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 D Iron Systems Part 4 Selected Systems from Cu-Fe-Si to Fe-N-U Editors G. Effenberg and S. Ilyenko
Authors Materials Science and International Team, MSIT®
ISSN
1615-2018 (Physical Chemistry)
ISBN
978-3-540-78643-6 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/11D4: 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 2008 Printed in Germany The use of general descriptive names, registered names, trademarks, etc. in this publication does not imply, even in the absence of a specific statement, that such names are exempt from the relevant protective laws and regulations and therefore free for general use. Product Liability: The data and other information in this handbook have been carefully extracted and evaluated by experts from the original literature. Furthermore, they have been checked for correctness by authors and the editorial staff before printing. Nevertheless, the publisher can give no guarantee for the correctness of the data and information provided. In any individual case of application, the respective user must check the correctness by consulting other relevant sources of information. Cover layout: Erich Kirchner, Heidelberg Typesetting: Materials Science International Services GmbH, Stuttgart Printing and Binding: AZ Druck, Kempten/Allgäu
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Editors: Associate Editor:
Günter Effenberg Svitlana Ilyenko Oleksandr Dovbenko
MSI, Materials Science International Services GmbH Postfach 800749, D-70507, Stuttgart, Germany http://www.matport.com
Authors: Materials Science International Team, MSIT® The present series of books results from collaborative evaluation programs performed by MSI and authored by MSIT®. In this program data and knowledge are contributed by many individuals and accumulated over almost twenty years, now. The content of this volume is a subset of the ongoing MSIT® Evaluation Programs. Authors of this volume are:
Christian Bätzner, Stuttgart, Germany
Tatiana Pryadko, Kyiv, Ukraine
Gabriele Cacciamani, Genova, Italy
James Robinson, Manchester, UK
Lesley Cornish, Randburg, South Africa
Lazar Rokhlin, Moscow, Russia
Olga Fabrichnaya, Freiberg, Germany
Elena Semenova, Kyiv, Ukraine
Mireille Harmelin, Paris, France
Larysa Shcherbak, Chernivtsi, Ukraine
Volodymyr Ivanchenko, Kyiv, Ukraine
Athanasios Stamou, Volos, Greece
Kostyantyn Korniyenko, Kyiv, Ukraine
Vasyl Tomashik, Kyiv, Ukraine
Viktor Kuznetsov, Moscow, Russia
Lyudmilla Tretyachenko, Kyiv, Ukraine
Nathalie Lebrun, Lille, France
Mikhail Turchanin, Kramatorsk, Ukraine
Evgeniya Lysova, Moscow, Russia
Tamara Velikanova, Kyiv, Ukraine
Hans Leo Lukas, Stuttgart, Germany
Andy Watson, Leeds, U.K.
