal section at 250°C
Axes: at.%
AuSn 4 20
80
AuSn2+AuSn4+τ AuSn2 AuSn2+AuSn+τ
40
AuSn4+τ +AuIn2
60
AuSn 60
40
AuSn+β 1+ε
(Au)+ζ+α 1
80
20
ε+γ+AuSn
AuSn+ζ +β
ζ
(Au)
Au
MSIT®
τ1
α1
AuSn+ +τ 1+AuIn
AuIn
AuIn2
γ' 40ψ γ '+ψ +AuIn 60 β 1 ε' AuIn γ +AuSn+AuIn τ 1+AuIn+AuIn2 2
20
80
In
Landolt-Börnstein New Series IV/11C3
B–Cr–Ni
153
Boron – Chromium – Nickel Anatoliy Bondar Introduction This ternary system is of interest because the Cr-Ni brazing alloys are used for joining silicon nitride ceramic bodies as well as for joining silicon nitride to Ni based superalloys, and small additions of B depresses the melting temperatures to below 1100°C. Data on the phase equilibria were presented as isothermal sections at 800°C [1972Che] and 1000°C [1960Kol, 1962Kol, 1974Lug], a partial liquidus surface projection within the composition region bounded by Cr, CrB, Ni2B and Ni [1985Omo], and a portion of the vertical section from Ni to CrB2 [1992Yan] (Table 1). By means of X-ray diffraction (XRD) and microstructural observation, [1972Che] determined the isothermal section at 800°C and established the existence of two ternary compounds, Cr3NiB6 and Cr2Ni3B6. The samples were prepared from powdered metals of electrolytic Cr (99.5%), Ni (99.9%) and B (99.3%), placed in Al2O3 crucibles, heated in a vacuum furnace to 1400°C, and then annealed at 800°C for at least 300 h. [1962Kol] determined the partial isothermal section for the field Cr-CrB-Ni2B-Ni at 1000°C using XRD, metallography, measurements of microhardness and chemical analysis. The alloys were prepared from pure elements and were then annealed at 1000°C for 250 h. [1974Lug] established the isothermal section at 1000°C by a combination of XRD and metallographic examinations. The samples were prepared from powders of Ni (purity of 99.97 or 99.5%) and Cr borides (purity of 99.8%) by induction melting in Al2O3 crucibles and then annealing at 1000°C for 200 h. At this temperature, two ternary compounds, Cr3NiB6 and Ni3Cr2B6 were confirmed. [1999Ohs] performed a numerical modeling for the transient liquid phase (TLP) bonding process of Ni using a B-Cr-Ni filler by combining thermodynamic calculation with diffusion analysis. Their calculated results agree well with their own experimental data. [1992Yan] constructed the Ni-CrB2 vertical section by using metallography, XRD, measurements of hardness and microhardness, and abrasive wear tests. The samples were prepared from powders of electrolytic Ni and CrB2 in ceramic crucibles under vacuum (0.1 - ~0.01 Pa) with the furnace temperature kept at 1250°C for 10 min. A thermodynamic assessment of the B-Cr-Ni system was performed by [2002Cam], who used the experimental phase equilibrium data at 1000°C [1974Lug] and 800°C [1972Che] for the optimization of the model parameters. Binary Systems The B-Cr phase diagram is based mainly on the experimental data of [1969Por, 1972Por]. The phase diagram from [1972Por] was accepted in the present work (Fig. 1). Subsequently, this binary system was reviewed by [1986Lia], and a thermodynamic assessment was conducted by [2002Cam]. [1969Por] reported that Cr2B and CrB2 possess remarkable homogeneity ranges with the characteristics of B and Cr deficiency, respectively. Such a spread could be expected for the homogeneity range of CrB also, as it is analogous to other borides with the same crystal structure and showing noticeable homogeneity ranges. Furthermore, a Cr2B3 boride was separated from the Al-B-Cr solid-liquid sample kept at 1500°C for 10 h [1987Oka]. Although the Cr2B3 crystals contained 3.2 mass% Al, EPMA measurement showed the Al solubility to be about the detection threshold estimated as 0.05 mass%. So, the Cr2B3 boride should be inserted in the phase diagram, as was the CrB4 boride, found after annealing at 1350 and 1400°C in [1968And] (after annealing at 1600°C it was absent). The B-Ni phase diagram is accepted after [1965Sch] who obtained detailed experimental data. In [1965Sch], as well as in later thermodynamic assessments [1993Tep, 1999Cam], all the nickel borides were presented as practically stoichiometric compounds, and Ni3B4 was treated as two phases Ni3+xB4–x (orthorhombic) and Ni3–xB4+x (monoclinic) with composition 41.4 and 43.6 at.% B (i.e. Ni3.1B2.9 and Ni2.95B4.05), respectively. Nevertheless, the lattice parameters of Ni3+xB4–x(o) reported in [1969Lun] are different for the Ni rich and B rich compositions, and [1965Sch] proposed its homogeneity field extension to be about 0.2 at.%. Landolt-Börnstein New Series IV/11C3
MSIT®
154
B–Cr–Ni
The Cr-Ni phase diagram was taken from [1995Udo], where a number of experimental data were carefully evaluated. The version of [1995Udo] nearly coincides with the phase diagram in [Mas2]. Solid Phases All the unary, binary and ternary phases are presented in Table 2. It seems that the “Cr2NiB4” ternary phase reported in [1956Pos] corresponds to the ternary compound Cr2Ni3B6 found in [1972Che, 1974Lug]. It is not impossible that all the ternary phases found are terminal solid solutions based on isostructural binary Cr5B6 and Cr2B3, taking into account the closeness of their lattice parameters and compositions. The solubilities of the third element in the chromium and nickel borides were accepted as negligibly small in [1960Kol, 1962Kol, 1972Che, 1974Lug]. Wet chemical analyses of extracted borides showed that the mutual solubility between Cr and Ni borides (Cr2B, Cr5B3, CrB, Ni3B, and Ni2B) is negligible, not exceeding 0.2 at.% [1960Kol]. However, [1977Spr] claimed that EPMA reveals 5 at.% Ni solubility in Cr2B, resulting in no change in the lattice parameters. [1960Kol, 1962Kol] reported that the solubility of B in the Ni based phase is about 0.02 to 0.04 at.% at 1000°C, which coincides with the computed solubility of 0.023 at.% B [1993Tep]. Quasibinary Systems Two quasibinary eutectic reactions should be added, based on the four-phase invariant equilibria shown in the liquidus given by [1985Omo] (Figs. 2 and 3). The first of them (e5(max)) is L (Ni) + Cr2B at a temperature between 1258°C (E1), and ~1390°C, which is the melting temperature of Cr~35Ni~65. The temperature for the second reaction (e6(max)), L Ni2B + CrB, is estimated to be about 1110°C based on the liquidus isotherms [1985Omo] (higher than 1096°C for E2 and lower than the melting point of Ni2B at 1156°C). Invariant Equilibria Invariant equilibria in the composition range bounded by Cr, CrB, Ni2B and Ni are presented in Table 3 [1985Omo]. Liquidus Surface Figure 3 presents the partial liquidus surface projection taken from [1985Omo]. The alternative version calculated in [2002Cam] for the whole concentration range proposes extended regions of primary separation of chromium borides Cr5B3, CrB, Cr3B4, and CrB2 up to less than 1 at.% Cr for Cr3B4 and CrB2. In the present assessment, this version was not adopted because it is in contradiction with the phase constituents of the Ni rich eutectic alloys. The widely used commercial alloys are well known to contain the three-phase eutectic (Ni) + Ni3B + CrB, rather than (Ni) + Ni3B + Cr5B3 suggested by [2002Cam]. The liquidus surface calculated in [2002Cam] also disagrees with the experimental liquidus temperatures [1974Lug] and results in an essentially larger extension of the primary regions of Ni3B and Ni2B as compared with [1985Omo]. Isothermal Sections Isothermal sections in the Ni rich region for temperatures from 1012 to 1200°C (Figs. 4a - 4e), calculated in [1999Ohs], show phase equilibria involving the melt, Ni3B and the (Ni) phase, which are of interest for treatment of brazing processes. The isothermal sections likely reflect the phase equilibria between the above mentioned phases. However, other phases are not included and the melt was presented as stable at 1012°C, which is noticeably lower that the most low-melting eutectic (Ni) + Ni3B + CrB, 1050°C after [1985Omo]. Figures 5 and 6 present isothermal sections at 1000 [1962Kol, 1974Lug] and 800°C [1972Che]. It is necessary to note that the data of the three studies are in good agreement with each other. For the 1000°C section, the results of [1974Lug] were proffered for high B contents. While in [1960Kol, 1962Kol], the equilibria with the (Ni) phase were studied more carefully using the extraction of borides by electrochemical dissolution of the metal matrix followed by XRD and wet chemical analyses. Phase equilibria with the MSIT®
Landolt-Börnstein New Series IV/11C3
B–Cr–Ni
155
participation of Ni4+xB3–x(o) and CrB4 were added as they are present in [1972Che] for 800°C, as well as equilibria with Cr2B3 (just in the Cr2B3 + Cr3NiB6 phase field there is a two-phase sample in [1974Lug]). Temperature – Composition Sections The Ni-CrB2 vertical section from 100 to 40 at.% Ni is presented in Fig. 7 after [1992Yan]. Thermodynamics Gibbs energies of the ternary phases Cr3NiB6 and Cr2Ni3B6 were optimized in [2002Cam] as (in J·mol–1): GCr3NiB6 = –37450 + 3.45·T + 0.3·GCr + 0.1·GNi + 0.6·GB; GCr2Ni3B6 = –34150 + 4.08575·T + 0.18·GCr + 0.27·GNi + 0.54·GB. Notes on Materials Properties and Applications A summary of studies of the materials properties and applications related to the B-Cr-Ni based brazing alloys is presented in Table 4. Miscellaneous [1991Har] studied reversible structural relaxation in amorphous alloys of compositions Cr14Ni68B18 and Cr14Ni68B18 using differential scanning calorimetry (DSC) and electrical resistance. [1966Yas] investigated the wetting of CrB2 by molten Ni under vacuum and Ar. The contact angle was measured to be 40°C at the temperature close to the Ni melting point. In [1997Hyo], an aluminium oxide intermediate layer was applied as a diffusion barrier to the brazing of a stainless steel using a filler of evaporated B-Cr-Ni film. It was found that the intermediate layer was effective in reducing the diffusion of Ni and B in the filler into the substrate. References [1956Pos] [1960Kol]
[1962Kol]
[1965Sch] [1966Yas]
[1968And] [1969Lun]
[1969Por]
Landolt-Börnstein New Series IV/11C3
Post, B., Pipitz, E., Herz, W.H., “A New Ternary Boride: Cr4NiB2”, Powder Met. Bull., 7(3-6), 149-152 (1956) (Crys. Structure, Experimental, 2) Kolomytsev, P.T., “Interaction of Boron with Chromium in the Ternary Nickel Based Alloys” (in Russian), Issled. po Zharoproch. Splavam, 6, 180-186 (1960) (Phase Diagram, Crys. Structure, Experimental, #, 4) Kolomytsev, P.T., “Investigation of the Structure of Alloys in the B-Cr-Ni System” (in Russian), Doklady Akad. Nauk SSSR, 144, 112-114 (1962) (Phase Diagram, Crys. Structure, Experimental, #, *, 6) Schöbel, J.D., Stadelmaier, H.H., “The Binary Nickel-Boron System” (in German), Z. Metallkd., 156, 856-859 (1963) (Experimental, Phase Diagram, Phase Relations) Yasinskaya, G.A., “The Wetting of Refractory Carbides, Borides and Nitrides by Molten Metals”, Sov. Powder Metall Met. Ceram., 43(7), 557-569 (1966), translated from Poroshk. Metall., (7), 53-56 (1966) (Experimental, Interface Phenomena, 5) Andersson, S., Lundström, T., “The Crystal Structure of CrB4”, Acta Chem. Scand., 22, 3103-3110 (1968) (Crys. Structure, Experimental, 8) Lunström, T., “Preparation and Crystal Chemistry of Some Refractory Borides and Phosphides”, Arkiv Kemi, Mineral. Geol., 31(19), 227-266 (1969) (Crys. Structure, Experimental, Review, 183) Portnoy, K.I., Romashov, V.M., Romanovich, I.V., “The Phase Diagram of the System Chromium-Boron” (in Russian), Poroshk. Metall, (4), 51-57 (1969) (Phase Diagram, Experimental, #, 19)
MSIT®
156 [1972Che]
[1972Por]
[1974Lug]
[1976Lun] [1977Spr]
[1981Cre]
[1985Omo]
[1986Lia]
[1987Oka]
[1991Har]
[1992Yan]
[1993Tep]
[1995Udo]
[1996Kay] [1996Tun]
[1997Hyo]
[1999Cam]
MSIT®
B–Cr–Ni Chepiga, M.V., Krivutskii, V.P., Kuz’ma, Yu.B., “The System Cr-Ni-B”, Inorg. Mater., 8(6), 928-932 (1972), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 8(6), 1059-1064, (1972) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, #, *, 16) Portnoy, K.I., Romashov, V.M., “Binary Phase Diagrams of Some Elements with Boron (Review)” (in Russian), Poroshk. Metall, (5), 48-56 (1972) (Phase Diagram, Phase Reltaions, Review, #, 44) Lugscheider, E., Knotek, O., Reimann, H., “The Ternary System Nickel-Chrom-Bor” (in German), Monatsh. Chem., 105(1), 80-90 (1974) (Phase Diagram, Crys. Structure, Experimental, #, *, 32) Lunström, T., Tergenius, L.E., “On the Solid Solution of Copper in -Rhombohedral Boron”, J. Less-Common Met., 47, 23-28 (1976) (Crys. Structure, Experimental, 10) Sprenger, H., Richter, H., Nickl, J.J., “Borides and Silicides as Reinforcing Phases in Eutectic High Temperature Composites” (in German), Z. Metallkd., 68(4), 241-252 (1977) (Phase Relation, Experimental, 52) Crespo, A.J., Tergenius, L.-E., Lundström, T., “The Solid Solution of 4d, 5d and Some p Elements in Rhomhedral Boron”, J. Less-Common Met., 77, 147-150 (1981) (Crys. Structure, Experimental, 12) Omori, Sh., Koyama, K., Hashimoto, Y., Yamada, K., “Liquidus Surface of B-Cr-Ni System” (in Japanese), J. Jpn. Inst. Met., 49(11), 935-939 (1985) (Phase Diagram, Experimental, #, *, 7) Liao, P.K., Spear, K.E., “The B-Cr (Boron-Chromium) System”, Bull. Alloy Phase Diagrams, 7(3), 232-237 (1986) (Crys. Structure, Phase Diagram, Thermodyn., Review, #, 39) Okada, S., Atoda, T., Higashi, I., “Structural Investigation of Cr2B3, Cr3B4, and CrB by Single-Crystal Diffractometry”, J. Solid State Chem., 68, 61-67 (1987) (Crys. Structure, Experimental, 9) Haruyama, O., Asahi, N., “Reversible Short-Range Ordering Due to the Structural Relaxation in Amorphous B-Cr-Ni Alloys”, J. Mater. Sci., 26, 1851-1855 (1991) (Morphology, Experimental, 22) Yanenskii, V.N., Guslienko, Yu.A., Fedorchenko, I.M., “Structure and Properties of Nickel-Chromium Diboride Alloys”, Sov. Powder Metall. Met. Ceram., (1), 31-35 (1992), translated from Poroshk. Metall., 1(349), 32-37 (1992) (Morphology, Phase Relations, Experimental, #, 8) Teppo, O., Taskinen, P., “Thermodynamic Assessment of Ni-B Phase Diagram”, Mat. Sci. Techn., 9, 205-212 (1993) (Thermodyn., Phase Diagram, Phase Relations, Calculation, Assessment, #, 41) Udovskii, A.L., “On the Theory of the Temperature Behavior of Heat Capacity in Two-component Two-phase Alloys: Application to the Ni-Cr System”, Phys. Dokl., 40(11), 574-578 (1995) (Thermodyn., Phase Diagram, Calculation, Assessment, #, 13) Kayser, G.F., Kayser, F.X., “Ni3B: Powder Diffraction Pattern and Lattice Parameters”, J. Alloys Compd., 233, 74-79 (1996) (Crys. Structure, Experimental, 13) Tung, S.K., Lim, L.C., Lai, M.O., “Solidification Phenomena in Nickel Base Brazes Containing Boron and Silicon”, Scr. Mater., 34(5), 763-769 (1996) (Morphology, Experimental, 16) Hyodo, T., Toyoda, T., Kibe, H., Kozaki, J., Kabasawa, M., Tamura, M., “Effect of Intermediate Layer of Diffusion Barrier on Tensile Strength of Brazed Joint of Stainless Steel Coated with B-Cr-Ni Alloy Filler” (in Japanese), Tetsu to Hagane, 83(3), 199-204 (1997) (Experimental, Mechan. Prop., 9) Campbell, C.E., Kattner, U.R., “A Thermodynamic Assessment of the Ni-Al-B System”, J. Phase Equilib., 20(5), 485-496 (1999) (Phase Diagram, Phase Relations, Assessment, Calculation, #, 50) Landolt-Börnstein New Series IV/11C3
B–Cr–Ni [1999Ohs]
[2002Cam]
[2003Lia]
157
Ohsasa, K., Shinmura, T., Narita, T., “Numerical Modeling of the Transient Liquid Phase Bonding Process of Ni Using Ni-B-Cr Ternary Filler Metal”, J. Phase Equilib., 20(3), 199-206 (1999) (Phase Diagram, Phase Relations, Calculation, Experimental, #, 21) Campbell, C.E., Kattner, U.R., “Assessment of the Cr-B System and Extrapolation to the Ni-Al-Cr-B Quaternary System”, Calphad, 26(3), 477-490 (2002) (Phase Diagram, Assessment, Calculation, #, 44) Liang, Y.N., Osendi, M.I., Miranzo, P., “Joining Mechanism in Si3N4 Bonded with a B-Cr-Ni Interlayer”, J. Eur. Ceram. Soc., 23(3), 547-553 (2003) (Morphology, Experimental, 25)
Table 1: Investigations of the B-Cr-Ni Phase Relations and Structures Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[2002Cam]
Thermodynamic calculation (Thermocalc software) using a regular solution model for melt and sublattice model for solid phases. The B-Cr-Ni parameters determined by the PARROT optimization and trial-and-error method.
The B-Cr-Ni liquidus surface and isothermal sections at 1000 and 800°C were calculated for the whole concentration region.
[1999Ohs]
Diffusion analysis and thermodynamic Isothermal sections in Ni rich region calculation with Thermocalc software. A join of showed phase equilibria of melt, Ni3B and (Ni) phase were presented. Ni blocks bonded with a Ni-15.2Cr-4.0B (mass%) filler metal of 30 m thick was studied by SEM/EPMA.
[1992Yan]
Optical metallography, XRD, measurements of hardness (HRC), microhardness and relative wear were applied for alloys melted in resistance furnace at 1250°C in vacuum using ceramic crucibles from powders of electrolytic Ni and CrB2.
Morphology and properties for alloys of the Ni-CrB2 section from 3 to 23 mass% CrB2, the Ni-CrB2 vertical section from 100 to 40 at.% Ni were presented.
[1985Omo]
As-cast and annealed (at 1000°C for 556 h) alloys were studied by optical metallography, XRD and DTA.
The liquidus surface projection was presented for the field Cr-CrB-Ni2B-Ni.
[1977Spr]
Samples melted in arc furnace and zone smelted Search of eutectics was carried out in were studied by DTA, XRD and metallography sections from Ni to chromium borides (optical microscopy and SEM/EPMA). Cr2B, Cr3B5 and CrB.
[1974Lug]
Optical metallography, XRD and DTA were applied for alloys melted in an induction furnace at 1700°C in Al2O3 crucibles and then annealed at 1000°C 20 to 200 h.
The isothermal section at 1000°C in the whole concentration range and temperatures of melting for alloys up to 35 at.% B and ~85 at.% Ni.
[1972Che]
XRD and optical metallography were applied for alloys melted in an arc furnace, heated to 1400°C in a resistance furnace (in Al2O3 crucibles) and then annealed at 800°C not less than 300 h.
The isothermal section at 800°C was presented for the whole concentration range.
Landolt-Börnstein New Series IV/11C3
MSIT®
B–Cr–Ni
158 Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1962Kol]
XRD, optical metallography and measurements The isothermal section at 1000°C was of microhardness were applied for alloys melted presented for the Cr-CrB-Ni2B-Ni field. from pure elements and then annealed at 1000°C 250 h. Borides extracted by electrochemical solution of metal matrix were examined by XRD and wet chemical analysis.
[1960Kol]
The same as [1962Kol]
The isothermal section at 1000°C was presented for the Ni rich portion (not exceeding 30 mass% Cr and 5 mass% B).
[1956Pos]
Powder XRD of samples hot pressed at 1500-1700°C or sintered at 1650°C in purified hydrogen.
Cr3B4 and new ternary phase Cr2NiB4 were identified.
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Cr) < 1863
cI2 Im3m W
a = 288.48
at 25°C [Mas2]
(Ni) < 1455
cF4 Fm3m Cu
a = 352.3
at 25°C [Mas2]
(B) < 2092
hR105 or hR111 B
a = 1093.02 c = 2381.66
pure B (99.9999%) [1976Lun]
a = 1092.65 0.04 arc melted crystalline B [1981Cre] c = 2380.96 0.13 Cr2B < 1850
oF40 Fddd Mn4B
a = 146.92 b = 739.9 c = 426.6
[1974Lug]
Cr5B3 < 1890
tI32 I4/mcm Cr5B3
a = 546.8 c = 1008.9
[1974Lug]
CrB < 2090
oC8 Cmcm CrB
a = 297.5 b = 786.6 c = 293.4
[1974Lug]
Cr3B4 < 2080
oI14 Immm Ta3B4
a = 294.6 b =1301.8 c = 296.4
[1974Lug]
MSIT®
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B–Cr–Ni
159
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
Cr2B3 1500
oC20 Cmcm V2 B 3
a = 302.64 0.05 b = 1811.5 0.4 c = 295.42 0.04
[1987Oka]
CrB2 < 2200
hP3 P6/mmm AlB2
a = 297.3 c = 307.0
[1974Lug]
CrB4 1400
oI10 Immm CrB4
a = 474.4 b = 547.7 c = 286.6
[1968And]
Ni3B < 1156
oP16 Pnma Fe3C
a = 521.95 0.05 b = 661.64 0.06 c = 439.12 0.04
[1996Kay]
Ni2B < 1125
tI12 I4/mcm Al2Cu
a = 499.1 c = 424.6
[Mas2, V-C2]
Ni4+xB3–x(o) < 1025
oP28 Pnma Ni4B3
a = 1195.3 b = 298.1 c = 656.9 a = 1197.3 b = 298.5 c = 658.4
x 0.1 Ni rich
B rich [1969Lun] x 0.05 [1969Lun]
Ni4–xB3+x(m) < 1031
mC28 C2/c Ni4B3
NiB < 1035
oC8 Cmcm CrB
a = 292.9 b = 739.2 c = 296.1
[Mas2, V-C2]
* -1, Cr3NiB6 1000
oC20 Cmcm V2 B 3
a = 303.4 0.3 b = 1811 2 c = 295.6 0.3
[1972Che]
a = 300.8 b = 1770.5 c = 295.6
[1974Lug]
Landolt-Börnstein New Series IV/11C3
a = 642.8 b = 488.0 c = 781.9 = 103.32°
MSIT®
B–Cr–Ni
160 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
* -2, Cr2Ni3B6 1000
oI14 Immm Ta3B4
a = 297.1 0.3 b = 2034 2 c = 301.1 0.3
[1972Che]
a = 293.1 b = 2063.1 c = 297.7
[1974Lug]
a = 596 b = 1267 c = 605
[1956Pos]
* -3, Cr2NiB4 1700
o** ordered Ta3B4 type
Table 3: Invariant Equilibria T [°C]
Reaction
Type
Phase
Composition (at.%) B
Cr
Ni
L (Ni) + Cr2B
1258 1390
e5(max)
L*
8.6
43.2
48.2
L (Ni) + (Cr) + Cr2B
1258
E1
L
4.3
57.9
37.8
L + Cr2B (Ni) + Cr5B3
1220
U1
L
12.3
35.3
52.4
L Ni3B + CrB
~1110
e6(max)
L*
32.5
15.5
52
L + Cr5B3 (Ni) + CrB
1096
U2
L
17.3
21.2
61.5
L Ni3B + Ni2B + CrB
1096
E2
L
35.6
14.7
49.7
L (Ni) + Ni3B + CrB
1050
E3
L
18.3
16.1
65.6
* - compositions are estimated from Fig. 3
Table 4: Investigations of the B-Cr-Ni Materials Properties Reference
Method/Experimental Technique Type of Property
[2003Lia]
XRD and SEM/EDX
A commercially available alloy foil Ni-14.2Cr-3.6B (mass%) named Ni-Flex95 (USA), consisting of (Ni) + Ni3B + CrB, was successfully used for joining a Si3N4 ceramic.
[1997Hyo]
SEM
An alloy Ni-15Cr-4B (mass%) exhibits increase in the tensile strength of brazed join of stainless steel as the filler thickness increases from 7 to 10 m. An intermediate Al2O3 coating the steel was found to be effective in increasing the tensile strength of the brazed join.
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161
Reference
Method/Experimental Technique Type of Property
[1996Tun]
SEM/EDS
A join of pure nickel plates brazed with an alloy Nicrobraz 150, Ni-15Cr-3.5B (mass%) contains the three-phase eutectic of nickel, nickel boride and chromium boride phases.
[1992Yan]
Optic metallography, XRD, measurements of hardness (HRC), microhardness and relative wear.
Morphology, Rockwell hardness and relative wear were studied for alloys the Ni-CrB2 section from 3 to 23 mass% CrB2.
Fig. 1: B-Cr-Ni. The B-Cr phase diagram
2250
2200 2090
2000
Temperature, °C
57
54
36 1863°C
31.5
1750
1890 1850 Cr5B3
83
1830 ~98
CrB2 (B)
CrB
1600 ~1
2092°C
2080 2050
16 Cr2B
1500
(Cr)
Cr2B3 CrB4
Cr3B4 1250
Cr
80
60
40
20
B
Cr, at.%
Landolt-Börnstein New Series IV/11C3
MSIT®
B-Cr
B-Cr-Ni
2080 p1 l + CrB2 Cr3B4
?
2050 e1 l CrB + Cr3B4
?
B-Ni
162
MSIT®
Cr-Ni
1890 p2 l + CrB Cr5B3 1850 p3 l + Cr5B3 Cr2B 1830 e2 l CrB2 + (B)
?
1600 e3 l (Cr) + Cr2B
1258-1390 e5(max) L (Ni) + Cr2B 1258
L (Ni) + (Cr) + Cr2B
(Cr)+(Ni)+Cr5B3 L+(Ni)+Cr5B3 1096
1220
Cr5B3+(Ni)+CrB
E1
L + Cr2B (Ni) + Cr5B3
Cr2B+(Ni)+Cr5B3
L + Cr5B3 (Ni) + CrB
B–Cr–Ni
1336 e4 l (Ni) + (Cr)
U2
U1
? L+Ni2B+CrB 1096
~1110 e6(max) L Ni3B + CrB
L Ni3B + Ni2B + CrB
L+(Ni)+CrB 1050
Ni3B+Ni2B+CrB L (Ni) + Ni3B + CrB (Ni)+Ni3B+CrB
Landolt-Börnstein New Series IV/11C3
Fig. 2: B-Cr-Ni. Partial reaction scheme
E2 E3 ?
1111 e7 l Ni3B + Ni2B 1093 e8 l Ni3B + Ni2B 1035 p4 l + (B) NiB
?
1025 p5 l+Ni4-xB3+xNi4+xB3-x
?
1018 e9 l Ni2B + Ni4+xB3-x
?
1018 e10 l Ni4-xB3+x+ NiB
B–Cr–Ni
163
B
Data / Grid: at.% Axes: at.%
Fig. 3: B-Cr-Ni. Partial liquidus surface projection 20
80
40
60
60
40
p2 p3
E2
1327°C
CrB
Ni B 2
e6 Cr5B3
80
e3
1127
e7
1127
1227
Ni3B
E3
20
e8
Cr2B U2 E1
(Cr) 20
Cr
40
e5
Fig. 4a: B-Cr-Ni. Phase equilibria of melt with (Ni), Ni3B and some other phases at 1200°C
1327 80
0.00 40.00 60.00
10
1227
(Ni)
60
e4 Cr Ni B
U1
1127
Ni
Data / Grid: at.% Axes: at.%
50
20
40
30
30
L 40
20
50
Cr Ni B Landolt-Börnstein New Series IV/11C3
60.00 40.00 0.00
10
50
60
70
80
90
(Ni)
Ni
MSIT®
B–Cr–Ni
164
Cr Ni B
Fig. 4b: B-Cr-Ni. Phase equilibria of melt with (Ni), Ni3B and some other phases at 1100°C
0.00 40.00 60.00
10
Data / Grid: at.% Axes: at.%
50
20
40
30
30
L 40
20
50
Cr Ni B
10
50
60.00 40.00 0.00
60
70
80
90
Ni
(Ni) Cr Ni B
Fig. 4c: B-Cr-Ni. Phase equilibria of melt with (Ni), Ni3B and some other phases at 1077°C
0.00 60.00 40.00
10
Data / Grid: at.% Axes: at.%
30
Cr10Ni60B30 Ni3B 20
20
L
30
Cr Ni B
MSIT®
40.00 60.00 0.00
10
70
80
90
(Ni)
Ni
Landolt-Börnstein New Series IV/11C3
B–Cr–Ni Cr Ni B
Fig. 4d: B-Cr-Ni. Phase equilibria of melt with (Ni), Ni3B and some other phases at 1027°C
165 0.00 60.00 40.00
10
Data / Grid: at.% Axes: at.%
30
Ni3B 20
20
L 30
Cr Ni B
10
70
40.00 60.00 0.00
80
Cr Ni B
Fig. 4e: B-Cr-Ni. Phase equilibria of melt with (Ni), Ni3B and some other phases at 1012°C
90
0.00 60.00 40.00
10
Ni
(Ni)
Data / Grid: at.% Axes: at.%
30
Ni3B 20
L
20
30
Cr Ni B Landolt-Börnstein New Series IV/11C3
40.00 60.00 0.00
10
70
80
90
(Ni)
Ni
MSIT®
B–Cr–Ni
166
B
Data / Grid: at.% Axes: at.%
Fig. 5: B-Cr-Ni. Isothermal section at 1000°C CrB4
CrB2
20
80
Cr3NiB6
Cr2Ni3B6
40
60
Cr2B3 Cr3B4
NiB Ni4-xB3+x(m) Ni4+xB3-x(o)
CrB 60
40
Cr5B3
Ni2B
Cr2B
Ni3B
80
Cr
20
20
(Cr)
40
60
80
B
Data / Grid: at.% Axes: at.%
Fig. 6: B-Cr-Ni. Isothermal section at 800°C CrB4 CrB2
20
Cr3NiB6
80
Cr2Ni3B6
Cr2B3 40 Cr3B4
60
NiB Ni4-xB3+x(m)
CrB
Ni4+xB3-x(o)
60
40
Cr5B3
Ni2B
Cr2B
Ni3B
80
Cr
MSIT®
Ni
(Ni)
(Cr)
20
20
40
60
80
(Ni)
Ni
Landolt-Börnstein New Series IV/11C3
B–Cr–Ni
167
1500
Fig. 7: B-Cr-Ni. Vertical section Ni-CrB2 from 100 to 40 at.% Ni
L L+(Ni)
Temperature, °C
1250
L+CrB
L+CrB+Ni3B
L+CrB+Ni2B 1096 1050
L+(Ni)+Ni3B
L+Ni3B
1000
(Ni)+Ni3B
(Ni)+CrB+Ni3B
CrB+Ni2B+Ni3B
Cr Ni B
Landolt-Börnstein New Series IV/11C3
20.00 40.00 40.00
60
80
Ni
Ni, at.%
MSIT®
168
Bi–In–Sb
Bismuth – Indium – Antimony Tamara Velikanova, Mikhail Turchanin, Hans Leo Lukas Introduction Scientific interest in phase relationships in the Bi-In-Sb system is mainly due to two reasons. The first one is a potential application for long wavelength infrared detector materials. The second one is the development of Pb free solders. The first systematic and most extensive studies on the Bi-In-Sb phase diagram are the three consecutive works of Peretti [1958Per, 1960Per, 1961Per] in which a number of vertical sections were determined using DTA and X-ray diffraction. [1979Ufi] presented the InSb-In2Bi vertical section. The Bi solubility limit in InSb was investigated by [1964Kos, 1972Jou, 1977Lat, 1979Ufi, 1980Zil1, 1981Lan, 2000Moh]. Experimental works of [1979Ufi, 2000Moh, 2000Dix, 2000Iwa] were devoted to studies of growth and structural properties of single crystals. Data on the growth and structural properties of thin films were obtained by [1978Zil, 1980Cad, 1980Zil1, 1980Zil2, 1997Lee1, 1999Wag, 1999Gao, 2001Osz]. The structure of rapidly quenched solid solutions of InSb-InBi alloys has been studied by [1991She]. Thermodynamic properties of liquid ternary alloys were investigated by [1976Vec, 1978Pre, 1982Erm, 1986Gor, 1997Kam]. Thermal effects of phase transitions have been studied by [1988Evg]. These and some more experimental works are summarized in Table 1. A calculation of the phase diagram by the Calphad method has been carried out by [2002Cui]. Binary Systems For the Bi-In, Bi-Sb and In-Sb systems the assessments of [Mas2] were accepted. Thermodynamic data sets of [2002Cui] for the Bi-In system, of [1992Feu] for Bi-Sb and of [1994Ans] for In-Sb were used to calculate diagrams and the reaction scheme presented in this assessment. Solid Phases No ternary compound has been reported. Table 2 summarizes data for the unary and binary phases. Negligible solubilities of the third component are found in all binary phases except InSb [1964Kos, 1958Per, 1960Per, 1961Per, 1977Lat, 1979Ufi, 1980Zil1, 2000Moh]. From the solid solutions AlxIn1–xSb, GaxIn1–xSb and InAsxSb1–x [2005Sch, 2005Luk, 2005Mis] it is known, that the homogenization of InSb solid solutions is very slow and even above the melting temperature of pure InSb needs more than 1000 h to be effective. Thus the measured composition in most cases is the frozen in equilibrium composition during growth conditions. For liquid phase epitaxy or single crystal synthesis by the methods of Bridgman or Czochralski these conditions depend mainly on the composition of the melt. The temperature of growth is a range just below the liquidus temperature of this composition. [1972Jou] varied the InBi content in InSb-InBi melts up to 70 mol% and from grown single crystals determined the solidus curve of InSb1–xBix (Fig. 2) along this section. [2000Moh] by the same method measured a solidus curve of similar shape but about 1 mol% larger solubility of InBi. [2000Dix] reported a higher value of x = 0.065 for the composition of a single crystal grown from a melt with 20 mol% InSb + 80 mol% InBi. Three other papers [1977Lat, 1979Ufi, 1981Lan] reported slightly lower x values than [1972Jou] for solid solutions grown from melts of the section InSb-InBi. These authors measured also samples grown from melts of the sections InSb-Bi and InSB-In2Bi and reported lower Bi contents for the first one, but slightly higher Bi contents for the second one. The generally accepted assumption, that the homogeneity range of InSb1–xBix is practically restricted to Sb-Bi exchange thermodynamically postulates the highest Bi solubility to be in the InSb-InBi section. In the dataset of [2002Cui] the homogeneity range of InSb1-xBix is neglected. Considering the possibility of Bi substitution on the Sb sublattice of this phase and adding a parameter oLIn:Bi = 5300 5T + 0.5(oG(In)In + oGBi) J#mol–1 to the dataset the calculation reproduces well the solidus line x = f(T) of InSb1–xBix given by [1972Jou]. MSIT®
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169
The lattice parameter variation of the InSb1-xBix solid solution up to x = 0.017 is given in Fig. 1 according to the measurements of [1972Jou]. It can be represented as a = 647.9 + 16x pm. These results correlate well with data of [2000Moh] where the lattice parameter was found to increase from 647.8 to 648.1 pm in the range x = 0 to x = 0.02. Metastable single-crystal epitaxial InSb1–xBix solid films with x = 0.12 (6 at.% Bi) were obtained by sputtering from independently controlled InSb and Bi targets [1980Zil1, 1980Zil2, 1980Cad]. The authors believe that these maximum values obtained in their experiments along the InSb-InBi section do not belong to a fundamental limit of the solid solubility. Quasibinary Systems The two vertical sections InSb-InBi and InSb-In2Bi are likely to be quasibinary ones because the three compounds InSb, InBi and In2Bi melt congruently. Indeed [1958Per, 1960Per, 1961Per, 1972Jou, 1979Ufi] reported the sections InSb-InBi and InSb-In2Bi to be quasibinary systems of eutectic type (degenerated to the compositions of InBi and In2Bi), respectively. No temperature maximum or minimum was found along the monovariant line of the liquid phase in the L + three-phase equilibrium [1958Per, 1960Per]. Thus the vertical section InSb-Bi is not a quasibinary system, which is easily understandable as in the binary systems the liquidus temperature increases monotonically from Bi towards Sb, but decreases monotonically from Bi towards the Bi + InBi eutectic. Calculations using the dataset of [2002Cui] confirm the quasibinary degenerated eutectic character of the two sections InSb-InBi and InSb-In2Bi (Figs. 2 and 3). In Fig. 3 the boundary of is not exactly in the plane of section, but at 50 at.% In. The line in the diagram is the projection into the section plane along constant Bi content. The section InSb-Bi is clearly not quasibinary. Invariant Equilibria Invariant equilibria calculated from the dataset of [2002Cui] are given in Table 3. In E1, U1, U2, E2, and E3 the liquid phase participates. In the publication [2002Cui] the authors reported the reactions E1 and E2 as U type reactions, the compositions given, however, show clearly the liquids to be located inside the triangles of the three solid phases and thus the reactions clearly are ternary eutectics. The temperatures of U1 and E3 reported by [2002Cui] do not exactly fit with the recalculated ones. Due to a ternary extension of Raoult’s law the difference between the binary Bi-In three-phase equilibrium temperatures and the ternary four-phase temperatures are strongly related to the Sb mole fractions of liquid in the four-phase equilibria. To show these differences the temperatures in Table 3 and Fig. 4 are given with two digits after the point. With the assumption of zero solubility of Sb in (In), , J phases coexisting in the solid-phase equilibrium D1 this equilibrium is totally degenerated and exactly equals to the binary eutectoid J + (In) with the phase being in equilibrium but not participating in the reaction. In reality it may be either below (E type) or above (U type) the binary three-phase reaction depending, if the solubility of Sb in J is greater or less than the mean of the Sb solubilities in and (In), respectively. The calculated reaction scheme is given in Fig. 4. The calculated equilibrium temperatures are in good agreement with the experimental ones measured by [1960Per, 1961Per]. Liquidus and Solidus Surfaces A liquidus surface projection of the Bi-In-Sb is given in Fig. 5, calculated from the dataset of [2002Cui] showing liquidus isotherms at temperature intervals of 50°C. The liquidus temperatures increase significantly with increasing Sb concentration whereas the solidus temperatures only slightly are suppressed, they contain the invariant equilibria summarized in Table 2. A liquidus surface projection drawn by [1960Per] shows the univariant curve L + + going exactly into the Bi corner, which is thermodynamically very unlikely. As [2002Cui] showed, the experimental points of [1960Per] fit equally well with the calculated vertical sections, meaning that this unlikely feature of the liquidus surface is not supported by the experimental data.
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Bi–In–Sb
The solidus surface of the Bi-In-Sb system in Fig. 6 is calculated using the dataset of [2002Cui] with the additional parameter oLIn:Bi mentioned above. Isothermal Sections Below the solidus temperature the phase (InSb1-xBix) is in equilibrium with all other solid phases. The corresponding isothermal sections can be constructed by connecting all boundaries of the binary phase diagrams by straight lines with the phase. This is demonstrated in Fig. 7 by the calculated isothermal section at 75°C. Isothermal sections above the lowest solidus temperature can be constructed similarly using additionally the corresponding liquidus isotherm shown in Fig. 5. Only the corner of the three-phase equilibrium L + + has to be extra calculated. It is somewhat more Sb rich than the corresponding solidus temperature of the binary Bi-Sb system. This is demonstrated by the calculated isothermal sections at 300 and 400°C in Figs. 8 and 9, respectively. Temperature – Composition Sections Seventeen vertical sections were measured by Peretti [1958Per, 1960Per, 1961Per]. Twelve of them were calculated by [2002Cui] and compared with Peretti’s experimental points. The fit with the experimental points is very good, although the drawings of Peretti contain some unlikely or even incorrect features. Only the calculated temperatures of secondary crystallization of the (Bi,Sb) solid solution ( phase) are significantly higher than the measured ones. This may be due to segregation as the first parts of the primarily crystallizing solid solution contain nearly pure Sb and do not equilibrate during the following crystallization. Thus the experimental points belong to secondary crystallization starting in a melt depleted by Sb at a lower temperature, whereas the calculation gives equilibrium temperatures where the L+ / L++ boundary passes the initial composition of the sample. In Figs. 10 and 11 two calculated vertical sections are shown as examples. Thermodynamics The enthalpy of mixing of Bi-In-Sb melts was determined by quantitative thermal analysis at 632°C for five sections with constant In/Bi ratios of 16/1, 2/1, 1/1, 1/2, and 1/9 by [1976Vec]. The error of the determination was estimated to be about 5%. The obtained results are represented graphically in Fig. 12 by isoenthalpy lines. Bi-In-Sb liquid alloys are characterized by positive and negative deviations from ideality. The maximum enthalpy of mixing 4.99 ± 0.20 kJ#mol–1 corresponds to a ternary composition with xIn = 0.441 and xBi = 0.059. The enthalpy of mixing along the InSb-Bi section was determined by [1978Pre] at 612°C using a high temperature calorimeter. The derived enthalpies of mixing were reported with reference not to the pure liquid elements as in Fig. 12, but referenced to the pure binary liquid InSb and liquid pure Bi, Fig. 13. They show a positive deviation from ideality with a maximum at about 0.56 kJ#mol–1 near xBi = 0.6. The authors compared their results with calculations using a Redlich-Kister description as well as an approach to the associated solution model. The thermodynamic activity of bismuth in dilute InSb-Bi liquid alloys with xBi < 0.014 has been determined by emf measurements (immersion method) at 678°C by [1982Erm]. For bismuth concentrations up to 1 at.% the deviation from ideality is negative. A transition to positive deviations from Raoult’s law occurs at bismuth concentrations above 1 at.%. The partial Gibbs energies of indium in liquid and heterogeneous mixtures were obtained by the emf method [1986Gor] for three sections with constant Bi:Sb ratios of 2:1, 1:1 and 1:2 in the temperature interval 497 to 527°C. From the temperature dependencies of the measured emf values partial enthalpies of mixing of In were derived. The results (Fig. 14) were used for calculation of the integral properties in the whole concentration range and of the liquidus surface. Thermal effects accompanied with the crystallization of liquid alloys belonging to the InSb-InBi and InSb-In2Bi sections were studied by quantitative thermal analysis by [1988Evg]. Their results are shown in Table 4. The concentration dependence of these thermal effects show regions near InSb in both the MSIT®
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171
InSb-InBi and InSb-In2Bi sections in which these thermal effects are larger than the heat of fusion of InSb. This is in accordance with the positive enthalpy of mixing shown in Fig. 13. Thermodynamic activities of indium in liquid alloys was studied between 613 and 906°C by [1997Kam] using the emf method with zirconia solid electrolyte. These investigations were carried out at fifteen alloys along the three sections In-Bi0.8Sb0.2, In-Bi0.5Sb0.5, and In-Bi0.2Sb0.8. The results are shown in Fig. 15. The mixing enthalpy and Gibbs energy of liquid alloys were estimated by [1991Pri] in the framework of their own empirical model. The enthalpies of mixing measured by [1978Pre], the activities of In in liquid from [1997Kam] and the partial enthalpy of mixing of In in liquid from [1986Gor] were used by [2002Cui] for a thermodynamic assessment of the system. The corresponding curves calculated from the assessed thermodynamic dataset are shown in Figs. 13 to 15. As the figures show, this thermodynamic dataset reproduces well the measured thermochemical quantities of liquid in the ternary system. As was outlined in the previous chapters the thermodynamic description proposed by [2002Cui] fits also well with most experimental data of phase relations in the system. This thermodynamic assessment can be improved by considering substitution of Sb by Bi in the phase and modeling the ternary solubility of this phase, which was neglected by [2002Cui]. As already pointed out above in “solid phases” the sublattice model In(Bi,Sb) with a single parameter oL o (In) + oG ) J#mol–1 added, the assessed dataset enables good reproduction In:Bi = 5300 5T + 0.5( G In Bi of the measurements of Bi solubility in by [1972Jou]. Notes on Materials Properties and Applications In the last years increasing efforts have been made to search for suitable Pb free solders. Since Bi, In and Sb are important elements for the development of Pb free solders, the Bi-In-Sb system is included in the present volume considering its potential technical perspective for solder applications. Table 5 summaries the available experimental works devoted to the study of physical properties of InSb1–xBix bulk alloys and thin films. Main emphasis in investigation of materials properties of Bi-In-Sb alloys was given to the InSb1–xBix phase. This material is very interesting as it is a semiconductor InSb modified by InBi with semi-metal behaviour. Hence the band-gap energy can be varied over a wide range. Consequently narrow-gap semiconducting compounds can be obtained, which are of highest interest as promising materials for photoelectronic devices operating in the long wavelength infrared range (7.4 to 16 m) [1997Lee1, 1997Lee2, 1998Lee, 2000Iwa]. The limited equilibrium solubility of Bi in InSb is the main obstacle in the growth of InSb1–xBix alloys with optimal electronic properties. The liquid phase epitaxy method (LPE) provides a maximum Bi alloy content of x = 0.01 to 0.025 with Eg 0.1 eV [1979Ufi, 1981Lan, 1997Akc] depending on the growth temperature, which is somewhat below the liquidus temperature of the selected liquid composition. Supersaturated thin films can be obtained by sputtering [1978Zil, 1980Zil1, 1980Zil2, 1980Cad] with Bi contents up to 6 at.% (x up to 0.12), by low-pressure metalorganic chemical vapor deposition (LP-MOCVD) [1997Lee1, 1997Lee2, 1998Lee, 1999Wag] with x up to 0.058, by quenching of liquid with a rate of about 106K#s–1 [1991She] (x = 0.02 to 0.04), or by flash evaporation [2001Osz] (x = 0.02). An increase of the electron concentration and a decrease of their mobility was observed with increasing Bi content of the solid solution. Figures 16 and 17 show these properties, measured by [1997Lee2] at 77 and 300 K, as functions of the Bi content. The metastable single-phase films show very good thermal stability. After annealing up to 120 h at temperatures below 229°C films with x = 0.12 remained single phase [1980Zil2]. Annealing 12 h at 266°C resulted in precipitation of Bi ( phase). After additional annealing for 12 h at 304°C finally the X-ray pattern of the InBi phase was observed. For films of lower supersaturation the annealing temperatures for decompositions into + or + + increase as shown in Fig. 18. The band gap Eg was calculated ab initio by the empirical pseudopotential method [1998Vyk, 1999Dei] as function of the Bi content (2x in InSb1–xBix) at three different temperatures, 0, 77 and 300 K. It is shown in Fig. 19a, the corresponding temperature coefficient in Fig. 19b. The results agree well with data from optical absorption experiments of [1969Jea]. Ab initio calculations of [1973Jea] gave a smaller decrease of the bandgap vs Bi content of 67% of that shown in Fig. 19a. [2000Dix] reported a bandgap of 113 meV at room temperature and a composition of x = 0.0654 (3.27 at.% Bi), which is higher than the values mentioned Landolt-Börnstein New Series IV/11C3
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Bi–In–Sb
before. Other measured properties of InSb1–xBix solid solutions are Hall coefficient [1969Jea, 1991She, 2000Dix], electric conductivity [1991She, 2000Dix] and microhardness (Fig. 20) [1977Lat, 2000Moh]. References [1958Per] [1960Per] [1961Per]
[1964Kos]
[1969Jea]
[1972Jou]
[1973Jea]
[1976Vec]
[1977Lat]
[1978Pre]
[1978Zil]
[1979Ger]
[1979Ufi]
[1980Cad]
MSIT®
Peretti, E.A., “The System InSb-InBi”, Trans. Metall. Soc. AIME, 212(1), 79 (1958) (Experimental, Phase Diagram, 3) Peretti, E.A., “Constitution Studies of the Antimony and Bismuth-Rich Portions of the Sb-Bi-In System”, Trans. ASM, 52, 1046-1058 (1960) (Experimental, Phase Diagram, 7) Peretti, E.A., “Constitution Studies of the Indium-Rich Portion of the System Antimony-Bismuth-Indium”, Trans. ASM, 54, 12-19 (1961) (Experimental, Phase Diagram, Phase Relations, 6) Kostur, I. L., Psaryev, V. I., “Solubility of Certain Elements in InSb and CdSb and Their Effects on their Effect on the Physical Properties”, Ukrain. Fiz. Zhur., 9(8), 900-907 (1964) (Crys. Structure, Electr. Prop., Experimental, Mechan. Prop., Phase Relations, Thermodyn., 6) Jean-Louis, A.M., Hamon, C., “Properties of the InSb1–xBix Alloys” (in French), Phys. Status Solidi, 34, 329-340 (1969) (Calculation, Electr. Prop., Electronic Structure, Experimental, 11) Joukoff, B., Jean-Louis, A.M., “Growth of InSb1–xBix Single Crystals by Czochralski Method”, J. Cryst. Growth, 12(2), 169-172 (1972) (Crys. Structure, Experimental, Phase Relations, 3) Jean-Louis, A.M., Duraffourg, G., “Properties of the InSb1–xBix Alloys. III. Theoretical Estimation of the Forbidden Band” (in French), Phys. Status Solidi B, 59, 495-503 (1973) (Calculation, Crys. Structure, Electronic Structure, 21) Vecher, A.A., Mechkovskii, L.A., Rubenchik, G.B., Savitskii, A.A., Skoropanov, A.S., “The Enthalpies of Formation of Indium-Bismuth-Antimony Melts”, Russ. J. Phys. Chem. (Engl. Transl.), 50(11), 1759-1760 (1976), translated from Zh. Fiz. Khim., 50, 2956-2958 (1976) (Experimental, Phase Relations, Thermodyn., 11) Latypov, Z.M., Saveliev, V.P., Averyanov, I.S., Faizullina, N.R., Prikhodtsev, M.N., “Investigation of the Region of Solid Solutions Based on InSb in the In-Sb-Bi System”, Izv. Akad. Nauk SSSR, Neorg. Mater., 13, 910-911 (1977) (Experimental, Phase Relations, Mechan. Prop., 0) Predel, B., Gerdes, F., “Thermodynamic of the Liquid Alloys of the InSb-Pd and InSb-Bi Systems”, J. Less-Common Met., 59, 153-164 (1978) (Experimental, Phase Relations, Thermodyn., 10) Zilko, J.L., Greene, J.E., “Growth of Metastable InSb1–xBix Thin Films by Multitarget Sputtering”, Appl. Phys. Lett., 33(3), 254-256 (1978) (Experimental, Interface Phenomena, Phys. Prop., 14) Gerdes, F., Predel, B., “Thermodynamic Properties of the Liquid Alloys of the InSb-Bi, GaSb-Bi and AlSb-Bi Systems” (in German), J. Less-Common Met., 65(1), 41-49 (1979) (Experimental, Thermodyn., 11) Ufimtsev, V.B., Zinov’ev, V.G., Raukhman, M.R., “Heterogeneous Equilibria in the System In-Sb-Bi, and Liquid-Phase Epitaxy of InSb-Based Solid Solutions”, Inorg. Mater. (Engl. Trans.), 15(10), 1371-1374 (1979), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 15(10), 1740-1743 (1979) (Experimental, Phase Diagram, Phase Relations, Phys. Prop., 11) Cadien, K.C., Zilko, J.L., Eltoukhy, A.H., Greene, J.E., “Growth of Single-Crystal Metastable InSb1–xBix and (GaSb)1-xGex Semiconducting Films”, J. Vac. Sci. Technol., 17(1), 441-444 (1980) (Electr. Prop., Experimental, Phase Relations, 14)
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb [1980Zil1]
[1980Zil2]
[1981Lan]
[1982Erm]
[1986Gor]
[1988Evg]
[1991Pri]
[1991She]
[1992Feu]
[1994Ans]
[1997Akc]
[1997Lee1]
[1997Lee2]
[1997Kam]
[1998Lee]
Landolt-Börnstein New Series IV/11C3
173
Zilko, J.L., Greene, J.E., “Growth and Phase Stability of Epitaxial Metastable InSb1–xBix Films on GaAs. I. Crystal Growth”, J. Appl. Phys., 51(3), 1549-1559 (1980) (Crys. Structure, Electr. Prop., Experimental, Optical Prop., 32) Zilko, J.L., Greene, J.E., “Growth and Phase Stability of Epitaxial Metastable InSb1–xBix Films on GaAs. II. Phase Stability”, J. Appl. Phys., 51(3), 1560-1564 (1980) (Electronic Structure, Experimental, 10) Lantsov, A.F., Akuchrin, R.Kh., Zinov`ev, V.G., “Distribution of Bismuth in Epitaxial Layers of InSb / Bi”, Inorg. Mater. (Engl. Trans.), 17, 1146-1148 (1981), translated from Izv. Akad. Nauk SSSR, Neorg. Mater, 17(9) 1550-1553 (1981) (Experimental, Transport Phenomena, 7) Ermakov, A.V., Egorkin, V.V., Ufimtsev, V.B., “Study of the Thermodynamic Properties of Dilute In-Sb-Bi Melts by the Immersion Method”, Russ. J. Phys. Chem. (Engl. Transl.), 56(6), 942-944 (1982) (Experimental, Thermodyn., 12) Goryacheva, V.I., Pashin, S.F. Geiderikh, V.A., “Thermodynamic Properties and Liquidus of the Indium-Bismuth-Antimony System”, Vestn. Mosk. Univ., Ser. 2: Khim., 26(2), 142-150 (1985) (Experimental, Phase Diagram, Thermodyn.) as quoted by [2002Cui] Evgen’ev, S.B., “Thermal Effects of Phase Transitions in the Ga-Sb-Bi and In-Sb-Bi Systems”, Inorg. Mater. (Engl. Trans.), 24, 454-457 (1988), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 24(4), 546-549 (1988) (Experimental, Phase Relations, Thermodyn., 6) Prikhod`ko, E.V., Garmash, L.I., “Effect of the Charged State of the Components on the Thermodynamic Properties of Metallic Melts”, Russ. Metall. (Engl. Transl.), 6, 209-213 (1991), translated from Izv. Akad. Nauk SSSR, Met., (6), 215-219 (1991) (Calculations, Thermodyn., 2) Shepelevich, V.G., “Structure and Electrical Properties of Rapidly Quenched InSb Foils and of Solid Solutions of the System InSb-InBi”, Inorg. Mater. (Engl. Trans.), 27(12), 2150-2152 (1991), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 27(12), 2505-2507 (1991) (Crys. Structure, Electr. Prop., Experimental, 7) Feutelais, Y., Morgant, G., Didry, J.R., Schnitter J., “Thermodynamic Evaluation of the System Bismith-Antimony”, Calphad, 16(2), 111-119 (1992) (Calculation, Phase Diagram, Thermodyn., 47) Ansara, I., Chatillon, C., Lukas, H.L., Nishizawa, T., Ohtani, H., Ishida, K., Hillert, M., Sundman, B., Argent, B.B., Watson, A., Chart, T. G., Anderson, T., “A Binary Database for III–V Compound Semiconductor Systems”, Calphad, 18(2), 172-222 (1994) (Calculation, Phase Diagram, Thermodyn.) Akchurin, R.Kh., Komarov, D.V., “Formation of Multilayer Strained-Layer Heterostructures by Liquid Epitaxy. II. Simulation of the Fabrication of Heterostructures Based on Indium-Arsenic-Antimony-Bismuth Solid Solutions”, Tech. Phys., 42(7), 762-768 (1997) (Crys. Structure, Experimental, Thermodyn., 11) Lee, J.J., Kim, J.D., Razeghi, M., “Long-Wavelength Infrared Photodetectors Based on InSbBi Grown on GaAs Substrates”, Appl. Phys. Lett., 71(16), 2298-2300 (1997) (Electr. Prop., Experimental, Optical Prop., 14) Lee, J.J., Kim, J.D., Razeghi, M., “Growth and Characterization of InSbBi for Long Wavelength Infrared Photodetectors”, Appl. Phys. Lett., 70(24), 3266-3268 (1997) (Electronic Structure, Experimental, Optical Prop., 16) Kameda, K., Yamaguchi, K., Kon T., “Activity of Indium in Molten In-Pb-Ag and In-Bi-Sb Alloys Measured by an EMF Method Using a Zirconia Electrolyte”, Nippon Kinzoku Gakkaishi, 61(5), 444-448 (1997) (Experimental, Thermodyn.) Lee, J.J., Kim, J.D., Razeghi, M., “Room Temperature Operation of 8-12 mym InSbBi Infrared Photodetectors on GaAs Substrates”, Appl. Phys. Lett., 73(5), 602-604 (1998) (Electr. Prop., Experimental, 10)
MSIT®
174 [1998Vyk]
[1999Dei]
[1999Gao]
[1999Wag]
[2000Dix]
[2000Iwa]
[2000Moh]
[2001Osz]
[2002Cui]
[2005Sch]
[2005Luk]
[2005Mis]
MSIT®
Bi–In–Sb Vyklyuk, J.I., Deibuk, V.G., “The Band Structure and Electron Density of InSb1–xBix Solid Solutions”, Acta Phys. Pol. A, 94(3), 611-616 (1998) (Electronic Structure, Experimental, 14) Deibuk, V. G., Viklyuk, Ya. I., Rarenko, I. M., “Calculating the Band Structure of InSb1–xBix Solid Solution”, Semicond., 33(3), 293-296 (1999) (Calculation, Electronic Structure, 14) Gao, Y.Z., Yamaguchi, T., “Liquid Phase Epitaxial Growth and Properties of InSbBi Films Grown from In, Bi and Sn Solutions”, Cryst. Res. Technol., 34(3), 285-292 (1999) (Electr. Prop., Experimental, Morphology, Optical Prop., 8) Wagener, M.C., Kroon, R.E., Botha, J.R., Leitch, A.W.R., “Analysis of Secondary Phases in InSbBi Thin Films”, Physica B, 273-274, 919-922 (1999) (Crys. Structure, Experimental, 9) Dixit, V.K., Rodrigues, B.V., Bhat, H.L., “Growth of InSb(1–x)Bix Crystals by Rotatory Bridgman Method and Their Characterization”, J. Cryst. Growth, 217, 40-46 (2000) (Crys. Structure, Electr. Prop., Experimental, Magn. Prop., 6) Iwanowski, R.J., Heinonen, M.H., Raczynska, J., Fronc, K., “XPS Analysis of Surface Compositional Changes in InSb1–xBix (111) Due to Low-Energy Ar(+) Ion Bombardment”, Appl. Surf. Sci., 153, 193-199 (2000) (Electronic Structure, Interface Phenomena, Experimental, 26) Mohan, P., Babu, S.M., Santhanaraghavan, P., Ramasamy, P., “Growth, Phase Analysis and Mechanical Properties of InSb1-xBix Crystals”, Mater. Chem. Phys., 66, 17-21 (2000) (Experimental, Mechan. Prop., Phase Relations, 14) Oszwaldowski, M., Berus, T., Szade, J., Jozwiak, K., Olejniczak, I., Konarski, P., “Structural Properties of InSbBi and InSbAsBi Thin Films Prepares by the Flash-Evaporation Method”, Cryst. Res. Technol., 36(8-10), 1155-1171 (2001) (Crys. Structure, Electr. Prop., Experimental, Optical Prop., 28) Cui, Y., Ishihara, S., Liu, X.J., Ohnuma, I., Kainuma, R., Ohtani, H., Ishida, K., “Thermodynamic Calculation of Phase Diagram in the Bi-In-Sb Ternary System”, Mater. Trans., JIM, 43(8), 1879-1886 (2002) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 26) Schmid-Fetzer, R., Lukas, H.L., “Aluminium-Indium-Antimony”, in Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology (New Series). Group IV: Physical Chemistry. Ed. W. Martiensen, “Ternary Alloy Systems. Phase Diagrams, Crystallographic and Thermodynamic Data”, Vol. 11C1, Effenberg, G., Ilyenko, S., (Eds.), Springer-Verlag, Berlin, Heidelberg, 144-153, (2005) (Phase Relations, Phase Diagram, Thermodyn., Crys. Structure, Assessment, 24) Lukas, H.L., “Gallium-Indium-Antimony”, in Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology (New Series). Group IV: Physical Chemistry. Ed. W. Martiensen, “Ternary Alloy Systems. Phase Diagrams, Crystallographic and Thermodynamic Data”, Vol. 11C1, Effenberg, G., Ilyenko, S., (Eds.), Springer-Verlag, Berlin, Heidelberg, 412-426, (2005) (Phase Relations, Phase Diagram, Thermodyn., Crys. Structure, Assessment, 68) Misra, S., Anderson, T., Ansara, I., Ivanchenko, V., Lakiza, S., “Arsenic-Indium-Antimony”, in Landolt-Boernstein, Numerical Data and Functional Relationships in Science and Technology (New Series). Group IV: Physical Chemistry. Ed. W. Martiensen, “Ternary Alloy Systems. Phase Diagrams, Crystallographic and Thermodynamic Data”, Vol. 11C1, Effenberg, G., Ilyenko, S., (Eds.), Springer-Verlag, Berlin, Heidelberg, 227-241, (2005) (Phase Relations, Phase Diagram, Thermodyn., Crys. Structure, Assessment, 85)
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb
175
Table 1: Investigations of the Bi-In-Sb Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1958Per]
Thermal analysis, X-ray analysis, optical microscopy
InSb-InBi, 100-600°C
[1960Per]
Thermal analysis, X-ray analysis, optical microscopy
8 vertical sections in Bi- InBi-InSb-Sb part of system, 100-600°C
[1961Per]
Thermal analysis, X-ray analysis, optical microscopy
8 vertical sections in InBi-InSb-In part of system, 50-600°C
[1964Kos]
X-ray analysis
Inx(InSb)yBiz as-cast ingots and single crystals, x = 0.7632-0.9208, y = 0.0765-0.2258, z = 0.0027-0.0110
[1972Jou]
Differential thermal analysis, X-ray analysis, liquid phase epitaxy
single crystal samples InSb1–xBix, x = 0-0.026
[1976Vec]
Quantitative thermographic analysis
Enthalpy of mixing of liquid, at 632°C in the whole composition triangle
[1977Lat]
X-ray analysis, microhardness measurements
InSb-InBi, InSb-In2Bi and InSb-Bi, 90-99.75 mol.% InSb
[1978Pre] [1979Ger]
High temperature calorimetry
InSb-Bi, 612°C
[1978Zil]
Thin film growth by multi target sputtering
metastable supersaturated InSb1–xBix (x0.12) solid solutions
[1979Ufi]
DTA, liquid phase epitaxy
InSb-(In2Bi, InBi or Bi), 300-500°C
[1979Ufi]
Differential thermal analysis, local X-ray-spectral analysis
InSb-InBi, InSb-In2Bi and InSb-Bi, 100-600°C
[1980Cad]
Wavelanght dispersive energy analysis, scanning transmission electron microscopy
polycrystalline metastable InSb1-xBix films, x = 0.07-0.12, 200-500°C
[1980Zil1]
Wavelength dispersive energy analysis, X-ray analysis
Single-crystal metastable InSb1-xBix films, x = 0.02-0.12, 200-500°C
[1980Zil2]
X-ray analysis
metastable InSb1–xBix films, x = 0-0.12 bulk alloys InSb1–xBix, x < 0.031, 90°C
[1981Lan]
Radioisotopic tracer method, local X-ray-spectral analysis
InSb-InBi, InSb-In2Bi InSb-Bi epitaxial layers, xBi < 0.015, 220-450°C
[1982Erm]
Emf immersion method
aBi, InSb-Bi, xBi < 0.014, 678°C
[1986Gor]
Emf salt electrolyte method
sections at Bi/Sb ratios 2/1 and 1/1 at 497°C, 1/2 at 527°C
[1988Evg]
Quantitative differential thermal analysis
crH, InSb-InBi, InSb-In2Bi
[1991She]
X-ray analysis
InSb-InBi rapidly quenched foils, 100-300°C
Landolt-Börnstein New Series IV/11C3
MSIT®
Bi–In–Sb
176 Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1997Kam]
Emf solid electrolyte method
aIn, sections In-Bi0.8Sb0.2, In-Bi0.5Sb0.5, In-Bi0.2Sb0.8, 613-906°C
[1997Lee1] [1997Lee2]
InSb1–xBix films, x = 0-0.058, 70-300 K Energy dispersive X-ray analysis, X-ray analysis of thin films grown by low pressure metalorganic chemical vapor deposition
[1999Gao]
Optical microscopy, electron-probe micro analysis of films grown by liquid phase epitaxy
[1999Wag]
InSb1–xBix films, 455°C Scanning electron microscopy, cross-sectional transmission electron microscopy, X-ray analysis of films grown by metalorganic vapor deposition
[2000Dix]
Energy-dispersive X-ray analysis, X-ray analysis, differential scanning calorimetry, infrared spectroscopy. Bridgman single crystal synthesis
InSb1–xBix single crystals, x = 0.0654
[2000Iwa]
X-ray photoelectron spectroscopy
single crystalline InSb1–xBix epitaxial layer, x = 0.005, 190-310°C
[2000Moh]
Differential thermal analysis, optical InSb1–xBix, x = 0 - 0.02, 50-700°C microscopy, X-ray analysis, scanning electron microscopy
[2001Osz]
X-ray analysis, scanning electron microscopy, X-ray microanalysis, secondary ion mass spectrometry. Metastable thin films prepared by flash evaporation
InSb1–xBix epilayers, (0.01 < x < 0.14)
InSb1–xBix films, x 0.2, 285-715°C
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, Bi1–xSbx
hR6 R3m As
(Bi) < 271.442 (Sb) < 630.755 (In), In1–xBix < 156.634
MSIT®
tI2 I4/mmm In
Lattice Parameters Comments/References [pm] 0 < x < 1, [Mas2] a = 454.613 c = 1186.152
pure Bi, 25°C [V-C2]
a = 430.84 c = 1127.4
pure Sb, 25°C [V-C2]
a = 325.3 c = 497.70
0 < x < 0.08 [Mas2] pure In, 25°C [Mas2]
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, InBi < 110.0
tP4 P4/nmm PbO
a = 501.5 c = 478.8
3, In5Bi3 < 88.9
tI32 I4/mcm Cr3B5
a = 854.4 c = 1268
, In2Bi < 89.5
hP6 P63/mmc Ni2In
a = 549.6 c = 657.9
J, In1–xBix 93.5 - 49
tI2 I4/mmm In
a = 347.2 c = 449.5
, InSb1–xBix < 525.7
cF8 F43m ZnS
177
Lattice Parameters Comments/References [pm] [Mas2] [V-C2] 62.2 to 62.66 at.% In [Mas2] [V-C2] 66.8 to 67.7 at.% In [Mas2] [V-C2] 0.08 x 0.12 [Mas2] 90 at.% In at 70°C [V-C2] 0 < x < 0.026 [1972Jou] at x = 0 [1972Jou]
a = 647.9
Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) In
Bi
Sb
L+
109.76
e3 (max)
L
49.992
49.992
0.016
L ++
109.19
E1
L
47.477
52.509
0.014
L + (In) + J
93.37
U1
L
82.518
17.445
0.037
L + + 3
88.84
U2
L
64.744
35.244
0.012
L+
88.16
e5 (max)
L
66.658
33.329
0.012
L+ +3
88.09
E2
L
65.888
34.100
0.012
L+J+
72.54
E3
L
77.778
22.210
0.012
J + + (In)
49.07
D1
-
-
-
-
Landolt-Börnstein New Series IV/11C3
MSIT®
Bi–In–Sb
178
Table 4: Thermodynamic Data of Reaction or Transformation Reaction or Transformation
T [°C]
Quantity, per mol of compound [kJ#mol–1]
Comments
L InSb
525.7
crH = 47.0 6.1
quantitative differential thermal analysis, [1988Evg]
L xInSb + (1–x)InBi
TL-TS
crH = 56.4 7.3
x = 0.99
TL-TS
crH = 37.8 4.9
x = 0.95
TL-TS
crH = 38.8 5.0
x = 0.80
TL-TS
crH = 26.2 3.4
x = 0.60
TL-TS
crH = 22.5 2.9
x = 0.38
TL-TS
crH = 31.2 4.0
x = 0.20
TL-TS
crH = 31.4 4.1
x = 0.05
Te = 110
crH = 10.0 1.3
Eutectic
TL-TS
crH = 37.8 4.9
x = 0.95
TL-TS
crH = 38.8 5.0
x = 0.80
TL-TS
crH = 26.2 3.4
x = 0.60
TL-TS
crH = 22.5 2.9
x = 0.38
TL-TS
crH = 31.2 4.0
x = 0.20
TL-TS
crH = 31.4 4.1
x = 0.05
Te = 90
crH = 10.0 1.3
Eutectic
L xInSb + (1–x)In2Bi
Table 5: Investigations of the InSb1–xBix Materials Properties Reference
Method/Experimental Technique
Type of Property
[1964Kos]
Compensating method In magnetic field 6800 e
Thermo emf Hall effect
In magnetic field
Electrical conductivity Hall effect
[1980Zil2]
Infrared spectrometry Van der Pauw technique
Absorption coefficient Hall effect
[1997Lee1, 1997Lee2]
Van der Pauw technique Infrared spectrometry X-ray spectrometry Scanning electron microscopy Nomarski differential interference contrast microscopy
Hall effect Photoconductive spectral response
[1969Jea]
MSIT®
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb
179
Reference
Method/Experimental Technique
Type of Property
[1998Lee]
Infrared spectrometry Van der Pauw technique X-ray spectrometry Energy dispersive X-ray analysis
Spectral voltage response Hall effect
[2000Dix]
Infrared spectrometry Van der Pauw technique
Absorption coefficient Hall effect
[2000Moh]
Vickers microhardness experiments
Microhardness
648.2
Fig. 1: Bi-In-Sb. Composition dependence of the lattice parameter of InSb1–xBix solid solutions
648.1
a, pm
648.0
647.9
647.8
InSb
5
10
15
x ⋅ 10
Landolt-Börnstein New Series IV/11C3
MSIT®
Bi–In–Sb
180
Fig. 2: Bi-In-Sb. Calculated quasibinary system InBi-InSb
500
L
Temperature, °C
400
L+β
300
200
β
100
γ+β
In 50.00 Bi 50.00 Sb 0.00
10
20
30
In 50.00 0.00 Bi Sb 50.00
40
Sb, at.%
600
Fig. 3: Bi-In-Sb. Calculated quasibinary system In2Bi-InSb
L
Temperature, °C
500
400
L+β
300
200
β 100
In2Bi+β 0
In 66.70 Bi 33.30 Sb 0.00
MSIT®
10
20
30
Sb, at.%
40
In 50.00 0.00 Bi Sb 50.00
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb
In-Sb
181
Bi-In-Sb
A-B-C
Bi-In
492.5 e1 lα+β 153.9 e2 l (In) + β
109.76 e3 (max) Lβ+γ Lα+β+γ
109.19
109.22 e4 lα+γ
E1
α+β+γ 93.37
93.58 p1 l + (In) ε
L + (In) β + ε
U1 88.86 p2 l+γχ
(In)+β+ε
L+γβ+χ
88.84
U2
γ+β+χ 88.16 e5 (max) Lβ+γ
l+β+χ 88.11 e6 lχ+δ
Lχ+β+δ
88.09
E2
χ+β+δ 72.58 e7 lε+δ 72.54
Lβ+ε+δ
E3
β+ε+ δ
49.07
ε (In) + δ, β
D1
49.07 e8 ε (In) + δ
(In)+β+δ Fig. 4: Bi-In-Sb. Reaction scheme
Landolt-Börnstein New Series IV/11C3
MSIT®
Bi–In–Sb
182
Sb
Data / Grid: at.% Axes: at.%
Fig. 5: Bi-In-Sb. Calculated liquidus surface projection 600 20
80
550
e1
α
40
60
500
60
450
40
500
β
400
450
80
20
350 400
300 350 300
e2 20
In
U1
40
e5 E U 2 2
E3
250°C 60
e3
80
Bi
250
E1
Sb
Data / Grid: at.% Axes: at.%
Fig. 6: Bi-In-Sb. Calculated solidus surface
450°C 20
400
80
360 340 40
60
320
β +α
β
α
310 300
60
40
290
ε+β +δ
(In)+β +ε 80 93.4°C
280
72.5°C 270
20
β +γ +χ α+β +γ 109.2°C
88.8°C
δ+β +χ 88.1°C
In
MSIT®
(In)
ε
20
δ
χ
40
γ
60
80
Bi
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb
183
Sb
Data / Grid: at.% Axes: at.%
Fig. 7: Bi-In-Sb. Calculated isothermal section at 75°C 20
80
40
60
α +β
α
β 60
40
80
20
β +γ +χ
ε+β (In)+ε+β
L+β L+δ+β
L+ε+β
In
ε
(In)
20
δ χ
L
α+β +γ
β +δ+χ
40
γ
60
80
Sb
Bi
Data / Grid: at.% Axes: at.%
Fig. 8: Bi-In-Sb. Calculated isothermal section at 300°C 20
80
α 40
60
α+β
β 60
40
L+β+α L+α 80
20
L+β L
In
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Bi
MSIT®
Bi–In–Sb
184
Sb
Data / Grid: at.% Axes: at.%
Fig. 9: Bi-In-Sb. Calculated isothermal section at 400°C
α 20
80
β+α
40
60
β
L+β+α
L+α
60
40
L+β
80
20
L
20
In
Fig. 10: Bi-In-Sb. Calculated vertical section at 10 at.% Sb
40
60
80
Bi
400
L
Temperature, °C
300
L+α L+β
200
β+α
L+β+α
α
L+β +InBi 100
InBi+β+α
β+ InBi+In5Bi3 In 90.00 0.00 Bi Sb 10.00
MSIT®
20
(In)+β +In2Bi
40
Bi, at.%
60
80
0.00 In Bi 90.00 Sb 10.00
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb
185
700
Fig. 11: Bi-In-Sb. Calculated vertical section InBi-Sb
600
L L+α
Temperature, °C
500
400
L+α +β 300
L+β
α +β
200
109.2 100
α +β +γ 0
In 50.00 Bi 50.00 Sb 0.00
20
40
60
80
Sb
Sb, at.%
Sb
Data / Grid: at.% Axes: at.%
Fig. 12: Bi-In-Sb. Mixing enthalpies of Bi-In-Sb liquid alloys at 632°C (kJ#mol–1) 20
80
40
60
-4.31+/-0.2
60
40
-3.77
80
-2.93
20
0 -0.42
-2.09 -1.26
In
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Bi
MSIT®
Bi–In–Sb
186
600
500
Enthalpy of mixing, H/J. mol–1
Fig. 13: Bi-In-Sb. Enthalpy of mixing along the InSb-Bi section at 612°C: Triangles experimental results of [1978Pre], solid line - calculated by [2002Cui]
400
300
200
100
0
0 InSb
20
40
60
80
100 Bi
80
100
Bi, at.%
0.0
-2.0
Partial mixing enthalpy of In, kJ. mol–1
Fig. 14: Bi-In-Sb. Partial enthalpy of mixing of In in liquid alloys along sections with constant Bi/Sb ratios: symbols experimental results [1986Gor] (triangles Bi/Sb = 1/2, 497°C; circles - Bi/Sb = 1/1, 497°C; crosses Bi/Sb = 1/2, 527°C; lines: calculated by [2002Cui] at same temperatures (solid Bi/Sb = 2/1, dashed – Bi/Sb = 1/1, dots – Bi/Sb = 1/2)
-4.0
-6.0
-8.0
-10.0
-12
20
40
60
In, at.%
MSIT®
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb
1.0 0.9 0.8
Activity of In in liquid
Fig. 15: Bi-In-Sb. Activity of In in liquid alloys at 727°C along sections with constant Bi/Sb ratios: symbols experimental results [1997Kam] (crosses In-Bi0.8Sb0.2, squares - InBi0.5Sb0.5, circles - In-Bi0.2Sb0.8; lines - calculated by [2002Cui] (dots In-Bi0.8Sb0.2, solid In-Bi0.5Sb0.5, dashed - In-Bi0.2Sb0.8)
187
0.7 0.6 0.5 0.4 0.3 0.2 0.1 0
20
40
60
80
100
Fig. 16: Bi-In-Sb. Electron concentration of InSb1–xBix vs Bi concentration at 77 and 300 K
Electron concentration, cm–3
In, at.%
1017
300 K
77 K 1016
0
1.0
2.0
3.0
4.0
5.0
6.0
Bi, at.%
Landolt-Börnstein New Series IV/11C3
MSIT®
Bi–In–Sb
188
Fig. 17: Bi-In-Sb. Electron mobility of InSb1–xBix vs Bi concentration at 77 and 300 K
Mobility, cm2/Vs
40000
30000
300 K
20000
10000
77 K
0 0
1.0
2.0
3.0
4.0
5.0
6.0
Bi, at.%
Temperature, °C
Fig. 18: Bi-In-Sb. Metastable phase map of (110) oriented InSb1–xBix films as a function of composition and annealing temperature
600
L 500
400
β +α +γ β +α
300
β
β
200
L+β
100
0
In 50.00 0.00 Bi Sb 50.00
MSIT®
2
4
Bi, at.%
6
8.00 In 8.00 Bi Sb 84.00
Landolt-Börnstein New Series IV/11C3
Bi–In–Sb
189
260
Fig. 19a:Bi-In-Sb. Composition dependence of the band gap EG at 0 (a), 77 (b) and 300 K (c) in InSb1–xBix solid solutions
240
0K 220
77 K
200
Eg, meV
180 160
300 K 140 120 100 80
InSb 0
0.5
1.0
1.5
2.0
2.5
3.0
x ⋅ 10 0.30
γ, meV/K
Fig. 19b:Bi-In-Sb. Composition dependence of the band gap temperature coefficient ( = – EG/ t) of InSb1–xBix solid solutions
0.28
0.26
0.24
0.22
InSb 0
0.5
1.0
1.5
2.0
x ⋅ 10
Landolt-Börnstein New Series IV/11C3
MSIT®
Bi–In–Sb
190
b a 210
ñ
Hv, kg.mm–2
Fig. 20: Bi-In-Sb. Composition dependence of microhardness of InSb1–xBix solid solutions along the InSb-Bi (a), InSb-InBi (b), and InSb-In2Bi (c) cross sections
190
170
99
96
93
InSb, mol%
MSIT®
Landolt-Börnstein New Series IV/11C3
Bi–In–Sn
191
Bismuth – Indium – Tin Olga Fabrichnaya, Viktor Witusiewicz, Joachim Gröbner Introduction The ternary Bi-In-Sn system was investigated for potential lead-free low-melting solder. First experimental work on this system using X-ray and microscopy methods was published by [1964Doo] follwed by experiments performed by [1967Sch]. Applying thermal, metallographic and X-ray analysis they found the temperature and liquid composition of an eutectic reaction involving (Sn), (Bi) and InBi. [1972Ste1] and [1972Ste2] studied 7 polythermal sections inside the In-InBi-Sn region using differential thermal analysis and X-ray diffraction. [1984Kab] found a eutectic at 59°C involving the phases In2Bi, , and . They solidified the alloys directionally and characterized them by metallographic methods. In [1986Kab] the section In2Bi-Sn was investigated and a projection of the liquidus surface is given. Specimens of eutectic Bi-InBi- at 77°C were unidirectionally solidified at very low speed and quenched to form a representative solid/liquid interface in the study of [1995Rug]. Precipitates observed in the InBi phase showed some solubility of Bi and Sn at the ternary eutectic temperature. In their review of low-temperature solders [1994Mei] showed a liquidus surface based on an unpublished work of Choongun Kim and J.W. Morris (Universtiy of California at Berkeley). The same diagram is given by [1997Hua]. [1999Yoo] investigated the ternary system by X-ray diffraction, energy-dispersive X-ray spectroscopy, optical and scanning electron microscopy. They compared 14 vertical sections with their own results and with those of [1964Doo], [1972Ste2], [1986Kab] and [1967Sch]. In the recent investigation of [2005Wit] measurements of enthalpy of melting were performed on 28 alloys by differential scanning calorimetry (DSC). The tie triangle (Bi)-InBi-(Sn) at 77°C, tie triangle In2Bi-- and tie line In2Bi- at 59°C and composition of liquid at eutectic at 59°C were measured using unidirectionally solidified samples which were derived from extremely slow pulling velocity and measured by energy and wavelength dispersive X-ray analysis (EDX and WDX). It was demonstrated that and phase could dissolve up to 11.4 and 21.2 at.% Sn, respectively A complete thermodynamic calculation of the ternary system is given by [1999Yoo, 2000Ohn, 2002Moe, 2003Moe, 2005Wit]. The mentioned assessments produce different liquidus surface. A summary of the experimental work can be found in [1999Yoo]. Binary Systems All three edge binary phase diagrams are in this evaluation accepted from [2002Moe, 2003Moe]. These systems are virtually equivalent to those of [Mas2]. Solid Phases The crystallographic data of the phases in the Bi-In-Sn system and their ranges of stability are listed in Table 1. Quasibinary Systems [1964Doo] constructed a simple eutectic section between InBi and Sn and assumed a quasibinary which divided the ternary phase diagram in two sections. [1997Hua] reported the eutectic temperature of this quasibinary section as 82°C. [1986Kab] concidered also the In2Bi - Sn section to be a quasibinary phase diagram. [1999Yoo] showed that both sections are not quasibinary. Invariant Equilibria The temperatures of the invariant reactions are listed in Table 2 along with phase composition. Those reactions which involve the liquid phase were calculated using the data sets developed by [2005Wit]. The
Landolt-Börnstein New Series IV/11C3
MSIT®
192
Bi–In–Sn
reaction scheme, following from this calculation is given in Fig. 1. In general, the calculated temperatures and compositions of invariant reactions approach the experimental values, even though the type of the reactions and the involved phases are different in some cases. It should be mentioned that the experimental works of [1967Sch, 1972Ste1, 1986Kab, 1999Yoo] are also mutually inconsistent regarding the phases and their nature. In the study of [2005Wit] the type and phases involved in E2 and U3 were experimentally verified. Liquidus, Solidus and Solvus Surfaces [1967Sch], [1986Kab] and [1994Mei] showed an experimentally derived liquidus surface. The first calculated liquidus surface is given by [1999Yoo]. However, in this calculation the phase (In9Bi) was missing. The liquidus projection presented in Fig. 2 is taken from yet unpublished calculations made by [2005Wit]. Isothermal Sections The isothermal sections at 76.45, 59.5 and 25°C were studied experimentally and calculated by [2005Wit], they are presented in Figs. 3, 4 and 5, respectively. Temperature – Composition Sections [1999Yoo] calculated 14 vertical sections and presented them along with the experimental metallurgical data measured by [1964Doo, 1972Ste2, 1986Kab, 1999Yoo]. [2005Wit] calculated 14 vertical sections at the same conditions as [1999Yoo] and 3 sections which were also studied experimentally. Figures 6 to 10 present 5 of those calculated vertical sections. The selected diagrams represent most of the experimentally studied sections, but the (Sn)/(Sn) transformation is not taken into account. Thermodynamics The first thermodynamic calculation of the whole ternary system, describing the Gibbs energies of all phases by the Calphad method was reported by [1999Yoo]. However, no experimentally measured thermodynamic data were available at that time. The only experimental thermodynamic data for the Bi-In-Sn system became available later in [2005Wit]. They are enthalpies of melting for 28 alloys measured by DSC technique. These enthalpies of the E1 and E2 invariant reactions are given in Table 3 as calculated from the thermodynamic database of [2005Wit], in a very good agreement with the DSC measurements. [2005Wit] gives a new data set for the ternary system considering the phase (In9Bi), which is in better agreement with the experimental data than the earlier calculation. The underlying data set was derived using all available data on phase relations and the experimentally determined enthalpies of melting. This is why many of the data and diagrams from that work have been included here. Notes on Materials Properties and Applications The lead - tin solder alloys are still commonly used. However there is growing need to develop Pb free solder alloys less damaging for the environment [1999Yoo, 2000Ohn, 2003Moe]. Several Pb free alloys with relatively low melting point have been considered [1999Ohn1, 1999Ohn2, 2000Ohn, 2003Moe]. A candidate lead free solder alloy must fulfill many requirements: wettability of the substrate, ability to form a strong chemical bond with the substrate, suitable melting and solidification behavior, good mechanical properties, corrosion resistance, electrical conductivity and it must not be harmful to health and environment [2003Moe]. The Bi-In-Sn alloys are among the most promising Pb free solders. [2002Wan] predicts the viscosity of the ternary Bi-In-Sn melt based on Seetharaman’s viscosity model and Chou’s geometric-thermodynamic model. The viscosity at 500 and 600°C is given for the ratio xBi/xIn = 1. Solute diffusion coefficients in the Bi-In-Sn system were measured in microgravity by [2004Gar], using the shear cell technique.
MSIT®
Landolt-Börnstein New Series IV/11C3
Bi–In–Sn
193
The microstructure and properties of a Bi-In-Sn eutectic alloy was studied in [2002Sen]. A ternary eutectic alloy obtained by continuous casting has more uniform microstructure contrary to the heavily segregated structure of the statically cast specimens. These differences in structure significantly affected the mechanical properties. Contrary to the statically cast specimens wires from continuous casting do exhibit a considerable ductility. References [1964Doo] [1967Sch]
[1972Ste1] [1972Ste2] [1984Kab]
[1986Kab] [1994Mei]
[1995Rug]
[1997Hua] [1999Ohn1]
[1999Ohn2]
[1999Yoo]
[2000Ohn]
[2002Moe]
[2002Sen]
[2002Wan]
Landolt-Börnstein New Series IV/11C3
Dooley, G.J., Peretti, E.A., “Phase Diagram of the System InBi-Sn”, J. Chem. Eng. Data, 9, 90-91 (1964) (Experimental, #, 0) Scherpereel, L.R., Peretti, E.A., “The Ternary Subsystem Indium Bismuthide-Bismuth-Tin”, J. Mater. Sci., 2, 256-259 (1967) (Experimental, Phase Diagram, #, 2) Stelmakh, S.I., Tsimmergakl, V.A., Sheka, I.A., “The In-Sn-InBi System”, Ukr. Khim. Zh., 38, 855-859 (1972) (Experimental, #, 13) Stelmakh, S.I., Tsimmergakl, V.A., Sheka, I.A., “The Compound In2Bi”, Progr. Sov. Chem., 38, 631-633 (1972) (Experimental, #, 9) Kabassis, H., Rutter, J.W., Winegard, W.C., “Microstructure of One of the Ternary Eutectic Alloys in the Bi-In-Sn System”, Metall. Trans. A, 15A, 1515-1517 (1984) (Experimental, Morphology, 12) Kabassis, H., Rutter, J.W., Winegard, W.C., “Phase Relationships in Bi-In-Sn Alloy System”, Mater. Sci. Technol., 2, 985-988 (1986) (Experimental, Phase Diagram, 8) Mei, Z, Vander Plas, H., Gleason, J., Baker, J., “Low-Temperature Solders“, Proc. Electronic Mater. Proces. Symp. Nov. 13-18, 485-495 (1994) (Review, Phase Diagram, #, 33) Ruggiero, M.A., Rutter, J.W., “Origing of Microstructure in 350 K Eutectic of Bi-In-Sn Ternay System”, Mater. Sci. Technol., 11(2), 136-142 (1995) (Experimental, Phase Relations, 22) Hua, F., Glazer, J., “Lead-Free Solders For Electronic Assembly”, Des. Reliab. Solders Solder Interconnect., Proc. Symp., 65-73 (1997) (Review, #, 63) Ohnuma, I., Liu, X.J., Ohtani, H., Ishida, K., “Thermodynamic Database for Phase Diagrams in Micro-Soldering Alloys”, J. Electron. Mater., 28(11), 1164-1171 (1999) (Calculation, Phase Relations, 22) Ohnuma, I., Liu, X.J., Ohtani, H., Ishida, K., “Phase Diagrams of Pb-Free Solder Alloys” (in Japanese), Materia Japan, 38, 923-926 (1999) (Calculation, Phase Diagram, Phase Relations, 13) Yoon, S.W., Rho, B.-S., Lee, H.M., Kim, C.-U., Lee, B.-J., “Investigation of the Phase Equilibria in the Sn-Bi-In Alloy System”, Metall. Mater. Trans. A, 30A, 1503-1515 (1999) (Experimental, Calculation, Phase Relations, #, *, 17) Ohnuma, I., Cui, Y., Liu, X.J., Inohana, Y., Ishihara, S., Ohtani, H., Kainuma, R., Ishida, K., “Phase Equilibria of Sn-In Based Micro-Soldering Alloys”, J. Electron. Mater., 29(10), 1113-1121 (2000) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 27) Moelans, N., “Calculation of Phase Diagram for Lead-free Solder Alloys Based on Bi-In-Sn-Zn System” (in Dutch), Ph.D. Thesis, Catholic University of Leuven, Leuven, Belgium (2002) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 44) Sengupta, S., Soda, H., McLean, A., “Microstructure and Properties of a Bismuth-Indium-Tin Eutectic Alloy”, J. Mater. Sci., 37, 1747-1758 (2002) (Experimental, Mechan. Prop., Morphology, 12) Wang, X., Bao, H., Li, W., “Estimation of Viscosity of Ternary-Metallic Melts”, Metall. Mater. Trans. A, 33(10), 3201-3204 (2002) (Calculation, Phys. Prop., 18)
MSIT®
Bi–In–Sn
194 [2003Moe]
[2004Gar]
[2005Wit]
Moelans, N, Hari Kumar, K.C, Wollants, P., “Thermodynamic Optimization of the Lead-Free Solder System Bi-In-Sn-Zn”, J. Alloys Compd., 360, 98-106 (2003) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 44) Garandet, J.P., Mathiak, G., Botton, V., Lehmann, P., Griesche, A., “Reference Microgravity Measurements of Liquid Phase Solute Diffusivities in Tin- and Aluminum-Based Alloys”, Int. J. Thermophys., 25(1), 249-272 (2004) (Experimental, Transport Phenomena, 43) Witusiewicz, V.T., Hecht, U., Boettger, B., Rex, S., “Thermodynamic Re-Optimisation of the Bi-In-Sn System on the Basis of New Experimental Data”, to be submitted (Experimental, Calculation, Phase Diagram, Phase Relations, Thermodyn., 23)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Bi) < 271.442
hR6 R3m As
a = 454.613 c = 1186.152
(In) < 156.63
tI2 I4/mmm In
a = 325.3 c = 494.70
(Sn) 293.9681 - 13
tI4 I41/amd Sn
a = 583.16 c = 318.15
(Sn) (l) < 13
cF8 Fd3m C (diamond)
a = 648.92
at 25°C [Mas2]
InBi < 110.0
tP4 P4/nmm PbO
a = 501.5 c = 477.1
[V-C2, Mas2]
In5Bi3 < 88.9
tI32 I4/mcm Cr5B3
a = 846 to 854 c = 1256 to 1268
62.2 to 62.66 at.% In [V-C2, Mas2]
In2Bi < 89.5
hP6 P63/mmc Ni2In
a = 549.6 c = 657.9
66.8 to 67.7 at.% In [V-C2, Mas2]
, In9–xSnxBi In9Bi 93.5 - 49
tI2 I4/mmm In
a = 347.2 c = 449.5
88 to 92 at.% In [Mas2] at 90 at.% In and 70°C [V-C2] 12 < x < 44 at.% Sn at x = 25 at.% Sn [Mas2, V-C2]
a = 345.9 c = 439.7
In1–xSnx ~140- ~–133 , In1–xSnx < 224
MSIT®
hP1 P6/mmm InBi
a = 320.8 c = 299.7
at 31°C [V-C2] ~ 0 at.% In 0 to 8 at.% Bi [Mas2] pure In 0 to 13.1 at.% Bi, [Mas2] pure Sn
72 < x < 98 at.% Sn at x = 80 at.% Sn [Mas2, V-C2]
Landolt-Börnstein New Series IV/11C3
Bi–In–Sn
195
Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Bi
In
Sn
L (Sn) + InBi
76.54
e5(max)
L (Sn) InBi
38.0 25.4 49.3
37.2 23.4 50.0
24.8 51.2 0.7
L (Bi) + (Sn) + InBi
76.40
E1
L (Bi) (Sn) InBi
39.3 99.9 26.1 49.4
36.0 0.00 22.9 50.0
24.7 0.01 51.0 0.6
L + (Sn) + InBi
76.25
U1
L (Sn) InBi
35.7 24.1 9.1 49.3
37.6 24.3 23.2 50.0
25.0 51.6 67.7 0.7
L + InBi + In5Bi3
62.26
U2
L InBi In5Bi3
24.4 48.9 21.7 36.6
57.8 50.0 43.2 62.5
17.8 1.1 35.1 0.9
L + In5Bi3 + In2Bi
61.50
U3
L In5Bi3 In2Bi
23.6 36.6 20.9 32.2
58.4 62.5 43.0 66.7
18.0 0.9 36.1 1.1
L + + In2Bi
59.28
E2
L In2Bi
20.7 16.3 8.8 32.0
60.2 40.4 66.8 66.7
19.1 43.3 24.4 13.1
Table 3: Thermodynamic Data of Reaction or Transformation Reaction or Transformation
T [°C]
Quantity, per mol of atoms [kJ, mol, K]
Comments
0.477(Sn)+0.5016In0.5Bi0.5+0.021(Bi) = In0.3599Sn0.2467Bi0.3933 (Liq)
76.40
H = 5149
[2005Wit] calculation + DSC
0.4313 In0.667Bi0.333 + 0.321 + 0.2473 = 59.28 In0.6021Sn0.1912Bi0.2067 (Liq)
H = 3478
[2005Wit] calculation + DSC
Landolt-Börnstein New Series IV/11C3
MSIT®
Bi-In
Bi-In-Sn
In-Sn
196
MSIT®
Bi-Sn
223.5 p1 L + (βSn) γ 139.5 p2 L + (In) β
138.8 e1 L (βSn) + (Bi)
118 e2 Lβ+γ
109.2 e3 L (Bi) + InBi 91.34 p3 L + (In) β 88.7 p4 L + InBi In5Bi3 87.8 e4 L In5Bi3 + In2Bi 76.4
L (βSn) + (Bi) + InBi
Bi–In–Sn
76.54 e5(max) L (βSn) + InBi E1
(βSn)+(Bi)+InBi 70.5 e6 L In2Bi + β
76.25
L + (βSn) γ + InBi
U1
(βSn)+γ+InBi 62.26
L + InBi γ + In5Bi3
U2
InBi+γ+In5Bi3 61.5
L + In5Bi3 γ + In2Bi
U3
In2Bi+γ+In5Bi3 59.28 Landolt-Börnstein New Series IV/11C3
L γ + β + In2Bi γ+β+In2Bi
Fig. 1: Bi-In-Sn. Reaction scheme
E2
Bi–In–Sn
197
Bi
Data / Grid: at.% Axes: at.%
Fig. 2: Bi-In-Sn. Liquidus surface projection
260°C 240
20
80
220 200
(Bi)
180
40
60
160
e3
140 60
p4
e4 e6 p3
80
In5Bi3
80
In2Bi
β
40
U1 (Sn)
U3 E2
γ 100
20
120 140
100
160
120
220 200
180 p2 20
In
e1
e5(max)
U2 80
(In) 140
120 E1
InBi
40
60
e2
Bi Fig. 3: Bi-In-Sn. Isothermal section at 76.45°C
p1
80
Sn
Data / Grid: at.% Axes: at.%
(Bi)
20
80
40
60
L+(Bi)+InBi InBi In5Bi3 60 In2Bi
L+(Sn)+InBi L+(Bi)+(Sn) L
L+(Sn)+InBi
L
80
20
L+(Sn)+γ
L+β (In)
In
Landolt-Börnstein New Series IV/11C3
40
(Bi)+(Sn)
β 20
L+β +γ 40
(Sn)+γ
γ 60
(Sn) 80
Sn
MSIT®
Bi–In–Sn
198
Bi Fig. 4: Bi-In-Sn. Isothermal section at 59.5°C
Data / Grid: at.% Axes: at.%
(Bi)
20
80
40
60
InBi
80
(Sn)+(Bi)+γ
(Bi)+InBi+γ
In5Bi3 60 γ+In In2Bi Bi+ In B 5 i 3
40
L
20
L+β +γ
β +In2Bi+L
β +γ
β
(In)
20
In
γ
40
60
80
Bi Fig. 5: Bi-In-Sn. Isothermal section at 25°C
(Sn)
Sn
Data / Grid: at.% Axes: at.%
(Bi)
20
80
40
60
InBi (Bi)+InBi+γ
In5Bi3 60 InBi+In5Bi3+γ In2Bi
40
(Sn)+(Bi)+γ
80
20
(In)+In2Bi+β
In2Bi+In5Bi3+γ In2Bi+β +γ
(In)
In
MSIT®
γ
β 20
40
60
80
(Sn)
Sn
Landolt-Börnstein New Series IV/11C3
Bi–In–Sn
Fig. 6: Bi-In-Sn. Vertical section at mass ratio Bi/In=95/5, plotted in at.%
199
L 200
Temperature, °C
L+(Bi) L+(Sn) (Sn) L+InBi+(Bi)
L+(Bi)+(Sn)
100
(Bi)+(Sn)
(Bi)+γ InBi+(Bi)+(Sn) InBi+(Bi)+γ 0
8.74 In Sn 0.00 Bi 91.26
Fig. 7: Bi-In-Sn. Vertical section at 50 mass% Sn, plotted in at.%
(Bi)+(Sn)+γ 20
40
60
80
Sn
Sn, at.%
175
L 150
L+(Sn)
Temperature, °C
125
L+(Bi)+(Sn) L+γ L+(Sn)+γ
100
L+γ+β 75
γ
γ+β 50
0
In 50.83 Sn 49.17 0.00 Bi
InBi+γ
(Sn)+γ
(Bi)+γ
In2Bi+γ+β
25
Landolt-Börnstein New Series IV/11C3
(Bi)+(Sn)
(Sn)
(Bi)+InBi+γ
In5Bi3+In2Bi+γ 50
55
In5Bi3+InBi+γ
Sn, at.%
60
0.00 In Sn 63.77 Bi 36.23
MSIT®
Bi–In–Sn
200
Temperature, °C
Fig. 8: Bi-In-Sn. Vertical section at mass ratio Sn/Bi=30/70, plotted in at.%
200
L
L+(Bi)
L+InBi 100
L+InBi+(Bi)
(Bi)+(Sn)
L+(Bi)+(Sn) InBi+(Bi)+(Sn)
InBi
(Bi)+γ
InBi+(Bi)+γ 0
In 49.50 Sn 0.00 Bi 50.50
Fig. 9: Bi-In-Sn. Vertical section InBi-Sn
10
20
30
0.00 In Sn 43.00 Bi 57.00
40
Sn, at.%
L 200
Temperature, °C
L+(Sn)
L+(Sn)+γ
L+γ L+(Bi)+InBi
(Sn) L+(Sn)
100
(Sn)+γ
L+InBi
(Bi)+(Sn)
76.4 (Bi)+InBi+(Sn)
InBi+(Sn)
(Bi)+γ
InBi+(Sn)+γ (Bi)+(Sn)+γ
InBi+(Bi)+γ 0
In 50.00 Sn 0.00 Bi 50.00
MSIT®
20
40
60
80
Sn
Sn, at.%
Landolt-Börnstein New Series IV/11C3
Bi–In–Sn
Fig. 10: Bi-In-Sn. Vertical section In2Bi-Sn, plotted in at.%
L+(Sn) L
Temperature, °C
200
L+(Sn)+γ
(Sn)
L+γ 100
L+InBi+In5Bi3 L+InBi
In2Bi+In5Bi3+γ 0
In 66.70 Sn 0.00 Bi 33.30
60
(Sn)+γ
L+InBi+γ
InBi+(Sn)+γ
L+InBi+γ
γ
L+In5Bi3+γ L+In2Bi+In5Bi3
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201
(Bi)+γ
InBi+γ
InBi+In5Bi3+γ
40
(Bi)+(Sn) 20
(Bi)+(Sn)+γ
Sn
In, at.%
MSIT®
202
Bi–Sn–Zn
Bismuth – Tin – Zinc Joachim Gröbner Introduction Experimental investigations of the Bi-Sn-Zn phase diagram date from the end of the nineteen century when composition of two coexisting liquids was determined by chemical analysis of separated layers [1891Wri, 1892Wri]. However these results should be considered as semi-quantitative, because later is was shown by [2000Mal1] that immiscibility can not take place above 600°C. A first part of the ternary miscibility gap for this solder system was given by [1949Jae1] at 650 and 750°C. [1949Jae2] showed a liquidus surface which was based on the earlier experimental work of [1923Muz]. [1923Muz] measured thermal arrests of 104 liquid alloys using cooling curves. The first and second recorded arrests could be identified as two- and three-phase equilibria, while the third one corresponded to eutectic equilibrium. Microscopic analysis was also applied by [1923Muz]. [1987Mal] reviewed the experimental data of the Bi-Sn-Zn based on activity determination from emf measurements of [1959Ole], [1966Pta] and calculations of [1977Pel]. A complete thermodynamic calculation of the ternary system is given by [2000Mal1]. This work included a detailed summary of the experimental work. The experimental works of the Bi-Sn-Zn system are summarized in Table 1. Binary Systems The phase diagrams for binary systems are taken from [1994Oht] for Bi-Sn, from [2000Mal2] for Bi-Zn and from [1999Oht] for Sn-Zn according to recommendation of [2000Mal1]. Phase diagram of the Sn-Zn is of simple eutectic type. The phase diagram of this binary system calculated in [1999Oht] is similar to the latest recommendation of [2002Per]. However, the temperature of eutectic reaction l (Zn) + (Sn) according to calculation of [1999Oht] is 6 degrees lower than in [2002Per]. The differences between phase diagram of [1999Oht] and [2002Per] are within uncertainty limit of experimental data on phase equilibria. Solid Phases In this system no ternary phase is known. The crystallographic data of the phases in the Bi-Sn-Zn system and their ranges of stability are listed in Table 2. Invariant Equilibria The ternary system contains only one invariant reaction: L (Bi) + (Sn) + (Zn) found experimentally at 130°C by [1923Muz] and calculated by [2000Mal1]. The compositions of the phases are listed in Table 3. It should be mentioned that there is an inconsistency between the results published in table 2 of [2000Mal1] and figure 7 published in the same work. It is proved in the present evaluation that published thermodynamic parameter reproduce data from table 2 of [2000Mal1], but give slightly different phase diagrams. Therefore, all figures presented in the present evaluation are re-calculated with thermodynamic parameters of [2000Mal1]. Liquidus Surface [1923Muz] presented the first experimental liquidus surface. The calculated liquidus surfaces are given by [1977Pel, 2000Mal1, 2003Moe] and [2004Oht]. The experimental and calculated liquidus surfaces are in a quite reasonable agreement. Liquidus surface calculated using thermodynamic parameters published by [2000Mal1] is presented in Fig. 1.