Pierre Perrot, Lille, France
Institutions The content of this volume is produced by MSI, Materials Science International Services GmbH and the international team of materials scientists, MSIT®. Contributions to this volume have been made from the following institutions:
The Baikov Institute of Metallurgy, Academy of Sciences, Moscow, Russia
Moscow State University, Department of General Chemistry, Moscow, Russia
Chernivtsi National University, Chemical Department, Chernivtsi, Ukraine
MSI, Materials Science International Services GmbH, Stuttgart, Germany
Donbass State Mechanical Engineering Academy, Kramatorsk, Ukraine
School of Chemical and Metallurgical Engineering, The University of the Witwatersrand, DST/NRF Centre of Excellence for Strong Material, South Afrika
I.M. Frantsevich Institute for Problems of Materials Science, National Academy of Sciences, Kyiv, Ukraine Institute for Semiconductor Physics, National Academy of Sciences, Kyiv, Ukraine
Technische Universität Bergakademie Freiberg, Institut für Werkstoffwissenschaft, Freiberg, Germany
G.V. Kurdyumov Institute for Metal Physics, National Academy of Sciences, Kyiv, Ukraine
Universita di Genova, Dipartimento di Chimica, Genova, Italy
Manchester Materials Science Centre University of Manchester and UMIST, Manchester, UK
Universite de Lille I, Laboratoire de Métallurgie Physique, Villeneuve d’ASCQ, France
Max-Planck-Institut für Metallforschung, Institut für Werkstoffwissenschaft, Pulvermetallurgisches Laboratorium, Stuttgart, Germany
University of Leeds, Department of Materials, School of Process, Environmental and Materials Engineering, Leeds, UK University of Thessaly, Volos, Greece
Preface
The sub-series Ternary Alloy Systems of the Landolt-Börnstein New Series provides reliable and comprehensive descriptions of the materials constitution, based on critical intellectual evaluations of all data available at the time and it critically weights the different findings, also with respect to their compatibility with today’s edge binary phase diagrams. Selected are ternary systems of importance to alloy development and systems which gained in the recent years otherwise scientific interest. In one ternary materials system, however, one may find alloys for various applications, depending on the chosen composition. Reliable phase diagrams provide scientists and engineers with basic information of eminent importance for fundamental research and for the development and optimization of materials. So collections of such diagrams are extremely useful, if the data on which they are based have been subjected to critical evaluation, like in these volumes. Critical evaluation means: there where contradictory information is published data and conclusions are being analyzed, broken down to the firm facts and re-interpreted in the light of all present knowledge. Depending on the information available this can be a very difficult task to achieve. Critical evaluations establish descriptions of reliably known phase configurations and related data. The evaluations are performed by MSIT®, Materials Science International Team, a group of scientists working together since 1984. Within this team skilled expertise is available for a broad range of methods, materials and applications. This joint competence is employed in the critical evaluation of the often conflicting literature data. Particularly helpful in this are targeted thermodynamic and atomistic calculations for individual equilibria, driving forces or complete phase diagram sections. Conclusions on phase equilibria may be drawn from direct observations e.g. by microscope, from monitoring caloric or thermal effects or measuring properties such as electric resistivity, electro-magnetic or mechanical properties. Other examples of useful methods in materials chemistry are massspectrometry, thermo-gravimetry, measurement of electro-motive forces, X-ray and microprobe analyses. In each published case the applicability of the chosen method has to be validated, the way of actually performing the experiment or computer modeling has to be validated as well and the interpretation of the results with regard to the material’s chemistry has to be verified. Therefore insight in materials constitution and phase reactions is gained from many distinctly different types of experiments, calculation and observations. Intellectual evaluations which interpret all data simultaneously reveal the chemistry of the materials system best. An additional degree of complexity is introduced by the material itself, as the state of the material under test depends heavily on its history, in particular on the way of homogenization, thermal and mechanical treatments. All this is taken into account in an MSIT® expert evaluation. To include binary data in the ternary evaluation is mandatory. Each of the three-dimensional ternary phase diagrams has edge binary systems as boundary planes; their data have to match the ternary data smoothly. At the same time each of the edge binary systems A-B is a boundary plane for many other ternary A-B-X systems. Therefore combining systematically binary and ternary evaluations increases confidence and reliability in both ternary and binary phase diagrams. This has started systematically for the first time here, by the MSIT® Evaluation Programs applied to the Landolt-Börnstein New Series. The degree of success, however, depends on both the nature of materials and scientists! The multitude of correlated or inter-dependant data requires special care. Within MSIT® an evaluation routine has been established that proceeds knowledge driven and applies both, human based expertise and electronically formatted data and software tools. MSIT® internal discussions take place in almost all evaluation works and on many different specific questions the competence of a team is added to the work of individual authors. In some cases the authors of earlier published work contributed to the knowledge
base by making their original data records available for re-interpretation. All evaluation reports published here have undergone a thorough review process in which the reviewers had access to all the original data. In publishing we have adopted a standard format that presents the reader with the data for each ternary system in a concise and consistent manner, as applied in the “MSIT® Workplace Phase Diagrams Online”. The standard format and special features of the Landolt-Börnstein compendium are explained in the Introduction to the volume. In spite of the skill and labor that have been put into this volume, it will not be faultless. All criticisms and suggestions that can help us to improve our work are very welcome. Please contact us via
[email protected]. We hope that this volume will prove to be as useful for the materials scientist and engineer as the other volumes of Landolt-Börnstein New Series and the previous works of MSIT® have been. We hope that the Landolt Börnstein Sub-series, Ternary Alloy Systems will be well received by our colleagues in research and industry. On behalf of the participating authors we want to thank all those who contributed their comments and insight during the evaluation process. In particular we thank the reviewers - Hans Leo Lukas, Marina Bulanova, Paola Riani, Lazar Rokhlin, Anatolii Bondar, Yong Du, Olga Fabrichnaya, Artem Kozlov, K.C. Hari Kumar, Viktor Kuznetsov, Ludmila Tretyachenko and Tamara Velikanova. We all gratefully acknowledge the dedicated scientific desk editing by Oleksandra Berezhnytska and Oleksandr Rogovtsov.
Günter Effenberg, Svitlana Ilyenko and Oleksandr Dovbenko
Stuttgart, November 2007
Foreword Can you imagine a world without iron and steel? No? I can’t either. The story of mankind is intimately linked to the discovery and successful use of metals and their alloys. Amongst them iron and steel - we could define steel as ‘a generally hard, strong, durable, malleable alloy of iron and carbon, usually containing between 0.2 and 1.5 percent carbon, often with other constituents such as manganese, Chromium, nickel, molybdenum, copper, tungsten, Cobalt, or silicon, depending on the desired alloy properties, and widely used as a structural material’, have shaped our material world. The story of iron takes us back to the period of the Hittite Empire around 1300 BC, when iron started to replace bronze as the chief metal used for weapons and tools. Until today the story remains uncompleted and the social and economic impact of the iron and steel industry is now beyond imagination. In the year 2005 1.13 billion tons of crude steel were produced. Compared to 2004 this is an increase of 6.8%. That same year the steel production in China increased from 280.5 to almost 350 million tons. Concerning stainless steel: according to the International Stainless Steel Forum (ISSF), the global production forecast for 2006 now stands at 27.8 million metric tons of stainless crude steel, up 14.3% compared to 2005. An English poem from the 19th century tells us Gold is for the mistress Silver for the maid Copper for the craftsman Cunning at his trade Good said the baron Sitting in his hall But iron, cold iron Is master of them all It is still actual and true. The list of different steel grades and related applications is impressive and still growing: low carbon strip steels for automotive applications, low carbon structural steels, engineering steels, stainless steels, cast irons, and, more recently: dual phase steels, TRIP-steels, TWIP-steels, maraging steels, … The list of applications seems endless: a wide range of properties from corrosion resistance to high tensile strength is covered. These properties depend on the percentage of carbon, the alloying elements, and increasingly on the thermo-mechanical treatments that aim at optimizing the microstructure. Yet many potential improvements remain unexplored, also due to the increasing complexity of the new steel grades. For instance, a recently patent protected new die steel for hot deformation has the following composition specifications: C 0.46 – 0.58; Si 0.18 – 0.40; Mn 0.45 – 0.75, Cr 0.80 – 1.20; Ni 1.30 – 1.70; Mo 0.35 – 0.65; V 0.18 – 0.25; Al 0.01 – 0.04; Ti 0.002 – 0.04; B 0.001 – 0.003; Zr 0.02 – 0.04; Fe remaining.