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Landolt-Börnstein New Series IV/11C3
Bi–Sn–Zn
203
Isothermal Sections Figures 2 to 4 show isothermal sections at 135°C, 170°C and 250°C calculated using thermodynamic description of [2000Mal1]. Experimental isothermal sections are not available so far. Temperature – Composition Sections Vertical sections at constant 5 mass% Sn and constant 40 mass% Sn calculated from thermodynamic data of [2000Mal1] are given in Figs. 5 and 6. They are in a good agreement with experimental data of [1923Muz]. Thermodynamics The activity of Zn in liquid phase of the Bi-Sn-Zn system was derived from emf measurements in works of [1959Ole, 1966Pta, 1967Glu, 1968Lou, 1985Mat] and from vapor pressure of Zn measured by Knudsen’s effusion method [1961Yok]. The Zn activities at 450, 551 and 641°C calculated from thermodynamic data of [2000Mal1] are reproduced in Figs. 7 to 9. They are in a good agreement with experimental data of [1959Ole] and [1966Pta]. [1986Zha1] and [1986Zha2] calculated the vertical section at constant 5 at.% Zn using the theorem of corresponding relations between neighbor phase regions. [1971Pta] calculated limited solubility of metals at 441, 484, 532 and 569°C based on measurements from [1966Pta]. First thermodynamic calculations of the entire ternary system by the Calphad method was made by [1977Pel], based on emf measurements of [1966Pta] for the ternary system and emf data of [1977Bal] for the quaternary Bi-Cd-Sn-Zn system. [1986Mat] calculated infinite dilution properties using a new approach, based on free volume theory and in extension of Kapoor’s model for binary substitutional molten metallic solutions. The calculated excess energies were in a good agreement with experimental data obtained by emf method in the same work [1986Mat]. [1999Ohn] and [2000Mal1] presented a thermodynamic calculation which reproduced all available experimental data. [2003Moe] and [2004Oht] present a Calphad-type calculation based on different thermodynamic descriptions of binary data. However, both thermodynamic descriptions [2003Moe] and [2004Oht] reproduce quite similar liquidus surface for the Bi-Sn-Zn system as thermodynamic parameters of [2000Mal1]. Notes on Materials Properties and Applications Numerous efforts have been undertaken in recent years to develop Pb-free substitutes for the lead-tin soldering alloys widely used in microelectronics. System containing Ag, Bi, Cu, In, Sb, Sn, Zn and other elements have been experimentally and thermodynamically studied [1977Pel, 1997Hua, 1997Yoo, 1999Ohn, 2003Moe, 2004Oht]. The candidates for Pb-free solders should fulfill many requirements such as wettability of substrate, formation of strong chemical bonds with substrate, low melting temperature, good mechanical properties, corrosion resistance, safe for health and environment and so on [2003Moe]. [2001Lee] examined the effects of Zn addition on the interfacial and mechanical properties and the of the solder joint with the Cu substrate. [2003Lin] investigated morphology and growth kinetics of this joint. Jointing properties including joint strength ant interfacial structure of Bi-Sn-Zn solders were given by [2003Sug]. [2004Sho] investigated tensile properties of low-melting solders of the Bi-Sn-Zn system. For the alloy Sn-8Zn-3Bi an approximately double tensile strength was found in the investigated strain-rate and temperature range compared with those of commercial Sn-37Pb alloy. Microhardness of the Sn-9Zn and Sn-8Zn-3Bi solders was experimentally measured and compared with Sn-37Pb solder in [2005Isl]. Micro-hardness of Sn-37Pb and Sn-9Zn (eutectic) decreases with increasing reflow temperature while it decreases for Sn-8Zn-3Bi (non-eutectic). It was shown by SEM investigation that the formation of Pb rich islands in Sn-37Pb, formation of Zn rod from spheroids in Sn-9Zn and precipitation of Bi rich phase in Sn-8Zn-3Bi are the important features that contribute to different hardness nature [2005Isl].
Landolt-Börnstein New Series IV/11C3
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204
Bi–Sn–Zn
References [1891Wri]
[1892Wri]
[1923Muz] [1949Jae1]
[1949Jae2]
[1959Ole]
[1961Yok] [1966Pta]
[1967Glu]
[1968Lou]
[1971Pta] [1977Bal]
[1977Pel]
[1985Mat]
[1986Mat]
[1986Zha1]
[1986Zha2]
MSIT®
Alder-Wright, C.R., Thompson, C., “On Certain Ternary Alloys. Part III. Alloys of Bismuth, Zinc, and Tin, and of Bismuth, Zinc, and Silver”, Proc. Roy. Soc., 49, 156-173 (1891) (Experimental, 2) Alder-Wright, C. R., “On Certain Ternary Alloys. Part V. Determination of Various Critical Curves and Their Tie-Lines and Limiting Points”, Proc. Roy. Soc., 50, 372-395 (1892) (Experimental, 11) Muzaffar, S.D., “CCLXV.-Equilibrium of the Ternary System Bismuth-Tin-Zinc”, J. Chem. Soc., 123, 2341-2354 (1923) (Experimental, Phase Relations, #, 1) Jaenecke, E., “Cd-Pb-Zn, Pb-Sn-Zn, Bi-Cd-Zn, Bi-Sn-Zn, Al-Cd-Sn, Al-Pb-Sn, Al-Bi-Sn, Al-Cd-Zn” (in German), Kurzgefasstes Handbuch aller Legierungen, 431-433 (1949) (Phase Diagram, Phase Relations, 0) Jaenecke, E.,“Ternaere Legierungen aus je Drei der Fuenf Metalle Bi-Cd-Pb-Sn-Zn“, Kurzgefasstes Handbuch aller Legierungen, 436-438 (1949) (Phase Diagram, Phase Relations, 1) Oleari, L., Fiorani M., “Thermodynamic Research on Metal Systems. XI. Determination of the Miscibility Gap by Computation though the Thermodynamic Quantities for the Ternary System Zn-Sn-Bi”, Ricerca Sci., 29, 2219-2229 (1959) (Experimental, Thermodyn., Phase Relations, 11) Yokokawa, T., Doi, A., Niwa, K., “Thermodynamic Studies on Liquid Ternary Zinc Solutions”, J. Phys. Chem., 65, 202-205 (1961) (Experimental, Thermodyn., 8) Ptak, W., Moser, Z., “Investigation of the Thermodynamic Properties in the Liquid Zn-Bi-Sn Alloys” (in Polish), Arch. Hutn., 11(4), 289-324 (1966) (Experimental, Phase Diagram, Thermodyn., 8) Gluck, J.V., Pehlke, R.D., “Solute Interactions with Zinc in Dilute Solution with Molten Bismuth: 1 - Third-Element Effects”, Trans. Met. Soc. AIME, 239, 36-47 (1967) (Experimental, 37) Louvet, R.L., Gluck, J.V., Pehlke, R.D., “Thermodynamic Interactions Between Zinc and Bismuth in Dilute Solution in Molten Tin”, Trans. Metall. Soc. AIME, 242(11), 2369-2371 (1968) (Experimental, Thermodyn., 21) Ptak, W., Moser, Z., “Range of Limited Solubility in Zn-Bi-Sn-Cd Quaternary Alloys”, Bull. Acad. Polon. Sci., Ser. Sci. Tech. 19, 1-6 (1971) (Experimental, Phase Relations, 8) Bale, C.W., Pelton, A.D., Rigaud, M. “The Measurement and Analytical Computation of Thermochemical Properties in a Quaternary System - Zn-Cd-Bi-Sn”, Z. Metallkd., 68, 69-74 (1977) (Experimental, Thermodyn., Calculation, 10) Pelton, A.D., Bale, C.W., Rigaud, M., “The Computer Calculation of a Quaternary Phase Diagram with Liquid-Liquid Immiscibility - The Zn-Cd-Bi-Sn System”, Z. Metallkd., 68, 135-140 (1977) (Calculation, 19) Mathur, V.N.S., Kapoor, M.L., “Thermodynamics of Dilute Solutions of Zinc in Binary Molten Bismuth-Tin Solvents at 773 K”, Z. Metallkd., 76, 16-23 (1985) (Experimental, Thermodyn., 39) Mathur, V.N.S., Kaushal, G.C., Kapoor, M.L., “An Approach of Thermodynamics of Dilute Solute in Ternary Randomly Distributed Substitutional Molten Metallic Solutions”, Z. Metallkd., 77, 536-544 (1986) (Thermodyn., Theory, 22) Zhao, M., Xu, B., Fan, X., Fu, Z., “Progress in the Theorem of Corresponding Relations in Isobaric Phase Diagrams and its Application”, High Temp. Sci., 22, 1-14 (1986) (Calculation, Phase Diagram, 6) Zhao, M., Fu, Z., “The Principle of Direct Calculation of the Vertical Section (Isopleth) in Phase Diagrams with the Aid of the Theorem of Corresponding Relations”, High Temp. Sci., 22, 15-25 (1986) (Calculation, Phase Diagram, Theory, 3)
Landolt-Börnstein New Series IV/11C3
Bi–Sn–Zn [1987Mal] [1994Oht]
[1997Hua] [1997Yoo]
[1999Ohn]
[1999Oht]
[2000Mal1]
[2000Mal2] [2001Lee]
[2002Per]
[2003Lin]
[2003Moe]
[2003Sug]
[2004Oht]
[2004Sho]
[2005Isl]
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Mallik, A.K., Marathe, K.V., “The Bismuth-Tin-Zinc System”, J. Alloy Phase Diagrams, 3(2), 128-132 (1987) (Phase Diagram, Review, 3) Ohtani, H., Ishida, K., “A Thermodynamic Study of the Phase Equilibria in the Bi-Sn-Sb System”, J. Electron. Mater., 23(8), 747-755 (1994) (Calculation, Experimental, Phase Diagram, Thermodyn., 36) Hua, F., Glazer, J., “Lead-Free Solders For Electronic Assembly”, Des. Reliab. Solders Solder Interconnect., Proc. Symp., 65-73 (1997) (Experimental, #, 63) Yoon, S.W., Soh, J.R., Lee, H.M., Lee, B.-J., “Thermodynamics-Aided Alloy Design and Evaluation of Pb-Free Solder, Sn-Bi-In-Zn System”, Acta Mater., 45(3), 951-960 (1997) (Calculation, Experimental, Phase Relations, #, 31) Ohnuma, I., Liu, X.J., Ohtani, H., Ishida, K., “Thermodynamic Database for Phase Diagrams in Micro-Soldering Alloys”, J. Electron. Mater., 28(11), 1164-1171 (1999) (Calculation, Phase Relations, #, 22) Ohtani, H., Miyashita, M., Ishida, K., “Thermodynamic Study of Phase Equilibria in the Sn-Ag-Zn System”, J. Jpn. Inst. Mat., 63, 685-694 (1999) (Calculation, Phase Relations, #, 68) Malakhov, D.V., Liu, X.J., Ohnuma, I., Ishida, K., “Thermodynamic Calculation of Phase Equilibria of the Bi-Sn-Zn System”, J. Phase Equilib., 21(6), 514-520 (2000) (Calculation, Phase Relations, Thermodyn., #, *, 32) Malakhov, D.V., “Thermodynamic Assessment of the Bi-Zn System”, Calphad 24, 1-14 (2000) (Calculation, Phase Relations, Thermodyn., 51) Lee, K.K., Kang, C.S., Lee, D.J., “Mechanical Properties of Sn-Bi-Zn Solder/Cu Joints” (in Japanese), J. Jpn. Inst. Met., 65(12), 1091-1095 (2001) (Experimental, Mechan. Prop., Morphology, 20) Perrot, P., Batista, S., Xing, X., “Sn - Zn (Tin - Zinc)”, MIST Binary Evaluation Program in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science Interantional Services GmbH, Stuttgart; Document ID: 20.11449.1.20 (2002) (Crys. Structure, Phase Diagram, Assessment, 10) Lin, W.H., Chuang, T.H., “Interfacial Reactions Between Liquid Sn-8Zn-3Bi Solders and Cu Substrates”, J. Mater. Eng. Perform., 12(4), 452-455 (2003) (Experimental, Morphology, 14) Moelans, N., Kumar, K.C.H., Wollants, P., “Thermodynamic Optimization of the Lead-Free Solder System Bi-In-Sn-Zn”, J. Alloys Compd., 360, 98-106 (2003) (Calculation, Phase Diagram, Thermodyn., #, 44) Sugizaki, T., Nakao, H., Kimura, T., Watanabe, T., “BGA Jointing Property of Sn-8.8 mass% Zn and Sn-8.0 mass% Zn-3.0 mass% Bi Solder on Electroless Nickel-Phosphorus/Immersion Gold Plated Substrates”, Mater. Trans., JIM, 44(9), 1790-1796 (2003), translated from J. Jpn. Inst. Met., 67(5), 232-238 (2003) (Crys. Structure, Experimental, Mechan. Prop., 11) Ohtani, H, Ono, S., Doi, K., Hasebe, M., “Thermodynamic Study of Phase Equilibria in the Sn-Ag-Bi-Zn Quaternary System”, Mater. Trans., JIM, 45(3), 614-624 (2004) (Calculation, Experimental, Phase Diagram, Thermodyn., 25) Shohji, I., Yoshida, T., Takahashi, T., Hioki, S., “Comparison of Low-Melting Lead-Free Solders in Tensile Properties with Sn-Pb Eutectic Solder”, J. Mater. Sci.: Mat. Electron., 15(4), 219-223 (2004) (Experimental, Mechan. Prop., Morphology, 16) Islam, R.A., Wu, B.Y., Alam, M.O., Chan, Y.C., Jillek, W., “Investigation on Microhardness of Sn-Zn Based Lead-Freee Solder Alloys as Replacement of Sn-Pb Solder”, J. Alloys Compd., 392(1-2), 149-158 (2005) (Experimental, Mechan. Prop., 13)
MSIT®
Bi–Sn–Zn
206
Table 1: Investigations of the Bi-Sn-Zn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1891Wri, 1892Wri]
chemical analysis of separate melts 600 to 800°C
[1923Muz]
thermal analysis, microscopy
104 alloys
[1959Ole]
emf
450, 500, 550°C, activity of Zn
[1961Yok]
Knudsen effusion method
317 to 387°C
[1966Pta]
emf
441 to 641°C, activity of Zn
[1967Glu]
emf
450 to 650°C, activity of Zn
[1968Lou]
emf
450 to 650°C, activity of Zn
[1971Pta]
emf
427 to 627°C, activity of Zn
[1985Mat]
emf
500°C
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Bi) < 271.442
hR6 R3m As
a = 454.613 c = 1186.152
at 31°C [Mas2]
(Sn) 231.9681 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
at 25°C [Mas2]
(Zn) < 419.58
hP2 P63/mmc Mg
a = 266.50 c = 494.70
at 25°C [Mas2]
Table 3: Invariant Equilibria Reaction L (Bi) + (Sn) + (Zn)
MSIT®
T [°C]
130
Type
E
Phase
L (Bi) (Sn) (Zn)
Composition (at.%) Bi
Sn
Zn
39.39 95.75 14.03 6.03 #10–4
54.35 3.44 85.41 4.23 # 10–4
6.25 0.81 0.56 99.999
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Bi–Sn–Zn
207
Bi Fig. 1: Bi-Sn-Zn. Calculated liquidus surface projection
Data / Grid: at.% Axes: at.%
e2
(Bi)
20
80
e'3 40
60
e4
60
40
(Zn)
50 0° C
80
E1
L 20
450 400
300
350
250
200 (Sn)
e3
20
Zn
40
60
80
Bi Fig. 2: Bi-Sn-Zn. Calculated isothermal section at 135°C
Sn
e1
Data / Grid: at.% Axes: at.%
(Bi)
20
80
40
60
(Bi)+L+(Sn) L
(Bi)+L+(Zn)
60
40
80
20
(Sn)+L+(Zn)
(Sn) (Zn)
Zn
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Sn
MSIT®
Bi–Sn–Zn
208
Bi Fig. 3: Bi-Sn-Zn. Calculated isothermal section at 170°C
Data / Grid: at.% Axes: at.%
(Bi)
20
80
40
60
(Zn)+L+(Bi)
L
60
40
80
20
(Sn)+L+(Zn) (Sn) (Zn)
20
Zn
40
60
80
Bi Fig. 4: Bi-Sn-Zn. Calculated isothermal section at 250°C
Sn
Data / Grid: at.% Axes: at.%
(Bi)
20
80
40
60
60
40
L 80
(Zn)
Zn
MSIT®
20
20
40
60
80
Sn
Landolt-Börnstein New Series IV/11C3
Bi–Sn–Zn
Fig. 5: Bi-Sn-Zn. Vertical section at constant 5 mass% Sn, plotted in at.%
209
600
L 500
Temperature, °C
L1+L2 400
L+(Zn) 300
200
L+(Bi)+(Zn) 100
(Sn)+(Bi)+(Zn) 0
Zn 97.18 Sn 2.82 0.00 Bi
10
20
30
40
50
60
70
80
Bi, at.%
0.00 Sn 8.48 Bi 91.52
90 Zn
400
Fig. 6: Bi-Sn-Zn. Vertical section at constant 40 mass% Sn, plotted in at.%
L
Temperature, °C
300
L+(Zn)
200
100
(Zn)+(Sn) (Zn)+(Bi)+(Sn) 0
Zn 73.14 Sn 26.86 0.00 Bi
Landolt-Börnstein New Series IV/11C3
10
20
Bi, at.%
30
40
Zn 0.00 Sn 53.99 Bi 46.01
MSIT®
Bi–Sn–Zn
210
1.0
Fig. 7: Bi-Sn-Zn. Calculated activity of Zn in liquid phase at 450°C at fixed x(Sn)/x(Sb) ratio
6 0.9
4
5 3
0.8
2 0.7
Activity of Zn
1 0.6
1 - Sn-Zn binary 2 - x(Sn)/x(Bi)=3 3 - x(Sn)/x(Bi)=1 4 - x(Sn)/x(Bi)=1/3 5 - x(Sn)/x(Bi)=1/9 6 - Bi-Zn binary
0.5 0.4 750
500
0
0
0.2
0.4
0.6
1.0
0.8
Zn, at.%
Bi
Data / Grid: at.% Axes: at.%
Fig. 8: Bi-Sn-Zn. Calculated isoactivity lines for Zn at 551°C 20
80
40
60
60
40
80
20
0.8 0.9
Zn
MSIT®
20
0.6 0.2
0.4 40
60
80
Sn
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Bi–Sn–Zn
211
Bi
Data / Grid: at.% Axes: at.%
Fig. 9: Bi-Sn-Zn. Calculated isoactivity lines for Zn at 641°C 20
80
40
60
60
40
80
20
0.8
0.9
Zn
Landolt-Börnstein New Series IV/11C3
20
0.4
0.6 40
60
0.2 80
Sn
MSIT®
212
Co–Cu–Mn
Cobalt – Copper – Manganese Lazar Rokhlin, Evgeniya Lysova, Nataliya Bochvar Introduction Based on the adjoining binary phase diagrams, [1909Jae] suggested a possible configuration for the Co-Cu-Mn phase diagram. It was characterized by the existence of a three-phase monovariant peritectic transformation region extending from the Cu-Mn side to the (Cu, Mn,Co) ternary solid solution and by the absence of an invariant transformation during crystallization. Such constitution of the Co-Cu-Mn phase diagram was confirmed by the theoretical work of [1942Hir], who used calculations of the free energy of the alloys. The Co-Cu-Mn phase diagram was studied experimentally by [1938Koe] using thermal, magnetic and microstructural analyses in the region from the Co-Cu side up to 40 mass% Mn. A partial liquidus surface and the monovariant line bounding the field of joint crystallization of the (Cu, Mn,Co) and (Co) solid solutions were constructed. In addition, the isolines of Curie temperature from 20 to 800°C were drawn assuming them as tie lines between the (Cu, Mn,Co) and (Co) solid solutions. [1938Koe] also constructed four vertical sections at 10, 20, 30 and 40 mass% Mn. [1958Top1, 1958Top2] studied Mn rich alloys with up to 50 mass% Co and up to 50 mass% Cu, which had been annealed at 500°C. The studied alloys contained (Mn) and ( Mn) solid solutions. The results of [1958Top1, 1958Top2] contradict, however, the Co-Mn phase diagram through the existence of an extended (Mn) solid solution [1979Cha]. [1982Has] investigated phase compositions of Co-Cu-Mn alloys at several temperatures in the range 877 to 1277°C. Based on their own experimental data, six isothermal sections of the phase diagram were presented. [1979Dri, 1979Cha] reviewed the Co-Cu-Mn phase diagram using data from [1938Koe] and [1958Top1, 1958Top2]. The methods and composition ranges studied in the different investigations are described in Table 1. Binary Systems The binary Co-Cu and Cu-Mn systems are accepted from the MSIT Evaluation Program after [2006Ans] and [2005Tur], respectively. The Co-Mn phase diagram is taken from [Mas2]. Solid Phases No ternary compound has been found in the Co-Cu-Mn system. Three binary ordered phases, 1, 2 and 3, have been established in the Cu-Mn system. There are extended solid solutions based on the separate allotropic forms of the pure metals. Characteristics of the solid phases of the Co-Cu-Mn system are shown in Table 2. Liquidus, Solidus and Solvus Surfaces The partial liquidus surface of the Co-Cu-Mn ternary system is shown in Fig. 1. The region adjoining the Co-Cu side is divided by a monovariant line into a surface corresponding to the primary crystallization of the (Co) and primary solid solutions. The monovariant line corresponds to the L + (Co) (Cu) peritectic reaction, the temperature of which falls from 1113.3°C at the Co-Cu binary down to ~800°C on the addition of ~45 at.% Mn. Its location is taken from [1938Koe] with a little correction necessary to meet the accepted Co-Cu binary diagram [2006Ans]. Taking into consideration the binary systems Co-Mn and Cu-Mn, the liquidus surface in the Mn rich region should include a surface of primary crystallization of the (Mn) and ( Mn) solid solutions. The location of these surfaces has not been established. Figure 2 shows the partial solidus surface in the region adjoining the Co-Cu side. It is divided by a double saturation line (which bounds the L + (Co) + three-phase region) into the surface of the joint crystallization of (Co) + and termination of the ternary solid solution. The double saturation line bounding the three-phase region is drawn after [1938Koe] with a small correction to meet the binary Co-Cu
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Co–Cu–Mn
213
[2006Ans]. In accordance with [1938Koe], tie lines corresponding to the isolines of Curie temperature are shown on the concentration field outlined by the double saturation line. Isothermal Sections Isothermal sections at different temperatures are presented in Figs. 3-8. They were taken from the calculations of [1982Has] with some corrections to meet the accepted binary phase diagrams. Temperature – Compositions Sections Two vertical sections are shown in Figs. 9 and 10. They are drawn according to [1938Koe] with some corrections to meet the accepted binary phase diagrams. In accordance with the accepted Cu-Mn phase diagram, the sections must cross the 3, (Mn) + 3 and (Mn) + 2 areas which are shown in the vertical sections only approximately. Notes on Materials Properties and Applications The Co-Cu-Mn system forms part of the commercial alloy SA04 (AP Cu-Mn-28Ni-29Co5-Si-B) used for the soldering of hot and corrosion resistant steels and alloys. The soldering temperature is 1050°C and its seal strength is 350 MPa. [1947Fre] discovered the propensity for alloys containing 2-4 mass% Co and up to 15 mass% Mn to age harden at temperatures of 480, 540 and 595°C. [1973Mas] established the anti-ferromagnetic behavior of Mn rich Co-Cu-Mn alloys and their Elinvar characteristics. Nonmagnetic ternary alloys of 32.8 - ~87.5% Mn and less than 67% Cu and 24% Co (mass%) were found to have pronounced Elinvar characteristics and fairly good corrosion resistance. [1976Ano] reported that Co-Cu-Mn alloys with a typical composition of 10%Co-58%Cu-31.5%Mn (mass%) can be successfully used as brazing materials for gas turbine applications in place of alloys containing expensive precious metals. [2003Cao] investigated the magnetic properties of the Mn layer-doped Co/Cu92Mn8/Co layered structures. Their behavior in magnetic fields was established. A summary of studies of the materials properties and applications related to the Co-Co-Mn alloys is presented in Table 3. References [1909Jae] [1938Koe] [1942Hir] [1947Fre] [1958Top1]
[1958Top2]
[1973Mas] [1976Ano]
Landolt-Börnstein New Series IV/11C3
Jänecke, E., “Ternary Alloys of Cu, Ag, Au; Cr, Mn; Fe, Co, Ni; Pd. Pt Metals” (in German), Z. Phys. Chem., 67, 668-688 (1909) (Experimental, Phase Diagram, Phase Relations, 23) Köster, W., Wagner, E., “The Ternary Co-Mn-Cu System” (in German), Z. Metallkd., 30, 352-353 (1938) (Experimental, Phase Diagram, Magn. Prop., #, 3) Hirone, T., Katayama, T., “On the Constitution of the Ternary Alloys”, Sci. Rep. Tohoku Imp. Univ., 30, 109-124 (1942) (Calculation, Phase Diagram, 9) Frederickson, J.W., “Age Hardening Copper - Cobalt - Manganese Alloys”, Trans. ASM, 38, 593-617 (1947) (Experimental, Mechan. Prop., 8) Topchiashvili, L.I., “The Influence of Iron, Cobalt, and Nickel on the Structure and Properties of Mn-Cu Alloys”, Russ. J. Inorg. Chem. (Engl. Transl.), 3, 253-256 (1958) (Experimental, Mechan. Prop., Electr. Prop., 0) Topchiashvili, L.I., Agladze, R.I., Mokhov, V.M., “Investigation of Mn-Cu-Co Alloys” (in Russian), Zh. Neorg. Khim., 3(11), 2537-2544 (1958) (Experimental, Mechan. Prop., Electr. Prop., 5) Masumoto, H., Sawaya S., Kikuchi, M., “Nonmagnetic Elinvar – Type Alloys in the Mn-Cu-Co System”, Trans. JIM, 14(1), 56-61, (1973) (Experimental, Mechan. Prop., 5) Anon, “Cu-Mn-Co Braze Alloy Beats Precious Metals”, Mater. Eng., 83(1), 20 (1976) (Mechan. Prop., 0)
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214 [1979Cha]
[1979Dri]
[1982Has]
[2003Cao]
[2005Tur]
[2006Ans]
Co–Cu–Mn Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Co-Cu-Mn”, INCRA Monograph Series 6 Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Systems, NSRD, Washington, 422-425 (1979) (Phase Diagram, Phase Relations, Review, 3) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Cu-Co-Mn” in “Binary and Multicomponent Copper - Base Systems” (in Russian), Abrikosov, N.X., (Eds.), Nauka, Moscow, 145-148 (1979) (Phase Diagram, Review, 2) Hasebe, M., Oikawa, K., Nishizawa, T., “Computer Calculation of Phase Diagrams of Co-Cu-Mn and Co-Cu-Ni Systems”, Nippon Kinzoku Gakkai Shi, 46(6), 584-590 (1982) (Calculation, Phase Diagram, #, 18) Cao, R., Ye, J., Zhao, Y., Jin, Q.Y., Wang, H., “Study of Magnetoresistance in Mn Layer-Doped Co/Cu/Co Multilayers”, Phys. Status Solidi A, 195(2), 434-439 (2003) (Experimental, Magn. Prop., 7) Turchanin, M., Agraval, P., Groeber, J., Matusch, D., Turkevich, V., “Cu-Mn (Copper-Manganese)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14136.1.20, (2005) (Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., Assessment, 25) Ansara I., Ivanchenko, V., Turchanin M., Agraval P., “Co-Cu (Cobalt-Copper)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2006) (Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., Assessment, 19)
Table 1: Investigations of the Co-Cu-Mn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1938Koe]
Thermal, magnetic, microstructural analyses
Alloys with constant Mn = 10, 20, 30, 40 mass%. Temperature - composition sections at constant Mn = 10, 20, 30, 40 mass%. Liquidus surface with monovariant peritectic line and miscibility gap in solid state.
[1958Top1, Melting in a high - frequency furnace 1958Top2] in pressed and sintered magnesite crucibles, casting in mild steel mold, homogenization at 900-800-700-600-500°C for 1 day at each temperature and added 500°C for 20 days in sealed quartz ampoules. Metallographic, hardness, resistivity, dilatometry.
Alloys in the Mn corner (100 - 50 mass% Mn) up to 50 mass% Co and 50 mass% Cu. These alloys have a two-phase structure in the annealed state: (Mn) and ( Mn,Co)
[1982Has]
Experimental results on the miscibility gap ( and f) between 877 and 1077°C and on the /L equilibrium between 1277 and 977°C. Six calculated isothermal sections at 1177, 1077, 977, 877, 777 and 577°C
MSIT®
Ternary diffusion couple technique, computer calculation
Landolt-Börnstein New Series IV/11C3
Co–Cu–Mn
215
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, (Cu, Mn,Co) < 1495
cF4 Fm3m Cu
Lattice Parameters Comments/References [pm]
(Cu) < 1084.62
a =361.46
pure Cu at 25°C [Mas2] dissolves 9.1 at.% Co at 1113,3°C [Mas2] and 100% Mn [2005Tur]
( Mn) 1138 - 1087
a = 386.0
pure Mn [Mas2] dissolves up to 5 at.% Co at 1138°C [Mas2] and 100 at.% Cu [2005Tur]
(Co) 1495 - 422
a = 354.47 a = 356.88
[Mas2] at 520°C [2006Ans] dissolves 11.7 at.% Cu at 1113.3°C and ~59 at.% Mn at 1161°C [Mas2]
(JCo) < 422
hP2 P63/mmc Mg
a = 250.71 c = 406.86
at 25°C [Mas2]
( Mn) 1246 - 1138
cI12 Im3m W
a = 308.0
dissolves 9.0 at.% Co at ~1185°C [Mas2] and 13 at.% Cu at 1097°C [2005Tur]
(Mn) 1087 - 707
cP20 P4132 Mn
a = 631.52
dissolves 41 at.% Co at 1161°C [Mas2] and 2.03 at.% Cu at 706°C [2005Tur]
(Mn) < 707
cI58 I43m Mn
a = 891.26
dissolves up to 3 at.% Co [Mas2] and 2.3 at.% Cu at 706°C [2005Tur]
3 (Cu-Mn phase) 700
c**
-
[2005Tur]
2, MnCu3 450
c**
-
[2005Tur]
1, MnCu5
c**
-
[2005Tur]
Landolt-Börnstein New Series IV/11C3
MSIT®
Co–Cu–Mn
216
Table 3: Investigations of the Co-Cu-Mn Materials Properties Reference
Method/Experimental Technique
Type of Property
[1947Fre]
Melting in a high - frequency induction furnace, cold - worked and solution - treated for 0.5 h at temperatures varying from 650 to 980°C, aging at temperatures from 315 to 650°C for times of 0.5, 1, 2, 4, 8, 16, 24 and 48 h
Time - hardness curves at various aging temperatures for cold - reduced and annealed Co-Cu-Mn alloys with 2, 3 and 4 mass% Co and 1 and 10 mass% Mn
[1973Mas]
Measurement of Young's modulus at –150 - ~400°C, rigidity modulus and hardness at room temperature. The resonance frequency 470 - 880 Hz. Annealing temperature - 900°C, cold working or water quenching and reheating for 1 h at 200, 400, 600 and 800°C
Elinvar characteristics for Co-Cu-Mn alloys of 0 - ~40.0 mass% Co, 0.1 - ~67.0 mass% Cu, 38.0 - ~88 mass% Mn subjected to various heat treatments and cold working. Nonmagnetic Elinvar - type alloys consisted of less than 67 mass% Cu, 24 mass% Co and 32.8 - ~87.5 mass% Mn and have fairly good corrosion resistance
[1976Ano]
Tensile, shear and fatigue strength
The alloy with10 mass% Co, 58 mass% Mn and 31.5 mass% Cu can be used as brazing materials in place of alloys of precious metals
[2003Cao]
Magnetic Properties
Giant magnetoresistance in Mn layer - doped Co/Cu/Co multilayers
Cu
Data / Grid: at.% Axes: at.%
Fig. 1: Co-Cu-Mn. Liquidus surface projection
p3
20
min 40
80
γ 60
(α Co) 60
40
p4 80
Mn
MSIT®
(δMn) p1
20
(β Mn) 20
p2 40
60
80
Co
Landolt-Börnstein New Series IV/11C3
Co–Cu–Mn
217
Cu
Data / Grid: at.% Axes: at.%
Fig. 2: Co-Cu-Mn. Solidus surface with isolines of Curie temperature 20
80
40
60
γ +(αCo) 60
40
20°C 200 80
600 400
800
20
γ
20
Mn
40
60
80
Cu
Co
Data / Grid: at.% Axes: at.%
Fig. 3: Co-Cu-Mn. Isothermal section at 1777°C 20
80
40
60
L
60
40
80
20
γ (δMn)
Mn
Landolt-Börnstein New Series IV/11C3
(β Mn)
20
40
60
80
Co
MSIT®
Co–Cu–Mn
218
Cu
Data / Grid: at.%
Fig. 4: Co-Cu-Mn. Isothermal section at 1077°C
Axes: at.%
(Cu)
20
80
40
60
L+(Cu)+γ L 60
40
80
20
γ
Mn
(β Mn)
20
40
60
80
Cu
Co
Data / Grid: at.% Axes: at.%
Fig. 5: Co-Cu-Mn. Isothermal section at 977°C
(Cu)
20
80
L+(Cu)+γ 40
60
L
60
40
γ
80
Mn
MSIT®
(β Mn) 20
40
20
60
80
Co
Landolt-Börnstein New Series IV/11C3
Co–Cu–Mn
219
Cu
Data / Grid: at.% Axes: at.%
Fig. 6: Co-Cu-Mn. Isothermal section at 877°C 20
80
L 40
tion ma for s n tra
60
γ
80
Mn
(β Mn) 20
20
40
60
80
Cu
Axes: at.%
20
60
80
(β Mn) 20
80
tion rma o f s tran tic gne a m
40
Landolt-Börnstein New Series IV/11C3
Co
Data / Grid: at.%
Fig. 7: Co-Cu-Mn. Isothermal section at 777°C
Mn
40
tic gne ma
60
γ
60
40
γ +(αCo)
40
60
20
80
Co
MSIT®
Co–Cu–Mn
220
Cu Fig. 8: Co-Cu-Mn. Isothermal section at 577°C
Data / Grid: at.% Axes: at.%
γ
20
80
tion rma nsfo a r t c neti mag
γ3 γ 40
60
Mn )+( βM n)+ γ
40
20
(α
80
60
(αCo)
(β Mn) 20
Mn (αMn)
Fig. 9: Co-Cu-Mn. Vertical section at 10 mass% Mn, plotted
in at.%
40
60
80
Co
1500
L 1250
Temperature, °C
L+(α Co) 1000
L+γ
L+(α Co)+γ
γ
750
(α Co)
mag neti c tra nsfo rma tion
500
γ3
250
γ+(α Co) 0
Mn 10.65 Co 89.35 Cu 0.00
MSIT®
20
40
Cu, at.%
60
80
Mn 11.39 Co 0.00 Cu 88.61
Landolt-Börnstein New Series IV/11C3
Co–Cu–Mn
Fig. 10: Co-Cu-Mn. Vertical section at 30 mass% Mn, plotted
Temperature, °C
in at.%
221
1500
L 1250
L+(α Co)
1000
L+γ (α Co) L+(α Co)+γ
750
γ
(α Co)+γ 500
(α Mn)+γ3 250
0
Mn 31.50 Co 68.50 Cu 0.00
Landolt-Börnstein New Series IV/11C3
magn etic tr ansfo rmatio n
10
20
(α Mn)+γ2 30
Cu, at.%
40
50
60
Mn 33.14 Co 0.00 Cu 66.86
MSIT®
222
Cr–Ni–P
Chromium – Nickel – Phosphorus Joachim Gröbner Introduction Electroless Ni-P has been widely used as an excellent protective coating in versatile applications due to its mechanical and chemical properties. Amorphous interlayer of Ni-Cr-P can be also used by transient liquid phase (TLP) metallic bonding for joining. The first investigation in the Cr-Ni-P system was carried out by Nowotny and Henglein [1948Now] who melted samples from pure powder of the elements up to 1200°C for 70 h. They assumed continuous solution between Cr2P and Ni2P but mentioned not complete reaction of the samples. [1962Lun] reported a ternary ) phase in the system Cr-Ni-P with a solubility range of ~3 at.% P. The samples were prepared by dropping pellets of red phosphorus into a melt of nickel. The resulting master alloys were arc-melted with Cr powder and heat treated at 1050°C for 20 to 30 d. An isothermal section at 797°C is given by [1984Ori]. The samples were prepared from element powders by press molding in steel crucibles at a pressure of 4.9 MPa. The resulting briquets were heat treated twice for 500 h at 797°C, in between re-melted in arc furnace under argon atmosphere and afterwards quenched in cold water. The samples were characterized by Debye-Scherrer powder method. One ternary phase with the composition Cr1.2Ni0.8P (-1) and large ternary solid solubilities of the phases CrP, Cr12P7, Cr3P and Ni2P were found. The ternary Sigma phase ()) reported by [1962Lun] was not confirmed by [1984Ori] at 797°C. [1990Muc] described the phase formation under combustion conditions and gave an isothermal section similar to [1984Ori]. Binary Systems The binary systems are accepted as given by [Mas2]. Solid Phases The solid phases are given in Table 1. The dependence of lattice parameters of the Cr3P, Cr12P7, and Ni2P phases after [1984Ori] is given in Figs. 1 to 3. Lattice parameters up to the maximal solubility of 0.7 mole fraction NiP in the CrP phase at 20°C by [1986Fje] is given in Fig. 4. Isothermal Sections The isothermal section at 797°C given in Fig. 5 is based on the data of [1984Ori]. Modifications have been introduced to the diagram of [1984Ori] to remove artificial homogeneity ranges of the binary phases (along the axis) which where not established experimentally and do not correspond to the accepted binary diagrams. At this temperature the binary phases CrP, Cr12P7, Cr3P and Ni2P dissolve large amounts of the third element (see Table 1). Cr is replaced by Ni on the same site in the crystal lattice and forms line compounds between the binary Cr-P and the corresponding Ni-P phases. A partial isothermal section at 1050oC by [1962Lun] is given in Fig. 6. The ternary Sigma phase ()) reported by [1962Lun] is most probably part of a ternary solid solution of the binary Ni3P phase, since both phases crystallize in the same crystal structure. Notes on Materials Properties and Applications [1997Yeh] describes transient liquid phase (TLP) metallic bonding for joining Inconel superalloy by amorphous interlayer of Ni-Cr-P at 1100°C. The structure and thermal stabilities of Ni-P-based ternary coating by RF magnetron sputtering were studied by [2003Wu], [2004Che1] and [2004Che2].
MSIT®
Landolt-Börnstein New Series IV/11C3
Cr–Ni–P
223
References [1948Now] [1962Lun] [1984Ori]
[1986Fje]
[1990Muc]
[1997Yeh]
[2003Wu]
[2004Che1]
[2004Che2]
Nowotny, H., Henglein, E., “Study of Ternary Alloys with Phosphorus” (in German), Monatsh. Chem., 79, 385-393 (1948) (Crys. Structure, Experimental, 18) Lundstroem, T., “A Ternary Sigma Phase in the System Cr-Ni-P”, Acta Chem. Scand., 16(1), 149-154 (1962) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 23) Orishchin, S.V., Kuz’ma, Yu.B., “Cr (W)-Ni-P Ternary Systems”, Inorg. Mater. (Engl. Trans.), 20(3), 360-365 (1984), translated from Izv. Akad. Nauk SSSR, Neorg. Metar., 20(3), 425-430 (1984) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, *, 12) Fjellvag, H., Kjekshus, A., “Solid Solution Phases with MnP Structure: T1–tNitP (T = Ti-Co)”, Acta Chem. Scand., Ser. A, A40, 8-16 (1986) (Crys. Structure, Experimental, 43) Muchnik, S.V., Lomnitskaya, Ya.F., Chernogorenko, V.B., Lynchak, K.A., “Phase Formation under Combustion Conditions in the Ni-Cr-P”, Inorg. Mater. (Engl. Trans.), 26(3), 393-396 (1990), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 26(3), 467-470 (1990) (Experimental, Phase Diagram, Phase Relations, 13) Yeh, M.S., Chuang, T.H., “Transient Liquid Phase Metallic Bonding of an Inconel 718SPF Superalloy”, Welding J., 76(12), S517-S521 (1997) (Interface Phemomena, Experimental, 13) Wu, F.-B., Duh, J.-G., “Mechanical Characterization of Ni-P-based Ternary Coatings by RF Magnetron Sputtering”, Thin Solid Films, 441, 165-171 (2003) (Interface Phemomena, Experimental, 21) Chen, W.-Y., Tien, S.-K., Duh, J.-G., “Thermal Stability and Microstructure Characterization of Sputtered Ni-P and Ni-P-Cr Coatings”, Surf. Coat. Technol., 188-189(2), 489-494 (2004) (Interface Phemomena, Experimental, 21) Chen, W.-Y., Duh, J.-G., “Thermal Stability of Sputtered Ni-P and Ni-P-Cr Coatings During Cycling Test and Annealing Treatment”, Surf. Coat. Technol., 177-178, 222-226 (2004) (Interface Phemomena, Experimental, 19)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(’Cr)
tI2 I4/mmm ’Cr
a = 288.2 c = 288.7
at 25°C, HP [Mas2]
(Cr) < 1863
cI2 Im3m W
a = 288.48
at 25°C [Mas2]
(Ni) < 1455
cF4 Fm3m Cu
a = 352.40
at 25°C [Mas2]
(P) (red) < 417
c*66
a = 1131
sublimation at 1 bar triple point at 576°C, > 36.3 bar; triple point at 589.6 at 1 atm [Mas2] [V-C2]
(P) (white) < 44.14
Landolt-Börnstein New Series IV/11C3
c** ? P (white)
a = 718
at 25°C [Mas2] common form of elemental P, probably less stable than P (red) at 25°C [Mas2] MSIT®
Cr–Ni–P
224 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(P) (black)
oC8 Cmca P (black)
a = 331.36 b = 1047.8 c = 437.63
at 25°C [Mas2]
’, CrNi2 < 590
oP6 Immm MoPt2
-
[Mas2]
(Cr1–xNix)3P < 1510
tI32 I4 Ni3P
a = 918.87 c = 455.93
0.8 x 1 [1984Ori] x = 1 [V-C2]
a = 913.1 c = 453.9
x = 0.8 [1984Ori]
Cr2P(HT) 1640
hP9 P62m Fe2P
-
[Mas2]
Cr2P(LT)
oP18 Pmmm ?
-
[Mas2]
(Cr1–xNix)12P7
hP26 P63/m Cr12P7
a = 898.1 c = 331.3
Cr12P7 (Cr1–xNix)P
a = 877.0 c = 342.7
oP8 Pnma MnP
CrP
a = 516.8 b = 327.8 c = 585.7
0 x 0.16 [1984Ori] x = 0.16 [1984Ori]
x = 0 [V-C2] 0 x 0.7 [1984Ori] x = 0.7 at 27°C [1984Ori] and [1986Fje]
a = 523 b = 334 c = 595
x = 0.7 at 1027°C [1986Fje]
a = 534.6 b = 310.7 c = 599.9
x = 0 [V-C2]
Cr2P3
?
-
[Mas2]
CrP2 1000 - 850
mC12 C2/m Ge2Os
a = 821.3 b = 303.4 c = 709.8 = 119.47°
[V-C2]
CrP4 1200 - 900
mC20 C2/c MoP4
a = 519.14 b = 1076.0 c = 577.12 = 110.65°
[V-C2]
MSIT®
Landolt-Börnstein New Series IV/11C3
Cr–Ni–P
225
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
Ni3P < 970
tI32 I4 Ni3P
a = 881.2 c = 437.3
[V-C2]
Ni5P2 < 1025
hP168 P3 ?
a = 1322.0 c = 2463.2
[V-C2]
Ni5P2 1170 - 1025
-
-
[Mas2]
Ni12P5 < 1025
tI34 I4/m Ni12P5
a = 864.6 c = 507.0
[V-C2]
Ni12P5 1125 - 1000
-
-
[Mas2]
(Cr1–xNix)2P
hP9 or P321 P62m Fe2P or Ni2P
a = 584.0 c = 351.4
Ni2P < 1100
0.5 x 1 [1984Ori] x = 0.5 [1984Ori]
a = 585.9 c = 338.2
x = 1 [V-C2]
Ni5P4
hP36 P63mc Ni5P4
a = 678.9 c = 1098.6
[V-C2]
Ni1.22P ~825 - 770
-
-
[Mas2]
NiP 850
oP16 Pcba NiP
a = 605.0 b = 488.1 c = 689.0
[V-C2]
NiP2
mC12 C2/c NiP2
a = 636.6 b = 561.5 c = 607.2 = 126.22°
[V-C2]
NiP3 700
cI32 Im3 CoAs3
a = 781.92
[V-C2]
* -1, Cr1.2Ni0.8P
oP12 Pnma Co2Si
a = 589.4 b = 349.3 c = 676.2
[1984Ori]
* ), Cr5NiP2 () phase)
tI32 I4 Ni3P
a = 913.1 c = 452.9
[1962Lun] at 1050°C Not confirmed by [1984Ori] at 797°C Probably solution of Ni3P
Landolt-Börnstein New Series IV/11C3
MSIT®
Cr–Ni–P
Fig. 1: Cr-Ni-P. Lattice parameters of the Cr3P phase at 797°C as function of composition at constant P content
c, pm
226
450
440
a, pm
920
910
0
0.10
0.20
0.30
Ni3P, mole fraction
340
c, pm
Fig. 2: Cr-Ni-P. Lattice parameters of the Cr12P7 phase at 797°C as function of composition at constant P content
330
900
a, pm
890
880
0
0.08
0.16
0.24
"Ni12P7", mole fraction
MSIT®
Landolt-Börnstein New Series IV/11C3
Cr–Ni–P
350
c, pm
Fig. 3: Cr-Ni-P. Lattice parameters of the Ni2P phase at 797°C as function of composition at constant P content
227
340
a, pm
590
580
0
0.30
0.15
0.45
Cr2P, mole fraction
605
c, pm
Fig. 4: Cr-Ni-P. Lattice parameters of the CrP phase at 20°C as function of composition at constant P content
595
585
b, pm
330
320
310
a, pm
535
525
515 0
0.20
0.40
0.60
0.80
1.00
"NiP", mole fraction
Landolt-Börnstein New Series IV/11C3
MSIT®
Cr–Ni–P
228
P
Data / Grid: at.% Axes: at.%
Fig. 5: Cr-Ni-P. Partial isothermal section at 797°C 20
80
NiP2 40
60
CrP+NiP2+Ni2P NiP2+Ni5P4+Ni2P
CrP τ 1+CrP+Cr12P7 Cr12P7
CrP+Ni2P
60
Ni5P4 40
τ 1+CrP+Ni2P
Cr2P
τ1
Cr3P
(Ni)+Ni3P+Ni2P
(Ni)+τ 1+Cr3P
80
Ni2P Ni12P5 Ni5P2 Ni3P 20
(Cr)+(Ni)+Cr3P (Cr) 20
Cr
40
60 (Ni)+τ +Ni P 80 1 2
Cr Ni P
Fig. 6: Cr-Ni-P. Partial tentative isothermal section at 1050°C
40.00 0.00 60.00
Ni
(Ni)
Data / Grid: at.% Axes: at.%
60
40
Cr3P 80
(αCr)+Cr3P
20
σ+Cr3P+? σ+Cr3P σ+L σ
Cr
MSIT®
20
σ+(Ni) 40
Cr Ni P
40.00 60.00 0.00
Landolt-Börnstein New Series IV/11C3
Cr–Ni–Si
229
Chromium – Nickel – Silicon Gautam Ghosh Introduction A summary of experimental studies of phase equilibria is given in Table 1. [1960Gua] determined a partial isothermal section at 1050°C, in the composition range of Ni-NiCr-NiSi, by means of metallography and XRD techniques. [1960Gup] determined the homogeneity range of ) phase at 1175°C using XRD technique. Later, [1963Gla1] investigated the phase equilibria of the entire system at 850°C. In addition, selected phase boundaries were also determined at 1050 and 1175°C. They prepared 185 alloys using 99.9% Cr, 99.95% Ni and silane Si (99.9% Si). The alloys were prepared in a high-frequency induction furnace under hydrogen atmosphere, and subsequently heat treated at 850°C for 200 to 1000 h under vacuum. The phase equilibria were determined by metallography and XRD techniques. [1963Gla1] reported the existence of five ternary phases. [1980Cha] calculated partial isothermal sections, in the composition range Cr-Cr3Si-Ni31Si12-Ni, at 427, 527, 627, 800 and 827°C by CALPHAD method. [1993Kod] determined an isothermal section at 900°C employing both diffusion couple data and phase analysis of bulk ternary alloys. They prepared several diffusion couples involving Ni silicides (Ni3Si, Ni5Si2 and Ni2Si), pure Si (99.99 at.%), pure Cr (99.9 at.%) and Cr-Ni alloys. Bulk ternary alloys were prepared by arc melting, and they were subsequently annealed at 900°C under a He atmosphere for 300 to 700 h. The phase relations were established using SEM and EPMA techniques. [1995Cec] reported two isothermal sections at 1050 and 1125°C. Both studies confirmed the existence of four of the five ternary phases reported earlier [1963Gla1]. Very recently, [2000Sch] reported the results of a very comprehensive study of phase equilibria. They prepared 35 ternary alloys using powders of Cr (99.95%), Ni (99.9%) and Si (99.9%). The as-cast alloys were encapsulated in quartz tubes under vacuum and then annealed at 900°C for 25 d. The solid phases were characterized by XRD, while the presence of invariant reactions involving the liquid phase was established by DTA. These experimental results were then used to derive a set of thermodynamic model parameters optimized by CALPHAD method. [2000Sch] summarized their results in terms of two isothermal sections (900 and 1050°C), two isopleths (10 and 40 at.% Cr), the liquidus surface and a reaction scheme. Atomic transport kinetics of Cr and Si in (Ni) was determined by [1982Joh]. Binary Systems The Cr-Ni and Ni-Si binary phase diagrams are accepted from [Mas2], the latter is based on the assessment of [1987Nas]. The Cr-Si binary system is accepted from [2006Leb]. Solid Phases The crystallographic data of the solid phases are listed in Table 2. The solubility of Cr in all Ni-Si intermetallics is less than 5 at.% at 850°C [1963Gla1]. However, at higher temperature the solubility of Cr in (Ni31Si12) and (Ni2Si) are much higher: 9.4 and 15.5 at.% at 900°C [1993Kod], 12.8 and 16.2 at.% at 1050°C [1995Cec], and 14.1 and 23.5 at.% at 1125°C [1995Cec], respectively. [1998Lan] carried out XRD of as-solidified (after levitation melting) (Ni2Si) alloyed with Cr, and their results suggest that (Ni2Si) dissolves about 33.3 at.% Cr. [1990Li] and [1991Zha] investigated the substitution behavior of Cr in 1 (Ni3Si) (cP4 or L12) at 900°C, but their results are rather confusing. [1990Li] reported that Cr atoms reside primarily in the Si-site, while [1991Zha] reported that Cr atoms reside in both Ni- and Si-site. A first-principles calculation of effective pair interaction (EPI) demonstrates that Cr has EPI with Ni and Si of nearly equal strength, suggesting that Cr does not have a particular preference for the sublattice occupancy [1995Slu]. These theoretical results seem to support the experimental results of [1991Zha].
Landolt-Börnstein New Series IV/11C3
MSIT®
230
Cr–Ni–Si
The solubility of Ni in Cr3Si, (Cr5Si3) and CrSi2 is less than 5 at.% at 850°C [1963Gla1]. On the other hand, the solubility of Ni in CrSi is about 30 at.% at 850 [1963Gla1] and 900°C [1993Kod], but it decreases to 17.8 at.% at 1050°C [1995Cec] and 10 at.% at 1125°C [1995Cec]. Furthermore, the dissolution of Ni in CrSi is associated with a decrease in its lattice parameter [1963Gla1, 1962Bur]. Five ternary phases have been reported: ) (Cr13Ni5Si2) [1957Aro, 1960Gup, 1962Gla1, 1962Gla2, 1962Gla3, 1963Gla1, 1993Kod, 1995Cec, 2000Sch], % (Cr3Ni5Si2) [1962Gla2, 1962Gla3, 1963Gla1, 1993Kod, 1995Cec, 2000Sch], -1 (Cr2Ni2Si) [1963Gla1, 1993Kod, 1995Cec, 2000Sch] and -2 (Cr3Ni3Si4) [1963Gla1, 1993Kod, 1995Cec, 2000Sch], T (Cr6Ni16Si7) [1961Gla2, 1963Gla1]. While the ) phase is not stable in binary Cr-Ni alloys, it is stabilized by the addition of about 6 at.% Si in Cr-Ni alloys, particularly at high temperatures [1957Aro, 1960Gup]. At 1175°C, the homogeneity range of ) phase extends from 35.5 to 43.5 at.% Ni and from 58 to 66.4 at.% Cr [1960Gup]. However, the homogeneity range is reduced considerably at lower temperatures [1963Gla1, 1993Kod, 1995Cec]. At 600°C, the ) phase is reported to decompose into (Ni31Si12) and (Cr) [1963Gla1]. [1962Gla3] calculated X-ray diffraction patterns of ) phase, at Cr19.5Ni7.5Si3, assuming two different site occupancy models. In Model-1, it was assumed that the atomic species are randomly distributed in five Wyckoff positions. In Model-2, it was assumed that all Cr and 16.7% of Si atoms reside in 4f, 8i(1) and 8j sites while all Ni and 83% of Si atoms reside in 2a and 8i(2) sites. Both models gave the same level of agreement between calculated and observed X-ray diffraction patterns. The Ni content in % phase lies in the range of 46 to 52 at.% at 800°C [1962Gla3] and 46 to 54 at.% at 1050°C [1995Cec] with an average Si content of 20 at.%. The % phase was not observed at 1125°C [1995Cec], and also it is not stable below 800°C [1963Gla1]. A comparison between calculated and observed X-ray diffraction patterns of % phase suggests a random mixing of Cr and Ni atoms in 12b and 4a sites [1962Gla3]. The composition of -1 phase may be approximated as Cr2Ni2Si, and it does not change appreciably with temperature. [1963Gla1] reported that -1 phase might have been stabilized by carbon; however, the presence of this phase was confirmed in all subsequent investigations [1993Kod, 1995Cec, 2000Sch]. The composition -2 phase may be approximated as Cr3Ni3Si4, but its crystal structure is not known. Even though the Si content in -2 remains around 42 at.%, the Cr content varies significantly with temperature. The following values of Cr content in -2 phase are extracted from the reported isothermal sections: 25.7 to 42.3 at.% at 900°C [1993Kod], 76.3 to 87.7 at.% at 1050°C [1995Cec] and 84.5 to 87.2 at.% at 1125°C [1995Cec]. The T phase was reported by Gladyshevsky et al [1961Gla1, 1961Gla2, 1962Gla2, 1962Gla4, 1963Gla1], but was not confirmed in the subsequent investigations [1993Kod, 1995Cec, 2000Sch]. Therefore, it is not considered to be an equilibrium phase in this assessment. Invariant Equilibria Based on thermal analysis of thirty-five ternary alloys, [2000Sch] established the presence of twenty invariant reactions involving the liquid phase; however, the compositions of participating phases were not reported. Twenty invariant reactions experimentally observed are listed in Table 3. In addition, [2000Sch] also detected two three-phase equilibria, L NiSi+CrSi (e13, max) and L NiSi+CrSi2 (e14, max). The observed invariant reactions are listed in Table 3. Despite the presence of four ternary eutectics, [1978Hao] did not observe any eutectic in this system; however, they did not report the composition range of investigation. The experimental values of the temperature, 1084 2°C [2000Sch] and 1077°C [1979Lug], for the eutectic reaction, L (Ni) + (Ni31Si12) + %, agree well with each other. A reaction scheme based on thermodynamic calculations [2000Sch] and valid up to 800°C is shown in Fig. 1. Here, it is important to note that the temperature of binary invariant reactions are slightly different from those in the accepted binaries as Fig.1 is based on CALPHAD modeling. In Fig.1, twenty-one invariant reactions (thirteen U type, four E type, two P type and two D type) participate in the primary solidification. Four ternary phases (), %, -1 and -2) were considered in thermodynamic calculations, and all are shown to participate in the primary crystallization during solidification. The homogeneity ranges of ) and -2 phases were considered in the thermodynamic modeling, but % and -1 phases were treated as stoichiometric with composition Cr3Ni5Si2 and Cr5Ni5Si3, respectively [2000Sch]. The thermodynamic MSIT®
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calculations can reproduce the temperature of all experimentally observed invariant reactions within 13°C. However, an invariant reaction L + Cr5Si3 Cr5Si3 + Cr3Si was predicted to occur at 1489°C, but it was not experimentally verified. The experimentally determined composition of the liquid phase in invariant reaction L (Ni) + -Ni31Si12 + % is 20 at.% Cr and 21 at.% Si [1979Lug], which agree well with the corresponding values obtained from thermodynamic calculations 19 at.% Cr and 22 at.% Si [2000Sch]. [1979Lug] proposed a partial reaction scheme, with four invariant reactions, for the solidification of Ni rich alloys. The invariant reactions PII, PIII and EIII of [1979Lug] correspond to D1, U2, E2, respectively, of [2000Sch]. Another invariant reaction PIV (L + ) (Ni) + %) postulated by [1979Lug] was not predicted in thermodynamic calculations. Liquidus Surface The liquidus temperature of several ternary alloys has been determined by DTA [1979Lug, 1990Kag, 1995Cec, 2000Sch]. The liquidus surface shown in Fig. 2 is based on thermodynamic calculations of [2000Sch]. In an earlier study, [1979Lug] reported the liquidus surface of Ni corner only showing the presence of two invariant reactions. The calculated liquidus surface in Fig. 2 shows melting grooves separating eighteen different areas of primary crystallization, including four ternary phases. The melting groove defining the composition range for the primary crystallization of % phase is very narrow. Selected liquidus isotherms are shown in Fig. 2. Isothermal Sections Figure 3 shows the isothermal section of Cr corner at 1175°C [1960Gup]. This study clearly established the homogeneity range of ) phase. Even though the adjoining two- and three-phase boundaries were not determined, the reported phase fields are consistent with the reaction scheme shown in Fig. 1. At 1175°C, the three-phase fields L+(Ni)+) and (Cr)+(Ni)+) originate from the invariant reaction U2, while L+)+Cr3Si and (Cr)+)+Cr3Si originate from the invariant reaction U3. Figures 4 and 5 show the isothermal section at 1125 and 1050°C [1995Cec], respectively. [1960Gua] also reported an isothermal section at 1050°C, but only in the Ni corner. The phase relations in [1960Gua] differ significantly from [1995Cec] due to following reasons: (i) [1960Gua] observed only one ternary phase -1, and (ii) [1960Gua] reported negligible solubilities of Cr in (Ni31Si12) and (Ni2Si), but they are found to be significant [1995Cec]. The % phase was not observed at 1125°C [1995Cec]. Thermodynamic calculations show that the % phase forms via a peritectic reaction (P2 in Fig. 1) at 1108°C [2000Sch], and it is not stable below 843°C due to the eutectoid reaction E5 (in Fig. 1). The latter temperature is in reasonably good agreement with an earlier report that % phase is not stable below 800°C [1963Gla1]. Similarly, thermodynamic calculations indicate that the ) phase is not stable below 525°C [2000Sch], while an experimental investigation found that it decomposes into (Ni31Si12) and (Cr) at 600°C. Figure 6 shows the isothermal section at 900°C [1993Kod], while Fig. 7 shows the isothermal section at 850°C [1963Gla1]. The isothermal section at 850°C reported by [1963Gla1] suffers from the following drawbacks: (i) the reported ternary phase T was not confirmed in any subsequent investigations, (ii) they did not establish the phase relations involving T and -1, and (iii) the reported three phase field %+)+(Ni) is inconsistent with the thermodynamic calculations. Due to these reasons, the following amendments are made in Fig. 7: (i) the T phase is not shown, and (ii) the phase fields involving %, -1 and ) are redrawn to make them consistent with the expected phase relations based on the reaction scheme in Fig. 1. For example, the three-phase field %+)+(Ni) reported by [1963Gla1] has been replaced by two three-phase fields, -1+)+(Ni) and %+-1+(Ni). The amended two- and three-phase fields involving %, -1 and ) are shown dashed in Fig. 7. [1980Cha] reported partial isothermal sections at 427, 527, 627 and 827°C based on thermodynamic calculations where only one ternary phase, ), was considered. The presence of a three phase (Cr)+(Ni)+) is predicted at all temperatures. In Fig. 3 to 7, adjustments are made along the binary edges to comply with the corresponding phase diagrams accepted in this assessment. Landolt-Börnstein New Series IV/11C3
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Temperature – Composition Sections Figures 8a and 9a show calculated isopleths along Cr0.1Ni0.9-Cr0.1Si0.9 and Cr0.4Ni0.6-Cr0.4Si0.6 section, respectively, based on thermodynamic calculations [2000Sch]. These authors also determined liquidus and solidus of several ternary alloys using DTA, and these results were used in the optimization of thermodynamic model parameters. To delineate the phase boundaries clearly in certain composition ranges, enlarged parts are shown in Figs. 8b and 9b. Thermodynamics [1977Ost] determined the enthalpy of solution of Si (up to 0.4 at.%) in a Ni-30Cr (at.%) alloy in the ∞ (in Ni-30Cr temperature range of 1550 to 1600°C by means of calorimetry. The experimental value of ΔH Si –1 (at.%)) is –102.1 ± 5.9 kJ#atom . [1968Che] determined the activity of Si (up to 2 mass%) in liquid Cr-Ni alloys at 1600°C by the emf method. Their results show that Cr has a strong tendency to increase the activity of Si. Thermodynamic modeling of the ternary system by [1980Cha] was restricted to the composition range of Cr-Cr3Si-Ni31Si12-Ni. More recently, [2000Sch] carried out a comprehensive thermodynamic assessment, and reported a set of self-consistent thermodynamic model parameters. Notes on Materials Properties and Applications A summary of experimental investigation of properties is given in Table 4. Both Cr-Ni and Cr-Ni-Si alloys are potential candidates for joining of Si3N4 ceramic parts by brazing and solid state bonding [1990Nak, 1991McD, 1993Kod, 1995Cec, 1996Tun, 2001Lin]. Consequently, these studies have investigated various physico-chemical and mechanical properties of Si3N4/Cr-Ni-Si system, such the contact angle [1995Cec], interfacial microstructure [1993Kod, 1995Cec, 1996Tun, 2001Lin], and the joint strength [1990Nak]. Mechanical, oxidation and electrochemical properties of three-phase alloys (Ni)+1(Ni3Si)+ (Ni31Si12) were reported by [2002Tak]. The high-temperature mechanical properties and oxidation resistance of these alloys are better then 1 (Ni3Si) while the corrosion resistance in sulphuric acid is comparable to 1 (Ni3Si). The wear resistance of Cr13Ni5Si2 () phase) was evaluated under dry sliding condition at room temperature [2004Tan]. The results show it has excellent wear resistance under dry sliding condition due to high hardness and strong atomic bonds. Also, due to same reasons, a two phase microstructure, Cr3Si+Cr13Ni5Si2, also exhibits good high temperature sliding wear resistance at 400, 500 and 600°C [2004Zha2]. When used as a coating on titanium alloys, Cr13Ni5Si2 based alloys exhibit excellent wear resistance under dry-sliding wear test conditions [2005Jia1, 2005Jia2]. The laser cladding of Cr-alloyed nickel silicide coating (Ni-22.6 mass% Si-5.6 mass% Cr) on 0.2% C steel has a rapidly solidified microstructure consisting of the Ni2Si primary cellular-dendrites and minor amount of interdendritic Ni2Si/NiSi eutectics [2003Cai]. The intermetallic coating has good wear resistance under dry sliding wear test conditions due to the high hardness, refined microstructure and strong intermetallic atomic bonds [2003Cai]. Alloying of Ni2Si/NiSi coatings (on 0.2% C steel) with 1.5% Cr and 5.6% Cr considerably enhanced the corrosion resistance under both anodic polarization and immersion corrosion test conditions [2004Cai]. [2001Don] and [2004Zha1] studied the electrical properties of amorphous thin films of Cr17Si80Ni3, on glass and n type Si (100) surface, in the temperature range 250 to 750°C. Crystallization leads to the formation of CrSi2 phase, and with its increasing fraction leads to the decrease in conductivity of the films. A proper mixture of amorphous and crystalline phases could result in a final zero temperature coefficient of resistance. Miscellaneous Atomic transport kinetics of Cr and Si in solid solutions has been studied by [1982Joh] at 1250°C. Eight diffusion couples, made of either Ni-Cr/Ni-Si or Ni-Cr-Si/Ni or Ni-Cr-Si/Ni-Si alloys, were prepared and then annealed at 1250°C for up to 180.5 h, and then the composition profiles were measured by EPMA.
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Ni Using the experimental composition profiles, [1982Joh] evaluated four interdiffusion coefficients, D CrCr , Ni Ni Ni D CrSi , D SiCr and D SiSi , in the composition range of 2.1 to 13.5 at.% Cr and 0.79 to 4.04 at.% Si. Their Ni Ni , follow a dilute solution model where results show that the indirect diffusion coefficients, D CrSi and D SiCr both are proportional to xCr and xSi, respectively, and approach zero as the respective concentrations Ni , was found to be consistent with the diffusion approach zero. The direct diffusion coefficient, D SiSi coefficient of Si in binary Ni-Si alloys, both in magnitude and Si concentration dependence. In contrast, Ni was found to be greater than the diffusion coefficient of Cr in binary Cr-Ni alloys and also a function D CrCr of Si concentration. The diffusivity data of [1982Joh] data was used to derive a set of optimized mobility parameters of Cr, Ni and Si in fcc solid solution within CALPHAD formalism [2001Du]. In a Ni-3.2Cr-4.91Si (mass%) alloy, the Cr and Si partitioning ratios (concentration in solid/ concentration in solid) at 1336°C were reported to be 1.13 and 0.59, respectively. Solid state reactions in Cr/Ni/Si thin films have been studied using a variety of techniques [1982Nau, 1984App, 1993Pim]. These results were reviewed by [1995Set]. The reaction products in thin film multilayers, though only binary silicides, depend primarily on the temperature. For example, [1984App] reported that Cr/Ni/Si layers react independently with Si to form the following products after 30 min anneal: Cr/Ni2Si/Si at 300°C, Cr/NiSi/Si at 400°C, CrSi2/NiSi/Si at 500°C, CrSi2/NiSi2/Si at 850°C. The formation of an amorphous NiSi around 320°C has also been reported [1993Pim]. No ternary phase was observed in any of these studies. Rapid solidification of Cr5Ni3Si2 by piston-and-anvil method leads to the formation of octagonal quasicrystals, which is closely related to the Mn structure [1987Wan, 1990Wan].
References [1957Aro] [1960Gua] [1960Gup]
[1961Gla1]
[1961Gla2]
[1962Bur]
[1962Gla1]
[1962Gla2]
[1962Gla3]
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Aronsson, A., Lundstrom, T., “Investigations on )-FeCrSi”, Acta Chem. Scand., 11, 365-372 (1957) (Crys. Structure, Experimental, 31) Guard, R.W., Smith, E.A., “Nickel-Chromium-Silicon System”, J. Inst. Met., 88, 373-374 (1960) (Experimental, Phase Diagram, Phase Relations, #, *, 4) 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) (Experimental, Phase Diagram, Phase Relations, #, *, 18) Gladyshevsky, E.I., Kripyakevich, P.I., Kuzma, Yu.B., “New Substitutes of the Structural Types of Mg6Cu16Si7 and Th6Mn23” (in Russian), Sov. Powder Metall. Met. Ceram. (Engl. Transl.), 769-770 (1961) (Crys. Structure, Experimental, 7) Gladyshevsky, E.I., Kripyakevich, P.I., Kuzma, Yu.B., Teslyuk, M.Yu., “New Representatives of Structures of Mg6Cu16Si7 and Th6Mn23”, Sov. Phys.-Crystallogr. (Engl. Transl.), 6, 615-616 (1961), translated from Kristallografiya, 6, 1961, 769-770 (Crys. Structure, Experimental, 7) Burger, K.O., Wittmann, A., Nowotny, H., “The Crystal Structure of Co3Al3Si4 and Co2AlSi2 and the Constitution of some Monosilicide Systemes of Transition Metals” (in German), Monatsh. Chem., 93, 9-14 (1962) (Crys. Structure, Experimental, 8) Gladyshevsky, E.I., Kripyakevich, P.I., “Intermetallic Compounds with a -U () phase) Type Structure” (in Russian), Vysokotemperaturnye Metallokeram. Materialy, 148-150 (1962) (Experimental, Crys. Structure, 8) Gladyshevsky, E.I., “Crystal Structure of Compounds and Phase Equilibria in Ternary Systems of two Transition Metals and Silicon”, Sov. Powder Metall. Met. Ceram. (Engl. Transl.), (4), 262-265 (1962) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 17) Gladyshevsky, E.I., Kripyakevich, P.I., Kuzma, Yu.B. “Crystal Structure of Ternary Compounds with Low Silicon Content in the Cr-Ni-Si and Cr-Co-Si Systems.”, J. Struct. Chem. (Engl. Transl.), 3(4), 402-410 (1962), translated from Zh. Strukt. Khim., 3(4), 414-423 (1962) (Calculation, Crys. Structure, Experimental, 29)
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[1963Gla1]
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Cr–Ni–Si Gladyshevsky, E.I., Markiv, V.Ya., Kuzma, Yu.B., “New Ternary Compounds with a Structure of the Mg6Cu16Si7 Type”, Dop. Akad. Nauk Ukr. RSR, (4), 481-483 (1962) (Crys. Structure, Experimental, 5) Gladyshevsky E.I., Borusevich, L.K., “The Ternary System Cr-Ni-Si.”, Russ. J. Inorg. Chem. (Engl. Transl.), 8(8), 997-1000 (1963), translated from Zhur. Neorg. Khim., 8(8), 1915 (1963) (Calculation, Crys. Structure, Experimental, Phase Diagram, Phase Relations, #, *, 18) Gladyshevsky, E.I., Kuzma, Yu.B., Kripyakevich, P.I., “Crystal Structure of Mn3Ni2Si, V3Ni2Si, Nb3Ni2Si, and of Cr and Ta Compounds of Similar Structure.”, J. Struct. Chem. (Engl. Transl.), 4(3), 343-349 (1963), translated from Zh. Strukt. Khim., 4, 327-379 (1963) (Crys. Structure, Morphology, Experimental, 19) Cherkasov, P.A., Averin, V.V., Samarin, A.M., “Activities of Silicon and Titanium in Molten Iron, Cobalt, and Nickel Containing Chromium”, Russ. J. Phys. Chem. (Engl. Transl.), 42(3), 401-404 (1968), translated from Zh. Fiz. Khim, 42(3), 767 (1968) (Experimental, Thermodyn., *, 19) Ostrovskii O.I., Stomakhin A. Ya, Dietrich E., Grigoryan V.A., “Heats of Solution of Aluminium, Silicon and Titanium in Iron-Chromium And Nickel-Chromium Melts” (in Russian), Vses. Konf. Kalorim., (Rasshir. Tezisy Dokl.), 59-63 (1977) (Experimental, Thermodyn., *, 11) Haour, G., Mollard, F., Lux, B., Wright I.G., “New Eutectics Based on Fe, Co and Ni. III Results Obtained for Ni-Base Alloys”, Z. Metallkd., 69(3), 149-154 (1978) (Experimental, Kinetics, Phase Diagram, Phase Relations, 14) Lugscheider, E., Knotek, O., Kloehn, K., “Melting Behaviour of Nickel-Chromium-Silicon Alloys”, Thermochim. Acta, 29, 323-326 (1979) (Experimental, Phase Diagram, Phase Relations, #, *, 6) Chart, T., Putland, F., Dinsdale, A., “Calculated Phase Equilibria for the Cr-Fe-Ni-Si System - I Ternary Equilibria”, Calphad, 4(1), 27-46 (1980) (Calculation, Phase Diagram, Phase Relations, *, 75) Johnston, G.R., “Diffusion of Chromium and Silicon in Nickel Solid-Solution Alloys of the Ni-Cr-Si System”, High Temp.-High Pressures, 14(6), 695-708 (1982) (Phase Relations, Experimental, Kinetics, *, 34) Naudé, M.O., Pretorius, R., Marais, D.J., “Bilayer Silicide Formation During the Interaction of thin Chromium, Nickel and Platinum with Silicon”, Thin Solid Films, 89(4), 339-348 (1982) (Crys. Structure, Experimental, Kinetics, 30) Appelbaum, A., Eizenberg, M., Brener, R., “Phase Separation on Layer Sequence Reversal During Silicide Formation With Ni-Cr Alloys and Ni-Cr Bilayers”, J. Appl. Phys., 55(4), 915-919 (1984) (Crys. Structure, Experimental, Kinetics, 17) Nash, P., Nash, A., “The Ni-Si (Nickel-Silicon) System”, Bull. Alloy Phase Diagrams, 8(1), 6-14 (1987) (Phase Diagram, Phase Relations, Crys. Structure, Review, Experimental, #, *, 59) Wang, N., Chen, H., Kuo, K.H., “Two-Dimensional Quasicrystal with Eightfold Rotational Symmetry”, Phys. Rev. Lett., 59(9), 1010-1013 (1987) (Crys. Structure, Review, Experimental, 20) Kagawa, A., Hirata, M., Sakamoto, Y., “Solute Partitioning on Solidification of Nickel-base Ternary Alloys”, J. Mater. Sci., 25(12), 5063-5069 (1990) (Experimental, Phase Diagram, Phase Relations, *, 14) Li, Y., Zhang, Z., Zheng, Z., Zhu, Y., “Solution Behaviour of Various Alloying Elements in Ni3Si”, Acta Metall. Sin. (China), 26(3), A172-A176 (1990) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 9) Nakamura, M., Peteves, S.D., “Solid-State Bonding of Silicon Nitride Ceramics with Nickel-Chromium Alloy Interlayers”, J. Am. Ceram. Soc., 73(5), 1221-1227 (1990) (Crys. Structure, Review, 29) Landolt-Börnstein New Series IV/11C3
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[1991McD]
[1991Zha]
[1993Kod]
[1993Pim]
[1995Cec]
[1995Set] [1995Slu]
[1996Tun]
[1998Lan]
[2000Sch]
[2001Don]
[2001Du]
[2001Lin]
[2002Tak]
[2003Cai]
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Wang, N., Kuo, K.H., “Transformation of the Octagonal Quasicrystal into the beta-Mn type Crystalline Structure”, Philos. Mag. Lett., 61(2), 63-68 (1990) (Crys. Structure, Experimental, 9) McDermid, J.R., Drew, R., “Thermodynamic Brazing Alloy Design for Joining Silicon-Carbide”, J. Am. Ceram. Soc., 74(8), 1855-1860 (1991) (Thermodyn., Interface Phenomena, 30) Zhang, T., Li, Y., Zheng, Z., Zhu, Y., “Alloying Behavior of Ni3Si and the 900°C Isotherms of Several Ni-Si-X Systems at Ni rich Corner”, Mater. Res. Soc. Symp. Proc.: High-Temp. Ordered Intermetallic Alloys IV, 213, 137-142 (1991) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, *, 7) Kodentsov, A.A., Gülpen, J.H., Schiepers, R.C.J., Kivilahti, J.K., van Loo, F.J.J., “Reaction Products at the Interface between Si3N4 and Ni-Cr Alloys”, Mater. Sci. Forum, 126-128, 289-292 (1993) (Experimental, Phase Diagram, Phase Relations, Interface Phenomena, #, *, 7) Pimentel, F., Zalar, A., Hofmann, S., Kohl, D., Panjan, P., “A Study of Thermally Activated Interfacial Reactions in a Ni/Cr/Si Multilayer Structure”, Thin Solid Films, 228(1-2), 149-153 (1993) (Crys. Structure, Experimental, Kinetics, 22) Ceccone, G., Nicholas, M.G., Peteves, S.D., Kodentsov, A.A., Kivilahti, J.K., van Loo, F.J.J., “The Brazing of Si3N4 with Ni-Cr-Si Alloys”, J. Eur. Ceram. Soc., 15(6), 563-572 (1995) (Experimental, Phase Diagram, Phase Relations, Interface Phenomena, #, *, 27) Setton M., “Ternary TM-TM-Si Reactions”, EMIS Datarev. Ser. 14, Properties of Metal Silicides, 14, 129-149 (1995) (Crys. Structure, Review, 78) Sluiter, M., Kawazone, Y., “Site Preference of Ternary Additions in Ni3Si”, High-Temperature Ordered Intermetallic Alloys VI, Mater. Res. Soc. Symp. Proc., 364(PT.2), 1064-1069 (1995) (Phase Relations, Crys. Structure, Experimental, Electronic Structure, *, 13) Tung, S.K., Lim, L.C., Lai, M.O., “Solidification Phenomena in Nickel Base Brazes Containing Boron and Silicon”, Scr. Mater., 34(5), 763-769 (1996) (Assessment, Phase Relations, 16) Landrum, G.A., Hoffmann, R., Evers, J., Boysen, H., “The TiNiSi Family of Compounds: Structure and Bonding”, Inorg. Chem., 37(22), 5754-5763 (1998) (Crys. Structure, Experimental, 34) Schuster, J.C., Du, Y., “Experimental Investigation and Thermodynamic Modeling of the Cr-Ni-Si System”, Metall. Mater. Trans. A, 31A, 1795-1803 (2000) (Assessment, Experimental, Phase Relations, Thermodyn., #, *, 37) Dong, X.P., Wu, J.S., “Study on the Crystallization of Amorphous Cr-Si-Ni Thin Films Using in-situ X-ray Diffraction”, J. Mater. Sci. Technol., 17, S43-S46 Suppl. 1 (2001) (Crys. Structure, Experimental, Electr. Prop., 10) Du, Y., Schuster, J.C., “Assessment of Diffusional Mobilities of Cr, Ni, and Si in fcc Cr-Ni-Si Alloys”, Z. Metallkd., 92(1) 28-31 (2001) (Assessment, Crys. Structure, Kinetics, 21) Lin, D.Y., Wu, W., Lin, C.H., Hsin-Hsin-Hsieh, “The Effect of Aging on the Intergranular Corrosion of a 24Cr-14Ni-0.7Si Stainless Steel for Welding in Architecture”, Steel Res., 72(7), 277-280 (2001) (Morphology, Experimental, Interface Phenomena, 7) Takasugi, T., Kawai, H., Kaneno, Y., “The Effect of Cr Addition on Mechanical and Chemical Properties of Ni3Si Alloys”, Mater. Sci. Eng. A, 329-331, 446-454 (2002) (Experimental, Interface Phenomena, Mechan. Prop., Phase Relations, 15) Cai, L.X., Wang, C.M., Wang, H.M., “Laser Cladding for Wear-resistant Cr-Alloyed Ni2Si-NiSi Intermetallic Composite Coatings”, Mater. Lett., 57(19), 2914-2918 (2003) (Crys. Structure, Experimental, Mechan. Prop., 8)
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[2004Tan]
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[2006Leb]
Cai, L.X., Wang, H.M., Wang, C.M., “Corrosion Resistance of Laser Clad Cr-alloyed Ni2Si/NiSi Intermetallic Coatings”, Surf. Coat. Technol., 182(2-3), 294-299 (2004) (Crys. Structure, Experimental, Interface Phenomena, Phys. Prop., 29) Tang, H.B., Fang, Y.L., Wang, H.M., “Microstructure and Dry Sliding Wear Resistance of a Cr13Ni5Si2 Ternary Metal Silicide Alloy”, Acta Mater., 52(7), 1773-1783 (2004) (Crys. Structure, Experimental, Mechan. Prop., 30) Zhang, Y.Q., Dong, X.P., Wu, H.S., “Microstructure and Electrical Characteristics of Cr-Si-Ni Films Deposited on Glass and Si (100) Substrates by RF Magnetron Sputtering”, Mater. Sci. Eng. B-Solid State Materials for Advanced Technology, 113(2), 154-160 (2004) (Experimental, Electr. Prop., 15) Zhang, L.Y., Wang, H.M., “Microstructure and High-temperature Sliding Wear Resistance of a Cr3Si/Cr13Ni5Si2 Intermetallic Alloy”, Rare Met. Mater. Eng., 33(5), 512-514 (2004) (Experimental, Phase Relations, Mechan. Prop., 6) Jian, L., Zhang, L.Y., Yu, R.L., Wang, H.M., “Microstructure and Wear Resistance of Laser Clad Cr13Ni5Si2 based Metal Silicide Coating on Titanium Alloy”, Rare Met. Mater. Eng., 34(6), 936-939 (2005) (Experimental, Mechan. Prop., 5) Jian, L., Wang, H.M., “Microstructure and Wear Behaviours of Laser-clad Cr13Ni5Si2 based Metal-Silicide Coatings on a Titanium Alloy”, Surf. Coat. Technol., 192(2-3), 305-310 (2005) (Morphology, Experimental, Mechan. Prop., 27) Lebrun, N., Dobatkina, T.V., Kuznetsov, V.N., Li, Changrong, “Cu-Si (Copper-Silicon)”, MSIT Binary Evoluation Program, in MSIT Workplace, Effenberg, G., (Ed.), MSI, Materials Science International Service GmbH, Stuttgart, to be published (2006) (Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., Assessment, 30)
Table 1: Investigations of the Cr-Ni-Si Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique Temperature/Composition/Phase Range Studied
[1960Gua]
Metallography and XRD
900-1050°C; 20-50 mass% Cr, 30 mass% Si, bal. Ni
[1960Gup]
Metallography and XRD
1175°C; 50-80 at.% Cr, 8-20 at.% Si, bal. Ni
[1961Gla1, 1961Gla2]
XRD
Cr6Ni16Si7
[1962Gla1]
XRD
Cr13Ni5Si2
[1962Gla2]
XRD
Cr3Ni5Si2
[1962Gla3]
Metallography and XRD
800-1100°C; 5-30 at.% Ni, 5-25 at.% Si, bal. Cr
[1963Gla1]
Metallography and XRD
850°C; entire composition range
[1963Gla2]
XRD
Cr3Ni2Si
[1968Che]
EMF
1600°C; 20 mass% Cr, 2 mass% Si, bal. Ni
[1977Ost]
EMF
1600°C; Cr: 0.1-1.3 mass%, Ni: 0.35-1.45 mass%, Cu: bal.
[1979Lug]
Thermal analysis
Up to 27 at.% Cr and up to 21 at.% Si
[1980Cha]
Thermodynamic modeling
427-827°C; up to 20 at.% Si
[1987Wan, 1990Wan]
TEM and HREM
Cr5Ni3Si2
MSIT®
Landolt-Börnstein New Series IV/11C3
Cr–Ni–Si
237
Reference
Method/Experimental Technique Temperature/Composition/Phase Range Studied
[1993Kod]
Diffusion couple, thermodynamic 900°C; entire composition range modelling
[1995Cec]
-
1050°C and 1125°C; entire composition range
[2000Sch]
DTA, XRD and thermodynamic modelling
900-1200°C; entire composition range
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Cr) < 1863
cI2 Im3m W
a = 288.48
pure Cr at 25°C [Mas2]
(Ni) < 1455
cF4 Fm3m Cu
a = 352.40
pure Ni at 25°C [Mas2]
(Si) < 1414
cF8 Fd3m C (diamond)
a = 543.06
pure Si at 25°C [Mas2]
CrNi2
oP6 Immm MoPt2
Cr3Si < 1780
cP8 Pm3m Cr3Si
[V-C2], 60 to 76.5 at.% Ni [V-C2] a = 252.4 b = 757.1 c = 356.8 a = 456.27 0.04
20.8 - 25.3 at.% Si [2006Leb] at 22.5 0.4 at.% Si [2006Leb] at 1200°C
a = 456.67 0.02
at 20.8 0.4 at.% Si [2006Leb] at 1600°C 37.5 - 37.7 at.% Si [2006Leb]
Cr5Si3 1666 - 1488
-
-
Cr5Si3 < 1519
tI32 I4mcm W5Si3
a = 917.0 c = 463.6
(Cr1–xNix)Si < 1424
cF8 P213 FeSi
a = 462.2 0.01
[V-C2], 0 x 0.3 [1962Bur] at CrSi
a = 459.24
at Cr40Ni10Si50 [1962Bur]
a = 458.64
at Cr37.7Ni12.3Si50 [1962Bur]
a = 456.92
at Cr30Ni20Si50 [1962Bur]
Landolt-Börnstein New Series IV/11C3
[V-C2] at 37.5 at.% Si
MSIT®
Cr–Ni–Si
238 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
CrSi2 < 1439
hP9 P6422 CrSi2
a = 442.83 0.01 c = 636.80 0.09
1, Ni3Si < 1035
cP4 Pm3m AuCu3
a = 350.6
2, Ni3Si 1115 - 990
mC16 C2/m GePt3(?)
3, Ni3Si 1170 - 1115
mC16 C2/m GePt3(?)
, Ni31Si12 < 1242
hP43 P321 Ni31Si12
a = 667.1 c = 1228
, Ni2Si 1306 - 825
hP6 P6322 Ni2Si
a = 383.6 to 380.2 37.5 - 43 at.% Si [V-C2] c = 494.8 to 486.3
, Ni2Si < 1255
oP12 Pnma Co2Si
J, Ni3Si2 < 830
oC80 Cmc21 Ni3Si2
(Ni1–xCrx)Si < 992
oP8 Pnma MnP
Lattice Parameters Comments/References [pm]
a = 697 2 b = 625 4 c = 507 8 = 48.74° a = 704 7 b = 626 4 c = 508 4 = 48.74°
[V-C2], 22.8 - 24.5 at.% Si at Ni3.04Si0.96 [Mas2], 24.5 - 25.5 at.% Si [1987Nas]
[Mas2], 24.5 - 25.5 at.% Si [1987Nas]
[V-C2]
a = 383.6 c = 494.8
at Ni1.86Si1.14 [V-C2]
a = 502.2 b = 374.1 c = 708.8
at 33.3 at.% Si [V-C2]
a = 566.72 b = 361.31 c = 688.72
[1998Lan], at equiatomic CrNiSi
a = 1222.9 b = 1080.5 c = 692.4 a = 519.0 b = 333.0 c = 562.8 a = 518.0 b = 336.0 c = 561.0
MSIT®
[V-C2], 66.3 - 68 at.% Si at 66.7 at.% Si
39.2 - 41 at.% Si [Mas2] at 33.3 at.% Si [V-C2]
0 x 0.05 [1962Bur] at NiSi [V-C2]
at Ni48Cr2Si50 [1962Bur]
Landolt-Börnstein New Series IV/11C3
Cr–Ni–Si
239
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
NiSi2 993 - 981
-
-
[Mas2]
NiSi2 < 981
cF12 Fm3m CaF2
a = 540.6
[V-C2]
* ), Cr13Ni5Si2
tP30 P42/mnm CrFe
a = 878.67 c = 457.02
[1962Gla1, 1962Gla2, 1962Gla3]
* %, Cr3Ni5Si2
cP20 P213 AlAu4
* -1, Cr3Ni2Si
a = 611.83
cF96 Fd3m NiTi2
a = 1062.0
* -2, Cr3Ni3Si4
-
-
T, Cr6Ni16Si7
cF116 Fm3m Mn23Th6 or Mg6Cu16Si7
[1962Gla2, 1962Gla3], 46 - 52 at.% Ni and 18 - 22 at.% Si at 800°C at Cr3Ni5Si2 [1963Gla2]
[1963Gla1]
a = 1111.0
[1961Gla1, 1961Gla2, 1962Gla2, 1963Gla1] not accepted in this assessment as a stable ternary phase
Table 3: Experimentally Observed Invariant Equilibria Reaction
T [°C]
Type
L + (Cr) Cr3Si + )
1294 2
U
L + (Cr) (Ni) + )
1294
U
L + Cr3Si (Ni2Si) + Cr5Si2
1210 2
U
L + CrSi + Cr5Si2 -2
1174 2
P
L Ni5Si2 + (Ni2Si)+Cr5Si2
1141 2
E
L + Cr3Si -1 + )
1140 2
U
L + ) (Ni) + -1
1122 4
U
L + Cr3Si (Ni31Si12) + )
1121 2
U
L + (Ni) + -1 %
1105 2
P
L + -1 (Ni31Si12) + %
1095 2
U
L (Ni) + (Ni31Si12) + %
1084 2
E
L + Cr5Si2 (Ni2Si) + -2
1063 2
U
L + (Ni2Si) (Ni2Si) + -2
964
U
L + (Si) CrSi2 + NiSi2
962 2
U
Landolt-Börnstein New Series IV/11C3
MSIT®
Cr–Ni–Si
240 Reaction
T [°C]
Type
L CrSi2+ NiSi + NiSi2
946 2
E
L + -2 (Ni2Si) + CrSi
944 2
U
L CrSi + NiSi, CrSi2
943 2
D
L (Ni2Si) + NiSi + CrSi
927 2
E
Table 4: Investigations of the Cr-Ni-Si Materials Properties Reference
Method/Experimental Technique Temperature/Composition/Phase Range Studied
[2001Don, 2004Zha1]
XRD, resistivity
250-750°C; Cr17Si80Ni3
[2003Cai]
Wear test
RT; Ni-22.6Si-5.6Cr (mass%) coatings on 0.2% C steel
[2004Cai]
Corrosion test
RT; Ni2Si/NiSi coatings (on 0.2% C steel) containing 1.5% and 5.6% Cr
[2004Tan]
Wear test
RT; Cr13Ni5Si2
[2004Zha1] XRD
Cr13Ni5Si2
[2004Zha2] Wear test
400-600°C; Cr3Si+Cr13Ni5Si2
[2005Jia1, 2005Jia2]
RT; Cr13Ni5Si2 based coatings on Ti base alloys
MSIT®
Wear test
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Cr-Ni
Ni-Si
Cr-Ni-Si
Cr-Si 1701 e1 L (Cr) + Cr3Si 1664 e2 L Cr3Si + βCr5Si3 1519 p1 L+βCr5Si3 αCr5Si3
1489 L+βCr5Si3αCr5Si3+Cr3Si U1
1488
βCr5Si3+αCr5Si3+Cr3Si
e3
βCr5Si3αCr5Si3+Cr3Si
1424 p2 L + αCr5Si3 CrSi 1408 e4 L CrSi + CrSi2
L+αCr5Si3+Cr3Si
L + (Cr) (Ni) + σ
1299
L+Cr3Si+σ
1240 e7 Lγ+δ 1205
(Cr)+Cr3Si+σ
L+ Cr3Si δ + αCr5Si3
L+αCr5Si3+δ 1170 U13 D1
U7 U5 E1 U10
Cr3Si+δ+ αCr5Si3
L+ CrSi + αCr5Si3 τ2
L+αCr5Si3+τ2
U4
CrSi+αCr5Si3+τ2
P1 L+CrSi+τ2 U14
D2
e6
241
MSIT®
Fig. 1a: Cr-Ni-Si. Reaction scheme, part 1
U3
1207 e8(max) L Cr3Si + σ
1199 p5 L + γ β3 1151 e9 L (Ni) + β3
L + (Cr) Cr3Si + σ
1297 L+δ+θ
U2
(Cr)+(Ni)+σ
L+(Ni)+σ 1251 p4 L+θδ
1328 e6 L CrSi2 + (Si)
1314 p3(max) L + (Cr) σ
Cr–Ni–Si
1345 e5 L (Cr) + (Ni)
Ni-Si
Cr-Ni-Si p4 p5 U e9 U U 4 2 3
Cr-Si
e7 e8
e4
e6
1141 e10(max) L Cr3Si + γ
1142 p6(max) L + Cr3Si τ1
L Cr3Si + γ + δ
1138 L + Cr3Si τ1 + δ
1132
P1
242
MSIT®
Cr-Ni
E1
Cr3Si+γ+δ
U5
Cr3Si+τ1+δ 1127 d1-d2 β3 β2, (Ni), γ
L + β3 β2, (Ni), γ
1127
1126
L + Cr3Si γ + τ1
D1 L+τ1+δ
U6
L+β2+(Ni)
Cr3Si+γ+τ1 L + σ (Ni) + τ1 σ+(Ni)+τ1
1108
U7
L+(Ni)+τ1
L + β2 (Ni) + γ
1119
Cr–Ni–Si
1126 L+γ+τ1
L+β2+γ
U8
L + (Ni) + τ1 π
P2
β2+(Ni)+γ
L+τ1+π L+(Ni)+π L+(Ni)+γ
L+γ+π L (Ni) + π + γ
1082
1076
1042 p7 (Ni) + β2 β1 Landolt-Börnstein New Series IV/11C3
1015 e11 β2 β1 + γ
L + τ1 γ + π
1086
Fig. 1b: Cr-Ni-Si. Reaction scheme, part 2
β2+β1+γ
U13
(Ni)+π+γ U10
αCr5Si3+δ+τ2
L+δ+τ2 U13
τ1+γ+π
E2
L + αCr5Si3 δ + τ2
U9
1018
(Ni) + β2 β1 + γ (Ni)+β1+γ
U11
E5
D2 U14 U12
Landolt-Börnstein New Series IV/11C3
Cr-Ni
Ni-Si 970 p8 L + (Si) NiSi2 966 e12 L θ + NiSi
Cr-Ni-Si U10
e6
e4
L + (Si) CrSi2 + NiSi2 U12 (Si)+CrSi2+NiSi2
L+CrSi2+NiSi2
L + δ θ + τ2
959
956 e14(max) L NiSi + CrSi2 L CrSi + NiSi, CrSi2
956
U13
L+θ+τ2
959 e13(max) L CrSi + NiSi
948 e15 L NiSi + NiSi2
P1 U9 P2 E2
p4
L+Si+CrSi2
L+CrSi+CrSi2 959
Cr-Si
D2
CrSi+NiSi+CrSi2 L + τ2 θ + CrSi
944
Cr–Ni–Si
L CrSi2+ NiSi + NiSi2
936
L+ θ+CrSi
U14 δ+θ+τ2
E3
CrSi2+NiSi+NiSi2 934
L θ + NiSi + CrSi
τ2+θ+CrSi
E4
θ+NiSi+CrSi 858 p9 θ + NiSi ε
858
θ + NiSi ε, CrSi θ+ε+CrSi
NiSi+ε+CrSi 855
CrSi + θ ε + τ2
CrSi+ε+τ2 844 θ+δ+ε 824 e16 θδ+ε
U15
θ+ε+τ2 θ + τ2 δ + ε
U16
τ2+δ+ε 843
π (Ni) + γ + τ1
E5
(Ni)+γ+τ1
243
MSIT®
Fig. 1c: Cr-Ni-Si. Reaction scheme, part 3
D3
Cr–Ni–Si
244
Si
Data / Grid: at.% Axes: at.%
Fig. 2: Cr-Ni-Si. Calculated liquidus surface projection e6
1300
20
80
1200
1100 e4 p2
p8
CrSi2
40
CrSi
p1 60
P1
β Cr5Si3
e2
αCr5Si3
U1
U5τ 1
Cr3Si
1500 20
Cr
σ
γ
π E2
U8
p4 e D1 7p 5
β3
20
e9
1200 1300 1400°C
(Ni) 60
e5 Cr Ni Si
Fig. 3: Cr-Ni-Si. Isothermal section of Cr corner at 1175°C
U9
40
p6
1400 U2 40
NiSi2 e15 D 2 e14 e13 NiSi e E4 12
U13 1200
U6
U7 P2
p3
U3
(Cr) 1800
e8 δ E1
e10
80
1650
U14
τ2
U10 U4
e1
60
U12 E3
80
50.00 0.00 50.00
Ni
Data / Grid: at.% Axes: at.%
60
40
70
30
Cr3Si 80
20
L+σ+Cr3Si (Cr)+σ+Cr3Si
90
L+(Ni)+σ
σ
(Cr)+σ (Cr)
Cr
MSIT®
L+σ
σ+Cr3Si
10
(Ni)+σ
(Cr)+(Ni) 10
20
30
40
Cr Ni Si
50.00 50.00 0.00
Landolt-Börnstein New Series IV/11C3
Cr–Ni–Si
245
Si
Data / Grid: at.%
Fig. 4: Cr-Ni-Si. Isothermal section at 1125°C
Axes: at.%
(Si)
20
80
CrSi2 40
60
CrSi
αCr5Si3
L
τ2
60
40
θ
δ γ β3
Cr3Si
τ1
80
L
20
σ (Ni) (Cr) 20
Cr
40
60
80
Si Fig. 5: Cr-Ni-Si. Isothermal section at 1050°C
Ni
Data / Grid: at.% Axes: at.%
(Si)
20
80
CrSi2 40
60
CrSi L
τ2
αCr5Si3 60
40
θ
δ γ
Cr3Si 80
τ1
β2 20
π
σ
(Ni)
(Cr)
Cr
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Ni
MSIT®
Cr–Ni–Si
246
Si Fig. 6: Cr-Ni-Si. Isothermal section at 900°C
Data / Grid: at.% Axes: at.%
(Si)
20
80
αNiSi2
CrSi2 40
60
CrSi NiSi
τ2
αCr5Si3 60
40
Cr3Si
θ δ
τ1
80
γ β1
20
π σ
(Ni)
(Cr) 20
Cr
40
60
80
Si Fig. 7: Cr-Ni-Si. Isothermal section at 850°C
Ni
Data / Grid: at.% Axes: at.%
(Si)
20
80
CrSi2
αNiSi2
40
60
CrSi NiSi
τ2
αCr5Si3 60
Cr3Si
40
γ
(Cr)
MSIT®
β1
20
π
σ
Cr
δ γ
τ1
80
θ
20
40
60
80
Ni
Landolt-Börnstein New Series IV/11C3
Cr–Ni–Si
247
1500
Fig. 8a: Cr-Ni-Si. Calculated vertical section at a constant Cr content of 10 at.%
1400
L+(Si)
L
Temperature, °C
1300
L+(Ni) L+δ 1200
(Ni) 1100
(Si)+CrSi2
L+CrSi
L+γ
γ+δ
L+τ2
(Ni)+γ
L+CrSi2 L+(Si)+CrSi2
1000
(Ni)+γ+π
Cr Ni Si
900
10.00 90.00 0.00
NiSi2+CrSi2 20
τ2+δ
40
CrSi+θ 60 NiSi+CrSi 80 2 τ2+θ NiSi+CrSi
Si, at.%
Cr Ni Si
10.00 0.00 90.00
Cr Ni Si
10.00 55.00 35.00
1000
Fig. 8b: Cr-Ni-Si. An enlarged part of Fig. 8a from 25 to 35 at.% Si and from 900 to 1000°C
(Ni)+γ
γ+δ
γ+π
Temperature, °C
δ +τ2
δ +Cr3Si
δ +α Cr5Si3
950
(Ni)+γ+π
γ+δ +Cr3Si γ+Cr3Si
γ1+τ1
Cr Ni Si
Landolt-Börnstein New Series IV/11C3
900
10.00 65.00 25.00
26
28
30
Si, at.%
32
34
MSIT®
Cr–Ni–Si
248
1500
Fig. 9a: Cr-Ni-Si. Calculated vertical section at a constant Cr content of 40 at.%
L 1400
L+α Cr5Si3 L+(Ni) L+CrSi
L+Cr3Si 1300
Temperature, °C
L+CrSi2
L+σ (Ni)
1200
1100
τ1+σ γ+Cr3Si
1000
δ+ +Cr3Si
τ2+CrSi CrSi+CrSi2
τ1+Cr3Si (Ni)+σ
Cr Ni Si
900
40.00(Cr)+(Ni) 60.00 0.00
20
40
δ +α Cr5Si3 τ2+α Cr5Si3
Si, at.%
1200
Fig. 9b: Cr-Ni-Si. An enlarged part of Fig. 8a from 48 to 52 at.% Si and from 900 to 1200°C
Cr Ni Si
40.00 0.00 60.00
Cr Ni Si
40.00 8.00 52.00
L+CrSi+CrSi2
L+CrSi+α Cr5Si3
L+CrSi L+CrSi+τ2
Temperature, °C
1100
CrSi+τ2 CrSi+α Cr5Si3
1000
CrSi+CrSi2 CrSi+τ2+α Cr5Si3
Cr Ni Si
MSIT®
900
40.00 12.00 48.00
49
50
Si, at.%
51
Landolt-Börnstein New Series IV/11C3
Cu–In–Sn
249
Copper – Indium – Tin Tamara Velikanova, Michail Turchanin, Olga Fabrichnaya Introduction The Cu-In-Sn system is of great importance for development of Pb free solders. The first extensive investigation of phase equilibria was carried out in the Cu rich part of the system by [1972Koe] at 400 - 900°C. On the base of the results of [1972Koe] the phase equilibria of system were critically assessed by [1979Dri]. The phase equilibria over the whole concentration range at 110 - 900°C were studied by [2001Liu]. Experimental works of [1956Kle, 2001Ita] were devoted to studing of thermodynamic properties of liquid alloys. The experimental works are summarized in Table 1. The thermodynamic assessment of the whole diagram in the framework of CALPHAD method was presented by [2001Liu]. Binary Systems Assessments of the Cu-In system by [2003Bah], of the Cu-Sn system by [Mas2], and of the In-Sn system by [Mas2] are accepted. In [1972Koe] binary systems were taken from [H]. For the Cu-Sn and In-Sn as well as for Cu-In system at temperatures above 400°C, they are in good agreement with [Mas2]. In the thermodynamic assessment of ternary system [2001Liu] the binary descriptions are accepted from the following sources: [2002Liu] for the Cu-In system, [2004Liu] for the Cu-Sn system, and [1996Lee] for the In-Sn system. These works are in reasonable agreement with [Mas2]. There is one exception for Cu-Sn system. Authors of [2001Liu] and [2004Liu] have recently reported that a two-stage ordering reaction A2-B2-D03 in the bcc phase region takes place, rather than a / 1 two phase equilibria. In [2001Liu] a simplified description of the Gibbs energy of the bcc phase was made, were the ordering reactions of the bcc phase was not considered. It should be mentioned that all phase diagrams of the ternary system included in the present evaluation are checked and corrected for consistency with the accepted binary systems [2003Bah, Mas2]. Solid Phases The crystallographic data on the Cu-In-Sn phases and their temperature ranges of stability when available are listed in Table 2. In and Sn are markedly dissolved in Cu, but solution of Cu in (In) and (Sn) is negligible. The high temperature bcc and hexagonal phases (NiAs type) form a continuous range of solid solutions between the Cu-In and Cu-Sn sides. A weak relative stabilization of solid solution in the ternary system is found. The phase decomposes in the ternary system via a eutectoid reaction at 568°C that is slightly lower than the eutectoid decomposition of the phase in both binary systems (e5 in Cu-Sn system is at 586°C and e6 in Cu-In system is at 576.5°C). The high temperature Cu-Sn and Cu-In based cubic 1 and 2 solid solutions essentially extend into the ternary forming two-phase equilibria in the central part of the concentration triangle. Dissolving a third component decreases the decomposition temperature of both phases down to 517 and 509°C for 1 and 2, respectively. As a result, 1 and 2 solid solutions exist in the ternary system as separate ternary phases in the temperature intervals of 517 to 520°C and 509 to 618°C, respectively, after [1972Koe]. The existence of the second ordering reaction of the bcc phase rather than the (bcc) + (D03) two-phase equilibrium in the Cu-Sn binary system is accepted by [2001Liu] (referring to own unpublished data) in contradiction to the Cu-Sn phase diagram accepted in the current assessment and by [1972Koe]. This question is open so far and further investigation is required. The Cu-Sn based and J phases as well as Cu-In based 2 phase exhibit the wide homogeneity ranges. The solubility of a third component essentially increases with decreasing temperature. The high temperature binary phase is stabilized by In additions down to 400°C and possibly below according to [1972Koe].
Landolt-Börnstein New Series IV/11C3
MSIT®
250
Cu–In–Sn
Below 582°C, the decomposition temperature of the phase in the Cu-Sn system, the ternary solid solution exists as a separate ternary phase like the above mentioned phases. The Cu-Sn based 1 phase exists in a narrow composition range close to the Cu-Sn side from ~590°C (temperature of the peritectoid reaction of its formation in the Cu-Sn system after [1972Koe]) to ~350°C (temperature of the eutectoid decomposition of the ternary 1 solid solution). Liu et al. [2001Liu] reported the Cu-Sn based phase in ternary alloys only at 600°C. Instead, a wide composition range of the 1 phase is proposed, for example, at 500°C up to about 16 mass% (10.5 at.%) In at 19 mass% (12.2 at.%) Sn in equilibrium with 2 and J phases. This contradicts the results of [1972Koe]. The authors of [2001Liu] did not discuss this contradiction and it will not be discussed in the current assessment because of the lack of information on the original experimental data in both works. The interpretation of phase equilibria proposed by [1972Koe] seems to be preferable since more alloys in this range of compositions were studied and an XRD method was used in addition to metallography and DTA (DSC in [2001Liu]) methods. A diffusion couple method for establishing an of equilibrium state seems to be less reliable compared to conventional metallurgical methods even if long-time equilibrated couples are used. Two ternary compounds, -1 and -2, exist in the system. Compound -1 (Cu11In2Sn) found by [1972Koe] forms by a peritectoid reaction at 585°C. At 400°C, its homogeneity range continues from 3 to 10 at.% Sn at almost constant content of Cu. This phase was confirmed by [1993Mor] and [2001Liu]. According to [2001Liu] the -1 phase has composition Cu16In3Sn and disappears at a temperature below 400°C. The decomposition reaction for the -1 phase based on the ternary phase diagram after [1972Koe] may be assumed as -1 (Cu) + 2 + . The -2 phase was reported by [1984Rom]. This phase was also found by [2001Liu] at 110°C. Its composition after [2001Liu] is Cu2In3Sn. According to calculations and DSC data of [2001Liu] this compound exists below 156°C. The crystal structure of both -1 and -2 phases is not established. Invariant Equilibria The data on the invariant equilibria are given in Table 3. Reaction scheme is given in Fig. 1. For the composition range from 50 to 100 at.% Cu equilibria U1 to U10, P1, E1, E2, and E3 are taken from the experimental work of [1972Koe]. For the copper poor part of the system four calculated invariant equilibria involving the liquid phase U11, U12, E6 and L + (In) 3 + (marked as P3 in [2001Liu]) are reported by [2001Liu]. However, reactions involving the -2 phase formation are not presented by [2001Liu]. It should be noted that the original experimental data on the phase composition of equilibrated at 110°C alloys, DSC data and vertical sections at 10, 20, and 30 mass% Cu of [2001Liu] indicate the equilibrium of the -2 phase with liquid. However, according to calculations of [2001Liu] liquid is not stable at 110°C. The -2 formation reaction may be assumed from the reaction scheme above 150°C as L + + (In) -2, P2. To connect the reaction P2 with the reaction E6 calculated by [2001Liu] two invariant equilibria U13 and U14 should be assumed. The equilibria U11, U12, and E6 are taken from [2001Liu]. To reach agreement between isothermal section at 110°C and vertical sections at 10, 20, and 30 mass% Cu presented by [2001Liu] the invariant reaction between solid phases + 3 -2 + at ~109°C, U15, should be supposed. The calculated temperatures of the invariant equilibrium of the liquid, , and 2 phases (651.6°C) and of the invariant equilibrium of the liquid, , J and phases (612.4°C) by [2001Liu] are close to the temperatures of the transition reactions U2 and U3, respectively, after [1972Koe]. However instead of the phase, the 1 phase participates in these reactions according to [1972Koe]. The eutectoid decomposition, E4, of the -1 ternary compound is proposed in the current assessment taking into account the phase equilibria at 400°C (and higher temperatures) after [1972Koe] and the experimental data of [2001Liu], who found, that the -1 phase is stable at 400°C and disappears in alloys equilibrated at 250°C. [1972Koe] considers the possibility of the eutectoid decomposition of the 1 phase via the reaction
1 (Cu) + + J, E5. The existence of such an invariant reaction is in good agreement with the isothermal section at 400°C after [1972Koe]. The temperature of this decomposition should be very close to the MSIT®
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temperature of the eutectoid reaction 1 (Cu) + J in the binary Cu-Sn, since the solubility of In in 1 is low. It was shown to be less than 1 at.% at 400°C. The phases , ’ and ’’ are undistinguished in discussing the ternary system, and the corresponding equilibria are not considered here and below, although they should exist at low temperatures. Liquidus Surface The partial liquidus surface projection for the copper rich part of the ternary diagram investigated by [1972Koe] (50 to 100 at.% Cu) is given in Fig. 2. Figure 3, taken from [2001Liu] does not present stable equilibria and may be considered reflecting the metastable crystallization of the alloys, excluding the formation of -2. It should be mentioned that, reaction U* in Fig. 3 (P3 according to [2001Liu]) is metastable, because experiments indicate the formation of -2 phase and therefore it should be involve in the reactions with the liquid phase. Isothermal Sections Experimental isothermal sections for Cu rich alloys at 900°C according to [2001Liu] and at 650, 600, 555, and 400°C after [1972Koe] are given in Figs. 4 to 8. It should be mentioned that equilibrium of phase with liquid is shown in Fig. 8 according to experimental data of [2001Liu]. Continuous solid solutions of phase at 650 and 600°C (Figs. 5 and 6) and phase at 400°C (Fig. 8) from the Cu-In to the Cu-Sn side are exposed. An important feature of the phase equilibria in this system is the high solubility of a third component in the Cu-In and Cu-Sn based phases 2, 1, 2, and (Figs. 5 to 8). Another feature is that the high temperature binary phases 2, 2, and are stabilized to lower temperature by picking up a third component, as demonstrated in Figs. 6 to 8. Very low solubility of In in the 1 phase is shown in Figs. 7 and 8 after [1972Koe] in contradiction with [2001Liu]. Experimental sections at 650, 600, 500, 400 and 250°C, in the whole concentration range, and at 180 and 110°C below 40 mass% Cu as well as the calculated ones at 650, 400, 180, and 110°C in the whole concentration range, are presented by [2001Liu]. The comparison of experimental data of [1972Koe] and [2001Liu] reveals several contradictions. According to [2001Liu] at 600°C the phase containing 1.9 mass% (1.25 at.%) In is in equilibrium with J and (instead of 1 after [1972Koe]). Also [2001Liu] considered the bcc based phases in the Cu-Sn system as a single phase field inconsistent with the two phase field + 1 accepted by [1972Koe] and in the current assessment. According to [2001Liu] the 1 phase presents extended solid solution at 500 and 400°C being in equilibrium with the (Cu), J, , and 2 phases that is inconsistent with the data of [1972Koe]. In the calculated sections at 180 and 110°C by [2001Liu], the 1 phase coexists with the phase in contradiction with data of [1972Koe] indicating a low solubility of In in 1. The results of [1972Koe] for a high temperature range, including 650, 600 and 400°C, are preferred because they are based on investigation of more alloys using an XRD method for identification of phases additionally to metallography and DTA (DSC in [2001Liu]) methods. Only the low Cu region of the isothermal sections at 180 and 110°C are accepted in the present assessment according to [2001Liu]. The rest can not be accepted because of the above mentioned contradictions with the data of [1972Koe]. The isothermal section at 180°C in the composition range 0 to 50 mass% Cu according to experimental data of [2001Liu] is shown in Fig. 9. The section at 108°C shown in Fig. 10 is derived from the experimental data of [2001Liu] for the temperature-composition sections at 10, 20, and 30 mass% Cu and for the isothermal section at 110°C. The section at 110°C after [2001Liu] differs from the one given in Fig. 10 in the central part of the composition range by the existence of the + 3 equilibrium instead of the alternative -2 + . According to the proposed reaction scheme, the section at 110°C falls into a very narrow temperature range between solid state reaction U15 (109°C) and the eutectic E6 (111.2°C). The isothermal section presented in Fig. 10 is more representative, because the topology will be preserved at wider temperature range below the U15 reaction. The tie line + 3 at 110°C is shown by a dashed line in Fig. 10. Difficulties in the realization of the equilibrium state especially for Cu poor alloys complicate investigation of phase equilibria very essentially. The disagreement in interpretations of the phase equilibria in different works arises from this fact. Further work is needed for clearing the situation not only near the In-Sn side of the system but in the whole composition range at low temperatures. Landolt-Börnstein New Series IV/11C3
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Cu–In–Sn
Temperature – Composition Sections Nine experimental temperature-composition sections are reported by [1972Koe]. Seven of them parallel to the In-Sn side in the whole concentration range are shown in Figs. 11 to 16. The section at 60 at.% Cu in the composition range 0 to 25 at.% In is given in Fig. 17. The wide homogeneity ranges of the 1, , J, -1, , and phases are cut by the sections. No vertical section of the ternary system is quasibinary, including ones through the and solid solutions. At 400°C the -1 phase is in equilibria with the (Cu) and phases in the section at 4 at.% In (Fig. 18) and in equilibria with (Cu) and 2 phases in the section at 2 at.% Sn (Fig. 19). The vertical sections at 10 and 30 mass% Cu are given in Figs. 20 and 21, respectively, partially based on the calculations of [2001Liu]. However, it is found in the present assessment that several phase boundaries are missed in the calculations of [2001Liu]. On the other hand, the topology of the vertical sections in the temperature range of 112 - 150°C obtained by [2001Liu] is in contradiction with the presented reaction scheme. Therefore, the phase relations in this temperature range are corrected in the present assessment to correspond to the accepted reaction scheme. Several data points obtained by DSC in [2001Liu] are in agreement with the vertical sections presented in Figs. 20-21. These sections demonstrate the coexistence of the -2 compound with the liquid together with , (In) and 3 phases. However it should be pointed out that the vertical sections presented in Figs. 20-21 are tentative in the temperature range 100-300°C and the composition range from 0 to 40 mass% Sn. Thermodynamics The mixing enthalpies of the liquid SnxCu1-x-SnxIn1-x alloys at x = 0.84 were measured callorimetrically by [1956Kle] by direct mixing of liquid SnxCu1–x and SnxIn1–x alloys in a calorimeter at 450°C. The results are presented in Table 4. The thermodynamic properties of the liquid solution between these terminal binary alloys are characterized by slight positive deviations from ideality. The thermodynamic activity of indium was measured by [2001Ita]. Measurements were carried out by emf method along xCu/xSn = 4, xCu/xSn = 1, and xCu/xSn = 0.25 sections at 727, 827, and 927°C, respectively. Results obtained are presented graphically in Figs. 22a, 22b and 22c. The positive deviations of thermodynamic properties of the liquid alloys from ideality increase with increasing copper contents. Results of [2001Ita] were used by [2001Liu] in the thermodynamic assessment of the system. The assessments of thermodynamic parameters in the ternary system are available from works [1997Lee, 2001Liu]. The isothermal section of the Cu-In-Sn ternary system at 250°C calculated using only the binary thermodynamic parameters is reported by [1997Lee]. Due to the lack of information on thermodynamics, the ternary compounds -1 and -2, as well as the homogeneity ranges of binary phases were not included in the calculation. This results in a serious disagreement of the calculations of [1997Lee] with the experimental and calculated phase diagram of [2001Liu] at 250°C. The assessment of [2001Liu] takes into account homogeneity ranges of solid phases as well as formation of -1 and -2 phases. However the calculated isothermal and vertical sections of [2001Liu] seem to be very contradictory (see chapters Isothermal Sections and Temperature–Composition Sections). Obviously more careful Calphad assessment is necessary to get quantitative character of isothermal and vertical sections at 110 - 250°C. Notes on Materials Properties and Applications The study of dispersion-hardening pastes containing copper, indium, and tin was carried out by [1981Mak]. It was reported that after heat treatment of the pastes at 200°C the alloys with a melt point of 600 - 800°C were obtained. The isothermal section of the Cu-In-Sn phase diagram at 250°C and metastable phase equilibria between (Cu) and liquid at the same temperature were calculated by [1997Lee] by CALPHAD method using the only binary thermodynamic parameters. The results obtained were used for prediction of interface reaction products between copper and In-Sn solder alloys. These predictions were extended by [2003Jeo] to the case of formation of ternary compounds in the system using thermodynamic assessment [2001Liu]. The growth kinetics of intermetallic compound layers formed between In-48 at.% Sn solder and Cu substrate were investigated by [2002Som] and [2005Kim]. In [2002Som] the diffusion reaction process in MSIT®
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Cu/In-48 at.% Sn/Cu joints were investigated between 180 and 400°C. Electron microprobe analysis revealed the presence of one or two intermetallic layers in the interaction zone. A layer of the phase formed below 200°C. Layers of the and phases formed above 200°C. According to [2005Kim], in the solder joint Cu/In-48 at.% Sn at 70 to 100°C, the intermetallic compound layer was composed of two phases: -2 adjacent to the solder and which was the dominant phase. A quantitative analysis of the intermetallic compound layer thickness as a function of time shows a constant growth rate at constant temperature [2002Som, 2005Kim]. The selective formation of intermetallic compounds in Sn-20In-0.8Cu (mass%) ball grid array solder joints with Au/Ni [2004Wu] and Au/Ni/Cu [2004Chi] surfaces was investigated. References [1956Kle] [1972Koe]
[1979Dri]
[1981Mak]
[1984Rom]
[1992Che]
[1993Mor]
[1996Lee]
[1997Lee]
[2001Ita]
[2001Liu]
[2002Som]
Landolt-Börnstein New Series IV/11C3
Kleppa, O.J., “A Calorimetric Investigation of Some Binary and Ternary Liquid Alloys Rich in Tin”, J. Phys. Chem., 60, 842 (1956) (Experimental, Thermodyn., 15) Köster, W., Gödecke, T., Heine, D., “The Constitution of the Copper-Indium-Tin System in the Range from 100 to 50 at.% Cu”, Z. Metallkd., 63, 802-807 (1972) (Experimental, Phase Diagram, #, 4) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, Rokhlin, L.L., Turkina, N.I., “Cu-In-Sn” (in Russian), in “Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 126-127 (1979) (Crys. Structure, Phase Diagram, Phase Relations, Review, 2) Makedontsev, M.A., Sarpov, N.I., Yuzhin, A.I., “Study of Dispersion-Hardening Pastes in the Indium-Tin-Copper System”, Russ. Metall., (2), 212-216 (1981) (Experimental, Morphology, Crys. Structure, 12) Romig, A.D., Jr, Yost, F.G., Hlava, P.F., “Intermetallic Layer Growth in Copper/Tin-Indium Solder Joints”, Microbeam Anal., 19th, 87-92 (1984) (Experimental, Morphology) cited from abstract Che, G.C., Ellner, M., “Powder Crystal Data for the High-Temperature Phases Cu4In, Cu9In4 (h) and Cu2In (h)”, Powder Diffr., 7(2), 107-108 (1992) (Crys. Structure, Experimental, 12) Morris, J.W. Jr., Goldstain, J.L., Freer, Mei, Z., “Microstructure and Mechanical Properties of Tin-Indium and Tin-Bismuth Solders”, JOM, 45(7), 25-27 (1993) (Experimental, Morphology, Mechan. Prop.) cited from abstract Lee, B.-E., Oh, C.-S., Shim, J.-H., “Thermodynamic Assessment of the Sn-In and Bi-In Binary Systems”, J. Electron. Mater., 25(6), 983-991 (1996) (Assessment, Phase Relations, Thermodyn., 50) Lee, B.-J., Hwang, N.M., Lee, H.M., “Prediction of Interface Reaction Products Between Cu and Various Solder Alloys by Thermodynamic Calculation”, Acta Mater., 45(5), 1867-1874 (1997) (Calculation, Phase Diagram, Phase Relations, Interface Phenomena, Thermodyn., 26) Itabashi, S., Kameda, K., Yamagauchi, K., Abstracts of the 128th Japan Institute of Metals Annual Meeting in Spring, Jim, Narashito, Japan (2001) (Experimental, Thermodyn.) as quoted by [2001Liu] Liu, X.J., Liu, H.S., Ohnuma, I., Kainuma, R., Ishida, K., Itabashi, S., Kameda, K., Yamaguchi, K., “Experimental Determination and Thermodynamic Calculation of the Phase Equilibria in the Cu-In-Sn System”, J. Electron. Mater., 300(9), 1093-1103 (2001) (Calculation, Experimental, Phase Relations, Phase Diagram, Thermodyn., 19) Sommadossi, S., Gust, W., Mittemeijer, E.J., “Characterization of the Reaction Process in Diffusion-Soldered Cu/In-48 at.% Sn/Cu Joints”, Mater. Chem. Phys., 77, 924-929 (2002) (Morphology, Experimental, Kinetics, 29)
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254 [2003Bah]
[2003Jeo]
[2002Liu]
[2004Chi]
[2004Wu]
[2004Liu]
[2005Kim]
Bahari, Z., Dichi, E., Legendre, B., Dugue, J., “The Equilibrium Phase Diagram of the Copper-Indium System: a New Investigation”, Thermochim. Acta, 401(2), 131-138 (2003) (Experimental, Phase Diagram, Thermodyn., 17) Jeong, S.W., Kim, J.H., Lee, H.M., “Thermodynamic Issues of Lead-Free Soldering in Electronic Packaging”, Mater. Sci. Forum, 426-432, 4081-4086 (2003) (Calculation, Morphology, Experimental, Phase Diagram, Phase Relations, Thermodyn., Interface Phenomena, 17) Liu, H.S., Liu, X.J., Ciu, Y., Wang, C.P., Ohnuma, I., Kainuma, R., Jin, Z.P., Ishida, K., “Thermodynamic Assessment of the Cu-In Binary System”, J. Phase Equilib., 23(5), 409-415 (2002) (Assessment, Phase Relations, Phase Diagram, Thermodyn., *, 36) Chiang, M.J., Chang, S.Y., Chuang, T.H., “Reflow and Burn-in of a Sn-20In-0.8Cu Ball Grid Array Package with a Au/Ni/Cu Pad”, J. Electron. Mater., 33(1), 34-39 (2004) (Morphology, Experimental, 17) Wu, H.-F., Chiang, M.-J., Chuang, T.-H., “Selective Formation of Intermetallic Compounds in Sn-20In-0.8Cu Ball Grid Array Solder Joints with Au/Ni Surface Finishes”, J. Electron. Mater., 33(9), 940-947 (2004) (Morphology, Experimental, 11) Liu, X.J., Wang, C.P., Ohnuma, I., Kainuma, R., Ishida, K., “Experimental Investigation and Thermodynamic Calculation of the Phase Equilibria in the Cu-Sn and Cu-Sn-Mn Systems”, Metall. Mater. Trans. A, 35a, 1641-1654 (2004) (Calculation, Crys. Structure, Experimental, Phase Diagram, Thermodyn., 55) Kim, D.-G., Jung, S.-B., “Interfacial Reactions and Growth Kinetics for Intermetallic Compound Layer Between In-48Sn Solder and Bare Cu Substrate”, J. Alloys Compd., 386, 151-156 (2005) (Morphology, Experimental, Interface Phenomena, Phys. Prop., Kinetics, 15)
Table 1: Investigations of the Cu-In-Sn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1956Kle]
High temperature calorimetry
450°C, SnxCu1–x - SnxIn1–x, x = 0.84
[1972Koe]
Optical microscopy, X-ray analysis; differential thermal analysis
xCu > 0.5, four isothermal sections at 650, 600, 555, 400°C; nine vertical sections at 83.5, 80, 77, 75, 70, 65, 60 at.% Cu, 4 at.% In, 2 at.% Sn
[2001Ita]
Admittedly emf method with zircona solid electrolyte
aIn, 727-927°C, xCu/xSn = 4, xCu/xSn = 1, xCu/xSn = 0.25
[2001Liu]
Diffusion couple method; differential scanning calorimetry (heating curve), metallography, energy dispersive X-ray spectroscopy, thermodynamic assessment using the CALPHAD method
Eight isothermal sections at 900, 650, 600, 500, 400, 250, 180, 110°C; eight vertical sections at 10, 20, 30, 40, 50, 60, 70, 80 mass% Cu.
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Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, (Cu)
cF4 Fm3m Cu
Cu100-x-yInxSny < 1084.62
Lattice Parameters Comments/References [pm] dissolves 11.0 at.% In at 576.5°C and y = 0 [2003Bah] dissolves 9.1 at.% Sn at 520-586°C and x = 0 [Mas2] a = 361.46
(In) < 156.634
tI2 I4/mmm In
(Sn) < 231.9681
tI4 I41/amd Sn
, Cu100–x–yInxSny 798 - 568 Cu13Sn3 798 - 586
cI2 Im3m W
a = 325.3 c = 497.70 a = 583.18
a = 302.61
1, Cu100–x–yInxSny, 755 - 517 Cu4Sn 755 - 520
cF16 F43m CuHg2Ti
1, Cu100–x–yInxSny, Cu41Sn11 590 - 350
cF416 F43m Cu41Sn11
, Cu100–x–yInxSny, 640 - 400
hP26 P63 Cu10Sn3
a = 611.76
Cu10Sn3 640 - 582
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dissolves 12 at.% Sn at 0°C [Mas2] pure In, 25°C [Mas2] dissolves several at.% In [Mas2] pure Sn, 25°C [Mas2]
13.1 y 16.5 at x = 0 [Mas2] 13.1 y 27.5 and 798-520°C at x = 0 in [2001Liu] at 14-25 at.% Sn, x = 0 and 710°C [V-C2]
18.5 x 23.7 at y = 0 [2003Bah] a = 298.0 to 299.0 at 16 x 18 and y = 0 [1992Che]
Cu100–xInx 711.3 - 576.5
J, Cu100–x–yInxSny, Cu3Sn < 676
pure Cu, 25°C [Mas2]
a = 1798.0
a = 733.0 c = 786.40 oC80 Cmcm Cu3Sn
a = 552.9 b = 4775.0 c = 432.3
15.5 y 27.5 at x = 0 [Mas2] at 14-25 at.% Sn, x = 0 and 710°C [V-C2] 20 y 21 at x = 0 [Mas2] [V-C2]
20.3 y 22.5 at x = 0 [Mas2] [V-C2]
24.5 y 25.9 at x = 0 [Mas2] [V-C2]
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256 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, Cu100–x–yInxSny, Cu6Sn5 415 - 186
hP4 P63/mmc NiAs
, Cu100–xInx 670.2 - 276.6
Lattice Parameters Comments/References [pm]
a = 419.2 c = 503.7 = 120° a = 426.9 c = 523.9 = 120°
43.5 y 45.5 at x = 0 [Mas2] [V-C2]
33.1 x 36.2 at y = 0 [2003Bah] at x = 35.6 [V-C2]
’, Cu100–ySny < 189
based on NiAs?
2, Cu100–x–yInxSny, 684.1 - 509 Cu9In4 684.1 - 617.8
cP52 P43m InMn3
2, Cu100–x–yInxSny, Cu7In3 < 632.2
aP40 P1 Cu7In3
’’, Cu100–xInx < 388.3
Ni2In
Q, Cu11In9 < 305.8
mC20 C2/m AlCu
3, In100–xSnx < 140.0
tI2 I4/mmm In
, In100–xSnx < 224
hP5 P6/mmm -
* -1, Cu11In2Sn 585 - (32575)°C
?
-
[1972Koe] Cu16In3Sn in [2001Liu]
* -2, Cu2In3Sn < 156°C
?
-
[2001Liu]
MSIT®
44.8 y 45.5, hexagonal [Mas2]
a = 925.03
a = 1007.1 b = 912.6 c = 672.4 = 90.22° = 82.84° = 106.81° -
a = 1281.4 b = 435.4 c = 735.3 = 54.49° a = 346.3 c = 440.4 a = 320.5 c = 299.5 = 120°
27.9 x 32.2 at y = 0 [2003Bah] at x = 29.6, y = 0 and 650°C [V-C2] 29.1 x 31.2 at y = 0 [2003Bah] at x = 30 and y = 0 [V-C2]
33.3 x 36.6 [2003Bah]
[2003Bah] at 280°C [V-C2]
12 x 44 [Mas2] at x = 25 [V-C2] 72 x 98.5 [Mas2] at x = 80 [V-C2]
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Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%)a) Cu
In
Sn
L + 1 + 2
676
U1b)
L
76 (62)
17 (25.5)
8 (12.5)
L + 2 1 +
648
U2b)
L
66.5 (51.9)
19.5 (27.7)
14 (20.4)
+ 1 + 2 -1
585
P1b)
-
-
-
-
L + 1 J +
583
U3
L
58.5 (43.2)
13 (17.4)
28.5 (39.4)
+ 2 1 + -1
578
U4
-
-
-
-
1 + J +
575
U5
-
-
-
-
+ 2 + -1
571
U6
-
-
-
-
(Cu) + 1 + -1
568
E1
-
-
-
-
1 + + 2
562
U7
-
-
-
-
1 + 2 + -1
552
U8
-
-
-
-
2 + + 2
551
U9
-
-
-
-
1 + -1 (Cu) +
540
U10
-
-
-
-
2 + 2 + -1
509
E2
-
-
-
-
1 (Cu) + 1 +
517
E3
-
-
-
-
-1 (Cu) + 2 +
325±50
E4
-
-
-
-
1 (Cu) + J +
~350
E5
-
-
-
-
L + (Sn) +
218.5
U11c)
L
1.18 (0.64) 0.02 (0) 56.17 (40.91) 0 (0)
5.08 (4.95) 1.56 (1.53) 12.47 (16.42) 2.93 (2.83)
93.74 (94.41) 98.42 (98.47) 31.36 (42.67) 97.07 (97.17)
0.57 (0.32) 55 (40.35) 60.61 (87.96) 0 (0)
99.43 (99.68) 45.0 (59.65) 35.13 (48.07) 99.999 (99.999)
1.4#10–3 (1.5#10–3) 0.0 (0.0) 4.26 (6.02) 0.001 (0.001)
(Sn) L + Q + (In)
154.1
U12c)
L Q (In)
L + + (In) -2
~150
P2
-
-
-
-
L + (In) -2 + 3
~140
U13
-
-
-
-
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L + -2 + 3 L+3+
T [°C]
Type
~112
U14
111.2
c)
E6
Phase
Cu
In
Sn
-
-
-
-
L
1.28 (0.75) 57.1 (36.64) 0 (0) 0 (0)
53.78 (53.23) 23.81 (31.68) 56.55 (55.73) 23.44 (22.85)
44.94 (46.02) 19.09 (26.26) 43.45 (44.27) 76.56 (77.15)
-
-
-
3 + 3 -2 +
~109
U15
Composition (at.%)a)
-
a) mass% are given in brackets; b) coordinates are taken from Fig. 2; c) calculated by [2001Liu].
Table 4: Thermodynamic Data of Reaction or Transformation Reaction or Transformation
T [°C]
Quantity, per mol of compound [J#mol–1]
Comments
SnxCu1–x(L) + SnxIn1–x(L)
450
83
x = 0.8354 y = 0.1123 z = 0.0523
450
73
x = 0.8466 y = 0.0962 z = 0.0572
SnxCuyInz(L)
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Landolt-Börnstein New Series IV/11C3
Cu-Sn
Cu-In-Sn
Cu-In
In-Sn
798 p1 l+αβ 755 p2 l + β γ1
711.3 p3 l+αβ
L + β γ1 + γ2
676
L + γ2 γ1 + η
648
U1
670.2 p4 l + γ2 η
β + γ1+ γ2
L + γ1+ γ2
637 p5 γ1 + ε ζ
U2 620.1 e3 γ2 β + δ2
γ1+ γ2 + η
617.8 e4 γ2 δ2 + η
L + γ1+ η
594 p6 γ1 + ζ δ1 586 e5 β α + γ1
β + γ1 + γ2 τ1
585 583
582 e7 ζ δ1 + ε
L + γ1 ε + η L+ε+η
575
U3
β+γ1+τ1
γ1+ε+η
γ1 + ε η + ζ
578
571 γ+ζ+η
Fig. 1a: Cu-In-Sn. Reaction scheme, part 1
U4 576.5 e6 β α + δ2
γ2+δ2+τ1
β + δ2 α + τ1
U6
α+β+τ1 E1
α+δ2+τ1
γ1+α+τ1 U10
β+γ2+τ1
U7
U7
U8
E4
E3 U9
259
MSIT®
E2 p 7 E5
β α + γ1 + τ 1
P1
β + γ2 δ2 + τ1
β + δ2 + τ1
U5
ζ+η+ε
568
γ1+γ2+τ1
Cu–In–Sn
640 e2 γ1 ε + l
679 e1 l β + γ2
Cu-In-Sn E1
U3 e 7 p6
U2
U5
Cu-In P1
U6
U4
In-Sn
260
MSIT®
Cu-Sn
e4
α + δ2 + τ1 γ1 + α + τ 1
γ1 + η ζ + γ2
562
γ1 + γ2 + ζ γ1 + γ2 ζ + τ1
552
γ1 + τ 1 α + ζ
540
γ1+α+ζ 517
415 p7 l+ ε η 350 e9 δ1 α + ε
U10
ζ + γ2 + η γ2 + η ζ + δ2
551
γ2 + δ2 + ζ
γ1 α + δ1 + ζ
E2
γ2 ζ + δ2 + τ1
509
δ1 α + ζ + ε
E3
ζ + δ2 + τ1 325+/-75
227 e10 l η + (βSn)
ζ + δ2 + η
α +τ1+ζ
α + δ1 + ζ
ca. 350
U9
Cu–In–Sn
520 e8 γ1 α + δ1
U8
U7
τ1 α + δ2 + ζ
E4
α + δ2 + ζ
E5
305.8 p8 l + η ϕ
α+ζ+ε 218.5 L + (βSn) η + θ U11 L+η+θ
θ + (βSn) + η
154.1 L + ϕ (In) + η U12
Landolt-Börnstein New Series IV/11C3
L+η+(Ιn) E6
Fig. 1b: Cu-In-Sn. Reaction scheme, part 2
P2
η+ϕ+(Ιn)
155.5 e11 l ϕ + (In)
224 p9 l + (βSn) θ
Landolt-Börnstein New Series IV/11C3
Cu-Sn
Cu-In-Sn
Cu-In
In-Sn
U12
U11
ca.150
L + η + (In) τ2
P2
η+(In)+τ2
L+τ2+η L+τ2+(In) ca.140
140 p10 l + (In) χ
L+ (In) τ2 + χ
(In)+τ2+χ
L+ τ2 + χ
Cu–In–Sn
L + τ2 η + χ
ca.112
U13
U14
L+η+χ 111.2
Lη + χ + θ
E6
120 e12 lχ+θ
τ2+η+χ
η+χ+θ ca. 109 τ2+χ+θ
U15
τ2+η+θ
261
MSIT®
Fig. 1c: Cu-In-Sn. Reaction scheme, part 3
η + χτ2 + θ
Cu–In–Sn
262
50.00 0.00 50.00
Cu In Sn
Fig. 2: Cu-In-Sn. Partial liquidus surface projection
Data / Grid: at.% Axes: at.%
e2,640 60
40
550
680 700 70
30
725
U3
750
600
80
20
p2,755
625
650
p1,798
U2
900
640
90
10
950 660
U1
1000
675 1050°C
Cu
Fig. 3: Cu-In-Sn. Partial liquidus surface projection near the In-Sn side calculated without considering -2 phase
10
e ,679 850 800 20 p3,711.3 1
β(Sn) e10
30
p4,670.2 40
Cu In Sn
50.00 50.00 0.00
Sn p9 U11
θ
e12 η
E6
150
χ
20 0
U*
p10 (In) U12
98
In, mass% MSIT®
99 p 8
ϕ
e11
In
Landolt-Börnstein New Series IV/11C3
Cu–In–Sn Cu In Sn
Fig. 4: Cu-In-Sn. Isothermal section at 900°C
263 50.00 0.00 50.00
Data / Grid: at.% Axes: at.%
60
40
70
30
80
20
L 90
10
(Cu)+L
(Cu) 10
Cu
20
30
Cu In Sn
Fig. 5: Cu-In-Sn. Isothermal section at 650°C
40
50.00 0.00 50.00
30
β +γ 1
20
L+γ 1+γ 2 L+γ2 L+γ 2+η
β +γ 1+γ 2 γ1+γ 2 β
10
γ 2+η γ2
(Cu)
Landolt-Börnstein New Series IV/11C3
L
ε +γ 1 γ1
Cu
η
β +γ 2 10
50.00 50.00 0.00
40
L+γ 1
ε
(Cu)+β
Cu In Sn
Axes: at.%
70
90
50.00 50.00 0.00
Data / Grid: at.%
60
80
Cu In Sn
20
30
L+η 40
MSIT®
Cu–In–Sn
264
50.00 0.00 50.00
Cu In Sn
Fig. 6: Cu-In-Sn. Isothermal section at 600°C
Data / Grid: at.% Axes: at.%
60
40
L+ε
L
ε+γ 1+L
70
30
L+γ 1
ζ
ε
80
L+γ 1+η
ε+γ 1
ζ+ε+γ 1 90
γ 1+η γ 1+γ 2+η
γ1 γ 1+γ 2
β+γ
(Cu)+β
1
γ2
β +γ 2
(Cu)
β +γ 1+γ 2 10
Cu
20
20
β
η
10
η+γ 2+δ2 30
40
δ 2 δ 2+η
50.00 0.00 50.00
Cu In Sn
50.00 50.00 0.00
Cu In Sn
50.00 50.00 0.00
Data / Grid: at.% Axes: at.%
L+ε
60
L+η
δ 2+γ 2
β +δ 2 β +δ2+γ 2 Cu In Sn
Fig. 7: Cu-In-Sn. Isothermal section at 555°C
γ 2+η
40
L 70
30
ε δ1
L+ε+η
80
ζ +ε
γ
1+
γ1 90
ζ
ζ+η
(Cu)+γ 1
γ2+ζ+η
γ 1+γ 2+τ 1 (Cu) (Cu)+γ 1+τ 1
Cu
MSIT®
20
ε +η ε +ζ +η
ζ
(Cu)+τ 1 10
γ 1+τ 1 τ 1+γ 2
τ1
γ2
τ 1+δ2+γ 2 (Cu)+δ2+τ1 20
γ 2+δ2 δ2
10
L+η
δ2+η η 30
40
Landolt-Börnstein New Series IV/11C3
Cu–In–Sn
265
Sn
Data / Grid: at.% Axes: at.%
Fig. 8: Cu-In-Sn. Isothermal section at 400°C 20
80
40
60
60
L
40
L+η
ε +ζ +δ 1 ε +η δ1 ε 80 ε+ζ+η
α+ζ+δ1
ζ α
Cu
20
δ 2+ζ +η η
τ1
α+τ 1+ζ α+δ2+τ 1
20
δ2
40
60
80
Sn
Data / Grid: at.%
(β Sn)
Fig. 9: Cu-In-Sn. Experimental partial isothermal section at 180°C in the Cu poor region
In
Axes: at.%
θ η+(β Sn)+θ η +θ
20
80
L+η+θ
40
60
L+η
60
40
L 80
20
η
L+ϕ+η
Cu
Landolt-Börnstein New Series IV/11C3
20
40
ϕ
60
80
In
MSIT®
Cu–In–Sn
266
Sn
Data / Grid: at.%
(β Sn)
Fig. 10: Cu-In-Sn. Isothermal section at 108°C in the Cu poor region. Dashed line indicates equilibria at 110°C
Axes: at.%
η+θ+(β Sn) θ
20
80
η+θ 40
60
η+(β Sn)
60
τ 2+χ +θ
40
τ 2+η +θ
η+τ 2
η
χ
τ 2+χ
τ2
80
20
τ 2+η+(In)
(In)+τ 2
(In)+τ 2+χ (In)
η+(In) 20
Cu
40
ϕ
60
80
η+(In)+ϕ
800
Fig. 11: Cu-In-Sn. Vertical section at 83.5 at.% Cu
L L+α
L+β L+α +β
Temperature, °C
700
γ1+β
β
α +β
600
In
γ1 α +β +δ 2
α +β +τ1
568
571
α +τ1+γ1
γ1+β +α γ1+δ 1 540
500
α +δ 2
ζ +α
α +ζ α +τ1+δ 2
Cu 83.50 In 16.50 Sn 0.00
MSIT®
10
Sn, at.%
517
α +δ 1 α+δ 1+ζ
α +τ1+ζ
400
γ1+ζ +α
Cu 83.50 0.00 In Sn 16.50
Landolt-Börnstein New Series IV/11C3
Cu–In–Sn
267
800
Fig. 12: Cu-In-Sn. Vertical section at 80 at.% Cu
L
L+α L+β +α
Temperature, °C
700
L+β
L+γ1 L+β +γ1
β +γ1
β 600
β +τ1 571
500
γ1
β +τ1+γ1
β +α
β +τ1+δ 2
α +τ1+β
α +δ 2
α +γ1
568
γ1+δ 1 γ1+ζ
α +τ1+γ1
γ1+ζ +δ 1
540
γ1+ζ +α
α +τ1
α +δ 1
α +ζ
α +τ1+δ 2
α +δ 1+ζ
α +τ1+ζ
400
Cu 80.00 In 20.00 Sn 0.00
Cu 80.00 0.00 In Sn 20.00
10
Sn, at.%
800
Fig. 13: Cu-In-Sn. Vertical section at 77 at.% Cu
L L+γ1 L+β +γ1
L+β +γ2
Temperature, °C
700
L+β β
β +γ2+δ 2 600
γ1
675
γ1+γ2+β
β +γ2 β +τ1+γ2
γ1+γ2+τ1 571 τ1+γ2
τ1+γ2+δ 2 γ1+τ1
ε+ζ +γ1
γ1+ζ
γ1+ζ +τ1
τ1+δ 2
500
ε+γ1
γ1+γ2
ε+ζ α +δ 2
400
Cu 77.00 In 23.00 Sn 0.00
Landolt-Börnstein New Series IV/11C3
τ1+ζ
τ1 α +τ1+δ 2 10
Sn, at.%
ζ
ε+δ 1
ε+ζ +δ 1 20
Cu 77.00 0.00 In Sn 23.00
MSIT®
Cu–In–Sn
268
800
Fig. 14: Cu-In-Sn. Vertical section at 75 at.% Cu
L
Temperature, °C
700
L+γ1+γ2
L+β
L+γ1
L+γ2 L+β +γ2
β +γ2+δ 2 600
γ2
β +γ2
β +δ 2
γ1
γ1+γ2 γ1+γ2+τ1
τ1+γ2
τ1+ζ +γ2
α +δ 2
ε+ζ
δ 2+ζ
α +τ1+δ 2 Cu 75.00 In 25.00 Sn 0.00
ε+ζ +η
10
Cu 75.00 0.00 In Sn 25.00
20
Sn, at.%
700
Fig. 15: Cu-In-Sn. Vertical section at 70 at.% Cu
L
L+γ2
Temperature, °C
γ2
L+γ2+η
648
γ2+η γ1+γ2+η
600
L+γ1+γ2
γ1+η+ζ γ1+η+ε γ2+δ 2+η
562
583 L+ε
η+ζ
δ 2+η
L+ε+η
ε+η+ζ
δ 2+η+ζ δ2 400
L+γ1+ε
γ2+η+ζ
552
Cu 70.00 In 30.00 Sn 0.00
L+γ1
L+γ1+η
γ1+η
γ2+δ 2
500
MSIT®
ε
η+ζ
η+δ 2+ζ
τ1+δ 2
400
ε+ζ +γ1
552 ζ +γ2 509
γ2+δ 2+τ1
500
ε+γ1
γ1+ζ γ2+δ 2+ζ ζ +η 562
τ1+β +γ2
571
γ1+ζ +η
ε+η
10
20
Sn, at.%
Cu 70.00 0.00 In Sn 30.00
Landolt-Börnstein New Series IV/11C3
Cu–In–Sn
269
700
Fig. 16: Cu-In-Sn. Vertical section at 65 at.% Cu
L
L+γ2 L+γ2+η
Temperature, °C
L+η
L+γ1
L+γ1+η
L+γ1+ε
600
583
η
L+ε+η L+ε
500
η+ζ ε+η δ 2+η+ζ
400
η+ε+ζ
δ 2+η
Cu 65.00 In 35.00 Sn 0.00
Temperature, °C
Fig. 17: Cu-In-Sn. Partial vertical section at 60 at.% Cu
10
20
30
Sn, at.%
Cu 65.00 0.00 In Sn 35.00
700
L
L+ε+γ1
L+γ1
600
L+γ1+η
583
L+η L+ε
L+ε+η 500
η
400
Cu 60.00 In 40.00 Sn 0.00
Landolt-Börnstein New Series IV/11C3
10
ε+η
20
Sn, at.%
30
Cu 60.00 0.00 In Sn 40.00
MSIT®
Cu–In–Sn
270
Fig. 18: Cu-In-Sn. Partial vertical section at 4 at.% In
800
L+γ1+β
L+α +β
β
700
Temperature, °C
L+α
L+β
L
β +γ1
L+γ1
α +β
γ1 γ1+ε
L+γ1+ε
α
γ1+ε+ζ
600
α +β +γ1 ε
L+ε
α +γ1 γ1+ζ
500
540
α +γ1+ζ α +ζ +τ1
L+ε+η
400
Cu 56.00 4.00 In Sn 40.00
Fig. 19: Cu-In-Sn. Partial vertical section at 2 at.% Sn
60
α +τ1
α +ζ
70
80
Cu 96.00 4.00 In Sn 0.00
90
Cu, at.%
800
L
L+α
L+β +γ2
700
Temperature, °C
ζ
ε+ζ
ε+η
L+α +β
L+β
L+γ2+η L+γ2
γ2+η γ2
δ 2+γ2+η
β +γ2
α +β
β
α
β +δ 2
β +γ2+δ 2
600
β +α +δ 2 γ2+δ 2
L+η
β +τ1+δ 2
δ2
β +τ1+α
τ1+γ2+δ 2 500
η
δ 2+η
α +τ1
τ1+δ 2 α +τ1+δ 2
400
Cu 58.00 In 40.00 Sn 2.00
MSIT®
60
70
80
Cu, at.%
90
Cu 98.00 0.00 In Sn 2.00
Landolt-Börnstein New Series IV/11C3
Cu–In–Sn
Fig. 20: Cu-In-Sn. Vertical section at 10 mass% Cu, plotted in at.%
271
L
Temperature, °C
500
L+η
250
L+η +ϕ L+τ2
(In)+ϕ (In)+τ2+η
L+τ2+η
χ +η τ2+χ (In)+τ2
(In)+η (In)+η +ϕ 0
Cu 16.72 In 83.28 Sn 0.00
(In)+τ2+χ
τ2+χ +η χ +η+θ τ2+χ +θ
20
τ2+θ
η+(β Sn)
τ2+η+θ
40
η+θ +(β Sn)
η+θ
L+η +θ
60
η+(β Sn)+(α Sn)
Sn, at.%
Fig. 21: Cu-In-Sn. Vertical section at 30 mass% Cu, plotted in at.%
(β Sn) (α Sn)
80 Cu 17.19
0.00 In Sn 82.81
L
Temperature, °C
500
L+η
L+η +(β Sn) 250
L+η +ϕ
L+η +θ
L+(In)+η
L+ϕ (In)+ϕ
L+τ2+η
0
Cu 43.64 In 56.36 Sn 0.00
Landolt-Börnstein New Series IV/11C3
10
(In)+τ2+η
20
η+(β Sn)
η+θ
τ2+η+χ η+χ η+χ +θ τ2+η τ2+η+θ
(In)+η
(In)+η +ϕ
L+η+χ
30
Sn, at.%
η+(β Sn)+θ η+(α Sn)
(β Sn)
40
50
Cu 44.46 0.00 In Sn 55.54
MSIT®
Cu–In–Sn
272
Fig. 22a:Cu-In-Sn. Activity of In in liquid alloys along xCu/xSn = 4 section
1.0
727°C 827°C 927°C
0.9 0.8 0.7
Activity of In
0.6 0.5 0.4 0.3 0.2 0.1 0
0
20
40
60
80
100
80
100
In, at.%
Fig. 22b:Cu-In-Sn. Activity of In in liquid alloys along xCu/xSn = 1 section
1.0 0.9 0.8 0.7
727°C
Activity of In
0.6
827°C
927°C
0.5 0.4 0.3 0.2 0.1 0
0
20
40
60
In, at.%
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–In–Sn
273
1.0
Fig. 22c: Cu-In-Sn. Activity of In in liquid alloys along xCu/xSn = 0.25
0.9 0.8 0.7
Activity of In
0.6
727°C 827°C
0.5
927°C 0.4 0.3 0.2 0.1 0
0
20
40
60
80
100
In, at.%
Landolt-Börnstein New Series IV/11C3
MSIT®
274
Cu–Mn–Ni
Copper – Manganese – Nickel Andy Watson, Sigrid Wagner, Evgeniya Lysova and Lazar Rokhlin, updated by Andy Watson Introduction The earliest recorded experimental studies were made by [1912Par, 1913Par] who determined the liquidus and solidus surfaces of the Cu-Mn-Ni phase diagram using thermal analysis. The surfaces were constructed for the full concentration range, assuming the existence of a continuous solid solution between the three components immediately below solidus surface. Resulting from this study, isotherms of the liquidus and solidus surfaces were drawn. Six vertical sections of the liquidus and solidus surfaces were also presented. The liquidus and solidus temperatures of the Cu-Mn-Ni phase diagram from [1912Par, 1913Par] were used by [1934Fis] to construct a ternary phase diagram in orthogonal coordinates. In the review by [1949Jae], the data of [1912Par, 1913Par] were reproduced without significant alteration. The changes consisted of drawing a tentative boundary outlining the part of the liquidus surface corresponding to the crystallization of the ( Mn) rich phase. [1951Zwi] studied the Mn rich part of the Cu-Mn-Ni phase diagram constructing partial isothermal sections at 1000, 700 and 500°C. The sections showed phase regions corresponding to existence of solid solutions based on the different allotropic forms of Mn. However, in the reviews of [1969Gue, 1979Cha, 1986Gup], the isothermal sections constructed by [1951Zwi] were found to be inconsistent with the reliable Mn-Ni binary phase diagram. [1958Chj] reinvestigated the Cu-Mn-Ni phase diagram in the Cu rich region, up to 35 mass% Mn and 35 mass% Ni using X-ray diffraction, thermal analysis, microscopy and hardness measurement. [1958Chj] confirmed existence of the fcc solid solution immediately below the solidus surface. A partial vertical section Cu-MnNi and two partial isothermal sections at 350 and 450°C were constructed. The vertical section Cu-MnNi was found to be quasibinary on the Cu side. With increasing temperature, the ( Mn,Ni,Cu)+MnNi phase field tends to narrow. The existence of a continuous fcc solid solution immediately under the solidus surface in the Cu-Mn-Ni phase diagram was also reported by [1958Top]. To a certain extent, the existence of the quasibinary section Cu-MnNi, at least partially below the solidus surface, can be inferred from the results of [1962Gla]. In this study, microhardness measurement was conducted on alloys with compositions lying along sections of constant Cu contents of 95 and 90 at.%. The alloys had been annealed at 900, 700 and 500°C. Each of the microhardness curves showed a distinct minimum at the points where the Cu-MnNi section was crossed. Such features in the microhardness curves was interpreted by the authors [1962Gla] as evidence of Mn and Ni interaction in the (Cu) solid solution corresponding to the MnNi compound. [1970Rol] studied effect of heat treatment on structure of alloys in the section Cu-MnNi using X-ray diffraction. Decomposition of the fcc solid solution at temperatures of about 350-370°C was observed. The decomposition was accompanied by the formation of the low-temperature modification of the MnNi compound. This can be interpreted as decrease in the solubility of MnNi in the fcc solid solution with decreasing temperature. [1974Sch1, 1974Sch2] studied the equilibria between the liquid and solid phases in the Cu rich and Ni rich alloys of the Cu-Mn-Ni system at temperatures between the liquidus and solidus surfaces. The alloys were held isothermally at selected temperatures in the solid-liquid state and quenched into water. The compositions of the liquid and solid phases were then determined by local analysis of structure. In the review [1986Gup], data from [1974Sch1, 1974Sch2] were compared with those given in [1912Par, 1913Par]. It was found that the data from [1974Sch1, 1974Sch2] disagreed significantly from those of [1912Par, 1913Par] in the Ni corner of the phase diagram above 1300°C but were in agreement in the Cu corner of the phase diagram below 1100°C. In the opinion of [1986Gup], the data from [1974Sch1, 1974Sch2] were more reliable than those of [1912Par, 1913Par] because a better experimental method was used in the former work.
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Mn–Ni
275
In the review [1979Cha], the liquidus surface according to [1912Par, 1913Par] was accepted with a small amendment showing the existence of a narrow region relating to the primary crystallisation of ( Mn). The isothermal section of the phase diagram at 350°C was presented after [1958Chj]. In the review [1979Dri], the liquidus and solidus surfaces of the Cu-Mn-Ni phase diagram after [1913Par] were accepted where it was assumed that there was a continuous solid solution between copper, nickel and the modification of manganese. Meanwhile, after [1951Zwi], [1979Dri] reported a face-centred tetragonal lattice for ( Mn) and noted the existence of a narrow two-phase region of the ( Mn,Ni,Cu) face-centred cubic solution and the face-centred tetragonal ( Mn) solid solution near the Mn corner. In addition to this, [1979Dri] accepted from [1958Chj], the partial isothermal section of the phase diagram at 450°C (up to 50 mass% Mn and 50 mass% Ni) with the existence of a two-phase region comprising of the fcc ( Mn,Ni,Cu) solid solution and the low-temperature modification ' of the compound MnNi. As with other reviews [1979Cha, 1979Dri], the review [1980Bra] accepted the liquidus surface from [1913Par]. Also, a simplified view of the main phase fields in the system for solid state was presented. The fields corresponded to the homogeneous areas of the ( Mn,Ni,Cu) fcc solid solution, with the phases on the base compounds MnNi3, MnNi or one of the Mn modifications, except ( Mn). [1984Wea] presented the temperature-composition section Cu-MnNi of the phase diagram, but without description of the performed experiments and their results in detail. The section presented is reasonable, in general, taking into consideration the boundary binary systems, but on the MnNi side the exact position of the phase boundaries is not confirmed by any experimental work and, therefore, may be considered only as tentative. In the review [1986Gup] the liquidus and solidus surfaces show isotherms accepted from [1912Par, 1913Par] with the addition of the monovariant line bounding the surface of the primary crystallization of ( Mn) following [1979Cha]. The interaction between the fcc solid solution ( Mn,Ni,Cu) and the MnNi based phase was characterized by the vertical section Cu-MnNi, taken mainly from [1958Chj] and solubility isotherms in ( Mn,Ni,Cu) at 350 and 450°C after [1958Chj]. [1989Udo] studied phases formed in alloys along the section Cu-MnNi at temperatures between 450 and 600°C in the composition range 10-80 at.% Cu. The boundary between the phase regions of the ( Mn,Ni,Cu) solid solution and ( Mn,Ni,Cu)+MnNi was constructed. More recently, [2003Mie] has performed a Calphad assessment on the ternary system using all available experimental phase equilibria and thermodynamic data. Despite introducing some simplifications into the study, the agreement between the assessed and experimental data is good. Details of the investigations are given in Table 1. Binary Systems The binary systems are taken from the MSIT Evaluation Program [2002Leb, 2005Tur, 2006Wat]. Solid Phases No ternary compounds were found in the Cu-Mn-Ni system. The unary and binary phases, including the fcc continuous solid solution ( Mn,Ni,Cu), are listed in Table 2. Quasibinary Systems There is no true quasibinary section in this ternary system, although the Cu-MnNi section is considered to be a partial quasibinary system at temperatures below the solidus surface [1958Chj]. Liquidus, Solidus and Solvus Surfaces Figure 1 shows the liquidus projection together with isotherms. It is constructed following [1986Gup] who, in their review of the system, modified the data of [1912Par, 1913Par] and considered the double saturation line connecting the invariant points of the ( Mn) L + ( Mn,Cu) and ( Mn) + L ( Mn,Ni) reactions in the Cu-Mn and Mn-Ni binary systems, respectively. The double saturation line divides the liquidus surface into two parts corresponding to the primary crystallization of the ( Mn,Ni,Cu) and ( Mn) solid solutions. Landolt-Börnstein New Series IV/11C3
MSIT®
276
Cu–Mn–Ni
Figure 2 shows the projection of the solidus surface together with isotherms. It is constructed after [1986Gup] who used the data of [1912Par, 1913Par]. The projection is supplemented by lines connecting the critical points of the ( Mn) L + ( Mn,Cu) and ( Mn) + L ( Mn,Ni) transformations in the Cu-Mn and Mn-Ni binary systems. The lines correspond to the compositions of the ( Mn) and ( Mn,Ni,Cu) solid solutions in the monovariant ternary equilibrium with the liquid phase in the ternary system. The locations of these lines were not established experimentally and, therefore, the lines are shown dashed and are considered tentative. Isothermal Sections No full isothermal section of the Cu-Mn-Ni phase diagram has been constructed. [1958Chj] constructed the partial isotherms at 350 and 450°C of the solvus surface giving the solubility of the low-temperature modification of MnNi in the ( Mn,Ni,Cu) solid solution. The isotherms are shown in Fig. 3 after [1986Gup] who used results of the experimental investigation of [1958Chj]. Temperature – Composition Sections Figure 4 shows the Cu-MnNi partial vertical section of the Cu-Mn-Ni phase diagram. It is constructed after [1986Gup] who took into account a few consistent experimental studies. The section was corrected only slightly to be consistent with the solvus isotherms presented in Fig. 3 and the accepted version of the Mn-Ni binary phase diagram [2006Wat]. The version of the same section by [1984Wea] was not taken into account as it was thought to be unreliable. Actually, the section Cu-MnNi shown gives the extension of the ( Mn,Ni,Cu) solid solution from the Cu corner and misses more complicated phase relations near the Mn-Ni side which could be expected in view of the Mn-Ni binary phase diagram. Notes on Materials Properties and Applications Cu-Mn-Ni alloys are characterized, at certain compositions, by a low temperature coefficient of electrical resistivity combined with high values of electrical resistivity and are attractive for applications as resistor materials. Properties and structure of this kind of alloy are studied and described in [1941Ave1, 1941Ave2, 1949Sam, 1954Agl, 1961Die]. Other applications are highloaded springs working at high temperatures, non-magnetic bearings and gears and highloaded screws [1961Die, 1973Gro]. One of the advantages of using Cu-Mn-Ni alloys is their high corrosion resistance [1966Lah]. Investigations [1976Kho] have shown that the change in the thermoelectromotive force of Cu-Mn-Ni alloys has large limits depending on the Mn and Ni contents. In [1984Bey], electrical and magnetic properties of spin glasses of the Cu-Mn-Ni alloys are reported. [1989Don] studied the application of Cu-Mn-Ni alloys as thermally sensitive thermobimetals and calculated their linear thermal expansion coefficient over the full concentration range. The values obtained were presented in the form of isoproperty contours. [1990Wak] studied the application of the Cu-Mn-Ni alloys, eventually with small additions of other alloying elements, as dental materials. Cu-Mn alloys are well-known for their vibration and noise dampening properties. Investigations have been carried out to study the effects of adding Ni [2003Yin]. Ni was found to enlarge and broaden the characteristic twin-boundary damping peak. The aging behavior of Cu-Mn-Ni ternary alloys with 12 mass% Mn has been studied as a function of temperature and Ni content by [1991God], which was found to be due to vacancy and ordering reactions. Details of some experimental works related to the Cu-Mn-Ni materials properties are given in Table 3. Miscellaneous In alloys containing about 25 - 60 at.% Cu, balance Mn and Ni by approximately even parts, very complicated phase transformations are observed during annealing at temperatures between 250 - 500°C. This is a result of solid solution decomposition [1972Fil1, 1972Fil2, 1973Gro, 1978Ron, 1982Mik]. The complicated phase transformations in the solid state were also observed in alloys with the Ni3Mn composition with the addition of about 0.5-9.3 at.% Cu [1973Lot1, 1973Lot2, 1973Lot3, 1973Lot4, 1977Lot]. In [1975Bel, 1977Pop], the crystal structure and distribution of components present in vacuum MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Mn–Ni
277
deposited films have been studied. [1980Yok, 1987Tak] studied the interdiffusion in Cu rich alloys of the Mn and D ˜ Mn , and the cross interdiffusion Cu-Mn-Ni system. The direct interdiffusion coefficients, D˜ NiNi CuCu coefficients showed strong composition dependence. The Kirkendall effect was observed by [1980Yok] with the marker interface always moving towards the lower Ni concentration side. The deformation mechanism of single crystals of the Cu-MnNi section were studied by [1991Pol]. Using electron microscopy, XRD and neutron diffraction, they found that the strengthening mechanism was due to the drag of dislocations around the MnNi precipitate particles. References [1912Par]
[1913Par]
[1934Fis] [1941Ave1]
[1941Ave2]
[1949Jae] [1949Sam]
[1951Zwi]
[1954Agl]
[1958Chj]
[1958Top]
[1961Die] [1962Gla]
[1966Lah]
Landolt-Börnstein New Series IV/11C3
Parravano, N., “The Ternary Nickel - Manganese - Copper Alloys” (in Italian), Gazz. Chim. Ital., 42(11), 385-394 (1912) (Phase Diagram, Phase Relations, Experimental, *, 7) Parravano, N., “The Ternary Alloys of Iron - Nickel - Manganese, Nickel - Manganese Copper, Iron - Manganese - Copper Systems” (in German), Int. Z. Metallogr., 4, 171-202 (1913) (Phase Diagram, Review, *, 14) Fischer, V., “The Equilibrium Diagrams for the Ternary Systems” (in German), Z. Metallkd., 26(4), 80-82 (1934) (Phase Diagram, Phase Relations, Review, Theory, *, 8) Averbach, B.L., “Electrolytic Manganese Alloyed with Copper and Nickel – I”, Metals and Alloys, 13, 730 -733 (1941) (Experimental, Review, Electr. Prop., Mechan. Prop., 16) Averbach, B.L., “Electrolytic Manganese Alloyed with Copper and Nickel – I”, Metals and Alloys, 14, 47-51 (1941) (Crys. Structure, Experimental, Review, Electr. Prop., Mechan. Prop.) Jaenecke, E., “The Short Book about the Alloys” (in German), Heidelberg, Carl Winter, Universitaetsverlag, 369-370 (1941) (Phase Diagram, Phase Relations, Review, *, 1) Samans, C.H., Brayton, C.C., Drake, H.L., Litchfield, L., “Dilatometric Effects of Hardening and Recrystallization in the 60 Copper - 20 Nickel - 20 Manganese Alloy”, Trans. Am. Soc. Met., 41, 961-983 (1949) (Crys. Structure, Phase Relations, Experimental, Mechan. Prop., 14) Zwicker, U., “About Structure of Manganese rich Manganese - Copper -Nickel Alloys” (in German), Z. Metallkd., 42, 331-335 (1951) (Phase Relations, Crys. Structure, Experimental, Review, *, 10) Agladze, R.I., Mokhov, V.M., Topchiashvili, L.I., “About the Alloys of Manganese with Copper and Nickel” (in Russian), Splavy Margantsa s Med'yu, Nikelem i Tsinkom, Tbilisi, Akad Nauk Gruz., 51-65 (1954) (Experimental, Phase Relations, Electr. Prop., Mehcan. Prop., 5) Chjan, B.C., “The Study of Ternary Copper Alloys of Copper - Nickel - Manganese”, Izv. Vyss. Uchebn. Zaved. Tsvetn. Metall., (5), 107-115 (1958) (Phase Relations, Experimental, *, 13) Topchiashvili, L.I., “The Influence of Iron, Cobalt and Nickel on Structure and Properties of Manganese - Copper Alloys”, Zh. Neorg. Khim., 3(3), 726-727 (1958) (Phase Relations, Experimental) Dies, K., “The Manganese Bronze” (in German), Metall, 15, 1161-1172 (1961) (Experimental, Electr. Prop., Mechan. Prop., 8) Glazov, V.M., Stepanova, M.V., “Chemical Interaction between Nickel and Manganese in Ternary Copper - Based Solid Solution at Various Temperatures” (in Russian), Dokl. Acad. Nauk SSSR, 144(3), 565-568 (1962) (Phase Relations, Experimental, Mechan. Prop., 12) Lahiri, A.K., Mukherjee, K.P., Banerjee, T., “Studies and Properties of Some Ternary Copper Alloys of the Cu-Mn-Zn System”, Trans. Indian Inst. Met., 19, 141-146 (1966) (Experimental, Electrochemistry, Mechan. Prop., 10) MSIT®
278 [1969Gue]
[1970Rol]
[1972Fil1]
[1972Fil2]
[1973Gro]
[1973Lot1]
[1973Lot2]
[1973Lot3]
[1973Lot4]
[1974Sch1]
[1974Sch2]
[1975Bel]
[1976Kho]
[1977Lot]
[1977Pop]
MSIT®
Cu–Mn–Ni Guertler, W., Guertler, M., Anastasiadias, E., “Copper-Manganese-Nickel”, in A Compendium of Constitutional Ternary Diagrams of Metallic Systems, WADC Technical Report 58-615, Project No 7351, Wright-Patterson Air Forse Base, Ohio, 564-567 (1969) (Phase Diagram, 7) Rolland, J., Priester, P., Whitwham, D., “Decomposition of Pseudobinary Cu-NiMn Alloys” (in French), Compt. Rend. Acad. Sci. Paris, ser. C, 270C, 1777-1780 (1970) (Phase Diagram, Crys. Structure, Experimental, 5) Filip, D.P., Carciuleanu, E., “60Cu-20Ni-20Mn Alloy at the Critical Temperature” (in French), Phys. Status Solidi, A11(2), K131-K134 (1972) (Experimental, Electr. Prop., 10) Filip, D., Chiriak, M., Chiriak,H., “On Order - Disordered Phenomena in Cu3NiMn”, An. Sliint. Univ., Al. I. Cusa, Iasi, 18, 49-54 (1972) (Phase Relations, Experimental, Electr. Prop., Magn. Prop., 9) Groma, G., “A Special High Temperature Cu-Ni-Mn Spring Material”, in “Dimensioning and Strength Calculation”, Proc. 4th Conf. on Dimensioning, Acad. Kiado, Budapest, 417-421, (1973) (Experimental, 9) Lotkov, A.I., Panin, V.E., Kolubaev, A.V., Astashkin, I.N., Fadin, V.P., “Neutronographic Investigation of the Characteristics of the Phase Transition in the Alloy Ni3Mn, Alloyed with Small Additions of Aluminum and Copper”, Sov. Phys. J., (16), 1493-1500 (1973) (Crys. Structure, Experimental, Electr. Prop., 26) Lotkov, A.I., Panin, V.E., Kolubaev, A.V., Astashkin, I.N., Fadin, V.P., “Neutronographic Investigation of the Phase Transition Peculiarities in the Alloy Ni3Mg, Alloyed with Small Additions of Aluminium and Copper” (in Russian), Izvest. Vyss. Uchebn. Zaved., Fiz., (11), 16-25 (1973) (Crys. Structure, Experimental, Electr. Prop., 26) Lotkov, A.I., Panin, V.E., Fadin, V.P., Sarksyan, V.V., “The Nature of Distribution of Atoms, and its Connection with Electron Structure in a Number of Ternary Alloys on Ni3Mn Basis”, Sov. Phys. J., (16), 91-97 (1973) (Crys. Structure, Electron Structure, Experimental, 28) Lotkov, A.I., Panin, V.E., Fadin, V.P., Sarksyan, V.V., “The Character of Atom Distribution, and its Connection with Electron Structure in a Number of Ternary Alloys on Ni3Mn Base” (in Russian), Izvest. Vyss. Uchebn. Zaved., Fiz., (1), 117-126 (1973) (Crys. Structure, Experimental, 28) Schuermann, E., Prinz, B., “Equilibria in Melted Nickel Rich and Copper Rich Copper-Manganese-Nickel Alloys. I” (in German), Z. Metallkd., 65(8), 535-539 (1974) (Phase Diagram, Experimental, 7) Schuermann, E., Prinz, B., “Equilibria in Melted Nickel rich and Copper Rich Copper-Manganese-Nickel Alloys. II” (in German), Z. Metallkd., 65(9), 593-598 (1974) (Phase Diagram, Experimental, 0) Belous, M.V., Bochvar, N.R., Lysova, E.V., Popov, V.I., Popova, A.A., “Investigation of the Structure Formation of Films Obtained by the Vacuum Evaporation of Cu-Mn-Ni Alloys” (in Russian), Fiz. Khim. Obrab. Mater., (6), 66-68 (1975) (Crys. Structure, Experimental, 2) Kholmyansky, V.A., Piterskaya, L.V., “Investigations of Thermo-EMF and Electrical Resistance of Alloys in the Cu rich Corner of the Cu-Ni-Mn Ternary System” (in Russian), Nauchn. Tr. Nauchno-Issled. Proectn. Inst. Splavov Obrab. Tsvetn. Met., (51), 35- 42 (1976) (Experimental, Electr. Prop., 7) Lotkov, A.I., Panin, V.E., Fadin, V.P., “Kinetics of Ordering in Ni3Mn Alloy with Additions of Copper”, Sov. Phys. J., (20), 306-309 (1977) (Crys. Structure, Experimental, Electr. Prop., 8) Popov, V.I., “Structural Characteristics of Vacuum Condensates of Copper-Base Alloys”, Met. Sci. Heat Treat., 19 (3-4), 214-217 (1977) (Crys. Structure, Experimental, 13)
Landolt-Börnstein New Series IV/11C3
Cu–Mn–Ni [1978Ron]
[1979Cha]
[1979Dri]
[1980Bra]
[1980Yok]
[1982Mik]
[1984Bey]
[1984Wea]
[1986Gup]
[1987Tak]
[1989Don]
[1989Udo]
[1990Wak] [1991God]
[1991Gok]
[1991Pol]
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279
Rondot, D., Mignot, J., “Initial Stage of the Structural Transformation in the 60Cu-20Ni-20Mn Alloy” (in French), Acta Metall., 26(2), 217-222 (1978) (Crys. Structure, Experimental, Mechan. Prop., 7) Chang, Y.A., Neumann, J.P., Mikula, A., Golgberg, D., “Cu-Mn-Ni”, INCRA Monograph Series 6 Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Systems, NSRD, Washington, Vol. 6, 538-542 (1979) (Phase Diagram, Phase Relations, Review, #, 18) Drits, M.E., Bochvar, N.R., Gusei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Copper-Manganese-Nickel” in “Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 168-169 (1979) (Phase Diagram, Phase Relations, Crys. Structure, Review, 7) Brandes, E.A., Flint, R.F., “Mn-Cu-Ni”, in “Manganese Phase Diagrams”, Manganese Centre, Paris, France, 109 (1980) (Phase Diagram, Phase Relations, Crys. Structure, Review, 4) Yokota, M., Harada, R., Mitani, H., “Interdiffusion in the Ni-Mn-Cu Ternary Alloy System at 1173 K”, Trans. Jpn. Inst. Met., 21(9), 573-579 (1980) (Experimental, Phase Relations, Transport Phenomena, 15) Miki, M., Hori, S., “Grain Boundary Reaction in a Cu-Mn-Ni Alloy” (in Japan), J. Jpn. Inst. Met., 46(3), 301-307 (1982) (Crys. Structure, Interface Phenomena, Morphology, Experimental, 25) Beylin, V.M., Serebrenik, L.A., Fradkov, Y.A., “Electrical and Magnetic Properties of Spin Glasses Cu-Mn and Cu-Mn-Ni”, Phys. Met. Metallogr., 58(5), 66-73 (1984), (Experimental, Magn. Prop., Electr. Prop., 4) Weatherill, A.E., Buckley, R.A., “Phase Transformations and Ageing Phenomena in Copper-Nickel-Manganese Alloys”, Mater. Res. Soc. Symp. Proc., 21, 531-536 (1984) (Phase Relations, Experimental, Kinetics, Mechan. Prop., 16) Gupta, K.P., Rajendraprasad, S.B., Jena, A.K., Sharma, R.S., “The Copper-Manganese-Nickel System”, J. Alloy Phase Diagrams, 2(3), 198-204 (1986) (Phase Diagram, Crys. Structure, Review, #, 31) Takahashi, T., Katoh, M., Minamino, Y., Yamane, T., “Interdiffusion in the Solid Solution of Cu-Ni-Mn System” (in Japanese), J. Jpn. Inst. Met.(Nippon Kinzoku Gakkai Shi), 51(8), 701-709 (1987) (Experimental, Kinetics, 35) Donets, L.I., Tret'yakov, B.N., “Optimization of Thermally Sensitive Characteristics of Thermobimetals Using Simplex-Lattice Planning”, Sov. Mater. Sci. Rev., 3(1-4), 199-206 (1989) (Theory, 5) Udovenko, V.A., Sanadze, V.V., Polyakova, N.A., Chichua, E.D., Gogua, L.D., “Equilibrium State Diagram of Aged Alloys of a Quasi-Binary Section of the Cu-Ni-Mn System”, Soobshch. Acad. Nauk Gruz. SSR, 134(1), 73-76 (1989) (Phase Diagram, Experimental, 5) Wakasa, K., Yamaki, M., “Dental Application of the 30Ni-30Cu-40Mn”, J. Mater. Sci., 1(1), 44-48 (1990) (Experimental, Review, Mechan. Prop., 25) Goedecke T., “Physical Measurements on Copper-Manganese Alloys” (in German), Z. Metallkd., 82(3), 198-208 (1991) (Experimental, Phase Relations, Electr. Prop., Kinetics, Phys. Prop., 11) Gokcen, N.A., “The Mn-Ni (Manganese - Nickel) System”, J. Phase Equilib., 12(3), 313-321 (1991) (Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., Review, *, 47) Polyakova, N.A., Udovenko, V.A., Chichua, E.D., “The Structure, Deformation Mechanism and Strength Properties of Ageing Cu-Ni-Mn Alloys”, Phys. Met. Metall., 71(1), 171-179 (1991), translated from Fiz. Metal. Metalloved., 71(1), 178-187, (1991) (Crys. Structure, Experimental, Mechan. Prop., 9)
MSIT®
Cu–Mn–Ni
280 [2002Leb]
[2003Mie]
[2003Yin]
[2005Tur]
[2006Wat]
Lebrun, N., “Cu-Ni (Copper - Nickel)”, MSIT Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID:20.14832.1.20, (2002) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 51) Miettinen, J., “Thermodynamic Description of the Cu-Mn-Ni System at the Cu-Ni Side”, Calphad, 27(2), 147-152 (2003) (Assessment, Phase Diagram, Phase Relations, Thermodyn., 18) Yin, F., Nagai, K., Watanabe, K., Kawahara, K., “The Damping Behavior of Ni Added Mn-Cu Damping Alloys”, Mater. Trans., JIM, 44(9), 1671-1674 (2003) (Experimental, Mechan. Prop., 14) Turchanin, M., Agraval, P., “Cu-Mn (Copper - Manganese)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; Document ID: 20.14136.1.20, (2005) (Crys. Structure, Phase Diagram, Assessment, 20) Watson, A., Wagner, Z., Lysova, E., Rokhlin, L., “Mn-Ni (Manganese - Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2006) (Crys. Structure, Phase Diagram, Assessment, 12)
Table 1: Investigations of the Cu-Mn-Ni Phase Relations, Structures and Thermodynamics Reference
Experimental Technique
Temperature/Composition/Phase Range Studied
[1912Par]
Thermal analysis
Full composition range, liquidus and solidus
[1913Par]
Thermal analysis
Full composition range, liquidus and solidus
[1941Ave2]
XRD
Diffraction pattern of 2%Cu49%Mn49%Ni alloy quenched from 950°C
[1949Sam]
Dilatometry, electrical resistance
60Cu20Mn20Ni, up to 800°C
[1951Zwi]
X-ray and microstructural examination
Mn rich alloys at 500, 700 and 1000°C
[1958Chj]
X-ray diffraction, thermal analysis, microstructural examination and hardness measurement
Cu rich region, up to 35 mass% Mn and 35 mass%
[1958Top]
Microstructural examination, dilatometry, electrical resistivity and hardness measurement
Samples quenched from 950-850°C, tempered at 500, 400 and 300°C
[1962Gla]
Microhardness
Phase relations in alloys along sections of constant Cu of 95 and 90%, quenched from 500, 700 and 900°C
[1970Rol]
XRD
Alloys in the Cu-MnNi section
[1973Lot1] [1973Lot2] [1973Lot3] [1973Lot4]
Neutron diffraction, resistivity measurement
Ordering in Ni3Mn-Cu
MSIT®
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Cu–Mn–Ni
281
Reference
Experimental Technique
Temperature/Composition/Phase Range Studied
[1974Sch1] [1974Sch2]
Isothermal annealing and microprobe analysis
Liquidus and solidus
[1977Lot]
Neutron diffraction, resistivity measurement
Ordering in Ni3(Mn,Cu) with Cu from 0.45 to 9.29 at.%. Samples tempered at 900°C.
[1978Ron]
XRD, metallography, mechanical kinetics of ordering property measurement
[1982Mik]
XRD, optical and electron microscopy, hardness measurement
[1984Wea]
XRD, dilatometry, metallography, Cu-MnNi quasibinary section DTA, hardness measurement
[1989Udo]
XRD
Cu-MnNi quasibinary section, 400-800°C
[2003Mie]
Calphad assessment
Phases considered limited to liquid, ( Mn,Ni,Cu), , J and phases
Cu-20Mn-20Ni and Cu-30Mn-30Ni (mass%) aged at 250-450°C
Table 2: Crystallographic Data of Solid Phases Phases/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
( Mn,Ni,Cu) < 1455
cF4 Fm3m Cu
( Mn) 1246 - 1138 (Mn) 1100 - 586
cI2 Im3m W cP20 P4132 Mn
Lattice Parameters Comments/References [pm]
a = 361.46
continuous ternary solid solution [1979Cha, 1986Gup] pure Cu at 25°C [Mas2]
a = 386.26
pure Mn at 1097°C [1991Gok]
a = 352.40
pure Ni at 25°C [1991Gok]
a = 308.13
dissolves up to 6 at.% Ni and 12.5 at.% Cu pure Mn at 1137°C [1991Gok]
a = 631.52
dissolves up to 18 at.% Ni pure Mn at room temperature [1991Gok]
a = 636.96
at 15 at.% Ni, room temperature [1991Gok]
(Mn) < 727
cI58 I3m Mn
a = 891.39
dissolves up to 9 at.% Ni pure Mn, room temperature [1991Gok]
, MnNi 911 - 675
cP2 Pm3m CsCl
a = 297.7
45-52 at.% Ni at 50 at.% Ni, 750°C [1991Gok]
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Mn–Ni
282 Phases/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
', MnNi 775 - 620
tP4 P4/mmm AuCu
'', MnNi < 480
t** ?
', MnNi3 < 520
cP4 Pm3m AuCu3
1, Mn3Ni < 430
t*
a = 369.8 c = 369.2
[1991Gok]
J, Mn2Ni 720 - 560
-
-
31.5-35.5 at.% Ni [1991Gok]
, MnNi2 710 - 580
t**
-
64-68.5 at.% Ni [1991Gok]
', MnNi2 < 440
t**
-
66.7-73 at.% Ni [1991Gok]
Cu5Mn 400
-
-
at ~14-18 at.% Mn [Mas2]
Cu3Mn 450
-
-
at ~23-27 at.% Mn [Mas2]
a = 373.1 c = 363.2
47-55.5 at.% Ni at 50 at.% Ni, room temperature [1991Gok]
-
46-54 at.% Ni [1991Gok]
a = 358.9
71-85 at.% Ni at 75 at.% Ni, room temperature [1991Gok]
Table 3: Investigations of the Cu-Mn-Ni Materials Properties Reference
Method/Experimental Technique
Type of Property
[1941Ave1, 1941Ave2]
Not specified
Resistivity and EMF with respect to temperature. Workability of alloys at 500°C
[1949Sam]
Rockwell hardness measurement
Softening and hardening characteristics during heat treatment
[1954Agl]
Resistivity measurement
Electrical resistivity
[1961Die]
Resistivity, mechanical property measurement.
Temperature coefficient of resistivity, tensile strength and hardness.
[1962Gla]
Microhardness measurement
Microhardness of alloys along sections of constant Cu of 95 and 90%, quenched from 500, 700 and 900°C
[1966Lah]
Combination of stress and ammonia Stress-corrosion cracking properties. Tensile environment. Mechanical property strength, elongation and hardness. tests.
[1972Fil1]
Double bridge
MSIT®
Resistivity with respect to temperature
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283
Reference
Method/Experimental Technique
Type of Property
[1972Fil2]
Dilatometry, electrical and magnetic measurement
Electrical and magnetic properties, ordering.
[1975Bel]
XRD, vacuum deposition
Structure of thin films
[1976Kho]
Thermoelectromotive force measurement, resistivity
Electrical properties
[1977Pop]
XRD, vacuum deposition
Structure of thin films
[1978Ron]
Mechanical property measurement Tensile strength, hardness
[1980Yok]
Diffusion couples
[1984Bey]
Induction method, 4-point method, Dynamic magnetic susceptibility, resistivity, vibrational magnetometer magnetization.
[1984Wea]
Hardness measurement
Hardness of MnNi-Cu quasibinary alloys with aging time at 450°C
[1987Tak]
Diffusion couples
Interdiffusion in Cu rich alloys
[1990Wak]
Hardness measurement
Hardness of 30Cu-40Mn-30Ni (mass%) experimental dental alloy
[1991God]
Dilatometry, resistivity measurement
Aging behavior
[1991Pol]
XRD, EM, neutron diffraction
Structure and deformation mechanisms
[2003Yin]
DMA
Young’s modulus, damping capacity
Landolt-Börnstein New Series IV/11C3
Interdiffusion and cross interdiffusion
MSIT®
Cu–Mn–Ni
284
Cu
Data / Grid: at.% Axes: at.%
Fig. 1: Cu-Mn-Ni. Projection of the liquidus surface 20
40
80
min.
60
871 900°C 1000
60
40
(γ Mn,Ni,Cu)
1200 1100
e 80
20
1300
(δ Mn)
1100
1200°C
1400 min. 1020 p
Mn
20
40
60
80
Cu
Ni
Data / Grid: at.% Axes: at.%
Fig. 2: Cu-Mn-Ni. Projection of the solidus surface 20
40
80
min. 871
60
900°C 1000
60
40
(γ Mn,Ni,Cu)
80
1100
bβ
(δ Mn)
Mn
MSIT®
1200
bγ
aγ aδ
20
1300
min. 1020 20
40
60
80
1400
Ni
Landolt-Börnstein New Series IV/11C3
Cu–Mn–Ni
285
Cu
Data / Grid: at.% Axes: at.%
Fig. 3: Cu-Mn-Ni. Partial isotherms at 350 and 450°C of the (Cu, Mn,Ni) solvus surface
10
(γ Mn,Ni,Cu)
90
20
80
30
40
70
350
60
50
Temperature, °C
Fig. 4: Cu-Mn-Ni. Partial vertical section MnNi-Cu
(γ Mn,Ni,Cu)+η''
10
Mn 60.00 Ni 0.00 Cu 40.00
450°C
20
50
30
40
50
Mn 0.00 Ni 60.00 Cu 40.00
L 1000
750
(γMn,Ni,Cu)
(Cu,γMn,Ni)+? 500
480°C
(γMn,Ni,Cu)+η ''
Mn 50.00 Ni 50.00 Cu 0.00
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Cu
Cu, at.%
MSIT®
286
Cu–Mn–Sn
Copper – Manganese – Tin Nathalie Lebrun Introduction Sn based soldering alloys are good candidates for microelectronics applications and are widely used in establishing electrical contacts between metals, of which one is invariably copper. The Cu rich part of the ternary system Cu-Mn-Sn has been extensively studied and experimental data are reported in Table 1. Only one ternary phase MnCu4Sn, called -, has been observed and disagreements are noticed concerning its composition. There is also disagreement in the experimental observations of solubility of Mn in the phase of the Cu-Sn binary system. According to [1933Ver, 1955Ada], the phase disappears at Mn contents of about 5 to 10 mass%, respectively, whereas [1942Car, 1953Fun, 1953Val, 1973Mey, 1987Leo] show a much wider phase region, up to 30 or 40 mass% Mn. [1942Car] mentioned that phase is stable at only elevated temperature but it can be retained above 625°C by quenching. [1954Bla] established experimentally partial polythermal sections for 5 to 20 mass% Mn. It was shown that the higher is Mn concentration, the lower is solid solubility of Sn in (Cu), varying above 600°C from 15.8 to 5.8 mass% when Mn concentration increases up to 20 mass%. Polythermal sections at 85 and 75 mass% Cu, and at 2 mass% Mn have been constructed experimentally by [1987Leo] and [1985Dri, 1987Leo], respectively. The liquidus and solidus curves measured by [1990Boc] are in good agreement with those reported previously. From diffusion couple measurements, [2004Liu] calculated isopleths at 16.7 at.% Mn and 16.7 at.% Sn, which disagree with those reported previously. From thermodynamic description, [2004Mie] calculated isopleths which fit well with experimental data of [1942Car, 1953Fun, 1953Val, 1954Bla, 1987Leo]. Isothermal sections at 350 to 750°C have been reported by [1953Fun, 1987Leo, 2004Liu]. Very good agreement of the results of [1953Fun] and [1987Leo] has been observed, in contradiction with data [2004Liu] who observed a A2/B2 order-disorder transition. The isothermal sections of the latter work have also been calculated using a CALPHAD method and present a congruent melting of the ternary phase - in disagreement with experimental work [1987Leo] suggesting formation of - via a eutectoid reaction involving the phase. Through thermodynamic considerations, [2004Mie] calculated isothermal sections which are in good agreement with those reported in the literature. In the compilation of [1995Vil], a liquidus surface done by [1955Ada] has been reported. [2004Mie] calculated a partial liquidus surface in the Cu rich corner which is not in agreement with the one given by [1955Ada]. A complete thermodynamic description of the ternary system in the Cu rich part has been performed by [2004Mie] using ThermoCalc software. [1961Die, 1979Cha, 1979Dri] undertook a review of the system Cu-Mn-Sn. Binary Systems The binary system Cu-Mn is taken from [2005Tur]. No reliable data on Cu-Sn and Mn-Sn are available in the literature after [Mas2]. Only the Cu rich part of the Cu-Sn phase diagram was re-investigated by [2004Liu] which concluded the presence of a two stage, second order reaction A2 > B2 > D03 in the bcc phase region, rather than a two-phase equilibrium between the disordered A2 bcc phase and the ordered D03 bcc phase as reported before. The existence of this second order transition has been detected by DSC technique at 649°C on a Cu-16.9 at.% Sn. This result has been confirmed using high temperature electron diffraction technique showing a D03 ordered structure at 550°C and a B2 ordered structure at 675°C in a Cu-16.1 at.% Sn alloy. These experimental results have been supported by a CALPHAD method in which the Gibbs energy of the bcc phase is described by the two-sublattice model in order to take into account the MSIT®
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A2/B2 ordering reaction. No significant modifications have been reported by this previous study. Consequently, the phase diagram proposed by [Mas2] is accepted in this assessment. Mn-Sn phase diagram is taken from [Mas2]. Solid Phases All crystallographic data for the unary, binary and ternary phases are reported in Table 2. A ternary phase has been detected by several authors with different referring: ’ [1954Bla] and [1953Fun, 1987Leo]. Various composition have been attributed: MnCu5Sn2 [1933Ver], MnCu2Sn [1954Bla] and MnCu4Sn [1961Gla, 1968Joh, 1973Mey, 1987Leo, 2004Liu]. MnCu2Sn was investigated by [1954Bla, 1972Gel, 1981Uhl]. It was found that this alloy melts at around 657°C and its L21 structure exits down to 497°C. Two structure changes have been observed [1961Tag]. Using annealed samples, it was shown that the stable phase MnCu2Sn at room temperature is not cubic. [1942Car] suggested a hexagonal structure. The binary phase presents a large domain of existence inside the ternary system. This solid solution is ordered and stable at temperatures higher than 400°C. It can be retained at room temperature through quenching [1963Oxl]. In fact, its domain strongly depends on the annealing treatment of the alloys. Using quenched samples, [1942Car] stipulated that consists of a body-centered structure with a face-centered superlattice and the lattice parameter was estimated to be a = 616.08 pm at a composition of MnCu2Sn. This value agrees well with the one reported later in [1953Val, 1972Gel, 1977Mac]. [1973Mey] determined the lattice parameters inside the domain of the phase and found a small variation of the lattice parameter varying from a = 606.6 0.3 to 617.6 0.9 pm). The same author detected a lower temperature MnCu4Sn type structure in all the alloys, the lattice parameter was estimated to be a = 698.2 0.3 pm in agreement with the value of a = 698.8 0.3 given by [1961Gla]. The most reliable composition of the ternary compound - seems to be MnCu4Sn which has been retained in this assessment. [2004Liu] suggested that - presents a small solubility range. Invariant Equilibria Invariant equilibria and reaction scheme are presented in Table 3 and Fig. 1. The values of the U6, U7, U8 and E1 given in the table are estimated from the diagram given by [1955Ada] and should be considered as tentative. For the ternary reactions U4 and U5, their positions have been deduced from the calculated partial liquidus surface presented by [2004Mie]. The phase compositions in other ternary reactions were directly deduced from Fig. 2 and should be also considered as tentative. Liquidus, Solidus and Solvus Surfaces [1995Vil] mentioned the liquidus surface constructed by [1955Ada] who found a narrow region of the primary phase. Recently [2004Mie] undertook a thermodynamic description of the ternary system and reported a partial calculated liquidus surface in the Cu rich corner. There is a clear disagreement concerning the stability of the phase. According to [1955Ada], the phase disappears at Mn content of about 10 mass% whereas other studies [1942Car, 1953Fun, 1953Val, 1972Mey, 1987Leo] show a much wider phase region reaching up to 30 or 40 mass% Mn. These last experimental results were confirmed by the thermodynamic calculation done by [2004Mie]. However, the phase region is not allowed to extend to such high tin content [1942Car]. Consequently, the phase region has been estimated in this assessment on the base of the partial liquidus surface proposed by [2004Mie]. The sequence of the ternary reactions proposed in the Cu rich corner by [1955Ada] have not been retained in this assessment since recent thermodynamic evaluation suggested different four-phase equilibria which are more reliable. The monovariant curve proposed by [1955Ada] in the Mn rich corner is not acceptable as different solid phases participate in the equilibria. The equilibria have been changed and a sequence of two ternary reactions has been proposed. Moreover, the equilibrium L + MnSn2 + + (Sn) is more likely of eutectic type since the
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Cu–Mn–Sn
solubilities of Mn and Cu in (Sn) are negligible. Then instead of the U7 reaction proposed by [1955Ada], it was replaced by a E reaction. The accepted liquidus surface, in accordance with the binary systems, is presented in Fig. 2 with the accepted data of [2004Mie] in the Cu rich corner and the schematic liquidus projection of [1955Ada] in the Sn corner. Tentative liquidus and isothermal curves have been indicated by dashed lines on the drawing. The calculated isothermal curves from [2004Mie] at 1000 and 900°C reproduce well those of [1955Ada]. Taking into account Alkemade’s rule, position of some invariant points have been modified in Fig. 2. Isothermal Sections Several isothermal sections have been constructed in the Cu rich corner [1953Fun, 1987Leo, 2004Liu]. All the experimental works are in good agreement. However the isothermal sections suggested by [2004Liu] have not been retained in this assessment since some inconsistencies have been observed with the polythermal sections proposed by [1954Bla, 1987Leo]. Indeed, the latter works detected a maximum temperature of 610°C for decomposition of the - phase, whereas [2004Liu] suggested its congruent melting at about 680°C using CALPHAD calculation. [2004Mie] calculated few isothermal sections in the Cu rich corner which are in quite good agreement with the experimental ones. Some of them are reported in Figs. 3 to 7. The diagrams have been slightly changed in accordance with the accepted binary diagrams. Temperature – Composition Sections Partial polythermal sections have been reported in the literature [1954Bla, 1987Leo, 1990Boc], the experimental data are in good agreement. The phase decomposes eutectoidally into ( Mn,Cu) and the ternary phase - [1987Leo, 1954Bla]. The temperature of the eutectoid reaction varies considerably with the Mn content, and has a maximum at 610°C and 13 mass% Mn [1954Bla]. These results are in contradiction with the calculated partial isopleths [2004Liu] who proposed a congruent melting of - at around 680°C. Since no experimental data supported the calculations, this latter work has not been accepted in this assessment. [2004Mie] also calculated isopleths which agree reasonably well with the experimental data [1953Fun, 1985Dri, 1987Leo, 1990Boc]. Some inconsistencies were observed with the polythermal sections proposed by [1955Ada] who did not detect the ternary phase -. Some examples of isopleths are reported in Figs. 8 to 13, in accordance with the binary edges. In order to be in agreement with the isothermal sections at 450 and 550°C, slight modifications have been done in the isopleth at 15 mass% Mn (Fig. 11) since the equilibrium - + (Mn) + Mn3Sn exists. The modifications are indicated as dashed lines. Thermodynamics The enthalpies of formation in the region have been measured by [1972Mey] ranging from 383 to 933 J#mol–1 closely the composition MnCu2Sn. Notes on Materials Properties and Applications Brief details of some experimental works on materials properties are reported in Table 4. Certain Cu-Mn-Sn alloys were ferromagnetic especially when the binary phase was present. The estimated value of the magnetization ) to be about 60.8 Am2#kg–1 [1983Bus] and 68.7 Am2#kg–1 [1975Wac] at 530°C for a MnCu2Sn alloy. These values are in the same range order as those reported previously by [1953Val]. For the same composition, [1942Car] observed a maximum value of ) = 76.65 Am2 # kg–1 at 660°C. [1914Heu] also measured the magnetic saturation which is found to be maximal (107 Am2#kg–1) for a composition of Mn2Cu9Sn. The Curie temperature of MnCu2Sn alloy has been extrapolated at 357°C [1981Uhl, 1982Uhl] where a decreasing of the magnetization occurs. This extrapolated value of the Curie temperature is far from the one calculated by [1975Wac] who found a value of 240°C. The ferromagnetic moment has been estimated to be OO = 4.00 B for a quenching temperature of 645°C. This value is in agreement with those of [1942Car, 1963Oxl]. On molten Cu-Sn alloys with Mn MSIT®
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impurities (25 at.%), [1975Gue] found a linear variation of the reciprocal magnetic susceptibility with temperature, varying from 350 to 900 g-atom#cm–3 in the temperature range 400 to 1400°C. These values are lower than those reported by [1975Wac] for the same composition range. Tempering quenched Cu alloys with 5-20 mass% Mn and 4-14 mass% Sn, results in extreme brittleness, owing to unfavorable distribution of the - phase [1954Bla]. Cold working the alloys before tempering leads to more homogeneous texture, i.e. dispersed particles - in a (Cu) matrix. Thus tensile strengths of 35-38 tons/in2 with elongations of 10-20% can be obtained. The electrical properties of some Cu based alloys and thin evaporated films containing additions of Mn and Sn have been measured by [1990Boc]. The electrical resistivity has been found for the alloys greater than values for the films. The electrical resistivity increases regularly from about 40#108 Ohmm at Cu-25%Mn to 83#108 Ohmm at Cu-25%Sn [1984Boc]. Miscellaneous In the case of amorphous MnCu2Sn, the only annealing temperature which produced ferromagnetism was 200°C [1982Tay]. At all temperatures this phase was mixed with a major second phase Mn2Cu3Sn. The magnitude and sign of the hyperfine field in MnCu2Sn has been measured at –196°C by means of Mössbauer spectroscopy and found to be +200 10 kOe [1972Gel]. Low temperature specific heat measurements were performed on MnCu2Sn and show a soft phonon distribution leading to an Einstein temperature of about –220°C [1991Fra]. References [1914Heu] [1933Ver]
[1942Car]
[1953Fun]
[1953Val]
[1954Bla]
[1955Ada]
[1961Die] [1961Gla]
[1961Tag] [1963Oxl]
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Heusler, F., “Study of the Magnetisation of the Manganese Alloys”, Z. Anorg. Chem., 88, 185-188 (1914) (Experimental, Magn. Prop., 3) Veroe, J., “Equilibrium Conditions of Further Alloyed Bronses, III. Cu-Mn-Sn Alloys Rich in Cu”, Mitt. Berg. Huett. Abt. Ung. Hochschule Berg. Forstwesen Sopron, 5, 128-155 (1933) (Experimental, Phase Diagram, 16) Carapella, L.A., Hultgren, R., “The Ferromagnetic Nature of the Beta Phase in the Copper-Manganese-Tin System”, Trans. Amer. Inst. Min. Met Eng., 147, 232-244 (1942) (Crys. Structure, Experimental, Magn. Prop., Phase Diagram, 36) Funk, C.W., Rowland, J.A., “ Solid-Solution Area of the Cu-Mn-Sn System”, Trans. Am. Inst. Min. Metall. Pet. Eng., 197, 723-725 (1953) (Experimental, Phase Diagram, Phase Relations, 18) Valentiner, S., “Into Knowledge of the Copper-Manganese-Tin System” (in German), Z. Metallkd., 44, 59-64 (1953) (Crys. Structure, Experimental, Magn. Prop., Phase Relations, 11) Blade, J.C., Cuthbertson, J.W., “The Structure and Mechanical Properties of Copper-Manganese-Tin Alloys”, J. Inst. Met., 82, 17-24 (1954) (Mechan. Prop., Phase Diagram, 14) Adachi, M., Sasakai, Y., “Invesigation of the Ternary Equilibrium Diagram of Copper-Manganese-Tin System”, Trans. Mining Metall. Aluminium Assoc., 12, 591-594 (1955) quoted in [1995Vil] Dies, K., “Manganese Bronses”, Metall, 15(12), 1161-1172 (1961) (Review, Experimental, 8) Gladyshevskii, E.I., Kripyakevich, P.I., Teslyuk, M.Y., Zarechnyuk, O.S., Kuz’ma, Y.B., “Crystal Structures of Some Intermetallic Compounds”, Sov. Phys.-Crystallogr., 6, 207-208 (1961), translated from Kristallografiya, 1960, 6(2), 267 (Review, Crys. Structure, 11) Taglang, P., Fournier, “On the Transformations of the Alloy MnSnCu2”, J. Phys. Radium, 22, 295-297 (1961) (Experimental, 7) Oxley, D.P., Tebble, R.S., Williams, K.C., “Heusler Alloys”, J. Appl. Phys., 34, 1362-1364 (1963) (Crys. Structure, 13) MSIT®
290 [1968Joh] [1972Gel]
[1972Mey] [1973Mey]
[1975Gue]
[1975Wac]
[1977Mac] [1979Cha]
[1979Dri]
[1981Uhl] [1982Tay]
[1982Uhl] [1983Bus]
[1984Boc]
[1985Dri]
[1987Leo]
[1990Boc]
MSIT®
Cu–Mn–Sn Johnston, G., “Neutron Diffraction Investigation of Ternary Manganese Alloys”, At. Energy Aust., 11, 18-24 (1968) (Crys. Structure, 15) Geldart, D.J.W., Campbell, C.C.M., Pothier, P.J., Leiper, W., “Moessbauer Effect Study of the Hyperfine Fields at the Sn Sites in Cu2MnSn”, Canad. J. Phys., 50, 206-212 (1972) (Experimental, Crys. Structure, 20) Meyers, M.A., Hepworth, M.T., “The Enthalpies of Formation of Ferromagnetic Cu-Mn-Sn Alloys”, Met. Trans. AIME, 3(9), 2544-2547 (1972) (Experimental, Thermodyn., 2) Meyers, M.A., Ruud, C.O., Barrett, C.S., “Observations on the Ferromagnetic Phase of the Cu-Mn-Sn System”, J. Appl. Crystallogr., 6, 39-41 (1973) (Crys. Structure, Magn. Prop., 10) Günther, H.H., Wachtel, E., Gerold, V., “Magnetic Properties of Molten Cu-Sn Alloys with Mn Impurities” (in German), Physica, 80B, 473-491 (1975) (Experimental, Magn. Prop., 15) Wachtel, E., Kruse, H., “Magnetic Investigations in Copper-Tin-Manganese Alloys with 25 at.% Mn (Solid and Liquid State)”, Z. Metallkd., 66(7), 431-437 (1975) (Experimental, Magn. Prop., 3) MacDonald, R.A., Stager, C.V., “Magnetic Properties of (NixCu1-x)2MnSn”, Canad. J. Phys., 55(17), 1481-1484 (1977) (Experimental, Crys. Structure, Magn. Prop., 9) Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Mn-Sn”, INCRA Monograph Series 6 Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Systems, NSRD, Washington, 6, 549-557 (1979) (Crys. Structure, Phase Diagram, Review, 18) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, M.P., Rokhlin, L.L., Turkina, N.I., “Cu-Mn-Sn”, in “Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 169-170 (1979) (Phase Diagram, Review, 9) Uhl, E., “Magnetic Properties of New Heusler Alloys (Cu1–xCox)2MnSn”, J. Magn. Magn. Mater., 25, 221-227 (1981) (Magn. Prop., 34) Taylor, R.C., Tsuei, C.C., “Amorphous Heusler Alloys: The Effect of Annealing on Structure and Magnetic and Transport Properties”, Solid State Commun., 41(6), 503-506 (1982) (Crys. Structure, Magn. Prop., 11) Uhl, E, “Magnetism in Mixed Heusler Alloys (Ni1–xCux)2MnSn”, Monatsh. Chem., 113(3), 275-284 (1982) (Magn. Prop., 26) Buschow, K.H.J., van Engen, P.G., Jongebreur, R., “Magneto-Optical Properties of Metallic Ferromagnetic Materials”, J. Magn. Magn. Mater., 38, 1-22 (1983) (Magn. Prop., Optical Prop., 23) Bochvar, N.R., Vigdorovich, V.N., Leonova, N.P., Lysova, E.V., “Resistivity of Alloys of Cu-Mn-Ge, Cu-Mn-Sn and Cu-Ge-Sn Systems” (in Russian), Metalloved. Term. Obrab. Met., 2, 61-64 (1984) (Experimental, Electr. Prop., 9) Drits, M.E., Bochvar, N.R., Lysova, E.V., Leonova, N.P., “A Study of the Nature of Component Interaction in Copper-Based Alloys Containing Manganese, Tin, and Germanium”, in “Stable and Metastable Phase Equilibria in Metallic Systems”, Nauka, Moscow, 55-59 (1985) (Experimental, Phase Diagram, 5) Leonova, N.P., Bochvar, N.R., Lysova, E.V., “Phase Equilibria in Copper-Rich Cu-Mn-Sn Alloys”, Russ. Metall., 4, 204-206 (1987) translated from Izv. Akad. Nauk SSSR, Met., 4, 203-205 (1987) (Experimental, Phase Diagram, 5) Bochvar, N., Lysova, E.V., “Correlation between Phase Diagrams and Electrical Properties of Cu-Alloys and Films Deposited via Evaporation in Vacuum”, User Aspects of Phase Diagrams, Proceedings Conf., Petten, Netherlands June 1990; Hayes, F.H. (Ed.), UMIST, Manchester (UK), 180-184 (1990) (Electr. Prop., Experimental, Phase Diagram, Phase Relations, 5)
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[1991Fra]
[1993Gok] [1995Vil] [2004Liu]
[2004Mie]
[2005Tur]
291
Goedecke, T., “Physical Measurements on Copper-Manganese Alloys. III. Influence of the Quenching Temperature on the Electrical Resistance and on the Dilatation of the Alloys” (in German), Z. Metallkd., 81, 826-835 (1990) (Experimental, Phase Diagram, Phase Relations, 17) Fraga, G.L.F., Brandao, D.E., Sereni, J.G., “Specific Heat of X2MnSn (X = Co, Ni, Pd, Cu), X2MnIn (X = Ni, Pd), and Ni2MnSb Heusler Compounds”, J. Magn. Magn. Mater., 102(1/2), 199-207 (1991) (Experimental, 27) Gokcen, N.A., “The Cu-Mn (Copper-Manganese) System“, J. Phase Equilib., 14(1), 76-83 (1993) (Review, Phase Diagram, Crys. Structure, Thermodyn., 52) Villars, P., Prince, A., Okamoto, H., Handbook of Ternary Alloy Phase Diagrams, ASM Int., Materials Park, OH, Vol. 8, 9694-9707 (1995) (Phase Diagram) Liu, X.J., Wang, C.P., Ohnuma, I., Kainuma, R., Ishida, K., “Experimental Investigation and Thermodynamic Calculation of the Phase Equilibria in the Cu-Sn and Cu-Sn-Mn Systems”, Metal. Mat. Trans. A, 35A, 1641-1654 (2004) (Experimental, Calculation, Phase Diagram, 55) Miettinen, J., “Thermodynamic Description of the Cu-Mn-Sn System in the Copper-Rich Corner”, Calphad, 28, 71-77 (2004) (Assessment, Calculation, Phase Diagram, Thermodyn., 28) Turchanin, M., Agraval, P., Gröbner, J., Matusch, D., Turkevich, V., “Cu-Mn (Copper Manganese)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14136.1.20 (2005) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Assessment, 25)
Table 1: Investigations of the Cu-Mn-Sn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1933Ver]
Thermal and microscopic methods
400-1100°C / Cu3Sn-Mn with 5 to 15 mass% Mn
[1942Car]
X-ray diffraction
640 to 715°C / Cu-based alloys with 0-40 mass% Mn and 0-30 mass% Sn
[1953Fun]
Metallographic analysis
350 to 750°C / Cu rich corner with 0 to 20 mass% Mn and 0 to 25 mass% Sn
[1953Val]
X-ray diffraction
MnCu2Sn
[1954Bla]
Metallographic analysis, DTA
400 - 700°C / Cu alloys with 5-20 mass% Mn and 4-14 mass% Sn
[1961Tag]
X ray diffraction
MnCu2Sn
[1968Joh]
X-ray and neutron diffraction
MnCu2Sn
[1972Gel]
X ray diffraction
Annealed at 640°C / MnCu2Sn
[1972Mey]
Liquid-metal solution calorimetry
25°C / ~ MnCu2Sn
[1973Mey]
X-ray analysis
25°C after annealing at 600°C for 3 h / MnCu2Sn and 50.1-64.8 at.% Cu with 24.8-32.3 at.% Mn and 8.7-25.1 at.% Sn
[1977Mac]
X ray diffraction
MnCu2Sn
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Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1985Dri]
Metallography
400 - 1200°C / Cu rich based alloys with 2 mass% Mn and 2-10 mass% Sn
[1987Leo]
microhardness measurements, DTA, metallographic analysis,
300 to 1000°C / Cu-Mn-Sn up to 30 mass% Sn and 30 mass% Mn
[1990Boc]
DTA, microscopic and X-ray analysis
700 - 1100°C / 2 mass% Mn and 0-25 mass% Cu
[2004Liu]
Diffusion couple method, DSC, electron diffraction and high temperature X-ray diffraction techniques
550 - 700°C / Cu based alloys with 0-30 at.% Mn and 0-30 at.% Sn
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
( Mn,Cu)
cF4 Fm3m Cu
Lattice Parameters Comments/References [pm]
a = 362.63
[Mas2] Melting point [1993Gok] at x = 0.04 [V-C2]
a = 375
at x = 0.80 [V-C2]
( Mn) 1138 - 707
a = 386
at x = 1 [Mas2]
(Cu) < 1084.62
a = 361.46
at x = 0 [Mas2]
( Mn) 1246 - 1138
cI2 Im3m W
a = 308.0
[Mas2]
(Mn) 1087 - 707
cP20 P4132 Mn
a = 631.52
[Mas2]
(Mn) < 707
cI58 I43m Mn
a = 891.26
at 25°C [Mas2]
( Sn)
tI2 ? Sn
a = 370 c = 337
at 25°C, 9.0 GPa [Mas2]
(Sn) 231.9681 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
at 25°C [Mas2]
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Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
[Mas2]
, Cu17Sn3 798 - 586
cI2 Im3m W
a = 302.61
from 13.1 to 16.5 at.% Sn and A2 type structure [Mas2, 2004Liu]
Ȗ, Cu3Sn(h) ~755 - 520
cF16 Fm3m
a = 611.76 0.10
from 15.5 to 27.5 at.% Sn [2004Liu] from 16.5 to 27.9 [Mas2, V-C2] D03 type structure [2004Liu]
or F43m BiF3 or CuHg2Ti ’ phase ~755 - 574
hP8 ? Na3As
a = 428.1 c = 784.2
from 15.5 to 23 at.% Sn [2004Liu] B2 type structure [2004Liu] metastable, martensitic form of quenched
J, Cu3Sn(r) < 676
oC80 Cmcm Cu3Sn
a =552.9 0.8 b = 477.5 0.6 c = 432.3 0.5
24.5 to 25.9 at.% Sn and Cu3Sn [Mas2, V-C2]
, Cu10Sn3 640 - 582
hP26 P63 Cu10Sn3
a = 733.0 0.4 c = 786.4 0.5
20.3 to 22.5 at.% Sn and Cu10Sn3 at 603°C [Mas2, V-C2]
, Cu41Sn11 582 - ~350
cF416 F43m Cu41Sn11
a = 1796.46 0.06 20 to 21 at.% Sn and Cu41Sn11 [Mas2, V-C2]
, Cu6Sn5(h) 415 - 186
hP4 P63/mmc NiAs
a = 419.2 0.2 c = 503.7 0.2
’, Cu6Sn5(r) < 189
h**
1, MnCu5 410
c**
-
[1990Goe]
J2, MnCu3 450
c**
-
[1990Goe]
3 (Cu-Mn phase) 700
c**
-
[1990Goe]
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43.5 to 44.5 at.% Sn and [Mas2, V-C2]
44.8 to 45.5 at.% Sn, ordered form with superlattice based on [Mas2]
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Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
Mn3Sn < 984
hP8 P63/mmc Ni3Sn
a = 567 c = 453
[V-C2] Mn19Sn6 in [2004Mie]
Mn2Sn < 884
hP6 P63/mmc InNi2
a = 437.7 0.1 c = 550.7 0.1
[V-C2]
MnSn2 < 549
tI12 I4/mcm Al2Cu
a = 665.9 0.3 c = 544.7 0.3
[V-C2]
* -, MnCu4Sn
cF24 F43m AuBe5
a = 698.8 0.3
[1961Gla]
* MnCu2Sn(h) 657 - 497
cF16 Fm3m BiF3
a = 617.6 0.9
[1973Mey]
* MnCu2Sn(r) < 497
hP2 P6/mmc Mg
a = 285.2 c = 401.4
[1954Bla]
Table 3: Invariant Equilibria Reaction
L + ( Mn) (Mn) + ( Mn,Cu)
T [°C]
Phase
Composition (at.%) Cu
Mn
Sn
U1
L
3.80
79.80
16.40
L + (Mn) Mn3Sn + ( Mn,Cu) 984-900
U2
L
4.10
76.50
19.4
L + ( Mn,Cu) + Mn3Sn
798-750
U3
L
62.94
28.90
8.16
L + Mn3Sn + Mn2Sn
700-650
U4
L
49.89
28.90
21.21
L + Mn2Sn +
650-600
U5
L
50.91
22.31
26.78
L + Mn2Sn + J
~ 510
U6
L
39.84
17.44
42.72
L + J Mn2Sn +
415-400
U7
L
12.22
6.02
81.76
L + Mn2Sn + MnSn2
400-231
U8
L
9.22
4.42
86.36
L + (Sn) + MnSn2
< 227
E1
L
0.46
0.46
99.08
MSIT®
1029-1000
Type
Landolt-Börnstein New Series IV/11C3
Cu–Mn–Sn
295
Table 4: Investigations of the Cu-Mn-Sn Materials Properties Reference
Method/Experimental Technique
Type of Property
[1914Heu]
Magnetometer
Magnetic properties
[1942Car]
Magnetometer
Magnetic properties
[1953Val]
Magnetometer
Magnetic properties
[1954Bla]
Hardness measurements
Tensile properties
[1963Oxl]
Standard Sucksmith technique
Magnetic properties
[1975Gue]
Susceptibility measurements
Magnetic properties
[1975Wac]
Susceptibility measurements
Magnetic properties
[1981Uhl]
Faraday magnetometer
Magnetic properties
[1982Uhl]
Faraday pendulum magnetometer
Magnetization and susceptibility coefficients
[1983Bus]
Spectroscopic technique, PAR vibrating sample magnetometer
Magnetization coefficient, magneto-optical Kerr rotation
[1984Boc]
Resistivity measurements
Electrical properties
[1990Boc]
Auger spectroscopy and electrical resistivity Electrical resistivity and electrical measurements resistivity temperature coefficient
Landolt-Börnstein New Series IV/11C3
MSIT®
296
MSIT®
Cu-Mn
Cu-Sn
Cu-Mn-Sn
Mn-Sn
1097 p1 l+(δMn) (γMn,Cu) 1029-1000
L + (δMn) (βMn) + (γMn,Cu)
1029 p2 l+(δMn) (βMn)
U1
984 p3 l+(βMn) Mn3Sn
(δMn)+(βMn)+(γMn,Cu) 984-900 798 p5 l+(γMn,Cu) β
L + (βMn)Mn3Sn + (γMn,Cu) U2
798-750 L + (γMn,Cu) Mn3Sn + β U3 β+Mn3Sn+(γMn,Cu)
755 p6 l+β γ
l+γ ε
p7
415 p9 l+ε η
Mn3Sn+β+Mn2Sn
650-600 L + β γ + Mn2Sn U5 β+γ+Mn2Sn 510
L+γε+Mn2Sn
U4
549 p8 l+Mn2Si MnSn2
U6
γ+Mn2Sn+ε 415-400
L + ε η + Mn2Sn U7
η+Mn2Sn+ε 400-231
L + Mn2Sn η + MnSn2
U8
MnSn2+Mn2Sn+η 227 e2 l η+(βSn)
< 227 L MnSn2 + η + (βSn) E1
Landolt-Börnstein New Series IV/11C3
MnSn2+(βSn)+η Fig. 1: Cu-Mn-Sn. Reaction scheme
231 e1 l MnSn2+(βSn)
Cu–Mn–Sn
700-650 L+Mn3Sn Mn2Sn+β 640
884 p4 l+Mn3Sn Mn2Sn
(βMn)+Mn3Sn+(γMn,Cu)
Cu–Mn–Sn
297
Sn e1 E1
Fig. 2: Cu-Mn-Sn. Liquidus surface projection
MnSn2
Axes: at.%
η
p9
U8
p8
Data / Grid: at.%
(β Sn) e2
20
500
80
U7
600 40
650
60
ε
Mn2Sn 750 700
p4 60 800
(β Mn) p2
1200
U6
40
γ
U5
p3
80
p7
U2 900
Mn3Sn U3
U1 1100 1000
(δMn) 20
Mn
20
β
U4
p1
900
(γ Mn,Cu) 40
60
1000 80
Mn 0.00 Cu 50.00 Sn 50.00
Fig. 3: Cu-Mn-Sn. Isothermal section in the Cu rich part at 750°C
p6 p5
Cu
Data / Grid: at.% Axes: at.%
10
40
20
30
L
γ +L
30
20
β +L+γ
γ β +γ
β
β +L 40
10
(γ Mn,Cu)+β (γ Mn,Cu) Mn 50.00 Cu 50.00 Sn 0.00 Landolt-Börnstein New Series IV/11C3
60
70
80
90
Cu
MSIT®
Cu–Mn–Sn
298
Mn 0.00 Cu 50.00 Sn 50.00
Fig. 4: Cu-Mn-Sn. Isothermal section in the Cu rich part at 700°C
Data / Grid: at.% Axes: at.%
10
40
L 20
30
L+γ
γ
30
L+ β
L+β
20
+γ
β +γ
β
40
β +(γ Mn,Cu)
10
(γ Mn,Cu) 60
Mn 50.00 Cu 50.00 Sn 0.00
70
80
90
Mn 0.00 Cu 50.00 Sn 50.00
Fig. 5: Cu-Mn-Sn. Isothermal section in the Cu rich part at 650°C
Data / Grid: at.% Axes: at.%
10
40
L+γ
20
30
ε
L+Mn2Sn+β
L+β +γ
γ +ε γ
30
L+β
Mn2Sn+β
Cu
20
β +γ β
40
10
(γ Mn,Cu)+β
γ 3+(γ Mn,Cu) Mn 50.00 Cu 50.00 Sn 0.00
MSIT®
60
70
80
γ3
(γ Mn,Cu) 90
Cu
Landolt-Börnstein New Series IV/11C3
Cu–Mn–Sn
299
Mn 0.00 Cu 50.00 Sn 50.00
Fig. 6: Cu-Mn-Sn. Calculated isothermal section in the Cu rich part at 550°C
Data / Grid: at.% Axes: at.%
10
40
20
30
τ +γ +Mn2Sn 30
δ
τ +Mn3Sn τ +Mn3Sn+(β Mn) τ +(β Mn)
?