Although many properties of steel are directly related to non-equilibrium states, it remains a fact that the equilibrium state creates the reference frame for all changes that might occur in any material - and consequently would effect its properties in use - that is actually not in its thermodynamic equilibrium state. This is what these volumes in the Landolt-Börnstein series stand for: they have collected the most reliable data on the possible phase equilibria in ternary iron based alloys. Therefore this first volume of data, as well as the other ones in a series of four to appear, is of immeasurable value for metallurgists and materials engineers that improve the properties of existing steels and develop new and more complex steel grades. It is about materials, it is about quality of life. The well-recognized quality label of MSIT®, the Materials Science International Team, also applies to the present volume of the Landolt-Börnstein series. It should be available for every materials engineer, scientist and student.
Prof. Dr. ir. Patrick Wollants Chairman - Department of Metallurgy and Materials Engineering Katholieke Universiteit Leuven Belgium
Contents IV/11D4 Ternary Alloy Systems Phase Diagrams, Crystallographic and Thermodynamic Data Subvolume D Iron Systems Part 4 Selected Systems from Cu-Fe-Si to Fe-N-U Introduction Data Covered ................................................................................................................................... XIII General............................................................................................................................................. XIII Structure of a System Report ........................................................................................................... XIII Introduction.............................................................................................................................. XIII Binary Systems ........................................................................................................................ XIII Solid Phases .............................................................................................................................XIV Quasibinary Systems................................................................................................................. XV Invariant Equilibria ................................................................................................................... XV Liquidus, Solidus, Solvus Surfaces........................................................................................... XV Isothermal Sections................................................................................................................... XV Temperature – Composition Sections ....................................................................................... XV Thermodynamics....................................................................................................................... XV Notes on Materials Properties and Applications....................................................................... XV Miscellaneous ........................................................................................................................... XV References............................................................................................................................. XVIII General References ..........................................................................................................................XIX
Ternary Systems Cu – Fe – Si (Copper – Iron – Silicon)................................................................................................. 1 Cu – Fe – Ti (Copper – Iron – Titanium) ........................................................................................... 15 Fe – H – Mn (Iron – Hydrogen – Manganese) ................................................................................... 32 Fe – H – Mo (Iron – Hydrogen – Molybdenum)................................................................................ 42 Fe – H – Ni (Iron – Hydrogen – Nickel) ............................................................................................ 50 Fe – H – O (Iron – Hydrogen – Oxygen) ........................................................................................... 65 Fe – H – P (Iron – Hydrogen – Phosphorus)...................................................................................... 72 Fe – H – Si (Iron – Hydrogen – Silicon) ............................................................................................ 76 Fe – H – V (Iron – Iron – Vanadium)................................................................................................. 82 Fe – La – Si (Iron – Lanthanum – Silicon)......................................................................................... 87 Fe – Mg – O (Iron – Magnesium – Oxygen)...................................................................................... 96 Fe – Mg – S (Iron – Magnesium – Sulphur) .................................................................................... 127 Fe – Mg – Si (Iron – Magnesium – Silicon)..................................................................................... 135 Fe – Mn – N (Iron – Manganese – Nitrogen)................................................................................... 148 Fe – Mn – Ni (Iron – Manganese – Nickel) ..................................................................................... 162 Fe – Mn – O (Iron – Manganese – Oxygen) .................................................................................... 178