τ
20
τ +γ
β +γ +τ (γM n,C u)+ β+ τ (γ Mn,Cu)+τ
(β Mn)+τ +(γ Mn,Cu)
40
δ+γ +ε
τ +Mn3Sn+Mn2Sn
γ +δ γ (γ Mn,Cu)+γ +β (γ Mn,Cu)+γ
β
10
(γ Mn,Cu)+β (γ Mn,Cu) 60
Mn 50.00 Cu 50.00 Sn 0.00
70
80
γ 3+(Cu,γ Mn)
Mn 0.00 Cu 50.00 Sn 50.00
Fig. 7: Cu-Mn-Sn. Calculated isothermal section in the Cu rich part at 450°C
Axes: at.%
40
20
τ +Mn2Sn+Mn3Sn
τ +Mn3Sn τ +Mn3Sn+(β Mn) 40
Mn 50.00 Cu 50.00 Sn 0.00 Landolt-Börnstein New Series IV/11C3
30
?
δ
20
τ
τ +(γ Mn,Cu)+δ
(β Mn)+τ (γ Mn,Cu)+(β Mn)+τ
(γ Mn,Cu)+(β Mn) (αMn)+(β Mn)+(γ Mn,Cu) (αMn)+(γ Mn,Cu) 60
Cu
Data / Grid: at.%
10
30
90
(γ Mn,Cu)+δ 10
τ +(γ Mn,Cu)
70 γ 3+(αMn) γ 3+(αMn)+(γ Mn,Cu)
(γ Mn,Cu) 80
γ3
90
Cu
MSIT®
Cu–Mn–Sn
300
Fig. 8: Cu-Mn-Sn. Isopleth at 2 mass% Mn in the Cu rich corner, plotted in at.%
1100
1000
L (γMn,Cu)+L
Temperature, °C
900
800
L+β +γ
β +L 700
(γMn,Cu)
L+γ
β
(γMn,Cu)+β
γ
β +γ
600
(γMn,Cu)+β +γ 500
400
(γMn,Cu)+τ
(γMn,Cu)+τ+δ
Mn 2.31 Cu 97.69 Sn 0.00
Fig. 9: Cu-Mn-Sn. Isopleth at 5 mass% Mn in the Cu rich corner, plotted in at.%
γ+τ+δ
10
τ+δ +ε
γ+ε+Mn2Sn
Sn, at.%
1000
L (γMn,Cu)+L
800
β +γ+L β +L 700
(γMn,Cu) (γMn,Cu)+β
γ+L
β β +γ
γ
600
(γMn,Cu)+β +τ (γMn,Cu)+γ+τ
500
(γMn,Cu)+τ 400
Mn 5.74 Cu 94.26 Sn 0.00
MSIT®
Mn 2.83 Cu 70.97 Sn 26.20
20
1100
900
Temperature, °C
γ+τ
(γMn,Cu)+γ
(γMn,Cu)+δ
γ+ε γ+τ+ε
(γMn,Cu)+τ+δ 10
γ+τ
β +γ+τ
γ+τ+ε τ+δ +ε 20
Sn, at.%
γ+ε
γ+τ+ε γ+ε+Mn2Sn
Mn 7.04 Cu 66.91 Sn 26.05
Landolt-Börnstein New Series IV/11C3
Cu–Mn–Sn
Fig. 10: Cu-Mn-Sn. Isopleth at 10 mass% Mn in the Cu rich corner, plotted in at.%
301
1100
1000
L
Temperature, °C
900
(γMn,Cu)+L 800
700
600
γ3
β +Mn2Sn+γ
(γMn,Cu)+β
β +γ
β
γ+L β +τ
γ3+(γMn,Cu)
(γMn,Cu)+τ
γ+τ+ε
(Cu,γMn)+τ+δ
τ+ε+Mn2Sn τ+δ +ε
Mn 11.39 Cu 88.61 Sn 0.00
10
γ+L+ +Mn2Sn γ+Mn2Sn
Mn 13.94 Cu 60.25 Sn 25.80
20
Sn, at.%
1100
1000
L
900
Temperature, °C
γ+τ+Mn2Sn
(Cu,γMn)+γ+τ
400
(γMn,Cu)+L
800
700
600
γ3
500
400
Mn 16.95 Cu 83.05 Sn 0.00
Landolt-Börnstein New Series IV/11C3
β +γ+τ γ+τ
(Cu,γMn)+β +τ
500
Fig. 11: Cu-Mn-Sn. Isopleth at 15 mass% Mn in the Cu rich corner, plotted in at.%
β +γ+L
β +L (γMn,Cu)
L+β (γMn,Cu)
(γMn,Cu)+β
β
τ+β+Mn3Sn β+τ
γ3+(γMn,Cu) (γMn,Cu)+β+τ (γMn,Cu)+τ
5
(γMn,Cu)+τ+(βMn) 10
Sn, at.%
τ+Mn3Sn+ +Mn2Sn
τ+γ+Mn2Sn τ+Mn3Sn τ+(βMn) Mn 19.63 τ+(βMn)+Mn 3Sn Cu 62.21 Sn 18.16
MSIT®
Cu–Mn–Sn
302
Fig. 12: Cu-Mn-Sn. Isopleth at 85 mass% Cu in the Cu rich corner, plotted in at.%
1100
L 1000
Temperature, °C
900
(γMn,Cu)+L
L+(γMn,Cu)+β
800
(γMn,Cu)+β 700
(γMn,Cu) 600
γ3
(γMn,Cu)+γ3
(γMn,Cu)+τ
500
Mn 16.95 Cu 83.05 Sn 0.00
Fig. 13: Cu-Mn-Sn. Isopleth at 75 mass% Cu in the Cu rich corner, plotted in at.%
Mn 0.00 Cu 91.37 Sn 8.63
5
Sn, at.%
1100
1000
L
Temperature, °C
900
(γMn,Cu)+L 800
L+β L+(γMn,Cu)+β
700
(γMn,Cu)
β
(γMn,Cu)+β
β+γ
600
(γMn,Cu)+β+τ (γMn,Cu)+γ
(γMn,Cu)+τ 500
Mn 27.83 Cu 72.17 Sn 0.00
MSIT®
5
(γMn,Cu)+δ 15Mn 0.00 (γMn,Cu)+γ+τ Cu 84.86 Sn 15.14
10
Sn, at.%
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
303
Copper – Nickel – Tin Gautam Ghosh Introduction The phase stability and phase equilibria of the Cu-Ni-Sn system are of significant scientific and technological interests due to following reasons: (i) the occurrence of spinodal decomposition followed by precipitation of ordered phases in Cu rich alloys [1974Sch1, 1974Sch2, 1979Bab, 1980Dit, 1983Mik1, 1984Mik1, 1986Kat, 1987Kra, 1988Sat, 1990Gou, 1994Zha, 1996Oel, 1998Zha1, 1998Zha2], (ii) the occurrence of discontinuous precipitation in Cu- and Ni rich alloys [1982Mik, 1983Mik2, 1984Mik2, 1990Mik, 1991Mik, 1994Mik], (iii) the occurrence of massive and martensitic transformations in 1-(Ni,Cu)3Sn alloys [1981Wat, 1983Mur, 1989Pak, 1990Pak, 1996Oel], (iv) there is a need to design and process Cu rich alloys with high-strength and high-conductivity properties for electrical components and electronic packaging [1982Sak, 1982Ste, 1986Zie], and (v) the usage of lead-free solders [2002Kat, 2002Lin, 2002Zen, 2003Che1] for electronic packaging. As a result, a large number of experimental studies have been carried out to determine the phase equilibria, and the results have been reviewed several times [1936Hum, 1948Ray, 1949Jae, 1969Gue, 1973Lev, 1979Cha, 1979Dri, 1988Gup, 2000Gup]. A summary of experimental studies of phase equilibria is given in Table 1. In earlier investigations [1928Pri, 1932Gui, 1932Ves, 1933Eas], the primarily interest was the solubilities of Ni and Sn in (Cu). A large number of alloys and a variety of experimental techniques, such as hardness [1928Pri, 1932Gui, 1932Ves], metallography [1928Pri, 1932Gui, 1932Ves, 1933Eas], X-ray diffraction (XRD) [1933Eas] and thermal analysis [1932Ves, 1933Eas] were employed to establish the solubility limits in (Cu). [1932Ves] prepared 61 ternary alloys using elements of following purity: 99.93% Cu, Mond’s Ni (containing 0.21% Fe, 0.05% Cu, 0.27% C) and 99.96% Sn. They reported six vertical sections and three partial isothermal sections. [1933Eas] prepared 92 ternary alloys using 99.96% Cu, but the purities of Ni and Sn were not mentioned. They [1933Eas] reported five vertical sections and five partial isothermal sections. The results of [1932Ves] and [1933Eas] show a good agreement with each other. [1940Glu] investigated the interfacial microstructure in diffusion couples made of Cu-Ni alloys and liquid Sn that were reacted at 630°C for up to 8 h. They observed Cu-Sn and Ni-Sn intermetallics in all diffusion couples. Mazzoleni [1953Maz1, 1953Maz2] carried out structural investigation of several ternary alloys, in the composition range of 6 to 38.5 at.% Ni and 36.1 to 58.7 at.% Sn, by XRD. While many of these alloys were found to have NiAs-type structure, presence of a Heusler phase (NiCu2Sn) was also established. [1971Bas] reinvestigated the phase equilibria of Cu corner, and reported the partitioning ratio of Ni and Sn between (Cu) and liquid, the liquidus and solidus isotherms, six vertical sections and four partial isothermal isothermal sections. They prepared a large number of ternary alloys using 99.994% Cu, 99.95% Ni and 99.99% Sn. The alloys were equilibrated at various temperatures, but in a solid-liquid two-phase field, and quenched in iced water. The partitioning ratio of Ni and Sn, and the tie lines between (Cu) and liquid phases were measured by electron-probe microanalysis (EPMA). [1984Wac] carried out magnetothermal analysis of Cu- and Ni rich alloys, and reported a partial liquidus surface, a partial isothermal section, and the Cu3Sn-Ni3Sn isopleth. The phase equilibria of CuxNi75–xSn25 alloys, with 0 x 30, have been studied extensively [1981Wat, 1989Pak, 1990Pak, 1996Oel, 1996Rud] using XRD, transmission electron microscopy (TEM), analytical electron microscopy (AEM), and differential thermal analysis (DTA). The vertical section along Ni2Cu-(Ni2Cu)1–xSnx section was investigated by [1995Bur] to determine the liquidus and solidus using DTA, while the solvus was determined using resistivity and XRD [1996Spr] and magnetometry [1997Spr]. Recently, Chen and co-workers [2002Lin, 2003Wan] have determined the phase equilibria using XRD, DTA and EPMA. They reported the liquidus surface [2002Lin], and isothermal sections at 800 [2003Wan], 240 [2002Lin] and 200°C [2003Che1]. [2002Lin] prepared 28 ternary alloys, containing up to 80 at.% Cu and 75 at.% Ni, by arc melting in an argon atmosphere using 99.98% Cu, 99.98% Ni and 99.95% Sn. After arc melting, the alloys were annealed at 300°C for 336 h and then at 240°C for 1695 h in a vacuum of 8 #
Landolt-Börnstein New Series IV/11C3
MSIT®
304
Cu–Ni–Sn
10–3 torr. The isothermal section at 240°C was constructed based on the results of EPMA and XRD. To establish the liquidus surface, heating/cooling experiments were carried out at 5°C#min–1. [2003Wan] prepared 54 ternary alloys by arc melting followed by annealing at 800°C for 30 d. They also prepared three types of diffusion couples, Sn-55 at.% Cu / Ni, Sn-65 at.% Cu / Ni and Sn-75 at.% Cu / Ni, that were reacted at 800°C for up to 20 min. The isothermal section at 800°C was constructed based on the composition analysis of phases by EPMA in bulk samples, and also in diffusion couples. [2003Che1] reported an isothermal section at 200°C based on the observed phases in diffusion couples. [2005Li] determined the phase relations of Sn corner at 240°C, with an emphasis on the phase equilibria involving (Sn), Cu6Sn5 and Ni3Sn4. More recently, [2005Li, 2006Li] reinvestigated the phase relations of Sn corner at 240°C, with an emphasis on the phase equilibria involving (Sn), Cu6Sn5 and Ni3Sn4. They prepared nine alloys, containing 10-25 at.% Cu and 10-30 at.% Ni, by arc melting in an argon atmosphere. Elements of following purity were used: 99.9% Cu, 99.99% Ni, and 99.999% Sn. The alloys were sealed in evacuated quartz tubes and heat treated at 240°C between 1222 to 1700 h. The phases were characterized by XRD and EPMA. The heat of mixing of liquid alloys was measured by [1979Poo1] and [2004Lue], and the atomic transport kinetics of Ni and Sn in (Cu) was determined by [1972Bas] and [1986Tak]. Thermodynamic modeling of phase equilibria within CALPHAD formalism was reported by [2003Mie]. Binary Systems The Cu-Ni binary system is accepted from [2002Leb]. The Cu-Sn binary system is accepted from the assessment of [1990Sau] which was later supplemented by thermodynamic modelling [1996Shi, 2000Moo, 2004Liu1] within CALPHAD formalism. In the assessment [1990Sau] and thermodynamic modelling [1996Shi, 2000Moo] of Cu-Sn system, the 1/ 1 phase relation is treated as a first-order phase transformation. In contrast, [2004Liu1] reported that a two-stage ordering transition 1B2 (CsCl type) 1 takes place in Cu rich alloys. They used diffusion couples in conjunction with high-temperature electron and X-ray diffraction techniques to study the two-stage ordering transitions. However, in the thermodynamic modelling of Cu-Sn system, [2004Liu1] did not consider the 1 phase. Since the two-stage ordering transition is yet to be verified by others, in this assessment the 1/ 1 phase relation is treated as a first-order phase transformation. The phase of Cu-Sn system is believed to be isotypic with B81-NiAs phase [1990Sau]; however, X-ray and electron diffraction studies have shown that both and ´ phases have monoclinic symmetry [1994Lar, 1995Lar]. The Ni-Sn binary system is accepted from [1985Nas] along with the thermodynamic modeling of [1999Gho, 2004Liu2]. Solid Phases The addition of Ni at a constant Sn content in Cu rich solid solutions decreases lattice parameter while that of Sn at a constant Ni content increases lattice parameter [1995Pal]. The formation of two transient ordered phases, D022 [1977Hel, 1979Bab, 1980Dit, 1983Mik1, 1984Kra, 1984Mik1, 1987Kra, 1994Zha, 1998Zha1, 1998Zha2] and L12 [1982Ray, 1983Mik1, 1994Zha, 1998Zha1], during aging of supersaturated solid solution is well established. An important fundamental issue that has been debated is whether chemical ordering, leading to D022 and/or L12 phases, precedes spinodal decomposition. Extensive and systematic transmission electron microscopy study [1998Zha1] showed that spinodal decomposition takes place prior to chemical ordering in a Cu-15Ni-8Sn (mass%) alloy. [1998Zha1] also showed that following spinodal decomposition, D022 ordering takes place prior to L12 ordering, and both L12 and D022 may co-exist in a Cu-15Ni-8Sn (mass%) alloy. However, only L12 phase has also been observed in a Cu-9Ni-6Sn (mass%) alloy aged at 350°C [1982Ray], and in a Cu-20Ni-8Sn (mass%) alloy aged at 400°C [1983Mik1, 1984Mik1] and only D022 phase has been observed after aging at 450°C [1983Mik1, 1984Mik1]. All these results suggest that the formation sequence and co-existence of L12 and D022 phases depend on both the alloy composition and aging temperature. 1-Cu3Sn dissolves substantial amount (up to 30 at.% Ni) of Ni [1937Rah, 1984Wac]. Dissolution of Ni in
-Cu41Sn11 phase increases its thermal stability [1953Maz2, 1977Boo] and lattice parameter [1977Boo]. In
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
305
both studies, NiCu9Sn3 alloy was quenched from 800°C yet its structure is found to be the same as
-Cu41Sn11. Experimental studies show that 1-Ni3Sn (D03) dissolves at least 25 at.% Cu above 977°C [1932Ves, 1981Wat, 1984Wac, 1989Pak, 1990Pak, 1996Oel]. Dissolution of Cu in 1-Ni3Sn enhances its thermal stability to temperatures well below the binary limit [1981Wat, 1989Pak, 1990Pak, 1996Oel], and also causes a decrease in lattice parameter [1981Wat]. At 600°C, 2-Ni3Sn (D019) dissolves about 5 at.% Cu with an increase of both a- and c-lattice constants [1981Wat]. Ni3Sn2 dissolves about 24 at.% Cu at 240°C [2002Lin] and 30 at.% Cu at 800°C [2003Wan]. In the temperature range of 125 to 240°C, Ni3Sn4 dissolves up to 7 at.% Cu [2001Che, 2001Gho, 2002Lin, 2003Che1, 2004Gho], while other studies report slightly higher values of 9 at.% at 240°C [2005Li, 2006Hsi, 2006Li], 9.4 at.% at 245°C [2005Jan], and 10.28 at.% at 265°C [2004Jan, 2005Jan]. In the temperature range of 125 to 240°C, Cu6Sn5 dissolves about 24 at.% Ni [2001Gho, 2002Lin, 2003Che1, 2004Gho, 2005Li, 2006Hsi, 2006Li]. At 265°C, Cu6Sn5 may dissolve up to 26.2 at.% Ni [2005Jan]. Three ternary phases, -1, -2 and -3, have been reported. The Heusler phase NiCu2Sn (-1) was first reported by [1953Maz2], and subsequently confirmed by [1979Sch, 1984Wac]. However, a detailed study of NiCu2Sn using DTA, X-ray diffraction and resistivity techniques show that the L21 structure is stable between 500 and 700°C [1979Sch]. It has been reported that between 245 and 400°C, the Heusler phase (-1) decomposes into two hcp structures (hcp1: a = 412 pm, c = 1260 pm; hcp2: a = 414 pm, c = 1240 pm) and an fcc structure (a = 363 pm) [1979Sch]. The Ni2CuSn alloy was also suspected to be a Heusler phase; however, Mössbauer spectrum (119Sn) shows an unusual splitting that is inconsistent with the cubic symmetry. This splitting is explained on the assumption of disordering in Ni and Sn sublattices, so that it is effectively a D03 structure [1971Dok]. Furthermore, transmission electron microscopy and X-ray diffraction investigations [1989Pak, 1990Pak] failed to make a distinction between D03 and L21 structures in Ni2CuSn. The ternary phase -2 (Ni5–xCuxSn2) with 0.8 x 1.41), isotypic with -Cu3Ti, was first reported by [1981Wat], and subsequently confirmed by others [1983Mur, 1984Mik1, 1989Pak, 1990Pak, 1996Oel, 1996Rud]. The stability of this phase clearly underscores the influence of Cu, as the binary Ni3Sn phase with -Cu3Ti structure is metastable [1973Pak]. Massive [1981Wat, 1983Mur] and martensitic transformations [1981Wat, 1983Mur, 1989Pak, 1990Pak, 1996Oel] in CuxNi75–xSn25 alloys with 0 x 30 have been studied by a variety of techniques, such as optical metallography, DTA, XRD and TEM. Slow cooling leads to massive ( 1 (D03) -2 (2H)) while fast cooling leads to martensitic ( 1 (D03) -2 (2H)) transformation [1983Mur]. Upon furnace cooling of alloys containing 11 to 22 at.% Cu from 1000°C, [1990Pak] reported another martensitic transformation 1 (D03) -3 (triclinic) where the product phase may be considered as a distorted structure of -2. The crystallographic details of all solid phases are listed in Table 2. Invariant Equilibria A partial reaction scheme of the system, shown in Fig. 1, is based on the results of [1984Wac] and [2002Lin]. While [1984Wac] reported the liquidus surface of Cu- and Ni rich alloys, [2002Lin] determined the liquidus surface of entire system by means of differential thermal analysis and microstructure characterization of 28 ternary alloys. Five U type transitional invariant reactions have been reported. Among these, U1 was first proposed by [1971Bas] and subsequently confirmed by [1984Wac] and [2002Lin]. The presence of U2 was reported by [1984Wac] and subsequently confirmed by [2002Lin]. The invariant reaction U4 was postulated by [1988Gup], but labelled as U3. In Fig. 1, the temperature of invariant reactions are only approximate, and they are estimated based on the liquidus data of [2002Lin] and the participating binary invariant reactions. The composition of liquid participating in invariant reactions is listed in Table 3; however, the composition of the solid phases are not known. Liquidus, Solidus and Solvus Surfaces Figure 2 shows the liquidus surface of the ternary system adopted from [1984Wac, 2002Lin]. There are eight regions of primary crystallization. The locus of monovariant line e2-U1-p1-p3 and the superimposed Landolt-Börnstein New Series IV/11C3
MSIT®
306
Cu–Ni–Sn
isotherms are taken from [1984Wac]. [1971Bas] also reported liquidus isotherms, but only for Cu rich alloys containing up to 7 mass% Ni and 14 mass% Sn. [1984Wac] was the first to report the presence of invariant reaction U2; however, the primary crystallization field of J-Cu3Sn in [1984Wac] is much smaller compared to [2002Lin]. The Ni content of liquid phase of the invariant reaction U2 is about 13.3 at.% [2002Lin] compared to 2.2 at.% reported by [1984Wac]. [2002Lin] observed J-Cu3Sn as the primary crystallization product in Cu-10Ni-40Sn (at.%) and Cu-10Ni-45Sn (at.%) alloys. [1971Bas] reported solidus isotherms, but only for Cu rich alloys containing up to 10 mass% Ni and 3 mass% Sn. [1975Ple] determined the solidus and spinodal temperatures of Cu rich alloys by measuring resistivity, and their data are presented in Table 4. It is not known if the spinodal temperature refers to chemical spinodal or coherent spinodal. Isothermal Sections Figures 3, 4 and 5 show isothermal sections of Cu corner at 1090, 1050 and 1025°C, along with the tie lines between and liquid phase [1971Bas]. Figure 6 shows the isothermal section at 800°C [2003Wan]. [2003Mie] calculated the isothermal section of Cu corner at 800°C by CALPHAD method by treating 1-(Ni,Cu)3Sn as a solution phase and 2-(Ni,Cu)3Sn as an ordered phase. An important difference between experimental [2003Wan] and calculated [2003Mie] isothermal sections is the presence of phase fields involving 1 that extends up to the Cu-Sn binary edge. However, in the corresponding accepted binary phase diagram the 1 phase is stable only below 798°C. In the Cu corner, experimental phase diagram shows the presence of a three-phase field (L+ + 1-(Ni,Cu)3Sn) while the calculated phase diagram shows the presence of two three-phase fields (L+1+ 1-(Ni,Cu)3Sn and +1+ 1-(Ni,Cu)3Sn). Therefore, it is necessary that the three-phase field L+1+ 1-(Ni,Cu)3Sn originates at a temperature higher than 800°C. However, the reaction scheme, which is drawn on the basis of the liquidus surface (Fig. 2), is inconsistent with this three-phase field. This implies that there must be, yet undetected by thermal analysis, a liquidus maxima/minima and/or an invariant reaction involving 1 phase occurring at a temperature higher than 800°C. Also, the thermodynamic stability of 1 and 1-(Ni,Cu)3Sn phases must be enhanced in the presence of Ni. Figure 7 shows experimentally determined isothermal section of Cu corner at 700°C [1971Bas]. [2003Mie] calculated the partial isothermal section by CALPHAD method and obtained a good agreement with the experimental data. It is to be noted that the reaction scheme in Fig. 1 is consistent with the presence of a +1+ 1-(Ni,Cu)3Sn three-phase field at 700°C. Figure 8 shows experimentally determined partial isothermal section at 550°C. Miki and co-workers [1983Mik1, 1984Mik1] were the first to report the partial isothermal section, and their results were subsequently verified by [1996Oel]. Extensive transmission and analytical electron microscopy results show that the Cu contents of phase in the tie-triangles +-2+ 2-Ni3Sn and +-2+ 1-(Ni,Cu)3Sn are higher than that reported by [1983Mik1, 1984Mik1]. Accordingly, the results of [1996Oel] are accepted here. Also, in constructing Fig. 8 the homogeneity range of -2 is taken from [1989Pak]. [2003Mie] calculated the partial isothermal section by CALPHAD method. In addition to the above three-phase fields, the calculated diagram shows the presence of a three-phase field, +J-Cu3Sn+ 1-(Ni,Cu)3Sn, which is yet to be experimentally verified. Figures 9 and 10 show the isothermal sections at 240 [2002Lin, 2005Li] and 200°C [2003Che1], respectively. The only difference in phase fields of these two isothermal sections is that in Sn corner the liquid phase at 240°C is replaced by (Sn) at 200°C. The three-phase fields involving the ternary phase -2 have not been determined experimentally, therefore, they are shown as dashed. Also, in these recent studies [2002Lin, 2003Che1, 2005Li] the -3 phase has not been identified, even though in earlier studies it was reported to be stable below 460°C [1989Pak, 1990Pak]. [1984Wac] reported a partial isothermal section at 647°C; however, as discussed below, their results are inconsistent with subsequent results. Consequently, the results of [1984Wac] are not accepted here.
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Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
307
Temperature – Composition Sections Experimental determination of vertical sections is summarized as follows: at 2, 4, 7, 10, 15 and 25 mass% Ni [1932Ves]; at 2, 3, 5, 10 and 20 mass% Ni [1933Eas]; along Ni75Sn25-Cu30Ni45Sn25 [1981Wat]; along Cu3Sn-Ni3Sn [1984Wac]; along Ni75Sn25-Cu22Ni53Sn25 [1989Pak, 1990Pak]; at 14 at.% Ni [1990Che]; along Cu33.3Ni66.7-Cu25Ni50Sn25 [1995Bur, 1995Kra]; along Cu40Ni60-Cu28Ni42Sn30 and Ni75Sn25-Cu35Ni40Sn25 [1996Oel]. Besides, [2003Mie] calculated several vertical sections by CALPHAD method. Using the results of [1928Pri, 1932Ves, 1933Eas, 1971Bas], the vertical sections at 15 mass% Ni [1998Zha1] and 7.5 mass% Ni [1998Zha2] were reconstructed. The experimental and calculated vertical section at 2 mass% Ni are compared in Fig. 11 [1932Ves] and Fig. 12 [2003Mie], respectively. Similarly, the experimental and calculated vertical section at 4 mass% Ni are compared in Fig. 13 [1932Ves] and Fig. 14 [2003Mie], respectively. In both cases, the calculated vertical sections show a good agreement with the experimental ones with respect to liquidus, solidus and solvus boundaries. Figures 15 and 16 show the vertical sections at 7 and 10 mass% Ni [1932Ves], respectively. Figure 17 shows the vertical section at 14 at.% Ni [1990Che]. The experimental and calculated vertical section at 15 mass% Ni are compared in Fig. 18 [1932Ves] and 19 [2003Mie], respectively, show only a modest agreement for the solidus and solvus boundaries. The experimental and calculated vertical section along Cu3Sn-Ni3Sn are compared in Fig. 20 [1989Pak, 1990Pak] and Fig. 21 [2003Mie], respectively. Along this section, the phase equilibria of Ni rich alloys containing up to 30 at.% Cu are well documented due to several comprehensive studies [1981Wat, 1989Pak, 1990Pak, 1996Oel]. However, the phase equilibria on Cu rich alloys reported by [1984Wac] remain doubtful and warrant further experimental investigation. Besides the corrections made by [2000Gup], a major discrepancy between the results of [1984Wac] and others is the presence of +-Cu6Sn5, ++-Cu6Sn5, +-Cu6Sn5, + 1-(Ni,Cu)3Sn+-Cu6Sn5 phase fields at low temperatures. In the phase equilibria studies using bulk samples and diffusion couples at 240 [2002Lin] and 200°C [2003Che1], these phase fields were not observed. Therefore, the solid-state equilibria below 550°C along Cu3Sn-Ni3Sn section reported by [1984Wac] are not accepted here. In the calculated vertical section, Fig. 21, the solubility of Cu in 2-Ni3Sn is not reproduced and the ternary phase -2 is treated as line compound. Also, the calculated solubility of Sn in (~5 at.%) is much lower than the experimental value of 9 at.% [1996Spr]. The experimental and calculated vertical section at a constant ratio xCu/xNi of 1:2 are compared in Fig. 22 [1995Bur, 1996Spr, 1997Spr] and Fig. 23 [2003Mie], respectively. [1995Bur] reported only the liquidus and the solidus boundaries. In subsequent investigations, the solvus boundary was determined using resistivity, magnetometry and XRD [1996Spr, 1997Spr]. At 800 and 660°C, the solvus boundary is located at 9 [1996Spr] and 7 at.% Sn [1997Spr], respectively. The presence of L+ + 1 and L+ 1 phase fields were not mentioned by [1995Bur], but proposed by [2000Gup] and verified by CALPHAD modelling [2003Mie]. Thermodynamics The heat of mixing of liquid alloys was measured by SETARAM high-temperature calorimeter at 1307°C [1979Poo1] and 1250°C [2004Lue]. [1979Poo1] measured heat of mixing of 163 compositions, of which 119 lie along eleven sections at a constant ratio of xCu:xSn (1:9, 1:2, 3:2, 5:1, 1:5, 2:3, 2:1, 9:1, 1:3, 1:1, 3:1), 33 lie along three sections at a constant ratio of xNi:xSn (1:1, 3:1, 1:3), and 11 lie along the section at a constant ratio of xCu:xNi (3:1). [2004Lue] carried out 90 measurements along three sections at a constant ratio of xCu:xNi (1:3, 1:1, 3:1). Figures 24 and 25 show the isoenthalpy (mixing) contours at 1307°C [1979Poo1] and 1250°C [2004Lue], respectively. The composition dependence of integral enthalpy of mixing of liquid alloys has been modelled using a variety of analytical functions with [1979Poo2, 1984Hoc1, 1984Hoc2, 1995Tom, 2003Mie, 2004Lue] and without [1987Hoc, 1996Heu] ternary interaction parameter(s). It is found that in order to obtain a satisfactory fit of the experimental data by analytical functions, it is necessary to introduce ternary interaction parameter(s). This is attributed to strong asymmetry of the composition dependence of enthalpy of mixing in Cu-Sn and Ni-Sn systems, and also due to strong ternary interaction between Cu, Ni and Sn atoms in the liquid phase. [1979Poo2] used a modified quasi-chemical model containing one ternary Landolt-Börnstein New Series IV/11C3
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Cu–Ni–Sn
interaction parameter, [1995Tom] employed “Thermodynamic Adapted Power” (TAP)-series with one ternary interaction parameter, and [2003Mie, 2004Lue] used Redlich-Kister-Muggianu polynomial containing three ternary interaction parameters to fit experimental data of enthalpy of mixing of liquid alloys. [2003Mie] carried out a thermodynamic assessment of the ternary system by the CALPHAD method, and demonstrated the fidelity of thermodynamic models by calculating phase equilibria up to 70 at.% Ni and 25 at.% Sn which agree very well with the experimental data. [2003Mie] treated (Ni,Cu)3Sn(h) ( 1) as a solid solution phase, (Ni,Cu)3Sn(r) ( 2) as an ordered intermetallic, and considered ternary interaction parameters in liquid, bcc, fcc and (Ni,Cu)3Sn(h) phases. The ternary phase -2 was represented as Cu3Ni27Sn10 instead of Ni5CuSn2. The calculated liquidus and solidus isotherms, partial isothermal sections at 800, 700 and 550°C, and isopleths at 2, 4, 15 mass% Ni, 25 at.% Sn, and at a constant ratio xCu:xNi of 1:2 as well as the calculated enthalpies of mixing at 1580 K were presented. [2005Wan] calculated metastable miscibility gaps in the fcc phase along Cu1–wSnw-Ni1–wSnw sections with Sn contents of 2.3, 4, 6 and 8 mass%. Notes on Materials Properties and Applications A summary of experimental investigation of properties is given in Table 5. The mechanical and electrical properties of Cu rich alloys have been studied extensively due to their combined properties of high-strength, high-conductivity and resistance to stress relaxation up to 250°C [1978Ger, 1984Sco, 1991And, 1996Leh, 1997Vir1]. Due to these properties, Cu rich ternary alloys find numerous applications as electrical and electronic interconnection material. There is a large body of mechanical properties data (hardness, Young’s modulus, yield stress, ultimate tensile stress, ductility, stress relaxation and fracture behavior) [1928Pri, 1932Gui, 1974Sch1, 1974Sch2, 1975Ple, 1978Ger, 1978Lef, 1984Kra, 1984Sco, 1986Kat, 1986Kha, 1987Kra, 1989Kra, 1990Dey, 1991Abd, 1991And, 1992Bog, 1992Hun, 1994Her, 1994Mik, 1996Leh, 1997Her, 1997Vir1, 1997Vir2, 1999Jun, 1999Mor] and electrical properties [1975Ple, 1984Sco, 1996Leh, 1997Vir2, 1999Jun, 1999Mor]. It has been demonstrated that by controlling thermomechanical processing Cu rich alloys with strength up to 1.4 GPa can be developed. [2004Li] obtained the bulk modulus of Ni2CuSn by means of an empirical relationship. [1992Kra] measured the steady state creep rates of Ni2CuSn in the temperature range of 600 to 900°C. Analysis of their data gives an activation energy of 229.3 kJ#mol–1 and a stress exponent of 3.9. [2003Che2] determined tensile properties of Ni/Sn-0.7 mass% Cu/Ni solder joints after aging at 200°C for up to 240 h. Solid state aging causes growth of Cu6Sn5 and Ni3Sn4 intermetallics at the solder/substrate interface which degrade the tensile properties of solder joints. Miscellaneous The decomposition of Cu rich solid solutions, containing up to 15 mass% Ni and 13 mass% Sn and in the temperature range of 150 to 850°C, has been studied extensively. For a given alloy, the mechanism of decomposition is strongly related to the aging temperature, and at least six types of decomposition processes have been identified [1998Zha1]: (i) grain boundary and intergranular precipitation of 1 phase, (ii) cellular or discontinuous precipitation of 1 phase, (iii) precipitation of metastable D022 ((Ni1–xCux)3Sn) phase, (iv) precipitation of metastable L12 ((Ni1–xCux)3Sn) phase, (v) spinodal decomposition, and (vi) precipitation of a ternary phase (-2). Among these, spinodal decomposition [1974Sch1, 1974Sch2, 1975Ple, 1977Hel, 1978Ger, 1978Hel, 1978Lef, 1979Bab, 1980Dit, 1980Ray, 1982Ray, 1984Kra, 1984Sco, 1985Nar, 1986Kat, 1986Kha, 1987Kra, 1988Col, 1988Sat, 1990Dey, 1990Gou, 1991Abd, 1992Hun, 1993Sat, 1994Her, 1994Zha, 1995Pal, 1999Kim, 1998Zha1, 1998Zha2, 1999Jun, 1999Rhu, 2004Sah], and cellular or discontinuous precipitation [1982Mik, 1983Mik1, 1983Mik2, 1984Mik1, 1984Mik2, 1990Mik, 1991Mik, 1994Mik, 1998Vir, 2004Lou] processes have received most attention. [1998Zha1] provided the time-temperature-transformation (TTT) curves of all decomposition processes in a Cu-15Ni-8Sn (mass%) alloy.
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Cu–Ni–Sn
309
Atomic transport kinetics in solid solutions has been studied by [1972Bas] and [1986Tak]. Based on the Cu , at 790°C are results of four diffusion couples, [1972Bas] reported that the interdiffusion coefficient, D SnNi –14 2 –1 –14 2 –1 m #s at 4.7 mass% Ni and 8 mass% Sn, and 1.23#10 m #s at 13.7 mass% Ni and 4.96#10 5 mass% Sn. In a subsequent and more comprehensive study, [1986Tak] measured interdiffusion coefficients in (Cu) solid solutions in the temperature range of 750 to 850°C, and confirmed the presence of zero-flux plane. They used twelve ternary diffusion couples containing up to 17.7 at.% Ni and 7.1 at.% Sn which were annealed at 750, 775, 800, 825 and 850°C. Their results show that the interdiffusion Cu , D Cu , D Cu Cu coefficients D NiNi NiSn SnNi and D SnSn are very sensitive to Ni and Sn concentrations. In particular, the interdiffusion coefficients in a Cu-6 at.% Ni-3.7 at.% Sn are expressed as: Cu = 1.4 # 10–5 exp (–189000/RT) in m2#s–1 D NiNi Cu = –3.1 # 10–9 exp (–120000/RT) in m2#s–1 D NiSn Cu = –5.8 # 10–7 exp (–153000/RT) in m2#s–1 D SnNi Cu = 2.1 # 10–5 exp (–175000/RT) in m2#s–1 D SnSn
where the activation energies are in J#mol–1, R is the universal gas constant and T is the temperature in K. One of the major hurdles of conventional processing of Cu rich alloys on a commercial basis is the severe segregation of Ni and Sn. Rapid solidification of such alloys minimizes chemical heterogeneity leading to uniform microstructure and properties [1988Col, 1992Bog, 1999Mor]. The efficacy of thermodynamic and kinetic approach as a predictive tool for homogenization of highly segregated alloys was demonstrated by [2001Pur] for a commercial alloy of Cu-15Ni-8Sn (mass%). The cellular or discontinuous reaction, ’+ 1-(Ni,Cu)3Sn, in ternary alloys has been studied extensively in the temperature range of 350 to 500°C [1982Mik, 1983Mik1, 1983Mik2, 1984Mik1, 1984Mik2, 1990Mik, 1991Mik, 1994Mik]. The discontinuous reaction may be suppressed by adding B and P which segregate in grain boundaries [1983Mik2, 1984Mik2]. The effect of Al, Cr, Fe, In, Mn, Si and Ti on cellular precipitation kinetics has also been investigated [1994Mik]. Among these, Al, Cr, Si and Ti are most effective in retarding the reaction kinetics. In particular, it has been shown that the precipitation of Ni31Si12 [1991Mik] and Ni3Ti [1994Mik] phase in grain boundaries are responsible for the retardation effect. In view of lead-free solders in electronic packaging, the growth kinetics of Cu3Sn and Cu6Sn5 intermetallic layers has been studied, in the temperature range of 80 to 170°C, using diffusion couples [2001You, 2004Yoo]. [2001You] prepared solder joints using Cu and Sn-2.7Cu-0.2Ni (mass%) alloy that were aged in the temperature range of 100 to 170°C for up to 360 h, while [2004Yoo] prepared solder joints of Cu and Sn-0.6Cu-0.05Ni (mass%) alloy that were aged in the temperature range of 80 to 150°C for up to 60 d. The activation energies for growth of Cu3Sn was reported to be 57.06 kJ#mol–1 [2001You], while that of Cu6Sn5 ranges from 66.35 kJ#mol–1 [2004Yoo] to 83.8 kJ#mol–1 [2001You]. Besides studies on thermal aging of solder joints, [2001Che] prepared Sn-0.7 mass% Cu/Ni/Sn-0.7 mass% Cu solder joints that were subjected to a current density of 500 A#cm–2 in the temperature range of 160 to 200°C. They found that in the absence of electrical current only Ni3Sn4 grows at both interfaces, while the passage of electric current causes growth of both Cu6Sn5 and Ni3Sn4 at the solder/Ni interface and only Ni3Sn4 at the Ni/solder interface. They also found that the direction of movement of electrons, Cu and Sn is the same at solder/Cu interface, and the growth rates of intermetallic layers are enhanced. Calculation of apparent effective charge, Za*, shows its decreasing magnitude with increasing temperature implying that the electromigration effect becomes insignificant at higher temperatures. [1971Bas] measured the partitioning ratio of Ni and Sn between (Cu) and liquid in alloys containing up to 9 mass% Ni and 9 mass% Sn. The Curie temperature of -(Ni2Cu)1–xSnx solid solution decreases linearly with Sn concentration up to 7 at.% [1997Spr], from 32°C at CuNi2 to –52°C at Cu31Ni62Sn7.
Landolt-Börnstein New Series IV/11C3
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310
Cu–Ni–Sn
References [1928Pri]
[1932Gui] [1932Ves]
[1933Eas]
[1936Hum] [1937Rah] [1940Glu]
[1948Ray] [1949Jae] [1953Maz1]
[1953Maz2]
[1969Gue]
[1971Bas]
[1971Dok]
[1972Bas]
[1973Lev] [1973Pak]
[1974Sch1]
[1974Sch2]
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Price, W.B., Grant, C. G., Phillips, A. J., “ Phase Boundary of Copper-Nickel-Tin System”, Trans. Amer. Inst Min. Met. Eng., 78, 503-517 (1928) (Experimental, Phase Diagram, #, *, 7) Guillet, L., Le Thomas, A., Ballay, M., “Properties of Cu-Ni Alloys” (in French), Compt. Rend. Acad. Sci. Paris, 194, 2102-2105 (1932) (Experimental, Phys. Prop., 1) Veszelka, J., “Investigation about the Balance Relations of Further Alloyed Bronze” (in German), Mitt. Berg-Huettenmaenn. Abt. Ungar. Hochschule Berg.-Forstw. Sopron, 4, 163-203 (1932) (Experimental, Phase Relations, #, *, 17) Eash, J.T., Upthegrove, C., “The Copper-Rich Alloys of the Cu-Ni-Sn System”, Trans. Am. Inst. Min. Metall. Pet. Eng., 104, 221-253 (1933) (Experimental, Phase Diagram, Phase Relations, #, *, 32) Hume-Rothery, W., “On the Theory of Equilibriun in Alloys.-I.”, Philos. Mag., 22, 1013 (1936) (Assessment, Crys. Structure, Experimental, Phase Diagram, Phase Relations, *, 24) Rahlfs, P., “About New Ternary Alloys with Superstructure Related to the Brass-Type” (in German), Metallwirtschaft, 16, 640-643 (1937) (Crys. Structure, Experimental, 8) Gluskin, D.Ya., “Phases, Formed by Interaction of Refractory Metals with Light-Melts Metals” (in Russian), Zh. Tekh. Fiz., 10(18), 1486 (1940) (Crys. Structure, Experimental, 15) Raynor, G.V., “XXVIII. Equilibrium Relationships in Ternary Alloys”, Philos. Mag., 39(290), 218-229 (1948) (Phase Diagram, Review, Theory, *, 12) Jänecke, E., “Cu-Ni-Sn” (in German), Kurzgefasstes Handbuch aller Legierungen, 517-518 (1949) (Phase Diagram) Mazzoleni, F., “Hume-Rothery Phases in the Alloys with more then two Components. II. Phases with the Ratio 7:4 in the System Ni-Cu-Sn” (in Italian), Metall. Ital., 45(10), 366-369 (1953) (Crys. Structure, Experimental, 13) Mazzoleni, F., “Hume-Rothery Phases in the Alloys with more then two Components. II. Phases with the Ratio 3:2 and 21:13 in the System Ni-Cu-Sn” (in Italian), Metall. Ital., 45(10), 370-373 (1953) (Crys. Structure, Experimental, 6) Guertler, W., Guertler, M., Anastasiadias, E., “Copper-Nickel-Tin”, A Comp. of Const. Ternary Diagr. Met. Systems, Isr. Pro. Sci. Tr., Jerusalem, 591-594 (1969) (Phase Diagram, Phase Relations, 8) Bastow, B.D., Kirkwood, D.H., “Solid/Liquid Equilibrium in the Cu-Ni-Sn System Determined by Microprobe Analysis”, J. Inst. Met., 99, 277-283 (1971) (Experimental, Phase Diagram, Phase Relations, #, *, 13) Dokusogus, H.Z., Stadelmaier, H.H., Bowen, L.H., “119Sn and 121Sb Mossbauer Effect Study of Ternary Nickel Borides and Heusler Alloys”, J. Less-Common Met., 23, 245-251 (1971) (Experimental, 21) Bastow, B.D., Kirkwood, D.H., “Binary and Ternary Diffusion in the Cu Corner of the Copper-Nickel-Tin System”, J. Inst. Met., 100, 24-28 (1972) (Experimental, Transport Phenomena, 16) Levine, E,D, “Cu-Ni-Sn (Copper-Nickel-Tin)”, Metals Handbook, 8, 426-427 (1973) (Phase Diagram, 3) Pak, H.R., Saburi, T., Nenno, S., “Martensitic and Massive Transformations in Ni3Sn Alloy” (in Japanese), J. Jpn. Inst. Met., 37, 1128-1134 (1973) (Experimental, Crys. Structure, 21) Schwartz, L.H., Mahajan, S., Plewes, J.T., “Spinodal Decomposition in A Cu-9Ni-6Sn (wt.%) Alloy”, Acta Metall., 22, 601-609 (1974) (Crys. Structure, Experimental, Phase Diagram, Phys. Prop., Thermodyn., 23) Schwartz, L.H., Plewes, J.T., “Spinodal Decomposition in A Cu-9Ni-6Sn (wt.%) Alloy II. A Critical Examination of Mechanical Strength of Spinodal Alloys”, Acta Metall., 22, Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
[1975Ple]
[1977Boo]
[1977Hel]
[1978Ger]
[1978Hel]
[1978Lef]
[1979Bab]
[1979Cha]
[1979Dri]
[1979Poo1]
[1979Poo2]
[1979Sch]
[1980Dit]
[1980Ray]
[1981Wat]
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911-921 (1974) (Crys. Structure, Experimental, Mechan. Prop., Morphology, Phase Relations, Electr. Prop., 33) Plewes, J.T., “High-Strength Cu-Ni-Sn Alloys by Thermomechanical Processing”, Metall. Trans. A, 6, 537-544 (1975) (Experimental, Phase Relations, Morphology, Mechan. Prop., 24) Booth, M.H., Brandon, J.K., Brizard, R.Y., Chieh, C., Pearson, W.B., “ -Brasses with F Cells”, Acta Crystallogr., Sect. B: Struct. Crystallogr. Crys. Chem., 33B(1), 30-36 (1977) (Crys. Structure, 22) Helmi, Fatma, M., Zsoldos, Lehel., “On the Thermal Decomposition of Cu-Ni-Sn after Prior Cold - Work”, Scr. Metall., 11(10), 899-901 (1977) (Crys. Structure, Experimental, Phase Relations, 6) Gerlach, U., Kreye, H., “Structural and Mechanical Properties of the Cu-19Ni-2Sn Alloy” (in German), Metall, 32(11), 1112-1115 (1978) (Experimental, Crys. Structure, Morphology, Mechan. Prop., 12) Helmi, F.M., Cziraki, A., Fogarassy, B., Zsoldos, L., Eigner, V., “Phase Transformations in Cold-Worked Cu-8Ni-5Sn Alloy During Thermal Ageing”, Acta Crystallogr., Sect. A: Found. Crystallogr., 34 (S4), S364 (1978) (Experimental, 6) Lefevre, B.G., D`Annessa, A.T., Kalish, D., “Age Hardening in Cu-15Ni-8Sn Alloy”, Metall. Trans. A, 9A(4), 577-586 (1978) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 29) Baburaj, E.G., Kulkarni, U.D., Menon, E.S.K., Krishnan, R., “Initial Stages of Decomposition in Cu-9Ni-6Sn”, J. Appl. Crystallogr., 12(5), 476-80 (1979) (Crys. Structure, Experimental, Phase Relations, 12) Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Ni-Sn”, INCRA Monograph Series 6. Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Systems, NSRD, Washington, 6, 604-609 (1979) (Phase Diagram, Phase Relations, Review, 13) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Cu-Ni-Sn” (in Russian), Binary and Multicomponent Copper-Base Systems”, Nauka, Moscow, 179 (1979) (Phase Diagram, Phase Relations, Review, 3) Pool, M.J., Arpshofen, I., Predel, B., Schultheiss, E., “Determination of Mixing Enthalpies of Liquid Alloys in the Cu-Ni-Sn Ternary System Using a SETARAM High Temperature Calorimeter” (in German), Z. Metallkd., 70(10), 656-659 (1979) (Experimental, Thermodyn., *, 6) Pool, M.J., Arpshofen, I., Predel, B., Schultheiss, E., “Calculation and Analysis of the Mixing Enthalpies in Liquid Cu-Ni-Sn Alloys using Various Models” (in German), Z. Metallkd., 70(11), 726-731 (1979) (Calculation, Thermodyn., 7) Schreiner, W.H., Pureur, P., Grandi, T.A., Kunzler, J.V., Brandas, D.E., “A Thermal X-Ray and Resistivity Study of the Heusler Alloy Cu2NiSn”, J. Therm. Anal., 17, 489-494 (1979) (Crys. Structure, Electr. Prop., Experimental, 11) Ditchek, B., Schwartz, L.H., “Diffraction Study of Spinodal Decomposition in Cu-10 wt.-% Ni-6 wt.% Sn”, Acta Metall., 28 (6), 807-822 (1980) (Experimental, Morphology, Crys. Structure, 27) Ray, R.K., Narayanan, S.C., Devaraj, S., “Precipitation and Recrystallization in a Cu-9Ni-6Sn Alloy”, Scr. Metall., 14, 1181-1184 (1980) (Electronic Structure, Experimental, 5) Watanabe, Y., Murakumi, Y., Kachi, S., “Martensitic and Massive Transformation and Phase Diagram in Ni30–xMxSn (M = Cu, Mn) Alloys” (in Japanese), Nippon Kinzoku Gakkai Shi, 45(6), 551-558 (1981) (Crys. Structure, Morphology, Phase Relations, Phase Diagram, Experimental, #, *, 25)
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312 [1982Mik]
[1982Ray]
[1982Sak] [1982Ste]
[1983Mik1]
[1983Mik2]
[1983Mur]
[1984Hoc1]
[1984Hoc2]
[1984Kra]
[1984Mik1]
[1984Mik2]
[1984Sco]
[1984Wac]
[1985Nas] [1985Nar]
[1986Kat]
MSIT®
Cu–Ni–Sn Miki, M., Sogabe, T., Hori, S., “Grain Boundary Reaction in a Cu-10Ni-8Sn (wt.%) Alloy”, Nippon Kinzoku Gakkai Shi, 46(3), 307-12 (1982) (Experimental, Interface Phenomena, Crys. Structure, 24) Ray, R.K., Narayanan, S.C., “Combined Recrystallization and Precipitation in a Cu-9Ni-6Sn Alloy”, Metall. Trans. A, 13, 565-573 (1982) (Electronic Structure, Experimental, 16) Sakamoto, T., “Technical Trends in Metallic Materials for Electronic Applications” (in Japanese), Denki Seiko (Electr. Furn. Steel), 53(2), 137-142 (1982) (Review) Steeb, J., Stuer, H., Durrschnabel, W., “New Spring Material Based on Copper, with High Strength and Good Electricity Conductivity” (in German), Metall, 36(11), 1185-1188 (1982) (Electr. Prop., Experimental, 8) Miki, M., Ogino, Y., “Precipitation in a Cu-20Ni-8Sn Alloy and the Phase Diagram of the Cu-Ni Rich Cu-Ni-Sn System. (Translation)”, Nippon Kinzoku Gakkai Shi, 47(7), 602-610 (1983) (Experimental, Crys. Structure, Morphology, Phase Diagram, Phase Relations, Mechan. Prop., #, *, 32) Miki, M., Ogino, Y., “Effect of the Addition of B and P on the Cellular Precipitation in Ni-Sn and Cu-Ni-Sn Alloys” (in Japanese), Nippon Kinzoku Gakkai Shi, 47(11), 983-990 (1983) (Experimental, Mechan. Prop., Phase Diagram, Morphology, Phase Relations, 25) Murakami, Y., Kachi, S., “On the Morphology of Massive and Martensitic Phases With 2H Structure in Ni67Cu8Sn25 Alloy”, Trans. Jpn. Inst. Met., 24(11), 741-747 (1983) (Morphology, Phase Diagram, 15) Hoch, M., Arpshofen, I., “A Modified Aggregation Model for the Calculation of Thermodynamic State Functions of Liquid Alloys” (in German), Z. Metallkd., 75(1), 23-40 (1984) (Thermodyn., Theory, 13) Hoch, M., Arpshofen, I., Predel, B., “Representation of Thermodynamic State Functions of Liquid Alloys Using a Modified Aggregation Model” (in German), Z. Metallkd., 75(1), 30-40 (1984) (Thermodyn., Theory, 15) Kratochvil, P., Mencl, J., Pesicka, J., Komnik, S.N., “The Structure and Low-Temperature Strength of the Age Hardened Cu-Ni-Sn Alloys”, Acta Metall., 32(9), 1493-1497 (1984) (Crys. Structure, Experimental, 20) Miki, M., Ogino Y., “Precipitation in a Cu-20Ni-8Sn Alloy and the Phase Diagram of the Cu-Ni Rich Cu-Ni-Sn System”, Trans. Jpn. Inst. Met., 25(9), 593-602 (1984) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, Phase Relations, Thermodyn., #, *, 32) Miki, M., Ogino, Y., “Effect of Addition of B and P on the Celluar Precipitation in Ni-Sn and Cu-Ni-Sn Alloys”, Trans. Jpn. Inst. Met., 25(9), 611-619 (1984) (Experimental, Mechan. Prop., 24) Scorey, C.R., Chin, S., Mhite, M.J., Livak, R.J., “Spinodal Cu-Ni-Sn Alloys for Electronic Applications”, J. Metals, 36(11), 52-54 (1984) (Electr. Prop., Experimental, Mechan. Prop., Phys. Prop., 12) Wachtel, E., Bayer, E., “Constitution and Magnetic Properties of the Ternary System Cu-Ni-Sn” (in German), Z. Metallkd., 75(1), 61-69 (1984) (Experimental, Phase Diagram, 47) Nash, P., Nash, A., “The Ni-Sn (Nickel-Tin) System”, Bull. Alloy Phase Diagrams, 6, 350-359 (1985) (Crys. Structure, Phase Diagram, Phase Relations, Review, #, *, 57) Narayanan, S.C., Ray, R.K., Devaraj, S., “Thermal Decomposition of a Cu-Ni-Sn Alloy With and Without Prior Cold-Work”, Trans. Indian Inst. Met., 38(1), 81-85 (1985) Experimental, Crys. Structure, 7) Kato, M., Katsuta, S., Okamine, S., Sato, A., “Deformation Behaviour and Microstructure of Cu-10Ni-6Sn Spinodal Alloy Single Crystals”, Mater. Sci. Eng. A, 77, 95-102 (1986) (Experimental, Mechan. Prop., 24)
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn [1986Kha]
[1986Tak]
[1986Zie] [1987Hoc]
[1987Kra]
[1988Col] [1988Gup]
[1988Sat] [1989Kra] [1989Pak] [1990Che]
[1990Dey]
[1990Gou]
[1990Mik]
[1990Pak]
[1990Sau]
[1991Abd]
[1991And]
Landolt-Börnstein New Series IV/11C3
313
Khamitov, Kh., Chumlyakov, Yu.I., Korotayev, A.D., “Deformation Mechanism of High-Duty Heterophase Cu-Ni-Sn Single Crystals”, Phys. Met. Metallogr., 62(2), 142-150 (1986) (Experimental, 19) Takahashi, T., Katoh, M., Minamino, Y., Yamane, T., “Ternary Diffusion and Thermodynamic Interactions in Alpha-Copper-Nickel-Tin Solid Solutions”, Trans. Jpn. Inst. Met., 27 (12), 949-959 (1986) (Experimental, Thermodyn., Phase Relations, Transport Phenomena, 37) Zeiger, H., “State of the Art and Development of Copper and Copper Alloys” (in German), Z. Werkstofftech., 17(3), 75-78 (1986) (Review, 32) Hoch, M., “Application of the Hoch-Arpshofen Model to the Thermodynamics of the Cu-Ni-Sn, Cu-Fe-Ni, Cu-Mg-Al and Cu-Mg-Zn Systems”, Calphad, 11(3), 237-246 (1987) (Thermodyn., 16) Kratochvil, P., Mandula, M., Mencl, J., Pesicka, J., Smola, B., “The Constitution of the Cu-9Ni-6Sn Alloy Determined by Changes of the Strength and by Structure Determination during Annealing”, Phys. Status Solidi A, 104(2), 597-602 (1987) (Crys. Structure, Experimental, 14) Collins, L.E., Barry, J.B., “Reduced Segregation in Rapidly Solidified Cu-Ni-Sn Alloys”, Mater. Sci. Eng., 98, 335-338 (1988) (Experimental, 4) Gupta, K. P., Rajendraprasad, S.B., Ramakrishna, D., Jena, A.K., “The Cu-Ni-Sn (Copper-Nickel-Tin) System”, J. Alloy Phase Diagrams, 4(3), 160-74 (1988) (Phase Diagram, Phase Relations, Review, 11) Sato, A., Katsuta, S.-I., Kato, M., “Stress Aging of a Cu-10Ni-6Sn Spinodal Alloy”, Acta Metall., 36(3), 633-640 (1988) (Experimental, Mechan. Prop., 19) Kratochvil, P., Pesicka, J., “Strengthening in the Cu-9Ni-1.2Sn Alloy”, J. Mater. Sci., 24, 2433-2436 (1989) (Experimental, Mechan. Prop., 19) Pak, J.C.L., Mukherjee, K., Inal, O.T., Pak, H.-R., “Phase Transformations in (Ni,Cu)3Sn Alloys”, Mater. Sci. Eng., 117A, 167-173 (1989) (Experimental, 15) Chepeleva, V.P., Yupko, L.M., Mgebrova, A.G., Cherepenina, E.S., “Phase Transformations and Structure of Eutectic Alloys in Cu-Sn-Ni and Cu-Sn-Ni-Cr Systems”, Sov. Powder Metall. Met. Ceram. (Engl. Transl.), 29(5), 357-360 (1990), translated from Poroshk. Metall., 5(329), 24-28, 1990 (Experimental, Phase Diagram, Phase Relations, Morphology, 7) Deyong, L., Tremblay, R., Angers, R., “Microstructural and Mechanical Properties of Rapidly Solidified Cu-Ni-Sn Alloys’”, Mater. Sci. Eng. A, 124, 223-231 (1990) (Experimental, Mechan. Prop., 11) Goudeau, Ph., Naudon, A., Welter, J.-M., “Anomalous Small-Angle X-Ray Scattering Study of the Early Stages of Decomposition in Cu-15Ni-8Sn (wt.%)”, J. Appl. Crystallogr., 23, 266-276 (1990) (Electronic Structure, Experimental, Phys. Prop., Theory, 24) Miki, M., Ogino, Y., “Effect of Si Addition on the Cellular Precipitation in a Cu-10Ni-8Sn Alloy.”, Mat. Trans., JIM, 31(11), 968-974 (1990) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, Review, 24) Pak, J.S.L., Mukherjee, K., Inal, O.T., Pak, H.R., “Crystal Structures and Microstructures of (Ni,Cu)3Sn Alloys”, Mater. Sci. Eng., A130, 83-92 (1990) (Experimental, Crys. Structure, 11) Saunders, N., Miodownik, A.P., “The Cu-Sn (Copper-Tin) System”, Bull. Alloy Phase Diagrams, 11, 278-287 (1990) (Phase Diagram, Phase Relations, Crys. Structure, Review, #, *, 57) Abdellatief, A.Y., Kratochvil, P., “Dynamic Softening of Spinodal Alloy Cu9Ni6Sn during Hot Deformation”, Mater. Sci. Eng. A, 137, 185-188 (1991) (Experimental, Mechan. Prop., 11) Andersson, C., Kamf, A., Sudberg, R., “Stress Relaxation of Some Copper Alloys for Connectors - Comparison Between Ni and Fe in Combination with Sn and Influence of MSIT®
314
[1991Mik]
[1992Bog]
[1992Hun]
[1992Kra]
[1993Sat]
[1994Her]
[1994Lar] [1994Mik]
[1994Zha]
[1995Bur]
[1995Kra]
[1995Lar] [1995Pal]
[1995Tom]
[1996Heu]
[1996Leh]
MSIT®
Cu–Ni–Sn Conductivity”, Mater. Issues Adv. Electron. Opto-Electron. Connectors, 29-42 (1991) (Electr. Prop., Experimental, 8) Miki, M., Ogino, Y., “Influence of Solution-Treatment Conditions on the Celluar Precipitation in Si-Doped Cu-10Ni-8Sn Alloy”, Mater. Trans., JIM, 32(12), 1135-1140 (1991) (Experimental, Mechan. Prop., 20) Bogel, A., Feind, J., “Material for Friction Bearings Made of Copper Alloys for Use in Automotive Industry” (in German), Metall (Berlin), 46(11), 1132-1135 (1992) (Experimental, 11) Hunnik, van E.W.J., Colijn, J., Schade van Westrum, J.A.F.M., “Heat Treatment and Phase Inter-Relationships of the Spray Cast Cu-15 wt.% Ni-8 wt.% Sn Alloy”, Mater. Sci. Forum, 102, 115-123 (1992) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, Phase Relations, 9) Kratochvil, P., Biegel, W., Sprusil, B., Chalupa, B., “Deformation and Phase Stability of Ni2CuSn Intermetallic Compound”, Phys. Status Solidi A, 131(2), 321-332 (1992) (Experimental, Mechan. Prop., Phase Diagram, Phase Relations, 12) Sato, A., Tamura, K., Ito, M., Kato, M., Mori, T., “In Situ Observation of Moving Dislocations in a Cu-10Ni-6Sn Spinodal Alloy”, Acta Metall. Mat., 41(4), 1047-1055 (1993) (Experimental, Mechan. Prop., 12) Hermann, Ph., Morris, D.G., “Relationship Between Microstructure and Mechanical Properties of a Spinodally Decomposing Cu-15Ni-8Sn Alloy Prepared by Spray Deposition”, Metall. Mater. Trans. A, 25, 1403-1412 (1994) (Experimental, Mechan. Prop., 30) Larsson, A.K., Stenberg, L., Lidin, S., “The Superstructure of Domain-Twinned ’-Cu6Sn5”, Acta Crystallogr., B50, 636-643 (1994) (Crys. Structure, Experimental, 14) Miki, M., Ogino, Y., “Effects of Doped Elements on the Celluar Precipitation in Cu-10Ni-8Sn Alloy”, Mater. Trans., JIM, 35(5), 313-318 (1994) (Experimental, Mechan. Prop., 13) Zhao, J.C., Notis, M.R., Tsubakino, H., “Spinodal Decomposition and Ordering Transformation in a Cu-15Ni-8Sn Alloy”, Proc. Int. Conf. Solid-Solid Phase Transformations in Inorganic Materials, PTM’94, Johnson, W.C., Howe, J.M., Laughlin, D.E., Soffa, W.A. (Eds.), 329-334, TMS, Warrendale, (1994) (Experimental, Crys. Structure, *, 10) Burianek, J., Sprusil, B., “DTA Investigation of the (Ni2Cu)1–xSnx System”, J. Alloys Compd., 227, L6-L8 (1995) (Experimental, Phase Diagram, Phase Relations, #, *, 6) Kratochvil, P., Sprusil, B., Oelgeschlager, D., Smola, B., Pesicka, J., “Dynamical Phase Analysis Along Pseudobinary Line Ni2Cu-Ni2CuSn”, Intermetallics, 3, 455-466 (1995) (Crys. Structure, Experimental, Phase Relations, #, *, 19) Larsson, A.K., Stenberg, L., Lidin, S., “Crystal Structure Modulations in -Cu5Sn4”, Z. Krystallogr., 210, 832-837 (1995) (Crys. Structure, Experimental, 15) Pal, H., Pradhan, S.K., De, M., “Microstructure and Phase-Transformation Studies of Cu-Ni-Sn Alloys”, Jpn. J. Appl. Phys., 34, 1619-1626 (1995) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 34) Tomiska, J., Wang, H., “On the Algebraic Evaluation of the Ternary Molar Heat of Mixing HE from Calorimetric Investigations”, Ber. Bunsen-Ges. Phys. Chem., 99(4), 633-640 (1995) (Thermodyn., Theory, 27) Heuzey, M.-C., Pelton, A.D., “Critical Evaluation and Optimization of the Thermodynamic Properties of Liquid Tin Solutions”, Metall. Mater. Trans. B, 27B(5), 810-828 (1996) (Assessment, Calculation, Phase Relations, Thermodyn., 156) Lehtinen, P., Tiainen, T., Laakso, L., “New Continuously Cast CuNiSn Alloys Provide Excellent Strength and High Electrical Conductivity”, Metall, 50(4), 267-271 (1996) (Electr. Prop., Experimental, Mechan. Prop., 8) Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn [1996Oel]
[1996Rud]
[1996Shi]
[1996Spr]
[1997Her]
[1997Spr]
[1997Vir1]
[1997Vir2]
[1998Vir]
[1998Zha1]
[1998Zha2]
[1999Gho]
[1999Jun]
[1999Kim]
[1999Mor]
[1999Rhu]
Landolt-Börnstein New Series IV/11C3
315
Oelgeschlaeger, D., Kratovchvil, P., Freyhardt, H.C., “Phase Formation and Kinetics in Melt-Spun and Sputtered Alloys Along the Line Cu40Ni60-Cu30Ni45Sn25”, Intermetallics, 4, 319-325 (1996) (Experimental, Phase Diagram, Phase Relations, Kinetics, #, *, 9) Rudajevova, A., Kratochvil P., “Phase Transformations in a Stoichiometric and Off-Stoichiometric Ni2CuSn”, Mater. Sci. Eng. A, 208, 122-125 (1996) (Experimental, Phase Relations, 9) Shim, J.-H., Oh, C.-S., Lee, B.-J., Lee, D.N., “Thermodynamic Assessment of the Cu-Sn System”, Z. Metallkd., 87, 205-212 (1996) (Phase Diagram, Phase Relations, Thermodyn., 36) Sprusil, S., Kratochvil P., Chalupa, B., “Resistometric and X-ray Investigation of (Ni2Cu)1–xSnx Alloys”, Czech. J. Phys., 46(5), 491-502 (1996) (Crys. Structure, Experimental, Phase Relations, 15) Hermann, Ph., Biselli, C., Morris, D.G., “Influence of Heat Treatments on Microstructure and Mechanical Properties of Cu-15Ni-8Sn Prepared by Spray Forming”, Mater. Sci. Technol., 13, 489-496 (1997) (Experimental, Mechan. Prop., Morphology, 22) Sprusil, S., Chalupa, B., “The Influence of Sn Alloying on the Magnetic Transformation in Ni2Cu”, J. Phase Equilib., 18(2), 152-156 (1997) (Calculation, Experimental, Magn. Prop., Phase Relations, Review, #, *, 24) Virtanen, P., Tiainen, T., “Stress Relaxation Behaviour in Bending of High Strength Copper Alloys in the Cu-Ni-Sn System”, Mater. Sci. Eng. A, 238, 407-410 (1997) (Experimental, Mechan. Prop., 13) Virtanen, P., T.Tianen, T., “Effect of Nickel Content on the Decomposition Behavior and Properties of CuNiSn Alloys”, Phys. Status Solidi A, 159, 305-316 (1997) (Crys. Structure, Experimental, Phase Relations, 19) Virtanen, P., Tiainen, T., Lepistoe, T., “Precipitation at Faceting Grain Boundaries of Cu-Ni-Sn Alloys”, Mater. Sci. Eng. A, 251, 269-275 (1998) (Experimental, Mechan. Prop., 39) Zhao, J.-C., Notis, M.R., “Spinodal Decomposition, Ordering Transformation, and Discontinuous Precipitation in a Cu-15Ni-8Sn Alloy”, Acta Mater., 46(12), 4203-4218 (1998) (Crys. Structure, Electr. Prop., Electronic Structure, Experimental, Morphology, Phase Relations, 45) Zhao, J.-C., Notis, M.R., “Microstructure and Precipitation Kinetics in a Cu-7.5Ni-5Sn Alloy”, Scr. Mater., 39(11), 1509-1516 (1998) (Experimental, Phase Relations, Morphology, Kinetics, Crys. Structure, 4) Ghosh, G., “Thermodynamic Modelling of the Nickel-Lead-Tin System”, Metall. Mater. Trans. A, 30A, 1481-1494 (1999) (Phase Diagram, Phase Relations, Thermodyn., Calculation, 109) Jung, Y.-C., Lee, J.-M., Han, S.-Z., Kim, C.-J., “Effect of Aging Temperature and Microstructure on Mechanical Properties of Cu-9Ni-6Sn Alloy” (in Japanese), J. Jpn. Inst. Met., 63(10), 1338-1347 (1999) (Crys. Structure, Experimental, Morphology, Mechan. Prop., 25) Kim, S.S., Rhu, J.C., Jung, Y.C., Han, S.Z., Kim, C.J., “Aging Characteristics of Thermochemically Processed Cu-9Ni-6Sn Alloy”, Scr. Mater., 40(1), 1-6 (1999) (Experimental, Mechan. Prop., 6) Morris, D. G., “Recent Research on Advanced Copper Alloys: Possibilities for Osprey Spray Deposition”, Powder Met., 42(1), 20-26 (1999) (Electr. Prop., Experimental, Mechan. Prop., Phase Relations, Review, 35) Rhu, J.C., Kim, S.S., Jung, Y.C., Han, S.Z., Kim, C.J., “Tensile Strength of Thermomechanically Processed Cu-9Ni-6Sn Alloys”, Metall. Mater. Trans. A, 30, 2649-2657 (1999) (Experimental, Mechan. Prop., Morphology, 11)
MSIT®
316 [2000Gup] [2000Moo]
[2001Che]
[2001Gho]
[2001Pur]
[2001You]
[2002Kat] [2002Leb]
[2002Lin]
[2002Zen]
[2003Che1]
[2003Che2]
[2003Mie]
[2003Wan]
[2004Gho]
[2004Jan]
MSIT®
Cu–Ni–Sn Gupta, K.P., “An Expanced Cu-Ni-Sn System (Copper-Nickel-Tin)”, J. Phase Equilib., 21(5), 479-484 (2000) (Crys. Structure, Phase Diagram, Phase Relations, Review, 8) Moon, K.-W., Boenttinger, W.J., Kattner, U.R., Biancaniello, F.S., Handwerker, C.A., “Experimental and Thermodynamic Assessment Sn-Ag-Cu Solder Alloys”, J. Electr. Mater., 29, 1122-1136 (2000) (Phase Diagram, Phase Relations, Thermodyn., Calculation, Experimental, 24) Chen, C.-M., Chen, S.-W., “Electromigration Effect upon the Sn-0,7 wt.% Cu/Ni and Sn-3,5 wt.% Ag/Ni Interfacial Reactions”, J. Appl. Phys., 90(3), 1208-1214 (2001) (Experimental, Phase Diagram, Phase Relations, Interface Phenomena, Transport Phenomena, 35) Ghosh, G., “Dissolution and Interfacial Reactions of Thin-Film Ti/Ni/Ag Metallization in Solder Joints”, Acta Mater., 49, 2609-2624 (2001) (Experimental, Crys. Structure, Interface Phenomena, 82) Purdy, G.R., Malakhov, D.V., Guha, A., “Homogenization of Multicomponent Alloys via Partial Melting”, J. Phase Equilib., 22(4), 439-450 (2001) (Theory, Phase Relations, Thermodyn., 49) Young, C.C., Duh, J.G., Tsai, S.Y., “Microstructural Evolution in the Sn-Cu-Ni and Pb-Sn Solder Joints with Cu and Pt-Ag Metallized Al2O3 Substrate”, J. Electron. Mater., 30(9), 1241-1248 (2001) (Electronic Structure, Experimental, Crys. Structure, Morphology, Transport Phenomena, 20) Kattner, U.R., “Phase Diagrams for Lead-Free Solder Alloys”, JOM, 12, 45-51 (2002) (Phase Diagram, Phase Relations, Review, 47) Lebrun, N., “Cu-Ni (Copper-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14832.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, #, *, 51) Lin, Ch.-H., Chen, S.-W., Wang, Ch.-H., “Phase Equilibria and Solidification Properties of Sn-Cu-Ni Alloys”, J. Electron. Mater., 31(9), 907-915 (2002) (Experimental, Morphology, Phase Diagram, Phase Relations, #, *, 12) Zeng, K., Tu, K.N., “Six Cases of Reliability Study of Pb-Free Solder Joints in Electronic Packaging Technology”, Mater. Sci. Eng. R, 38, 55-105 (2002) (Mechan. Prop., Morphology, Phase Diagram, Phase Relations, Review, Thermodyn., 122) Chen, S.-W., Wu, S.-H., Lee, S.-W., “Interfacial Reactions in the Sn-(Cu)/Ni, Sn-(Ni)/Cu, and Sn/(Cu,Ni) Systems”, J. Electron. Mater., 32(11), 1188-1194 (2003) (Experimental, Phase Diagram, Phase Relations, #, *, 19) Chen, S.-W., Lee, S.-W., Yip, M.-C., “Mechanical Properties and Intermetallic Compound Formation at the Sn/Ni and Sn-0,7wt.%Cu/Ni Joints”, J. Electron. Mater., 32(11), 1284-1289 (2003) (Experimental, Mechan. Prop., 22) Miettinen, J., “Thermodynamic Description of the Cu-Ni-Si System at the Cu-Ni Side”, Calphad, 27(3), 309-318 (2003) (Assessment, Phase Diagram, Phase Relations, Thermodyn., #, *, 41) Wang, C.-H., Chen, S.-W., “Isothermal Section of the Ternary Sn-Cu-Ni System and Interfacial Reactions in the Sn-Cu/Ni Couples at 800°C”, Metall. Mater. Trans. A, 34(10), 2281-2287 (2003) (Experimental, Phase Diagram, Phase Relations, Morphology, Kinetics, #, *, 12) Ghosh, G., “Reactive Interdiffusion between a Lead-free Solder abd Ti/Ni/Ag Thin-Film Metallizations”, J. Electron. Mater., 33, 229-240 (2004) (Experimental, Crys. Structure, Kinetics, 23) Jang, G.-Y., Huang, C.-S., Hsiao, L.-Y., Duh, J.-G., Takahashi, H., “Mechanism of Interfacial Reaction for the Sn-Pb Solder Bump with Ni/Cu Under-Bump Metallization in Flip-Chip Technology”, J. Electon. Mater., 33(10), 1118-1129 (2004) (Experimental, Phase Relations, 24) Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn [2004Li]
[2004Liu1]
[2004Liu2]
[2004Lou]
[2004Lue]
[2004Sah]
[2004Yoo]
[2005Jan]
[2005Li]
[2005Wan]
[2006Hsi]
[2006Li]
Landolt-Börnstein New Series IV/11C3
317
Li, C., Chin, Y.L., Wu, P., “Correlation between Bulk Modulus of Ternary Intermetallic Compounds and Atomic Properties of Their Constituent Elements”, Intermetallics, 12, 103-109 (2004) (Electronic Structure, Thermodyn., 24) Liu, X.-J., Wang, C.P., Ohnuma, I., Kainuma, R., Ishida, K., “Experimental Investigation and Thermodynamic Calculation of the Phase Equilibria in the Cu-Sn and Cu-Sn-Mn System” (in English) Metall. Mater. Trans. A, 35A, 1641-1654 (2004) (Experimental, Phase Diagram, 55) Liu, H.S., Wang, J., Jin, Z.P., “Thermodynamic Optimization of Ni-Sn Binary System”, Computer Coupling of Phase Diagrams and Thermochemistry, 28, 363-370 (2004) (Phase Diagram, Thermodyn., 17) Lourenco, N., Santos, H., “Study of Transformations in a Cu-9Ni-6Sn Alloy by Differential Scanning Calorimetry (DSC)”, Mater. Sci. Forum, 455-456, 262-266 (2004) (Experimental, Phase Relations, 6) Luef, Ch., Flandorfer, H., Ipser, H., “Lead-free Solder Materials: Experimental Enthalpies of Mixing in the Ag-Cu-Sn and Cu-Ni-Sn Ternary Systems”, Z. Metallkd., 95(3), 151-163 (2004) (Experimental, Thermodyn., *, 11) Sahu, P., Pradhan, S.K., De, M., “X-Ray Diffraction Studies of the Decomposition and Microstructural Characterization of Cold-Worked Powders of Cu-15Ni-Sn Alloys by Rietveld Analysis”, J. Alloys Compd., 377, 103-116 (2004) (Crys. Structure, Experimental, 42) Yoon, J.-W., Lee, Y.-H., Kim, D.-G., Kang, H.-B., Suh, S.-J., Yang, C.-W., Lee, C.-B., Jung, J.-M., Yoo, C.-S., Jung, S.-B., “Intermetallic Compound Layer Growth at the Interface Between Sn-Cu-Ni Solder and Cu Substrate”, J. Alloys Compd., 381, 151-157 (2004) (Experimental, Morphology, 21) Jang, G.-Y., Duh, J.-G., “The Effect of Intermetallic Compound Morphology on Cu Diffusion in Sn-Ag and Sn-Pb Solder Bump on the Ni/Cu Under-Bump Metallization”, J. Electon. Mater., 34(1), 68-79 (2005) (Experimental, Phase Relations, Kinetics, Morphology, 25) Li, C.-Y., Duh, J.-G., “Phase Equilibria in the Sn-rich Corner of the Sn-Cu-Ni Ternary Alloy System at 240°C”, J. Mater. Res., 20(11), 3118-3124 (2005) (Experimental, Phase Relations, #, *, 25) Wang, C.P., Liu, X.J., Jiang, M., Ohnuma, I., Kainuma, R., Ishida, K., “Thermodynamic Database of the Phase Diagrams in Copper Base Alloy Systems”, J. Phys. Chem. Solids, 66(2-4): 256-260 (2005) (Theory, Thermodyn., 14) Hsiao, L.-Y., Kao, S.-T., Duh, J.-G., “Characterizing Metallurgical Reaction of Sn3.0Ag0.5Cu Composite Solder by Mechanical Alloying with Electroless Ni-P/Cu Under-Bump Metallization after Various Relow Cycles”, J. Electron. Mater., 35(1), 81-88 (2006) (Experimental, Phase Relations, 23) Li, C.-Y., Chiou, G.-J., Duh, J.-G., “Phase Distribution and Phase Analysis in Cu6Sn5, Ni3Sn4 and the Sn-rich Corner in the Ternary Sn-Cu-Ni Isotherm at 240°C”, J. Electron. Mater., 35(2), 343-352 (2006) (Experimental, Phase Relations, #, *, 31)
MSIT®
Cu–Ni–Sn
318
Table 1: Investigations of the Cu-Ni-Sn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1928Pri]
Optical microscopy, hardness
200 - 800°C; 8.9 - 76.5 mass% Ni and 6.5 - 22 mass% Sn
[1932Gui]
Optical metallography, hardness 150 - 450°C; 50 - 70 mass% Ni and 0.5 - 15 mass% Sn
[1932Ves]
Thermal analysis, optical metallography
[1933Eas]
Thermal analysis, XRD, optical < 1140°C; 1-20 mass% Ni and 1-31 mass% Sn microscopy
[1937Rah]
XRD
Cu45Ni30Sn25 (D03)
[1940Glu]
Optical metallography, XRD
630°C; (Cu-Ni) alloy/liquid-Sn
[1953Maz1]
XRD
900°C; 26.8 - 34.9 at.% Ni and 54-58.7 at.% Sn
[1953Maz2]
XRD
800 - 900°C; 6-38.5 at.% Ni and 36.1-55.5 at.% Sn
[1971Dok]
Mössbauer spectroscopy
NiCu2Sn (L21)
[1974Sch1], [1974Sch2]
TEM and XRD
350°C; 9 mass% Ni and 6 mass% Sn
[1975Ple]
Resistivity
300 - 825°C; 3.5-12 mass% Ni and 2.5-8 mass% Sn
[1977Hel]
XRD
200 - 450°C; 8 mass% Ni and 2-5.46 mass% Sn
[1978Ger]
TEM, XRD and hardness
330 - 780°C; 9 mass% Ni and 1-6 mass% Sn
[1978Hel]
TEM and XRD
8 mass% Ni and 5 mass% Sn
[1978Lef]
TEM and XRD
300 - 550°C; 15 mass% Ni and 8 mass% Sn
[1979Bab]
TEM
400 - 825°C; 9 mass% Ni and 6 mass% Sn
[1979Poo1]
SETARAM high-temperature calorimeter
1307°C; xCu:xSn=1/9, 1/5, 1/3, 1/2, 2/3, 1/1, 2/1, 3/2, 3/1, 5/1, 9/1; xNi:xSn=1/1, 3/1, 1/3; xCu:xNi=3/1
[1979Sch]
DTA, XRD and resistivity
700°C; NiCu2Sn
[1980Dit]
TEM and XRD
350°C, 10 mass% Ni and 6 mass% Sn
[1980Ray]
TEM and XRD
350 - 600°C; 9 mass% Ni and 6 mass% Sn
[1981Wat]
DTA, optical microscopy and XRD
300 - 850°C; CuxNi75–xSn25 with 5 x 30
[1982Mik]
Optical microscopy, TEM and XRD
350 - 500°C; 10 mass% Ni and 8 mass% Sn
[1982Ray]
TEM and XRD
350 - 600°C; 9 mass% Ni and 6 mass% Sn
MSIT®
< 1050°C; 2-22.5 mass% Ni and 5-40 mass% Sn
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
319
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1983Mik1]
Optical microscopy, TEM and XRD
350-850°C; 10-90 mass% Ni and 8-40 mass% Sn
[1983Mik2]
Optical microscopy and SEM
450-850°C; 10-80 mass% Ni and 8 mass% Sn
[1983Mur]
TEM
< 877°C; Cu8Ni67Sn25
[1984Kra]
TEM
300-400°C; 9 mass% Ni and 2-5 mass% Sn
[1984Mik1]
Optical microscopy, TEM and XRD
[1984Mik2]
Optical microscopy and SEM
450-850°C; 10-80 mass% Ni and 8 mass% Sn
[1984Sco]
Optical microscopy and TEM
350-400°C; 15 mass% Ni and 8 mass% Sn
[1985Nar]
TEM
350-850°C; 8.9 mass% Ni and 6 mass% Sn
[1986Tak]
SEM
750-850°C; 3.5-17.7 at.% Ni and 2.4-7.1 at.% Sn
[1987Kra]
TEM
300-500°C; 9 mass% Ni and 6 mass% Sn
[1988Col]
Optical microscopy and SEM
300-500°C; 10-15 mass% Ni and 2-8 mass% Sn
[1988Sat]
TEM
350°C; 10 mass% Ni and 6 mass% Sn
[1989Pak]
Optical microscopy, DTA, TEM < 1000°C; CuxNi75–xSn25 with 10 x 22 and XRD
[1990Che]
Optical microscopy, DTA and XRD
[1990Dey]
Optical microscopy, EPMA and 400-825°C; 10-15 mass% Ni and 6-12 mass% Sn XRD
[1990Gou]
XRD
350°C; 15 mass% Ni and 8 mass% Sn
[1990Pak]
DTA, TEM and XRD
CuxNi75–xSn25 with 11 x 22
[1991Mik]
Optical microscopy
450°C; 10 mass% Ni and 8 mass% Sn
[1992Hun]
Optical microscopy and SEM
350-775°C; 15 mass% Ni and 8 mass% Sn
[1992Kra]
XRD and TEM
900°C; Ni2CuSn (D03)
[1994Her]
Optical microscopy and XRD
400°C; 15 mass% Ni and 8 mass% Sn
[1994Mik]
Optical microscopy, TEM and XRD
450-850°C; 10 mass% Ni and 8 mass% Sn
[1995Bur]
DTA
850-1400°C; (Ni2Cu)1–xSnx with 0.11 x 0.22
[1995Kra]
TEM/AEM and XRD
490-800°C; 50.2-65.9 at.% Ni and 3.4-24.8 at.% Sn
[1995Pal]
XRD and SEM
450°C; 5-22 mass% Ni and 2.7-13 mass% Sn
[1996Oel]
AEM/TEM and XRD
850°C; (Cu0.33Ni0.67)1–xSnx, (Cu0.4Ni0.6)1–xSnx 0 x 0.3
[1996Rud]
Dilatometry and XRD
850°C; Ni2CuSn and (CuNi2)0.775Sn0.225
Landolt-Börnstein New Series IV/11C3
< 1100°C; 14 at.% Ni and 10-22 at.% Sn
MSIT®
Cu–Ni–Sn
320 Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1996Spr]
Electrical resistivity, XRD
< 900°C; (Ni2Cu)1–xSnx with 0 x 0.143
[1997Spr]
Magnetometry
< 700°C; (Ni2Cu)1–xSnx with 0 x 0.07
[1997Vir2]
Electrical resistivity, TEM/AEM 310-350°C; 2-8 mass% Ni and 7-9 mass% Sn
[1998Vir]
SEM and TEM
800-900°C; 6-9 mass% Ni and 6-7 mass% Sn
[1998Zha1]
TEM
225-840°C; 15 mass% Ni and 8 mass% Sn
[1998Zha2]
TEM
150-800°C; 7.5 mass% Ni and 5 mass% Sn
[1999Jun]
Resistivity
300-850°C; 9 mass% Ni and 6 mass% Sn
[1999Kim]
TEM
350°C; 9 mass% Ni and 6 mass% Sn
[1999Rhu]
Optical microscopy, SEM and TEM
350-850°C; 9 mass% Ni and 6 mass% Sn
[2001Che]
Diffusion couples
160-200°C
[2001You]
Diffusion couples
100-170°C
[2002Lin]
DTA, EPMA and XRD
240-300°C; 0.5-70.5 at.% Ni and 20.7-99.5 at.% Sn
[2003Che1]
Diffusion couples
200°C
[2003Wan]
Optical metallography and EPMA
800°C; < 80 at.% Cu and < 75 at.% Ni
[2004Lou]
Differential scanning calorimetry
830°C; 9 mass% Ni and 6 mass% Sn
[2004Lue]
SETARAM high-temperature calorimeter
1250°C; xCu:xNi=1/3, 1/1, 3/1
[2004Sah]
XRD
400°C; 15 mass% Ni and 5-13 mass% Sn
[2004Yoo]
Diffusion couples
80-150°C
[2005Li] [2006Li]
XRD, EPMA
240°C; 10-25 at.% Cu, 10-30 at.% Ni, bal. Sn
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
321
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, (Ni,Cu)
cF4 Fm3m Cu
(Ni) 1455 (Cu) 1084.62
Lattice Parameters [pm]
Comments/References
a = 352.32
pure Ni at 25°C [V-C2]
a = 361.46 a = 362.25
pure Cu at 25°C [V-C2] as-quenched Cu-5Ni-2.7Sn(mass%), at 25°C [1995Pal] as-quenched Cu-15Ni-2.7Sn(mass%), at 25°C [1995Pal] as-quenched Cu-22Ni-2.7Sn(mass%), at 25°C [1995Pal] as-quenched Cu-15Ni-5Sn(mass%), at 25°C [1995Pal] as-quenched Cu-15Ni-9Sn(mass%), at 25°C [1995Pal] as-quenched Cu-15Ni-13Sn(mass%), at 25°C [1995Pal] as-quenched Cu-65.9Ni-3.4Sn(at.%), at 25°C [1996Spr] as-quenched Cu-62.0Ni-5.5Sn(at.%), at 25°C [1996Spr]
a = 360.90 a = 359.87 a = 362.04 a = 363.56 a = 365.20 a = 358.06 a = 360.10 (Sn) 231.5 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
pure Sn at 25°C [V-C2]
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
pure Sn [V-C2]
1, Cu17Sn3 798 - 586
cI2 Im3m W
a = 297.81 to 298.71 0.131 xSn 0.158 [V-C2]
J, Cu3Sn(r) < 676
oC80 Cmcm Cu3Sn
, Cu41Sn11 590 - 350
cF416 F43m Cu41Sn11
a = 1798 a = 1801
, Cu10Sn3 640 - 582
hP26 P63 -
a = 733.0 a = 784.0
Landolt-Börnstein New Series IV/11C3
a = 552.9 b = 4775.0 c = 432.3
0.245 xSn 0.259 [V-C2] at 25 at.% Sn [V-C2]
0.203 xSn 0.208 [V-C2], at 20.5 at.% Sn NiCu9Sn3, at 25°C [1977Boo] 0.209 xSn 0.225 [V-C2], at 23.1 at.% Sn [V-C2]
MSIT®
Cu–Ni–Sn
322 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, Cu6Sn5(h) 415 - 186
mP36 P21/c Cu5Sn4
mC54 C2 Cu5Sn4 ’, Cu6Sn5(r) < 186
mP44 C2/c Cu6Sn5
2, Ni3Sn(r) < 977
hP8 P63/mmc NiSn3
1, Ni3Sn2(h) 1264 - 600
hP4 P63/mmc NiAs
2, Ni3Sn2(r) < 794.5
oP20 Pnma Ni3Sn2
Ni3Sn4 < 794.5
mC14 C2/m Ni3Sn4
MSIT®
Lattice Parameters [pm]
a = 983.0 b = 727.0 c = 983.0 = 62.5° a = 1260.0 b = 727.0 c = 1020.0 = 90° a = 1103.6 b = 728.8 c = 984.0 = 98.81° a = 530.5 c = 425.4
Comments/References
0.435 xSn 0.445 [Mas2], superstructure of B81-NiAs type, at 44.44 at.% Sn [1995Lar]
superstructure of B81-NiAs type, at 44.44 at.% Sn [1995Lar]
0.435 xSn 0.445 [Mas2], superstructure of B81-NiAs type, at 44.44 at.% Sn [1994Lar]
0.241 xSn 0.260 [Mas2], at 25 at.% Sn [V-C2]
a = 530.5 c = 425.4
Cu1.92Ni73.08Sn25 quenched from 600°C, at 25°C [1981Wat]
a = 530.5 c = 425.4
Cu3.99Ni71.01Sn25 quenched from 600°C, at 25°C [1981Wat]
a = 530.5 c = 425.4
Cu5Ni70Sn25 quenched from 600°C, at 25°C [1981Wat]
a = 414.6 c = 525.3 a = 711.2 b = 521.1 c = 823.2 a = 1221.4 b = 406.02 c = 521.93 = 105.01°
0.388 xSn 0.425 [Mas2], at 40 at.% Sn [V-C2] 0.241 xSn 0.260 [Mas2], at 40 at.% Sn [V-C2]
0.555 xSn 0.57 [Mas2], [V-C2]
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
1, (Ni,Cu)3Sn
cF16 Fm3m BiF3
Cu3Sn(h) 755 - 520
323
Lattice Parameters [pm]
a = 606.05 to 611.76 0.165 xSn 0.279 [V-C2]
Ni3Sn(h) 1174 - 850
a = 594.6
Cu45Ni30Sn25, at 25°C [1937Rah]
a = 598.2
0.2325 xSn 0.2725 [Mas2], at 25 at.% Sn [V-C2]
a = 592.93
Cu20.95Ni54.05Sn25 quenched from 600°C, at 25°C [1981Wat] Cu21.98Ni53.02Sn25 quenched from 600°C, at 25°C [1981Wat] Cu22.98Ni52.02Sn25 quenched from 600°C, at 25°C [1981Wat] Cu25Ni50Sn25 quenched from 600°C, at 25°C [1981Wat]
a = 593.04 a = 593.20 a = 593.41 * -1, NiCu2Sn 700 - 500
* -2, Ni5CuSn2 < 740
* -3, Ni2CuSn < 460
Comments/References
cF16 Fm3m AlCu2Mn
oP8 Pmmn Cu3Ti
a = 579.7
[1953Maz2]
a = 596.4
[1979Sch]
a = 596.0
[1984Wac] CuxNi5–xSn2, with 0.8 x 1.41 [1981Wat] at x = 1 [1989Pak]
a = 449.5 b = 538.0 c = 428.5
aP?
a = 453.0 b = 531.0 c = 434.0 = 85° = 86° = 84°
[1990Pak]
Table 3: Invariant Equilibria Reaction
L + 1 + 1-(Ni,Cu)3Sn
T [°C]
Type
Phase
Composition (at.%) Cu
Ni
Sn
~780
U1
L
77.8
5.5
16.7
L + 1-(Ni,Cu)3Sn J-Cu3Sn + ~620 Ni3Sn2(h)
U2
L
44.5
12.7
42.8
L + J-Cu3Sn + Ni3Sn2(r)
U3
L
12.6
2.6
84.8
Landolt-Börnstein New Series IV/11C3
~400
MSIT®
Cu–Ni–Sn
324 T [°C]
Reaction
Type
Phase
Composition (at.%) Cu
Ni
Sn
L + Ni3Sn4 (Sn) + Ni3Sn2(r)
~230
U4
L
1.4
2.7
95.9
L + Ni3Sn2(r) (Sn) +
~228
U5
L
2.3
1.1
96.6
Table 4: Solidus and Spinodal Temperatures of Cu Rich Alloys Alloy Composition (mass%)
Solidus Temperature [°C] (±5)
Spinodal Temperature [°C] (20)
Cu
Ni
Sn
93.57
3.41
3.02
617
360
90.41
4.89
4.70
692
410
86.14
5.98
7.88
770
450
85.22
9.03
5.75
740
464
84.93
10.41
4.66
751
450
81.32
11.46
7.22
816
490
79.01
14.04
6.95
780
480
Table 5: Investigations of the Cu-Ni-Sn Materials Properties Reference
Method/Experimental Technique Type of Property
[1974Sch1], [1974Sch2]
Tensile test
Tensile properties of spinodal microstructure
[1975Ple]
Tensile test
Tensile properties of spinodal microstructure
[1978Lef]
Tensile test
Tensile properties of spinodal microstructure
[1984Kra]
Tensile test
Yield stress
[1984Sco]
Electrical and mechanical tests
Young’s and shear moduli, yield and ultimate tensile stress, electrical conductivity
[1986Kat]
Tensile test
Tensile properties of single crystal at 25°C
[1986Kha]
Tensile test
Tensile properties of single crystal at 25°C
[1989Kra]
Tensile test
Tensile properties
[1991Abd]
Tensile test
Hot deformation behavior between 570 and 850°C
[1991And]
Tensile test
Tensile properties, stress relaxation
[1993Sat]
Tensile test
Tensile properties
[1992Bog]
Electrical, mechanical and thermal tests
Hardness, Young’s modulus, thermal conductivity, thermal expansion coefficient
[1992Kra]
Creep tests
Creep between 300 and 900°C
[1994Her]
Tensile test
Hardness and tensile properties
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
325
Reference
Method/Experimental Technique Type of Property
[1996Leh]
Electrical and mechanical tests
Tensile properties and electrical conductivity
[1997Her]
Tensile test
Tensile properties
[1997Vir1]
Tensile test
Stress relaxation
[1997Vir2]
Electrical
Electrical resistivity
[1999Jun]
Electrical and mechanical tests
Tensile properties and electrical conductivity
[1999Kim]
Tensile test
Tensile properties
[1999Mor]
Electrical and mechanical tests
Tensile properties and electrical conductivity
[1999Rhu]
Tensile test
Tensile properties
[2001Che]
Electrical test
Electromigration
[2003Che2]
Tensile test
Tensile properties of solder joints
Cu-Sn
Cu-Ni-Sn
A-B-C
Ni-Sn 1160 e1 L γ1 + λ1 1130 e2 L (Ni) + γ1
798 p1 L + (Cu) β1 ca.780 755 p3 L + β1 γ1
L + γ β1 + γ1
794.5 p2 L + λ1 Ni3Sn4
U1
γ + β1 + γ1
L+β1+γ1
640 e3 γ1 L + ε ca.620
L + γ1 ε + λ1
γ'+ε+λ1
415 p4 L + ε η
ca.400
U2 L+ε+λ1
L + ε η + γ2
U3 231.15 e4 L Ni3Sn4+(Sn)
ε+η+γ2 ca.230 L + Ni3Sn4 (Sn) + γ2
L+γ2+η
γ2+Ni3Sn4+(Sn)
ca.228 227 e5 L (Sn) + η
L+(Sn)+η
U4
L+γ2+(Sn)
L + Ni3Sn2 (Sn) + η
U5
(Sn)+η+γ2
Fig. 1: Cu-Ni-Sn. Reaction scheme for the solidification of Cu-Ni-Sn alloys Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Ni–Sn
326
Sn Fig. 2: Cu-Ni-Sn. Liquidus surface projection
e4
U4
(β Sn) e5 U5
Ni3Sn4 p2
Axes: at.%
η p4
U3
20
Data / Grid: at.%
80
727 827
ε
40
9 9 27 10 77 2 10 7 77
60
87 7
727
27 11
60
γ 1 1077
1127
80
1027
e2
1327
1427 °C 20
Ni
γ
p3
927
U1
20
β1
p1
1127 60
Ni Cu Sn
977
1227
40
Fig. 3: Cu-Ni-Sn. Isothermal section of Cu corner at 1090°C
40
U2
Ni3Sn2 1177 e1
e3
827
80
0.00 88.00 12.00
Cu
data curves & grid: at.% axes scaling: at.%
10
5
5
10
L (Cu) Ni Cu Sn
MSIT®
12.00 88.00 0.00
90
95
Cu
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn Ni Cu Sn
Fig. 4: Cu-Ni-Sn. Isothermal section of Cu corner at 1050°C
327 0.00 88.00 12.00
data curves & grid: at.% axes scaling: at.%
10
5
5
L
10
(Cu) Ni Cu Sn
90
12.00 88.00 0.00
95
Ni Cu Sn
Fig. 5: Cu-Ni-Sn. Isothermal section of Cu corner at 1025°C
Cu
0.00 88.00 12.00
data curves & grid: at.% axes scaling: at.%
10
5
5
L
10
(Cu) Ni Cu Sn Landolt-Börnstein New Series IV/11C3
12.00 88.00 0.00
90
95
Cu MSIT®
Cu–Ni–Sn
328
Sn
Data / Grid: at.% Axes: at.%
Fig. 6: Cu-Ni-Sn. Isothermal section at 800°C 20
80
L
40
60
60
λ1
40
λ 1+γ 1+L
λ 1+γ 2+γ 1
γ2 80
γ1
γ +γ 2+γ 1
20
L+γ +γ 1
γ 20
Ni
40
60
Ni Cu Sn
Fig. 7: Cu-Ni-Sn. Partial isothermal section at 700°C
80
0.00 30.00 70.00
10
Cu
Data / Grid: at.% Axes: at.%
60
20
50
30
40
40
30
γ1
50
20
γ +γ 1+β 1 β1
60
10
γ Ni Cu Sn
MSIT®
70.00 30.00 0.00
40
50
60
70
80
90
Cu
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
329
Sn
Data / Grid: at.% Axes: at.%
Fig. 8: Cu-Ni-Sn. Partial isothermal section at 550°C 20
80
40
60
60
40
τ2
γ 2+τ 2 80
γ1
γ 1+τ 2
γ2 γ +γ 2 γ 20
Ni
20
γ +γ 1+τ 2 γ +γ 2+τ 2
40
γ +τ 2
γ +γ 1
60
80
Sn Fig. 9: Cu-Ni-Sn. Isothermal section at 240°C
Data / Grid: at.% Axes: at.%
L
20
40
Ni3Sn4
λ2
60
80
60
L+Ni3Sn4+η
λ 2+Ni3Sn4+η
η 40
γ 2+τ 2+λ 2
τ2
γ2
τ 2+ε+λ 2
λ 2+ε+η ε
80
Landolt-Börnstein New Series IV/11C3
20
20
τ 2+ε+γ
γ 2+τ 2+γ
Ni
Cu
40
60
80
γ
Cu
MSIT®
Cu–Ni–Sn
330
Sn
20
λ2 γ 2+τ 2+λ 2
80
60
λ 2+Ni3Sn4+η
60
η 40
τ2
γ2
τ 2+ε+λ 2
λ 2+η+ε ε
80
Temperature, °C
20
τ 2+ε+γ
γ 2+τ 2+γ
20
Ni
Axes: at.%
Ni 3 Sn 4 +(β Sn) +η
40
Ni3Sn4
Fig. 11: Cu-Ni-Sn. Vertical section at a constant Ni content of 2 mass%, plotted in at.%
Data / Grid: at.%
(β Sn)
Fig. 10: Cu-Ni-Sn. Isothermal section at 200°C
40
60
80
γ
Cu
1100
L 1000
L+γ 900
L+β 1+γ1 800
γ
γ+β 1
L+β 1 γ1+β 1+L
β1 γ1+β 1
700
L+γ1
γ1
600
2.16 Ni Cu 97.84 Sn 0.00
MSIT®
γ+γ1
γ+γ1+β 1
10
20
Sn, at.%
2.65 Ni Cu 71.10 Sn 26.25
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
331
1100
Fig. 12: Cu-Ni-Sn. Calculated vertical section at a constant Ni content of 2 mass%, plotted in at.%
L
Temperature, °C
1000
L+γ 900
L+β 1 L+β1+γ1
800
γ
γ+β 1
700
β1
γ+γ1+β 1
γ1+β 1
γ1
γ+γ1
600
2.16 Ni Cu 97.84 Sn 0.00
10
2.65 Ni Cu 71.10 Sn 26.25
20
Sn, at.%
Fig. 13: Cu-Ni-Sn. Vertical section at a constant Ni content of 4 mass%, plotted in at.%
Temperature, °C
L 1000
L+γ
L+γ+β 1
γ 750
γ+β 1 γ+γ1+β 1
L+γ1+β 1
L+β 1
β1
L+γ1
γ1+β 1 γ1
γ+γ1
4.32 Ni Cu 95.68 Sn 0.00
Landolt-Börnstein New Series IV/11C3
10
20
Sn, at.%
5.30 Ni Cu 68.51 Sn 26.20
MSIT®
Cu–Ni–Sn
332
1200
Fig. 14: Cu-Ni-Sn. Calculated vertical section at a constant Ni content of 4 mass%, plotted in at.%
1100
Temperature, °C
L 1000
L+γ 900
L+γ+β 1 800
L+γ1+β 1
γ
L+γ1
γ+γ1+β 1 700
γ+γ1
600
4.32 Ni Cu 95.68 Sn 0.00
Temperature, °C
Fig. 15: Cu-Ni-Sn. Vertical section at a constant Ni content of 7 mass%, plotted in at.%
γ1+β 1 γ1
10
Sn, at.%
1100
L
1000
L+γ 900
γ
γ1+L
L+γ+γ1
800
γ+γ1 700
7.54 Ni Cu 92.46 Sn 0.00
MSIT®
5.30 Ni Cu 68.51 Sn 26.20
20
10
γ1 20
Sn, at.%
9.24 Ni Cu 64.64 Sn 26.12
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
Temperature, °C
Fig. 16: Cu-Ni-Sn. Vertical section at a constant Ni content of 10 mass%, plotted in at.%
333
1200
1100
L 1000
L+γ 900
γ1+L
γ
L+γ+γ1
800
γ1
γ+γ1 700
Ni 10.74 Cu 89.26 Sn 0.00
Temperature, °C
Fig. 17: Cu-Ni-Sn. Vertical section at a constant Ni content of 14 at.%
10
Ni 13.17 Cu 60.80 Sn 26.04
20
Sn, at.%
1200
1100
L
1000
L+γ 900
γ
L+γ1
L+γ+γ1 800
γ+γ1
γ1
700
Ni 14.00 Cu 86.00 Sn 0.00
Landolt-Börnstein New Series IV/11C3
10
20
Sn, at.%
Ni 14.00 Cu 61.00 Sn 25.00
MSIT®
Cu–Ni–Sn
334
Temperature, °C
Fig. 18: Cu-Ni-Sn. Vertical section at a constant Ni content of 15 mass%, plotted in at.%
1200
L 1100
1000
L+γ
900
L+γ+γ1
L+γ1
800
γ1 700
Ni 16.04 Cu 83.96 Sn 0.00
5
10
Ni 18.60 Cu 63.00 Sn 18.40
15
Sn, at.%
1200
Fig. 19: Cu-Ni-Sn. Calculated vertical section at a constant Ni content of 15 mass%, plotted in at.%
L 1100
Temperature, °C
L+γ 1000
900
L+γ+γ1 L+γ1
γ
800
γ1
700
600
Ni 16.04 Cu 83.96 Sn 0.00
MSIT®
γ+γ1+τ2 γ+τ2
10
20
Sn, at.%
Ni 19.65 Cu 54.44 Sn 25.91
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Sn
335
1250
Fig. 20: Cu-Ni-Sn. Vertical section at a constant Sn content of 25 at.%
L
1000
L+γ1
Temperature, °C
γ1 γ1+γ2
750
γ1+τ2
500
γ2
γ1+τ2+τ3
250
τ2 γ1+τ3 0
Ni 75.00 Cu 0.00 Sn 25.00
10
20
Cu, at.%
Ni 51.00 Cu 24.00 Sn 25.00
1200
Fig. 21: Cu-Ni-Sn. Calculated vertical section at a constant Sn content of 25 at.%
L 1100
Temperature, °C
1000
900
γ1-Cu3Sn+γ1-Ni3Sn 800
γ1-Cu3Sn+τ2+γ1-Ni3Sn γ1-Cu3Sn
700
γ+γ1-Ni3Sn 600
γ+τ2+γ1-Ni3Sn τ2
ε+γ1-Cu3Sn
τ2+γ1-Ni3Sn 500
Ni 75.00 Cu 0.00 Sn 25.00
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40
Cu, at.%
60
0.00 Ni Cu 75.00 Sn 25.00
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336
1400
Fig. 22: Cu-Ni-Sn. Vertical section at a constant ratio x(Cu):x(Ni)=1:2
L
1300
Temperature, °C
1200
L+γ
1100 1000
L+λ 1 L+γ1
L+γ+γ1
γ
L+γ1+λ 1
γ+γ1
900
γ1+λ 1
800 700
γ+γ2
600 500
Ni 66.70 Cu 33.30 Sn 0.00
Fig. 23: Cu-Ni-Sn. Calculated vertical section at a constant ratio x(Cu):x(Ni)=1:2
10
γ2+λ 1
Ni 46.70 Cu 23.30 Sn 30.00
20
Sn, at.%
1400
L
1300
Temperature, °C
1200 1100
L+γ 1000
γ
900
L+γ1 L+γ+γ1
800 700
γ1 γ+γ1
600 500
Ni 66.70 Cu 33.30 Sn 0.00
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10
20
Sn, at.%
Ni 50.00 Cu 25.00 Sn 25.00
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Cu–Ni–Sn
337
Sn
Data / Grid: at.% Axes: at.%
Fig. 24: Cu-Ni-Sn. Integral enthalpy of mixing of liquid alloys at 1307°C, in kJ/mol of atoms
0
20
80
-4
40
60
60
40
-20 -16 -12
80
20
-8 -4 0 20
Ni
40
60
80
Sn
Cu
Data / Grid: at.% Axes: at.%
Fig. 25: Cu-Ni-Sn. Integral enthalpy of mixing of liquid alloys at 1250°C, in kJ/mol of atoms
0 20
-2
80
40
60
60
40
-20 -18 -16 -14 80
-12
-10 -8
20
-6 -4 -2 0
Ni
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60
80
Cu
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338
Cu–Ni–Zn
Copper – Nickel – Zinc Nathalie Lebrun, Pierre Perrot Introduction In the first telephone exchanges built in 1879, the cables were made of copper and the switches of “nickel silver” (Cu-Ni-Zn alloy) because of its excellent wear and corrosion resistance [1990Seg]. The first liquidus of the Cu-Ni-Zn system was proposed by [1908Taf]. Micrographic observations were reported by [1929Bau2], and isothermal sections at 400, 600 and 800°C were given by [1929Bau1, 1930Bau]. The older information on the metallurgical properties of the alloys was obtained by [1912Car, 1913Gui, 1922Voi, 1924Pri, 1925Gui, 1925Ost, 1926Sma]. The ternary phase diagram has been investigated by [1930Bau, 1934Yam], later reproduced by [1949Jae], then more extensively by [1935Sch]. These diagrams are very similar to those generally accepted and were later confirmed through their intensive use as a basis for diffusion studies at 775°C [1970Tay, 1971Coa, 1977Sis, 1977Wir, 1979Day, 1985Kan]. The most complete work on the phase equilibria in the ternary diagram is that of [1935Sch] whose results were taken into account in the thermodynamic assessment of [2003Mie, 2005Jia]; the isothermal sections at 650°C and 775°C proposed by [1973Lev, 1979Dri] are those of [1935Sch] modified according to the metallographic observations of [1952Haw]. The isothermal section at 775°C was widely confirmed by [1979Day, 1984Kim] which used it to put into evidence the existence of zero flux and flux reversal planes in ternary diffusion experiments, and by [1985Kan] who used it to extend this observation to the quaternary systems. [1936Sch] proposed a low temperature isothermal section accepted by [1956Fre, 1973Lev, 1979Dri]. The equilibria between and solid solutions have been investigated by [1930Bau] at 600 and 800°C, [1952Haw] at 672°C and more extensively by [1979May] at 600°C. Investigations related to phase equilibria, structure and thermodynamics are reported in Table 1. Binary Systems The binary systems Cu-Zn and Ni-Zn present three isotypical phases , and in which copper and nickel may replace each other. The phases present two polymorphic varieties. The label is used for the cP2 structure which corresponds to the high temperature form for (Ni,Zn) and to the low temperature form for (Cu,Zn). After [Mas2], the binary phase diagram Cu-Zn has been calculated by [1993Kow] and assessed by [1994Mio, 2001Her, 2003Mie]. [1993Kow] incorporates in its thermodynamic evaluation an explicit description of the ordering in the phase and a four sublattice model applied to the phase. More realistic solid phase boundaries than [1994Mio] have been proposed for the and phases and have been retained in its assessment. The Cu-Zn phase diagram accepted from [2006Leb] is a compilation of [1993Kow, 1994Mio, 2001Her]. The binary phase diagrams Cu-Ni and Ni-Zn were respectively accepted from [2002Leb] and [1991Nas]. Solid Phases Crystallographic data of solid phases are reported in Table 2. X-ray diffraction measurements of (Ni,Cu,Zn) alloys show the appearance of superstructure lines with a maximum for the composition NiCu2Zn indicating a long-range order whose theory has been presented by [1994Fon, 1996Alt]. The long-range order, overwhelming from measurements of physical properties such as electrical resistivity, Hall coefficient [1967Liv], heat capacity, thermal expansion and elastic modulus was also observed by [1965Hir] from neutron diffraction study of a single crystal of NiCu2Zn. The superstructure is L12 (AuCu3 type), stable between 325 and 501°C, in which Cu and Ni atoms are statistically distributed in the equivalent positions (1/2, 1/2, 0), whereas zinc atoms occupy the equivalent (0,0,0) position [1969Bar, 1981Weg]. Thermal neutron scattering and electron microscopy measurements [1976Vri] show that, below 450°C, ordering of NiCu2Zn occurs on four interpenetrating simple cubic sublattices: one with Zn, one with Ni and the other two with Cu. [1979Hos] uses the interatomic potentials to predict a value of ordering energy and
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Landolt-Börnstein New Series IV/11C3
Cu–Ni–Zn
339
the critical temperature in the NiCu2Zn alloy taking into account nearest neighbor interactions. Tetragonal distortion of the lattice upon ordering appears very small. [1960Cha] proposes for the lattice parameter of the solid solution the following expression a (pm) = 360.75 + 20.5 xZn – 11.2 xNi (xZn < 0.23, xNi < 0.33) Below 325°C, a fully ordered superstructure of the type L10 is stable around the composition NiCu2Zn [1981Roo] in which Zn and Ni atoms occupy two different sublattices whereas Cu occupies the remaining two. Long range order parameter of Zn as function of temperature in NiCu2Zn is determined by neutron diffraction study [1983Weg] and compared with order parameters calculated by Kikuchi’s Cluster Variation Method. Using anomalous scattering of synchrotron radiation, [1985Has] pointed out a short range ordering between Ni and Zn atoms and also between Zn and Cu atoms while a correlation of the opposite direction exists between Cu and Ni atoms. These results are consistent with the latter calculations of [1987Has] and with the prediction of ordering energies given by [1980Roo]. Both ordered structures L10 (AuCu type) and L12 (AuCu3 type) were shown to have important consequences on the mechanical properties: yield stress, hardening and stacking fault energy of the NiCu2Zn alloy [1984Hos]. Monte-Carlo simulations based on effective interactions obtained from first principles calculations [1998Sim] reveal the existence of three ordered phases in NiCu2Zn: a modified L10 (0-600 K), L12 (600-850 K) and a partially ordered (Cu,Zn)(Ni,Cu)-L10 (850-1200 K) which contradicts the generally accepted picture which assumes the existence of only two ordered phases. An increase of resistivity was also measured during slow cooling of the alloy Cu3Ni3Zn2, attributed to a long-range order, which was also detected by X-ray methods [1969Sch]. The alloy Cu3Ni3Zn2 clearly falls in the region and a long-range order of the type A3B is assumed where Cu and Ni atoms are on the A sites and Zn atoms on the B sites. However, no crystallographic evidence is provided. -NixCu1–xZn alloys undergo a martensitic transformation from the CsCl type ordered cubic (B2 structure) to the AuCuI type tetragonal (L10 structure) [1976Shi]. The martensite L10 phase appears at room temperature when x < 0.325. The B2-L10 transition is also observed on the temperature elastic constants curves in -NixCu1–xZn. Liquidus Surface The liquidus surface shown Fig. 1 is based on the works of [1934Yam, 1935Sch]. It was adjusted to be consistent with the accepted binary systems and the Calphad assessments of [2003Mie] in the domain 0.3 < xZn < 0.7 and [2005Jia] in the zinc rich corner. No information concerning the monovariant lines in the zinc rich corner is available. The invariant equilibria from [1979Cha] are given in Table 3. Isothermal Sections The isothermal section at 850°C given in Fig. 2 is taken from [2003Mie]. The isothermal section at 775°C shown in Fig. 3 is a composite proposed by [1979Cha], based on the 775°C diagram given by [1935Sch] and on the isopleths given by [1934Yam]. The phase equilibria at 775°C are confirmed by the diffusion measurements carried out at this temperature in the single-phase domain [1979Day] and in the two-phase + domain [1970Tay, 1971Coa, 1977Sis, 1977Wir]. The isoactivity lines in the domain are from [1977Sis]. The isothermal section at 650°C presented in Fig. 4 is from [1973Lev] and takes into account the Calphad assessment of [2003Mie] carried out for xZn < 0.6. At low temperatures, alloys undergo an ordering towards the composition NiCu2Zn. The ( + L12) two-phase field calculated at 427°C is shown in Fig. 5. Below 335°C for the composition Ni26Cu47Zn27 (325°C for the stoichiometric composition NiCu2Zn) appears the L10 ordered phase as shown in Fig. 6. Figure 6 represents the -L12-L10 equilibria at 327°C from [1982Roo], but data above 30 at.% Zn are not shown here, as they are not consistent with the accepted binary diagrams. Also it should be mentioned that the location of the L10 phase field in Fig. 6 covers the stoichiometric composition NiCu2Zn, but according to the resistivity measuments reported in the same work [1982Roo] at this temperature the L10 phase field should be rather around the composition Ni26Cu47Zn27. The isothermal sections calculated by [1982Roo] at 227and 127°C show the miscibility gap
Landolt-Börnstein New Series IV/11C3
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Cu–Ni–Zn
340
observed in the solid solution at low temperatures. However these sections are not reproduced in the present evaluation, because they are in disagreement with the accepted binary diagrams Cu-Zn and Ni-Zn. Temperature – Composition Sections The cross section Cu0.5Ni0.5–xZnx (0.15 < x < 0.35) showing -L12-L10 equilibria and given in Fig. 7 was obtained through electrical resistivity measurements [1981Roo] and calculated with the tetrahedron approximation of the Kikuchi’s cluster variation method [1980Roo]. The calculated equilibria where ordering occurs are in reasonable agreement with electrical resistivity measurements. Maximum critical temperature for the L12- (A3B-fcc) transition is 784 K (511°C) at an overall composition of Cu47Ni26Zn27. The critical point for the L10-L12 (A2BC-A3B) transition has the same composition and occurs at 608 K (335°C). These critical temperatures are 10 K higher than the critical temperatures of the stoichiometric composition [1981Weg, 1982Roo]. Thermodynamics The activity coefficient of zinc in solid solution (xNi < 0.4, xZn < 0.35) was determined by vapor pressure measurements [1961Cha] and the partial molar quantities calculated at 727°C. The results confirm the existence of a short-range order in copper-zinc alloys and suggest that a long-range order exists in the Cu-Ni-Zn alloys associated with the composition NiCu2Zn. Based on these measurements and available thermodynamic data for the three binary systems, [1977Sis] calculated the activities of Cu, Zn and Ni and the integral Gibbs energy of mixture at 775°C in the and domains. However, they used for their calculation the 775°C isotherm given by [1935Sch] which does not show the order-disorder transformation in the region. Therefore, the activities in the region are doubtful. The activity of Zn in liquid Cu-Ni-Zn alloys have been measured at 1100 and 1200°C by an isopiestic method in the composition range xZn < 0.09 and xNi < 0.08 [1986Sug]. The first and second interaction parameters of Zn in liquid copper were determined as follows:
Ni İ Zn = lim
Ni ȡ Zn =
xZn , xNi
§ ∂ ln ȖZn ¨ → 0 ¨© ∂xNi
· ¸ = − 3.75 at 1150°C and − 3.76 at 1200°C ¸ ¹
§ ∂ 2 ln ȖZn 1 ¨ lim xZn , xNi → 0 ¨© ∂xNi2 2
Ni, Zn ȡ Zn = lim
xZn , xNi
§ ∂ 2 ln ȖZn ¨ → 0 ¨© ∂xNi ∂xZn
· ¸ = 4.68 at 1150°C and 4.47 at 1200°C ¸ ¹
· ¸ = 15.1 at 1150°C and 1200°C ¸ ¹
By equilibrating Cu-Zn solutions at 1200°C with Cu-Ni-Zn alloys, [1987Ryc1, 1987Ryc2] proposes:
Ni İ Zn = lim
xZn , xNi
§ ∂ ln ȖZn ¨ → 0 ¨© ∂xNi
· ¸ = 2.4 ± 2.8 at 1200°C ¸ ¹
However, the assumption of equilibrium between solid and liquid phases is questionable.
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Landolt-Börnstein New Series IV/11C3
Cu–Ni–Zn
341
Notes on Materials Properties and Applications Cu-Zn alloys are the base of materials offering exceptional combination of strength, formability, conductivity and resistance to softening and relaxation, and may be used for high current carrying connectors in the automotive industry [1987Puc]. Nickel acts as a copper replacement element in brass [1912Car, 1926Sma] and is well known to enhance mechanical properties of Cu-Zn alloys such as yield strength [1968Bar], cold working [1912Gui, 1913Gui, 1991Mar] and resilience [1920Gui2] as well as surface properties [1980Ble]. Analogies have been observed between the mechanical behavior of austenitic steel and Cu-Ni-Zn alloys [1914Tho]: decrease of the grain size by increasing the nickel content [1924Pri, 1926Sma], recrystallization after cold working. Annealing improves mechanical properties of Nickel brasses (Cu50Zn45Ni5) [1920Gui1, 1925Gui]; however, a loss of ductility is observed by overheating [1914Tho]. Tensile tests, compression tests, hardness, electrical resistance, cooling curves of as cast and annealed samples have been carried out by [1916Tho]. Structural hardening is always observed after annealing and quenching followed or not by a tempering [1925Gui, 1925Ost]. Alloys close to the composition NiCu2Zn were found to show an anomalous hardening effect [1966Phi] on quenching from 600°C and aging at 400°C. This effect is attributed to the lattice strains caused by clustering rather than to a long-range order hardening which is known to occur by annealing between 300 and 350°C [1980Ito, 1988Ito]. Ni reduces also the quantity of constituent without affecting the hot-working properties of ordinary - brasses [1926Sma]. [1935Kih] investigated tensile properties like fluidity, shrinkage, pressure tightness and established correlations between composition, oxidation resistance, hardness and liquidus temperatures for alloys containing up to 30% Ni and 50% Zn. [1974Mar] pointed out a nonuniformity of the anodic etching of polycrystalline bars of nickel brasses. The rate of the anodic process on the (001) face is 5 times greater than on the (111) face and 3 times greater than on the (101) face. / duplex alloys present a superplastic behavior during tensile straining [1978Liv], whence nucleation and growth of cavities at the / boundaries leading to brittle superplastic fracture. New technologies depend on a large extent on the electrical properties of copper and its alloys with Ni and Zn which play an important role in all communication systems [1990Seg]. These properties have been extensively investigated since [1922Voi, 1957Koe, 1961Phi, 1963Koe] who measured resistivity, Hall effect, thermoelectric effect and electrochemical properties. The combination of high electrical and thermal conductivity, corrosion resistance makes Cu-Ni-Zn alloys useful in various fields such as electronic, electrotechnic and computer industry [1988Pry]. The ordered phase is characterized by higher values of the electrical resistance and the temperature coefficient than disordered state. In both heating and cooling from an initial disordered state, the ordering process is accompanied by an increase of the electrical resistance, known as the K effect [1963Koe]. The electrical resistivity of Cu-Ni-Zn alloys increases during a low temperature anneal, phenomenon also attributed at the appearance of a “K state” that is a special short range ordering [1972Tho, 1973Hor, 1981Ara]. The electrical resistance may change in different way, depending on the conditions of heat treatment and the result may be interpreted by the superposition of two reversible kinetic processes: ordering-disordering and the deposition of a second phase. A maximum of order is observed for the composition Cu5Ni4Zn3.5 [1988Gur]. Electrical conductivity, heat conductivity, heat capacity measurement have been used on Cu65Ni12Zn23 alloys [1973Hor] to explore the influence of the short range ordering on the Debye temperature. Various observations have been reported in solid solutions. Stacking faults induced by cold working were investigated by [1977Hal, 1991Mar] from the detailed analysis of X-ray peak shift, peak asymmetry and peak broadening. In situ analysis of (Ni,Cu,Zn) alloys [1977Sch] allow to identify precipitated particles with a cuboidal shape in a semi-coherent relationship with the matrix. These particles, whose lattice parameter is a = 847 pm are identified with an oxide of spinel structure. Disordered specimens are more sensible to fire cracking than ordered ones [1978Sch]. However, the chemical composition of the alloy and the presence of impurities such as lead are more important than ordering phenomena.
Landolt-Börnstein New Series IV/11C3
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342
Cu–Ni–Zn
Miscellaneous Owing the large extent of the , and solid solutions, the ternary Cu-Ni-Zn system was extensively used for diffusion studies in ternary systems, first by [1940Glu] who investigated the Zn diffusion in Cu-Ni alloys at 630°C, then by [1977Sis, 1979Day, 1984Kim] who used the experimental observations of the diffusion paths in the Cu-Ni-Zn system to propose a general theory of the diffusion in ternary solid solutions. The Zener relaxation in solid solutions provides a powerful tool to study atomic mobility at low temperatures. The Zener relaxation time -r is controlled by the rate of atomic movement and the temperature following: -r = -o exp (–Qr/RT). For alloys of nominal composition Cu70Ni10Zn20, Cu60Ni10Zn30, Cu60Ni20Zn20, the preexponential factor is -o 10–15.50.5 s and the activation energy Qr 190 10 kJ#mol–1 which is also the experimentally measured activation energies for the diffusion of Cu, Ni or Zn into the alloy [1976Ban]. Long range order near the NiCu2Zn composition have been described by two processes [1979Van]: a fast one involving the redistribution of the most mobile constituent, which is the usual Zener effect and a slow one involving a redistribution in the ordered phase. Using first principles together with the density functional theory of atomic short-range order, [1996Alt] calculates, for NiCu2Zn, two ordered states at low temperatures in good agreement with experimental measurements. However, for Ni2CuZn [1997Alt], the quality of the fit is much poorer. Some results regarding solderability, surface roughness, brightness and bendability of tin coated copper alloys strips are reported by [1988Puc]. [1982Ste, 1986Zei] compare hardness of a classical Cu62Ni18Zn20 alloy with the best copper alloys. [1974Sht] investigates partial relaxation process in ternary monophased Cu-Ni-Zn alloys by measuring internal friction and elastic deformation coefficients. The bulk modulus B of the Cu25+xNi25–xZn50 may be evaluated by an empirical formula B = 2V [2004Li] with an error of 16 %, where is the electron density of the lattice and V is the volume of the lattice in m3#mol–1. A corrective parameter taking into account the enthalpy of formation may be introduced, which allows the evaluation of B with an error of 8%. The experimental values are B = 107.8 GPa, 103.1 GPa and 99.3 GPa for x = respectively 0, 10 and 20. [1977Sis] investigated diffusion paths in the and monophased domains at 775°C. Intrinsic diffusion and interdiffusion coefficients were calculated from tracer diffusion coefficients and thermodynamic data. By using diffusion couples in the monophased domain, [1979Day] identified zero flux planes, that are planes in which the interdiffusion flux of a component goes to zero. The composition of zero flux planes correspond to composition points of intersection of diffusion paths and isoactivity lines drawn to the terminal alloys of the couple. [1985Kan] by measuring diffusion profiles, calculation of interdiffusion fluxes, identified zero flux planes for individual components in the domain of the quaternary Cu-Ni-Zn-Mn system. Diffusion in the + region at 775°C was investigated by [1970Tay] who described the morphology of the layers, and measured the diffusion paths by microprobe analysis. Morphological stability of the / interface at 775°C and diffusion coefficients in the composition range of interest (35 to 50 Zn and 0 to 10% Ni) were provided by [1971Coa]. Diffusion paths between a common cubic terminal alloy and a set of alloys were also measured by [1977Wir]. All sets developed a (bcc) phase with two interface: a planar / interface and an / interface showing transitions from planar to non planar. The composition of either side of planar interface were consistent with those based on equilibrium tie lines. References [1908Taf]
[1912Car]
[1912Gui]
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Tafel, V.E., “Studies on the Constitution of Zinc-Copper-Iron alloys and Binary Systems Copper-Nickel, Zinc-Copper and Zinc-Nickel” (in German), Metallurgie, Z. Ges. Huettenkde, 5(12-13), 343-352, 375-383 (1908) (Phase Diagram, Phase Relations, Experimental, #, 27) Carpenter, H.C.H., “The Effect of Other Metals on the Structure of the Constituent in Copper-Zinc Alloys”, J. Inst. Met., 7(1), 59-73 (1912) (Experimental, Morphology, Phase Relations, 2) Guillet, L., “On Copper-Zinc-Nickel Alloys” (in French), Compt. Rend. Acad. Sci. Paris, 155(26), 1512-1514, (1912) (Experimental, Morphology, 1)
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Zn [1913Gui] [1914Tho] [1916Tho] [1920Gui1] [1920Gui2] [1922Voi] [1924Pri]
[1925Gui]
[1925Ost]
[1926Sma] [1929Bau1]
[1929Bau2]
[1930Bau]
[1934Yam] [1935Kih]
[1935Sch]
[1936Sch]
[1940Glu]
[1949Jae] [1952Haw]
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Guillet, L., “Nickel Brasses” (in French), Mem. Sci. Rev. Metall., 10, 1130-1141 (1913) (Experimental, 3) Thompson, F.Ch., “The Metallography of German Silver (Cu-Ni-Zn Alloys)”, J. Chem. Soc., 105, 2342-2349, (1914) (Experimental, Morphology, 8) Thompson, F.Ch., “The Annealing of Nickel Silver (Cu-Ni-Zn Alloys), Part I”, J. Inst. Met., 15, 230-263 (1916) (Experimental, Phase Relations, 10) Guillet, L., “On Copper-Zinc-Nickel Alloys” (in French), Compt. Rend. Acad. Sci. Paris, 170, 460-462, (1920) (Experimental, 2) Guillet, L., “New Investigations on Nickel Brasses” (in French), Mem. Sci. Rev. Metall., 17, 484-493 (1920) (Experimental, 2) Voigt, W., “Investigations on New Silvers (Cu-Ni-Zn Alloys)” (in German), Z. Anorg. Allg. Chem., 120, 309-319 (1922) (Experimental, Phase Relations, 6) Price, W.B., Grant, C.G., “Some Low Copper-Nickel Silvers (Cu-Ni-Zn Alloys)”, Trans. AIME, 70, 328-341 (1924) (Experimental, Morphology, Phase Relations, Phase Diagram, #, 2) Guillet, L., “New Researches on Various Nickel Brasses and their Thermal Treatments” (in French), Mem. Sci. Rev. Metall., 22, 383-394 (1925) (Experimental, Morphology, Phase Relations, 4) Ostroga, F.M., “Annealing, Quenching and Tempering of Some Industrial Nickel Brasses” (in French), Mem. Sci. Rev. Metall., 22, 776-786 (1925) (Experimental, Morphology, Phase Relations, 3) Smalley, O., “Special Nickel Brasses”, Trans. AIME, 73, 799-833 (1926) (Experimental, Phys. Prop., 0) Bauer, O., Hansen, M., “Influence of a Third Metal on the Constitution of Brasses. IIInfluence of Nickel” (in German), Z. Metallkd., 21(11), 357-367 (1929) (Experimental, Phase Diagram, Phase Relations, #, 31) Bauer, O., Hansen, M., “Influence of a Third Metal on the Constitution of Brasses. IIInfluence of Nickel” (in German), Z. Metallkd., 21(12), 406-411 (1929) (Experimental, Morphology, Phase Diagram, Phase Relations, 1) Bauer, O., Hansen, M., “Influence of a Third Metal on the Constitution of Brasses. IIInfluence of Nickel” (in German), Mitt. Mater., 23(7), 7-25 (1930) (Experimental, Phase Diagram, Morphology, #, 43) Yamaguchi, K., Nakamura, K., (in Japanese) Rikwagaku Kenkyu-jo Iho 13, 89-117 (1934) (Experimental, Morphology, Phase Diagram, #, 10) Kihlgren, T.E., Pilling, N.B., Wise, E.M., “Physical and Casting Properties of the Nickel Silvers (Cu-Ni-Zn Alloys)”, Trans AIME, 117, 279-312 (1935) (Experimental, Phase Diagram, Mechan. Prop., Morphology, 18) Schramm, J., “Copper-Nickel-Zinc Alloys” (in German), Dissertation Thesis, Konrad Triltsch Ed., Wuerzburg, Germany, 128 pp. (1935) (Experimental, Phase Diagram, Phase Relations, Morphology, #, *, 37) Schramm, J., Vaupel, O., “Roentgenographic Investigation on the Ni-Cu-Zn Ternary System” (in German), Metallwirtschaft, 15, 723-726, (1936) (Phase Diagram, Experimental, #, 22) Gluskin, D.Ya., “Nature of Phases Obtained by Reaction between Refractory Metals and Alloys with Low Melting Point Metals” (in Russian), Zh. Tekh. Fiz., 10, (18), 1486-1501, (1940) (Experimental, 15) Jaenecke, E., “Cu-Ni-Zn” (in German), Kurzfasstes Handbuch aller Legierungen, Carl Winter Universitaetverlag, Heidelberg, 504-508 (1949) (Phase Diagram, Review, #) Haworth, J.B., Hume-Rothery, W., “The Effect of Four Transition Metals on the / Brass Type of Equilibrium”, Philos. Mag., (7), 43(341), 613-629 (1952) (Phase Diagram, Experimental, 23)
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344 [1956Fre]
[1957Koe]
[1960Cha]
[1960Koe]
[1961Cha]
[1961Phi] [1963Koe]
[1965Hir]
[1966Phi]
[1967Liv]
[1968Bar]
[1969Bar] [1969Sch] [1970Tay] [1971Coa] [1972Tho]
[1973Hor]
[1973Lev]
[1974Mar]
MSIT®
Cu–Ni–Zn French, A.R., “The Metallurgigal Control of Quality in the Production of Copper-Base Alloy Castings”, J. Inst. Met., 85, 293-317 (1956) (Electr. Prop., Mechan. Prop., Phase Diagram, Phys. Prop., Review, 124) Koester, W., Schuele, W., “Conductivity and Hall Constant. V: Copper-Nickel-Zinc-Alloys” (in German), Z. Metallkd., 48(11), 595-600 (1957) (Electr. Prop., Experimental, 10) Chadwick, G.A., Argent, B.B., “The Lattice Parameters of the Solid Solutions of Copper-Zinc-Nickel and Copper-Zinc-Manganese”, J. Inst. Met., 88, 318-319 (1960) (Crys. Structure, Experimental, 8) Koester, W., “Conductivity and Hall Constant. XII. Proof of the Ordering the Copper-Nickel-Zinc Mixed Crystalls” (in German), Z. Metallkd., 51(12), 716-721 (1960) (Electr. Prop., Experimental, Mechan. Prop., 14) Chadwick, G.A., Argent, B.B., “Thermodynamic Properties of Solid Solutions, Part 4: Copper-rich Solid Solutions of Copper + Zinc + Nickel”, Trans. Faraday Soc., 57, 2138-2142 (1961) (Thermodyn., Experimental, 21) Phillips, V.A., “Anomalous Electrical Resistivity of Ordered Cu2NiZn at Low Temperatures”, Acta Metall., 9, 976-978 (1961) (Electr. Prop., Experimental, 7) Köster, W., Stoering, R., “Conductivity and Hall Constant. XXV. Overview About the Behavior of the Ternary Copper-Nickel-Zinc Mixed Crystals” (in German), Z. Metallkd., 54(3), 182-186 (1963) (Electr. Prop., Experimental, Phase Relations, 12) Hirabayashi, M., Hoshino, S., Sato, K., “Neutron Diffraction Study of the Long Range Order in a Single Crystal Cu2NiZn”, J. Phys. Soc. Jpn, 20(3), 381-388, (1965) (Crys. Structure, Experimental, 20) Phillips, V.A., Roberts, B.W., “Neutron-Diffraction Evidence Suggesting Clustering in Commercial “Nickel Silver” Close to the Cu2NiZn Composition”, Trans. Metall. Soc. AIME, 236, 1012-1014 (1966) (Crys. Structure, Mechan. Prop., Experimental, 26) Livshits, B.G., Tchupyatova, L.P., Lileev, A.S., “Investigation on Ordering in Alloys of Ternary Systems” (in Russian), Izv. Vyss. Uchebn. Zavedn., Chern. Metall., (1), 123-127, (1967) (Experimental, 6) Bartsch, G., “Yield Point and Plastic Deformation Behaviur of Copper-Nickel-Zinc Alloys with Different Degree of Short Range and Long Range Order” (in German), Z. Metallkd., 59(9), 729-735 (1968) (Experimental, Mechan. Prop., 12) Bartsch, G., “Neutron Diffraction Study of Cu-Zn-Ni Alloys” (in German), Z. Metallkd, 60(2), 139-142, (1969) (Crys. Structure, Experimental, 13) Schuele, W., Colella, R., “Increase in Electrical Resistivity due to Long-Range Order in Cu3Ni3Zn2”, J. Inst. Met., 97, 270-273 (1969) (Experimental, Electr. Prop., 9) Taylor, C.W., Dayananda, M.A., Grace, R.E., “Multiphase Diffusion in Cu-Zn-Ni Alloys”, Met. Trans., 1(1), 127-131 (1970) (Experimental, #, 5) Coates, D.E., Kirkaldy, J.S., “Morphological Stability of - Phases Interfaces in the Cu-Ni-Zn System at 775°C”, Met. Trans., 2(12), 3467-3477, (1971) (Experimental, 31) Thomas, H., “On the Electrical Resistivity of Copper-Nickel-Zinc Alloys and the Influence of Low-Temperature Deformation” (in German), Z. Metallkd., 63(2), 106-109 (1972) (Electr. Prop., Experimental, 21) Horch, R., Bruemmer, O., “Study of the K-State by Determining Thermal Properties of New Silver (Cu-Ni-Zn Alloys)” (in German), Krist. Tech., 8(6), 717-728 (1973) (Experimental, Electr. Prop., 16) Levine, E.D., “Cu-Ni-Zn (Copper-Nickel-Zinc)”, Metals Handbook, 8th Ed., Metallography, Structures and Phase Diagrams, ASM Int., Metals Park, Ohio, 427-428, (1973) (Phase Diagram, Review, #, 2) Marshakov, I.K., Karavaeva, A.P., Vavresyuk, I.V., “Anodic Solution of Different Faces of a Crystal of the Solid Solution in the System Cu-Zn-Ni”, Prot. Met., 10(4), 367-369, (1974), translated from Zashchita Metallov, 10(4), 399-401, (1974) (Experimental, 7) Landolt-Börnstein New Series IV/11C3
Cu–Ni–Zn [1974Sht]
[1976Ban] [1976Shi]
[1976Vri]
[1977Hal]
[1977Sch]
[1977Sis]
[1977Wir] [1978Liv]
[1978Sch]
[1979Cha]
[1979Day] [1979Dri] [1979Hos] [1979May] [1979Van]
[1980Ble] [1980Ito] [1980Roo] [1981Ara]
Landolt-Börnstein New Series IV/11C3
345
Shtrakhmann, K.M., Piguzov, Yu.V., Logvinienko, Yu.S., “Partial Relaxation Process Parameters in Ternary Monophased Cu-Zn-Ni Alloys” (in Russian), Relaxation of Metals in the Solid-State, Akad. Nauk. Litovsk. SSR, Kaunas, 158-162 (1974) (Experimental, 9) Banerjee, K.G., Mills, B., “The Zener Relaxation and Diffusion in the Copper-Nickel-Zinc System”, Phys. Status Solidi A, 33A, 707-714, (1976) (Experimental, 20) Shimuzu, S., Murakami, Y., Kachi, S., “Lattice Softening and Martensitic Transformation in Cu-Ni-Zn Phase Alloys”, J. Phys. Soc. Jpn., 41(1), 79-84 (1976) (Crys. Structure, Experimental, 24) Vrijen, J., Bronsveld, P.M. van der Veen, J., Radelaar, S., “Long Range Order in Cu2NiZn, Studied by Means of Thermal Neutron Scattering and Electron Microscopy”, Z. Metallkd, 67(7), 473-478, (1976) (Crys. Structure, Experimental, 15) Halder, S.K., De, M., Sen Gupta, S.P., “An X-Ray Diffraction Study on the Microstructures of Cold-Worked fcc Cu-Ni-Zn Alloys”, J. Appl. Phys., 48(8), 3560-3565 (1977) (Crys. Structure, Experimental, 23) Schlaepfer, H.-W., Form, W., “Transmission Electron Microscopy Analysis of a Spinel in Copper-Nickel-Zinc Alloys” (in German), Z. Metallkd. 68(1), 62-68 (1977) (Crys. Structure, Experimental, 20) Sisson, R.D., Dayananda, M.A., “Diffusional and Thermodynamic Interactions in the Cu-Zn-Ni System at 775°C”, Met. Trans. A, 8A(12), 1849-1856 (1977) (Thermodyn., Experimental, #, *, 41) Wirtz, L.E., Dayananda, M.A., “Diffusion Paths and Structures in Multiphase Cu-Ni-Zn Couples at 775°C”, Met. Trans. A, 8A(4), 567-575 (1977) (Experimental, #, 10) Livesey, D.W., Ridley, N., “Superplastic Deformation, Cavitation and Fracture of Microduplex Cu-Ni-Zn Alloys”, Met. Trans. A, 9A(4), 519-526 (1978) (Experimental, Mechan. Prop., 12) Schlaepfer, H.-W., Form, W., “Fire-Cracking and Ordering in Copper-Nickel-Zinc Alloys with Lead” (in German), Z. Metallkd., 69(3), 143-148 (1978) (Crys. Structure, Experimental, 26) Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Ni-Zn”, in “INCRA Monograph, Serie 6: Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metal Systems”, NSRD, Washington, 6, 620-630, (1979) (Phase Diagram, Review, #, *, 15) Dayananda, M.A., Kim, C.W., “Zero-Flux Planes and Flux Reversals in Cu-Ni-Zn Diffusion Couples”, Met. Trans A, 10A(9), 1333-1339 (1979) (Experimental, #, 38) Dritz, M.E., “Copper, Nickel, Zinc” (in Russian), in “Binary and Multicomponent Copper-Based Systems”, Nauka, Moscow, 185-186 (1979) (Review, #, 1) De Hosson, J.Th.M., “Atomistic Approach to Ordering in Cu 2NiZn”, AIP Conf., Proc., 53, 146-148 (1979) (Theory, 10) Mayall, O.S., Mathew, A., “The Lines in a Two-Phase Region in the Cu-Ni-Zn System”, J. Appl. Cryst., 12, 360-364, (1979) (Phase Diagram, Experimental, #, 4) Van der Veen, J., Jansen, W., de Groot, J., van Royen, E., Bronsveld, P., de Hosson, J., “Zener-Relaxation Effect in -CuNiZn Alloys”, Z. Metallkd., 70(7), 454-458 (1979) (Experimental, Mechan. Prop., Phase Relations, 18) “Copper-Nickel-Zinc Alloys (German Silver)” (in German), Blech Rohre Profile, 27(12), 873-874 (1982) (Phys. Prop., #) Ito, T., Nakayama, Y., “The Effect of Annealing on the Mechanical Properties of Cu-Ni-Zn Alloys”, Trans. Jpn. Inst. Met., 21(11), 745-752 (1980) (Experimental, Mechan. Prop., 17) de Rooy, A., van Royen, E.W., Bronsveld, P.M., de Hosson, J.Th.M., “The Coherent Phase Diagram of Cu-Ni-Zn”, Acta Met., 28, 1339-1347 (1980) (Phase Diagram, Theory, #, 28) Arato, P., Kedves, F.J., Kajdi, K., Gergely, L., “The Effect of Cold Rolling on Decomposition of K State and on Lattice Distorsion in Cu-Ni-Zn Alloys”, Prace Inst. Metali Niezel., 10(4), 165-167 (1981) (Experimental, Mechan. Prop., 12) MSIT®
346 [1981Roo]
[1981Veg]
[1981Weg]
[1982Roo]
[1982Ste]
[1983May] [1983Weg]
[1984Hos]
[1984Kim]
[1985Has]
[1985Kan]
[1986Sug]
[1986Zei] [1987Has]
[1987Puc]
[1987Ryc1] [1987Ryc2]
MSIT®
Cu–Ni–Zn de Rooy, A., van der Wegen, G.J.L., Bronsveld, P.M., de Hosson, J.Th.M., “The Quasi-Binary Cross Section in the Ternary System Cu-Ni-Zn. Part II: Electrical Resistivity Measurements”, Scr. Metall., 15(12), 1362-1364 (1981) (Experimental, Electr. Prop., #, 6) van der Vegt, W.H.M., van der Wegen, G.J.L., Bronsveld, P.M., de Hosson, J.Th.M., “The Lattice Parameter of Cu2NiZn”, Acta Crystallogr., Sect. A, 34A, C-101 (1981) (Crys. Structure, 2) van der Wegen, G.J.L., de Rooy, A., Bronsveld, P.M., de Hosson, J.Th.M., “The Order-Disorder Transition in the Quasi-Binary Cross Section Cu50Ni50–xZnx. Part I: Transmission Electron Microscopic Observations”, Scr. Metall., 15(12), 1359-1361 (1981) (Experimental, Crys. Structure, #, 8) de Rooy, A., Bronsveld, P.M., de Hosson, J.Th.M., “The Ternary Phase Diagram of Cu-Ni-Zn Investigated with Electrical Resistyivity Measurements”, Z. Metallkd., 73(10), 610-615 (1982) (Phase Diagram, Electr. Prop., Experimental, #, *, 14) Steeb, J., Stueer, H., Duerrschnabel, W., “New Spring Copper-based Spring Material with Higher Strength and Higher Electrical Conductivity” (in German), Metallwissenschaft und Technik, 36(11), 1185-1188 (1982) (Electr. Prop., Mechan. Prop., Experimental, 8) Mayall, O.S., “The fcc + Tetragonal Two-Phase Region of the Cu-Ni-Zn System”, J. Appl. Cryst., 16, 99-102 (1983) (Crys. Structure, #, 4) van der Wegen, G.J.L., Helmholdt, R., Bronsveld, P.M., de Hosson, J.Th.M., “Single Crystal Neutron Diffraction Study of the Long Range Order in Cu2NiZn”, Z. Metallkd., 74(9), 592-597 (1983) (Crys. Structure, Experimental, 13) de Hosson, J.Th.M., “Superlattice Dislocations in Ternary Alloys: a Transmission Electron Microscopic Study on the Microstructure of Cu2NiZn”, Res. Mech., 11, 97-138 (1984) (Crys. Structure, Electronic Structure, Experimental, Mechan. Prop., 75) Kim, C.W., Dayanada, M.A., “Zero-Flux Planes and Flux Reversals in the Cu-Ni-Zn System at 775°C”, Metall. Trans. A, 15A(4), 649-659 (1984) (Calculation, Experimental, Phase Relations, Transport Phenomena, 40) Hashimoto, S., Iwazaki, H., Ohshima, K., Harada, J., Sakata, M., Terauchi, H., “Study of Local Atomic Order in a Ternary Cu0.47Ni0.29Zn0.24 Alloy Using Anomalous Scattering of Synchrotron Radiation”, J. Phys. Soc. Jpn., 54(10), 3796-3807 (1985) (Crys. Structure, Experimental, 21) Kansky, K.S., Dayananda, M.A., “Quaternary Diffusion in the Cu-Ni-Zn-Mn System at 775°C”, Met. Trans A, 16A(6), 1123-1132 (1985) (Experimental, Transport Phenomena, 19) Sugino, S., Hagiwaraa, H., “Effects of Aluminum and Nickel on the Activity of Zinc in Molten Copper” (in Japanese), J. Jpn. Inst. Met., 50(12), 1068-1074 (1986) (Thermodyn., Experimental, 19) Zeiger, H., “State of Art and Development of Copper and Copper Alloys” (in German), Z. Werkstofftech., 17, 75-78 (1986) (Review, 32) Hashimoto, S., “Intensity Expression for Short-Range-Order Diffuse Scattering as a Function of Ordering Energies in a Ternary Alloy System”, J. Appl. Crystallogr., 20(3), 182-186 (1987) (Calculation, Transport Phenomena, 10) Puckert, F., “Recently Developed Copper-based Material for High Current Carrying Connectors and Frames for Semiconductor” (in German), Metall, 41(11), 1116-1119 (1987) (Review, 9) Rychlewski, M., “Iron, Cobalt and Nickel Interactions with Zinc in Dilute Solution with Molten Copper”, Z. Metallkd., 78(3), 214-217 (1987) (Thermodyn., Experimental, 10) Rychlewski, M., Pomianek, T., “On the Influence of Tin, Germanium and the Iron Group Metals on the Zinc Activity in Liquid Copper Alloys” (in Polish), Metalurgia I Odlewnictwa, 109(1138), 55-61 (1987) (Thermodyn., Experimental, 19)
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Zn [1988Gur]
[1988Ito] [1988Pry] [1988Puc] [1990Seg] [1991Mar]
[1991Nas] [1993Ham] [1993Kow] [1994Cha]
[1994Fon]
[1994Mio]
[1996Alt]
[1997Alt]
[1997Von] [1998Sim]
[2001Her] [2002Leb]
[2003Mie] [2004Li]
Landolt-Börnstein New Series IV/11C3
347
Gurevich, R.L., Shvartsman, A.B., “Ordering and Breakdown in Cu-Zn-Ni Alloys”, Russ. Metall., (2), 95-100 (1988), translated from Izv. Akad. Nauk SSSR, Metally, (2), 99-104 (1988) (Electr. Prop., Experimental, 9) Ito, T., Nakayama, Y., “Dislocation Structures in Deformed Cu-Ni-Zn Alloy Single Crystals”, J. Mater. Sci., 23, 2174-2180 (1988) (Crys. Structure, Experimental, 23) Prym-Werke, W., “Properties and Applications of the Most Useful Copper Alloys” (in German), Metall, 42(8), 821-822 (1988) (Review, Experimental, 0) Puckert, F., Duerrschnabel, W., “Interactions Between Solderability and Bendability of Copper Alloys” (in German), Metall, 42(3), 254-258 (1988) (Experimental, 3) Segal, A., “Copper in Communication”, Met. Mater., 428-430 (1990) (Electr. Prop., Optical Prop., Review, 0) Martinova, Z., Antonova, N., Zaidel, T., Dimitrov, L., “Influence of Phase Composition and Impurities on Cracking of Nickel Brass During Hot Rolling” (in Russian), Metallurgia (Sofia), 46(1), 12-16 (1991) (Experimental, Mechan. Prop., 7) Nash, P., Pan, Y.Y., “Ni-Zn (Nickel-Zinc)”, in “Phase Diagrams of Binary Nickel Alloys”, ASM International, Materials Park, OH, 382-390 (1991) (Review, #, 54) Hammond, L., Wright, D., “XRD Pattern Fitting as a Tool to Control Plating Parameters in Zinc-Nickel Electroplates”, Adv. X-ray Anal., 36, 315-325 (1993) (Experimental, 11) Kowalski, M., Spencer, P.J., “Thermodynamic Reevaluation of the Cu-Zn System”, J. Phase Equilib., 14(4), 432-438 (1993) (Thermodyn., #, 36) Chakrabarti, D.J., Laughlin, D.E., Chen, S.W., Chang, Y.A., “Cu-Ni (Copper - Nickel)” in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 276-286 (1994) (Review, #, 85) de Fontaine, D., “Cluster Approach to Order-Disorder Transformations in Alloys”, Solid State Phys., 47, 33-176 (1994) (Crys. Structure, Phase Diagram, Review, Theory, Thermodyn., 213) Miodownik, A.P., “Cu-Zn (Copper - Zinc)” in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 487-496 (1994) (Review, 98) Althoff, J.D., Johnson, D.D., Pinski, F.J., Staunton, J.B., “Electronic Origins of Ordering in Multicomponent Metallic Alloys: Application to the Cu-Ni-Zn System”, Phys. Rev. B: Condens. Matter, 53, 10610-10625 (1996) (Electronic Structure, Theory, 44) Althoff, J.D., Johnson, D.D., “The Electronic Origins of Atomic Short-Range Order in Disordered fcc Cu-Ni-Zn Ternary Alloys”, J. Phase Equilib., 18(6), 567-572 (1997) (Electronic Structure, Review, 19) Voncken, J.H.L., Verkroost, Th.W., “Powder Diffraction of Cubic -Brass”, Powder Diffr., 12(4), 228-229, (1997) (Experimental, Crys. Structure, 7) Simak, S.I., Ruban, A.V., Abrikosov, I.A., Skriver, H.L., Johansson, B., “Ordered Phases in Cu2NiZn: A First-Principles Monte-Carlo Study”, Phys. Rev. Lett., 81(1), 188-191, (1998) (Crys. Structure, Theory, 20) Hertz, J., “What Will be Done in the Future with O.J. Kleppa’s Enthalpy Data Set?”, J. Alloys Compd., 321, 201-222 (2001) (Experimental, Thermodyn., #, 50) Lebrun, N., “Cu-Ni (Copper - Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.14832.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, 51) Miettinen, J., “Thermodynamic Description of the Cu-Ni-Zn System Above 600°C”, Calphad, 27(2), 263-274 (2003) (Assessment, Thermodyn., #, 41) Li, C., Chin, Y.L., Wu, P., “Correlation Between Bulk Modulus of Ternary Intermetallic Compounds and Atomic Properties of their Constituent Elements”, Intermetallics, 12, 103-109 (2004) (Electronic Structure, Thermodyn., 24) MSIT®
Cu–Ni–Zn
348 [2005Jia]
[2006Leb]
Jiang, M., Wang, C.P., Liu, X.J., Ohnuma, I., Kainuma, R., Vassilev, G.P., Ishida, K., “Thermodynamic Calculation of Phase Equilibria in the Cu-Ni-Zn System”, J. Phys. Chem. Solids, 66(2/4), 246-250 (2005) (Assessment, Phase Relations, Phase Diagram, Thermodyn., #, 21) Lebrun, N., Perrot, P., “Cu - Zn (Copper - Zinc)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), Materials Science International Services, GmbH, Stuttgart; to be published (2006) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 18)
Table 1: Investigations of the Cu-Ni-Zn Phase Relations, Structures and Thermodynamics Reference
Experimental Technique
Temperature/ Composition/ Phase Range Studied
[1908Taf]
Thermal analysis
The whole diagram, 0-1200°C
[1934Yam]
Thermal analysis
The whole diagram, 400-1300°C
[1935Sch]
Thermal analysis
The whole diagram, 0-1450°C
[1936Sch]
X-Ray investigation
The whole diagram, 298 K
[1952Haw]
- equilibrium, annealing and microscopic examination
4-12 at.% Ni, 36-48 at.% Zn, 672°C
[1960Cha]
Lattice parameters measurements
solid solution (< 24 at.% Zn, < 25 at.% Ni), 298 K
[1960Koe]
Ordering determination by electrical Ni/Zn = 1, 35-70 at.% Cu, 400-700°C resistance, Hall constant and X-Ray
[1961Cha]
Zinc activities determined from zinc solid solutions (< 24 at.% Zn, < 25 at.% Ni), vapor pressure measurements 1000 K
[1965Hir]
Long range order by neutron diffraction
NiCu2Zn annealed at 300°C
[1966Phi]
Long range order by neutron diffraction
NiCu2Zn quenched from 600°C and annealed at 400°C
[1969Bar]
Long range order by neutron diffraction
Ni/Zn = 1, 40-60 at.% Cu, 200-600°C
[1969Sch]
Long range order by electrical and X-Ray measurements
Ni3Cu3Zn2 annealed at 600°C
[1979May]
Tie-lines determination by X-Ray measurements
(+) region: 30-50 at.% Zn, < 35 at.% Ni, 600°C
[1981Roo]
Order-disorder transition by electrical resistance measurements
50 at.% Cu, 15-35 at.% Zn, 250-500°C
[1981Veg]
X-Ray, parameter measurements, order-disorder transition
NiCu2Zn, 200-800°C
[1981Weg]
Order-disorder transition by electron 50 at.% Cu, 20-30 at.% Zn, microscopy observations 350-500°C
[1982Roo]
Order-disorder transition by electrical resistance measurements
MSIT®
domain, > 30 at.% Cu, 10-40 at.% Zn, 125-500°C
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Zn
349
Reference
Experimental Technique
Temperature/ Composition/ Phase Range Studied
[1983May]
Tie-lines determination by X-Ray measurements
( + 1) region: 25-50 mass% Zn, < 55 mass% Cu, 600°C
[1983Weg]
Long range order by single crystal neutron diffraction
NiCu2Zn annealed between 300 and 500°C
[1985Has]
Atomic order by anomalous scattering of synchrotron radiation
Cu0.47Ni0.29Zn0.24 annealed at 600°C
[1986Sug]
Zinc activity measurement in liquid Liquid Cu (< 9 at.% Zn, < 8 at.% Ni) alloy by an isopiestic method 1100-1150°C
[1987Ryc1, 1987Ryc2]
Zinc activity measurement in liquid Liquid Cu (< 12 at.% Zn, < 10 at.% Ni), alloy by an isopiestic method 1200°C
[2003Mie]
Calphad assessment
< 70 mass% Zn, 600-1400°C
[2005Jia]
Calphad assessment
The whole diagram, 600-1400°C
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, (Ni,Cu,Zn)
cF4 Fm3m Cu
Cu0.5Ni0.25Zn0.25 1100 - 501 (Cu) < 1084.62
Cu0.5Ni0.5 (Ni) < 1455 , (Zn) < 419.58
hP2 P63/mmc Mg
, (Ni,Cu)Zn
cP2 Pm3m CsCl
NiZn 1040 - 675 CuZn < 468
Landolt-Börnstein New Series IV/11C3
Lattice Parameters Comments/References [pm]
a = 363.9
[1981Veg]
a = 361.46
a = 369.61
pure Cu at 25°C [1994Cha] 38.27 at.% Zn in solid (Cu) at 454°C [1997Von] 35.84 at.% Zn in solid (Cu) at 300°C
a = 356.6
[1994Cha]
a = 352.410
pure Ni at 25°C [Mas2] 48.3 at.% Zn in solid Ni at 1040°C [Mas2]
a = 266.50 c = 494.70
pure Zn at 25°C Dissolves 2.83 at.% Cu at 425°C [Mas2] Dissolves 0.2 at.% Ni at 418.5°C [Mas2]
a = 290.83
a = 295.9
High temperature phase of Ni-Zn Lattice parameter from [V-C2] at 890°C Low temperature phase of Cu-Zn Lattice parameter at 49.5 at.% Zn [V-C2] MSIT®
Cu–Ni–Zn
350 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
1, NiZn < 810
tP2 P4/mmm AuCu
a = 274.1
', CuZn 902 - 454
cI2 Im3m W
a = 295.39
cI52 I43m Cu5Zn8
a = 893.57
57.7 to 70.6 at.% Zn in the Cu-Zn binary Lattice parameter from [V-C2]
a = 893.57
70 to 85 at.% Zn in the Ni-Zn binary Lattice parameter for Ni4Zn22 [1993Ham] at 500°C
Cu5Zn8 < 835 Ni4Zn22 < 881
, CuZn3 700 - 560
hP3 P6 CuZn3
a = 427.5 c = 259.0
J, CuZn4 < 598
hP2 P63/mmc Mg
a = 274.18 c = 429.39
, (Ni-Zn) < 490
mC6 C2/m CoZn13
45% Zn [1983May] Ordered form of Ni-Zn (L10 superstructure) High temperature form of -CuZn (labelled in binary diagrams) Lattice parameter at 47.5 at.% Zn [V-C2]
High temperature phase Lattice parameters from [V-C2] at 500°C 78 to 88 at.% Zn [Mas2] Lattice parameters from [V-C2] ~89 at.% Zn (labelled in the Ni-Zn binary) Lattice parameters for Ni3Zn22 [1993Ham] at 500°C
a = 1337.6 b = 751.1 c = 762.7 = 113.256
* NiCu2Zn(h) 501 - 325
cP4 Pm3m AuCu3
a = 363.7
at 400°C [1981Veg] Ordered form (L12 superstructure) of the phase [1979Hos, 1981Weg]
* NiCu2Zn(r) < 325
tP2 P4/mmm AuCu
a = 363.5 c 363.5
at 300°C [1981Veg] Ordered form (L10 superstructure) of the phase [1979Hos, 1981Roo]
Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Ni
Cu
Zn
L + + '
935
P
L
22.67
26.93
50.4
L + ' +
860
U
L
24.88
11.99
63.13
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Zn
351
Zn e2
Fig. 1: Cu-Ni-Zn. Liquidus surface projection
p2
Data / Grid: mass%
0 70 0 80
p4
85 0
20
p80 5
γ
U
e1
Axes: mass%
p3 600
β
40
p6
60
900
p1
110 0
10 00
β'
P
α
60
40
p7
1200
80
130 0
20
1400 °C 20
Ni
40
60
80
Zn
Cu
Data / Grid: at.% Axes: at.%
Fig. 2: Cu-Ni-Zn. Isothermal section at 850°C 20
80
40
L
γ+β β
60
L+β
α +β
60
40
α 80
Ni
Landolt-Börnstein New Series IV/11C3
20
20
40
60
80
Cu
MSIT®
Cu–Ni–Zn
352
Zn
Data / Grid: at.% Axes: at.%
Fig. 3: Cu-Ni-Zn. Isothermal section at 775°C showing isoactivity lines in solid solution from [1977Sis]
L L+γ
20
β +β 1 β1
60
γ +β ' β
β 1+α
β'
β +β '
α +β
60
0.5
γ
γ +β
40
Cu Ni Zn
80
0.4
α +β ' 0.1
0.2
40
0.3
0.1 0.08 0.05 0.03
0.6 80
20
0.7 0.8
0.02
0.9
Ni
0.01 0.005
0.2 0.3
0.4
20 0.5
400.6
0.7 60
0.8
Zn
0.002
0.9
80
Cu
Data / Grid: at.% Axes: at.%
Fig. 4: Cu-Ni-Zn. Isothermal section at 650°C
L L+γ
20
80
δ
γ 40
γ+β+β 1
60
γ+β β
β1
α+β
60
40
α+β+β1 α
80
Ni
MSIT®
20
40
20
60
80
Cu
Landolt-Börnstein New Series IV/11C3
Cu–Ni–Zn
353
Zn
Data / Grid: at.% Axes: at.%
Fig. 5: Cu-Ni-Zn. Calculated partial isothermal section at 427°C 20
80
40
60
60
40
α+L12
L12
80
20
α
20
Ni
40
60
80
Zn
Cu
Data / Grid: at.% Axes: at.%
Fig. 6: Cu-Ni-Zn. Calculated partial isothermal section at 327°C 20
80
40
60
60
40
α+L12
80
L10
L12 20
α
Ni
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Cu
MSIT®
Cu–Ni–Zn
354
Fig. 7: Cu-Ni-Zn. Vertical section Cu0.5Ni0.5–xZnx (0.15 < x < 0.35)
500
α
Temperature, °C
α +L12 L12 (Cu2NiZn(h))
400
α +L12
300
α +L10+L12 α +L10
L10 (Cu2NiZn(r)) L10+L12
200
Ni 35.00 Cu 50.00 Zn 15.00
MSIT®
20
30
Zn, at.%
Ni 15.00 Cu 50.00 Zn 35.00
Landolt-Börnstein New Series IV/11C3
Cu–P–Sn
355
Copper – Phosphorus – Tin Peter Rogl Introduction High mechanical strength, wear resistance and sea-water corrosion resistance of Cu-Sn based bronces have triggered early scientific interest in the constitution of the Cu-P-Sn system. Although the entire diagram still awaits further investigation, several research teams collected experimental phase diagram information on the Cu rich corner of the phase diagram [1910Hud, 1911Lev, 1926Gla1, 1926Gla2, 1933Ver, 1937Ham, 1954Mur, 1987Tak1, 1987Tak2]. Whilst [1933Ver] discussed earlier results of [1910Hud, 1911Lev, 1926Gla1, 1926Gla2], the authors of [1987Tak1, 1987Tak2] seemed to be unaware of previous extensive work in the system. A summary of the constitution of the Cu-P-Sn system was given by [1979Dri], presenting a liquidus projection and partial isothermal sections at 600 and 300°C for the Cu rich region. A thermodynamic assessment for the Cu rich region was recently published [2001Mie]. Cu4SnP10 is the only ternary compound hitherto reported for the ternary system [1979Hoe, 1980Hoe], however, phase relations involving the compound are still unknown. The various experimental activities related to the constitution of the ternary Cu-P-Sn system are summarized in Table 1. Binary Systems The binary systems Cu-P and P-Sn are used from the MSIT Binary Evaluation Program [2002Per, 2002Wat]. The Cu-Sn system is adopted from [1994Sau]. Solid Phases Crystallographic and melting data pertinent to the Cu-P-Sn system are given in Table 2. Only one ternary compound, Cu4SnP10, for which a {P10}6– polyanion of decaphospha-adamantane structure and a {Cu3Sn} four-center 2-electron cluster is remarkable, has so far been reported for the ternary system [1979Hoe, 1980Hoe]. Invariant Equilibria Early investigators [1910Hud, 1926Gla1, 1926Gla2, 1937Ham] claimed the existance of a ternary eutectic at Cu80.7Sn14.8P4.5 mass% and 628°C, L (Cu) + + Cu3P. Experiments by [1911Lev, 1933Ver] indicated that the “ternary eutectic” should rather be a transition reaction, L + (Cu) + Cu3P, for which [1933Ver] gave composition and temperature: Cu79.6Sn16.2P4.2 mass% and 637°C. It has to be noted, that the composition of the liquid reported is close to the tie line Cu3P - but still within the triangle (Cu)+Cu3P+, thus rather indicating a ternary eutectic. Metallographic inspection and DTA-data from [1954Mur, 1987Tak1, 1987Tak2] were however, resolved in terms of a transition reaction at 655°C [1954Mur] or at 642°C [1987Tak1, 1987Tak2], respectively: L + (Cu) + Cu3P. The composition of the liquid was reported at significantly lower Cu- and P-contents, Cu77.5Sn20.7P1.6 mass% [1954Mur] on a practically straight line connecting the binary eutectic, L (Cu)+Cu3P, with the binary peritectic point, L + (Cu) . The composition of the ternary liquid L would fall on the line connecting Cu3P with (Cu17Sn3), assuming a minor change in composition of the phase (Cu76.2Sn23.8 mass%) from only a small phosphorus solubility. Essentially based on DTA and metallographic inspection (no details on X-ray data and EMPA were given), [1954Mur] reported a composition of Cu81.5Sn17.8P0.7 mass% (point B in figure 7 of [1954Mur]) for presumably the vertex of the four-phase reaction plane at the ternary phase, 655°C. Similarly, the vertex at the (Cu) solid solution was located at a surprisingly low Sn content (Cu97.8Sn2P1.2 mass%) with respect to a maximal Sn solubility in (Cu) at 650°C of 15.6 mass% Sn. A redetermination of the vertices of the isothermal reaction plane will thus be necessary in order to resolve the doubts in the compositions given as well as to resolve the discrepancies between the data of [1910Hud, 1926Gla1, 1926Gla2, 1933Ver, 1937Ham] and the results of [1954Mur, 1987Tak1, 1987Tak2]. Although Landolt-Börnstein New Series IV/11C3
MSIT®
356
Cu–P–Sn
somewhat doubtful but as the only experimental information available, the parameters for the ternary reaction taken at 642-655°C after [1954Mur] are listed in Table 3. It should be furthermore noted that the phase (Cu3Sn) was not observed [1954Mur, 1987Tak1, 1987Tak2] probably due to insufficiently fast quenching rates. This fact devaluates somewhat the vertical sections presented by [1954Mur] (at 1 mass% P, 6 and 10 mass% Sn) as well as those given by [1987Tak1, 1987Tak2] for 3, 5, 7 mass% P and 10, 15 mass% Sn). The phase was retained in the experiments by [1933Ver], however, the different position of the ternary reaction isotherm at 637°C makes a comparison difficult. All experiments agree on a eutectoid ternary decomposition of the phase (and phase ?): from vertical sections the authors of [1954Mur, 1987Tak1, 1987Tak2] concluded a eutectoid decomposition of the phase at 515°C (Cu75.5Sn23.7P0.8 mass%), ( Cu3P + (Cu) + ), whereas [1933Ver] seemed to indicate the decomposition: Cu3P + (Cu) + at 587°C. The data of [1954Mur, 1987Tak1, 1987Tak2] have to be taken with care as these authors were unable to quench and observe the phase (BiF3 type). The proper reaction thus will be Cu3P + (Cu) + as indicated by [1933Ver] at about 500°C. No details on the composition of the corresponding phases are available. A tentative Schultz-Scheil diagram summarizing the invariant reactions in the Cu rich part of the diagram is given in Fig. 1. Liquidus, Solidus and Solvus Surfaces A liquidus projection for the Cu rich region was presented by [1933Ver, 1954Mur, 1987Tak1, 1987Tak2] superseeding older data by [1926Gla1, 1926Gla2] on the melting trough running from the binary Cu-Cu3P eutectic to the binary Cu-Sn peritectic L+ (Cu) (see also discussion in section Invariant Equilibria). The liquidus projection is shown in Fig. 2 accepting an invariant ternary transition reaction. Isothermal Sections Four isothermal sections at 300, 550, 600 and 660°C were prepared by [1933Ver] for the Cu rich region up to 27 mass% Sn and 9 mass% P. All solid state phase equilibria are characterized by tie lines from Cu3P to the various binary Cu-Sn phases. For the 300°C section a tie line Cu3P - “Cu4Sn” ( phase, Cu41Sn11) was indicated [1933Ver], although more recent phase diagrams assume decomposition of Cu41Sn11 below about 350°C. The isothermal section for the Cu rich region at 550°C is presented in Fig. 3 showing also the location of the ternary phase -1, Cu4SnP10. The corresponding isothermal section at 600°C reveals similar phase relations, however, all binary tin-phosphides are already liquid (see Fig. 4). Temperature – Composition Sections [1933Ver] derived a set of isopleths with constant contents of 1, 2 and 3 mass% P (from 0 to 25 mass% Sn) as well as with constant content of 5, 10 and 20 mass% Sn (from 0 to 5 mass% P). Independently [1954Mur, 1987Tak1, 1987Tak2] determined isopleths with constant contents of 6, 10, 15 mass% Sn (from 0 to 8 mass% P) as well as with constant content of 1, 3, 5 and 7 mass% P (from 0 to 15 mass% Sn). Comparison of the two sets of data reveals liquidus temperatures of [1933Ver] to be generally lower by about 30°C and shifted to lower P contents of 4.2 mass% P for the monovariant liquid trough. In both data sets we encounter a series of inconsistencies with the solid state equilibria: (i) [1954Mur, 1987Tak1, 1987Tak2] have not observed the phase; (ii) their vertical section is inconsistent with the intersections of the L + solid phase lines on the four phase plane; (iii) although the phase was included in the sections of [1933Ver] and measurements were carried out on a finer sample grid, there are some misinterpretations of the TA data including several violations of the phase rule; (iv) whilst [1933Ver] shows invariant solid state reactions at 587°C (involving the phase) and 500°C (involving the phase), [1954Mur, 1987Tak1, 1987Tak2] only report an invariant temperature of 500-515°C neglecting the phase completely. The isopleth at 6 mass% Sn (up to 5 mass% P) and the isopleth at 10 mass% Sn (up to 5 mass% P) are shown in Fig. 5 and Fig. 6, respectively. A major point of discrepancy is the location of the four-phase plane near the Cu-Sn boundary for the transition reaction at 642-655°C. MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–P–Sn
357
Thermodynamics No experimental thermodynamic data are presently available for the ternary system. A thermodynamic CALPHAD type description of the Cu rich part of the system up to to 32 mass% Sn and 14 mass% P is due to [2001Mie]. The calculation includes a thermodynamic modelling of the Cu-P and P-Sn systems up to 30 and 50 at.% P, respectively. For the ternary system [2001Mie] based their calculation exclusively on experimental phase relations as described by [1987Tak1]. Serious discrepancies were encountered between calculated and experimental extents of phase regions. Notes on Materials Properties and Applications Technical use of Cu based friction bearings for automotive industry was summarized by [1992Bog]. A special band-casting technology for various Cu based alloys (including alloying constituents such as Sn, P) was patented by [1993Due]. The technical use of copper alloys in electric and electronic appliences and communication systems have been addressed by [1992Sch, 1990Seg, 1982Sak]. Development and properties of a high-conductivity/high strength cold worked copper alloy containing a submicron dispersion of intermetallic compounds in a Cu-matrix (Mg and P in the ratio of Mg3P2) revealing high resistance to heat softening has been reported and was compared to Cu95Sn5Pmin0.03 mass% and Cu92Sn8Pmin0.03 mass% alloys [1970Fis]. The dispersion pins dislocations and grain boundaries thus retaining high temperature strain-hardening. Effects of phosphorus on the friction and wear characteristics of a Cu95Sn5P1-5 at.% alloy were studied using a pin-disc apparatus [1979Yag, 1979Tag]. The coefficient of friction and the rate of wear decreased with increasing amount of matrix-Cu3P. The P surface content increases after heating to 300°C (P atoms at 400°C dissolved in -ss, whereas after cooling to room temperature the top surface layer was covered by a Cu3P film). [1956Fre] studied the influence of basic composition, microstructure and impurities on the properties of copper tin bronzes as well as the control of metal and casting quality. Examination of the influence of Sn and P contents as well as the effect of casting on the hot-working capacity (including hot forging, hot extrusion and hot rolling) of phosphorized tin bronzes showed the difficulties in handling these materials [1950Sho]. Miscellaneous The ability to form amorphous nanosize products (5-10 nm) was investigated via ballmilling 250 mesh elemental powder blends for the compositional range Cu86–xSnxP14 (x = 2-15) [1999Zha]. After 28 h of mechanical alloying Cu84Sn2P14 essentially amorphized with some amount of crystallized phase still present. Increased milling time increased the crystalline phase. The periods of complete amorphization of Cu86-xSnxP14 for x = 4, 5, 8, 10 were given as 28, 20, 12 and 32 h, respectively. No amorphous phase forms if the Sn content exceeds 15 mass% Sn. The glass forming range was identified within 4-14 mass% Sn. Explanation for the ability to form amorphous products in Cu-P-Sn alloys was attempted by the authors in the framework of the Zhang-Bangwei theory. The effect of Sn on the microstructure and solidus temperatures of Cu-P bronces was studied by [1989Kra]. References [1910Hud] [1911Lev] [1926Gla1]
[1926Gla2]
Landolt-Börnstein New Series IV/11C3
Hudson, O.F., Law, E.F., “A Contribution to the Study of Phosphor Bronze”, J. Inst. Met., 3, 161-186 (1910) (Phase Diagram, Review, 3) Levi- Malvano, M. L., Orofino, F. S., Gazz. Chim. Ital., 41(11), 269, 297 (1911) as quoted in [1933Ver] Glaser, L.C., Seemann, H.J., “On The Phosphor Bronze Based on Thermal Investigations in the Copper-Phosphor-Zinc System” (in German), Z. Techn. Physik, 7, 42-46 (1926) (Phase Diagram, Review, 11) Glaser, L.C., Seemann, H.J., “Description about Phosphor Bronze Based on Thermal Investigations in the Copper-Phosphor-Zinc System” (in German), Z. Techn. Physik, 7, 90-92 (1926) (Optical Prop., Phase Diagram, 7)
MSIT®
358 [1933Ver] [1937Ham] [1950Sho] [1954Mur] [1956Fre]
[1970Fis]
[1979Dri]
[1979Hoe]
[1979Tag]
[1979Yag] [1980Hoe]
[1982Sak] [1987Tak1]
[1987Tak2]
[1989Kra] [1990Seg] [1992Bog]
[1992Sch]
[1993Due]
MSIT®
Cu–P–Sn Veroe, J., “About the Phase Diagram of the Cu-Sn-P-Alloys” (in German), Z. Anorg. Chem., 213, 257-272 (1933) (Crys. Structure, Phase Diagram, 10) Hamazumi, M., J. Jpn. Inst. Met., 1, 165 (1937) (Experimental, Phase Diagram) as quoted in [1954Mur] Showell, D.W.D., “The Hot Working of Tin Bronzes”, J. Inst. Met., 76, 527-540 (1950) (Experimental, Phase Diagram, Review, Theory, 13) Muromachi, S., Watanabe, H., Tomimoto, S., “Metallographic Study of Phosphor Bronse”, Nippon Kinzoku Gakkai-Si, 18, 309 (1954) (Experimental, Phase Relations, 8) French, A.R., Member, F.I.M., “The Metallurgigal Control of Quality in the Production of Copper-Base Alloy Castings”, J. Inst. Met., 85, 293-317 (1956) (Electr. Prop., Mechan. Prop., Phase Diagram, Phys. Prop., Review, 124) Fisher, H.J., Hay, D.A., Finlay, W.L., “The Development and Properties of a New High-Conductivity/High-Strength Cu Alloy”, J. Inst. Met., 98, 368-375 (1970) (Electr. Prop., Experimental, Mechan. Prop., 4) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, Rokhlin, L.L., Turkina, N.I., “Cu-Sn-P” (in Russian), in “Binary and Multicomponent Copper-Base Systems”. Nauka, Moscow, 192-194 (1979) (Phase Diagram, Review, 3) Hoenle, W., von Schnering, H.G., “New Compounds Containing Ag6(4+) Cluster Units”, Solid Compounds of Transition Elements VI, Int. Conf., Stuttgart, 101-102 (1979) (Crys. Structure, 3) Taga, Y., Nakajima, K., “Effect of Phosphorus on the Friction and Wear Characteristic of Cu-Sn-P Alloys”, J. Lubric. Tech. - Trans. ASME, 101(2), 201-207 (1979) (Electronic Structure, Experimental, Mechan. Prop., 19) Yaga, Y., Najajima, K., “Effect of Phosphorus on the Friction and Wear Characteristics of Cu-Sn-P Alloys”, Mechan. Eng., 101(1), 95 (1979) (Abstract, Mechan. Prop., 0) Hoenle, W., von Schnering, H.G., “Cu4SnP10, a Compound with Decaphospha-Adamantane Anions P10 and Four-Center (SnCu3) Clusters”, Z. Kristallogr., 153, 339-350 (1980) (Crystal Structure, 27) Sakamoto, T., “Technical Trends in Metallic Materials for Electronic Applications” (in Japanese), Denki Seiko (Electr. Furn. Steel), 53(2), 137-142 (1982) (Review, 0) Takemoto, T., Okamoto, I., Matsumura, J., “Phase Diagrams of Copper-Silver-Phosphorus and Copper-Tin-Phosphorus Ternary Brazing Filler Metals - Copper Phosphorus Brazing Filler Metals with Low Melting Temperature (II)”, Trans. JWRI, 16(2), 301-307 (1987) (Experimental, Phase Diagram, 8) Takemoto, T., Okamoto, I., Matsumura, J., “Phase Diagrams of Copper-Silver-Phosphorus and Copper-Tin-Phosphorus Ternary Filler Metals Copper Phosphorus Brazing Filler Metals with Low Melting Temperatures (Report II)” (in Japanese), Yosetsu Gakkai Ronbunshu, 5(1), 81-86 (1987) (Phase Diagram, 13) Kraft, V.V., Uchitel, V.A., “Effect if Tin on the Solidus Temperature and Microstructure of Copper-Phosphorus Brazes”, Svar. Proizvod., (7), 21 (1989) Segal, A., “Copper in Communication”, Met. Mater., 428-430 (1990) (Electr. Prop., Optical Prop., 0) Bogel, A., Feind, J., “Material for Friction Bearings Made of Copper Alloys for Use in Automotive Industry” (in German), Metallwiss. Techn. Metall (Berlin), 46(11), 1132-1135 (1992) (Experimental, 11) Schleicher, K., Duerschnabel, W., Boegel, A., “Copper Alloy Materials as Base Materials for Components in Electrical Engineering and Electronics” (in German), Metallwissenschaft und Technik, Metall, 46(11), 1116-1118, 1120 (1992) (Experimental, Review, 4) Duerrschnabel, W., Mueller, H., Kehse, G., Boegel, A., Breu, M., “Continuous Band Casting of Precipitation- or Segregation-Prone or Stress-Sensitive Copper Alloys” (in
Landolt-Börnstein New Series IV/11C3
Cu–P–Sn
[1994Sau]
[1994Sub]
[1999Zha]
[2001Mie] [2002Per]
[2002Wat]
359
German), German Patent DE4126079A1, Wieland-Werke AG, Germany, 1-8 (1993) (Experimental, 0) Saunders, N., Miodownik, A.P., “Cu-Sn (Copper-Tin)”, in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 412-418 (1994) (Phase Diagram, Crys. Structure, Thermodyn., Review, # 57) Subramanian, P. R., Laughlin, D. E., “Cu-P (Copper-Phosphorus)”, in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 295-300 (1994) (Review, #, *, 24) Bangwei, Z., Haowen, X., Shuzhi, L., “Amorphous Forming Ability in the Ternary Cu-Sn-P System by Mechanical Alloying”, J. Mat. Proc. Tech., 89-90, 378-384 (1999) (Experimental, Thermodyn., 29) Miettinen, Y., “Thermodynamic Description of Cu-P-Sn System in the Copper-Rich Corner”, Calphad, 25(1), 67-78 (2001) (Calculation, Phase Relations, Thermodyn., 17) Perrot, P., Batista S., Xing, X., “Cu - P (Copper - Phosphorus)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.16300.1.20, (2002) (Phase Diagram, Crys. Structure, Assessment, 7) Watson, A., Wagner, S. Lysova, E. Rokhlin, L. “P - Sn (Phosphorus - Tin)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.29529.1.20, (2002) (Phase Diagram, Crys. Structure, Assessment, 7)
Table 1: Investigation of the Cu-P-Sn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1910Hud]
presumably melting of Cu4Sn + Cu3P + Cu(Sn).
Ternary eutectic claimed at Cu81Sn14.2P4.8 mass% 620°C.
[1926Gla1] [1926Gla2]
melting of 50 g alloys from electrolytic Sn, Cu and Cu74.5P15.4 master alloy in graphite crucibles under reducing CO-containing atmosphere. Thermal analysis, chemical analysis, metallography.
Determination of melting trough from binary Cu-Cu3P eutectic to peritectic (Cu)+L; ternary eutectic point Cu80.7Sn14.8P4.5 at 628°C.
[1933Ver]
melting of electrolytic Cu, Sn-ingots and a Cu89.9P10.1 (mass%) master alloy in ceramic (Schamott=fireclay) crucibles under charcoal. Thermal analyses with calibrated Pt/PtRh thermocouples. Casting of 12mm diameter rods in cold iron crucibles. Chemical analyses. Annealing of alloys packed on charcoal powder in wire-wound furnace. Annealing times: 660°C (24h for <15% Sn, 10 min > 15% Sn); 600°C (24h for <15% Sn, 2 h > 15% Sn); 550°C (60h for <15% Sn, 24 h > 15% Sn); 300°C (120 h for <15% Sn, 12 h > 15% Sn). Metallography.
Determination of isopleths with constant contents of 1, 2 and 3 mass% P and from 0 to 25 mass% Sn. Determination of isopleths with constant content of 5, 10 and 20 mass% Sn from 0 to 5 mass% P. Determination of liquidus and solidus isotherms in the Cu rich region (<25 mass% Sn, <9 mass% P). Determination of partial isothermal sections in the Cu rich region (<25 mass% Sn, <9 mass% P) at 650, 600, 550, 300°C.
Landolt-Börnstein New Series IV/11C3
MSIT®
360
Cu–P–Sn
Reference
Method/Experimental Technique
[1937Ham]
Alloys were melted from 99.9 min purity Determination of liquidus and solidus elements in graphite crucibles under argon isotherms in the Cu rich region (< 35 and cast into mild steel molds. Chemical mass% Sn, < 15 mass% P). analyses, TA.
[1950Sho]
no details given on alloy preparation but chemical analyses on 7 standard refined grades of tin are given. Various copper contents < 0.129 mass% Cu. Cast cylinders of bronzes containing 0.1 mass% P, 0.5 mass% P.
Examination of the influence of Sn and P contents as well as the effect of casting on the hot-working capacity (including hot forging, hot extrusion and hot rolling) of phosphorized tin bronzes.
[1956Fre]
Various methods of casting and investigation techniques.
Study of the influence of basic composition, microstructure and impurities on the properties of copper tin bronzes. Control of metal and casting quality.
[1954Mur]
Alloys were melted from 99.9 Sn, electro Cu and Cu15P under argon and cast into green sand moulds molds at 1100°C and annealed for 24 h in vacuum. Above 550°C annealed in salt bath and quenched in ice-water. Chemical analyses, DTA, TA, metallography were employed.
Determination of isopleths with constant contents of 6 mass% Sn, 10 mass% Sn and from 0 to 4 mass% P. Determination of isopleths with constant content of 1 mass% P from 0 to 25 mass% Sn. Determination of the liquidus surface in the Cu rich region (<35 mass% Sn, <14 mass% P).
[1979Hoe] [1980Hoe]
Single crystals were obtained from a reaction of Cu, P in Sn-flux. The compound can also be obtained from elements in a sealed quartz tube (1000°C, 24 h). X-ray single crystal data, R=0.047.
Crystal structure of Cu4SnP10.
[1987Tak1] [1987Tak2]
Alloys were melted from 99.99 min purity elements in graphite crucibles under argon and cast into mild steel molds. Chemical analyses, DTA, EMPA, metallography and X-ray powder diffraction were employed.
Determination of isopleths with constant contents of 10 mass% Sn, 15 mass% Sn and from 0 to 8 mass% P. Determination of isopleths with constant content of 3, 5 and 7 mass% P from 0 to 15 mass% Sn. Determination of the liquidus surface in the Cu rich region (<25 mass% Sn, <10 mass% P). Determination of partial isothermal sections in the Cu rich region (<25 mass% Sn, <9 mass% P) at 650, 600, 550, 300°C.
MSIT®
Temperature/Composition/Phase Range Studied
Landolt-Börnstein New Series IV/11C3
Cu–P–Sn
361
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1999Zha]
Ball milling of elemental powder blends. Characterization by XPD and TEM.
Investigation of amorphous forming ability and determination of glass forming region in the system.
[2001Mie]
Calphad technique. Experimental data for the ternary refer only to phase relations published by [1987Tak1, 1987Tak2].
Thermodynamic description of the Cu rich corner up to 32 mass% Sn and 14 mass% P. Vertical sections were calculated for constant 10 and 15 mass% Sn (0 to 10 mass% P); and for 5, 7 mass% P (up to 20 mass% Sn). Partial isopleths in the Cu rich corner were calculated for 714, 678, 648, 644, 585, 519°C.
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, (Cu) < 1084.62
cF4 Fm3m Cu
a = 361.46
0 to 3.5 at.% P [Mas2] 0 to 9.1 at.%Sn [Mas2] melting point [1994Sub]
(P)(I)
cP1 Pm3m Po
a = 237.7 a = 229
high pressure phase, above ~10 GPa [V-C2]
(P)(II)
hR6 R3m As
a = 337.7 c = 880.6
high pressure phase, 5 to 11.1 GPa [V-C2]
(P)(red) < 417
c*66
a = 1131
Sublimation at 1 bar, triple point at 576°C, >36.3 bar; triple point at 589.6 at 1 atm [Mas2], [V-C2]
(P) (white)
c** ? P(white)
a = 718
at 25°C [Mas2]
(P) (black)
oC8 Cmca P(black)
a = 331.36 b = 1047.8 c = 437.63
at 25°C [Mas2]
( Sn)
tI2 ? Sn
a = 370 c = 337
at 25°C, 9.0 GPa [Mas2]
(Sn) 231.96 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
[Mas2]
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–P–Sn
362 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Sn) < 13
cF8 Fd3m C(diamond)
a = 648.92
Cu3P < 1022
hP8 P3m1 Cu3P
a = 409.2 c = 718.6
[Mas2]
congruent melting at 1 bar, 25 to 31 at.% P [2002Per], at 560°C [V-C2]
Cu3P
hP24 P63cm Cu3P
a = 695.93 c = 714.3
Low temperature phase, [2002Per]
CuP2 < 891
mP12 P21/c CuP2
a = 580.04 b = 480.63 c = 752.63 = 112.70°
Congruent melting at 15.2 bar [2002Per]
Cu2P7 < 850
mC72 C2/m Cu2P7
a = 1265.8 b = 725.6 c = 1463.0 = 107.46°
[2002Per]
Sn4P3 559
hR21 R3m Bi3Se4
a = 396.3 0.2 c = 3530 2
[2002Wat]
Sn3P4 ~ 560
hR21 R3m
-
[2002Wat]
SnP3
hR24 R3m SnP3
a = 737.85 0.05 [2002Wat] c = 1051.25 0.11
SnP, metastable?
hP12 P3m1 CdI2-derivative
a = 439.22 0.07 c = 604.0 0.3
[2002Wat]
SnP, metastable?
hP16 P3m1 ?
a = 878.0 0.1 c = 598.0 0.1
[2002Wat]
SnP, metastable?
cF8 Fm3m NaCl
SnP, metastable?