Fe – Mn – P (Iron – Manganese – Phosphorus) ............................................................................... 200 Fe – Mn – S (Iron – Manganese – Sulfur)........................................................................................220 Fe – Mn – Si (Iron – Manganese – Silicon) ..................................................................................... 246 Fe – Mn – Ti (Iron – Manganese – Titanium).................................................................................. 280 Fe – Mn – V (Iron – Manganese – Vanadium) ................................................................................ 291 Fe – Mn – Zr (Iron – Manganese – Zirconium) ............................................................................... 303 Fe – Mo – N (Iron – Molybdenum – Nitrogen)................................................................................ 311 Fe – Mo – Ni (Iron – Molybdenum – Nickel) .................................................................................. 321 Fe – Mo – O (Iron – Molybdenum – Oxygen) ................................................................................. 341 Fe – Mo – Si (Iron – Molybdenum – Silicon) .................................................................................. 355 Fe – N – Nb (Iron – Nitrogen – Niobium) ....................................................................................... 371 Fe – N – Ni (Iron – Nitrogen – Nickel)............................................................................................ 381 Fe – N – Si (Iron – Nitrogen – Silicon)............................................................................................ 398 Fe – N – Ti (Iron – Nitrogen – Titanium) ........................................................................................ 409 Fe – N – U (Iron – Nitrogen – Uranium) .........................................................................................421
Introduction
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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 quasibinary sections; data on invariant equilibria are generally given in the form of tables. The world literature is thoroughly and systematically searched back to the year 1900. Then, the published data are critically evaluated by experts in materials science and reviewed. Conflicting information is commented upon and errors and inconsistencies removed wherever possible. It considers those, and only those data, which are firmly established, comments on questionable findings and justifies re-interpretations made by the authors of the evaluation reports. In general, the approach used to discuss the phase relationships is to consider changes in state and phase reactions which occur with decreasing temperature. This has influenced the terminology employed and is reflected in the tables and the reaction schemes presented. The system reports present concise descriptions and hence do not repeat in the text facts which can clearly be read from the diagrams. For most purposes the use of the compendium is expected to be self-sufficient. However, a detailed bibliography of all cited references is given to enable original sources of information to be studied if required. Structure of a System Report The constitutional description of an alloy system consists of text and a table/diagram section which are separated by the bibliography referring to the original literature (see Fig. 1). The tables and diagrams carry the essential constitutional information and are commented on in the text if necessary. Where published data allow, the following sections are provided in each report: 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. 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.
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Introduction
Fig. 1.
Structure of a system report
Phases with the same formulae but different space lattices (e.g. allotropic transformation) are distinguished by: – small letters (h), high temperature modification (h2 > h1) (r), room temperature modification (1), low temperature modification (l1 > l2) – Greek letters, e.g., ε, ε' – Roman numerals, e.g., (I) and (II) for different pressure modifications. In the table “Solid Phases” ternary phases are denoted by * and different phases are separated by horizontal lines. Quasibinary Systems
Quasibinary (pseudobinary) sections describe equilibria and can be read in the same way as binary diagrams. The notation used in quasibinary systems is the same as that of vertical sections, which are reported under “Temperature – Composition Sections”.
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Fig. 2.
Typical reaction scheme
Introduction
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Introduction
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).
Fig. 3.
Hypothetical liqudus surface showing notation employed
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Fig. 4.
5
Hypotheticcal isothermal section showing notation employed
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 quasibinary phase diagrams. Thermodynamics
Experimental ternary data are reported in some system reports and reference to thermodynamic modeling 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.
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Fig. 5.
Introduction
Hypothetical vertical section showing notation employed
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. References
The publications which form the bases of the assessments are listed in the following manner: [1974Hay] Hayashi, M., Azakami, T., Kamed, M., “Effects of Third Elements on the Activity of Lead in Liquid Copper Base Alloys” (in Japanese), Nippon Kogyo Kaishi, 90, 51–56 (1974) (Experimental, Thermodyn., 16) This paper, for example, whose title is given in English, is actually written in Japanese. It was published in 1974 on pages 51–56, volume 90 of Nippon Kogyo Kaishi, the Journal of the Mining and Metallurgical Institute of Japan. It reports on experimental work that leads to thermodynamic data and it refers to 16 cross-references. Additional conventions used in citing are: # to indicate the source of accepted phase diagrams * to indicate key papers that significantly contributed to the understanding of the system. Standard reference works given in the list “General References” are cited using their abbreviations and are not included in the reference list of each individual system.