Sn7P10, metastable
MSIT®
tI4 I4mm GeAs h**
a = 553.59
a = 383.1 0.1 c = 596.3 0.1 a = 443.30 c = 283.94
High pressure [2002Wat] High pressure [2002Wat] [2002Wat]
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363
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
, Cu17Sn3 798 - 586
cI2 Im3m W
a = 302.61
, Cu3Sn 755 - 520
cF16 Fm3m BiF3
13.1 to 16.5 at.% Sn [Mas2] at 710°C and 15 at.% Sn [V-C2] 15.5 to 27.5 at.% Sn [Mas2]
a = 611.76
at 710°C [V-C2]
, Cu41Sn11 ~585 - 350
cF416 F43m Cu41Sn11
a = 1798.0
20 to 21 at.% Sn [Mas2, V-C2]
, Cu10Sn3 ~640 - 582
hP26 P63 Cu10Sn3
a = 733.0 c = 786.4
20.3 to 22.5 at.% Sn [Mas2, V-C2]
J, Cu3Sn
oC80 Cmcm Cu3Sn
a = 552.9 b = 4775 c = 432.3
24.5 to 25.9 at.% Sn [Mas2, V-C2]
, Cu6Sn5(h) 415 - 186
hP4 P63/mmc NiAs
a = 419.2 c = 503.7
43.5 to 45.5 at.% Sn [Mas2, V-C2]
’, Cu6Sn5(r) < 189
hP*
-
44.8 to 45.5 at.% Sn, ordered form with superlattice based on [Mas2]
Cu7Sn13
hP2 P6m2 WC
a = 318.8 c = 297.6
Metastable, rapidly quenched from the liquid phase [V-C2]
’, ~Cu7Sn2
hP8 Na3As
a = 428.1 c = 784.2
Metastable, martensitic form of quenched [V-C2]. The ’ martensite is found in two composition ranges between 15 and 25 at.% Sn [Mas2]
* -1, Cu4SnP10
cF60 F43m Cu4SnP10
a = 1028.3
[1979Hoe]
Two other forms of metastable martensite obtained from quenching the , Cu17Sn3 phase and a series of metastable phases quenched from the vapor phase are also reported [Mas2].
Landolt-Börnstein New Series IV/11C3
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Cu–P–Sn
364 Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Cu
Sn
P
L + + Cu3P
642-655
U1
L Cu3P
84 98.51 75 88.14
12 1.08 0 10.31
4 0.41 25 1.55
L + Cu3P +
?
U2
-
-
-
-
+ Cu3P +
587
E1
84.05
14.12
15.15
+ Cu3P +
~510
E2
-
-
-
-
Cu-Sn
A-B-C
Cu-P
Cu-P-Sn A-B-C
798 p1 L + (Cu) β 756 p2 L+βγ 714 e1 L (Cu) + Cu3P L + (Cu) β + Cu3P
~655
?
L + β Cu 3P + γ
U1
U2
L +Cu3P+γ 586 e2 β γ + (Cu)
β (Cu) + γ + Cu3P
~585
E1
γ+δ+Cu3P
520 e3 γ δ + (Cu) 500
γ (Cu) + δ + Cu3P
E2
(Cu)+δ+Cu3P
Fig. 1: Cu-P-Sn. Tentative reaction scheme for the Cu rich region
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Cu–P–Sn Cu Sn P β Cu3P
Fig. 2: Cu-P-Sn. Liquidus and solidus surfaces for the Cu rich region
365 75.00 0.00 25.00
80
Data / Grid: at.% Axes: at.%
20
e1,714°C
90
10
U1
α 10
Cu
β Cu Sn P
Fig. 3: Cu-P-Sn. Isothermal section for the Cu rich region at 550°C
p2,755
p1,798
70.00 0.00 30.00
20
Cu Sn P
75.00 25.00 0.00
Cu Sn P
70.00 30.00 0.00
Data / Grid: at.% Axes: at.%
Cu3P
80
20
90
10
α
Cu
Landolt-Börnstein New Series IV/11C3
10
γ
20
δ
ε
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Cu–P–Sn
366
Cu Sn P
Fig. 4: Cu-P-Sn. Isothermal section for the Cu rich region at 600°C
70.00 0.00 30.00
Data / Grid: at.% Axes: at.%
Cu3P
80
20
90
10
(Cu)+β +Cu3P
α 10
Cu
Fig. 5: Cu-P-Sn. Isopleth at 6 mass% Sn (up to 5 mass% P), plotted in at.%
β
δ
20
Cu Sn P
ε
ζ
70.00 30.00 0.00
1000
L L+α
Temperature, °C
750
L+α +β
L+α +Cu3P
α +β
α +β +Cu3P
α 500
α +Cu3P
250
α +δ +Cu3P
0
Cu 96.70 Sn 3.30 0.00 P
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4.0
8.0
P, at.%
10 Cu 86.85
Sn 3.13 P 10.01
Landolt-Börnstein New Series IV/11C3
Cu–P–Sn
Fig. 6: Cu-P-Sn. Isopleth at 10 mass% Sn (up to 5 mass% P), plotted in at.%
367
L 1000
L+α
Temperature, °C
750
L+α +β
α +β
500
α
L+α +Cu3P
α +β +Cu3P
α +Cu3P α +δ +Cu3P
250
α +δ 0
Cu 94.39 Sn 5.61 0.00 P
Landolt-Börnstein New Series IV/11C3
4.0
8.0
P, at.%
10 Cu 84.48
Sn 5.32 P 10.20
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368
Cu–Pb–Sn
Copper – Lead – Tin Kostyantyn Korniyenko Introduction The nature of the phase relations in the Cu-Pb-Sn system has been of interest for many years, in particular, because Cu rich Cu-Pb-Sn alloys make excellent bearing metal for use in adverse conditions. More recently, interest in the phase relations of this system is due to the reaction between Pb-Sn alloys and copper through the soldering process; it being one of the oldest metallurgical processes for joining metal parts, in particular owing to the wide possibilities of the use of solder in modern microelectronic technology. For example, in circuit board manufacturing, solder is used to connect the leadframe of a chip to the circuit board. In flip chip technology, solder joins the chip directly to the board [2001Tu]. The constitution of the Cu-Pb-Sn system was first investigated by [1910Gio]. They determined a miscibility gap in the Cu rich corner by extrapolation of the isotherms for the separation of two liquids. Their results were in accordance with a later study of [1931Bri] but differed from the data of [1929Gue]. [1963Dav] conducted an investigation of lead rich alloys and presented the liquidus surface projection for the Pb corner. A similar study was conducted by [1983Mar] in conjunction with a determination of the temperature dependence of the monovariant process L + Cu3Sn Cu6Sn5 in the ternary system. [1983Pik] investigated the liquidus surface in the range of compositions bordering the Pb-Sn system. Critical assessments of the publications relating to the constitution of the system were carried out by [1932Blu, 1970Hof, 1979Cha, 1994Cha]. A review of studies of interfacial reactions was carried out by [2000Kat]. Thermodynamic calculations of isothermal sections were carried out by [1997Lee, 2001Tu, 2002Zen], and of the liquidus surface by [2001Her] and [2001Tu]. Experimental details of different studies are presented in Table 1. Information concerning the phase equilibria in this system is so far incomplete. In particular, the liquidus surface in the Sn corner, the solidus surface over the whole concentration range and also, isothermal and polythermal sections need further experimental investigation. Acquiring this knowledge would extend the practical application of the Cu-Pb-Sn alloys. Binary Systems The Cu-Pb system is accepted from [2006Sch]. Data relating to the Cu-Sn and Pb-Sn systems are taken from [Mas2]. Solid Phases No ternary phases have been reported. Crystallographic data concerning the known unary and binary phases are listed in Table 2. Invariant Equilibria Temperatures, types of reactions and compositions of the phases taking place in the invariant equilibria are listed in Table 3. This table is compiled on the basis of [1931Bri, 1963Dav, 1979Cha, 1983Mar] and [1983Pik] with the conversion of compositions from mass to atomic per cent. [1979Cha] assigned E type for the invariant reactions U3, U4, U5 in his table. However, in the drawing of the liquidus surface these reactions were shown as U type, which were retained in the present evaluation. [1983Pik] also reported the existence of a transition reaction L + J + (Pb) at 250°C, with a composition of the liquid of about 0.79Cu-54.40Pb-44.81Sn (at.%). However, these data are in doubt because the authors indicated the position of the primary crystallization of the J phase in conflict with the binary diagrams accepted in this assessment. According to the constitution of the Cu-Sn binary system, the ' phase (instead of the phase) seems to take part in the invariant equilibrium L (Sn) + (Pb) + '. A partial reaction scheme is presented in Fig. 1. However, it should be considered as tentative only, in particular in the Pb corner, where the
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Cu–Pb–Sn
369
presence of the phase primary crystallisation field at the liquidus surface seems to be very doubtful, because it does not exist in equilibria with the liquid phase in the binary system; and there is no data concerning the formation of this phase in the ternary system. Liquidus, Solidus and Solvus Surfaces The liquidus surface projection and isotherms of the miscibility gap of the system are shown in Figs. 2 and 3, respectively. They are based on the results of experimental investigations by [1910Gio], [1931Bri], [1983Mar] (as quoted by [1986Zak]) and [1983Pik] as well as the reviews of [1932Blu], [1970Hof], [1979Cha] and [1994Cha]. The miscibility gap in the ternary system is wider than in the binary Cu-Pb system, and it extends to high temperatures - from 980.5°C in the binary system up to ~1130°C for the c1 critical point in the ternary. The results of [1931Bri] indicate that the miscibility gap closes around 1130 to 1140°C and that the critical point lies at a composition of 60Cu-26Pb-14Sn (at.%). As quoted in the review of [1932Blu], the authors of [1913Bor] and [1929Bau] predicted a much higher temperature of closing, probably due to misinterpretation of their electrical conductivity measurements of the melt at high temperatures. Liquidus and solidus temperatures of lead rich alloys were determined by [1983Mar] using thermal analysis. Liquidus temperatures for alloys up to about 3.2 at.% Cu are listed in Table 4. These data qualitatively agree with the steep liquidus surface in Fig. 2. The reported solidus temperature of the alloys is quite consistent with the eutectic temperature E1 (Fig. 2). The liquidus surface projection was calculated from thermodynamic parameters by [2001Her] and [2001Tu]. The positions of the fields of primary crystallization of the ,Jandphases need further experimental verification. Also, the existence of the
primary crystallisation field needs experimental verifcation. The calculated position of the L’ + L’’ region in the ternary system is shifted to the Pb corner in comparison with experimental data. Isothermal Sections Figure 4 presents the isothermal section at 955°C which is close to the monotectic temperature in the Cu-Pb binary system (952.7°C), compiled by [1994Cha] on the basis of experimental results from [1910Gio] and [1931Bri] and the review of [1932Blu]. The miscibility gap and the monotectic temperature (952.7°C compared with 955°C in [1994Cha]) are consistent with the Cu-Pb binary phase diagram [2006Sch]. Isothermal sections for the whole composition range at temperatures of 400, 350, 283°C (Figs. 5 to 7 respectively) and 200, 170°C (Figs. 9 and 10 respectively) were calculated from thermodynamic parameters by [2001Tu]. The isothermal section at 250°C calculated by [1997Lee] is presented in Fig. 8. The isothermal section for 100°C proposed by [1994Cha] is shown in Fig. 11. Figure 12 shows the partial isothermal section in the Sn rich corner at 220°C, calculated by [2002Zen]. All of these sections are in a good agreement with accepted binary systems. On comparing with figures from the original work, compositions were converted from mass to atomic per cent (Figs. 5, 6, 7, 9, 10, 12), and some stretching of narrow phase fields is marked (for example, J + (Cu), + J etc). The (Pb) solid solution takes part in equilibria with all phases and components of the Cu-Sn binary system at temperatures of 250°C and below, but at higher temperatures equilibrium between (Pb) and the phase does not occur because this phase is in equilibrium with the liquid of the Pb-Sn system. Thermodynamics Using the transpiration method, [1970Aza] (as quoted by [1979Cha]) measured the vapor pressure of Pb over Cu-Pb-Sn melts at 1100°C. Isoactivity plots of Cu, Pb and Sn, as well as the excess Gibbs energy, also shown by [1970Aza], are derived from these data (probably by Gibbs-Duhem integration) and shown in Figs. 13 to 16, respectively. The approximate position of the L’ + L’’ region at 1100°C is presented in Fig. 13 (by a dotted line), on the basis of the experimentally determined temperature dependence of the liquid miscibility gap (see Fig. 3). Comparing this with the data presented by [1970Aza], the position of the L’ + L” loop is somewhat shifted. However, there is an inconsistency in the original data of [1970Aza]
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370
Cu–Pb–Sn
because the smoothly curved isoactivity lines were also shown inside the gap; thermodynamics require straight (tie) lines. The slope of these tie lines, estimated from the entry and exit points to the gap are also inconsistent with the location of the critical points in Fig. 2. The values for the activity coefficient of Pb and its compositional dependence in the Cu corner at 1100°C, reported by [1970Aza] as well as the values of the Pb vapor pressure over Cu rich alloys containing up to 10 at.% Sn and 10 at.% Pb obtained by [1974Hay] using the Knudsen effusion technique at 1130°C are listed in Table 5. It was noted in [1974Hay], that the activity of Pb in binary dilute solutions showed a remarkably positive deviation from Raoult’s law and obeyed Henry’s law to at least 0.05 at.% of Pb. The calculated Pb isoactivity curves in the Pb rich liquid phase at 1100°C show good consistency with the experimental data of [1974Hay]. The influence of Sn on the activity of Pb in dilute solid solutions has been measured by [1987Ryc1, 1987Ryc2] using an improved isopiestic method. The technique of [1987Ryc2] involved total minimum error in the temperature range 1100 to 1200°C. The first-order interaction parameter is given by [1987Ryc2] as: JPbSn = 8720/T – 1.42 based on their experimental results, combined with a value of ln ( Pb(Cu)0) = 3360/T – 0.6 for the activity coefficient of Pb at infinite dilution derived from data for binary liquid Cu-Pb alloys. An analysis of the temperature dependence of the JPbSn interaction parameter calculated by [1987Ryc1] on the basis of experimental results and using Krupkowski’s method, has shown that the parameter value decreases with increasing temperature. Indirectly derived activity data on the basis of analysis of the phase relations in the Pb rich corner over the temperature interval from 327 to 527°C, have been reported as activity coefficients of Cu and Sn [1963Dav]. Thermodynamic calculations of the liquidus surface [2001Her, 2001Tu] and of isothermal sections at 250°C [1997Lee], 400, 350, 283, 200 and 170°C [2001Tu] as well as at 200°C [2002Zen] were conducted using Calphad type calculations with a description of the liquid solution phase without ternary interactions. The calculations of [2001Her] show also a comparison with the activity data [1970Aza], misprinted as [1974Hay], however, a comprehensive Calphad assessment including all of the experimental thermodynamic and phase diagram data is still missing. Notes on Materials Properties and Applications Information relating to the investigation of properties of the Cu-Pb-Sn alloys is listed in Table 6. Cu-Pb-Sn alloys are applied, in particular, to cases where a metal is exposed to friction (and therefore to wear) under high specific pressure, with shocks and vibrations and at elevated temperatures. So, friction bearings made of copper alloys are used in automotive applications on a large scale, and the Cu-Pb-Sn alloy system is one of the most common [1992Boe]. New trends in development in this area involve high-strength alloys such as Cu-Ni-Sn bronzes, in particular, by the alloys (mass%) 79Cu-20Pb-1Sn and 84Cu-15Pb-1Sn (in at.% they are 92.2Cu-7.2Pb-0.6Sn and 94.2Cu-5.2Pb-0.6Sn, respectively). Another important aspect of the application of Cu-Pb-Sn alloys is related to the reactions between Pb-Sn alloys and copper [2001Tu, 2002Zen] during soldering. Solders are widely used to connect chips to their packaging substrates in flip chip technology as well as in surface mount technology. Miscellaneous In the review of [1991But], two models were proposed which can be used to predict the occurrence of liquid immiscibility in ternary metallic alloys. The Cu-Pb-Sn system is denoted as type III in their classification, where liquid immiscibility first appears at elevated temperatures within the ternary diagram, and as the
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Cu–Pb–Sn
371
temperature decreases, the region of liquid immiscibility expands until eventually it intersects one or more of the constituent binaries. References [1910Gio]
[1913Bor]
[1929Bau]
[1929Gue] [1931Bri] [1932Blu] [1932Isa]
[1934Bug]
[1939Hom]
[1939Isa]
[1944Ray]
[1954Tyz] [1957Bag]
[1957Kno] [1959Har]
[1960Cah] [1963Dav]
Landolt-Börnstein New Series IV/11C3
Giolitti, F., Marantonio, M., “Investigation of Bronze Special-I-Bronze” (in Italian), Gazz. Chim. Ital., 40(1), 51-77 (1910) (Phase Diagram, Experimental) as quoted by [1932Blu] Bornemann, E., Wagenmann, K., “The Electric Conductivity of Alloys in Liquid State” (in German), Ferrum, 11, 276-282 (1913) (Phase Relations, Experimental, Electr. Prop.) as quoted by [1932Blu] Bauer, O., Hansen, M., “The Influence of Third Elements on the Constitution of Brass” (in German) in “Mitteilungen der Deutschen Materialpruefungs-Anstalten, Sonderheft IX”, 3-22 (1929) (Phase Relations, Experimental, Electr. Prop.) as quoted by [1994Cha] Guertler, W., Leitgebel, W., “From the Earth to the Metal Workshop” (in German), Leipzig, 391 (1929) (Phase Diagram, Phase Relations, Experimental) as quoted by [1932Blu] Briesemeister, S., “The Miscibility Gap in the Systems Pb-Cu and Pb-Cu-Sn” (in German), Z. Metallkd., 23(8), 225-230 (1931) (Phase Diagram, Phase Relations, Experimental, *, 14) Blumenthal, B., “The Constitution of the Copper-Lead-Tin System”, Met. Alloys, 3, 181-188 (1932) (Phase Diagram, Review, *, 9) Isaichev, I.V., Kurdyumov, G.V., “Transformation in Copper-Tin Eutectoid Alloys. I” (in Russian), Zh. Fiz., 5, 21-26 (1932) (Crys. Structure, Experimental) as quoted by [1957Bag] Bugakov, V., Isaichev, I.V., Kurdyumov, G.V., “Transformation in Copper-Tin Eutectoid Alloys. II” (in Russian), Zh. Fiz., 5, 22-30 (1934) (Crys. Structure, Experimental) as quoted by [1957Bag] Homer, C.E., Plummer, H., “Embrittlement of Tin at Elevated Temperatures and its Relationship to Impurities”, J. Inst. Met., 64, 169-200 (1939) (Phase Relations, Experimental) as quoted by [1994Sau] Isaichev, I.V., “Transformation in Copper-Tin Eutectoid Alloys. IV. Decomposition of the Phase on Annealing” (in Russian), Zh. Techn. Fiz., 9, 1867-1872 (1939) (Phase Relations, Crys. Structure, Experimental) as quoted by [1994Sau] Raynor, G.V., “The Cu-Sn Phase Diagram”, Annotated Equilibrium Diagram Series, 2, The Institute of Metals, London (1944) (Phase Diagram, Phase Relations, Review) as quoted by [1994Sau] Tyzack, C., Raynor, G.V., “The Lattice Spacings of Lead rich Substitutional Solid Solutions”, Acta Crystallogr., 7, 505-510 (1954) (Crys. Structure, Experimental, 18) Bagariatskii, Yu.A., “About Crystal Structure of the Metastable Phase Formed During the Tempering of Cu-Sn Alloys Containing 24-27 % Sn” (in Russian), Kristallografiya, 2(2), 283-286 (1957) (Crys. Structure, Experimental, 10) Knoedler, H., “The Structure Relation between - and J-Phase in the System Copper-Tin” (in German), Acta Crystallogr., 10, 86-87 (1957) (Crys. Structure, Experimental, 6) Hartmann, H., Ensslin, F., Wunderlich, E., “Experiments on the Removal of Cu from Pb in the Presence of Sn, As and Sb” (in German), Z. B. Metallhuttenw., 12, 374-381 (1959) (Phase Relations, Experimental) as quoted by [1994Cha] Cahn, J.W., Theaftis, H.N., “The Solubility of Tin in Solid Lead”, Trans. AIME, 218(2), 376-377 (1960) (Crys. Structure, Experimental, 11) Davey, T.R.A., “Phase Systems Concerned with Copper Drossing of Lead”, Bull. Inst. Min. Met., 72, 553-620 (1963) (Phase Diagram, Phase Relations, Theory) as quoted by [1979Cha]
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372 [1963Der]
[1967Deb]
[1970Aza]
[1970Hof] [1970Tok]
[1971Van] [1973Gan]
[1973Van] [1974Hay]
[1975Bra]
[1977Boo]
[1979Cha]
[1980Zak]
[1983Kuw]
[1983Mar]
[1983Pik]
[1983Wat]
MSIT®
Cu–Pb–Sn Deruytterre, A., “Phase Transformations Occurring at 300°C in an Alloy of Copper plus 16.5 at.% Tin” (in French), Mem. Sci. Rev. Metall., 60(5), 359-370 (1963) (Crys. Structure, Thermodyn., Experimental, 15) De Bondt, M., Deruyttere, A., “Pearlite and Bainite Formation in a Cu-16.5 at.% Sn Alloy”, Acta Metall., 15, 993-1005 (1967) (Crys. Structure, Phase Relations, Experimental) as quoted by [1994Sau] Azakami, T., “Activity Measurements in Liquid Cu-Pb-Bi and Cu-Pb-Sn Alloys at 1100°C. Thermodynamic Studies of Liquid Copper Alloys (6th Report)”, Nippon Kogyo Kaishi, 86, 865-869 (1970) (Thermodyn., Experimental, 11) “Lead-Copper-Tin”, in “Lead and Lead Alloys”, Hofmann, W. (Ed.), Springer-Verlag, New York, 151-153 (1970) (Phase Diagram, Review, *, 6) Tokseitova, P.K., Baimbetov, N., Chernousova, K.T., Presnyakov, A.A., “On Solubility of Tin in Copper”, in “Properties of Non-ferrous Metals and Alloys: Proceedings of Institute for Nuclear Physics of Kazakh. SSR” (in Russian), Vol. 2, Academy of Sciences of Kazakh. SSR, Alma-Ata, 103-106 (1970) (Phase Relations, Experimental, 10) as quoted by [1997Boc] Vandermeulen, W., Ph. D. Thesis, Leuven University, Belgium (1971) (Crys. Structure, Experimental) as quoted by [1973Van] Gangulee, A., Das, G.S., Bever, M.B., “An X-Ray Diffraction and Calorimetric Investigation of the Compound Cu6Sn5”, Metall. Trans., 4, 2063-2066 (1973) (Crys. Structure, Experimental, Thermodyn., 21) Vandermeulen, W., Deruyttere, A., “An 7 Phase in the Cu-16.5 At. Pct Sn Alloy”, Metall. Trans., 4(7), 1659-1664 (1973) (Crys. Structure, Experimental, 12) Hayashi, M., Azakami, T., Kameda, M., “Effects of Third Elements on the Activity of Lead in Liquid Copper Base Alloys” (in Japanese), J. Jpn. Inst. Met., 90, 51-56 (1974) (Thermodyn., Experimental, 16) Brandon, J.K., Pearson, W.B., Tozer, D.J.N., “A Single Crystal Diffraction Study of the Bronze Structure, Cu20Sn6”, Acta Crystallogr., Sect. B., 31, 774-779 (1975) (Phase Relations, Crys. Structure, Experimental, 13) Booth, M.H., Brandon, J.K., Brizard, R.Y., Chieh, C., Pearson, W.B., “ -Brasses with F Cells”, Acta Crystallogr., Sect.B., 33B(1), 30-36 (1977) (Crys. Structure, Experimental, 22) Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Pb-Sn”, INCRA Monograph Series 6 Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Systems, NSRD, Washington, 631-642 (1979) (Phase Diagram, Phase Relations, Review, *, 14) Zakharova, M.I., Dudchenko, G.N., “Phase Structure in Quenched and Aged Copper-Tin and Copper-Tin-Aluminium Alloys” (in Russian), Phys. Met. Metall., 49(1), 174-177 (1980) (Crys. Structure, Experimental, 9) Kuwano, N., Wayman, C.M., “Precipitation Processes in a -Phase Cu-15 at.% Sn Shape Memory Alloy”, Trans. Jpn. Inst. Met., 24(8), 561-573 (1983) (Crys. Structure, Experimental, 22) Marcotte, V.C., Schroder, K., “The Cu-Pb-Sn Ternary System: Low Cu Additions to Pb-Sn Binary System”, in “Alloy Phase Diagrams. Symp., Boston, Mass., Nov. 1982”, New York, 403-410 (1983) (Phase Relations, Experimental, *) as quoted by [1986Zak] Pikunov, M.V., Volkov, V.A., Verbov, N.E., “On Phase Diagram of the Sn-Pb-Cu System” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., 5, 128-129 (1983) (Phase Diagram, Experimental, 1) Watanabe, Y., Fujinaga, Y., Iwasaki, H., “Lattice Modulation in the Long Period Superstructure of Cu3Sn”, Acta Crystallogr., Sect.B., 39, 306-311 (1983) (Phase Relations, Crys. Structure, Experimental) as quoted by [1994Sau]
Landolt-Börnstein New Series IV/11C3
Cu–Pb–Sn [1984Sau]
[1986Zak]
[1987Ryc1]
[1987Ryc2]
[1989Fel]
[1991But]
[1991Tep]
[1992Boe]
[1994Cha] [1994Sau]
[1995Dae]
[1997Boc]
[1997Lee]
[1999Mor]
[2000Kat]
Landolt-Börnstein New Series IV/11C3
373
Saunders, N., Ph. D. Thesis, “Phase Formation in Co-Deposited Alloy Thin Films”, Univ. of Surrey, U.K. (1984) (Phase Relations, Crys. Structure, Experimental) as quoted by [1994Sau] Zakharov, A.M., “Cu-Pb-Sn”, in “Phase Diagrams of Metallic Systems Published in 1985” (in Russian), Petrova, L.A. (Ed.), Vol. 30, Part 2, VINITI, Moscow, 504-506 (1986) (Phase Diagram, Review, *, 1) Rychlewski, M., Koperski, S., “Comparison of Calculation Results of a Component Activity in a Ternary Solution by Krupkowski’s Method with Experimental Results of the Copper-Tin-Lead Diluted Solution Investigations”, Arch. Hutn., 32(1), 9-25 (1987) (Thermodyn., Experimental, 21) Rychlewski, M., Koperski, S., Pomianek T., “Influence of Tin on the Lead Activity in Cu-Pb-Sn Dilute Solutions”, Trans. Jpn. Inst. Met., 28(10), 781-787 (1987) (Thermodyn., Experimental, 15) Felton, L.E., Selsley, A.D., Ficalora, P.J., “Intermetallic Structure of Laser Reflowed Cu/Pb-Sn Solder Joints”, Appl. Phys. Lett., 54(21), 2074-2075 (1989) (Electronic Structure, Experimental, 7) Butt, M.T.Z., Bodsworth, C., “Liquid Immisicibility in Ternary Metallic Systems”, Mater. Sci. Technol., 7(9), 795-802 (1991) (Experimental, Phase Diagram, Phase Relations, Review, 39) Teppo, O., Niemelae, J., Taskinen, P., “The Copper-Lead Phase Diagram”, Thermochim. Acta, 185(1), 155-169 (1991) (Phase Diagram, Phase Relations, Thermodyn., Review, #, 17) Boegel, A., Feind, J., “Material for Friction Bearings Made of Copper Alloys for Use in Automotive Industry” (in German), Metall (Berlin), 46(11), 1132-1135 (1992) (Experimental, 11) Chattopadhyay, S., Srikanth, S., “The Cu-Pb-Sn System”, J. Phase Equilib., 15(5), 553-557 (1994) (Phase Diagram, Review, #, 17) Saunders, N., Miodownik, A.P., “Cu-Sn (Copper-Tin)”, in “Phase Diagrams of Binary Copper Alloys”, Subramanian, S.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 412-418 (1994) (Phase Diagram, Crys. Structure, Thermodyn., Review, #, 55) Daebritz, S., Hauffe, W., “Characterization of Cu-Sn/Pb Diffusion Zones of Microelectronic Contacts by Means of Electron Probe Microanalysis and Ion Beam Sputtering”, Fresenius’ Z. Anal. Chem., 353, 271-277 (1995) (Electr. Prop., Transport Phenomena, Experimental, 6) Bochvar, N.R., “Cu-Sn. Copper-Tin”, in “Phase Diagrams of Binary Metallic Systems” (in Russian), Lyakishev, N.P., (Ed.), Vol. 2, Mashinostroenie, Moscow, 323-326 (1997) (Phase Diagram, Review, *, 10) Lee, B.-J., Hwang, N.M., Lee, H.M., “Prediction of Interface Reaction Products Between Cu and Various Solder Alloys by Thermodynamic Calculation”, Acta Mater., 45(5), 1867-1874 (1997) (Calculation, Phase Diagram, Phase Relations, Thermodyn., Interface Phenomena, 26) Morris, D.G., “Recent Research on Advanced Copper Alloys: Possibilities for Osprey Spray Deposition?”, Powder Met., 42(1), 20-26 (1999) (Phase Relations, Experimental, Electr. Prop., Mechan. Prop., Review, 35) Kattner, U.R., Erikson, G., Hahn, I., Schmid-Fetzer, R., Sundman, B., Swamy, V., Kussmaul, A., Spenser, P.J., Anderson, T.J., Chart, T.G., Costa e Silva, A., Jansson, B., Lee, B.-J., Schalin, M., “Applications of Computational Thermodynamics: Groups 4 and 5: Use of Thermodynamic Software in Process Modelling and New Applications of Thermodynamic Calculations”, Calphad, 24(1), 55-94 (2000) (Calculation, Phase Relations, Review, Thermodyn., 202)
MSIT®
Cu–Pb–Sn
374 [2001Her] [2001Hir]
[2001Lee]
[2001Tu]
[2002Li]
[2002Pra]
[2002Zen]
[2006Sch]
Hertz, J., “What Will be Done in the Future with O.J. Kleppa’s Enthalpy Data Set?”, J. Alloys Compd., 321, 201-222 (2001) (Phase Relations, Review, Thermodyn., 50) Hirose, A., Fujii, T., Imamura, T., Kobayashi, K.F., “Influence of Interfacial Reaction on Reliability of QFP Joints with Sn-Ag Based Pb Free Solders”, Mater. Trans., JIM, 42(5), 794-802 (2001) (Phase Relations, Experimental, Phys. Prop., Interface Phenomena) cited from abstract Lee, J.-H., Park, J.-H., Shin, D.-H., Lee, Y.-H., Kim, Y.-S., “Kinetics of Au-Containing Ternary Intermetallic Redeposition at Solder/UBM Interface”, J. Electron. Mater., 30(9), 1138-1144 (2001) (Experimental, Kinetics, Mechan. Prop., Interface Phenomena, Morphology, 21) Tu, K.N., Zeng, K., “Tin-Lead (Sn-Pb) Solder Reaction in Flip Chip Technolody”, Mater. Sci. Eng. R, R34, 1-58 (2001) (Phase Relations, Review, Calculation, Thermodyn., 109) Li, M., Zhang, F., Chen, W.T., Zeng, K., Tu, K.N., Balkan, H., Elenius, P., “Interfacial Microstructure Evolution between Eutectic SnAgCu Solder and Al/Ni(V)/Cu Thin Films”, J. Mater. Res., 17(7), 1612-1621 (2002) (Phase Relations, Morphology, Interface Phenomena, Experimental, 26) Prakash, K.H., Sritharan, T., “Textured Growth of Cu/Sn Intermetallic Compounds”, J. Electron. Mater., 31(11), 1250-1255 (2002) (Experimental, Morphology, Phase Relations, 18) Zeng, K., Tu, K.N., “Six Cases of Reliability Study of Pb-Free Solder Joints in Electronic Packaging Technology”, Mater. Sci. Eng. R, 38, 55-105 (2002) (Mechan. Prop., Morphology, Phase Diagram, Phase Relations, Review, Thermodyn., Interface Phenomena, 122) Schmid-Fetzer, R., Shcherban, O., Tomashik, V., Jialin, Y., “Cu-Pb (Copper-Lead)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, to be published (2006) (Crys. Structure, Phase Diagram, Assessment, #, 7)
Table 1: Investigations of the Cu-Pb-Sn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1910Gio]
Thermal analysis
810 - 1084°C, 75 to 100 at.% Cu, 0 to 28 at.% Pb, 0 to 43 at.% Sn
[1913Bor]
Electrical conductivity measurements
> 1140°C, Cu rich corner
[1929Bau]
Electrical conductivity measurements
> 1140°C, Cu rich corner
[1929Gue]
Thermal analysis
0 to 72 at.% Sn
[1931Bri]
Chemical analysis of the samples from the melt, thermal analysis, metallography
800 - 1200°C, 0 to 63 at.% Sn
[1959Har]
Determination of Cu and Sn solubility in liquid Pb
Pb rich corner
[1963Dav]
Determination of Cu and Sn solubility in liquid Pb. Metallography, thermal analysis
325 - 625°C, Pb rich corner (< 2 at.% Cu, < 2 at.% Sn)
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Pb–Sn
375
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1970Aza]
Transportation method
1100°C
[1974Hay]
Knudsen effusion-mass loss analysis
0 to 10 at.% Pb, 0 to 10 at.% Sn
[1983Mar]
Thermal analysis, metallography, electronic microprobe analysis
270 - 610°C, Pb rich corner, the alloys containing Cu6Sn5 phase
[1983Pik]
Thermal analysis, metallography
180 - 500°C, 0 to 3.8 at.% Cu
[1987Ryc1] Modified isopiestic method, 1100 - 1200°C, 0 to 10 at.% Sn, 0 to 10 at.% Pb chemical analysis, thermodynamic calculation (Krupkowski’s method) [1987Ryc2] Modified isopiestic method, thermodynamic calculation (Krupkowski’s method)
1100 - 1200°C, 2 to 10 at.% Sn, 2 to 10 at.% Pb
[1989Fel]
Laser reflowed Cu/Pb-Sn solder joints investigation, scanning electron microscopy, energy dispersive X-ray analysis, X-ray diffraction
Room temperature, whole range of compositions
[1995Dae]
Diffusion in soldered microelectronic contacts between Cu and Sn/Pb solders, ion beam etching
150 - 500°C
[1997Lee]
Thermodynamic calculation (Calphad-type)
Whole range of compositions
[2001Her]
Thermodynamic calculation (Calphad-type)
Whole range of compositions
[2001Hir]
Interfacial reactions of quad flat packages (QFP) joints, exposure test
125°C, 37 at.% Pb
[2001Lee]
Redeposition kinetics of intermetallics at the solder Au/Ni/Cu interface during the aging treatment, under bump metallurgy (UBM) structure
230°C, 2.9Cu-34.3Pb-62.8Sn (at.%)
[2001Tu]
Thermodynamic calculation (Calphad-type)
Whole range of compositions
[2002Li]
220°C Interfacial microstructure of eutectic PbSn solders on Al/Ni(V)/Cu thin films, under bump metallurgy (UBM) structure
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Pb–Sn
376 Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[2002Pra]
Growth kinetics of the Cu-Sn intermetallic compounds at the molten Pb-Sn solder/Cu interface, X-ray diffraction
193 - 310°C, whole range of compositions
[2002Zen]
Thermodynamic calculation (Calphad-type)
Sn rich corner (< 9 at.% Cu, < 47 at.% Pb)
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype
Comments/References
(Cu) < 1084.62
cF4 Fm3m Cu
pure Cu at 25°C [Mas2]
at 798°C x = 0, 0 y 0.077 [1944Ray, 1994Sau] at 700°C x = 0, 0 y 0.087 [1970Tok, 1997Boc] at 250°C x = 0, 0 y 0.057 [1970Tok, 1997Boc] at 200°C x = 0, 0 y 0.007 [1944Ray, 1994Sau] at 952.7°C y = 0, 0 x 0.004 [1991Tep] at 326.5°C y = 0, 0 x 0.000046 [1991Tep] at 183°C x + y = 1, 0 x 0.013, 0 y 0.281 [Mas2] a = 361.46 to 370.46 at x = 0, 0 y 0.091 [1994Sau]
Cu1–x–yPbxSny
(Pb) < 327.502
a = 361.46
cF4 Fm3m Cu
a = 495.02
at 25°C [Mas2]
at 183°C x = 0, 0 y 0.293 [1954Tyz] at 173°C x = 0, 0 y 0.264 [1960Cah] at 25°C x = 0, 0 y 0.02 [1960Cah] at 326.5°C y = 0, 0 x < 0.0002 [1991Tep]
CuxPb1–x–ySny
(Pb) HP
hP2 P63/mmc Mg
a = 326.5 c = 538.7
at 25°C, 10.3 GPa [Mas2]
(Sn) 231.9681 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
at 25°C [Mas2]
(Sn), CuxPbySn1–x–y
MSIT®
at y = 0, 0 x 0.0001, [1939Hom, 1994Sau]
Landolt-Börnstein New Series IV/11C3
Cu–Pb–Sn
377
Phase/ Temperature Range [°C]
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype
Comments/References
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
[Mas2]
,Cu1–x–yPbxSny 798 - 586
cI2 Im3m W
x = 0, 0.131 y 0.165 [1994Sau] a = 297.81 to 298.71 x = 0, 0.134 y 0.157 [1994Sau] a = 298.1 to 299.1 x = 0, 0.152 y 0.172 [1939Isa, 1994Sau]
, Cu3Sn 755 - 520
cF16 Fm3m BiF3
a = 606.05 to 611.76 15.5-27.5 at.% Sn at 0 at.% Pb [1994Sau] 16.6-25.0 at.% Sn at 710°C [1994Sau] a = 611.6 0.6
, Cu41Sn11 590 -~ 350
cF416 F43m Cu41Sn11
20-21 at.% Sn at 0 at.% Pb [1994Sau] labelled as “Cu31Sn8” [1979Cha] a = 1798.0
, Cu10Sn3 640 - 582
J, Cu3Sn < 676
, Cu6Sn5(h) 415 - 186
hP26 P63 Cu10Sn3
oC80 Cmcm Cu3Sn
hP4 P63/mmc NiAs
a = 733.0 c = 786.4
a = 552.9 b = 477.5 c = 432.3 a = 419.0 c = 508.6 a = 419.2 0.2 c = 503.7 0.2
Landolt-Börnstein New Series IV/11C3
at 25 at.% Sn, 700°C [1957Kno]
at 20.5 at.% Sn, single crystal selected from the alloy subsequently annealed at 710°C (2 d) and 560°C (5.5 d) and then quenched in water [1977Boo] 20.3-22.5 at.% Sn at 0 at.% Pb [1994Sau] at 23.1 at.% Sn, single crystal selected from the alloy annealed at 6022°C (6 d) and rapidly quenched [1975Bra] 24.5-25.9 at.% Sn at 0 at.% Pb [1994Sau] at 25 at.% Sn [1983Wat, 1994Sau]
43.5-45.5 at.% Sn at 45.45 at.% Sn [1994Sau]
single crystal with 45.45 at.% Sn selected from the alloy annealed at 210°C (49 d) and quenched into iced water [1973Gan]
MSIT®
Cu–Pb–Sn
378 Phase/ Temperature Range [°C]
Pearson Symbol/ Lattice Parameters Space Group/ [pm] Prototype
Comments/References
', Cu6Sn5(r) < 186
h*
~ 45.45 at.% Sn at 0 at.% Pb; superlattice based on NiAs-type structure [1994Sau] single crystal with 45.45 at.% Sn selected from the alloy annealed at 210°C (49 d) and at 175°C (49 d) and then cooled to room temperature for 14 d [1973Gan]
a = 2087.0 c = 2508.1
' (Cu-Sn)
hP2 P63/mmc Mg a = 2087.0 c = 2508.1 a = 2087.0 c = 2508.1
7 (Cu-Sn)
hP12
a = 421 c = 1110 X (Cu-Sn)
hP9
a = 728 c = 258 a = 740 c = 260 2 (Cu-Sn)
ordered rhombic a = 1273 b = 424 c = 600
'(Cu-Sn)
ordered cubic a = 899
MSIT®
metastable ~8-11.5 at.% Sn at 0 at.% Pb [1963Der, 1967Deb, 1994Sau] at ~ 8 at.% Sn [1963Der]
at 11.5 at.% Sn in the alloy with 16.5 at.% Sn heated to 570°C for 10 min, and then aged at 100°C and 130°C [1984Sau, 1994Sau] labelled as “ ’ ” [1973Van] metastable, 15-16 at.% Sn at 0 at.% Pb [1971Van, 1973Van, 1980Zak] in single-crystal with 15.5 at.% Sn quenched from 700°C [1980Zak] metastable on tempering of a quenched single-phase alloy [1932Isa, 1934Bug, 1957Bag] at ~15 at.% Sn [1932Isa, 1957Bag]
in a 15 at.% Sn alloy with memory of shape, annealed at 100°C and quenched into water; labelled as “L” [1983Kuw] metastable in a 15.5 at.% Sn alloy quenched from 700°C and aged during 15 years [1980Zak] metastable in a 19.5 at.% Sn alloy, vapor quenched, below 200°C [1984Sau, 1994Sau]
Landolt-Börnstein New Series IV/11C3
Cu–Pb–Sn
379
Table 3: Invariant Equilibria T [°C]
Reaction
Type
Phase
Composition (at.%) Cu
Pb
Sn
L, L’, L’’
~1130
c1, critical L
~60
~26
~14
L’ + (Cu) L’’ +
772 [1979Cha]
U1
L’ (Cu) L’’
~ 83.33 ~ 6.15 -
~ 3.29 ~ 90.56 -
~ 13.38 ~ 3.29 -
L’ + (Cu) L’’ +
734 [1979Cha]
U2
L’ (Cu) L’’
~ 80.22 ~ 3.08 -
~ 1.84 ~ 88.69 -
~ 17.94 ~ 8.23 -
L’, L’’, ?
~600
c2, critical L
~29
~22
~49
L + + (Cu)
577 [1979Cha]
U3
-
-
-
-
L + + (Cu)
500 [1979Cha]
U4
-
-
-
-
L + J + (Cu)
> 350 on the basis of [1963Dav]
U5
-
-
-
-
L (Pb) + J + (Cu)
326 [1963Dav]
E1
L
0.19
99.78
0.03
L (Sn) + (Pb) + ’
182.2 [1963Dav] E2 or 180 [1983Pik]
L
0.45
26.60
72.95
Table 4: Liquidus temperatures of the Pb rich alloys [1983Mar, 1986Zak] Composition (at.%)
Liquidus Temperature [°C]
Cu
Pb
Sn
0.32
98.81
0.87
481
1.60
96.70
1.70
555
1.58
94.18
4.24
562
1.55
90.13
8.32
560
3.16
95.15
1.69
592
3.08
88.68
8.24
800
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Pb–Sn
380
Table 5: Thermodynamic Properties of Pb in the Cu rich Corner. Standard State Pb (liquid) Thermodynamic Quantity
Data of [1970Aza], at 1100°C
Data of [1974Hay], at 1130°C
Pb0
6.7
4.3
JPbPb JPbSn 'PbSn
–3.6
0.0
3.3
7.0
0
–15
Table 6: Investigations of the Cu-Pb-Sn Materials Properties Reference
Method/Experimental Technique
Type of Property
[1913Bor]
Electrical conductivity measurements
Electrical conductivity
[1929Bau]
Electrical conductivity measurements
Electrical conductivity
[1995Dae]
Ion beam slope cut method, scanning electronic microscopy
Specific electric resistance, growth constants of the Jand' phases
[1999Mor]
Review
Young’s modulus, yield stress, tensile stress, elongation, electrical conductivity
MSIT®
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Cu-Pb
Pb-Sn
Cu-Pb-Sn
Cu-Sn
~1130 (c1)
952.7 e1 l' (Cu) + l"
798 p1 l + (Cu) β
L' + (Cu) L" + β
772
U1 755 p2 l + β γ
L'+L"+β L' + β L" + γ
734
U2 L+β+(Cu)
L'+L''+γ
640 p3 γ ε + L
L+β+γ ~600 (c2) L + β γ + (Cu)
577
326.5 e2 l (Cu) + (Pb)
L + δ ε + (Cu) L+(Pb)+ε 326
L+ε+(Cu)
415 p4 l + ε η
γ+δ+(Cu)
L+δ+(Cu)
L+δ+ε
U4
U5
δ+ε+(Cu)
L (Pb) + ε + (Cu)
E1 227 e3 l ε + (βSn)
(Pb)+ε+(Cu) 183 e4 l (βSn) + (Pb)
Cu–Pb–Sn
L + γ δ + (Cu)
500
350
β+γ+(Cu)
L+γ+(Cu)
L+δ+γ
U3
182.2 or 180 L(βSn)+(Pb)+η'
E2
(βSn)+(Pb)+η'
381
MSIT®
Fig. 1: Cu-Pb-Sn. Partial reaction scheme (tentative)
Cu–Pb–Sn
382
Sn Fig. 2: Cu-Pb-Sn. Liquidus surface projection
Data / Grid: at.% Axes: at.%
e3 p4
γ
η
20
80
δ
ε
β
e4 E2
ε
U3 U 4
40
60
Pb
e2
c2,~600°C
p3
γ
E1
U5
?
60
40
L'+L'' p2 80 p1
U'2
β
20
734 772 U'1
1000
U''2
900°C 950
(Cu) 20
Cu
c1 ,~1130 °C
U''1 980.540
e'1
e''1 60
E1
80
Sn
e2
Pb
Data / Grid: at.% Axes: at.%
Fig. 3: Cu-Pb-Sn. The liquid miscibility gap at various temperatures 20
80
40
60
c2,~600 60
40
955
0 106
80
20
c1,~1130
1110 1005
Cu
MSIT®
20
e'1
40
e''1 60
80
Pb
Landolt-Börnstein New Series IV/11C3
Cu–Pb–Sn
383
Sn
Data / Grid: at.% Axes: at.%
Fig. 4: Cu-Pb-Sn. Isothermal section at 955°C 20
80
40
60
L 60
40
80
20
L'+L''
(Cu)
L+(Cu) 20
Cu
40
60
80
Sn
Pb
Data / Grid: at.% Axes: at.%
Fig. 5: Cu-Pb-Sn. Isothermal section at 400°C L
20
L+η
80
L+ε+η 40
60
η
L+ε
60
40
ε δ
80
δ+ε+L
20
δ+(Cu)+L (Cu)
Cu
Landolt-Börnstein New Series IV/11C3
(Cu)+L 20
40
60
80
Pb
MSIT®
Cu–Pb–Sn
384
Sn
Data / Grid: at.% Axes: at.%
Fig. 6: Cu-Pb-Sn. Isothermal section at 350°C
L 20
80
L+η 40
60
η 60
L+ε+η
40
L+ε
ε 80
20
ε+L+(Cu) (Cu)
(Cu)+L 20
Cu
40
60
80
Sn Fig. 7: Cu-Pb-Sn. Isothermal section at 283°C
Data / Grid: at.% Axes: at.%
L
20
80
L+η
40
Pb
60
η 60
40
L+η+ε
ε L+ε+(Pb)
80
L+ε 20
ε+(Pb)
Cu
MSIT®
(Pb)
ε+(Cu)+(Pb)
(Cu) 20
(Cu)+(Pb)
40
60
80
Pb
Landolt-Börnstein New Series IV/11C3
Cu–Pb–Sn
385
Sn Fig. 8: Cu-Pb-Sn. Isothermal section at 250°C
Data / Grid: at.% Axes: at.%
L
20
80
L+η 40
60
η 60
40
L+η+(Pb)
ε
η+ε+(Pb)
80
η+(Pb)
20
(Pb)
ε+(Pb) (Cu)+ε+(Pb)
(Cu) 20
Cu
40
60
80
Sn
Data / Grid: at.%
(β Sn)
Fig. 9: Cu-Pb-Sn. Isothermal section at 200°C
Axes: at.%
20
L+ η
+( β
Sn )
η+(β Sn)
Pb
80
L L+η 40
60
L+η+(Pb)
η 60
ε
40
η+ε+(Pb)
80
η+(Pb)
20
(Pb)
ε+(Pb) ε+(Cu)+(Pb)
(Cu)
Cu
Landolt-Börnstein New Series IV/11C3
(Cu)+(Pb)
20
40
60
80
Pb
MSIT®
Cu–Pb–Sn
386
Sn
Data / Grid: at.%
(β Sn)
Fig. 10: Cu-Pb-Sn. Isothermal section at 170°C
Axes: at.%
η'+(β Sn) 20
80
η'+(β Sn)+(Pb) 40
60
η' 60
40
η'+ε+(Pb)
ε
η'+(Pb)
80
ε+(Pb)
ε+(Cu)+(Pb) 20
Cu
40
60
(Pb)
80
Sn
Pb
Data / Grid: at.%
(β Sn)
Fig. 11: Cu-Pb-Sn. Isothermal section at 100°C
20
Axes: at.%
(β Sn)+η' 20
80
40
60
(β Sn)+η'+(Pb)
η' 60
ε
40
η'+ε+(Pb)
80
η'+(Pb)
ε+(Cu)+(Pb)
(Pb)
(Cu)
Cu
MSIT®
20
ε+(Pb) 20
40
60
80
Pb
Landolt-Börnstein New Series IV/11C3
Cu–Pb–Sn
387
Sn
Data / Grid: at.%
(β Sn)
Fig. 12: Cu-Pb-Sn. Partial isothermal section at 220°C
Axes: at.%
L+(β Sn)
(β Sn)+η 10
90
L+(β Sn)+η
20
80
L+η
L
30
70
40
60
L+η+(Pb)
Cu Pb Sn
10
50.00 0.00 50.00
20
30
L+(Pb)
40
Sn
Cu Pb Sn
0.00 50.00 50.00
Data / Grid: at.% Axes: at.%
Fig. 13: Cu-Pb-Sn. Activity of Cu at 1100°C. Standard state Cu (liquid) 20
80
0.1
0.2
40
60
0.3 0.4 60
40
0.5 0.6 80
20
0.7 L'+L''
0.8 0.9
Cu
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Pb
MSIT®
Cu–Pb–Sn
388
Sn
Data / Grid: at.%
Fig. 14: Cu-Pb-Sn. Activity of Pb at 1100°C. Standard state Pb (liquid)
Axes: at.%
20
80
0.1 0.2 0.3 40
60
0.4 0.5 60
40
0.6 0.7 0.75
80
20
L'+L'' 0.8 0.9 20
Cu
40
60
80
Sn
Pb
Data / Grid: at.% Axes: at.%
Fig. 15: Cu-Pb-Sn. Activity of Sn at 1100°C. Standard state Sn (liquid)
0.9 20
0.8
80
0.7 0.6
40
60
0.5 0.4
60
40
0.3 0.2
80
20
L'+L''
Cu
MSIT®
20
0.1
40
60
80
Pb
Landolt-Börnstein New Series IV/11C3
Cu–Pb–Sn
389
Sn
Data / Grid: at.% Axes: at.%
Fig. 16: Cu-Pb-Sn. Excess Gibbs energy G (kJ/mol) at 1100°C. Standard states are Cu (liquid), Pb (liquid)
20
80
40
60
60
40
-1 -2
0
1 2
80
3 L'+L''
20
4 5
Cu
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Pb
MSIT®
Cu–Pd–Sn
390
Copper – Palladium – Tin Viktor Kuznetsov Introduction The Cu-Pd-Sn system gained interest in the search for new lead free solders because the performance of Pd-coated components with lead-free alloys is equivalent to, or better than SnPb. This is due to the slightly greater dissolution rates of Pd in high tin alloys. Generally, lead-free alloys perform adequately on OSP (Organic Solderability Preservatives) coated boards, but have improved solderability on metallic coatings such as Sn, Ag or Pd. The intermetallics of the system occur in nature as mineral species. The natural specimens have also varying content of Pt (5% and more) as well as admixtures of Ni, Sb and Pb (see reviews in [1970Yus], [1973Ras]). The only experimental investigation of the system [1980Evs] (Table 1) is performed for rationalization of phase relations and compositions of those minerals. The samples were prepared from powders of components of special purity, melted in the ampoules (material not specified) and subjected to DTA. After that the samples were studied by optical microscopy, X-ray analysis and EPMA. The transition temperatures in the solid state were determined by high-temperature X-ray analysis. X-ray patterns were taken at room temperature before and after the experiment and 100°C below solidus after 1 h annealing. During heating and cooling with 10°C#min–1 rate the changes of the diffraction patterns were monitored in the 2 region of 35 to 42°. In addition, a separate samples with several compositions were synthesized by hydrothermal method (300 to 700°C, p(H2O) = 1 kbar) in presence of HCl. Authors noticed that phase relations in the Pd-Sn-Cu-HCl system differ from those in dry Cu-Pd-Sn system, so these results are not included here; for further details see original work [1980Evs]. Solid Phases At higher temperatures the solid solution of Cu in the Pd3Sn phase extends for all the compositions studied (till ~30 at.% Cu). Six ternary phases are claimed to exist in the low-temperature part; they were characterized mainly by optical properties (color, bireflexivity, anisotropy of reflexivity). In particular, the latter was taken as evidence for symmetry lowering as compared to Pd3Sn compound. X-ray patterns for ternary phases are tabulated, but no structural information seems to be extracted, except qualitative observation of splitting of some lines of Cu3Au structure; in particular, no hkl are ascribed to the lines in the table. A suggested scheme of ordering of the high-temperature Pd3Sn based phase seems to be purely speculative. The phases of the studied part of the Pd3Sn-Cu3Sn section are listed in Table 2. Temperature – Composition Sections The partial Pd3Sn-Cu3Sn section is presented in Fig. 1. This is based on the data of [1980Evs], but with some modifications. First, the spurious two-phase region in declared field of stability of the solid solution of Cu in Pd3Sn, is removed; the reasons of its appearing in the figure are quite unclear from the text. Second, the congruent maximum of formation of the “Pd5Sn2Cu” phase is tentatively added at 550°C instead of a peritectoid reaction, given in the original figure, where two participating phases would be the same. Finally, the liquidus and solidus lines are adjusted to meet the accepted melting temperature of Pd3Sn (Table 1). The low-temperature part of the section can not be considered firmly established and seems to need more detailed investigations, in particular for invariant reactions shown in the figure. References [1970Yus]
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Yushko-Zakharova O.Ye., Avdonin A.S., Bykov V.P., Kulagov E.A., Lebedeva S.I., Chernyaev L.A., Yurkina K.V. “About Composition of Platinum-Based Minerals in
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Cu–Pd–Sn
[1973Ras]
[1980Evs]
391
Copper-Nickel Ores of Talnakh and Norilsk Fields” in “Mineralogicheskiye Issledovaniya (Mineralogical Studies)”, (2), Moscow, Nauka, 58-69 (1972) (Review, 8) Rasin L.V., Begazov V.D., Meshchankina V.I., “Notes about Mineralogy of Platinoid Metals in Talnakh Field” (in Russian), in “Trudy ZNIGRI”, (108), 96-151 (1973) (Review, 15) Evstigneeva, T.L., Nekrasov, I.Ya., “Conditions of Synthesis of Phases and Phase Relations in Pd3Sn-Cu3Sn and Pd-Sn-Cu-HCl Systems” (in Russian), Ocherki Fiz.-Khim. Petrol., 9, 20-35 (1980) (Phase Relations, Experimental, 11)
Table 1: Investigations of the Cu-Pd-Sn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1980Evs]
DTA, optical microscopy, X-ray analysis, EPMA
Pd3Sn-Cu3Sn section, 0 to 66 mol.% of Cu3Sn
Table 2: Crystallographic Data of Solid Phases of the Pd 3Sn-Cu3Sn Section Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Pd,Cu)3Sn
cP4 Fm3m AuCu3
a = 398
a = 385
at 16.7 mol.% Cu3Sn, 1040°C [1980Evs] at 25.0 mol.% Cu3Sn, 1004°C [1980Evs] at 33.3 mol.% Cu3Sn, 1000°C [1980Evs] at 40.0 mol.% Cu3Sn, 1000°C [1980Evs] at 50.0 mol.% Cu3Sn, 890°C [1980Evs]
a = 397.6
[V-C2, Mas2]
a = 394 a = 390 a = 388
Pd3Sn < 1326 Cu3Sn < 676
oC80 Cmcm Cu3Sn
a = 552.9
[V-C2, Mas2]
* Pd5Sn2Cu < 550
-
-
[1980Evs]
* Pd9Sn4Cu3 < 320
-
-
[1980Evs]
* Pd2SnCu < 198
-
-
[1980Evs]
* Pd9Sn5Cu6 < 100
-
-
[1980Evs]
* Pd3Sn2Cu3 < 100
-
-
[1980Evs]
* PdSnCu2
-
-
[1980Evs]
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Cu–Pd–Sn
392
Fig. 1: Cu-Pd-Sn. Partial vertical section Pd3Sn-Cu3Sn
1250
L
Temperature, °C
1000
750
500
Pd5Sn2Cu
550
320 Pd9Sn4Cu3 250
198
Pd2SnCu
Pd3Sn2Cu3
<100 <100 0
Pd 75.00 Cu 0.00 Sn 25.00
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20
PdSnCu2 40
Pd9Sn5Cu6
Cu, at.%
60
Pd 0.00 Cu 75.00 Sn 25.00
Landolt-Börnstein New Series IV/11C3
Cu–Si–Zn
393
Copper – Silicon – Zinc Nataliya Bochvar, Tatyana Dobatkina Introduction [1906Gui, 1912Car, 1930Gou, 1930Vad1, 1930Vad2, 1931Vad1, 1931Vad2, 1942Mas, 1945Smi, 1974Smi] investigated the Cu rich part of the Cu-Si-Zn system containing up to 40 mass% Zn and 6 mass% Si. The boundaries of the (Cu) field have been determined at 800°C by [1930Gou], at 750°C by [1930Vad1, 1930Vad2] and at 400, 600, 800°C by [1945Smi]. [1974Smi] shows the variation of the limiting solubility of the (Cu) region as a function of the composition at 0, 5, 10, 15, 20, 25, 30 (mass%) Zn and up to 5 mass% Si and temperature from 1050 to 300°C. At 800°C, the boundaries of the + and phase fields have also been determined [1930Gou]. These authors report lower solubility of Si in the Cu-Zn alloys than in later investigations [1957Mim, 1962Pop, 1964Pop1]. A part of the Cu-Si-Zn diagram, up to 30 at.% Si and 90 at.% Zn, has been investigated in detail by [1956Mim, 1957Mim, 1960Mim1, 1960Mim2, 1964Mim1, 1964Mim2]; another part, up to 20 at.% Si and 50 at.% Zn, has been studied by [1962Pop, 1964Pop1]. The projection of the liquidus surface has been constructed by [1956Mim]. The range near solid solution has been investigated by [1960Mim1, 1960Mim2] in detail. Isothermal sections at 700, 600, 400°C were constructed by [1957Mim, 1960Mim1, 1960Mim2] and by [1962Pop, 1964Pop1] at 847, 760, 600 and 482°C. [1985Far] quotes a ternary eutectic between “Cu4Si and two unidentified ternary phases” at 19 mass% Zn, 7 mass% Si and 765°C. This is very near to the quasibinary peritectic minimum L + found by [1957Mim] at 21.5 mass% Zn, 6.9 mass% Si and 775°C, which contains only two solid solution phases and . [1969Gue, 1979Dri, 1979Cha] give the review of the Cu-Si-Zn system. [2003Bor] used the thermodynamic calculation for construction of isothermal sections at 670, 720, and 760°C in the composition range up to 60 mass% Zn and 10 mass% Si and vertical section at 1 mass% Si. Phase diagram information from [1964Pop1] and the personal experimental data were used by [2003Bor] in optimization. Experimental methods used in different works are summarized in Table 1. Binary Systems The binary systems Cu-Zn, Cu-Si and Si-Zn are accepted from [2006Leb], [2002Leb] and [1996Jac] respectively. The , , and J phases of the Cu-Si binary system [2002Leb] are designated here as 1, 1, and J1, respectively. Solid Phases Ternary phases do not form. The phase of the binary Cu-Si and Cu-Zn systems forms continuous series of solid solutions at high temperatures [1956Mim, 1957Mim, 1962Pop, 1964Pop1]. Reported by [1956Mim, 1957Mim, 1964Pop1] the formation of the continuous solid solution between phase of Cu-Zn system and high temperature 1 phase of Cu-Si system is not accepted, because they have different crystal structure. The phase has very wide homogeneity range towards Cu-Si system up to alloy with the composition of 1.7 at.% Zn and 17.3 at.% Si in equilibrium with 1 phase at 760°C. The solubility of Zn in the phase is significant; it reaches ~16 at.% Zn at 700°C. The solubility of Zn in the phase is 2.9 at.% at the eutectic temperature, 530°C. In the J1 and 1 phases it is about 2 at.% Zn. The data on the solid phases are summarized in Table 2. Quasibinary Systems Based on metallographic and thermal analyses, [1956Mim] established that along the - Si section at 721°C, 58 mass% Zn and 3.7 mass% Si, a eutectic L + (Si) occurs. However this section can not considered as quasibinary because phase forms by peritectic reaction.
Landolt-Börnstein New Series IV/11C3
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394
Cu–Si–Zn
Invariant Equilibria Four four-phase and one three-phase invariant equilibria (E2, U6, U10, U14, e4) with the participation of liquid have been found by [1956Mim]. The quasibinary peritectic minimum L + at 775°C determined by [1956Mim] is doubtful, because the and 1 phases do not form a continuous solid solution. One can suppose that there is the transaction reaction L + + 1(U1) at 775°C instead of the peritectic minimum. According to [1985Far] in this point exist three phases: Cu4Si and two unidentified ternary phases. The composition of Cu4Si phase is close to 1 phase. One of the ternary phases can be ternary solid solution and another - phase with dissolved Si. Moreover five hypothetical invariant equilibria with the participation of the 1 phase (U2, U3, U4, U5, E1) were given by us supplementary. In solid state seven invariant equilibria have been established by [1957Mim]. However, two of them contradict with the isothermal sections [1964Pop1]. A four-phase reaction + (Cu) + after [1957Mim] occurs at 625°C, whereas after [1962Pop, 1964Pop1] this reaction is below the isothermal section given for 600°C. To obtain the three-phase equilibria [1962Pop, 1964Pop1] at 482°C three four-phase reactions are necessary between 600 and 482°C, whereas the section given by [1957Mim] for 400°C, showing the three-phase equilibria (Cu) + 1 + and J1 + 1 + , implies only one four-phase equilibrium in this range. Invariant reactions in points U7, U9, U4 and E5 are given according to the phase equilibria after [1964Pop1] at the temperatures 600°C and below. The reaction scheme is given in Fig. 1. The hypothetical invariant equilibria are shown by dashes lines. Determined by [1956Mim] eutectic (E4) and transition (U6, U10) reaction temperatures as 559, 667, 598°C, were corrected according to accepted binary Cu-Zn system [2002Leb] as ~548, ~665 and ~574°C. The compositions of the phases are given in Table 3. The compositions of the liquid phase in the invariant equilibria at 597 and 424°C have been corrected to a non-zero Si content, which was neglected by [1956Mim]. At 600, 500, and 480°C [1960Mim1, 1960Mim2] reported identical compositions for and ' or ' and ''. This contradicts the rules of heterogeneous equilibria and therefore it is tentatively corrected in Table 3. Liquidus Surface The projection of the liquidus surface presented in Fig. 2 is based on [1956Mim] and on the assumption of the existence of the primary crystallization field of the 1 phase. Nine fields of primary crystallization, of the (Cu), , , 1, , J, , (Si) and (Zn) phases, are separated from each other by the lines of monovariant equilibria. The hypothetical monovariant equilibria are shown by dashed lines. Isothermal Sections The isothermal sections constructed by [1957Mim] at 600°C and below disagree with those of [1962Pop, 1964Pop1]. [1962Pop, 1964Pop1] identified the phases by micrography and X-ray diffraction, whereas [1957Mim] used only micrography, which is less secure for the identification of phases than X-ray diffraction. Therefore, the results by [1962Pop, 1964Pop1] are preferential. Five isothermal sections constructed at 847, 760°C by [1964Pop1], at 700 by [1957Mim] and 600, 482°C by [1964Pop1] are displayed in Figs. 3 to 7. They are corrected to satisfy the accepted binary systems [2002Leb, 2006Leb]. Some details, not given by [1957Mim] or [1964Pop1] but following from the reaction scheme and the liquidus surface, are given by dashed lines. Temperature – Composition Sections A vertical section from Cu + 20 mass% Zn to Cu + 10 mass% Si has been constructed by [1942Mas]. However, this section disagrees with the binary Cu-Zn [2006Leb] and ternary systems [1957Mim, 1962Pop, 1964Pop1]. Several vertical sections at constant content of Si (2 and 4 mass%) and constant content of Zn (5 and 20 mass%) were constructed by [1957Mim]. Only one section with 2 mass% Si agrees with phase equilibria obtained by [1962Pop, 1964Pop1]. This section constructed in at.% is shown on Fig. 8. Some correction was made in solid state according to the isothermal sections [1964Pop1]. Also, the temperatures of invariant reactions on this section were corrected according to the reaction scheme (Fig. 1). Other vertical sections contradicted isothermal sections at temperatures 600°C and below obtained in [1962Pop, MSIT®
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Cu–Si–Zn
395
1964Pop1] and are not shown in the present assessment. The vertical section at 1 mass% Si calculated by [2003Bor] is not accepted here. It shows the increase the temperature of peritectic reaction from Cu-Zn system when Si is added, whereas this temperature must be lower according to the liquidus surface. Notes on Materials Properties and Applications [1930Gou] studied the effect of Si on the hardness of ternary Cu-Si-Zn alloys in the composition range from ~13 to ~15 mass% Zn and 3.9; 5.1 mass% Si and in the composition range from ~37 to 40 mass% Zn and 0.8 mass% Si. The hardness of alloys was measured after casting and after quenching at 850°C followed by annealing at 750, 600, 550, 500, 450, 350 and 250°C. [1956Fre] made review of copper-base alloys most commonly used for making castings in UK. The constitution and the effect of composition and structure on properties are discussed. Mechanical properties for sand-casting, chill-casting, extruded into rods and annealed at different temperatures of Cu-Si-Zn alloys were studied in many works [1930Vad1, 1930Vad2, 1931Vad1, 1931Vad2, 1942Mas, 1974Smi, 2003Ois]. These alloys with high silicon content, especially when it is all retained in solid solution, have valuable properties. They can be worked hot as well as cold. [1972Wie, 1974Sol, 1977Wie, 1984Mur] investigated the shape memory behavior in the phase region of ternary Cu-Si-Zn alloys containing from 33.4 to 35.9 at.% Zn and from 0.1 to 2.2 at.% Si. Miscellaneous [1964Mim1, 1964Mim2] using thermal analysis, microscopic examination and also partially X-ray diffraction and resistivity measurements studied the influence of Si on the ' and 1 transformations. The ' transformation temperature depends only slightly on the additions of Si. The 1 transformation temperature is lowered by Si or Cu additions. This transformation proceeds by a diffusionless martensitic mechanism. The twin-like figure which appears in the 1 phase of the slowly cooled specimens is attributed to the transformation from to 1. [1970Mim] investigated the phase transformation in high strength brass containing 35-40 mass% Zn, 1-5 mass% Si and about 2 mass% Mn by DTA, electron probe X-ray micro-analysis, microscopic examination. Four phases were observed in the specimens slowly cooled from 850 to 550°C and quenched into ice water, that is (Cu), , (Mn, Si)-intermetallic compound and Si-bearing Cu-Zn solid solution. The morphology and the structure of martensite phases have been examined by [1967Pop, 1968Pop, 1974Cor, 1976Gil, 1980Fuk, 1982Gar]. The investigations were carried out on alloys with 33 to 36 at.% Zn and from 0.6 to 2.2 at.% Si. The martensitic phases are composed of a lamellar mixture of two close-packed structures (orthorhombic and cubic) with different stacking sequence. The martensitic transformation temperature depends on the composition and the temperature of quenching. The mechanism of stabilization of Cu-Si-Zn martensite has been studied by [1991Dut, 1993Dut]. Cu-Si-Zn alloys show a significant tendency to martensite stabilization after ageing at room temperature. After continuous heating (20 K#min–1) of stabilized Cu-Si-Zn martensite to 300°C no significant structural changes were found. [1991Udo] investigated structure of diffusion layer which is formed during thermochemical treatment of Cu-39 mass% Zn alloy at 800 - 860°C for 4 h in powder mixture containing Si. [1991Udo] shown that wear rate and friction coefficient are depended from structure of surface layer. Porous layer has better properties. [2002Yos] investigated season cracking susceptibility of Cu-1Si-30Zn (mass%) alloy after exposure in ammonical environment using optical micrographs, SEM and tensile testing methods. [2003Ois] determined the grain size of specimens Cu-10 mass% Zn-(0.3 to 2.07) mass% Si using the intercept method with optical micrographs and showed that the grain size of these specimens becomes smaller with increasing silicon concentration.
Landolt-Börnstein New Series IV/11C3
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396
Cu–Si–Zn
References [1906Gui] [1912Car] [1930Gou]
[1930Vad1]
[1930Vad2]
[1931Vad1] [1931Vad2] [1942Mas]
[1945Smi]
[1956Mim]
[1956Fre]
[1957Mim]
[1960Mim1]
[1960Mim2]
[1962Pop]
[1964Mim1]
[1964Mim2]
MSIT®
Guillet, L., “General Investigations of Brasses” (in French), Rev. Metall., 3, 159-204 (1906) (Review, 1) Carpenter, H.C.H., “The Effect of other Metals on the Structure of the Constituent in Copper - Zinc Alloys”, J. Inst. Met., 8, 59-73 (1912) (Experimental, Morphology, 2) Gould, H.W., Ray, K.W., “Effects of Silicon on the Properties of Brass”, Metals and Alloys, 1(10), 455-457 (1930) (Phase Diagram, Phase Relations, Experimental, Mechan. Prop., 4) Vaders, E., “A New Silicon - Zinc - Copper Alloy”, J. Inst. Met., 44, 363-383 (1930) (Experimental, Phase Diagram, Phase Relations, Mechan. Prop., Morphology, Phys. Prop., 6) Vaders, E., “A New Silicon-Zinc- Copper Alloy”, Met. Ind. London, 37, 278, 315-318, 365-368 (1930) (Phase Diagram, Phase Relations, Experimental, Mechan. Prop., Morphology, Phys. Prop., 6) Vaders, E., “A New Silicon-Zinc- Copper Alloy”, Met. Ind., 29(3), 108-110 (1931) (Phase Diagram, Mechan. Prop., Phys. Prop., 5) Vaders, E., “A New Silicon-Zinc-Copper Alloy”, Met. Ind., 29(4), 155-156 (1931) (Experimental, Mechan. Prop., Phys. Prop.) Masing, G., Wallbaum, H.J., “On the Structure Transformations in a Silicon - Zinc - Copper Alloy” (in German), Z. Metallkd., 34, 87-89 (1942) (Experimental, Phase Diagram, Crys. Structure, Morphology, Mechan. Prop., *, 4) Smiryagin, A.P., “Phase Diagrams of Copper-Silicon, Copper-Silicon-Zinc and Copper-Silicon-Manganese” (in Russian), in “Spesial Bronzes and Brasses”, Nemilov, V.A. (Ed.), Metallurgiya, Moscow, 5-20 (1945) (Phase Diagram, Phase Relations, Mechan. Prop., 30) Mima, G., Hasegawa, M., “The Diagram of Cu-Si-Zn Alloy (First Report)”, Technol. Repts. Osaka Univ., 6, 313-321 (1956) (Experimental, Phase Diagram, Phase Relations, Morphology, #, *, 5) French, A.R., Member, F.I.M., “The Metallurgigal Control of Quality in the Production of Copper-Base Alloy Castings”, J. Inst. Met., 85, 293-317 (1956) (Review, Phase Diagram, Mechan. Prop., Phys. Prop., Electr. Prop., 124) Mima, G., Hasegawa, M., “The Diagram of Cu-Si-Zn Alloy – II”, Technol. Repts. Osaka Univ., 7, 385-397 (1957) (Experimental, Morphology, Phase Diagram, Phase Relations, #, *, 1) Mima, G., Hasegava, M., “Phase Diagram of the System Cu-Si-Zn in the Phase Region” (in Japanese), Nippon Kinzoku Gakkai- Shi, 23, 585-587 (1959) (Experimental, Phase Diagram, Morphology, 14) Mima, G., Hasegawa, M., “The Diagram of Cu-Si-Zn Alloy – III”, Technol. Repts. Osaka Univ., 10, 157-168 (1960) (Experimental, Phase Diagram, Phase Relations, Morphology, #, *, 8) Pops, H., “The Contributuon of the Primary Solid Solubility Isothermal Limit and Some Intermediate Phase in Cu-Rich Cu-Si-Zn Alloys”, Dissertation Univ. Pittsburg, 1-138 (1962) (Crys. Structure, Phase Diagram, Experimental, 30) Mima, G., Hasegawa, M., “Study of the Diagram of Cu-Si-Zn Alloy in the Range Near Solid Solution” (in Japanese), Nippon Kinzoku Gakkai-Shi., 27, 370-376 (1963) (Crys. Structure, Experimental, Morphology, Electr. Prop., Phase Diagram, Phase Relations, 15) Mima, G., Hasegawa, M., “The Diagram of Cu-Si-Zn Alloy - IV”, Technol. Repts. Osaka Univ., 14, 623-633 (1964) (Experimental, Phase Diagram, Electr. Prop., Morphology, Phase Relations, #, *, 14)
Landolt-Börnstein New Series IV/11C3
Cu–Si–Zn [1964Pop1]
[1964Pop2] [1967Pop]
[1968Pop]
[1969Gue]
[1970Mim]
[1972Wie]
[1974Cor]
[1974Smi]
[1974Sol]
[1976Gil]
[1977Wie]
[1979Cha]
[1979Dri]
[1980Fuk]
[1982Gar]
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397
Pops, H., “The Constitution of Copper-Rich Copper - Silicon - Zinc Alloys”, Trans. Met. Soc. AIME, 230, 813-820 (1964) (Experimental, Phase Diagram, Phase Relations, Morphology, #, *, 14) Pops, H., “Lattice Parameters of Phase in Copper - Silicon - Zinc Alloys”, Trans. Met. Soc. AIME, 230, 267-268 (1964) (Crys. Structure, Experimental, 5) Pops, H., Kittl, J.E., “Influence of Inhomogeneity on the Martensitic Transformation of a Cu-Zn-Si -Phase Alloy”, Trans. Met. Soc. AIME, 239, 1668-1670 (1967) (Experimental, Phase Relations, Morphology, 7) Pops, H., Delaey, L., “Electron Microscopy of Cu-Zn-Si Martensitic”, Trans. Met. Soc. AIME, 242, 1849-1851 (1968) (Experimental, Phase Relations, Crys. Structure, Morphology, 11) Guertler, W., Guertler, M., Anastasiadias, E., “14-29-30”, A Comp. of Const. Ternary Diagr. Met. System, Isr. Pro. Sci. Tr., Jerusalem, 19(8), 555-561 (1969) (Review, Phase Diagram, Mechan. Prop., Morphology, Phase Relations, 6) Mima, G., Yamaguchi, M., Sato, J., On the Phase Transformation Observed in the High Strength Brass During Heat Treatment” (in Japanese), Shindo Gijutsu Kenkyukai-Shi, 9, 182-184 (1970) (Phase Diagram, Experimental, Morphology, 4) Wield, D.V., Gillam, E., “Shape Memory Effect and Pseudoelasticity in Cu-Zn-Si”, Scr. Metall., 6(12), 1157-1160 (1972), (Experimental, Phys. Prop., Mechan. Prop., Morphology, 15) Cornelis, I., Wayman, C.M., “Direct Observation of the 3R-Stacking Sequence in Cu-Zn-Si Martensite”, Mater. Res. Bull., 9(8), 1057-1062 (1974) (Experimental, Crys. Structure, Morphology, 8) Smiryagin, A.P., Smiryagina, N.A., Belova, A.V., “Silicon Containing Brasses” (in Russian), in “Industrial Non-Ferrous Metals and Alloys”, Metallurgiya, Moscow, 106-109 (1974) (Review, Phase Diagram, Mechan. Prop.) Solov’ev, L.A., Khachin, V.N., “Influence of Internal Stresses on the Process of Phase Transformation in Titanium-Nickel and Copper- Zinc-Silicon Alloys”, Phys. Met. Metallogr., 37(5), 184-186 (1974), translated from Fiz. Met. Metalloved. 37(5), 1095-1097 (1974) (Experimental, Phys. Prop., 12) Gillam, E., Wield, D.V., “Premartensitic Instability in beta-Cu-Zn-Si Alloys “, Scr. Metall., 10(11), 965-970 (1976) (Experimental, Crys. Structure, Morphology, Phase Relations, Phase Diagram, 22) Wield, D.V., Gillam, E., “Deformation Behavior of Cu-Zn-Si Allys Close to their Martensitic Transformation Temperature”, Acta Metall., 25, 725-733 (1977) (Experimental, Mechan. Prop., Morphology, 18) Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Si-Zn”, INCRA Monograph Series. Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Systems, NSRD, Washington, Vol. 6, 669-675 (1979) (Review, Phase Diagram, Phase Relations, 16) Drits, M.E., Bochvar, N.R., Guzei, L.S., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Ci-Si-Zn” (in Russian), Binary and Multicomponent Copper-Base Systems, Nauka, Moscow, 158-159 (1979) (Review, Phase Diagram, 11) Fukamachi, M., Kajiwara, S., “Lattice Imaging Study of Boundaries in Martensite in Shape Memoruy Alloys”, Jpn. J. Phys., 19(8), L479-L482 (1980) (Experimental, Phys. Prop., Morphology, 10) Garshina, M.N, Agapitova, N.V., Evsyukov, V.A., “Effect of Loading on the Crystal Structure of Cu-Zn-Si Alloy Martensite” (in Russian), Phase Transformation in Solid State (Fazovye Prevrasheniya v Tverdykh Telakh), Voronezh, 63-68 (1982) (Experimental, Crys. Structure, 2)
MSIT®
398 [1984Mur]
[1985Far]
[1991Dut]
[1991Udo]
[1993Dut]
[1996Jac]
[2002Yos]
[2002Leb]
[2003Ois]
[2003Bor]
[2006Leb]
Cu–Si–Zn Murakami, K., Murakami, Y., Mishima, K., Ikai, Y., “Degradation and Impruvement of Shape Memory Effect of Cu-Base Alloys” (in Japanese), J. Jpn. Inst. Met., 48(2), 115-121 (1984) (Experimental, Mechan. Prop., Phys. Prop., 13) Farkas, D. Birchenall, C.E., “New Eutectic Alloys and Their Heats of Transformation”, Metall. Trans., 16A, 323-328 (1985) (Experimental, Calculation, Morphology, Thermodyn., 18) Dutkiewicz, J., Kostorz, G., “High-Resolution Study of Stabilised Martensite in Cu-Zn-Si Alloys”, Phys. Status Solidi A, 123(1), 63-70 (1991) (Experimental, Crys. Structure, Morphology, 15) Udovitskii, I.V., Fedotov, I.A., “Intensification of Thermochemical Treatment in the Cu-Zn-Si System During Liquid Phase Formation” (in Russian), Izv. Akad. Nauk SSSR, Met., (1), 135-137 (1991) (Interface Phenomena, Experimental, Morphology, Mechan. Prop., 4) Dutkiewicz, J., Chandrasekaran, M., Cesari, E., “Stabilisation of Martensite in Cu-Zn-Si”, Scr. Metall. Mater., 29(1), 19-24 (1993) (Experimental, Crys. Structure, Phase Relations, Morphology, 16) Jacobs, M.H.G., Spencer, P.J., “A Critical Thermodynamic Evolution of the Systems Si-Zn and Al-Si-Zn”, Calphad, 20(3), 307 (1996) (Phase Diagram, Phase Relations, Thermodyn., Calculation, 33) Yoshimura, Y., Kanai, K., Kita, K., Inoue, A., “Influence of Additional Elements on the Season Cracking Susceptibility of Cu-30 Mass%Zn Alloy”, (in Japanese), J. Jpn. Inst. Met., 66(9), 865-868 (2002) (Experimental, Mechan. Prop., Morphology, 11) Lebrun, N., Dobatkina, T.V., Kuznetsov, V.N., Li, Changrong, “Cu-Si (Copper-Silicon)”, MSIT Binary Evoluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Service GmbH, Stuttgart, Dokument ID: 20.12505.1.20 (2002) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 23) Oishi, K., Sasaki, I., Otani, J., “Effect of Silicon Addition on Grain Refinement of Copper Alloys”, Mater. Lett., 57, 2280-2286 (2003) (Experimental, Morphology, Electr. Prop., Phys. Prop., 13) Borggren, U., Selleby, M., “A Thermodynamic Database for Special Brass”, J. Phase Equilib., 24(2), 110-121 (2003) (Phase Diagram, Phase Relations, Thermodyn., 18) Lebrun, N., “Cu-Zn (Copper-Zinc)”, MSIT Binary Evoluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Matrials Scince International Service GmbH, Stuttgart, to be published (2006) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 18)
Table 1: Investigations of the Cu-Si-Zn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
[1912Car]
Microscopic examination. Melting 51.05 mass% Cu, 47.70 mass% Zn, 1.19 mass% Si from copper-silicon, copper and zinc
[1930Gou]
Microscopic examination, hardness test. Melting in gas-fired crucible furnace, annealing at 725 to 825°C for 1 to 12 h and slowly cooling in the furnace
MSIT®
Temperature/Composition/Phase Range Studied
Alloys: 85 mass% Cu-15 mass% Zn up to 10 mass% Si and 60 mass% Cu-40 mass% Zn -up to 10 mass% Si. Annealing at 825°C and quenched in ice brine mixture. Phase boundaries (Cu)/(Cu) +/ at 800°C. Hardness after annealing at 750, 650, 550, 450, 350 and 250°C
Landolt-Börnstein New Series IV/11C3
Cu–Si–Zn
399
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1930Vad1] [1930Vad2] [1931Vad1] [1931Vad2]
Microscopic examination, tensile testing. Melting in a high frequency induction furnace, annealing at 750°C for 8 h and slowly cooling in the furnace
Alloys up to 70-90 mass% Cu, up to mass 5% Si, Zn to balance. Phase boundary of (Cu) at 750°C. Density, melting points. Tensile strength, elongation, Brinell hardness for sand-cast, chill-cast, extruded into rods and annealed at 600, 700, 750°C alloys.
[1942Mas]
Thermal, microscopic, X-ray analyses
The vertical section between alloys 80 mass% Cu-20 mass% Zn and 90 mass% Cu-10 mass% Si
[1945Smi] [1974Smi]
Thermal, microscopic analyses, hardness and electrical resistivity. Tensile testing, hardness. Heat treatment at 400 - 800°C for 12 to 240 h in atmosphere of nitrogen
Alloys up to 30 mass% Zn, up to 5 mass% Si, Cu to balance. Phase boundaries of (Cu) 400, 600 and 800°C. Region of (Cu) in Cu-Si-Zn as a function of the composition and temperature. Mechanical properties.
[1956Mim]
Thermal and microscopic analyses. Melting in electric furnace
Cu-Zn alloys with up to 20 mass% Si. Liquidus surface. Compositions and temperatures of invariant reactions
[1957Mim]
Thermal and microscopic analyses. Melting in electric furnace
Vertical sections at 5 and 20 mass% Zn. Isothermal sections at 400, 600, 700°C. Mono- and invariant reactions in liquid and solid states
[1960Mim1] Thermal and microscopic analyses, Alloys in the range near phase. Projection of [1960Mim2] measurement of electrical resistance. transformation of ’ and ’ ’’ Melting in electric furnace [1964Mim1] Thermal and microscopic analyses, [1964Mim2] electrical resistance measurements. Melting in electric furnace. Homogenizing treatment at 650°C
Alloys in the range near phase. Annealing at 400°C for 12 h and 300°C for 24 h, air-cooling on water quenching. Transformation of phase and order-disorder transformation of solid solution
[1962Pop] [1964Pop1]
X-ray diffraction and microscopic observations. Melting, homogenization in helium filled quartz tubes at 760°C for 10 days
Cu corner in the range up to 50 mass% Zn and 20 mass% Si Annealing at 847, 760, 600, 482°C for 1, 5, 10, 25 days, respectively. Then quenching in to an ice-brine solution. Isothermal sections at 847, 760, 600, 482°C and vertical section through the phase field.
[1964Pop2]
X-ray method. Melting and casting 70 to 80 at.% Cu, 2 to 30 at.% Zn, 1,5 to 11 at.% Si. under reduced pressures of helium in Lattice parameters values for (Cu) phase quartz capsules. Homogenization at 760°C for 10 days.
[1985Far]
DTA, EMA, calculation. Melting in Measurement and calculation of eutectic induction furnace composition, temperature, heat of fusion
[2003Bor]
X-ray spectroscopy analyses, thermodynamic calculation. Melting in high-frequency induction furnace, homogenization treatment at 760°C for 4 days
Landolt-Börnstein New Series IV/11C3
66 to 72 mass% Cu, 0.7 to 1.8 mass% Si, Zn to balance. Calculation of isothermal sections at 670, 720, 760°C
MSIT®
Cu–Si–Zn
400 Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Cu) < 1084.62
cF4 Fm3m Cu
Lattice Parameters [pm]
Comments/References
a = 361.46
Dissolves up to 11.25 at.% Si at 842°C [2002Leb] and up to 38.27 at.% Zn at 454°C [2002Leb], pure Cu, 25°C [Mas2]
a = 369.612
at 35.84 at.% Zn and 300°C [2006Leb]
a = 365.49 to 365.90 at 17.5 at.% Zn and from 2 to 6.5 at.% Si [1964Pop2] a = 363.05 to 363.30 at 5 at.% Zn and from 6 to 11 at.% Si [1964Pop2] a = 364.06 to 364.21 at 10 at.%Zn and from 5 to 7 at.% Si [1964Pop2] at 30 at.% Zn, 1.5 at.% Si, 68.5 at.% Cu a = 368.47 [1964Pop2] (Si) < 1414
cF8 Fm3m C (diamond)
a = 543.06
pure Si, 25°C [Mas2]
(Zn) < 419.58
hP2 P63/mmc Mg
a = 266.50 c = 494.70
pure Zn, 25°C [Mas2]
, Cu10+7xZn10–10xSi3x
cI2 Im3m W
0<x<1 36.1 to 55.8 at.% Zn [2006Leb] 14.2 to 16.2 at.% Si [2002Leb]
CuZn(h) < 930 - 454
a = 259.39
at 47.5 at.% Zn [2006Leb]
~ Cu6Si 853 - 787
a = 285.4
at 14.9 at.% Si [2002Leb]
', CuZn(r) < 468
cP2 Pm3m CsCl
a = 295.9
at 49.5 at.% Zn [2006Leb]
, Cu5Zn8 < 835
cI52 I43m Cu5Zn8
a = 886.9
57 to 70 at.% Zn [2006Leb]
, CuZn3 700 - 560
hP3 CuZn3
MSIT®
a = 427.5 c = 259.0
72.45 to 76 at.% Zn [2006Leb] at 600°C, Cu0.7Zn2 [2006Leb]
Landolt-Börnstein New Series IV/11C3
Cu–Si–Zn
401
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
J, CuZn4 < 574
hP2 P63/mmc Mg
a = 274.18 c = 429.39
78 to 88 at.% Zn [2006Leb]
, Cu3Si(h2) 859 - 558
hR* or t** R3m
a = 247 = 109.74°
23.4 to 24.9 at.% Si [2002Leb]
a = 726.7 c = 789.2 ', Cu3Si(h1) 620 - 467
hR* R3
a = 472 = 95.72°
23.2 to 25.2 at.% Si [2002Leb]
'', CuSi3(r) < 570
o**
a = 7676 b = 700 c = 2194
23.3 to 24.9 at.% Si [2002Leb]
J1, Cu15Si4 < 800
cI76 I43d Cu15Si4
a = 961.5
at 21.2 at.% Si [2002Leb]
, Cu7Si 842 - 552
hP2 P63/mmc Mg
1, Cu5Si(h) 824 - 711
t**
1, Cu5Si(r) < 729
cP20 P4132 Mn
11.05 to 14.5 at.% Si at 12.75 at.% Si [2002Leb]
a = 256.05 c = 418.46
17.6 to 19.6 at.% Si annealing at 700°C [2002Leb]
a = 881.5 c = 790.3 a = 619.8
17.15 to 17.6 at.% Si [2002Leb]
Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Cu
Si
Zn
L + + 1
775
U1
L
66.2
14.5
19.3
L + (Si)
721
e4
L (Si)
37.2 39.4 0
8.1 2.7 100.0
54.7 57.9 0
L + + (Si)
683
E2
L (Si)
41.0 60.5 72.8 0
18.4 14.8 22.1 100.0
20.6 24.7 5.1 0
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Si–Zn
402 Reaction
T [°C]
Type
Phase
Composition (at.%) Cu
Si
Zn
L + + (Si)
~665
U6
L
(Si)
19.6 31.2 27.6 0
0.7 0.7 0 100.0
79.7 68.1 72.4 0
J1 + + '
600
U9
J1 '
78.7 76.3 68.3 76.7
20.4 22.4 15.2 22.0
0.9 1.3 16.5 1.3
L + J + (Si)
~574
U10
L
J (Si)
~11.5 24.0 22.0 0
0.7 0 0 100.0
87.8 76.0 78.0 0
' + + (Si)
530
E4
' (Si)
73.5 73.9 84.8 0
23.6 23.6 12.4 100.0
2.9 2.5 2.8 0
1 + ' '' + (Si)
480
U13
1 ' '' (Si)
57.6 73.3 73.6 0
11.1 23.9 23.6 100.0
31.3 2.8 2.8 0
L + J (Zn) + (Si)
424
U14
L J (Zn) (Si)
1.5 12.8 2.8 0
~0.7 ~0 ~0 100.0
97.8 87.2 97.2 0
MSIT®
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Cu-Zn
Cu-Si-Zn
A-B-C
Cu-Si
903 p1 l + (Cu) β
Si-Zn
853 p2 l + (Cu) β 842 p3 β + (Cu) κ
835 p4 l+βγ
824 p5 l + β δ1 820 e1 l η + δ1 802 e2 l η + (Si) 800 p6 η + δ1 ε1
β+γ+δ1
L+γ+δ1 775
787 e3 β δ1 + κ
U1
L + δ1 γ + η
U2
721 e4 l γ + (Si)
γ+η+ δ1 δ1 + η γ + ε1
?
δ1 + β γ + κ
?
U3
?
~700 683
665 p8 l+γδ
L γ + η + (Si) γ+η+(Si)
β+γ+δ1
δ1 + κ γ + γ1
δ1+ε1+γ
η+γ+ε1
U4
U5
δ1+γ+γ1 δ1 ε1 + γ + γ1
729 p7 δ1 + κ γ1
Cu–Si–Zn
L + β γ + δ1
775
711 e5 δ1 ε + γ1
E1
E2 ~665
L + γ δ + (Si)
U6
ε1+γ+γ1
γ+δ+(Si) L+δ+(Si)
Fig. 1a: Cu-Si-Zn. Reaction scheme, part 1
E3
U10
U7
403
MSIT®
U8 U9
E4
Cu-Si-Zn U6 E2 <600
p3
U4
Cu-Si
U5
γ + γ1 κ + ε1
625?
U8 ε1 + η γ + η'
600
L + δ ε + (Si)
U10
U7 620 p9 η + ε1 η'
U9
ε1+γ+η'
570 p11 η' + ε1 η"
L+ε+(Si) U14
κ + γ1 (Cu) + ε1
? ~548
δ γ + ε + (Si)
E3 (Cu)+κ+ ε1
η+γ+η' ηη' + γ + (Si)
U11
(Cu)+γ1+ε1
γ+ε+(Si)
530
552 e8 κ γ1 + (Cu)
Cu–Si–Zn
548 e6 δε+γ
558 e7 η (Si) + η'
γ+κ+ε1
γ1+κ+ε1
δ+ε+(Si)
Si-Zn
E1
β+(Cu)+γ
E5 ~574
U3
κ + β Cu + γ κ+Cu+γ
574 p10 l+δε
A-B-C
E4
?
U8
κ (Cu) + γ + ε1
E5
(Cu)+γ+ε1 500
γ + η' η'' + (Si)
480 424 p12 l + ε (Zn)
U10 ~424
Landolt-Börnstein New Series IV/11C3
L + ε (Zn) + (Si) ε+Si+(Zn)
Fig. 1b: Cu-Si-Zn. Reaction scheme, part 2
ε1 + η' γ + η'' η''+γ+η''
(Si)+γ+η'
γ+η''+(Si)
404
MSIT®
Cu-Zn
U12
ε1+γη'' U13 η'+η''+(Si)
U14 L+(Zn)+(Si)
467 e9 η' η'' + (Si) 419 e10 l (Zn) + ζ
Cu–Si–Zn
405
Si
Data / Grid: at.% Axes: at.%
Fig. 2: Cu-Si-Zn. Liquidus surface projection 20
80
40
60
(Si) 60
40
e2 e1 80 p5 p2
η
E2
δ U1
(Cu)
U2
e4
β
20
Cu
20
γ p1 40
p4 60 Cu Zn Si
Fig. 3: Cu-Si-Zn. Partial isothermal section at 847°C
U6 80 p 8
50.00 0.00 50.00
U14
U10 p10
p12
e10
Zn
Data / Grid: at.% Axes: at.%
60
40
70
30
80
20
L
90
10
β (Cu)
Cu
Landolt-Börnstein New Series IV/11C3
10
20
30
40
Cu Zn Si
50.00 50.00 0.00
MSIT®
Cu–Si–Zn
406
Cu Zn Si
Fig. 4: Cu-Si-Zn. Partial isothermal section at 760°C
50.00 0.00 50.00
Data / Grid: at.% Axes: at.%
60
40
70
30
L+η+(Si)
ε 1+η +γ η ε1 δ1+ε1+γ 80 δ1 δ 1+κ +γ δ1+κ +β
20
γ +η+L
L
κ
90
(Cu)+κ +β
10
γ
β
(Cu) 10
Cu
20
30
Cu Zn Si
Fig. 5: Cu-Si-Zn. Partial isothermal section at 700°C
50.00 0.00 50.00
50.00 50.00 0.00
Cu Zn Si
50.00 50.00 0.00
Axes: at.%
40
70
30
L+(Si)+η
ε1+η+γ η ε1 80
L
20
η+L+γ
κ +γ 1+γ 90
Cu Zn Si
Data / Grid: at.%
60
ε 1+γ 1+γ γ1
40
κ +γ +β κ
γ +β γ
10
(Cu)+κ +β
β
(Cu)
Cu
MSIT®
10
20
30
40
Landolt-Börnstein New Series IV/11C3
Cu–Si–Zn Cu Zn Si
Fig. 6: Cu-Si-Zn. Partial isothermal section at 600°C
407 50.00 0.00 50.00
Data / Grid: at.% Axes: at.%
60
40
70
30
η
η+γ +(Si)
ε1
80
ε 1+η +γ
γ1 κ
90
κ +γ 1+ε 1
κ +γ +ε1
20
κ +γ +β
(Cu)
(Cu)+κ +β
10
20
10
γ β
Cu
30
Cu Zn Si
Fig. 7: Cu-Si-Zn. Partial isothermal section at 482°C
50.00 0.00 50.00
γ1
90
Cu Zn Si
50.00 50.00 0.00
40
30
ε1+η''+γ
80
(Cu)+γ 1+ε1
50.00 50.00 0.00
Axes: at.%
70
ε1
Cu Zn Si
Data / Grid: at.%
60
η''
40
γ +η''+(Si) 20
(Cu)+γ +ε1 10
γ (Cu) (Cu)+γ +β
Cu
Landolt-Börnstein New Series IV/11C3
10
20
30
40
β
MSIT®
Cu–Si–Zn
408
Fig. 8: Cu-Si-Zn. Vertical section at 2 mass% Si, plotted in at.%
L 1000
Temperature, °C
L+(Cu) L+β
L+(Si) L+(Si)+γ 750
β
L+γ ~665 L+(Si)+δ
(Cu)
(Si)+γ
~548°C
γ
γ+δ
500
(Cu)+γ
(Si)+ε
L+(Si)+ε 424
Cu 0.51 Zn 94.96 4.53 Si
MSIT®
20
(Si)+γ+ε
40
60
(Cu)+γ+β
Cu, at.%
80
Cu 95.86 Zn 0.00 4.14 Si
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Ti
409
Copper – Tin – Titanium Chunlei Liu, Ulrich E. Klotz, Jörg F. Löffler Introduction Phase equilibria data concerning the Cu-Sn-Ti system are mainly based on the experimental work of [1999Kob, 2001Nak1]. The first extensive investigation of the Cu-Sn-Ti ternary system is concerned with the determination of the isothermal section at 397°C [1999Kob]. In that work 123 alloys were prepared by arc melting of commercial purity elements (Ti: 99.93 mass%; Cu: 99.96 mass%, Sn: 99.99 mass%) in argon atmosphere. The phase equilibria of annealed alloy samples were investigated by XRD and electrical resistivity measurements. [2001Nak1] determined the liquidus surface of the Cu-Sn-Ti ternary system. By arc melting under purified argon gas, Cu-Sn-Ti ingots were prepared from pure elements (Cu, Sn and Ti with 99.9 mass% purity), and all samples were annealed at 800°C for 168 h. The liquidus surface of the Cu-Sn-Ti ternary system was then established by determining the primary crystallizing phases as well as the phase reactions involving the liquid phase. A partial reaction scheme was also proposed by [2001Nak1]. More recent experimental works about the Cu-Sn-Ti ternary system are summarized in Table 1. Binary Systems The three binary systems are characterized by a series of peritectic reactions and by a number of stable and metastable intermetallic phases. Assessments of the Cu-Sn system by [2005Fra], of the Cu-Ti system by [2002Ans] and of the Sn-Ti system by [2005Liu], are accepted in the present work. They are based on previous assessment from [1996Shi, 2000Moo] for Cu-Sn, [1996Kum, 2002Can] for Cu-Ti and [1987Mur, 1998Kup] for Sn-Ti. Solid Phases The crystallographic data of the Cu-Sn-Ti phases and their temperature ranges of stability are listed in Table 2. All the Cu-Ti intermetallic phases are included in the review by [2002Ans] except the Ti3Cu phase because the literature data was only up to 1999 in his review. In accordance with [2002Can] the differential thermal analysis (DTA) of the Cu-80 at.% alloy, Ti3Cu is stable below 905°C. This result confirmed those reported by [1951Kla] which would lead to a peritectoid Ti2Cu + (Ti) Ti3Cu and an eutectoid reaction (Ti) (Ti) + Ti3Cu instead of an eutectoid reaction (Ti) (Ti) + Ti2Cu indicated in [2002Ans]. The Sn-Ti phase diagram [1987Mur] shows four intermetallic phases, Ti3Sn, Ti2Sn, Ti5Sn3 and Ti6Sn5. Recently, a new stable phase, Ti2Sn3, was reported by [1998Kup] and its structure was determined by XRD and Mössbauer spectroscopy [2000Kue, 2003Obr] Ti6–xCuxSn5 (0.0 < x < 0.1) was treated as a solid solution according to the experimental data of [1999Kob], which is an extension of the binary phase Ti6Sn5 into the ternary. Four ternary phases, namely Ti5CuSn3(-1), TiCuSn(-2), TiCu2Sn(-3) and Ti3CuSn(-4) have been identified in this system. TiCuSn(-2) was identified as a stoichiometric phase with LiGaGe structure type and P63mc space group [1999Kob], but [2000Wei] found that TiCuSn(-2) crystallizes with the space group P63/mmc and is isostructural to InNi2. TiCu2Sn(-3) was initially reported as having CsCl type structure by [1962Hei], which is the only ternary intermetallic phase included in the [V-C2]. However, [1967Hof] found that TiCu2Sn(-3) has L21-MnCu2Al type of structure, and this was confirmed by [1999Kob]. The peritectoid formation of the TiCu2Sn(-3) phase can be attributed to the difficulty to obtain this phase in the experimental work reported by [2001Nak1]. Recently, [2005Hua] reported a new ternary compound in the Cu-Sn-Ti system, Ti3CuSn(-4). The alloy, Cu-23Ti-17Sn (at.%), was prepared by mixing and milling the appropriate quantity of Cu powder, Sn powder and TiH2 powder. The samples were annealed at 900°C for 10 h. It was found that Ti3CuSn(-4) is thermodynamically stable at 900°C and its crystal structure is of the Ni3Sn2 type (space group of P63/mmc). The type of reactions by which the -3 and -4 phases form is unknown and suggested for future studies.
Landolt-Börnstein New Series IV/11C3
MSIT®
410
Cu–Sn–Ti
The most conflicting data about ternary intermetallic phases concern the Ti5CuSn3(-1) phase. Based on XRD [1999Kob] determined the change in unit cell volume of Ti5CuxSn3 as a function of the mole fraction of Cu. [1999Kob] concluded that Ti5CuxSn3 is an extended solid solution of Cu in Ti5Sn3 and it is isostructural to Mn5Si3. However, [2000Sch] reported that the Ti5CuSn3 phase is a ternary intermetallic phase, crystallizing with a space group of P63/mcm and the structure that can be described either as a filled D88-Mn5Si3 type or as a Ti5Ga4 type. At the same time, [2000Wei] confirmed that Ti5CuSn3 is a ternary stoichiometric compound by full profile Rietveld analysis of the X-ray powder diffractogram. [2001Ham, 2005Hua] also found that Ti5CuSn3 is related to the Ti5Ga4 type structure and that it has only a narrow homogeneity range. Therefore, Ti5CuSn3 is treated as a ternary stoichiometric compound in this assessment. Quasibinary Systems [1999Ham] proposed a quasibinary section between (Cu) and the Ti5CuSn3 compound in this system, which was confirmed by [2001Nak1]. Quasibinary eutectic involving (Cu) and Ti5CuSn3 is located at 878°C. The eutectic liquid composition is close to Cu-9Sn-15Ti (at.%). The diagram is given in Fig. 1 according to the experimental data of [1999Ham, 2001Nak1]. Invariant Equilibria The data concerning invariant equilibria are given in Table 3. They are mainly based on the experimental data of [2001Nak1]. A partial reaction scheme is given in Fig. 2, adapted from [2001Nak1]. Due to the limited experimental data, only the phase reactions involving the liquid phases are included. Ti2Sn3 was also included with appropriate modification to the eutectic reaction E2. Liquidus Surface The projection of the liquidus surface, Fig. 3, is taken from [2001Nak1] with modifications applied to the binary edges and to some of the invariant reactions according to the critically assessed binary systems. However, the liquidus surface (Fig. 3) and the reaction scheme (Fig. 2) proposed by [2001Nak1] do not show equilibrium of the -4 phase with the liquid phase. Such an equilibrium was proposed by [2005Hua] at 900°C, but the nature of the phase formation is yet unknown. This reveals an inconsistency which requires further experimental investigations. Due to the highly reactive nature of the element Ti, it is difficult to avoid oxidation and side reactions between sample and crucible, especially at high temperatures completely. To reduce possible reactions between the sample and the crucible [2001Nak1] employed differential scanning calorimetry with a heating rate of 10°C#min–1 when they recorded the phase transitions. Also care has been taken the measurements of the liquidus and/or solidus of Cu-Sn-Ti system may not be very accurate. Isothermal Sections The isothermal sections at 400 and 900°C are given in the Figs. 4 and 5, respectively. These were constructed on the basis of the experimental data obtained by [1999Kob, 2001Ham, 2005Hua] and amended to agree with the accepted binary systems and stability regimes of the ternary intermetallic phases. According to [2005Hua] -4 exists in equilibrium with the liquid phase, however this does not agree with the earlier investigations of the liquidus surface [2001Nak1], therefore the three-phase equilibrium Ti3CuSn+Ti5CuSn3+L shown by dashed lines in Fig. 5 seems to be questionable. The position of this three-phase region at 900°C was refined based on the data of [2005Hua] and the proposed quasibinary system (Cu)--1. Notes on Materials Properties and Applications Alloys of the Cu-Sn-Ti system belong to the group of high strength age hardening alloys [1961Saa, 1987Kiz]. In recent years, Cu-Sn-Ti alloys are widely used as brazing filler metals for diamond abrasive MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Ti
411
tools, because Cu-Sn-Ti alloys wet diamond well because of their high adhesion activity with regard to diamond; they produce strong joints of diamond tools, exhibit a low melting temperature, and have a satisfactory wear resistance [1975Kiz, 1999Ham, 2001Nak2, 2004Kha]. These are some reasons for the increasing interest in Cu-Sn-Ti alloys. References [1951Kla] [1961Saa]
[1962Hei]
[1967Hof] [1975Kiz]
[1987Kiz] [1987Mur]
[1994Ali]
[1994Sub]
[1996Kum]
[1996Shi]
[1998Kup]
[1999Ham]
[1999Kob]
[2000Kue]
[2000Moo]
Landolt-Börnstein New Series IV/11C3
Klarsson, N., “An X-ray Study of the Phases in the Copper Titianium System”, J. Inst. Met., 79(11), 341-405(1951) (Crys. Structure, Experimental, 15) Saarivirta, M.J., “Development of Copper Base High Strength Medium Conductivity Alloys Cu-Ti-Sn and Cu-Ti-Sn-Cr”, Trans. Metall. Soc. AIME, 221, 596-606(1961) (Morphology, Phase Relations, Experimental, 6) Heine, W., Zwicker, U., “Phases of the B-2 Type (CsCl) in Ternary Systems Containing Cu and Ni” (in German), Naturwissenschaften, 49(17), 391 (1962) (Crys. Structure, Experimental, 1) Hofer, G., Stadelmaier, H. H., “Co, Ni and Cu Phases of the Ternary MnCu2Al-Type” (in German), Monatsh. Chem., 98, 408-411 (1967) (Crys. Structure, Experimental, 9) Kizikov, E.D., Lavrinenko, I.A., “Investigations of Cu-Sn-Ti Alloys Used for Bonding Diamond Abrasive Tools”, Met. Sci. Heat Treat., 17(1-2), 61-65 (1975) (Morphology, Phase Relations, Experimental, 3) Kizikov, E.D., Kebko, V.P., “Microadditions to Alloys of the System Cu-Sn-Ti”, Met. Sci. Heat Treat., 29(1-2), 68-71 (1987) (Morphology, Phase Relations, Experimental, 8) Murray, J.L., “The Sn-Ti (Tin-Titanium) System” in “Phase Diagrams of Binary Titanium Alloys”, Murray, J.L., (Ed.), ASM Int., Materials Park, OH, 294-299 (1987) (Crys. Structure, Phase Diagram, Assessment, 22) Alisova, S.P., Lutskaya, N.V., Kobylkin, A.N., Budberg, P.B., “TiFe-Ti2Cu Section of the Ti-Fe-Cu System. Conditions of the Formation of Ti2Fe Compound”, Russ. Metall., 5, 121-123 (1994) (Experimental, Phase Diagram, Phase Relations, 8) Subramanian, P.R., “Cu (Copper)”, in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E., (Eds.), ASM International, Materials Park, OH, 1-3 (1994) (Crys. Structure, Thermodyn., Review, 16) Kumar, K.C.H., Ansara, I., Wollants, P., Delaey, L., “Thermodynamic Optimisation of the Cu-Ti System”, Z. Metallkd., 87(8), 666-672 (1996) (Crys. Structure, Phase Diagram, Assessment, 58) Shim, J.H., Oh, C.S., Lee, B.J., Lee, D.N., “Thermodynamic Assessment of the Cu-Sn System”, Z. Metallkd., 87(3), 205-212 (1996) (Crys. Structure, Phase Diagram, Assessment, 36) Kuper, C., Peng, W., Pisch, A., Goesmann, F., Schmid-Fetzer, R., “Phase Formation and Reaction Kinetics in the System Ti-Sn”, Z. Metallkd., 89(12), 855-862 (1998) (Phase Diagram, Experimental, Phase Relations, 12) Hamar-Thibault, S., Allibert, C.H., Tillmann, W., “Phase Constitution of Cu77Sn8Ti14Zr1 as a Binder for Diamond Tools”, EUROPM 99, Turin, Italy, 1999, 57-64 (1999) (Phase Relations, Experimental, 10) Koblyuk, N.O., Akselrud, L.G., Skolozdra, R.V., “Interaction Between the Components in the Ti-Cu-Sn System at 670 K”, Pol. J. Chem., 73, 1465-1471 (1999) (Crys. Structure, Phase Diagram, Experimental, Phase Relations, *, 15) Kuennen, B., Jeitschko, W., Kotzyba, G., Mosel, B.D., “Crystal Structure and Properties of the Titanium Stannide Ti2Sn3”, Z. Naturforsch., 55B, 425-430 (2000) (Crys. Structure, Experimental, 21) Moon, K.W., Boettinger, W.J., Kattner, U.R., Biancaniello, F.S., Handwerker, C.A., “Experimental and Thermodynamic Assessment of Sn-Ag-Cu Solder Alloys”, J. Electron. Mater., 29(10), 1122-1136 (2000) (Phase Diagram, Assessment, #, 24) MSIT®
412 [2000Sch]
[2000Wei]
[2001Ham]
[2001Nak1]
[2001Nak2]
[2002Ans]
[2002Can]
[2002Oka] [2002She]
[2003Obr]
[2004Kha]
[2005Hua]
[2005Fra]
[2005Liu]
MSIT®
Cu–Sn–Ti Schuster, J.C., Naka, M., Shibayanagi, T., “Crystal Structure of CuSn3Ti5 and Related Phases”, J. Alloys Compd., 305, L1-L3 (2000) (Crys. Structure, Experimental, Phase Relations, *, 5) Weitzer, F., Perring, L., Shibayanagi, T., Naka, M., Schuster, J.C., “Determination of the Crystal Structure of CuSnTi by Full Profile Rietveld Analysis”, Powder Diffr., 15(2), 91-93 (2000) (Crys. Structure, Experimental, *, 9) Hamar-Thibault, S., Allibert, C.H., “New Phases in the Ternary Cu-Ti-Sn System”, J. Alloys Compd., 317-318, 363-366 (2001) (Crys. Structure, Phase Relations, Experimental, *, 14) Naka, M., Schuster, J.C., Nakade, I., Urai, S., “Determination of the Liquidus of Ternary System Cu-Sn-Ti”, J. Phase Equilib., 22(3), 352-356 (2001) (Experimental, Phase Relations, *, 6) Naka, M., Nakade, I., Shibayanagi, T., Maeda, M., Schuster, J.C., Urai, S., “The Investigation of Cu-Sn-Ti System for Brazing of Ceramics”, Brazing, High Temp. Brazing Diff. Welding, Aachen, 2001, 170-172 (2001) (Phase Diagram, Experimental, Phase Relations, 12) Ansara I., Ivanchenko, V., “Cu-Ti (Cooper-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.11457.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, 26) Canale, P., Servant, C., “Thermodynamic Assessment of the Cu-Ti System Taking into Account the New Stable Phase CuTi3”, Z. Metallkd., 93(4), 273-276 (2002) (Crys. Structure, Phase Diagram, Assessment, 29) Okamoto, H., “Cu-Ti (Copper-Titanium)”, J. Phase Equilib., 26(3), 549-550 (2002) (Crys. Structure, Phase Diagram, Assessment, #, 6) Shelyapin, M.G., Koblyuk, N., Romaka, L., Stadnyk, Y., Bodak, O., Hlil, E.K., Wolfers, P., Fruchart, D., Tobol, J., “Electronic Structure of MCuSn Compounds with M (Ti, Zr and Hf)”, J. Alloys Compd., 347, 43-51 (2002) (Crys. Structure, Electronic Structure, Experimental, 10) Obrien, J.W., Dunlap, R.A., Dahn, J.R., “A Mossbauer Effect and X-ray Diffraction Investigation of Ti-Sn Intermetallic Compounds: I. Equilibrium Phases”, J. Alloys Compd., 353, 60-64 (2003) (Crys. Structure, Experimental, 14) Khalid, F.A., Klotz, U.E., Elsener, H.-R., Zigerlig, B., Gasser, P., “On the Interfacial Structure of Brazed Diamond Grits”, Scr. Mater., 50, 1139-1143 (2004) (Morphology, Experimental, 20) Huang, S.-F., Tsai, H.-L., Lin, S.-T., “Crystal Structure and X-ray Diffraction Pattern of CuSnTi3 Intermetallic Phase”, Intermetallics, 13, 87-92 (2005) (Crys. Structure, Experimental, *, 16) Franke, P. Neuschütz, D., SGTE, “Cu-Sn”, in Landolt-Börnstein. Numerical Data and Functional Relationships in Science and Technology (New Series). Martiensen, W. (Ed. in Chief), Group IV. Physical Chemistry, Springer-Verlag, Berlin, Heidelberg, Vol. 19-B3, 1-4 (2005) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Review, Assessment, #, 3) Liu, C., Klotz, U.E., Uggowitzer, P.J., Löffler, J.F., “Thermodynamic Assessment of the Sn-Ti System”, Chem. Monthly, 136(11), 1921-1930 (2005) (Phase Diagram, Assessment, #, 32)
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Ti
413
Table 1: Investigations of the Cu-Sn-Ti Phase Relations and Structures Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1975Kiz]
EPMA1,
0 to 30 mass% Sn, 0-20 mass% Ti
[1999Kob]
XRD, electrical resistivity
397°C, 123 samples in whole composition region
[1999Ham]
XRD, electron microscopy
600 to 950°C, quasibinary between (Cu) and Ti5CuSn3
[2000Sch]
XRD, EPMA
800°C, 10Cu-30Sn-60Ti (mass%), 11.1Cu-33.3Sn-66.7 Ti (mass%), and 15Cu-30Sn-55Ti (mass%), crystal structure of Ti5CuSn3
[2000Wei]
Full profile analysis of the XRD
45Cu-30Sn-25Ti (mass%), TiCuSn
XRD2
[2001Nak1] EPMA, XRD, DTA/DSC
800°C, 7.5 to 60 at.% Sn, 5 to 90 at.%Ti liquidus surface
[2001Nak2] EPMA, XRD, DTA/DSC3
Cu, 15 at.% Sn, 0 to 25 at.% Ti, liquidus and solidus
[2001Ham]
EPMA, XRD
900°C, compostion range: Cu-Ti3Sn and Cu-Ti6Sn5
[2002She]
Crystal structure, XRD
800°C, TiCuSn
[2005Hua]
EPMA, XRD, TEM4
900°C, Cu-17 at.%Sn-23 at.%Ti, Ti3CuSn
1
EPMA: Electron Probe Micro Analysis; XRD: X-ray diffraction; 3 DTA/DSC: Differential Thermal Analysis and Differential Scanning Calorimetry. 4 TEM: Transmission Electron Microscopy 2
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Cu) < 1084.62
cF4 Fm3m Cu
a = 361.46
at 25°C [Mas2] melting point [1994Sub]
(Sn) 231.9681 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
at 25°C [Mas2]
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
[Mas2]
(Ti) 1670 - 882
cI2 Im3m W
a = 330.65
[Mas2]
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Sn–Ti
414 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Ti) < 882
hP2 P63mmc Mg
a = 295.06 c = 468.35
at 25°C [Mas2]
(CuSn)
cI2 Im3m W
a = 302.6
[2005Fra, Mas2]
(CuSn) 757 - 520
cF16 Fm3m AlFe3
a = 611.7
Cu3Sn; [2005Fra, Mas2, V-C2]
Cu41Sn11 590 - 350
cF416 F43m Cu41Sn11
a = 1798.0
[2005Fra, Mas2, V-C2]
Cu10Sn3 640 - 582
hP26 P63 Cu10Sn3
a = 733.0 c = 786.4
[2005Fra, Mas2, V-C2]
Cu3Sn < 678
oC80 Cmcm Cu3Sn
a = 552.9 b = 4775.6 c = 432.3
[2005Fra, Mas2, V-C2]
Cu6Sn5 416 - 198
hP4 P63/mmc NiAs
a = 419.2 c = 503.7
[2005Fra, Mas2, V-C2]
’Cu6Sn5 < 198
hR22
a = 864.8 c = 1067.5
[2005Fra, V-C]
Ti2Cu < 1012
tI6 I4/mmm MoSi2
a = 295.3 c = 1073.4
[Mas2, V-C2, 1994Ali]
TiCu < 982
tP4 P4/nmm TiCu
a = 310.8 to 311.8 48 to 52 at.% Cu [Mas2, V-C2] c = 588.7 to 592.1
Ti3Cu4 < 925
tI14 I4/mmm Ti3Cu4
a = 313.0 c = 1994
[Mas2, V-C2]
Ti2Cu3 < 875
tP10 P4/nmm Ti2Cu3
a = 313 c = 1395
[Mas2, V-C2]
TiCu2 890 - 870
oC12 Amm2 VAu2
a = 436.3 b = 797.7 c = 447.8
[Mas2, V-C2]
TiCu4 885 - ~400
oP20 Pnma ZrAu4
a = 452.5 b = 434.1 c = 1295.3
~78 to ~80 at.% Cu [Mas2, V-C2]
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Ti
415
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
TiCu4 500
tI10 I4/m MoNi4
-
~78 to ~80 at.% Cu [Mas2]
Ti3Cu < 905
tP4 P4/mmm
a = 415.8 c = 359.4
[V-C2]
oP* Pbam
[2002Oka]
Ti3Sn < 1676
hP8 P63/mmc Ni3Sn
a = 591.6 c = 476.4
[Mas2], [2003Obr]
Ti2Sn < 1550
hP6 P63/mmc InNi2
a = 465.3 c = 570
[Mas2], [2003Obr]
Ti5Sn3 < 1506
hP16 P63/mcm Mn5Si3
a = 804.9 c = 545.4
[Mas2], [2003Obr]
Ti6Sn5(l) < 686
oI44 Immm Nb6Sn5
a = 1693 b = 914.4 c = 573.5
[Mas2], [2003Obr]
Ti2Sn3 < 767
oC* Cmca V2GaSn2
a = 596.74 b = 1995 c = 701.3
[2000Kue]
oC* Cmca ?
a = 596 b = 1994 c = 702
[2003Obr]
Ti6–xCuxSn5 1488 - 686
hP22 P63/mmc Ti6Sn5
a = 917.4 to 922.3 0.0x0.1 [1999Kob] c = 563.6 to 569.1
* -1, Ti5CuSn3
hP16 P63/mcm Mn5Si3 or Ti5Ga4
a = 816.2 c = 557.4
[2000Sch]
* -2, TiCuSn
hP6 P63mc LiGaGe or hP6 P63/mmc InNi2
a = 439.72 c = 601.68
[1999Kob]
a = 439.555 c = 601.505
[2000Wei]
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Sn–Ti
416 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
* -3, Ti1–xCu2+xSn
cP2 Pm3m CsCl or cF* Fm3m MnCu2Al
a = 592
hP6 P63/mmc Ni3Sn2
a = 463.6 c = 522.9
* -4, Ti3CuSn
[1962Hei]
a = 590.4 to 592.6 0.0x0.3 [1999Kob] [2005Hua]
Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Cu
Sn
Ti
L + Ti5Sn3 + Ti2Sn -1
> 1300
P1
L Ti5Sn3 Ti2Sn -1
~11.1 0.0 0.0 11.1
~33.3 37.5 33.3 33.3
~55.6 62.5 66.7 55.6
L + Ti2Sn -1 + Ti3Sn
?
U1
-
-
-
-
L + Ti5Sn3 -1 + Ti6Sn5
?
U2
-
-
-
-
L + (Ti) Ti3Sn +Ti2Cu
< 989
U3
-
-
-
-
L + Ti2Cu Ti3Sn + TiCu
959
U4
-
-
-
-
L + Ti3Sn -1 + TiCu
859
U5
-
-
-
-
L + -1 + Ti6Sn5 -2
827
P2
-
-
-
-
L + TiCu -1 + Ti3Cu4
?
U6
-
-
-
-
L + Ti3Cu4 -1 + TiCu2
?
U7
-
-
-
-
l (Cu) + -1
878
e6 (max)
l
76
9
15
L + (Cu) -1 + TiCu4
860
U8
-
-
-
-
L + -1 + Ti6Sn5 -2
827
P2
-
-
-
-
L -1 + TiCu2 + TiCu4
?
E1
-
-
-
-
L + (Cu) (CuSn) + -1
772
U9
L (Cu) (CuSn) -1
77.7 96.7 11.1
17.6 1.3 33.3
4.7 3.0 55.6
L + (CuSn) -1 + (CuSn)
748
U10
-
-
-
-
L + -1 (CuSn) + -2
732
U11
-
-
-
-
L + (CuSn) + -2 Cu3Sn
687
P3
-
-
-
-
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Ti T [°C]
Reaction
Type
417 Phase
Composition (at.%) Cu
Sn
Ti
L + -2 Cu3Sn + Ti6Sn5
572
U12
-
-
-
-
L + Cu3Sn Cu6Sn5 + Ti6Sn5
407
U13
-
-
-
-
L + Ti2Sn3 Ti6Sn5 + (Sn)
230
U14
-
-
-
-
E2
-
-
-
-
L Cu6Sn5 + Ti6Sn5 + (Sn) ~226 .
Fig. 1: Cu-Sn-Ti. Quasibinary system Cu-Ti5CuSn3(-1)
1500
L
Temperature, °C
1250
L+Ti5CuSn3
1000
878°C (Cu) 750
(Cu)+Ti5CuSn3
500
Ti 55.55 Cu 11.11 Sn 33.33
Landolt-Börnstein New Series IV/11C3
20
40
60
80
Cu
Cu, at.%
MSIT®
Cu-Ti
Cu-Sn-Ti
Sn-Ti
418
MSIT®
Cu-Sn
1585 e1 l (ETi) + Ti3Sn 1550 p1 l + Ti3Sn Ti2Sn 1506 p2 l + Ti2Sn Ti5Sn3
> 1300 L + Ti5Sn3 + Ti2Sn W1
1488 e2 l Ti5Sn3 + Ti6Sn5
P1
Ti5Sn3+Ti2Sn+τ1 L+Ti5Sn3 W1+Ti6Sn5
L + Ti2Sn Ti3Sn W1
U2
U1
Ti2Sn+Ti3Sn+τ 1 Ti5Sn3+Ti6Sn5+τ1
< 989
L +(ETi) Ti3Sn +Ti2Cu
960 e4 l TiCu +Ti2Cu
U3
(βTi)+Ti3Sn+Ti2Cu 959
L + Ti2Cu Ti3Sn + TiCu
U4
Ti2Cu+Ti3Sn+TiCu 859
925 p3 l + TiCu Ti3Cu4
U5
Ti3Sn+TiCu+τ 1
890 p4 l + Ti3Cu4 TiCu2
L + TiCu W1 + Ti3Cu4
U6
Ti3Cu4+TiCu+τ 1
885 p5 l + (Cu) TiCu4 Landolt-Börnstein New Series IV/11C3
P2 U8 Fig. 2a: Cu-Sn-Ti. Reaction scheme, part 1
L + Ti3Sn W1 + TiCu
L + Ti3Cu4 W1 + TiCu2
E1
U7
Ti3Cu4+TiCu2+τ1
Cu–Sn–Ti
1005 e3 l (ETi) + Ti2Cu
Landolt-Börnstein New Series IV/11C3
Cu-Sn
Cu-Ti
Cu-Sn-Ti U2
p5
878 e6(max) l (Cu) + W1
860
L + (Cu) W1 + TiCu4
875 e5 l TiCu4 + TiCu2
Sn-Ti U7
U8
(Cu)+TiCu4+τ1
827 797.3 p6 l + (Cu) E(CuSn) 756.6 p7 l + E(CuSn) J(CuSn)
L + W1 + Ti6Sn5 W2
P2
L W1 + TiCu2 + TiCu4
Ti6Sn5+τ1+τ2
L + (Cu) E(CuSn) + W1
772
TiCu2+TiCu4+τ1
U9
767 p8 l + Ti6Sn5 Ti2Sn3
(Cu)+β(CuSn)+τ1
748
L + E(CuSn) W1 + J(CuSn)
E1
U10
Cu–Sn–Ti
(CuSn)+γ(CuSn)+τ1
732
L + W1 J(CuSn) + W2
γ(CuSn)+τ1+τ2 638.5 p9 J(CuSn) Cu3Sn+l 572
687
U11
L + J(CuSn) + W2 Cu3Sn
L + W2 Cu3Sn+ Ti6Sn5 U12
415.8 p10 l+Cu3SnKCu6Sn5
P3
(CuSn)+Cu3Sn+τ2
Ti6Sn5+Cu3Sn+τ2
407
L+Cu3Sn KCu6Sn5+Ti6Sn5
U13
Cu3Sn+ηCu6Sn5+Ti6Sn5
231 e7 l Ti2Sn3 + (Sn)
~230 L + Ti2Sn3 Ti6Sn5 + (Sn) U14 226.9 e8 l KCu6Sn5 +(Sn)
~226 L KCu6Sn5 +Ti6Sn5 + (Sn)
E2
Ti2Sn3+Ti6Sn5+(Sn)
Fig. 2b: Cu-Sn-Ti. Reaction scheme, part 2
419
MSIT®
ηCu6Sn5+Ti6Sn5+(Sn)
Cu–Sn–Ti
420
Sn e7
Fig. 3: Cu-Sn-Ti. Liquidus surface projection
Ti2Sn3 p8
Data / Grid: at.%
U14 e8
Axes: at.%
ηCu6Sn5 E2 p10 U13
20
80
Cu3Sn 40
60
U12 Ti5Sn3 e2 Ti2Sn 60 p2 p1
80
Ti6Sn5 p9 P1
U2 P2
τ1
e1
(Ti)
U3 U4 20
Ti
e3
γ (CuSn)
P3
U1
Ti3Sn
40
τ2
40
U11 U10
U8
p4 e5 80 TiCu 60 p3 Ti3Cu4TiCu2TiCu4
e4 Ti2Cu
20
U9 e6(max)
U6 U7 E1
U5
p7
Sn
p6 (Cu)
p5
Cu
Data / Grid: at.% Axes: at.%
Fig. 4: Cu-Sn-Ti. Isothermal section at 400°C
L
20
80
40
60
Ti2Sn3 Ti6Sn5
ηCu6Sn5
Ti5Sn3 60 Ti2Sn
40
τ2 τ1
Ti3Sn
τ3
Cu3Sn Cu41Sn11
80
20
(α Ti)
Ti
MSIT®
β (CuSn)
(Cu) 20
Ti3Cu
Ti2Cu 40
TiCu Ti Cu 60 Ti2Cu3 3 4
80 TiCu
4
Cu
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Ti
421
Sn
Data / Grid: at.% Axes: at.%
Fig. 5: Cu-Sn-Ti. Isothermal section at 900°C L
20
80
40
60
Ti6Sn5 Ti5Sn3 60 Ti2Sn
40
τ1
Ti3Sn
τ4
80
?
20
(β Ti)
Ti
Landolt-Börnstein New Series IV/11C3
(Cu) 20
Ti3Cu Ti2Cu 40
TiCu
Ti3Cu4
60
80
Cu
MSIT®
422
Cu–Sn–Zn
Copper – Tin – Zinc Nataliya Bochvar, Evgeniya Lysova, Lazar Rokhlin, Volodymyr Ivanchenko Introduction Despite the fact that for the ternary Cu-Sn-Zn system plentiful experimental data are available, no complete phase diagram has been reported. One may mark the following milestones: the first systematic research of microstructures and mechanical properties of brasses alloyed by Sn with the structure (Cu) and (Cu) + [1905Gui]; the first measurements of mutual solubility between Sn and Zn in solid Cu [1905Gui, 1912Car, 1912Joh]; the determination of liquidus surface and limit of saturation of the ternary (Cu) phase by means of differential thermal analysis (DTA) and microscopic examinations of the ternary (Cu) phase up to 30 mass% Sn and 50 mass% Zn [1913Hoy1, 1913Hoy2, 1924Tam]. [1915Hud, 1915Hoy, 1919Hoy, 1920Cam, 1921Gui, 1930Bau1, 1930Bau2, 1931Bau, 1933Yam, 2004Vil] were successful in attempts to show the phase constitution under equilibrium conditions in wide temperature ranges within 50 to 70 mass% Cu and up to 7 mass% Sn region that includes practically all the useful Sn brasses. According to these researchers the solubilities of Sn in (Cu) and phases are functions of Zn content. High Sn contents are responsible for an eutectoid structure with the formation of a hard phase 1, as a consequence of reaction (Cu) + 1. [1915Hud] has established the isothermal section at 425°C, and [1920Cam] constructed the section basing on the alloys annealed at 650°C and then cooled slowly. However, these isothermal sections contradicted in some details with the phase rule. [1930Bau1, 1930Bau2] determined seven isothermal sections from 800 to 400°C and established the phase equilibria between (Cu), and 1 phases. [1933Yam] presented the isothermal sections at 650, 600 and 550°C, the results being in accordance with [1930Bau1, 1930Bau2]. In [2004Vil] the equilibrium phases and their compositions at 550, 450, and 350oC have been determined for Cu contents between 55.4 and 67.5 mass% and Sn contents up to 5.3 mass% by electron probe microanalysis (EPMA). The liquidus surface of the Cu rich part of the Cu-Sn-Zn system was investigated by [1933Yam]. In addition, several vertical sections in Cu rich part were constructed by [1913Hoy1, 1913Hoy2, 1924Tam, 1930Bau2, 1933Yam]. Phase equilibria in the Sn rich part of the system were only focused on by [1987Zhe]. [1997Pen] reported the liquidus temperatures for 40 alloys along four vertical sections with constant Cu:Sn ratios of 1:3, 1:2, 1:1 and 2:1 (in atomic fraction). The published liquidus temperatures were taken from the distinct break in the emf values versus temperature curves at the liquidus. The accuracy was estimated to be 7 K. [1937Kon] studied the homogeneity range of (Cu) solid solution at 500, 350 and 310°C by means of lattice parameters and compared the results with previous investigations. Beginning from early works to nowadays, the Cu-Sn-Zn alloys were melted in air using a gas-fired furnace with blast [1913Hoy1] to an induction furnace [1976Mac, 2004Vil]. Covered clay crucibles were used in the work due to [1920Cam]. As a guard against oxidation a cover of borax or charcoal gave very good results. [1926Han, 1930Bau1] reviewed some of the earlier investigations, whilst the later reviews of the system have been made by [1936Hum, 1949Jae, 1973Smi, 1979Cha, 1979Dri]. The thermodynamic descriptions of the Cu-Sn-Zn system have been presented by [1998Jan, 2002Mie]. The isothermal section at 250°C was calculated by [2003Jeo], and the liquid phase was modelled by [2001Her], whose assessment can reproduce the measured partial Gibbs energies of Zn and the liquidus temperatures published by [1997Pen]. Binary Systems The binary systems Cu-Sn, Sn-Zn and Cu-Zn are accepted from [1994Sau, Mas2], [2002Per] and [2006Leb], respectively.
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Zn
423
Solid Phases Table 1 summarizes the crystal structure data for the unary and binary phases. No ternary compounds were observed. The system shows a complete solid solution between isostructural binary phases of the Cu-Sn and Cu-Zn systems at high temperatures. The phase 1 dissolves up to 28 mass% Sn and the phase dissolves up to 9 mass% Zn at 600°C [1933Yam]. The phase shows the quite small homogeneity range and dissolves up to 5 mass% Zn at 550°C [1933Yam]. Invariant Equilibria Eight invariant four-phase reactions, five in the Cu rich part [1933Yam] and three in Sn rich one [1987Zhe], have been established experimentally. They are given in Table 2. The compositions of the equilibrium phases in Cu rich part of system are estimated from the established diagrams of [1933Yam]. According to [1933Yam], also three three-phase invariant equilibria exist in the Cu rich part of the Cu-Sn-Zn system. The reaction scheme in the Cu-Sn-Zn system is proposed in Fig. 1. In addition to the experimentally established invariant reactions, five speculative invariant four-phase reactions (U2, U4, U6, U8, E3) are also presented and shown in the diagram by dashed lines. The derived five invariant reactions correlate with experimental data [1933Yam, 1987Zhe] and also with isothermal section at 250°C [2003Jeo]. Liquidus Surface The liquidus surface in Cu rich part of the Cu-Sn-Zn system is shown in Fig. 2. The presented liquidus surface and liquidus isotherms are taken from the evaluation of [1979Cha], which is based on the investigations of [1913Hoy1, 1913Hoy2, 1924Tam, 1930Bau1, 1933Yam]. These investigations are in good agreement with each other. The liquidus surface and liquidus isotherms in Sn rich part of the system are according to [1987Zhe]. The liquidus isotherms in the central part of the Cu-Sn-Zn system are given by [1997Pen]. The calculated liquidus temperatures for the two vertical sections with ratio of Cu:Sn = 1:1 and 2:1 [2001Her] are in agreement with the experimental data from [1997Pen]. The calculated liquidus surfaces in the whole system presented by [1998Jan] agree well with the experimental data [1913Hoy1, 1924Tam, 1930Bau2] only in region of (Cu)-liquidus surface. The and J1 liquidus surfaces calculated by [1998Jan] differ significantly from the experimental data [1930Bau2, 1933Yam, 1987Zhe]. The calculation for the Cu rich part of the Cu-Sn-Zn system presented by [2002Mie] is in reasonable agreement with the experimental data critically reviewed by [1979Cha]. Isothermal Sections The partial isothermal section at 600°C shown in Fig. 3 is from the works [1933Yam, 1979Cha]. The partial isothermal sections at 450 and 350°C [2004Vil] are presented in Figs. 4 and 5, respectively. The compositions of the (Cu), ’ and 1 phases in the three-phase region given by [2004Vil] differ from those obtained by [1930Bau2, 1933Yam]. According to [2004Vil], the compositions of both (Cu) and 1 are almost constant at 450 and 350°C. Furthermore, [2004Vil] did not observe the three-phase region (Cu) + ’ + 1 at 550°C, whereas this region was reported by [1930Bau2] to extend up to 550°C. The results of [2004Vil] are preferable because [2004Vil] used a more accurate technique (EPMA) to determine the phase compositions. The isothermal section at 250°C calculated by [2003Jeo] is presented in Fig. 6. Figure 7 shows the homogeneity range of (Cu) at 500 and 300°C presented by [1979Cha], who used the data from [1937Kon, 1933Yam] as the base for the assessment. Temperature – Composition Sections Figure 8 shows the vertical section between the binary Cu-60Zn (mass%) to binary Cu-31Sn (mass%) alloys. It is drawn actually after [1933Yam] and agrees with the accepted binaries Cu-Sn and Cu-Zn. The temperature of the E2 reaction was shown in [1933Yam] as 511 instead of 528°C; this was corrected in Fig. 8. Meanwhile, intersection of four lines into one point on the Cu-Zn side is only an approximation. According to the accepted Cu-Zn phase diagram such a case could have been only at 59.8 mass% Zn as Landolt-Börnstein New Series IV/11C3
MSIT®
424
Cu–Sn–Zn
compared with 60 mass% Zn of the presented vertical section. As far as this difference between concentrations is very small, the indicated intersection of four lines into one point is quite justified. Figure 9 presents the part of the vertical section between binary Cu-60Zn (mass%) and binary Cu-33.6Sn (mass%) alloy within 0 to 20 mass% Zn [1933Yam]. Both sections were somewhat corrected to be consistent with the established isothermal sections. Thermodynamics The influence of small addition of Cu on the thermodynamic properties of a dilute solution of Zn in molten Sn has been studied using an electrochemical method in the temperature range of 450 to 650°C [1974Ngu]. The thermodynamic activities of Zn in Cu-Sn-Zn melts with compositions up to 6 at.% for both Cu and Zn between 430 and 600°C were measured by [1978Hag] employing emf technique. Using the torsion-effusion technique, the effect of small additions of Cu on the activity coefficient of Zn in dilute Sn-Zn melts at 530°C was investigated by [1974Tef]. Results of [1974Ngu, 1974Tef, 1978Hag] are very close to each other. And the Wagner interaction parameter JZnCu was estimated to be 2.901.20 – (3245900)/T [1974Ngu]. [1987Ryc] using the isopiestic method has determined the Wagner interaction parameters JZnSn in the dilute solution based on Cu at 1200°C. Following the same technique, [1986Sug] measured the activity of Zn in Cu-Sn-Zn melts at 1150 and 1100°C and in the composition range up to 8 at.% for both Zn and Sn. The activity of Zn in the ternary liquid phase has been measured at 700 [1996Kar] and 750°C [1997Pen] for four vertical sections with Cu:Sn ratio of 1:3, 1:2, 1:1 and 2:1. Both groups of authors utilized the emf method. From the measured emf values, the thermodynamics properties, such as activities of Zn, the partial Gibbs energy, partial enthalpy and partial entropy of the components, were derived. [1998Pen] performed a comparison between the calculated and measured thermodynamic properties reported in [1996Kar, 1997Pen]. The summaries for the work from [1998Pen] and [1998Jan] are given in [1997Din] and [1999Kau], respectively. Notes on Materials Properties and Applications [1930Bau2] measured the Brinell hardness of alloys containing 35 to 45 mass% Zn and up to 5 mass% Sn. A Cu based alloy containing 38.2 mass% Zn and 3.1 mass% Sn was subjected to severe plastic deformation at 400°C using the procedure of equal-channel angular pressing (ECAP) [2001Nei, 2004Nei]. The elongation achieved in this alloy amounted up to 900% at the slowest strain rate of 1.0 # 10–4 s–1 at 400°C. This elongation value was significantly higher than the value of 200% achieved in similar quasi-superplastic Cu-Sn-Zn alloys with about 2 mass% Pb in the temperature range of 500 to 600°C [1999Mat]. [2002Sho] investigated the interfacial reaction and mechanical properties of Sn-Zn alloy / Cu joints under the thermal exposure conditions. According to [2002Sho] the Vickers hardness decreased as a result of the transformation from Zn into Cu-Zn compounds. Miscellaneous The term “kalhoid” was used by early investigators [1913Hoy1, 1913Hoy2, 1919Hoy] to designate the ternary Cu-Sn-Zn alloys in order to avoid confusion with the names related to binary brass and bronze. Currently, this term is not in use. When Sn bearing brasses are compared with bronzes, it is shown that the chief difference from the view point of structure lies in the fact that in the brasses the phase remains stable down to room temperature, whereas in the bronzes it breaks down in the neighbourhood of 500°C into a mechanical mixture of (Cu) and or the eutectoid as has been proposed by [1969Rao]. [1956Fre] discussed the effects of composition and structure on properties of industrial brasses and bronzes used for making castings. Some problems associated with freezing brasses and bronzes have been discussed by [1971She]. Formations of interface phases under different conditions of interdiffusion have been examined by [1940Glu, 1976Urq, 1987Tak, 1998Lee]. The morphology and morphological stability of large precipitates formed in Cu-Sn-Zn alloys were investigated by [1967Mal] using EPMA technique.
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Zn
425
[1987Tak] determined the interdiffusion coefficients for (Cu) phase of the Cu-Sn-Zn system at 800°C by measuring the composition profiles of the solid-solid diffusion couples using the EPMA technique. [1997Lee, 1998Lee, 2000Kat, 2002Sho, 2003Jeo] studied the interface reactions between Cu substrate and Sn-Zn binary eutectic in lead-free solder alloys. Basing on the calculation of metastable equilibria among initial phases a new approach has been suggested for predicting the occurrence of the compound, which forms first on the substrate/solder interface during the soldering process. [1984Mur, 1995Fra, 1995Mor, 2002Miu] investigated martensite transformation in aged Cu-Sn-Zn shape memory alloys. For the Cu-Sn-Zn parent phase, two stages of transformation during ageing at 200°C were observed [1995Fra]. The stages were established to be connected with redistribution of Sn atoms. [2002Miu] found the shape memory effect below 268 K, the superelasticity above 278 K and about 4% maximum pseudoelastic strain in Cu-4.3 Sn-33.0 Zn (mass%) polycrystals. [1984Mur] evaluated the resistance against degradation due to repeated martensite transformation in Cu-24Zn-6Sn (mass%) alloy through optical reflectivity measurements. References [1905Gui] [1912Car] [1912Joh] [1913Hoy1] [1913Hoy2] [1915Hoy] [1915Hud] [1919Hoy] [1920Cam] [1921Gui] [1924Tam] [1926Han] [1930Bau1]
[1930Bau2]
[1931Bau]
[1933Yam]
Landolt-Börnstein New Series IV/11C3
Guillet, M.L., “The Examination of Brasses Alloyed with Bronzes” (in French), Rev. Metall., 2, 97-120 (1905) (Experimental, Phase Diagram) Carpenter, H.C.H., “The Effect of Other Metals on the Structure of the Constitutent in Copper-Zinc Alloys”, J. Inst. Met., 8, 59-73 (1912) (Experimental, Phase Diagram, 5) Johnson, F., “The Influence of Tin and Lead on Microstructure of Brass”, J. Inst. Met., 7, 201-217 (1912) (Experimental, Phase Diagram, 3) Hoyt, S.L., “On the Copper-Rich Kalhoids (Copper-Tin-Zinc Alloys)”, Engineering, 96, 667-706 (1913) (Phase Diagram, Experimental, *, 30) Hoyt, S.L., “On the Copper-rich Kalhoids (Copper-Tin-Zinc Alloys)”, J. Inst. Met., 10, 235-274 (1913) (Phase Diagram, Experimental, *, 30) Hoyt, S.L., Brinton, P.H.M.P., “Notes of the Copper-Rich Kalhoids”, J. Inst. Met., 10, 178-188 (1915) (Experimental, Phase Diagram, *, 2) Hudson, O.F., Jones, R.M., “The Constitution of Brasses Containing Small Percentage of Tin”, J. Inst. Met., 14(2), 98-115 (1915) (Phase Diagram, Experimental, 5) Hoyt, S.L., “Constitution of Tin Bronzes”, Trans. Am. Inst. Min. Metall. Eng., 60, 198-205 (1919) (Experimental, Phase Diagram, 3) Campbell, W., “A Note on the Constitution of Certain Tin-Bearing Brasses”, Proc. Amer. Soc. Test. Mater., 20(11), 104-114 (1920) (Experimental, Phase Diagram, *, 4) Guilet, M.L., “The Alloying of Brasses with Tin” (in French), Rev. Metall., 18, 445-458 (1921) (Experimental, Phase Diagram, 3) Tammann, G., Hansen, M., “On the Ternary System Cu-Sn-Zn” (in German), Z. Anorg. Allg. Chem., 138, 137-161 (1924) (Phase Diagram, Experimental, *, 22) Hansen, M., “The Constitution of Copper Alloys” (in German), Z. Metallkd., 18, 347-349 (1926) (Phase Diagram, Review, *, 16) Bauer, O., Hansen, M., “Effect of the Third Metal on the Structure of Brasse. III The Influence of Tin” (in German), Z. Metallkd., 22, 387-391 (1930) (Phase Diagram, Experimental, #, 18) Bauer, O., Hansen, M., “Effect of the Third Metal on the Structure of Brasse. III The Influence of Tin” (in German), Z. Metallkd., 22, 405-411 (1930) (Phase Diagram, Experimental, #, 0) Bauer, O., Hansen, M., “Effect of the Third Metal on the Structure of Brasse. III The Influence of Tin” (in German), Z. Metallkd., 23, 19-22 (1931) (Phase Diagram, Experimental, #, 0) Yamaguchi, K., Nakamura, I., “Investigation of the Phase Equilibria in Cu-based Cu-Sn-Zn Alloys” (in Japanese), Bull. Inst. Phys. Chem. Res., 12, 1330-1352 (1933), (Phase Diagram, Experimental, *, #, 16) MSIT®
426 [1936Hum] [1937Kon]
[1940Glu]
[1949Jae] [1956Fre]
[1967Mal]
[1969Rao] [1971She] [1973Smi] [1974Ngu]
[1974Tef]
[1976Mac] [1976Urq] [1978Hag]
[1979Cha]
[1979Dri]
[1984Mur]
[1986Sug]
[1987Ryc]
MSIT®
Cu–Sn–Zn Hume-Rothery, W., “On the Theory of Equilibrium in Alloys.-I”, Philos. Mag., 22, 1013 (1936) (Assessment, Crys. Structure, Experimental, Phase Diagram, 24) Konobeevski, S.T., Tarasova, V.P., Stepanova, A.A., “X-Ray Determination of Maximum Solubility of alpha-Phase in the Ternary System Copper-Zinc-Tin at Low Temperatures by the Method of X-Ray Analysis”, Acta Physicochimica URSS, 6, 799-814 (1937), translated from Zh. Fiz. Khim. 9, 693-703 (1937) (Experimental, Phase Diagram, *, #, 12) Gluskin, D.Ya., “The Nature of Phases Forming under Interaction of Refractory Metals and Alloys with Eutectic Solders” (in Russian), Zh. Tekh. Fiz., 10, 1486-1501 (1940) (Crys. Structure, Experimental, 15) Jaenecke, E., “Cu-Sn-Zn” (in German), in “Kurzgefasstes Handbuch aller Legierungen”, Winter, Heidelberg, 562-563 (1949) (Review, 5) French, A.R., Member, F.I.M., “The Metallurgical Control of Quality in the Production of Copper-Base Alloy Castings”, J. Inst. Met., 85, 293-317 (1956) (Elect. Prop., Mechan. Prop., Phys. Prop., Phase Diagram, Review, 124) Malcolm, J.A., Purdy, G.R., “The Morphology and Morphological Stability of Large Precipitates Formed in CuZn and CuZnSn”, Trans. Metall. Soc. AIME, 239, 1391-1399 (1967) (Crys. Structure, Experimental, 34) Rao, S.S., Anantharaman, T.R., “Constitution of Brasses Below 500°C”, Z. Metallkd., 60, 312-315 (1969) (Crys. Structure, Experimental, *, 16) Sheremetev, G.F., Lebedev, K.P., “Solidification of Sn Bronzes” (in Russian), Tr. Leningrad. Politechn. Inst., (319), 107-113 (1971) (Experimental, Mechan. Prop., 8) Smith, C., Levine, E.D., “Copper-Tin-Zinc”, in “Metals Handbook”, ASM Int., Metals Park, OH, 8, 431 (1973) (Review, 6) Nguyen-Duy, P., Rigaud, M., “Activity of Zinc in the Tin-Rich Corner of the Ternary Alloy of Zinc, Copper and Tin in the Temperature Range 723 to 923 K”, J. Chem. Thermodyn., 6(10), 999-1004 (1974) (Thermodyn., Experimental, 2) Tefelske, T., Chang, Y.A., “The Effects of Copper and Gold on the Activity Coefficient of Zinc in Dilute Liquid Zinc Tin Alloys at 803 K”, Mater. Sci. Eng., 14(3), 211-216 (1974) (Thermodyn., Experimental, 16) Mack, J., Yang, W.J., “The ' Eutectoid in the Copper-Zinc System”, Metall. Trans., 7A, 1798-1800 (1976) (Crys. Structure, Experimental, *, 8) Urquhart, A.W., “Interdiffusion Studies on Contact Plating Materials”, Electr-Contacts, Proc. Ann. Holm. Semin., 22, 185-189 (1976) (Experimental, Interface Phenomena, 7) Hagiwara, H., Sugino, S., Takahashi, T., “Interaction Parameters of Zinc with Silver, Copper and Gold in Liquid Tin”, (in Japanese), Denki Kagaku oyobi Kogyo Kadaku, 16(8), 455-462 (1978) (Thermodyn., Experimental, 20) Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Sn-Zn”, in “INCRA Monograph Series 6. Phase Diagrams and Thermodunamic Properties of Ternary Copper-Metal Systems”, NSRD, Washington, 6, 678-683 (1979) (Phase Diagram, Crys. Structure, Review, #, 11) Drits, M.E., Bochvar, N.R., Gusei, L.S., Lysova, E.V., Padezhnova, Rohlin, L.L., Turkina, N.I., “Cu-Sn-Zn”, in “Binary and Multicomponent Copper-Base Systems” (in Russian), Nauka, Moscow, 193-194 (1979) (Phase Diagram, Review, 4) Murakami, K., Murakami, Y., Mishima, K., Ikai, Y., “Degradation and Improvement of Shape Memory Effect of Cu-Base Alloy”, J. Jpn. Inst. Met., 48(2), 115-121 (1984) (Experimental, Mechan. Prop., 13) Sugino, S., Nakumura, S.., Hagiwara, H., “Effect of Tin on the Activity of Zinc in Molten Copper” (in Japanese), J. Jpn. Inst. Met., 50(5), 462-466 (1986) (Themodyn., Experimental, 10) Rychlewski, M., Pomianek, T., “Effect of Iron-Group Metals as well as Tin and Germanium on the Zinc Activity in Dilute Copper Solutions” (in Polish), Zesz. Nauk. Akad. Gorn.-Hutn. im. Stanislawa Stazhika, Metal. Odlew, (109), 55-61 (1987) (Thermodyn, Experimental, 19) Landolt-Börnstein New Series IV/11C3
Cu–Sn–Zn [1987Tak]
[1987Zhe]
[1994Sau]
[1995Fra]
[1995Mor] [1996Kar]
[1997Din] [1997Lee]
[1997Pen] [1998Jan]
[1998Lee]
[1998Pen]
[1999Kau] [1999Mat]
[2000Kat]
[2001Her] [2001Nei]
[2002Mie]
Landolt-Börnstein New Series IV/11C3
427
Takahashi, T., Kato, M., Minamino, Y., Yamane, T., “Interdiffusion in alpha-Solid Solutions of the Copper-Zinc-Tin System”, J. Mater. Sci., 22, 3194-3202 (1987) (Calculation, Experimental, Transport Phenomena, 43) Zeng, C., Wang, Z., Pan, F., Ying, L., Ye, Y., “An Investigation on the Liquidus of the Tin-Rich Corner of the Copper-Zinc-Tin System” (in Chinese), Acta Phys. Sin., 3(4), 435-439 (1987) (Phase Diagram, Experimental, #, 4) Saunders, N., Miodownik, A.P., “Cu-Sn (Copper-Tin)”, in “Phase Diagram of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM International, Materials Park, OH, 412-418 (1994) (Phase Diagram, Review, 42) Frackowiak, J.E., Plotkowski, K., Salamon, A., “Positron Annihilation Study of the Aging Effects in the Cu-Zn-Sn Parent Phase”, Acta Phys. Pol. A, 88(2), 327-332 (1995) (Experimental Phys. Prop., 6) Morgiel, J., Dutkiewicz, J., “Martensitic Transformation in Aged CuZnSn Shape Memory Alloys”, J. Physique (Paris), 5(C8), 991-996 (1995) (Experimental, Mechan. Prop., 11) Karlhuber, S., Peng, M., Mikula, A., “Thermodynamic Properties of Ternary Liquid Zinc-Alloys”, J. Non-Cryst. Solids, 205-207, 421-424 (1996) (Experimental, Thermodyn., 11) Dinsdale, A.T., “Summary of the Proceedings of the Calphad XXV Meeting”, Calphad, 21(1), 105-135 (1997) (Review, 154) Lee, B.-J., Hwang, N.M., Lee, H.M., “Prediction of Interface Reaction Products between Cu and Various Solder Alloys by Thermodynamic Calculation”, Acta Mater., 45(5), 1867-1874 (1997) (Calculation, Thermodyn., Phase Diagram, 26) Peng, M., Mikula, A., “Thermodynamic Properties of Liquid Cu-Sn-Zn Alloys”, J. Alloys Compd., 247, 185-189 (1997) (Thermodyn., Experimental, 20) Jantzen, T., Spencer, P.J., “Thermodynamic Assessment of the Cu-Pb-Zn and Cu-Sn-Zn Systems”, Calphad, 22(3), 417-434 (1998) (Phase Diagram, Thermodyn., Calculation, *, 33) Lee, H.M., Yoon, S.W., Lee, B.-J., “Thermodynamic Prediction of Interface Phases at Cu/Solder Joints”, J. Electron. Mater., 27(11), 1161-1166 (1998) (Experimental, Phase Diagram, Calculation, 25) Peng, M., Qiao, Z., Mikula, A., “Comparison between Calculated and Measured Thermodynamic Data of Liquid (Ag, Au, Cu)-Sn-Zn Alloys”, Calphad, 22(4), 459-468 (1998) (Thermodyn., Calculation, 19) Kaufman, L., Dinsdale, A.T., “Summary of the Proceedings of the Calphad XXVII Meeting”, Calphad, 23 (3-4), 265-303 (19997) (Review, Phase Diagram, Thermodyn., 155) Matsubara, R., Ashie, N., Nakamura, K., Miura, S., “Superplastic Forming of Corrosion Resistance CuZnSn Alloys”, Mater. Sci. Forum, 304-306, 753-758 (1999) (Experimental, Mechan. Prop., 4) Kattner, U.R., Eriksson, G., Hahn, I., Schmid-Fetzer, R., Sundman, B., Swamy, V., Kussmaul, A., Spencer, P.J., Anderson, T.J., Chart, T.G., Costa e Silva, A., Jansson, B., Lee, B.-J., Schalin, M., “Applications of Computational Thermodynamics: Goups 4 and 5: Use of Thermodynamic Software in Process Modeling and New Applications of Thermodynamic Calculations”, Calphad, 24(1), 55-94 (2000) (Review, Phase Diagram, Thermodyn., 201) Hertz, J., “What Will be Done in the Future with O.J.Kleppa’s Enthalpy Data Set?”, J. Alloys Compd., 321(2), 201-222 (2001) (Review, Phase Diagram, Thermodyn., 50) Neishi, K., Uchida, T., Yamauchi, A., Nakamura, K., Horita, Z., Langdon, T.G., “Low-Temperature Superplasticity in a Cu-Zn-Sn Alloy Processed by Server Plastic Deformation”, Mater. Sci. Eng. A., 307, 23-28 (2001) (Experimental, Mechan. Prop., 27) Miettinen, J., “Thermodynamic Description of the Cu-Al-Zn and Cu-Sn-Zn Systems in the Copper-Rich Corner”, Calphad, 26(1), 119-139 (2002) (Calculation, Phase Diagram, Thermodyn., 20) MSIT®
Cu–Sn–Zn
428 [2002Miu] [2002Per]
[2002Sho]
[2003Jeo]
[2004Nei]
[2004Vil]
[2006Leb]
Miura, S., “Shape-Memory and Supererelasticity in CuZnSn Alloy and Its Application”, Mater. Sci. Forum, 394-395, 399-402 (2002) (Experimental, Mechan. Prop., 6) Perrot, P., Batista S., Xing, X., “Sn-Zn (Tin-Zinc)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.11449.1.20, (2002) (Phase Diagram, Crys. Structure, Assessment, 9) Shohji, I., Nakamura, T., Mori, F., Fujiuchi, S., “Interface Reaction and Mechanical Properties of Lead-Free Sn-Zn Alloy/Cu Joints”, Mater. Trans., JIM, 43(8), 1797-1801 (2002) (Experimental, Mechan. Prop., 11) Jeong, S.W., Kim, J.H., Lee, H.M., “Thermodynamic Issues of Lead-Free Soldering in Electronic Packaging”, Mater. Sci. Forum, 426-432, 4081-4086 (2003) (Calculation, Phase Relations, Thermodyn., 17) Neishi, K., Uchida, T., Ashie, N., Yamauchi, A., Nakamura, K., Horita, Z., Langdon, T.G., “Superplasticity of a Cu-Zn-Sn Alloy Processed by Equal-Chanel Angular Pressing”, Mater. Sci. Forum, 447-448, 483-488 (2004) (Experimental, Mechan. Prop., 8) Vilarinho, C., Soares, D., Castro, F., “Contribution to the Knowledge of the Cu-Sn-Zn System for Compositions Close to Brass Alloys”, J. Alloys Compd., 379, 161-165 (2004) (Experimental, Phase Diagram, 18) Lebrun, N., “Cu-Zn (Copper-Zinc)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2006) (Phase Diagram, Crys. Structure, Assessment, 18)
Table 1: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Cu) < 1084.62
cF4 Fm3m Cu
a = 361.46
pure Cu [V-C2] dissolves up to 9.1 at.% Sn at 586°C [Mas2, 1994Sau] and up to 38.27 at.% Zn at 454°C [2006Leb]
a = 370.46
at 9.1 at.% Sn [1994Sau]
a = 369.612 0.014
at 35.84 at.% Zn [2006Leb]
(Sn) 231.9681 - 13
tI4 I41/amd Sn
a = 583.18 c = 318.18
pure Sn at 25°C [Mas2, 1994Sau], dissolves 0.6 at.% Sn at 198.5°C [2002Per]
(Sn) < 13
cF8 Fd3m C(diamond)
a = 648.92
[Mas2, 1994Sau]
(Zn) < 419.58
hP2 P63/mmc Mg
a = 266.50 c = 494.70
pure Zn at 25°C [Mas2], dissolves 0.0039 at.% Sn at 198.5°C [2002Per]
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Zn Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
, {(Cu17Sn3),(CuZn)} cI2 Im3m Cu17Sn3 W 798 - 586 CuZn 903 - 454 755 - 520
cF16 Fm3m BiF3
, Cu10Sn3 640 - 582
hP26 P63 Cu10Sn3
Lattice Parameters Comments/References [pm] 13.1 to 16.5 at.% Sn [Mas2, 1994Sau] a = 297.81 to 298.71
at 13.4 to 15.7 at.% Sn [1994Sau] 36.1 to 55.8 at.% Zn [2006Leb]
a = 295.39
at 47.5 at.% Zn [V-C2, 2006Leb]
a = 606.05 to 611.76
cF416 F43m Cu41Sn11
J, Cu3Sn < 676
oC80 Cmcm Cu3Sn
, Cu6Sn5(h) 415 - 186
hP4 P63/mmc AsNi
', Cu6Sn5(r) < 189
h**
15.5 to 27.5 at.% Sn [1994Sau] at 16.6 to 25 at.% Sn [1994Sau] 20.3 to 22.5 at.% Sn [1994Sau], the phase is superstructure based on , AgZn [V-C2, 1994Sau]
a = 733.0 c = 786.4
, Cu41Sn11 590 - ~350
429
a = 1798.0
a = 552.9 b = 477.5 c = 432.3 a = 419.0 c = 508.6
a = 2087.0 c = 2508.1
at 23.1 at.% Sn [1994Sau] 20 to 21 at.% Sn [1994Sau] at 20.5 at.% Sn [1994Sau] 24.5 to 25.9 at.% Sn [1994Sau] at 25 at.% Sn [1994Sau]
43.5 to 44.5 at.% Sn [1994Sau] at 45.45 at.% Sn [1994Sau] 45 at.% Sn superlattice based on AsNi-type structure[1994Sau] [1994Sau]
', CuZn (r) < 468
cP2 Pm3m CsCl
a = 295.9
at 49.5 at.% Zn [2006Leb]
1, Cu5Zn8 < 835
cI52 I43m Cu5Zn8
a = 886.9
57 to 70 at.% Zn [2006Leb] [2006Leb]
1, CuZn3 665 - 548
hP3 P6 CuZn3
a = 427.5 c = 259.0
J1, CuZn4 < 574
hP2 P63/mmc Mg
a = 274.18 c = 429.39
Landolt-Börnstein New Series IV/11C3
72.45 to 76 at.% Zn [2006Leb] [V-C2, 2006Leb] 78 to 88 at.% Zn [2006Leb] [2006Leb]
MSIT®
Cu–Sn–Zn
430 Table 2: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) Cu
Sn
Zn
L + 1 +
705
U1
L 1
~69 ~73 ~72.4 ~70.2
~19.6 ~14.6 ~16.9 ~16.9
~11.4 ~12.4 ~10.7 ~12.8
+ +J
588
U3
J1
>77.7 ~79.0 ~79.0 >75.2
~22.1 ~19.9 ~20.4 ~24.7
<0.2 ~1.1 ~0.6 <0.1
(Cu) + + 1
571
E1
(Cu) 1
~75.6 ~78.2 ~72.5 ~74.7
~11.7 ~3.0 ~16.0 ~15.3
~12.7 ~18.8 ~11.5 ~10.0
+ J + 1
550
U5
J1
1
~72.2 >75.2 ~75.0 ~72.6
~17.1 ~24.7 ~19.3 ~16.7
~10.7 <0.1 ~5.7 ~10.7
(Cu) + + 1
528
E2
(Cu)
1
~73.6 ~78.6 ~75.0 ~75.3
~15.3 ~3.0 ~19.3 ~15.3
~11.1 ~18.4 ~5.7 ~9.4
L + J J1 +
358
U7
L
10.0
85.0
5.0
L + J1 + (Sn)
218
U9
L
~0.3
93.5
~6.2
L (Zn) + (Sn) + J1
190
E4
L
0.3
84.5
15.2
MSIT®
Landolt-Börnstein New Series IV/11C3
Landolt-Börnstein New Series IV/11C3
Cu-Sn
Cu-Sn-Zn
A-B-C
Cu-Zn
798 p3 l + (Cu) β
903 p1 l + (Cu) β
755 p4 l+βγ
835 p2 l + β γ1
ζ+γδ+ε
588
582 e5 ζδ+ε
550
U3 γ+γ1+ε
γ + ε δ + γ1
δ+ε+γ1
L+γ1+ε
γ+γ1+δ
? E2
L + ε η + ε1
U7
U4
L+ε1+γ1
δ1+ε1+γ1
L + γ1 ε + ε1
U6
548 e6 δ1 γ1 + ε1 424 p9 l + ε1 (Zn)
γ1+ε1+ε
L+ε+ε1
358
574 p7 l + δ1 ε1
E1
L + δ1 ε1 + γ1
ca. 560
U5
γ (Cu) + δ + γ1
528
β (Cu) + γ + γ1
571
γ+δ+ε
538 e7(max) γ (Cu) + δ
592 e2(max) β (Cu) + γ1
β+γ+γ1
Cu–Sn–Zn
ζ+δ+ε
227 e10 l η + (Sn)
665 p5 l + γ1 δ1
U2
591 e3(max) β (Cu) + γ
586 e4 β (Cu) + γ
350 e9 δ (Cu) + ε
L + γ γ1 + ε
ca. 600
590 p8 γ+ζδ
415 p10 l+εη
U1
l+γ+γ1
640 p6 γ+εζ
520 e8 γ (Cu) + δ
l + β γ + γ1
705
640 e1 γε+l
Sn-Zn
(Cu)+δ+γ1
ε+ε1+η δ (Cu) + ε + γ1
?
E3
L+η+ε1
(Cu)+ε+γ1 218
L + η (Sn) + ε1 η+(Sn)+ε1
?
U9
(Cu) + γ1 β' + ε (Cu)+β'+ε
L+(Sn)+ε1
190
U8
γ1+β'+ε
L (Sn) + (Zn) + ε1
E4
198.5 e11 l (Zn) + (Sn)
Fig. 1: Cu-Sn-Zn. Reaction scheme
431
MSIT®
(Sn)+(Zn)+ε1
Cu–Sn–Zn
432
Sn
Data / Grid: at.%
e10
Fig. 2: Cu-Sn-Zn. Liquidus surface projection
Axes: at.%
U9 U7
p10
e11 E4
400
20
80
450 500
0 55
40
60
U to
60
6
600
e1 U2
40
650
p4
γ1
γ
80
p3
700 U1
β
950 1000 1050° C
750
800
U4
850
(Cu) 20
Cu
20
p1
40
p2 60 Cu Zn Sn
Fig. 3: Cu-Sn-Zn. Cu corner of the isothermal section at 600°C
80
50.00 0.00 50.00
p7
p5
p9
Zn
Data / Grid: at.% Axes: at.%
60
40
70
30
ε ζ
γ +ζ+ε
80
20
γ γ1 β +γ +γ 1
90
10
(Cu)
Cu
MSIT®
10
20
30
40
β
Cu Zn Sn
50.00 50.00 0.00
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Zn Cu Zn Sn
Fig. 4: Cu-Sn-Zn. Cu corner of the isothermal section at 450°C
433 50.00 20.00 30.00
Data / Grid: at.% Axes: at.%
60
20
70
10
γ1
(Cu)+β +γ 1
β
(Cu) Cu Zn Sn
30
80.00 20.00 0.00
40
Cu Zn Sn
Fig. 5: Cu-Sn-Zn. Cu corner of the isothermal section at 350°C
50.00 20.00 30.00
Cu Zn Sn
50.00 50.00 0.00
Cu Zn Sn
50.00 50.00 0.00
Data / Grid: at.% Axes: at.%
60
20
70
10
γ1
β (Cu)+β +γ 1
(Cu) Cu Zn Sn Landolt-Börnstein New Series IV/11C3
80.00 20.00 0.00
30
40
MSIT®
Cu–Sn–Zn
434
Sn Fig. 6: Cu-Sn-Zn. Calculated isothermal section at 250°C
Data / Grid: at.% Axes: at.%
(Sn)
20
80
40
60
η+(Sn)+γ 1 η 60
ε
40
ε+η+γ 1
(β Sn)+(Zn)+ε1
80
20
ε +γ 1+β ' (Zn)
Cu
20
(Cu)
40
Cu Zn Sn
Fig. 7: Cu-Sn-Zn. Solubility isotherms of Sn and Zn in (Cu) at 300 and 500°C
60
β'
(Cu)+ε+β '
γ1
80 ε 1 γ 1+(β Sn)+ε1
60.00 0.00 40.00
70
Zn
Data / Grid: at.% Axes: at.%
30
80
20
90
10
500°C 300 (Cu)
Cu
MSIT®
10
20
30
Cu Zn Sn
60.00 40.00 0.00
Landolt-Börnstein New Series IV/11C3
Cu–Sn–Zn
Temperature, °C
Fig. 8: Cu-Sn-Zn. Vertical section from 69Cu-31Sn to 40Cu-60Zn (mass%), plotted in at.%
435
900
L 800
L+β +γ1
L+γ L+β
L+β +γ
β +γ1
700
β +γ
705
γ
β +γ+γ1 (Cu)+β +γ1
γ1
γ+γ1
600
592
(Cu)+γ 571
γ+δ
(Cu)+γ+γ1
528
500
(Cu)+γ1
Cu 40.69 Zn 59.31 Sn 0.00
Temperature, °C
Fig. 9: Cu-Sn-Zn. Partial vertical section from 66.4Cu-33.6Sn to 40Cu-60Zn, up to 20% Zn (mass%), plotted in at.%
(Cu)+γ+δ
(Cu)+δ +γ1
Cu 80.61 Zn 0.00 Sn 19.39
10
Sn, at.%
800
L L+β 700
L+γ
L+β +γ L+β +γ1
705 L+γ+γ 1
γ γ
γ+ζ +ε γ+ζ
600
γ+ε
584
δ +ζ +ε
550
γ+ζ +δ δ
γ+δ +ε ε+γ1
ε+δ +γ1
500
δ +ε Cu 80.45 Zn 19.55 Sn 0.00
Landolt-Börnstein New Series IV/11C3
10
Sn, at.%
20
Cu 78.68 Zn 0.00 Sn 21.32
MSIT®
436
Cu–Ti–Zr
Copper – Titanium – Zirconium Tamara Velikanova, Mikhail Turchanin Introduction The investigation of the partial isothermal section at 750°C in the Ti rich corner of the system by [1960Bun, 1961Enc] was the first work undertaken on the Cu-Ti-Zr ternary system. The interest in the system was related to the development of a new class of Ti based alloys, which exhibit high strength at elevated temperatures. A number of experimental works [1983Dut, 1988Woy, 1990Che, 1992Kov] on the phase relations in the ternary Cu-Ti-Zr system were stimulated by the discovery of glass formation through rapid quenching of liquid alloys, over a wide concentration range. [1983Dut] reported preliminary results on the liquidus surface in the central portion of the phase diagram. [1988Woy] carried out a more detailed investigation of the system in the range 30 to 100 at.% Cu. A partial liquidus surface and isothermal section at 703°C were presented. The existence of a ternary compound reported previously by [1965Tes, 1969Tes] was confirmed in this work as well as in the investigation of the ZrCu-TiCu section by [1992Kov]. As-cast, annealed and rapidly quenched alloys along the Zr2Cu-Ti2Cu section were studied by [1990Che]. The temperature-composition section and the diagram of metastable crystallization of amorphous alloys at 33 at.% Cu were presented. The experimental works devoted to the study of phase relations in the Cu-Ti-Zr system are summarized in Table 1. A thermodynamic assessment of the system was made by [2003Arr]. The calculation results agreed well enough with the experimental data available, except for some regions in the ternary isothermal section at 703°C. The optimized and self-consistent thermodynamic description of the Cu-Ti-Zr system by [2003Arr] was used in the present assessment for the generation of the majority of phase diagrams. Some figures were corrected because of differences with respect to the accepted Cu-Ti binary system (see Binary Systems and other sections below). Binary Systems The Cu-Zr and Ti-Zr systems are accepted from the thermodynamic assessments of [1994Zen] and [1994Kum], respectively. These diagrams are generally consistent with [2006Sem, Mas2]. Reactions associated with the high-temperature Zr13Cu24 phase from [1986Kne] as well as the polymorphism of Zr2Cu described by [2006Sem] are not taken into account in this assessment as there is no consensus of opinion concerning the reliability of the information. The Cu-Ti diagram is accepted from [1996Kum], but with corrections to the Ti rich part. The congruent melting of Ti2Cu is accepted from [2002Ans], and the nature of the reaction between , and liquid is accepted as eutectic rather than peritectic, as given in [1996Kum]. In [1988Woy], the binary phase diagram of the Cu-Ti system was taken from [1983Mur], and for the Cu-Zr system from [1979Dri]. According to the latter work, the ZrCu phase remains stable at low temperatures. This peculiarity should be taken into account when considering the results of [1988Woy]. Solid Phases The crystallographic data of the solid phases of the Cu-Ti-Zr system and their temperature and concentration ranges of stability are listed in Table 2. [1961Enc] found that the solubility of Cu in (Ti) diminishes with the addition of Zr. In agreement with this work, a decrease in the solubility of copper in (Ti) with increasing Zr concentration, as well as a decrease in the solubility of copper in (Zr) with increasing Ti content is shown by the thermodynamic computation of [2003Arr]. The maximum content of copper in the solid solution in the ternary system is not more than 12 at.%. The extension of the phase in the ternary system in equilibrium with the and phases (enriched by Ti) is less than ~6 at.% Cu at 750°C. The solubility of Zr in the (Ti,Cu) solid solution is less than 0.01 at.%, as was shown by the calculation of [1961Enc].
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
437
The formation of a continuous solid solution ( phase) between the isostructural binary compounds, Zr2Cu and Ti2Cu, was established by [1990Che] for temperatures below 800°C. The lattice parameters of the solid solution increase linearly from Ti2Cu to Zr2Cu, while the c/a ratio decreases [1990Che], as shown in Fig. 1. Previously, the possibility of the dissolution of zirconium in Ti2Cu at 750°C was remarked by [1961Enc]. Analysis of the thermodynamic model for the phase presented by [2003Arr] predicts a miscibility gap with a critical point at 612°C. The existence of extremely narrow two-phase regions, Zr3Cu8 + TiCu4, Zr3Cu8 + Zr14Cu51, Zr3Cu8 + Zr7Cu10, Zr3Cu8 + Ti2Cu3 (in the Cu rich region) and Zr7Cu10 + Ti3Cu4, Ti3Cu4 + -1, TiCu + Ti3Cu4, ZrCu + -1 (in the central portion of the phase diagram) allows one to suppose rather limited homogeneity ranges for the binary phases in the ternary system, not exceeding 5 at.% [1988Woy]. The TiCu and Zr7Cu10 phases, however, dissolve up to nearly 10 at.% of the third component at 703°C [1988Woy]. A solubility of 4 at.% Zr in TiCu at 842°C was reported by [1992Kov]. The ternary -1 phase having the Laves type MgZn2 structure was first reported by [1965Tes]. Its composition, reported by [1965Tes, 1969Tes] as ZrTiCu4, was later revised by [1988Woy, 1992Kov] and an average stoichiometry of ZrTiCu2 was proposed. This phase was confirmed in the as-cast and annealed (at 703°C) alloys by [1988Woy]. At 703°C, the extension of the homogeneity range of -1 is 25 to 30 at.% Zr and 45 to 50 at.% Cu. According to [1992Kov], at 50 at.% Cu the homogeneity range is 21 to 37 at.% Zr at ~830°C and 22.5 to 29 at.% Zr at 630°C. The phase was established by [1988Woy, 1992Kov] to melt congruently. A melting temperature 867°C at 25 at.% Zr was found by [1992Kov]. From the thermodynamic calculation, the melting temperature is given as 883°C, which is in good agreement with experiment. This phase was observed as a metastable phase in as-cast and melt-spun alloys with compositions along the Zr2Cu-Ti2Cu section having Zr contents between 27 and 36 at.% Zr. It transformed into the phase under annealing at 700°C. Thorough experimental investigation of the solubility ranges of ternary -1 phase, the binary phases as well as their stability in the ternary system, is needed. It is worth noting that only the , and phases were treated as ternary solid solution phases in the thermodynamic assessment of [2003Arr]. Quasibinary Systems It would have been assumed that quasibinary systems exist in the ternary system because quasibinary eutectics formed between the congruently melting compounds have been reported. Quasibinary eutectic maxima were reported to exist by [1988Woy] on the congruent monovariant liquid lines of joint crystallization of the -1 phase and ZrCu, Zr2Cu as well as the TiCu phase. However, the maximum point on the monovariant line corresponding to the L -1 + TiCu eutectic equilibrium was not confirmed by calculation. At the same time, the ZrCu--1 section was shown to be quasibinary of eutectic type above the temperature of the ZrCu eutectoid decomposition, as shown in Fig. 2. Below this temperature, the three-phase Zr7Cu10 + -1 + field exists in the section. So the section ZrCu--1 can be considered as only partially quasibinary. The eutectic composition, 50Cu-9.5Ti-40.5Zr (at.%), and a eutectic temperature of 841°C calculated by [2003Arr] are in good agreement with the experimental solidus and liquidus points for alloys along the ZrCu-TiCu section in the composition range 30 to 50 at.% Zr, obtained by [1992Kov]. Additionally, both the -1 and TiCu phases form quasibinary eutectics with the highly stable congruently melting Zr14Cu51 compound. The calculated -1-Zr14Cu51 and TiCu-Zr14Cu51 sections are shown to be quasibinary systems (Figs. 3 and 4, respectively). For the -1-Zr14Cu51 system, the eutectic composition and temperature are 55.0Cu-20.7Ti-24.3Zr (at.%) and 879°C. For the TiCu-Zr14Cu51 system, they are 61.4Cu-30.4Ti-8.2Zr (at.%) and 885°C. The latter is only 3°C higher than the solidus temperature of the adjoining TiCu + Zr14Cu51 + Ti3Cu4 three-phase field. Invariant Equilibria Tables 3 and 4 list the invariant reactions for the ternary system, taken mostly from the thermodynamic calculation of [2003Arr]. For the thermodynamic assessment of the ternary system, datasets for the binary systems were taken from [1994Kum, 1994Zen, 1996Kum]. In the calculation, all of the solid phases, except for , and , were considered to be stoichiometric. The reaction scheme in Fig. 5 shows an incongruent Landolt-Börnstein New Series IV/11C3
MSIT®
438
Cu–Ti–Zr
three-phase invariant equilibrium l + at 1005°C, calculated by [1996Kum]. This is changed to a congruent reaction l + at the same temperature in order to agree with the accepted Cu-Ti binary phase diagram. Agreement was obtained between the calculated temperatures and the data reported by [1988Woy] for the central part of the ternary system. Two invariant equilibria in reaction scheme, E4 and E5, out of the three proposed by [1988Woy] were confirmed. The eutectic reactions involving the -1 phase were determined by [1988Woy] using the Transient Liquid Phase (TLP) bonding technique. In this procedure, two phases, A and B, differing in composition are brought into contact and assembled into a sandwich structure, such as A/B/A or B/A/B, at a temperature higher than the transformation temperature involving these phases. If a eutectic reaction between them takes place, inter-diffusional processes eventually lead to the formation of a liquid layer at the interfaces. Eutectic regions were identified in the quenched microstructure and the composition of the liquid at the eutectic point was determined using EMPA. Diffusion couples involving -1 and either Zr7Cu10, Zr2Cu, ZrCu or TiCu were used. The sandwich structures were held at temperatures of around 867°C for 30 min. From these experiments, the three ternary eutectic reactions involving the -1 phase were determined: L -1 + ZrCu + Zr2Cu (E4), L -1 + + Zr2Cu (E5) and L -1 + Ti2Cu + TiCu. The temperature of the L -1 + Ti2Cu + TiCu reaction was reported by [1992Kov] to be 842°C. This reaction was not confirmed by calculation. Instead, two invariant equilibria with compositions very close to those of the invariant liquid phases and their temperatures, U4 and E1, at 856 and 855°C, respectively, were found. Two degenerated invariant equilibria, D1 and D2, including the TiCu2 phase were obtained by recalculation in the current assessment, instead of the P and U types found by [2003Arr]. The degeneracy is caused neglecting the possible Zr solubility in the Cu-Ti binary phases participating in the equilibria. Additionally, a ternary eutectoid reaction ZrCu -1 + + Zr7Cu10 at 720°C (E6) and transition type reaction TiCu4 + Zr14Cu51 ZrCu5 + Ti2Cu3 (U9) between 703 and 750°C are proposed based on the preceding monovariant equilibria and comparison of the phase equilibria at 703 and 750°C. Liquidus and Solidus Surfaces The liquidus projection, given in Fig. 6a for the whole concentration range of the Cu-Ti-Zr system, as well as an enlarged portion in the Cu rich part of the diagram, Fig. 6b, were recalculated using the thermodynamic data set of [2003Arr]. In Figs. 6a, 6b the substitution of incongruent point p (accepted by [2003Arr]) by congruent e1 at 30 at.% Cu for l + equilibrium is shown. The results of the calculation of the monovariant and of the invariant transformation temperatures are in satisfactory agreement with the experimental partial liquidus projection for the central part of the system reported by [1988Woy], except for some details. For example, the proposed compositions of the temperature maxima of the quasibinary eutectics formed by -1 with the TiCu, ZrCu and phases cannot be accepted because of inconsistencies with geometric thermodynamic rules. According to calculation, the liquidus temperature of the phase decreases with increasing Cu content up to the point E5, which corresponds to the minimum liquidus temperature of the ternary system. The main feature of liquidus surface in the central portion of the ternary system is the existence of a primary crystallization surface for the -1 ternary phase with a melting temperature lower than the melting temperatures of the coexisting binary phases. Such a feature creates an overall liquidus depression in the central portion of the ternary system, (Figs. 8 to 11) which explains the wide range of glass formation. Between 925 and 875°C, the liquid phase field branches into two topologically non-connected regions, as shown in Fig. 9. One of them corresponds to the liquidus depression observed in the Zr2Cu-Ti2Cu composition line according to the results of [1990Che]. The elongated region of liquid parallel to the Cu-Ti edge coincides with location of the lowest liquidus in any of the copper containing binaries bounding the ternary system. The solidus projection, Fig. 7, is constructed based on the thermodynamic calculation of [2003Arr]. Five tie lines of maximum solidus temperatures exist in the system: TiCu + Zr14Cu51, -1 + ZrCu, -1 + Zr14Cu51, -1 + a) and -1 + b), where a) and b) are phase enriched with Zr or with Ti, respectively, which correspond to three-phase invariant equilibria of eutectic type. Four of the tie-lines show equilibria involving the ternary compound -1, and their temperatures, in some cases (-1 + Zr14Cu51 and -1 + b) are MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
439
noticeably higher than those of the adjoining three phase planes. The three-phase -1 + a) + alloys demonstrate the lowest solidus temperature in the ternary system. Isothermal planes TiCu4 + Ti2Cu3 + TiCu2 at 869°C and Ti2Cu3 + TiCu2 + Ti3Cu4 at 874°C are not exposed because of their degeneration to lines due to the assumption of zero content of Zr in the Cu-Ti binary phases. Isothermal Sections Calculated isothermal sections for temperatures from 703 to 925°C are given in Figs. 8 to 14. Isothermal sections for 703, 875 and 925°C, Figs. 14, 9 and 8, respectively, are taken from [2003Arr]. The others are calculated in the present assessment on the basis of the thermodynamic assessment of [2003Arr]. Calculated and experimental sections at 750°C are compared in Figs. 13a and 13b. The partial isothermal section at 750°C was investigated by [1960Bun, 1961Enc]. Twenty one ternary alloys containing up to 30 mass% Zr and up to 40 mass% Cu were prepared in a non-consumable arc furnace under argon. The alloys were homogenized for 48 h at 825°C prior to annealing at 750°C for 5 d and then additionally for 7 d in order to check that the alloys had reached equilibrium. Specimens were studied by metallographic examination and X-ray diffraction. The existence of the three-phase + + Ti2Cu field was established in the Ti rich corner of the section, Fig. 13b. Satisfactory agreement between the calculated and experimental data can be seen. The isothermal section at 703°C and 50 to 100 at.% Cu was studied by [1988Woy]. The alloys were annealed at 703°C for 14 d and quenched into iced brine. The samples were analyzed using optical microscopy and EMPA. The phase relations were determined by the fitting of the corresponding three-phase triangles because the two-phase regions were very narrow. The authors report that the extent of phase penetration into the ternary system from the respective binary phases is rather limited and generally, does not exceed 5 at.%. However, TiCu and Zr7Cu10 are exceptions, in that they remain stable in the ternary system up to nearly 10 at.% of the third component. It is worth noting that, besides these two binary phases, homogeneity ranges can be suggested only indirectly for the Zr3Cu8, ZrCu, and Ti3Cu4 binary phases. The other binary intermediate phases have no extension into the ternary system. One of the dominant features of the 703°C section is the presence of the ternary -1 phase, which coexists in equilibria with most of the stable phases at this temperature. The -1 + + and -1 + + three-phase fields reported by [1988Woy] contradict the experimental results of [1961Enc] and the calculation of [2003Arr]. The results of [1988Woy] suggest that equilibrium was not achieved in their samples (taking into account that -1 + + phase field exists at higher temperatures, as one can see, for example, for 827 and 810°C in Figs. 11, 12, respectively). There is no agreement between the experimental data of [1988Woy] and the calculations of [2003Arr] in the Cu rich corner of the phase diagram, in which the existence of Zr7Cu10 + Ti2Cu3 and Zr7Cu10 + Ti3Cu4 fields were reported by [1988Woy] at 703°C instead of the alternative TiCu + Zr14Cu51, Zr14Cu51 + -1 or Zr3Cu8 + -1 obtained by the calculation. The isothermal ternary section for 703°C as calculated by [2003Arr] is shown in Fig. 14. The calculated isothermal sections at 750 and 703°C differ in the compositions of the , and phases in the respective three-phase equilibria, and by the existence of Zr14Cu51 + TiCu4 equilibria instead of ZrCu5 + Ti3Cu2. Temperature – Composition Sections Two vertical sections, Zr2Cu-Ti2Cu and ZrCu-TiCu, are given in Figs. 15 and 16 as results of calculations performed in the current assessment. The Zr2Cu-Ti2Cu vertical section given in Fig. 15 differs from the corresponding one by [2003Arr] near the Zr2Cu and Ti2Cu ordinates. The phase equilibria in the Zr rich portion of the section presented by [2003Arr] which shows a transgression of the phase rule is corrected. The Ti rich part the diagram was corrected taking into account the congruent melting of Ti2Cu, and the topology of the Zr rich part was similar to that in the present calculations. This part is shown by a dashed line. The Zr2Cu-Ti2Cu and ZrCu-TiCu sections were studied experimentally by [1990Che] and [1992Kov], respectively. [1990Che] studied thirteen alloys along Zr2Cu-Ti2Cu section, both in the as-cast and annealed (at 700°C) conditions. It was established that the mutual substitution of Ti and Zr in Ti2Cu and Zr2Cu leads to a decrease in the liquidus temperature down to 830°C at 36.7 at.% Zr. The precipitation of a third phase was observed in as-cast alloys at Zr contents of 27 to 36 at.% Zr, which was identified as a MgZn2 type Landolt-Börnstein New Series IV/11C3
MSIT®
440
Cu–Ti–Zr
Laves phase. A continuous solid solution was found below 830°C. Generally, good agreement was achieved between the experimental liquidus and solidus points determined by [1990Che] and the calculated solidus and liquidus lines. The composition-temperature range in which the -1 ternary phase was observed by [1990Che] corresponds to that where the -1 phase coexists with other phases of the ternary system below the solidus over some temperature intervals according to the calculation of [2003Arr]. The minimum solidus temperature after [2003Arr] is about 805°C. The Zr2Cu-Ti2Cu temperature composition section calculated by [2003Arr], Fig. 15, is in acceptable agreement with the experimental data of [1990Che]. The cut of the solidus isotherms of the U8 and E5 invariant reactions are seen at the middle of the section. The presence of the phase together with solid solution in the calculated section is connected with the fact that the composition of the Zr2Cu and Ti2Cu compounds are not strictly stoichiometric. It is important to note that the calculated phase fraction in the + phase field is very small, from 10–3 to 10–9. The experimental detection of this phase would have been too difficult owing to the small amount. [1992Kov] investigated twenty one as-cast and annealed (at 700°C) alloys of the ZrCu-TiCu section and the corresponding temperature-composition section was reported. Satisfactory agreement is observed between the experimental liquidus points detected by [1992Kov] and calculated liquidus lines in the vertical section. According to [1992Kov], the -1 phase (defined as a 1 phase) forms congruently at 867°C and 25 at.% Ti and has a homogeneity range along the ZrCu-TiCu section, which extends between 30 at.% Ti (at 842°C) and 12.5 at.% Ti (at ~820°C). [1992Kov] reports a limiting solubility of 4 at.% Zr in TiCu at solidus. Unfortunately, the interpretation of the phase relations of the section by [1992Kov] is in contradiction with the experimental data of [1988Woy] as well as with the calculation of [2003Arr]. Figure 16 shows the ZrCu-TiCu temperature-composition section across the whole concentration range calculated according to the thermodynamic assessment of [2003Arr]. Enlarged parts of the calculated ZrCu-TiCu section are shown in Figs. 17 and 18. Thermodynamics Experimental information on the thermodynamic properties of the ternary phases is not available. A thermodynamic assessment of the Cu-Ti-Zr system was carried out by [2003Arr] using the Calphad approach. Binary interaction parameters for the Cu-Ti, Cu-Zr and Ti-Zr systems were taken from previous work [1996Kum, 1994Zen, 1994Kum]. The ternary parameters for the liquid and phases and the thermodynamic description for the - and ternary -1 phases were determined from the experimental data on phase relations after [1988Woy, 1990Che]. The ternary parameters for the solution phases (liquid and ) take into account their mixing enthalpy only. The Gibbs free energy of the phase was described in the framework of the two sublattice model. The ternary -1 phase and all of the binary phases (except , TiCu and TiCu4) were represented by [2003Arr] as stoichiometric compounds. For the TiCu and TiCu4 phases, the homogeneity ranges in the binary systems only were taken into account in the calculation. Good agreement was obtained between the calculated temperatures for the invariant reactions in the central part of the system with the data reported by [1988Woy] based on the Transient Liquid Phase bonding technique. However, some discrepancies between the calculated and experimental ternary isothermal sections at 703°C after [1988Woy] are observed: continuous solid solution according to calculation (in agreement with [1990Che]) instead of -1 + + or -1 + + phase fields after [1988Woy]; phase equilibria -1 + Zr3Cu8 instead of Ti3Cu4 + Zr7Cu10 or Ti2Cu3 + Zr7Cu10 after [1988Woy] and others. It is worth noting that the three-phase equilibria in the composition region 0 to 50 at.% Cu assumed by [1988Woy] do not correspond to the equilibrium conditions at 703°C if taking into account the results of [1990Che]. More experimental information on the area enriched by Cu is needed. It also appears to be necessary to assess the reliability of the experimental data at 703°C for the ternary system at high Cu contents. Experimental investigations of the thermodynamic properties of the ternary -1 phase, phase and liquid alloys are required. Notes on Materials Properties and Applications The Cu-Ti-Zr system is well known as a basic system for the preparation of ternary and multicomponent amorphous alloys by rapid quenching from the liquid. Figure 19 shows the experimentally determined glass forming range in Cu-Ti-Zr alloys compiled from different literature sources. It can be seen that glass MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
441
forming range occurs in the central portion of composition triangle. Preparation of amorphous alloys by different techniques is considered in [1988Mas, 1990Che, 1992Kov, 1999Lee, 2001Ino, 2002Kas, 2002Lou, 2004Par]. In [1988Mas, 1990Che, 1992Kov, 2001Ino, 2002Kas, 2002Lou, 2004Par], thin ribbon alloy samples were prepared by a melt spinning technique. A pulsed laser quenching technique was used by [1988Mas]. The amorphous Cu90–xZrxTi10 (x = 20-80), Cu80–xZrxTi20 (x = 30-50), Cu90–xZrxTi10 (x = 30-40) and Cu40Zr20Ti40 alloys were successfully synthesized by mechanical alloying of mixtures of crystalline Cu, Ti, and Zr powders using a high energy ball milling technique [1999Lee]. These ranges are similar to those for amorphous alloys produced by melt spinning or argon atomization techniques. Ternary bulk amorphous alloys were prepared from the melt by an injection casting technique [1992Kov, 2004Par]. [2001Ino] reported that Cu based bulk glassy alloys of the composition Zr30Ti10Cu60 with a rod diameter of 4 mm can be formed by copper mold casting. The alloys had a fracture strength of 2050 to 2160 MPa and plastic elongation of 0.8 to 1.7%. However, only X-ray diffraction was used to characterize the samples. [2002Lou] further studied the microstructure of samples annealed at various stages using TEM. The formation of nanocrystals in the annealed samples was observed. An electron microscopy study by [2002Kas] on both as-cast and as-spun Zr30Ti10Cu60 samples clearly demonstrated that as-prepared samples contain a significant volume fraction (about 5-10%) of nanocrystals with diameters ranging from 5 to 15 nm. Thus, the as-prepared ternary samples can by classified as a nanocomposite: nanocrystals embedded in an amorphous matrix. In [2003Jia], microscopic structures and heat release under heating of as-prepared and annealed Zr30–xTi10+xCu60 alloys (x = 0-10) were studied by TEM, X-ray diffraction and DSC measurements. The first crystalline phase formed during constant rate heating (0.33 K#s–1 under flow of purified argon) is a Zr14Cu51-like phase with nanometer-sized grains. The first exothermic peak found on DSC curves corresponds to the amorphous-to-nanocrystalline Zr14Cu51-like phase transition. The second crystalline phase is also hexagonal, space group P63/mmc, and is related to the second exothermic peak found in the DSC curves. Amorphous alloys of the Cu-Ti-Zr system are widely used as filler materials for the brazing of titanium and titanium alloys [1994Bot, 1998Hir, 2001Kun] as well as for the brazing of ceramic-metal composite materials [1992Nak]. Couples of Pd-Ti/Pd-Ti (ASTM grade 7) and 6Al-Ti-4V/6Al-Ti-4V alloys (mass%) were brazed in a vacuum furnace using rapidly solidified amorphous Zr25Ti25Cu50 brazing foil [1994Bot]. The tensile strength, fatigue resistance and microstructure of the joints were studied by X-ray diffraction analysis, energy-dispersive spectroscopy and scanning microscopy. The tensile strength of the joint material is close to that of base metal. The fatigue properties of Pd-Ti joints do not differ from those of this base metal. The microstructure and mechanical properties of the brazed joints depend on the brazing cycle conditions: a fine lamellar eutectic joint microstructure consisting of (Ti) and the phase is observed after brazing a 6Al-Ti-4V alloy at 900°C for 10 min followed by rapid cooling. This brazing operation results in high strength joints. Brazing at temperatures higher than 900°C and/or with a relatively low cooling rate results in a coarse dendritic microstructure consisting of the phase and hexagonal -1 phase. Fine precipitates of the -1 phase in the phase matrix were also observed in the transition area between the base metal and the joint. Joints with such microstructures are brittle and have a low strength. It was shown that rapid cooling suppresses formation of the brittle -1 phase, thus resulting in high mechanical properties of the brazed joint. Commercial purity Ti and Zr90Ti10 alloy were brazed with amorphous 25Ti-25Zr-50Cu filler metal using Ar gas shielding [1998Hir]. To obtain a joint strength of 400 MPa, the Ti brazes required a brazing temperature of 920°C and a holding time of 120 s. Under these conditions, fracture occurred in the base metal. The Zr90Ti10 joint brazed at 880°C and 900°C had a joint strength of approximately 400 MPa for a holding time of 300 ks and were fractured in the base metal after 600 s. This difference in the brazing properties was explained by the microstructural evolution of the brazing filler metals during brazing on the basis of the liquidus projection of the Cu-Ti-Zr ternary system. The investigations were carried out using SEM, X-ray diffractometry and TEM. The reaction mechanism between Ti and Ni under brazing using an amorphous Cu-Ti-Zr alloy filler metal was investigated by [2000Oht]. The dissolution of the base metal in the liquid filler metal was completed rapidly: in about 60 s at 852°C and 877°C, and for 120 s at 927°C. The dissolution of the titanium base metal is three times that of nickel base metal. The maximum width of the bonded interlayer was related to the Landolt-Börnstein New Series IV/11C3
MSIT®
442
Cu–Ti–Zr
bonding temperature. The Cu-Ti and Ni-Ti intermetallic compound layers formed at the interface of the titanium base metal and the bonded interlayer. A Ni-Ti shape memory alloy was brazed to commercial pure Ti with Cu-Ti-Zr amorphous filler [2001Kun]. The brazing temperature was higher than 920°C. The effect of brazing condition on reaction layer formation as well as on the joint strength was investigated. Cu-Ti-Zr amorphous foil was used as insert metal to bond rod shaped specimens with diameters between 10 and 50 mm to alumina ceramics and stainless steel by a diffusion bonding method. Bonding of the materials was possible at 860-880°C. [1990Che] studied the properties of amorphous alloys along the Zr2Cu-Ti2Cu section. Thirteen alloys were obtained in a glassy state across the whole composition range along the section. The variation with composition of electroresistance ', microhardness H, embrittlement temperature Tb, and critical thickness dc of the amorphous alloys along the Zr2Cu-Ti2Cu section is shown in Fig. 20. Amorphous alloys with 18-44 at.% Zr have the maximum values of dc = 200-250 m. The concentration dependence of critical thickness of the amorphous alloys along the ZrCu-TiCu section was investigated by [1992Kov]. The results of the study are shown in Fig. 21. The break in the dc dependence corresponds to the composition of the bulk amorphous alloy Zr40Ti10Cu50. The influence of wheel speed on the glass forming ability and thermal behavior of Zr20Ti20Cu60 amorphous ribbons was studied by [2004Rev]. The DSC measurements revealed that the higher the wheel speed, the higher the glass transition temperature Tg and crystallization temperature. Independent of wheel speed, crystallization takes place in a two-stage process. The microhardness exhibits linear dependence on the crystallized volume fraction. A perfect solute mixtures model of defect-free nanoparticles embedded in an amorphous matrix was used to account for this strengthening mechanism. The crystallization behavior, the thermal stability, and the microhardness of ZrxTi40–xCu60 amorphous alloys (x = 15, 20, 22, 25, 30) prepared by melt-spinning have been studied by means of differential scanning calorimetry, DTA, X-ray diffraction and microhardness studies [2004Con1]. The addition of Zr does not affect the microhardness of the amorphous alloys, although they exhibited good mechanical properties. The effect of continuous heating and isothermal heat treatment of ductile Zr20Ti20Cu60 amorphous ribbons was investigated by [2004Con2]. Upon continuous heating, the alloy exhibited a glass transition, followed by a change to a supercooled liquid region and two exothermic crystallization stages. The decomposition of the amorphous phase was also observed. The first crystallization stage resulted in the formation of a nanocomposite structure with hexagonal Cu51Zr14 particles embedded in the amorphous matrix, while in the second crystallization stage, a hexagonal Cu2TiZr-like phase precipitated. The released enthalpies were 19 J#g–1 and 30 J#g–1 for both of the crystallization stages. The experimental studies of materials properties of the Cu-Ti-Zr system are summarized in Table 5. Miscellaneous The diagram of metastable crystallization of the amorphous alloys of Zr2Cu-Ti2Cu section, Fig. 22, was constructed by [1990Che]. The crystallization of alloys containing less than 20 at.% Zr or Ti runs in one stage from the amorphous to the phase. The range of crystallization of the metastable Laves phase -1' occupies the central part of the diagram at 20-44 at.% Zr. In this concentration range, the amorphous alloys have highest crystallization temperature Tx = 478°C. At the same time, the maximum temperature of recrystallization of the -1' phase to the phase is observed at 600°C and 30 at.% Zr. The kinetics of the amorphous-to-Zr14Cu51 phase transformation in an as-cast Zr20Ti20Cu60 rod has been investigated by DSC by [2003Cao]. The relative volume fraction of the transformed crystalline phase as a function of annealing time, obtained for 440, 443,450, 455, and 460°C, has been analyzed in detail using a number of nucleation and growth models together with the Johnson-Mehl-Avrami model. A steady-state nucleation rate of the order of 1022-1023 nuclei m–3#s–1 in the temperature range 440-460°C and activation energy about 550 kJ#mol–1 for the phase transformation in the as-cast Zr20Ti20Cu60 rod were detected. The crystallization kinetics of ductile Zr20Ti20Cu60 ribbons was studied from the isothermal treatment and was analyzed using the Johnson-Mehl-Avrami-Kolmogorov model by [2004Rev, 2004Con2].
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
443
References [1960Bun]
[1961Enc]
[1965Tes]
[1969Tes]
[1979Dri]
[1983Dut]
[1983Mur] [1986Kne]
[1988Mas]
[1988Woy]
[1990Che]
[1991Sin]
[1992Kov]
[1992Nak]
[1994Ari]
Landolt-Börnstein New Series IV/11C3
Bunshah, R.F., Osterberg, D., Ence, E., Margolin, H., “Further Studies on Active-Eutectoid Alloys of Titanium”, U.S. Dept Comm. Office Tech. Serv., PB Rept., 161964, 1-73 (1960) (Crys. Structure, Experimental, Mechan. Prop., Phase Diagram, Phase Relations, 26) Ence, E., Margolin, H., “A Study of the Ti-Cu-Zr System and the Structure of Ti2Cu”, Trans. Metall. Soc. AIME, 221, 320-322 (1961) (Crys. Structure, Experimental, Phase Diagram, 9) Teslyuk, M.Yu., “Ternary Intermetallic Compounds with Structure of Laves Phases” (in Russian), Abstract Ph. D. Theses, L’viv University, L’viv, 16 pp. (1965) (Experimental, Crys. Structure, Theory) cited from abstract Teslyuk, M.Yu., “Intermetallic Compounds with Structure of Laves Phases” (in Russian), in “Intermetallic Compounds with Structure of Laves Phases”, Nauka, Moscow, 1-138 (1969) (Experimental, Crys. Structure, Theory) Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, E.M., Rokhlin, L.L., Turkina, N.I., “Copper-Niobium-Tin” in “Binary and Multicomponent Copper-Base System” (in Russian), Nauka, Moscow, 67 (1979) (Review, Phase Diagram, Phase Relations) Dutkiewicz, J., Massalski, T.B., “Search for New Amorphous Materials in Ternary Metal Systems”, Konf. Metalozn., (Mater. Konf.), 11th, 308-312 (1983) (Phase Diagram, Experimental) cited from abstract Murray, J.L., “The Cu-Ti (Copper-Titanium) System”, Bull. Alloy Phase Diagrams, 4(1), 81-95 (1983) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., 85) Kneller, E., Khan, Y., Corres, U., “The Alloy System Copper-Zirconium. Part I. Phase Diagram and Structural Relations”, Z. Metallkd., 77, 43-48 (1986) (Experimental, Phase Diagram, Crys. Structure, 26) Massalski, T.B., Woychik, C.G., Dutkiewicz, J., “Solidification Structures in Rapidly Quenched Copper-Titanium-Zirconium Alloys”, Metall. Trans. A, 19A(7), 1853-1860 (1988) (Phase Relations, Phase Diagram, Morphology, Experimental) cited from abstract Woychik, C.G., Massalski, T.B., “Phase Diagram Relationships in the System Cu-Ti-Zr”, Z. Metallkd., 79(3), 149-153 (1988) (Experimental, Phase Diagram, Phase Relations, Morphology, 19) Chebotnikov, V.N., Molokanov, V.V., “Structures and Properties of Amorphous and Crystalline Alloys in the Ti2Cu-Zr2Cu Section in the Ti-Zr-Cu System”, Inorg. Mater. (Engl. Trans.), 26(5), 808-811 (1990) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 7) Singh, R., Lawley, A., Freidman, S., Murty, Y, “Microstructure and Properties of Spray Cast Cu-Zr Alloys”, Mater. Sci. Eng., 145A, 243-255 (1991) (Experimental, Crys. Structure, Phys. Prop., 30) Kovneristyi, Yu.K., Pashkovskaya, A.G., “Bulk Amorphization of Alloys in the Intermetallic-Containing System Ti-Cu-Zr” (in Russian), in “Amorphous Glassy Metallic Materials”, Ros. Akad. Nauk, Baikov Inst. Metallurgy, 153-157 (1992) (Experimental, Morphology, Phase Diagram, Phase Relations, 4) Nakamura, N., “Development of Ceramic-Metal Composite Materials”, Kenkyu Hokoku Fukuoka-ken Kogyo Gijutsu Senta, 1991(2), 32-36 (1992) (Morphology, Experimental) cited from abstract Arias, D., Abriata, J.P., “Cu-Zr (Copper-Zirconium)” in “Phase Diagrams of Binary Copper Alloys”, Subramanian, P.R., Chakrabarti, D.J., Laughlin, D.E. (Eds.), ASM Materials Park, OH, 497-502 (1994) (Assessment, Phase Diagram, Phase Relations, Thermodyn., 50)
MSIT®
444 [1994Bot]
[1994Kum]
[1994Zen]
[1996Kum]
[1998Bra]
[1998Hir]
[1999Lee]
[2000Oht]
[2001Ino]
[2001Kun]
[2002Ans]
[2002Kas]
[2002Lou]
[2003Arr]
[2003Cao]
[2003Jia]
MSIT®
Cu–Ti–Zr Botstein, O., Rabinkin, A., “Brazing of Titanium-Based Alloys with Amorphous 25%Ti-25%Zr-50%Cu Filler Metal”, Mater. Sci. Eng. A, A188(1-2), 305-315 (1994) (Morphology, Phase Relations, Experimental) cited from abstract Kumar, H., Wollants, P., Delaey, L., “Thermodynamic Assessment of the Ti-Zr System and Calculation of the Nb-Ti-Zr Phase Diagram”, J. Alloys Compd., 206(1), 121-127 (1994) (Phase Relations, Phase Diagram, Thermodyn., Assessment, 32) Zeng, K.J., Haemaelaeinen, M., Lukas, H. L., “A New Thermodynamic Description of the Cu-Zr System”, J. Phase Equilib., 15(6), 577-586 (1994) (Phase Relations, Phase Diagram, Assessment, 55) Kumar, H., Ansara, I., Wollants, P., Delaey, L., “Thermodynamic Optimisation of the Cu-Ti System”, Z. Metallkd., 87(8), 666-672 (1996) (Assessment, Phase Relations, Thermodyn., 55) Braga, M.H., Malheiros, F.M., Castro, F., Soares, D., “Experimental Points and Invariant Reactions in the Cu-Zr System”, Z. Metallkd., 89(8) 541-545 (1998) (Experimental, Phase Diagram, 31) Hirose, A., Nojiri, M., Ito, H., Kobayashi, K., “Brazing of Ti Alloys with Ti-Zr-Cu Amorphous Filler Metal”, Int. J. Mater. Prod. Techn., 13(1-2), 13-27 (1998) (Morphology, Phase Relations, Experimental) cited from abstract Lee, P.Y., Lin, C.K., Chen, G.S., Louh, R.F., Chen, K.C., “Formation of Cu-Zr-Ti Amorphous Alloys with Significant Supercooled Liquid Region by Mechanically Alloying”, Mat. Sci. Forum, 312-314, 67-72 (1999) (Morphology, Experimental) cited from abstract Ohta, S., Inoue, K., Nishida, M., Araki, T., “Reaction Mechanism Between Ti and Ni on Brazing Process”, Yosetsu Gakkai Ronbunshu, 18(2), 198-207, (2000) (Morphology, Kinetics, Experimental) cited from abstract Inoue, A., Zhang, W., Zhang, T, Kurosaka, K., “High-Strength Cu-Based Bulk Glassy Alloys in Cu-Zr-Ti and Cu-Hf-Ti Ternary Systems”, Acta Mater., 49(14), 2645-2652 (2001) (Morphology, Phase Relations, Mechan. Prop., Experimental, 36) Kunimasa, T., Seki, M., Yamamoto, H., Nojiri, M., Uenishi, K., Kobayashi, K. F., “Brazing of Ti-Ni Shape Memory Alloy with Pure Titanium”, Zairyo, 50(11), 1218-1222 (2001) cited from abstract Ansara I., Ivanchenko, V., “Cu-Ti (Cooper-Titanium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.11457.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, 26) Kasai, M., Saida, J., Matsushita, M., Osuna, T., Matsubara, E., Inoue, A., “Structure and Crystallization of Rapidly Quenched Cu-(Zr or Hf)-Ti Alloys Containing Nanocrystalline Particles”, J. Phys.: Condens. Matter, 14, 13867-13877 (2002) (Crys. Structure, Experimental, Morphology, 13) Louzguine, D.V., Inoue, A., “Nanocrystallization of Cu-(Zr or Hf)-Ti Metallic Glasses”, J. Mater. Res., 17(8), 2112-2120 (2002) (Crys. Structure, Phase Relations, Morphology, Experimental, 32) Arroyave, R., Eagar, T.W., Kaufman, L., “Thermodynamic Assesment of the Cu-Ti-Zr System”, J. Alloys Compd., 351, 158-170 (2003) (Assessment, Calculation, Phase Relations, Thermodyn., *, #, 21) Cao, Q.P., Zhou, Y.H., Horsewell, A., Jiang, J.Z., “Amorphous-to-Cu51Zr14 Phase Transformation in Cu60Ti20Zr20 Alloy”, J. Phys.: Condens. Matter, 15(50), 8703-8712 (2003) (Phase Relations, Kinetics, Experimental, 40) Jiang, J.Z., Kato, H., Ohsuna, T., Saida, J., Inoue, A., Saksl, K., Franz, H., Stahl, K., “Origin of Nandetectable X-Ray Diffraction Peaks in Nanocomposite CuTiZr Alloys”, Appl. Phys. Lett., 83(16), 3299-3301 (2003) (Crys. Structure, Experimental, Phase Relations, 27)
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr [2004Con1]
[2004Con2]
[2004Par]
[2004Rev]
[2006Sem]
445
Concustell, A., Revesz, A., Surinach, S., Varga, L.K., Heunen, G., “Microstructural Evolution During Decomposition and Crystallization of the Cu60Zr20Ti20 Amorphous Alloy”, J. Mater. Res., 19(2), 505-512 (2004) (Crys. Structure, Experimental, Kinetics, Morphology, Phase Relations, 34) Concustell, A., Zielinska, M., Revesz, A., Varga, L.K., Surinach, S., Baro, M. D., “Thermal Characterization of Cu60ZrxTi40–x Metallic Glasses (x=15, 20, 22, 25, 30)”, Intermetallics, 12(10-11), 1063-1067 (2004) (Morphology, Structure, Kinetics, Mechan. Prop., Experimental, 34) Park, E.S., Chang, H.J., Kim, D.H., Kim, W.T., Kim, Y.C., Kim, N.J., Kim, Y.W., “Formation of Amorphous Phase in Melt-Spun and Injection-Cast Cu60Zr30Ti10 Alloys”, Scr. Mater., 51(3), 221-224 (2004) (Crys. Structure, Experimental, Morphology, Phase Relations, 16) Revesz, A., Concustell, A., Varga, L.K., Surinach, S., Baro, M.D., “Influence of the Wheel Speed on the Thermal Behavior of Cu60Zr20Ti20 Alloys”, Mater. Sci. Eng. A, A375-A377, 776-780 (2004) (Morphology, Crys. Structure, Kinetics, Mechan. Prop., Experimental, 34) Semenova, E., Sidorko, V., “Cu-Zr (Cooper-Zirconium)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2006) (Crys. Structure, Phase Diagram, Assessment, 31)
Table 1: Investigations of the Cu-Ti-Zr Phase Relations and Structures Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1961Enc]
Optical microscopy, X-ray analysis
< 50 mass% Ti, isothermal section at 750°C
[1983Dut]
Optical microscopy, EMPA
liquidus surface projection, low melting eutectics
[1988Woy]
Optical microscopy, EMPA, Transient Liquid Phase bonding technique
isothermal section at 703°C and >30 at.% Cu, eutectic reactions in central portion of phase diagram
[1990Che]
Optical microscopy, X-ray analysis, Temperature-composition Ti2Cu–Zr2Cu section differential thermal analysis, measurements of microhardness
[1992Kov]
Optical microscopy, X-ray analysis, Temperature-composition TiCu-ZrCu section differential thermal analysis
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Ti–Zr
446 Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Lattice Parameters Comments/References [pm] Space Group/ Prototype
(Cu), ZrxTiyCu1–x–y < 1084.87
cF4 Fm3m Cu
, Zr1–x–yTixCuy 1854 - 600
cI2 Im3m W
a = 361.46
pure Cu at 25°C [Mas2] x = 0.0012 at 968°C and y = 0 [1994Zen] y = 0.08 at 896°C and x = 0 [1996Kum] 0 < x < 1 in Zr-Ti system [1994Kum, Mas2]
(Zr) 1854 - 863
a = 360.90
pure Zr(h) at T > 863°C [Mas2] y = 0.037 at 997°C and x = 0 [1994Zen]
(Ti) 1668 - 882
a = 330.65
pure Ti(h) at T > 882°C [Mas2] y = 0.159 at 1000°C in Cu-Ti system [1996Kum]
, Zr1–x–yTixCuy (Zr) < 863
hP2 P63/mmc Mg
(Ti) < 882
a = 323.16 c = 514.75
0 < x < 1 in Zr-Ti system [Mas2, 1994Kum] pure Zr(r) at 25°C [Mas2] y = 0.002 at 816°C and x = 0 [1994Zen]
a = 295.06 c = 468.35
pure Ti(r) at 25°C [Mas2] y = 0.014 at 799°C in Cu-Ti system [1994Kum]
ZrCu5 < 968
cF24 F43m AuBe5
a = 687.0
[1994Zen, 1991Sin, 1986Kne]
Zr14Cu51 < 1113
hP68 P6/m Gd14Ag51
a = 1124.44 c = 828.15
[1994Zen, V-C2]
Zr3Cu8 < 891
oP44 Pnma Hf3Cu8
a = 786.93 b = 815.47 c = 998.48
[1994Zen, V-C2]
Zr7Cu10, (Zr1–xTix)7Cu10 < 935
oC68 C2ca Zr7Ni10
ZrCu 936-704
cP2 Pm3m CsCl
MSIT®
a = 1267.29 b = 931.63 c = 934.66 a = 325.87
0 < x < 0.24 at 703°C [1988Woy] [1994Zen, 1994Ari, 1998Bra]
[1994Zen, 1994Ari, 1998Bra]
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
447
Phase/ Temperature Range [°C]
Pearson Symbol/ Lattice Parameters Comments/References Space Group/ [pm] Prototype
, (Zr1-xTix)2Cu Zr2Cu < 1011
tI6 I4/mmm MoSi2
Ti2Cu < 1012
a = 322.04 c = 1183.2
0 < x < 1 [1990Che] [1994Zen, 1994Ari]
a = 324.26 c = 1127.2
[1986Kne, 1998Bra]
a = 295.3 c = 1073.4
[1996Kum, V-C2] [2002Ans]
TiCu4 < 896
oP20 Pnma ZrAu2
a = 452.5 b = 434.1 c = 1295
78.6-80.8 at.% Cu [1996Kum, Mas2, V-C2]
TiCu2 870 - 890
oC12 Cmm2 VAu2
a = 436.3 b = 797.7 c = 447
[1996Kum, V-C2]
Ti2Cu3 < 874
tP10 P4/nmm Ti2Cu3
a = 313 c = 1395
[1996Kum, V-C2]
Ti3Cu4 < 933
tI14 I4/mmm Ti3Cu4
a = 313.0 c = 1994
[2002Ans, V-C2]
TiCu, ZrxTi1–xCu < 975
tP4 P4/nmm TiCu
a = 312.5 c = 591.5
[V-C2] 48.5-51.7 at.% Cu at x = 0 [1996Kum] 0 < x < 0.2 at 703°C [1988Woy] 0 < x < 0.08 at 842°C [1992Kov]
* -1, ZrTiCu2
hP12 P63/mmc MgZn2
a = 513 c = 825
ZrTiCu4 (?) [1965Tes, 1969Tes]
< 867
< 883
Landolt-Börnstein New Series IV/11C3
25 at.% Zr, 25 at.% Ti, 50 at.% Cu at 867°C; 21 to 37 at.% Zr at 50 at.% Cu at ~ 830°C; 22.5 to 29 at.% Zr at 50 at.% Cu at ~ 630°C [1992Kov] 25 to 30 at.% Zr and 45 to50 at.% Cu at 703°C [1988Woy] calculated by [2003Arr]
MSIT®
Cu–Ti–Zr
448 Table 3: Invariant Four-Phase Equilibria Reaction
L + TiCu Ti3Cu4 + Zr14Cu51
T [°C]
Type
Phase
Composition (at.%) Cu
Zr
Ti
U1
L TiCu Ti3Cu4 Zr14Cu51
64.5 51.3 57.1 78.5
6.8 0 0 21.5
28.7 48.7 42.9 0
L + Ti3Cu4 + TiCu2 Ti2Cu3 874
D1(P)
L Ti3Cu4 TiCu2 Ti2Cu3
73.4 57.1 66.7 60.0
25.2 0 0 0
1.4 42.9 33.3 40.0
L + TiCu2 TiCu4 + Ti2Cu3
869
D2(U)
L TiCu2 TiCu4 Ti2Cu3
73.7 66.7 78.7 60.0
2.0 0 0 0
24.3 33.3 21.3 40.0
L + (Cu) TiCu4 + ZrCu5
863
U2
L (Cu) TiCu4 ZrCu5
80.1 93.5 81.2 83.3
4.0 0.1 0 16.7
15.9 6.4 18.8 0
L + Zr14Cu51 Zr3Cu8 + -1
861
U3
L Zr14Cu51 Zr3Cu8 -1
56.5 78.5 72.7 50.0
33.3 21.5 27.3 25.0
10.2 0 0 25.0
L + -1 + TiCu
856
U4
L -1 TiCu
50.7 33.3 50.0 49.6
14.0 4.1 25.0 0
35.3 62.6 25.0 50.4
L + ZrCu5 Zr14Cu51 + TiCu4
855
U5
L ZrCu5 Zr14Cu51 TiCu4
75.7 83.3 78.5 79.6
4.2 16.7 21.5 0
20.1 0 0 20.4
L TiCu + -1 + Zr14Cu51
855
E1
L TiCu -1 Zr14Cu51
52.7 49.8 50.0 78.5
14.2 0 25.0 21.5
33.1 50.2 25.0 0
L + Ti3Cu4 Ti2Cu3 + Zr14Cu51
854
U6
L Ti3Cu4 Ti2Cu3 Zr14Cu51
73.7 57.1 60.0 78.5
4.2 0 0 21.5
22.1 42.9 40.0 0
L TiCu4 + Ti2Cu3+ Zr14Cu51
852
E2
L TiCu4 Ti2Cu3 Zr14Cu51
74.2 79.0 60.0 78.5
4.1 0 0 21.5
21.7 21.0 40.0 0
MSIT®
882
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr Reaction
T [°C]
Type
449 Phase
Composition (at.%) Cu
Zr
Ti
L + Zr3Cu8 -1 + Zr7Cu10
845
U7
L Zr3Cu8 -1 Zr7Cu10
55.8 72.7 50.0 58.8
36.8 27.3 25.0 41.2
7.4 0 25.0 0
L -1 + ZrCu + Zr7Cu10
839
E3
L -1 ZrCu Zr7Cu10
52.3 50.0 50.0 58.8
39.9 25.0 50.0 41.2
7.8 25.0 0 0
L + -1 +
835
U8
L -1
34.4 33.3 50.0 11.6
28.8 20.6 25.0 25.3
36.8 46.1 25.0 63.1
L -1 + + ZrCu
833
E4
L -1 ZrCu
44.8 50.0 33.3 50.0
41.9 25.0 58.7 50.0
13.3 25.0 8.0 0
L -1 + +
827
E5
L -1
35.8 50.0 33.3 12.2
33.7 25.0 35.8 31.0
30.5 25.0 30.9 56.7
ZrCu -1 + + Zr7Cu10
720
E6
ZrCu -1 Zr7Cu10
50.0 50.0 33.3 58.8
50.0 25.0 63.8 41.2
0 25.0 2.9 0
U9
-
-
-
-
Phase
Composition (at.%)
TiCu4 + Zr14Cu51 ZrCu5 + 703-750 Ti2Cu3
Table 4: Invariant Two- and Three-Phase Equilibria Reaction
T [°C]
Type
Cu
Zr
Ti
L -1
883
congruent
L -1
50.0 50.0
25.0 25.0
25.0 25.0
L TiCu + Zr14Cu51
885
e9 (max)
L TiCu Zr14Cu51
61.4 50.8 78.5
8.2 0 21.5
30.4 49.2 0
L -1 + Zr14Cu51
879
e10 (max)
L -1 Zr14Cu51
55.0 50.0 78.5
24.3 25.0 21.5
20.7 25.0 0
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Ti–Zr
450 T [°C]
Reaction
Type
Phase
Composition (at.%) Cu
Zr
Ti
L -1 + b)
869
e11 (max)
L -1 b)
43.6 50.0 33.3
18.1 25.0 7.0
38.3 25.0 59.7
L -1 + ZrCu
841
e12 (max)
L -1 ZrCu
50.0 50.0 50.0
40.5 25.0 50.0
9.5 25.0 0
L -1 + a)
835
e13 (max)
L -1 a)
42.0 50.0 33.3
39.9 25.0 56.4
18.1 25.0 10.3
a) and b) denote phase enriched by Zr or by Ti, respectively
Table 5: Investigations of the Cu-Ti-Zr Materials Properties Reference
Method/Experimental Technique
Type of Property
[1983Dut]
X-ray analysis, DSC
Tx, Tg of amorphous alloys,
[1988Mas]
X-ray analysis, DSC
Tg of amorphous alloys,
[1990Che]
X-ray analysis, electroresistance tests, microhardness tests
electroresistance ', microhardness H, embrittlement temperature Tb, and critical thickness dc of the amorphous alloys of Zr2Cu-Ti2Cu section
[1992Kov]
X-ray analysis, DTA
Critical thickness dc and critical cooling rates Rc of the amorphous alloys of ZrCu-TiCu section
[1994Bot]
Energy dispersive X-ray spectroscopy, X-ray analysis, SEM, microhardness tests
Microstructure and mechanical properties of brazed joints
[2001Ino]
DSC, mechanical properties tests
Structure of Zr30Ti10Cu60 bulk amorphous alloy
[2001Kun]
X-ray analysis, mechanical properties tests
Microstructure and mechanical properties of brazed joints
[2002Kas]
Structure of the as-cast and as-spun Zr30Ti10Cu60 X-ray analysis, anomalous X-ray scattering measurements, high resolution amorphous alloy TEM, energy dispersive X-ray spectroscopy
[2002Lou]
X-ray analysis, TEM, energy dispersive Structure and crystallization of Zr30Ti10Cu60 X-ray spectroscopy, DSC amorphous alloy
[2003Jia]
X-ray analysis, DSC
Structure and crystallization of Zr30–xTi10+xCu60 amorphous alloys
[2003Cao]
DSC
Kinetic of crystallization of Zr20Ti20Cu60 bulk amorphous alloy
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr Reference
Method/Experimental Technique
451
Type of Property
[2004Con1] X-ray analysis, DSC, DTA, microhardness tests
Crystallization, thermal stability and mechanical properties of ZrxTi40–xCu60 (x = 15, 20, 22, 25, 30) amorphous alloys
[2004Con2] X-ray analysis, DSC, synchrotron radiation transmission, high resolution TEM
Crystallization and heat release of Zr20Ti20Cu60 amorphous alloy
[2004Rev]
Kinetic of crystallization of Zr20Ti20Cu60 amorphous alloy
DSC
3.70 3.60
c/a
Fig. 1: Cu-Ti-Zr. Lattice parameters of phase as function of composition along the Zr2Cu-Ti2Cu section
c/a
3.50 3.40
c 11.10
a
11.00
a, D
10.90 3.00
c, D
3.20
10.80 10.70
Ti2Cu
10
20
30
40
50
60
Zr2Cu
Zr, at%.
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Ti–Zr
452
890
Fig. 2: Cu-Ti-Zr. Calculated partially quasibinary system ZrCu--1
875
Temperature, °C
L
L+τ1
L+ZrCu
850
841°C
τ1+ZrCu
835
Zr 25.00 Ti 25.00 Cu 50.00
Fig. 3: Cu-Ti-Zr. Calculated quasibinary system -1-Zr14Cu51
30
Zr 50.00 0.00 Ti Cu 50.00
40
Zr, at.%
1100
Temperature, °C
L
1000
L+Zr14Cu51
900
879°C L+τ1
Zr 25.00 Ti 25.00 Cu 50.00
MSIT®
τ1+Zr14Cu51 20
10
Ti, at.%
Zr 21.50 0.00 Ti Cu 78.50
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
Fig. 4: Cu-Ti-Zr. Calculated quasibinary system Zr14Cu51-TiCu
453
Temperature, °C
1100
L
1000
Zr14Cu51+L
L+TiCu 900
885°C Zr14Cu51+TiCu
Zr 78.50 0.00 Ti Cu 21.50
10
20
30
0.00 Zr Ti 50.00 Cu 50.00
40
Ti, at.%
Cu-Ti
Cu-Ti-Zr
A-B-C
Cu-Zr 1015
1005 e1 lβ+γ
p1
l+Zr14Cu51ZrCu5
996 e2 lγ+β 968 e3 l ΤiCu + γ
968 e4 l (Cu) + ZrCu5
933 p2 l +TiCu Τi3Cu4
924 e5 l ZrCu + γ 922
896 p4 l + (Cu) ΤiCu4
p3
l+Zr14Cu51Zr3Cu8
891
e6
lZr3Cu8+Zr7Cu10
890 887 p5 l+Ti3Cu4ΤiCu2
e7
lZr3Cu8+ZrCu
885 e8 l ΤiCu4+ TiCu2
885
e9(max)
LTiCu+Zr14Cu51
882 L+TiCuTi3Cu4+Zr14Cu51 U1 TiCu+Ti3Cu4+Zr14Cu51 879 e10(max) Lτ1+Zr14Cu51 U8 U4 U2 D1 D2 E1 U3
U3 U2 E1 U6 U5 U7
E3 E4 E5
Fig. 5a: Cu-Ti-Zr. Reaction scheme, part 1 Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Ti–Zr
454
Cu-Ti
Cu-Ti-Zr
A-B-C
Cu-Zr
e10 U1 p1 p3 e6 e4 e7 e5 e2
e1 e3 p4 p5 e8 e9 e10
874 L+Ti3Cu4+TiCu2Ti2Cu3 D1 TiCu2+Ti3Cu4+Ti2Cu3 869 e11(max) L τ1 + γ
869 L+TiCu2TiCu4+Ti2Cu3 D2 TiCu2+TiCu4+Ti2Cu3 863 L+(Cu) TiCu4+ZrCu5 U2 TiCu4+(Cu)+ZrCu5 861 L+Zr14Cu51Zr3Cu8+τ1 U3 τ1+Zr14Cu51+Zr3Cu8 856
L + γ τ1 + TiCu
L+γ+TiCu
U4
τ1+γ+TiCu
855 L TiCu+ Zr14Cu51+τ1 E1 855 L+ZrCu5Zr14Cu51+TiCu4 U5 TiCu+τ1+Zr14Cu51 854 L+Ti3Cu4Ti2Cu3+Zr11Cu51 U6 852 LTiCu4+Ti2Cu3+ Zr14Cu51 E2 Ti2Cu3+Ti3Cu4+Zr11Cu51 845 L+Zr3Cu8τ1+ Zr7Cu10 U7 841 e12(max) Lτ1+ ZrCu
τ1+Zr3Cu8+Zr7Cu10 839 L τ1+ZrCu+ Zr7Cu10 E3
835
L + γ τ1 + β
835 e13(max) Lτ1+ γ
U8
τ1+γ+β 833
L τ1 + γ + ZrCu
E4 827
L τ1 + γ + β
E5 725
720 ZrCu τ1+γ+Zr7Cu10 E6
τ1+γ+β
e14
CuZrγ+Zr7Cu10
τ1+γ+Zr7Cu10 703
Zr14Cu51+ZrCu5+Ti2Cu3
ZrCu5+ZrCu4+Ti2Cu3
Fig. 5b: Cu-Ti-Zr. Reaction scheme, part 2
MSIT®
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
455
Cu
Data / Grid: at.% Axes: at.%
Fig. 6a: Cu-Ti-Zr. Calculated liquidus projection with liquidus isotherms
e4 ZrCu5 p1
(Cu) U2
p4
20
80
1075 Zr14Cu51
Zr7Cu10 p
3
40 Zr3Cu8 e6 e7
ZrCu e5
925 E3
60
γ
TiCu D2 e 4 p8 D1 5 Ti3Cu4 p2 e9 60
U5 E2 975 U6 U1 E1
U7U3 e12
U4 TiCu
e10
τ1 e13
E4
E5 U8
e2
e11
γ
875
925
975
e3 40
e1 1075
80
20
β
1375 1475 1575°C
20
Zr
40
60
Zr Ti Cu
Fig. 6b: Cu-Ti-Zr. Enlarged part of the liquidus projection
80
0.00 14.00 86.00
Ti
Data / Grid: at.% Axes: at.%
(Cu)
U2
ZrCu5
80
p4 TiCu4
10
U5 Zr14Cu51
E2
Ti2Cu3 D2
U6 Ti3Cu4 Zr Ti Cu Landolt-Börnstein New Series IV/11C3
14.00 14.00 72.00
20
e8 D1
TiCu2 p5 Zr Ti Cu
0.00 28.00 72.00
MSIT®
Cu–Ti–Zr
456
Cu
Data / Grid: at.%
Fig. 7: Cu-Ti-Zr. Solidus surface projection
Axes: at.%
(Cu) (Cu)+ZrCu5 ZrCu5 20
Zr14Cu51 Zr3Cu8
863 80
855
TiCu4
852
861
854
879 40
Zr7Cu10
Ti2Cu3
855
845
60
885 841
ZrCu
Ti3Cu4
τ 1,883
882
TiCu
γ +ZrCu 60
833 835
γ ,Zr2Cu
869
827
40
γ +τ 1
γ ,Ti2Cu
β +γ
80
20
β +τ 1
β 20
Zr
γ +TiCu
856
40
60
80
Cu Fig. 8: Cu-Ti-Zr. Calculated isothermal section at 925°C
Ti
Data / Grid: at.% Axes: at.%
(Cu) L+(Cu)+ZrCu5 ZrCu5 L+ZrCu5+Zr14Cu51 20 Zr14Cu51
80
40
60
ZrCu
γ ,Zr2Cu
TiCu L+γ +TiCu
L
60
Ti3Cu4 TiCu+Ti3Cu4+L
40
γ ,Ti2Cu
L+β +γ
80
20
L+β +γ
β
Zr
MSIT®
20
40
60
80
Ti
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
457
Cu Fig. 9: Cu-Ti-Zr. Calculated isothermal section at 875°C
Data / Grid: at.% Axes: at.%
(Cu)
L+(Cu)+ZrCu5 ZrCu5
L+(Cu)+TiCu4
L+ZrCu5+Zr14Cu51 20 Zr14Cu51 Zr3Cu8 L+τ 1+Zr14Cu51 L+Zr3Cu8+Zr14Cu51 L+Zr3Cu8+Zr7Cu10
TiCu4 80
L
L+TiCu4+TiCu2 L+Ti3Cu4+Zr14Cu51 TiCu2 TiCu+Ti3Cu4+Zr14Cu51
40
60
Zr7Cu10 ZrCu L+γ +ZrCu
τ1
60
Ti3Cu4 L+TiCu+Zr14Cu51 TiCu L+γ +TiCu 40
L
γ ,Zr2Cu
γ ,Ti2Cu
L+β +γ
L+β +γ 80
20
β 20
Zr
40
60
80
Cu Fig. 10: Cu-Ti-Zr. Calculated isothermal section at 850°C
Data / Grid: at.% Axes: at.%
(Cu)
ZrCu5+Zr14Cu51+TiCu4
ZrCu5
(Cu)+TiCu4+ZrCu5 TiCu4
Zr14Cu51 20 τ 1+Zr14Cu51+Zr3Cu8 Zr3Cu8 L+τ 1+Zr3Cu8 L+Zr3Cu8+Zr7Cu10 Zr7Cu10 L+ZrCu+Zr7Cu10
80
TiCu4+Ti2Cu3+Zr14Cu51 Ti2Cu3+Ti3Cu4+Zr14Cu51 Ti2Cu3
40
60
ZrCu
τ1
L
Ti3Cu4 TiCu+Ti3Cu4+Zr14Cu51 TiCu τ 1+γ +TiCu
τ 1+TiCu+Zr14Cu51
L+γ +τ 1
60
L+γ +ZrCu γ ,Zr2Cu
Ti
40
γ ,Ti2Cu
L+β +γ
L+β +γ
80
20
β
Zr
Landolt-Börnstein New Series IV/11C3
α
20
40
60
80
α
Ti
MSIT®
Cu–Ti–Zr
458
Cu Fig. 11: Cu-Ti-Zr. Calculated isothermal section at 827°C
Data / Grid: at.% Axes: at.%
(Cu) (Cu)+TiCu4+ZrCu5
ZrCu5 ZrCu5+Zr14Cu51+TiCu4
TiCu4
Zr14Cu51 20
80
τ 1+Zr14Cu51+Zr3Cu8
TiCu4+Ti2Cu3+Zr14Cu51
Zr3Cu8 τ 1+Zr3Cu8+Zr7Cu10 Zr7Cu10
Ti2Cu3+Ti3Cu4+Zr14Cu51 Ti2Cu3
40
60
τ 1+ZrCu+Zr7Cu10
τ1
ZrCu 60
Ti3Cu4 TiCu+Ti3Cu4+Zr14Cu51 τ 1+TiCu+Zr14Cu51
τ 1+γ +TiCu
τ 1+γ +ZrCu
TiCu 40
L
γ ,Zr2Cu
L+τ 1+γ
γ ,Ti2Cu
τ 1+β +γ
L+β +γ
80
20
L+β +τ 1
β
Zr
α
20
40
60
Cu Fig. 12: Cu-Ti-Zr. Calculated isothermal section at 810°C
Axes: at.%
(Cu)+TiCu4+ZrCu5
ZrCu5 ZrCu5+Zr14Cu51+TiCu4
TiCu4
Zr14Cu51 20
80
τ 1+Zr14Cu51+Zr3Cu8
TiCu4+Ti2Cu3+Zr14Cu51
Zr3Cu8
Ti2Cu3+Ti3Cu4+Zr14Cu51 Ti2Cu3
τ 1+Zr3Cu8+Zr7Cu10 Zr7Cu10 40
60
Ti3Cu4 TiCu+Ti3Cu4+Zr14Cu51
τ 1+ZrCu+Zr7Cu10 ZrCu
τ1
τ 1+γ +ZrCu
τ 1+γ +TiCu
τ 1+β +γ
γ ,Zr2Cu
TiCu 40
τ 1+β +γ
γ ,Ti2Cu
80
20
α+β +γ
Zr
MSIT®
α
Ti
Data / Grid: at.%
(Cu)
60
α
80
β 20
40
60
80
α
Ti
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
459
Cu Fig. 13a: Cu-Ti-Zr. Calculated isothermal section at 750°C
Data / Grid: at.% Axes: at.%
(Cu) (Cu)+TiCu4+ZrCu5
ZrCu5 ZrCu5+Zr14Cu51+TiCu4
TiCu4
Zr14Cu51 20 τ 1+Zr14Cu51+Zr3Cu8 Zr3Cu8 τ 1+Zr3Cu8+Zr7Cu10 Zr7Cu10
80
TiCu4+Ti2Cu3+Zr14Cu51 Ti2Cu3+Ti3Cu4+Zr14Cu51 Ti2Cu3
40
60
τ 1+ZrCu+Zr7Cu10
τ 1+TiCu+Zr14Cu51
ZrCu
τ1
τ 1+γ +ZrCu
60
Ti3Cu4 TiCu+Ti3Cu4+Zr14Cu51
τ 1+γ +TiCu
TiCu 40
γ ,Ti2Cu
γ ,Zr2Cu 80
20
α+β +γ
α+β +γ β
Zr
α
20
40
60
Zr Ti Cu
Fig. 13b: Cu-Ti-Zr. Experimental partial isothermal section at 750°C
α
80
0.00 60.00 40.00
Ti
Data / Grid: at.% Axes: at.%
γ 10
γ+α
30
20
20
β+γ
30
10
α+β +γ
β Zr Ti Cu Landolt-Börnstein New Series IV/11C3
40.00 60.00 0.00
70
α +β 80
α 90
Ti
MSIT®
Cu–Ti–Zr
460
Cu Fig. 14: Cu-Ti-Zr. Calculated isothermal section at 703°C
Data / Grid: at.% Axes: at.%
(Cu)
ZrCu5+Zr14Cu51+Ti2Cu3
(Cu)+TiCu4+ZrCu5
ZrCu5
TiCu4
Zr14Cu5120 τ 1+Zr14Cu51+Zr3Cu8 Zr3Cu8
80
TiCu4+Ti2Cu3+ZrCu5 Ti2Cu3+Ti3Cu4+Zr14Cu51
τ 1+Zr3Cu8+Zr7Cu10
Ti2Cu3
40
60
Zr7Cu10
Ti3Cu4 TiCu+Ti3Cu4+Zr14Cu51
τ 1+TiCu+Zr14Cu51 τ 1+γ +Zr7Cu10
τ1
60
τ 1+γ +TiCu
TiCu 40
γ ,Ti2Cu
γ ,Zr2Cu 80
20
α+β +γ α+β +γ
Zr
Fig. 15: Cu-Ti-Zr. Calculated temperature-composition Zr2Cu-Ti2Cu section
β 20
α
40
60
α
80
Ti
1000
Temperature, °C
L L+γ
L+γ
900
L+γ+β L+γ+β
L+β L+τ1+β
U8
β +γ+τ1 β +τ1
E5 800
β +γ+τ1 γ+β
Zr 66.70 0.00 Ti Cu 33.30
MSIT®
10
20
30
Ti, at.%
40
50
60
0.00 Zr Ti 66.70 Cu 33.30
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
Fig. 16: Cu-Ti-Zr. Calculated temperature-composition ZrCu-TiCu section
461
1100
1000
L
Temperature, °C
900
L+τ1 L+τ1
L+ZrCu
839
L+TiCu
855
856 800
τ1+ZrCu τ1+TiCu+Zr14Cu51
720°C 700
τ1+γ+Zr7Cu10 600
500
Zr 50.00 0.00 Ti Cu 50.00
20
30
0.00 Zr Ti 50.00 Cu 50.00
40
Ti, at.%
880
875
L
Temperature, °C
Fig. 17: Cu-Ti-Zr. Enlarged part of calculated temperature-composition ZrCu-TiCu section at the Ti composition 32 - 40 at.%
10
L+TiCu
L+τ1 L+γ+TiCu
L+γ L+τ1+γ
856°C 855 L+τ1+TiCu
τ1+Zr14Cu51+TiCu
850
Zr 18.00 Ti 32.00 Cu 50.00
Landolt-Börnstein New Series IV/11C3
34
36
Ti, at.%
38
Zr 10.00 Ti 40.00 Cu 50.00
MSIT®
Cu–Ti–Zr
462
870
L+TiCu
Temperature, °C
Fig. 18: Cu-Ti-Zr. Enlarged part of the calculated temperature-composition ZrCu-TiCu section at the Ti composition 42 - 50 at.%
L+TiCu+Zr14Cu51
L+γ+TiCu
856
855°C L+τ1+TiCu TiCu+Zr14Cu51
850
τ1+TiCu+Zr14Cu51 845
8.00 Zr Ti 42.00 Cu 50.00
44
46
Ti, at.%
Cu
Data / Grid: at.% Axes: at.%
Fig. 19: Cu-Ti-Zr. Glass forming range in Cu-Ti-Zr alloys 20
80
40
60
60
40
80
Zr
MSIT®
0.00 Zr Ti 50.00 Cu 50.00
48
20
20
40
60
80
Ti
Landolt-Börnstein New Series IV/11C3
Cu–Ti–Zr
ρ, μOhm.sm
Fig. 20: Cu-Ti-Zr. Variation of properties of amorphous alloys with composition along the Zr2Cu-Ti2Cu section: a - electroresistance; b - microhardness; c - embrittlement temperature; d - critical thickness of amorphous alloys.
463
200
180
(a)
Tb, °C
H, MPa
160 6000
4000 377
(b) (c)
327 277
d c, μm
400 300 200 100
(d)
0
Zr2Cu
20
40
Ti2Cu
Ti, at.%
Fig. 21: Cu-Ti-Zr. Concentration dependence of critical thickness of amorphous alloys along the ZrCu-TiCu section
1600
d c, μm
1200
800
400
0
ZrCu
10
20
30
40
TiCu
Ti, at.%
Landolt-Börnstein New Series IV/11C3
MSIT®
Cu–Ti–Zr
464
627
γ
τ1 527
Tx, °C
Fig. 22: Cu-Ti-Zr. Diagram of metastable crystallization of amorphous alloys along the Zr2Cu-Ti2Cu section
Amorphous 427
327
Zr2Cu
20
40
Ti2Cu
Ti, at.%
MSIT®
Landolt-Börnstein New Series IV/11C3
In–Sn–Zn
465
Indium – Tin – Zinc Hans Leo Lukas Introduction In the first phase diagram work a sketch of the liquidus surface was given by [1954Spe] as part of investigation of the quinary system Cd-Ge-In-Sn-Zn. He reported a ternary eutectic at 108°C and 52.2In-46Sn-1.8Zn (mass%), but did not yet distinguish between (In), (Sn) and the binary In-Sn phases. A detailed investigation of the phase diagram was presented by [1998Xie]. These authors measured liquidus temperatures and secondary arrests by DTA at 40 samples along four isopleths (Sn/Zn = 2, 1, 0.5 and constant 10 at.% In). Liquidus temperatures were also measured by [1998Yuj] using DTA along 10 sections from 0 to 50 at.% Zn and along constant 1 at.% Zn. [2001Sab1] measured liquidus temperatures as well as lower phase transitions by scanning calorimetry. The data of the three papers agree very well. Phase boundaries in completely solid state between (In) + (Zn), (In) + + (Zn), + (Zn), + + (Zn), + (Zn), + (Sn) + (Zn) and (Sn) + (Zn) could not be exactly located, as the diffusion is too slow at these low temperatures [1998Xie]. The vapor pressure of Zn was measured by the Knudsen cell method at 9 samples at about 5, 10 and 15 at.% Zn and In/Sn ratios of 0.334, 0.920 and 3.00. The results are presented as log p = A + BT equations and activity of Zn plots at 352°C. In another paper [2003Wnu] the influence of small concentrations of In on the Zn-activity of dilute solutions of Zn in (Sn) was determined and expressed by interaction parameters. They summarized the results by the equation ln fZn = 1.268 – 8.621xZn – 3.19xIn, valid for 750 K, where fZn is the activity coefficient of Zn. The partial enthalpies of mixing of In, Sn and Zn in ternary In-Sn-Zn melts were measured by dropping small amounts of the solid elements into the calorimetric bath, starting with a binary liquid of equal amounts of moles of the other two elements. Preparation of the calorimetric bath in at least 10 steps was used for calibration. By summing up consecutive measurements integral enthalpies of mixing were calculated. These measurements are part of similar measurements in the quaternary system In-Pb-Sn-Zn. They were represented by tables and isoenthalpy plots smoothed by the Redlich-Kister formalism with a single ternary term and by Darken’s formalism distinguishing “solvent” and “solutes” in the ternary system, and thus getting three ternary parameters. Using the same method [2000Anr] measured enthalpies of mixing of liquid at 70 samples. The results are tabulated as integral enthalpies, but can be recalculated to mean partial enthalpies of Zn or Sn in sections of different In/Sn (0.333, 1.0, 3.0) or Zn/In ratios (0.333, 1.0), respectively. The temperature was 339 to 341°C, for Zn in In/Sn = 0.333 also 634°C. In 1997 three review papers [1997Lee, 1997Hua, 1997Yoo] collected data on In-Sn-Zn alloys and some other elements for the use as Pb-free solders. The above mentioned experimental phase diagram investigations [1998Xie, 1998Yuj, 2001Sab1] as well as measurements of surface, wetting and corrosion properties [1999Pra, 2000Yu, 2001Sab2] are also mainly dedicated to the knowledge of Pb-free solder materials. There are three independent assessments of datasets for thermodynamic calculation of the In-Sn-Zn ternary system [2001Cui, 2001Xie, 2003Moe]. Phase diagrams calculated by the different sets do not deviate more than 5°C and agree well with the experimental data. The figures of this report are calculated from the dataset of [2003Moe]. Details of the experimental works on the In-Sn-Zn phase relations, structures and thermodynamics are given in Table 1. Binary Systems The binary systems are accepted as calculated from the binary datasets implemented into the ternary dataset of [2003Moe]: In-Sn modified by [2003Moe]; In-Zn [1996Lee]; Sn-Zn [1998Fri]. These systems are virtually equivalent to those of [Mas2].
Landolt-Börnstein New Series IV/11C3
MSIT®
466
In–Sn–Zn
Solid Phases There are no ternary phases in the system. Except along the binary In-Sn system ((In), , ) solubilities in the solid phases are small. The crystallographic data of unary and binary solid phases are given in Table 2. Invariant Equilibria The reaction scheme is given in Fig. 1 as calculated from the dataset of [2003Moe]. Details of the invariant reactions are given in Table 3. In the other calculations [2001Cui, 2001Xie] the invariant temperatures agree within 5°C. The experimentally determined and tabulated values [1954Spe, 1998Xie, 1998Yuj, 2001Sab1] of the invariant temperature of E1 all are within 107 and 109°C. The other tabulated experimental temperatures agree similarly well, as far as they can be clearly attributed to invariant reactions. Liquidus Surface The liquidus surface projection is shown in Fig. 2, calculated from the dataset of [2003Moe]. It is dominated by the area of primary crystallization of (Zn). Isothermal Sections The calculated isothermal section at 125°C is given in Fig. 3. The solubilities of Zn in the phases and are not determined but they are known to be small. In the calculation they are ignored. However, they may be of similar magnitude as the solubilities of Zn in (In) and (Sn). Thus in Fig. 3 the part along the binary In-Sn system must be taken as tentative. Temperature – Composition Sections The four temperature composition sections investigated by [1998Xie] are shown in Figs. 4 to 7. The curves are calculated from the dataset of [2003Moe], they reproduce very well the experimental points of [1998Xie]. Thermodynamics In the liquid phase enthalpies of mixing were experimentally investigated by two groups [1997Fio, 2000Anr]. In Fig. 8 the integral enthalpy as well as the partial enthalpy of Zn are shown along the section with In/Sn ratio of 1/1 vs xZn, calculated from the dataset of [2003Moe]. The datasets of [2001Cui, 2001Xie] give curves deviating less than 0.1 kJ#mol–1 from those of [2003Moe]. The curves fit the experimental data of [1997Fio, 2000Anr], except the partial enthalpies of Zn on the xIn = xSn line, where the deviation slightly exceeds the scatter of the experimental points, the calculated partial enthalpies are about 1 kJ#mol–1 smaller than the measured ones. The activity of Zn in nine alloys with less than 15 at.% Zn was determined by [1961Yok]. Calculations fit well with these data. Another activity determination was performed by [2003Wnu], who compared the Zn vapor pressure in dilute binary solutions of Zn in (Sn) with dilute ternary solutions of Zn and In in (Sn). The derived formula ln Zn = 1.268 8.621xZn 3.19xIn disagrees with the datasets [2001Cui, 2001Xie, 2003Moe], which give a much smaller dependence of the Zn activity aZn on the In content xIn. An argument against a strong dependence of aZn on xIn is the following consideration: in the two-phase field L + (Zn) the phase (Zn) is nearly pure Zn. Therefore the liquidus isotherms of the field of primary crystallisation of Zn are exactly isoactivity lines of Zn. As Fig. 2 shows, the ratio of the Zn content belonging to the same Zn activity between the In-Zn side and the Sn-Zn side is more than 0.5. This ratio is the factor, by which the activity coefficient of Zn increases by increasing the In mole fraction from zero to one. The above formula, however, predicts a decrease by the large factor exp(3.19) = 0.041 for a higher temperature, although activity coefficient usually decreases with increasing temperature.
MSIT®
Landolt-Börnstein New Series IV/11C3
In–Sn–Zn
467
Notes on Materials Properties and Applications In-Sn-Zn alloys are promising candidates to replace the toxic Sn-Pb alloys for soldering. Near the invariant reaction U1 the liquidus temperature is nearly the same as in the standard Sn-Pb solder alloy. Wetting behavior, with flux [2000Yu] and without flux [1999Pra], mechanical [1997Hua, 1999Pra] and corrosion [2001Sab2] properties were investigated. Due to the Zn content soldering preferably should be done under inert gas atmosphere to avoid formation of zincoxide. References [1954Spe]
[1961Yok] [1996Lee]
[1997Fio]
[1997Hua]
[1997Lee] [1997Yoo]
[1998Fri]
[1998Xie]
[1998Yuj]
[1999Pra]
[2000Anr]
[2000Yu]
[2001Cui]
Landolt-Börnstein New Series IV/11C3
Spengler, H., “A Study of Two- and Multi-Component Systems of the B-Metals: The Quinary System Cadmium-Germanium-Indium-Tin-Zinc” (in German), Metall, 8, 936-939 (1954) (Experimental, Phase Diagram, 19) Yokokawa, T., Doi, A., Niwa, K., “Thermodynamic Studies on Liquid Ternary Zinc Solutions”, J. Phys. Chem., 65, 202-205 (1961) (Experimental, Thermodyn., 9) Lee, B.-J., “Thermodynamic Assessment of the Sn-Zn and In-Zn Binary Systems”, Calphad., 20(4), 471-480 (1996) (Assessment, Calculation, Phase Diagram, Thermodyn., 51) Fiorani, J.M., Naguet, C., Hertz, J., Bourkba, A., Bouirden, L., “The Enthalpy of Mixing of the Quarternary (In,Pb,Sn,Zn) Liquid Homogeneous Phase”, Z. Metallkd., 88(9), 711-716 (1997) (Experimental, Thermodyn., 8) Hua, F., Glazer, J., “Lead-Free Solders For Electronic Assembly” in “Design and Reliability of Solders and Solder Interconnections”, Mahidhara, M.K., Frear, D.R., Sasiry, S.M.L., Murty, K.L., Liaw, P.K., Witerbottom, W., (Eds.), The Minerals, Metals and Materials Soc., 65-73 (1997) (Mechan. Prop., Phase Diagram, Review, 63) Lee, N.C., “Getting Ready for Lead-Free Solders“, Soldering Surf. Mount Technol., 9(2), 65-69 (1997) (Mechan. Prop., Review) Yoon, S.W., Soh, J.R., Lee, H.M., Lee, B.-J., “Thermodynamic Aided Alloy Design and Evaluation of Pb-free Solder, Sn-Bi-In-Zn System”, Acta Mater., 45(3), 951-960 (1997) (Calculation, Experimental, Phase Relations, 31) Fries, S., Lukas, H.L., “System Sn-Zn” in “COST 507 Thermochemical Database for Light Metal Alloys”, Ansara, I., Dinsdale, A.T., Rand, M.H., (Eds.), Vol. 2, European Community, Luxembourg, 288-289 (1998) (Assessment, Calculation, Phase Diagram, Thermodyn.) Xie, Y., Schicketanz, H., Mikula, A., “The In-Sn-Zn System“, Ber. Bunsen-Ges. Phys. Chem., 102(9), 1334-1338 (1998) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 13) Yujun, Z., Qiyun, Z., “Phase Diagram of Sn-In-Zn System and its Use for Lead-Free Solder”, Acta Phys. Sin. (Chin. J. Phys.), 14(12), 1098-1103 (1998) (Experimental, Phase Diagram, 11) Prasad, L.C., Xie, Y., Mikula, A., “Lead Free Solder Materials In-Sn-Zn System“, J. Non-Cryst. Solids, 250-252, 316-320 (1999) (Experimental, Phase Diagram, Interface Phenomena, 16) Anres, P., Alaoui-Elbelghiti, M., Gambino, M., Bros, J.P., “Enthalpy of Formation of the (In-Sn-Zn) Liquid System”, Thermochim. Acta, 346, 49-56 (2000) (Experimental, Thermodyn., 17) Yu, S.-P., Liao, C.-L., Hon, M.-H., Wang, M.-C., “The Effects of Flux on the Wetting Characteristics of Near-Eutectic Sn-Zn-In Solder on Cu Substrate”, J. Mater. Sci., 35, 4217-4224 (2000) (Electronic Structure, Experimental, Mechan. Prop., Interface Phenomena, 19) Cui, Y., Liu, X.J., Ohnuma, I., Kainuma, R., Ohtani, H., Ishida, K., “Thermodynamic Calculation of the In-Sn-Zn Ternary System”, J. Alloys Compd., 320(2), 234-241 (2001) (Assessment, Calculation, Phase Diagram, Thermodyn., 16) MSIT®
In–Sn–Zn
468 [2001Sab1]
[2001Sab2]
[2001Xie]
[2003Moe]
[2003Wnu]
Sabbar, A., Zrineh, A., Gambino, M., Bros, J.P., “About Phase Equilibria in the Ternary System Indium-Tin-Zink” (in French), Thermochim. Acta, 369, 125-136 (2001) (Experimental, Phase Diagram, Thermodyn., 28) Sabbar, A., El Hajjaji, S., Ben Bachir, A., Alaoui-El Belghiti, M., Zrineh, A., “Evaluation of Corrosion Behaviour of a New Class of Pb-Free Solder Materials (Sn-In-Zn)”, Mater. Corr., 52, 298-301 (2001) (Electrochemistry, Interface Phenomena, Phys. Prop., Experimental, 11) Xie, Y., Qiao, Z.Y., Mikula, A., “The Sn-In-Zn System. Application of CALPHAD Technique to Phase Diagram Measurement”, Calphad, 25(1), 3-10 (2001) (Assessment, Calculation, Phase Diagram, Thermodyn., 9) Moelans, N., Hari Kumar, K.C., Wollants, P., “Thermodynamic Optimization of the Lead-Free Solder System Bi-In-Sn-Zn”, J. Alloys Compd., 360, 98-106 (2003) (Assessment, Calculation, Phase Diagram, Thermodyn., 44) Wnuk, G., Romanowska, J., Pomianek, T., “Influence of In on the Activity of Zinc in Liquid (Tin+Zinc+Indium)”, J. Chem. Thermodyn., 35(5), 717-717 (2003) (Experimental, Thermodyn., 14)
Table 1: Investigations of the In-Sn-Zn Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1954Spe]
DTA
first sketch of liquidus surface
[1961Yok]
Knudsen cell effusion
vapor pressure of Zn, 9 alloys, xZn 0.15, 352°C
[1997Fio]
drop calorimetry
partial and integral enthalpies of mixing of liquid, 347 or 383°C, 40 measurements
[1998Xie]
DTA, X-ray, microprobe
4 isopleths, xSn/xZn = 2, 1, 0.5 and xIn = 0.1
[1998Yuj]
DTA
liquidus surface < 50 mass% Zn
[2000Anr]
drop calorimetry
partial and integral enthalpies of mixing of liquid, 440 and 634°C, 70 measurements
[2001Sab1]
DSC
6 isopleths, xIn/xSn = 0.05, 0.15, 0.50, 0.52, 0.67 or 0.85
[2003Wnu]
isopiestic vapor pressure with liquid 477°C, aZn of dilute In + Zn in liquid Sn Sn-Zn
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(In) < 156.634
tI2 I4/mmm In
(Sn) 231.9681 - 13
MSIT®
tI4 I41/amd Sn
Lattice Parameters Comments/References [pm]
a = 325.20 c = 494.70
a = 583.18 c = 318.18
dissolves 16.3 at.% Sn and 1.8 at.% Zn [2003Moe] pure In [Mas] dissolves 3.6 at.% In and 0.6 at.% Zn [2003Moe] pure Sn, at 25°C [Mas2]
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469
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(Sn) < 13
cF8 Fd3m C (diamond)
a = 648.92
(Zn) < 419.58
hP2 P63/mmc Mg
pure Sn [Mas2]
dissolves 0.7 at.% In and 0.2 at.% Sn [2003Moe] pure Zn [P]
a = 266.49 c = 494.70
, In1–xSnx 144
tI2 I4/mmm In
a = 346.3 c = 440.4
, InxSn1–x < 224
hP1 P6/mmm BiIn
a = 320.6 c = 299.7
0.19 x 0.57 25 at.% Sn, 24°C [V-C2] 0.06 x 0.24 80 at.% Sn, 10°C [V-C2]
Table 3: Invariant Equilibria Reaction
T [°C]
Type
Phase
Composition (at.%) In
Sn
Zn
L + (Sn) + (Zn)
183.3
U1
L (Sn) (Zn)
9.9 3.6 6.5 0.02
79.5 95.7 93.5 0.03
10.6 0.6 0.0 99.95
L + (In) + (Zn)
121.5
U2
L (In) (Zn)
76.1 81.8 80.7 0.06
21.0 16.3 19.3 0.01
2.9 1.9 0.0 99.93
L + + (Zn)
107.4
E1
L (Zn)
52.5 57.0 23.6 0.03
45.0 43.0 76.4 0.01
2.5 0.0 0.0 99.96
Landolt-Börnstein New Series IV/11C3
MSIT®
In-Sn-Zn
A-B-C
In-Zn
223.4 p1 l + (Sn) γ
Sn-Zn
470
MSIT®
In-Sn
198.5 e1 l (Sn) + (Zn) L + (Sn) γ + (Zn)
183.3
U1
γ + (Sn) + (Zn) 143.6 e2 l (In) + (Zn) 139.5 p2 l + (In) β 121.5
U2 L+γ+(Zn)
β + (In) + (Zn)
L+ β+(Zn)
118.0 e3 lβ+γ 107.4
L β + γ + (Zn)
Landolt-Börnstein New Series IV/11C3
β + γ + (Zn)
Fig. 1: In-Sn-Zn. Reaction scheme
E1
In–Sn–Zn
L + (In) β + (Zn)
In–Sn–Zn
471
Sn
Data / Grid: at.% Axes: at.%
Fig. 2: In-Sn-Zn. Liquidus surface projection
p1 (Sn)
e1 20
80
U1
200
175
γ 40
150
60
225 250 e3 E1
275
60
40
β
300 325 350
U2
80
20
(Zn)
p2 (In)
375°C 20
Zn
40
60
Sn Fig. 3: In-Sn-Zn. Isothermal section at 125°C
(Sn)+(Zn)
e2
80
In
Data / Grid: at.% Axes: at.%
(Sn)
(Sn)+(Zn)+γ
γ
20
80
(Zn)+γ
L+γ
40
60
L+(Zn)+γ L
60
40
β
L+(Zn)
80
20
L+(In)+(Zn)
Zn
Landolt-Börnstein New Series IV/11C3
20
40
60
(In)+(Zn) 80
(In)
In
MSIT®
In–Sn–Zn
472
Fig. 4: In-Sn-Zn. Isopleth at Sn/Zn = 2/1
300
Temperature, °C
L
L+(Zn)
L+(Sn)+(Zn) 200
L+(In) L+(In)+(Zn)
L+γ+(Zn)
100
γ+(Zn)
L+β +(Zn)
(In)
107.4°C
β +(In)+(Zn)
γ+(Sn)+(Zn)
(Sn)+(Zn)
Zn 33.30 0.00 In Sn 66.70
20
(In)+(Zn)
β +(Zn)
β +γ+(Zn) 40
60
80
In
In, at.%
400
Fig. 5: In-Sn-Zn. Isopleth at Sn/Zn = 1/1
L
Temperature, °C
300
L+(Zn) 200
L+(Sn)+(Zn)
L+γ+(Zn)
L+(In)+(Zn)
L+(In)
L+β +(Zn) 100
(Sn)+(Zn)
Zn 50.00 0.00 In Sn 50.00
MSIT®
(In)
β +(In)+(Zn)
γ+(Sn)+(Zn) β +γ+(Zn) γ+(Zn) 20
(In)+(Zn)
β +(Zn) 40
60
80
In
In, at.%
Landolt-Börnstein New Series IV/11C3
In–Sn–Zn
473
400
Fig. 6: In-Sn-Zn. Isopleth at Sn/Zn = 1/2
L
Temperature, °C
300
L+(Zn) L+(Sn)+(Zn) 200
L+(In)+(Zn)
L+(In)
L+β +(Zn) L+γ+(Zn) 100
γ+(Zn)
β +(In)+(Zn) β +γ+(Zn)
Zn 66.70 0.00 In Sn 33.30
20
(In) (In)+(Zn)
β +(Zn) 40
60
80
In
In, at.%
400
Fig. 7: In-Sn-Zn. Isopleth at constant 10 at.% In
L
Temperature, °C
300
L+(Zn)
200
L+γ
L+(Zn)+γ
(In)+(Zn) 100
(Zn)+γ
(Zn)+β (Zn)+γ+β
Zn 90.00 In 10.00 Sn 0.00
Landolt-Börnstein New Series IV/11C3
20
(Zn)+(Sn)+γ 40
Sn, at.%
60
80
Zn 0.00 In 10.00 Sn 90.00
MSIT®
In–Sn–Zn
474
Fig. 8: In-Sn-Zn. Integral and Zn-partial enthalpy of mixing of liquid at 445°C
10.0
8.0
Hliq, kJ.mol-1
ΔHliqZn 6.0
4.0
2.0
ΔmixHliq
0
In 50.00 Sn 50.00
MSIT®
0.2
0.4
0.6
0.8
Zn
Mole fraction Zn
Landolt-Börnstein New Series IV/11C3
Mn–Ti–Zr
475
Manganese – Titanium – Zirconium Volodymyr Ivanchenko Introduction The Mn-Ti-Zr system is of high importance for a number of practical applications. Firstly, Zr and Mn are often used as alloying elements in the titanium industries. Secondly, alloys of this system are considered to be promising candidates for reversible hydrogen storage [1978Oes, 2000Bob] as well as for hydrogen storage tanks in proton exchange membrane fuel cell technology [2004Tai]. Thirdly, promising candidate solders may emerge from that system. Eutectics reactions with relatively low melting temperatures in the Mn-Ti and Ti-Zr systems plus a minimum in the fusibility diagram of the Ti-Zr system have draw the attention to this ternary system in developing new solders. Possible applications are reaction brazing of Ti-based alloys and, especially, titanium aluminide based materials [2004Kho]. The knowledge of the constitution of the Mn-Ti-Zr system therefore is important for developing new advanced materials. Nevertheless, the information on the equilibrium phases in this system is restricted to a few publications only. [1980Jac, 1987Bla] showed the formation of continuous solid solutions between isostructural TiMn2 and ZrMn2 compounds and [2006Iva1, 2006Iva2] studied the phase equilibria in composition range Ti-TiMn2-ZrMn2-Zr. The same authors constructed experimentally the solidus surface in the Ti rich part of the system, projections of solidus and liquidus surfaces as well as some polythermal sections [2006Iva1, 2006Iva2]. A section with metastable phases formed in as cast brazing alloys was presented by [2004Kho]. The hydrogen sorption and desorption properties of (Ti1–xZrx)Mn2–y base materials have been studied by [1978Oes, 1980Jac, 2000Bob, 2004Tai]. The investigations of phase relations, structures and thermodynamics of the Mn-Ti-Zr alloys are summarized in Table 1. Binary Systems The Ti-Zr binary systems is accepted as published by [Mas2] as well as the Mn-Ti binary system, Fig. 1. An amendment has been applied to the temperature of the peritectic formation of TiMn which was definied to be 1225°C by [2004Iva] in comparison with value ~1200°C presented by [Mas2]. The Mn-Zr system is accepted from [1999Sch] where the homogenity region of ZrMn2 is determined more accurately then in [Mas2]. The temperature of the eutectic reaction L (Zr)+ZrMn2 however has been corrected to 1125 7°C, Fig. 2. This is 35°C higher then presented by [1999Sch] and results from the experiments on the binary and ternary alloys performed by the author of the present evaluation. The liquidus temperature at 50 at.% Mn - 50 at.% Zr was measured as 1420°C. In the light of the results of [1999Sch, 2006Iva2] the present evaluation does not assume that an MnZr intermetallic compound exist. Solid Phases The known solid phases are listed in Table 2 with continious solid solutions being formed between TiMn2 and MnZr2 intermetallic compounds. In the Ti-Zr binary system the (Zr) and (Ti) form a continuous solid solution () which can dissolve a limited amount of Mn (1-2 at.%). The (Zr) and (Ti) also form a series of continious solid solutions () dissolving Mn from 10 at.% at the Zr reach side to 30 at.% at Ti rich part. In the ternary system the TiMn phase dissolves up to 5 at.% Zr. The TiMn phase in binary alloys was detected in as cast conditions only after dead melting during 30 min in the solid-liquid state at a temperature of 1190°C, which is a little higher than the eutectic temperature (1175°C) but lower than the temperature of the peritectic reaction (1225°C) [2004Iva]. The same situation with the TiMn phase formation was detected in the ternary alloys study [2006Iva2]. A phase isotypic with QHfRe had also been found in the system Mn-Ti at approximately 50.15 at.% Mn (TiMn) [1962Wat]. In analogy with the cell reported for QHfRe , the TiMn phase has also been tentatively indexed with a tetragonal cell with a = 819 pm and c = 1281 pm. As it was later shown by [1986Cen], it can be indexed with a new Zr21Re25 type rhombohedral
Landolt-Börnstein New Series IV/11C3
MSIT®
476
Mn–Ti–Zr
structure containing 1200 pm-thick infinite MgZn2 (Laves)-type columns. [1986Cen] expected that the still unknown structure of TiMn, which appears at an intermediate composition between Ti21Mn25 and TiMn2, also contains Friauf polyhedra [1986Cen]. The new intermetallic phase is described here as metastable in accordance with [1999Sch], although it may in fact be part of the equilibrium diagram. The tentatively assigned stoichiometry is Zr21Mn25 [1999Sch], by analogy with a similar phase in the Mn-Ti system [1986Cen]. Quasibinary Systems According to the study of [2006Iva2] there must be two quasibinary sections in the Mn-Ti-Zr system. One of them, TiMn2-ZrMn2, is characterized by continuous solid solutions between these Laves phases with the maximum of melting point at 2 at.% Ti. The phase diagram of this quasibinary system is presented in Fig. 3. The second one may be presented by the temperature-composition section oriented along the tie-line connecting the solid solution with the (Zr,Ti)Mn2 solid solution. The end points of this section are for (Zr,Ti)Mn2: the point “max” at 1470°C in Fig. 6 and for the phase: the minimum point at 1040°C on the solidus line in Fig. 5. Invariant Equilibria The data on invariant ternary reactions are tabulated according to [2006Iva2] in Table 3. Liquidus, Solidus and Solvus Surfaces The solidus surface in the Ti corner of the phase diagram was constructed by [2006Iva1] based on DTA measurements. Due to the absence of ternary phases, this part of the solidus surface should be smooth and may be described by a power function as: TS = 1670 – 29.053 xMn + 0.4246 xMn2 – 5.692 xZr + 0.0566 xZr2 + 0.6478 xZr xMn – 0.0307 xZr2 xMn – 0.0065 xZrxMn2. The calculated isothermal lines at temperatures of 1175-1540°C are presented in Fig. 4. The solidus surface and liquidus surface projections of the partial Ti-TMn2-ZrMn2-Zr system are presented in Fig. 5 and Fig. 6, respectively, according to [2006Iva2]. Temperature – Composition Sections Two vertical sections along the Ti48Mn52-Zr50Mn50 and Ti61Mn39-Zr67.5Mn32.5 lines reflect the smooth transition from the quasibinary TiMn2-ZrMn2 section with maximum on the fusibility diagram to the binary Ti-Zr system with minimum on the solidus and liquidus curves. They are shown in Fig. 7 and Fig. 8. As it was noted by [2004Iva], cooling with the rate of 0.7 K#s–1 did not lead to the TiMn phase formation on eutectic crystallization in the binary system and this situation is retained in ternary alloys. The phase composition realised in as cast alloys cooled with the rate above 0.1 K#s–1 is presented in Fig. 9 in accordance with [2004Kho], where phase equilibria involving the TiMn intermetallic compound are missing. Notes on Materials Properties and Applications The experimental investigations of materials properties are summarised in Table 4. Microhardness of the (ZrxTi1–x)Mn2 solid solutions lineary varies from 5817 MPa for TiMn2 to 7750 MPa for ZrMn2 [1987Bla]. Butt brased joints in TiAl (47XD) were made by the vacuum brazing method, using adhesive active filler metals of the Mn-Ti-Zr system. Short-time strength of the butt brazed joints is on the level of ~690 MPa at 20°C and ~310 MPa at 700°C [2004Kho]. The maximum absorption capacity NH of the (ZrxTi1–x)Mn2 system steeply increased from NH = 0.3 H atoms (f.u.)–1 for x = 0 to NH = 3.65 H atoms (f.u.)–1 for x = 0.4; it stays practivally constant at approximately 4 H atoms (f.u.)–1 for 0.6 < x < 1 [1978Oes, 1980Jac]. The sorption-desorption characteristics are strongly dependent on the Mn content in the intermetallic compound [2000Bob, 2004Tai].
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Landolt-Börnstein New Series IV/11C3
Mn–Ti–Zr
477
References [1961Wat]
[1962Wat]
[1976Sve]
[1978Oes] [1980Jac]
[1981Mur1] [1981Mur2] [1986Cen]
[1987Bla]
[1997Kod]
[1999Sch] [2000Bob]
[2002Mit]
[2004Iva]
[2004Kho]
[2004Tai]
[2006Iva1]
Landolt-Börnstein New Series IV/11C3
Waterstrat, R.M., “Identification of Intermediate Phase in the Manganese-Titanium System”, Trans. Metall. Soc. AIME, 221(8), 686-690 (1961) (Experimental, Phase Relations, 11) Waterstrat, R.M., Das, B.N., Beck, P.A., “Phase Relationships in the Titanium-Manganese”, Trans.Metall. Soc. AIME, 224(6), 512-518 (1962) (Crys. Structure, Experimental, Phase Relations, 17) Svechnikov, V.N, Petkov, V.V., “Laves Phase in Alloys of Mn with Transition Metals of Groups IVA-VA” (in Russian), Izv. Akad. Nauk Ukr. SSR, Metallofizika, 64, 24-29 (1976) (Crys. Structure, Experimental, Phase Relations, 7) Oesterreicher, H., Bittner, H., “Studies of Hydride Formation in Ti1–xZrxMn2”, Mat. Res. Bull., 13, 83-88 (1978) (Crys. Structure, Experimental, Calculation, Thermodyn., 13) Jacob, I., Stern, A., Moran, A., Shaltiel, D., Davidov, D., “Hydrogen Absorption in (ZrxTi1–x)B2 (BCr, Mn) and the Phenomenological Model for the Absorption Capacity in Pseudo-Binary Laves-Phase Compounds”, J. Less-Common Met., 73, 369-376 (1980) (Crys. Structure, Experimental, Morphology, 13) Murray, J.L., “The Ti-Zr (Titanium - Zirconium) System”, Bull. Alloy Phase Diagrams, 2(2), 197-201 (1981) (Crys. Structure, Phase Diagram, Review, 31). Murray, J.L., “The Mn-Ti (Manganese – Titanium) System”, Bull. Alloy Phase Diagrams, 2(3), 334-343 (1981) (Crys. Structure, Phase Diagram, Review, 50). Cenzual, K., Parthe, E., Waterstrat, R.M., “Zr21Re25, a New Rhombohedral Structure Type Containing 12 Å-Thick Infinite MgZn2 (Laves)-Type Columns”, Acta Crystallogr., C42, 261-266 (1986) (Crys. Structure, Experimental, Phase Relations, 24) Blazina, Z., Trojko, R., “On Friauf-Laves Phases in the Zr1–xAlxT2, Zr1–xSixT2 and Zr1–xTixT2 (T = Mn, Fe, Co) Systems”, J. Less-Common Met., 133(2), 277-286 (1987) (Crys. Structure, Experimental, 10) Kodama, T., “An Attempt to Estimate the Extent of the Single Phase Region of the ZrMn2 Phase by Means of the X-Ray Diffraction-Profile Halfwidths”, J. Alloys Compd., 256, 263-268 (1997) (Crys. Structure, Experimental, Phase Relations, 24) Schiesinger, M.E., “The Mn-Zr (Manganese-Zirconium) System”, J. Phase Equilib., 20(1), 79-83 (1999) (Crys. Structure, Phase Diagram, Phase Relations, Review, 42). Bobet, J.-L., Chevalier, B., Darriet, B., “Crystallographic and Hydrogen Sorption Properties of TiMn2 Based Alloys”, Intermetallics, 8, 359-363 (2000) (Crys. Structure, Experimental, Phase Relations, 16) Mitrokhin, S.V., Bezuglaya, T.N., Verbetsky, V.N., “Structure and Hydrogen Sorption Properties of (Ti,Zr)-Mn-V Alloys”, J. Alloys Compd., 330-332, 146-151 (2002) (Crys. Structure, Experimental, Phase Relations, 12) Ivanchenko, V.G., Gavrylenko, I.S., Pogorelaya, V.V., Nychyporenko, V.I., Pryadko, T.V., “Investigation of Phase Equilibria in Alloys of the Ti-Mn System.” (in Ukrainian), Metallurgy and Processing of Metals (Metaloznavstvo ta Obrobka Metaliv), 4, 16-20 (2004) (Crys. Structure, Experimental, Phase Relations, 14). Khorunov, V.F., Maksymova, S.V., Ivanchenko, V.G., “Obtaining of Titanium Aluminide Brazed Joints and their Properties Investigation“ (in Russian), Adhesion of Melts and Brazing of Materials (Adgeziya Rasplavov i Payka Metallov), 37, 88-95 (2004) (Experimental, Phase Relations, Phys. Prop., 9) Taizhong, H., Zhu, W., Xuebin, Y., Jinzhou, C., Baojia, X., Tiesheng, H., Naixin, X., “Hydrogen Absorption-Desorption Behavior of Zirconium-Substituting Ti-Mn Based Hydrogen Storage Alloys”, Intermetallics, 12, 91-96 (2004) (Crys. Structure, Experimental, Calculation, Thermodyn., 11) Ivanchenko, V.G., Nichiporenko, V.I., Pryadko, T.V., “The Determination of Solidus Surface in Titanic Angle of Diagram tates of System Ti-Zr-Mn” (in Russian), Metal Phys. MSIT®
Mn–Ti–Zr
478
[2006Iva2]
Adv. Technol., (to be published), 27 (2006) (Experimental, Calculation, Phase Diagram, Phase Relations, 2) Ivanchenko, V.G., Pryadko, T.V., Gavrylenko, I.S., Pogorelaya, V.V., “Phase Equilibria in the Ti – TiMn2 – ZrMn2 – Zr Partial System”, J. Alloys Compd., (to be published) (2006) (Crys. Structure, Experimental, Calculation, Phase Diagram, Phase Relations, 10)
Table 1: Investigations of the Mn-Ti-Zr Phase Relations, Structures and Thermodynamics Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1978Oes]
X-ray diffraction, Sieverts method
room temperature, ZrxTi1–xMn2, x = 0; 0.2; 0.4; 0.6; 0.8; 1.0, C14 Laves phase
[1980Jac]
X-ray diffraction, hydrogen absorption capacity
room temperature, ZrxTi1–xMn2, x = 0; 0.2; 0.4; 0.6; 0.8; 1.0, C14 Laves phase
[1987Bla]
X-ray diffraction, microhardness
annealed at 800 to 1500°C for 24 h, quenched at rate 100°C#h–1, Zr1–xTixMn2, x = 0; 0.2; 0.4; 0.6; 0.8; 1.0, hexagonal MgZn2
[2000Bob]
X-ray diffraction, electron microprobe analysis
annealed at 800°C for 7 days, Zr0.05Ti0.95Mn2–x x = –0.05; 0; 0.05; 0.15; 0.35, C-14 Laves phase
[2004Kho]
X-ray diffraction, DTA investigation, light optical microscopy, electron microprobe analysis, mechanical properties
annealed at 1250°C, quenched at rate >0.1°C#s–1, metastable polythermal section of Ti60Mn40 - Zr67.5Mn32.5
[2004Tai]
X-ray diffraction, induction coupled plasma spectroscopy, entropy and enthalpy of hydrogen desorption
annealed at 960°C for 6 h, room temperature, ZrxTi1–xMn2, x = 0; 0.1; 0.2; 0.3; 0.4, C14 Laves phase
[2006Iva1]
Optical light microscopy, DTA investigation
as cast and annealed 900°C for 35 h, Ti-(Ti-25 at.%Mn)-(Ti-20 at.% Zr-10 at.% Mn), (Ti), +(Zr,Ti)Mn2
[2006Iva2]
X-ray diffraction, DTA investigation, light optical microscopy, electron microprobe analysis
annealed at 950°C for 1 week, water quenching. Ti-TiMn2-ZrMn2-Zr (Zr,Ti), (Zr,Ti)Mn2, TiMn
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Landolt-Börnstein New Series IV/11C3
Mn–Ti–Zr
479
Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
( Mn) 1246 - 1138
cI2 Im3m W
a = 308.0
pure Mn [Mas2]
( Mn) 1138 - 1100
cF4 Fm3m Cu
a = 386.0
pure Mn [Mas2]
(Mn) 1100 - 727
cP20 P4132 Mn
a = 631.52
pure Mn [Mas2]
(Mn) < 727
cI58 I43m Mn
a = 891.26
pure Mn at 25°C [Mas2]
, Zr1–x–yTiyMnx
cI2 Im3m W
a = 337.7 a = 343.7 a = 349.7
0 < x < 0.3, 0 < y < 1, x = 0, y = 0.3 x = 0, y = 0.5, x = 0, y = 0.7 [1981Mur1]
(Zr) 1855 - 863
a = 360.90
pure Zr, [Mas2]
(Ti) 1670 - 882
a = 330.65
pure Ti, [Mas2]
, (Zr,Ti)
hP2 P63/mmc Mg
a = 295.1 to 323.2 0-100 at.% Zr, at 25°C c = 468.4 to 515.0 [1981Mur1]
(Zr) < 863
a = 323.16 c = 514.75
pure Zr at 25°C, [Mas2]
(Ti) < 882
a = 295.06 c = 468.35 a = 295.11 c = 468.43
pure Zr, [Mas2] 0-0.4 at.% Mn [P]
TiMn < 1200
-
-
52 at.% Mn [1962Wat] 48-52 at.% Mn, [2004Iva]
TiMn < 950
t*58
a = 819 c = 1281
at 50.15 at.% Mn [Mas2] [1962Wat]
TiMn3 < 1250
o** P212121
a = 481.2 c = 789.5
75.5 at.% Mn, at 1000°C [1981Mur2]
TiMn4 1230 - 930
hR159 R3m Co5Cr2Mo3
a = 1100.03 c = 1944.6
[1961Wat] [VC2] Temperature interval from [1962Wat]
Landolt-Börnstein New Series IV/11C3
MSIT®
Mn–Ti–Zr
480 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
(ZrxTi1–x)Mn2y
hP12 P63/mmc MgZn2
a = 482.1 to 504.1 0 < x < 1, y = 0, [1987Bla] c = 791.4 to 824.9 a = 483.3 + 38.6 x 0 < x < 1, y = 0, [2006Iva2] – 19 x2 0.2 c = 792.3 63.3 x – 31.2 x2 0.4
ZrMn2 < 1450
a = 505.2 to 498.5 c = 831.5 to 819.0 a = 504.5 c = 828.0 a = 498.0 c = 817.6
TiMn2 < 1325
a = 485.7 to 481.8 64-70 at.% Mn, at 1310°C [1976Sve] c = 801.7 to 87.9 hR276 R3c Zr21Re25
Zr21Mn25 metastable
-
60-72 at.% Mn, at 1510°C, [1976Sve], 64.3 at.% Mn, (Zr) - saturated, 77.3 at.% Mn, (Mn) - saturated, [1999Sch]
[1997Kod] 54.4 at.% Mn [1999Sch]
Table 3: Invariant Equilibria T (°C)
Reaction
Type
Phase
Composition (at.%) Mn
Ti
Zr
l (Zr,Ti)Mn2
1470
max
l (Ti,Zr)Mn2
~ 67 ~ 67
~2 ~2
~ 31 ~ 31
l + TiMn + (Zr,Ti)Mn2
1090
U
l TiMn (Ti,Zr)Mn2
~ 22 ~ 50 ~ 27 ~ 57
~ 51 ~ 45 ~ 69 ~ 23
~ 27 ~5 ~4 ~ 20
l + (Zr,Ti)Mn2
1040
e (min)
l (Ti,Zr)Mn2
~ 21 ~ 15 ~ 57
~ 47 ~ 54 ~8
~ 32 ~ 31 ~ 35
Table 4: Investigation of the Mn-Ti-Zr Materials Properties Reference
Method/Experimental Technique
Type of Property
[1978Oes]
Sieverts method
Desorption isotherms of hydrogen
[1987Bla]
Vickers microhardness measurement
Microhardness under a load of 100 g
[2000Bob]
Sieverts method
Hydrogen sorption capacity
[2002Mit]
Sieverts method
Hydrogen sorption capacity
[2004Tai]
A gas reaction controller
Sorption-desorption isotherms
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Landolt-Börnstein New Series IV/11C3
Mn–Ti–Zr
Fig. 1: Mn-Ti-Zr. Phase diagram of the Mn-Ti binary system
1750
L 1500
L+TiMn2 L+(ρ TiMn)
L+(β Ti)
Temperature, °C
481
1250
~30
1175
β TiMn+TiMn2
1246°C 1138°C (γMn) 1100°C (β Mn)
~930 TiMn2
727°C (β Ti)+α TiMn
500
TiMn4
~950
750
1148
TiMn3
40
(β Ti)+β TiMn (α Ti)+(β Ti)
(δ Mn)
1250
1225 (β Ti)
1000
1330
(α Mn)
550
(α Ti) (α Ti)+α TiMn
α TiMn+β TiMn
250 20
Ti
40
60
80
Mn
Mn, at.%
2000
Fig. 2: Mn-Ti-Zr. Phase diagram of the Mn-Zr binary system
1855°C 1750
Temperature, °C
1500
1450
1250
ZrMn2
1125
(β Zr)
1160 1140
1000
863°C
(δ Mn)
1246°C
(γMn)
1138°C 1100°C
(β Mn)
790 730
750
500
(α Zr)
(α Mn)
250
0
Zr
20
40
60
80
Mn
Mn, at.%
Landolt-Börnstein New Series IV/11C3
MSIT®
Mn–Ti–Zr
482
1600
Fig. 3: Mn-Ti-Zr. Quasibinary section Zr33.3Mn66.7 Ti33.3Mn66.7
1500
L
1470
Temperature, °C
1450°C 1400
1330°C
L+(Zr,Ti)Mn2 1300
(Zr,Ti)Mn2 1200
1100
Zr 33.30 0.00 Ti Mn 66.70
10
20
0.00 Zr Ti 33.30 Mn 66.70
30
Ti, at.%
Zr 0.00 Ti 50.00 Mn 50.00
Fig. 4: Mn-Ti-Zr. Solidus isotherms at 1180°C (1), 1210°C (2), 1280°C (3), 1340°C (4), 1410°C (5), 1540°C (6). Dashed line shows the projection of maximal solubility line of the solid solution
Data / Grid: at.% Axes: at.%
10
40
20
30
1 2
30
20
3 4 40
10
5 6
β Zr 50.00 Ti 50.00 Mn 0.00
MSIT®
60
70
80
90
Ti
Landolt-Börnstein New Series IV/11C3
Mn–Ti–Zr
483
Mn
Data / Grid: at.% Axes: at.%
Fig. 5: Mn-Ti-Zr. Solidus surface of the Ti - TiMn2 - ZrMn2 Zr partial system 20
80
(Zr,Ti)Mn2
40
60
MnTi+Mn2Ti MnTi
1090
β +MnTi
60
40
1100 1040 80
1055
1080
20
β +(Zr,Ti)Mn2 β +TiMn+(Zr,Ti)Mn2 β 20
Zr
40
60
80
Mn
Ti
Data / Grid: at.% Axes: at.%
Fig. 6: Mn-Ti-Zr. Liquidus surface of the Ti - TiMn2 ZrMn2 - Zr partial system
20
80
max
40
60
(Zr,Ti)Mn2 p, 1225°C 60
40
e, 1175°C
1125°C, e U 80
Landolt-Börnstein New Series IV/11C3
20
20
e1 (min) 1040
β
Zr
β TiMn
40
60
80
Ti
MSIT®
Mn–Ti–Zr
484
Temperature, °C
Fig. 7: Mn-Ti-Zr. Vertical section Zr50Mn50 -Ti48Mn52
1500
L
L+(Zr,Ti)Mn2 1250
L+β +(Zr,Ti)Mn2 L+TiMn+(Zr,Ti)Mn2 TiMn 1000
β +(Ti,Zr)Mn2 Zr 50.00 0.00 Ti Mn 50.00
Fig. 8: Mn-Ti-Zr. Vertical section Zr67.5Mn32.5 Ti61Mn39
β +TiMn+(Zr,Ti)Mn2 10
20
30
0.00 Zr Ti 48.00 Mn 52.00
40
Ti, at.%
1300
L
1200
Temperature, °C
L+(Zr,Ti)Mn2
L+TiMn+(Zr,Ti)Mn2
L+TiMn L+β +TiMn
1125°C 1100
1090
β +TiMn
L+β +TiMn2 L+β +(Zr,Ti)Mn2 1000
β +TiMn+(Zr,Ti)Mn2
β +(Zr,Ti)Mn2
900
Zr 67.50 0.00 Ti Mn 32.50
MSIT®
20
40
Ti, at.%
0.00 Ti 61.00 Mn 39.00
60 Zr
Landolt-Börnstein New Series IV/11C3
Mn–Ti–Zr
Temperature, °C
Fig. 9: Mn-Ti-Zr. Metastable Zr67.5Mn32.5 Ti60Mn40 section under the cooling rate above 0.1 K#s–1
485
1300
L
1200
L+(Zr,Ti)Mn2 L+β +(Zr,Ti)Mn2
1100
1000
β +(Zr,Ti)Mn2
900
Zr 67.50 0.00 Ti Mn 32.50
Landolt-Börnstein New Series IV/11C3
20
40
Ti, at.%
0.00 Zr Ti 60.00 Mn 40.00
MSIT®
486
Ni–Pd–Si
Nickel – Palladium – Silicon Elena L. Semenova Introduction Data on the Ni-Pd-Si phase equilibria, obtained from the investigation of bulk alloys, are restricted to [1976Wop] and [2004Loo]. [1976Wop] studied the ternary alloys of about 60 compositions in the region of 20-75 at.% Si with powder diffraction method using Guinier camera. The alloys were prepared from components of 99.9% purity by induction melting under argon atmosphere (those with less than 40 at.% Si) or in arc furnace (those with more than 40 at.% Si). Their homogenization at 980-700°C for 16 h - 4 d was followed by water quenching. Isothermal section at 800°C was constructed. Two ternary phases were found to exist - NiPd2Si and Ni18Pd7Si9. They are in equilibrium and they take part in a number of equilibria with some binary phases. The powder diffraction pattern of the NiPd2Si phase was not analyzed (it was supposed that the crystal structure of the phase is similar to that of Pd3Si, of Fe3C type) whereas the crystal structure of the Ni18Pd7Si9 phase was determined as of Pd25Ge9 type. Equiatomic intermetallics, NiSi and PdSi, were shown to be completely dissolved while a miscibility gap was observed at this temperature between Ni2Si and Pd2Si phases. The isothermal section at 800°C demonstrates an equilibrium of the monosilicide solid solution, NiSi2 phase and (Si); a vertex of the tie-triangle for the first phase is placed at about 12 at.% Pd. An evidence of this value was not given in [1976Wop] as well as of other concentration points. Some recent experimental works on the phase equilibria of the Ni-Pd-Si system were carried out in order to provide a necessary basis for understanding of the processes connected with destabilizing of NiSi2 in ternary system [2004Loo, 2003Wan]. The processes are important for technology of the Ni-Pd-Si films. [2004Loo] studied the Ni-Pd-Si alloys in the region of 50 to 100 at.% Si with the methods presented in Table 1. The concentration of Pd in the monosilicide phase that can effect the decomposition of the disilicide phase was determined. It was shown that this quantity increases with the rising of temperature. The thermal stability of NiSi phase in the presence of Pd, intercalated between Ni and Si in the film was investigated in [2003Wan]. Using of the Pd layer contributed in enhancing of the thermal stability of the monosilicide phase. A number of studies were devoted to the investigation of the Ni-Pd-Si amorphous alloys as prospective basis for development of modern amorphous materials [1990Rab, 1988Ino, 1985Ino, 1966Tsu] that can be used in production of solders. They searched for the regions where the ternary alloys could be amorphizied easily, studied the techniques of the preparing of glasses, processes of their crystallization and the properties of the amorphous samples. In the study of the glass forming ability characteristics of the Ni-Pd-Si amorphous alloys by the methods pointed out in Table 1, [1990Rab] determined some constitutional data on rapidly solidified samples such as solidus and liquidus temperatures as well as the temperature of glass transition and crystallization for some compositions. The influence of hydrogen on physical and chemical properties of the amorphous alloys was discussed by [1984Fil, 1991Yos]. Binary Systems The binary phase diagrams of Ni-Pd and Ni-Si [1991Nas1, 1991Nas2] as well as Pd-Si [1993Oka] were accepted. Solid Phases [1993Mas] using DTA, EMPA and XRD methods analyses showed many new phases in the Pd-Si system. The determination of their crystal structures was not carried out. [1993Oka] questioned the existence of some phases, especially four of them, related to Pd2Si. Data on solid phases are given in Table 2 including all the phases listed in [1993Oka].
MSIT®
Landolt-Börnstein New Series IV/11C3
Ni–Pd–Si
487
Ni2Si dissolves nickel up to about 50 at.%. Both lattice parameters of the hexagonal solid solution decrease appreciably with increasing of nickel content along the section 33 at.% Si. On extending of the NiSi based solid solution into the Ni-Pd-Si ternary system, its lattice parameters measured at 800 and 700°C increase (Table 2). Quasibinary Systems Based on the new version of the Pd-Si system [1993Oka], where the PdSi monosilicide exists in the very restricted temperature interval, 908 - 888°C, it can be accepted that complete solubility of the PdSi and NiSi phases exists in the narrow interval at subsolidus area. Here the section between monosilicides will be quasibinary. Liquidus and Solidus Surfaces According to DTA data of [1990Rab] obtained on amorphous samples of the Ni-Pd-Si system the liquidus and solidus temperature of the samples belonging to the section between the alloys Ni-85Pd-15Si (at.%) and 78Ni-Pd-22Si (at.%) drops with increased nickel content in the alloys, that are close to the Pd-Si side of the ternary phase diagram (5-10 at.% Ni). Isothermal Sections To make an agreement with the last version of the Pd-Si phase diagram [1993Oka], showing a lack of the PdSi phase at 800°C, the isothermal section constructed by [1976Wop] had to be corrected. Taking into account that the ternary alloy with 20Ni - 12Pd - 68Si (at.%) is a two-phase one, (Ni,Pd)Si + (Si), at 800°C [1976Wop], it can be assumed that the monosilicide solid solution exists at this temperature at least up to this alloy and the composition of the monosilicide that is in equilibrium with Pd2Si and Si phases should be somewhat richer in Pd. An arrangement of other three-phase fields at the isothermal section at 800°C is given in Fig. 1 according to [1976Wop]. Mentioned in [1993Mas, 1993Oka] a group of phases close to the hexagonal Pd2Si phase is not shown here. A series of three phase equilibria observed by [1976Wop] is the following: Ni18Pd7Si9 + 1 + ; + + Pd2Si; Pd2Si + + Ni18Pd7Si9; Pd2Si + Pd3Si + NiPd2Si; + Pd2Si + J; Pd2Si +J +(Ni,Pd)Si; Pd2Si + NiPd2Si + Ni18Pd7Si9; (Ni,Pd)Si + NiSi2 + (Si). Considering that the compositions of the (Ni,Pd) solid solution in equilibria with the Pd3Si and NiPd2Si, NiPd2Si and Ni18Pd7Si9; Ni18Pd7Si9 and Ni3Si phases are unknown, the equilibria are shown tentatively by three triangles. Temperature – Composition Sections Figure 2 demonstrates the vertical section passing through binary alloys 85Pd-15Si and 78Ni-22Si (at.%) in the temperature range of the alloy crystallization. The alloys of the section in the composition range of 5-10 at.% Ni have the lowest melting temperature and the narrowest melting interval [1990Rab]. Below 888°C, the temperature of eutectoid decomposition of the PdSi phase, PdSi (Si) + Pd2Si, the section through the monosilicides cannot be quasibinary one because the compositions of conjunctive phases, Pd2Si and (Si), and their tie lines are off the plane of the section. Since the (Ni,Pd)Si phase was observed at 800°C in the alloy with 20 at.% Pd [1976Wop] the decomposition temperature of it should be lower than 800°C. With considering this the section through the monosilicides is built as a general guide (Fig 3). Notes on Materials Properties and Applications The amorphous (Ni,Pd)82Si18 alloys with 29-82 at.% Pd revealed a high ability to form wire of circular cross section, as well as high strength combined with good ductility, and good corrosion resistance and can be used in practice as the tip material in various kinds of brushes and pens [1985Ino]. [1988Ino] managed to prepare the amorphous spherical particles with diameters of 0.5 to 2 mm by injection of molten Ni-Pd-Si alloy containing 1 at.% B into stirred cold water. So large-scale amorphous balls are important for various Landolt-Börnstein New Series IV/11C3
MSIT®
488
Ni–Pd–Si
practical use including the soldering balls for a connection, for which a high dimensional accuracy is required, while in a ball-point pen the balls in bearing as well. The hardness, crystallization parameters for large particles found to retain practically the same as for the thick ribbons. A number of the ternary alloys compositions containing 22 at.% Si may be used as materials for step brazing because their melting range fits precisely any particular temperature interval between 770 and 1000°C. Vacuum tube brazing is the first successful application where this advantage of Ni-Pd base alloys has been used [1990Rab]. Due to oxidation resistance, ductility, relatively high melting points the Ni-Pd based alloys were recommended to use in brazing ceramics (SiN) to metals [1992Sel]. The Ni-Pd-Si system is considered as a prospective base for production of materials for integrated circuits [2004Loo, 2003Wan, 2000Tsa]. [2004Loo] found that an addition of Pd as an interlayer enhanced thermal stability of the Ni-Si system, where Ni is a film and Si - substrate. [2000Tsa] studied the processes occurring in the Ni-Pd-Si films, where Pd is an interlayer, by measuring the total force per unit width of the film during isochronal annealing. It was shown, that introduction of Pd layer retarded the formation of Ni2Si and NiSi phases due to the increase of the formation temperature. Miscellaneous The resistance and thermoelectric power of the Ni-Pd-Si amorphous hydrogenized alloys were increased monotonously towards positive value with increasing the hydrogen pressure. It points out that sites accessible for hydrogen exist even at hydrogen pressure of about 2 GPa. Superconductivity at temperatures down to 1.5 K in the Ni-Pd-Si alloys was absent [1984Fil]. The solubility of hydrogen in the amorphous ternary 8.8Ni - Pd - 18Si alloy is decreased in comparison with the Pd82Si18 amorphous binary alloy [1991Yos]. The crystallization process in the amorphous Ni-Pd-Si alloys with 10-30 at.% Ni and 50-70 at.% Pd is greatly influenced by nickel: the onset time of crystallization retarded on increasing of nickel content [1980Uts]. Amorphous structure was obtained in the alloy 15Ni-65Pd-20Si (at.%) by quenching from the liquid using “piston and anvil” technique [1966Tsu]. [1978Chu] studied angular correlations of both electron irradiated and quenched amorphous alloy with 6Ni-77.5Pd-16.5Si (at.%) and concluded that the electron irradiation does not induce vacancies in the material. References [1953And] [1966Tsu] [1976Wop] [1978Chu]
[1980Uts] [1984Fil]
[1985Ino]
MSIT®
Anderko, K., Schubert, K., (in German), Z. Metallkd., 44, 307-312 (1953) (Crys. Structure, Experimental) Tsuei, C.C., Duwez, P., “Metastable Amorphous Ferromagnetic Phases in Palladium-Base Alloys”, J. Appl. Phys., 37, 435 (1966) (Experimental, 6) Wopersnow, W., Schubert, K., “On the Mixture Nickel-Palladium-Silicon” (in German), Z. Metallkd., 67(12), 807-810 (1976) (Experimental, Crys. Structure, Phase Diagram, 13) Chuang, S.Y., Tseng, P.K., Jan, G.J., Chen, H.S., “An Experimental Study of the Angular Correlation of Positron Annihilation Radiations in Electron Irradiated Pd-Ni-Si Metallic Glass”, Phys. Status Solidi A, 48, K181-K183 (1978) (Experimental, 12) Utsumi, K., Kawamura, K., “Studies on Crystallization Processes of Amorphous Pd-Ni-Si Alloys”, Trans. Jpn. Inst. Met., 21(5), 269-274 (1980) (Experimental, 15) Filipek, S.M., Szafranski, A.W., Duhaj, P., “Some Properties of Amorphous Pd-Si-and Pd-Ni-Si Alloys under Hydrogen Pressures of up to 2 GPa”, J. Less-Common Met., 101, 299-304 (1984) (Experimental, 20) Inoue, A., Masumoto, Y., Yano, N., Kawashima, A., Hashimoto, K., Masumoto, T., “Production of Ni-Pd-Si and Ni-Pd-P Amorphous Wires and their Mechanical and Corrosion Properties”, J. Mater. Sci., 20, 97-104 (1985) (Experimental, 17)
Landolt-Börnstein New Series IV/11C3
Ni–Pd–Si [1988Ino]
[1990Rab]
[1991Yos]
[1991Nas1]
[1991Nas2]
[1992Sel]
[1993Mas] [1993Oka] [2000Tsa] [2003Wan]
[2004Loo]
489
Inoue, A., Ekimoto, T., Masumoto, T., “Production of Large Amorphous Spheres by Melt Ejection into Stirred Water of Pd-Ni-P-B, Pd-Ni-Si-B and Pd-Cu-Si-B Alloys and their Thermal Stability”, Int. J. Rapid Solidif., 3(3), 237-252 (1988) (Experimental, 13) Rabinkin, A., “Melting Characteristics, Glass-Forming Ability and Stability of Ni-Pd-Si and Ni-Pd-TM-Si Alloys in the Concentration Range 10-22 at.% Si”, Mater. Sci. Eng. A, A124(2), 251-258 (1990) (Experimental, Phase Relations, 18) Yoshinary, O., Kircheheim, R., “Solubility and Diffusivity of Hydrogen in Amorphous Pd73.2X8.8Si18 (X = Ag, Cu, Cr, Fe, Ni) Alloys”, J. Less-Common Met., 172-174, 890-898 (1991) (Experimental, 7) Nash, P., “Ni-Pd (Nickel-Palladium)” in “Phase Diagrams of Binary Nickel Alloys”, Nash, P. (Ed.), ASM, Metals Park, OH, 247-256 (1991) (Review, Phase Diagram, Phase Relations, #, 36) Nash, P., Nash, A., “Ni-Si (Nickel-Silicon)”, in “Phase Diagrams of Binary Nickel Alloys”, Nash, P. (Ed.), ASM, Metals Park, OH, 299-306 (1991) (Review, Phase Diagram, Phase Relations, #, 60) Selverian, J.H., Kang, S., “Ceramic-to-Metal Joints Brazed with Palladium Alloys”, Welding J., 71(1), 25-33 (1992) (Experimental, Interface Phenomena, Phys. Prop., 13) cited from abstract Massara, R., Feshotte, P., “The Binary System Pd-Si”, J. Alloys Compd., 190, 249-254 (1993) (Experimental, Phase Diagram, 6) Okamoto, H., “Pd-Si (Palladium-Silicon)”, J. Phase Equilib., 14(4), 536-538 (1993) (Assessment, Phase Diagram, #, 3) Tsai, C.J., Chung, P.L., Yu, K.H., “Stress Evolution of Ni-Pd-Si Reaction System under Isochronal Annealing”, Thin Solid Films, 365, 72 (2000) (Experimental, 38) Wang, R.N., Feng, J.Y., “Comparison of the Thermal Stabilities of NiSi Films in Ni-Si, Ni-Pd-Si and Ni-Pt-Si Systems”, J. Phys.: Condens. Matter, 15, 1935-1942 (2003) (Experimental, Phase Relations, 20) Loomans, M.E., Chi, D.Z., Chua, S.J., “Monosilicide-Disilicide-Silicon Phase Equilibria in the Nickel-Platinum-Silicon and Nickel-Palladium-Silicon Systems”, Metall. Mater. Trans. A, 35(10), 3053-3061 (2004) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, 31)
Table 1: Investigation of the Ni-Pd-Si Phase Relations, Structures and Thermodynamics Reference
Experimental Technique
[1976Wop]
Induction melting and arc furnace melting, powder 20-75 at.% Si diffraction with using Guinier camera
[1990Rab]
Jet cast and planar flow cast to obtain amorphous foils; DSC, DTA, XRD, plasma emission spectroscopy
5-68 at.% Ni-10-79.5 at.% Pd-10-22 at.% Si
[2004Loo]
Arc melting with following vacuum casting Annealing at 730-902°C for 28 - 11 days EDS, optical microscopy, SEM
Ni-2.83Pd-65.6Si Ni-1.57Pd-66.9Si
Landolt-Börnstein New Series IV/11C3
Temperature/Composition/Phase Range Studied
MSIT®
Ni–Pd–Si
490 Table 2: Crystallographic Data of Solid Phases Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Ni,Pd)
cF4 Fm3m Cu
(Ni) < 1455 (Pd) < 1555
Lattice Parameters Comments/References [pm]
a = 352.40
pure Ni at 25°C [Mas2]
a = 389.03
pure Pd at 25°C [Mas2]
a = 543.06
at 25°C [Mas2]
a = 350.6
22.8 to 24.5 at.% Si [Mas2] [V-C2]
(Si) < 1414
cF8 Fd3m C
1, Ni4Si < 1035
cP4 Pm3m AuCu3
2, Ni3Si 1115 - 990
mC16 C2/m GePt3
~24.5 to 25.5 at.% Si [Mas2]
3, Ni3Si 1170 - 1115
mC16 C2/m GePt3
~24.5 to 25.5 at.% Si [Mas2]
, Ni31Si12 < 1242
hP43 P321 Ni31Si2
a = 666.7 0.2 c = 1227.7 0.2
[V-C2], hP14 in [Mas2] at 27.9 at.% Si [Mas2] dissolves ~7 at.% Pd at 800°C [1976Wop]
, Ni2Si 1306 - 825
hP6 P6322 Ni2Si
a = 383.6 0.1 c = 494.8 0.1
[V-C2] 33.4 to 41 at.% Si [Mas2]
, Ni2Si < 1255
oP12 Pnma Co2Si
J', Ni3Si2(h) 845 - 800
-
J, Ni3Si2(r) < 830
oC80 Cmc21 Ni3Si2
NiSi2 993 - 981
-
-
[Mas2], [1991Nas2]
NiSi2 < 981
cF12 Fm3m CaF2
a = 538.3
[V-C2] dissolves 1.43 at.% Pd at 882°C; 1.39 at.% Pd at 854°C; 1.18 at.% Pd at 801°C; 1.07 at.% Pd at 759°C; [2004Loo]
MSIT®
a = 502.2 0.1 b = 374.1 0.1 c = 708.8 0.1 -
[V-C2] at 33.3 at.% Si [Mas2]
39.2 to 41 at.% Si [Mas2] oP80 in [Mas2] 39.0 to 41 at.% Si [Mas2] dissolves ~7 at.% Pd at 800°C [1976Wop]
Landolt-Börnstein New Series IV/11C3
Ni–Pd–Si
491
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters Comments/References [pm]
Pd21Si5 < 811
-
-
Pd5Si 856 - 809
mP24 P21 Pd5P
Pd14Si3 < 795
-
Pd9Si2 819 - 764
oP44 Pnma Pd9Si2
Pd15Si4 792 - 753
aP20 P1 ?
at 21.1 at.% Si, forms by peritectoid reaction [1993Oka]
Pd3Si < 1074
oP16 Pnma Fe3C
[1993Oka] dissolves ~8 at.% Ni at 800°C at 25 at.% Si [1976Wop]
Pd2Si < 1404
a = 846.5 b = 748.5 c = 555.5 = 100.7° -
a = 941.4 b = 741.88 c = 905.48
a = 574.0 0.1 b = 755.6 0.1 c = 526.1 0.1
hP9 P62m Fe2P a = 649.3 0.5 c = 342.7 0.5
at 16 to 16.3 at.% Si, forms by peritectoid reaction [1993Oka] at 16.7 at.% Si [1993Oka] [V-C2] prototype from [Mas2]
at 17.6 at.% Si, forms by peritectoid reaction [1993Oka] at 18.2 at.% Si [1993Oka] [V-C2]
at 33.5 at.% Si [1993Oka] dissolves ~50 at.% Ni at 800°C [1976Wop] at 33.3 at.% Si [1953And]
a = 647.7 0.1 c = 343.2 0.1
at 2.7 at.% Ni
a = 619.6 0.1 c = 325.0 0.1
at 49.2 at.% Ni at 790°C [1976Wop]
Pd2Si' < 1053
-
-
at 33.3 at.% Si by peritectoid reaction [1993Oka]
Pd2Si'' < 1080
-
-
at 33.9 at.% Si by peritectoid reaction [1993Oka]
Pd2Si''' < 1072
-
-
at 34.5 at.% Si by peritectic reaction [1993Oka]
Landolt-Börnstein New Series IV/11C3
MSIT®
Ni–Pd–Si
492 Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Ni,Pd)Si
oP8 Pnma MnP
PdSi 908 - 888
Lattice Parameters Comments/References [pm]
NiSi < 992
a = 561.7 0.3 b = 338.6 0.2 c = 614.9 0.4
[1993Oka] [1976Wop]
a = 562.8 0.2 b = 519.0 0.1 c = 333.0 0.1
at 50 at.% Si at 800°C complete solid solution with PdSi at temperature >888°C dissolves >20 at.% Pd at 800°C at 20 at.% Pd at 800°C
a = 584.3 0.6 b = 538.1 0.5 c = 332.1 0.3 a = 562.4 0.1 b = 519.2 0.1 c = 333.1 0.1
at 20 at.% Pd at 700°C [1976Wop]
* Ni18Pd7Si9
? ? Pd25Ge9
a = 683.5 0.1 c = 991.6 0.1
quenched from 800°C [1976Wop]
* NiPd2Si
?
?
~69-75 at.% Ni [1976Wop]
Si
Data / Grid: at.% Axes: at.%
Fig. 1: Ni-Pd-Si. Isothermal section at 800°C 20
80
αNiSi2 40
60
NiSi 60
ε γ
40
δ
Pd2Si
β1
80
Ni18Pd7Si9
Pd3Si
NiPd2Si
20
Pd9Si2 Pd21Si4
(Ni,Pd)
Ni
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20
40
60
80
Pd
Landolt-Börnstein New Series IV/11C3
Ni–Pd–Si
Temperature, °C
Fig. 2: Ni-Pd-Si. Liquidus and solidus temperature of the alloys along the vertical section 85Pd15Si - 78Ni22Si
493
1200
1100
1000
Liquidus
900
Solidus 800
700
Ni 78.00 Pd 0.00 Si 22.00
10
20
30
40
50
60
70
Pd, at.%
80
0.00 Ni Pd 85.00 Si 15.00
1000
Fig. 3: Ni-Pd-Si. Vertical section NiSi-PdSi
992°C
Temperature, °C
900
(Ni,Pd)Si
800
(Ni,Pd)Si+(Ni,Pd)2Si+(Si) 700
600
Ni 50.00 Pd 0.00 Si 50.00
Landolt-Börnstein New Series IV/11C3
10
20
30
Pd, at.%
40
0.00 Ni Pd 50.00 Si 50.00
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