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General References [C.A.] Chemical Abstracts – pathways to published research in the world's journal and patent literature – http://www.cas.org/ [Curr. Current Contents – bibliographic multidisciplinary current awareness Web resource – http:// Cont.] www.isinet.com/products/cap/ccc/ [E] Elliott, R.P., Constitution of Binary Alloys, First Supplement, McGraw-Hill, New York (1965) [G] Gmelin Handbook of Inorganic Chemistry, 8th ed., Springer-Verlag, Berlin [H] Hansen, M. and Anderko, K., Constitution of Binary Alloys, McGraw-Hill, New York (1958) [L-B] 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. [Mas] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, ASM, Metals Park, Ohio (1986) [Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) [P] Pearson, W.B., A Handbook of Lattice Spacings and Structures of Metals and Alloys, Pergamon Press, New York, Vol. 1 (1958), Vol. 2 (1967) [S] Shunk, F.A., Constitution of Binary Alloys, Second Supplement, McGraw-Hill, New York (1969) [V-C] Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, ASM, Metals Park, Ohio (1985) [V-C2] Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Index of alloy systems
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Index of alloy systems Index of Ternary Iron Alloy Systems Cu-Fe-Si to Fe-Ti-Y Cu–Fe–Si (Copper – Iron – Silicon) Cu–Fe–Ti (Copper – Iron – Titanium) Fe–H–Mn (Iron – Hydrogen – Manganese) Fe–H–Mo (Iron – Hydrogen – Molybdenum) Fe–H–Ni (Iron – Hydrogen – Nickel) Fe–H–O (Iron – Hydrogen – Oxygen) Fe–H–P (Iron – Hydrogen – Phosphorus) Fe–H–Si (Iron – Hydrogen – Silicon) Fe–H–V (Iron – Hydrogen – Vanadium) Fe–La–Si (Iron – Lanthanum – Silicon) Fe–Mg–O (Iron – Magnesium – Oxygen) Fe–Mg–S (Iron – Magnesium – Sulfur) Fe–Mg–Si (Iron – Magnesium – Silicon) Fe–Mn–N (Iron – Manganese – Nitrogen) Fe–Mn–Ni (Iron – Manganese – Nickel) Fe–Mn–O (Iron – Manganese – Oxygen) Fe–Mn–P (Iron – Manganese – Phosphorus) Fe–Mn–S (Iron – Manganese – Sulfur) Fe–Mn–Si (Iron – Manganese – Silicon) Fe–Mn–Ti (Iron – Manganese – Titanium) Fe–Mn–V (Iron – Manganese – Vanadium) Fe–Mn–Zr (Iron – Manganese – Zirconium) Fe–Mo–N (Iron – Molybdenum – Nitrogen) Fe–Mo–Ni (Iron – Molybdenum – Nickel) Fe–Mo–O (Iron – Molybdenum – Oxygen) Fe–Mo–Si (Iron – Molybdenum – Silicon) Fe–N–Nb (Iron – Nitrogen – Niobium) Fe–N–Ni (Iron – Nitrogen – Nickel) Fe–N–Si (Iron – Nitrogen – Silicon) Fe–N–Ti (Iron – Nitrogen – Titanium) Fe–N–U (Iron – Nitrogen – Uranium) Fe–Ti–Y (Iron – Titanium – Yttrium)
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Cu–Fe–Si
1
Copper – Iron – Silicon Nathalie Lebrun, Pierre Perrot, Mireille Harmelin
Introduction Fe and Si are known for hardening Cu alloys due to the precipitation of iron silicides [1927Cor] and the Cu-Fe-Si alloys are commonly formed during melting in treatment process of domestic wastes incinerators [1999Hin]. The phase diagram presented by [1949Jae] takes into account the previous determination of [1933Rol] in the iron rich corner and [1934Han] in the copper rich corner. The actual knowledge of the ternary diagram rests on a thorough investigation by [1953Vog, 1999Hin] of the Cu-Fe-FeSi-Cu3Si domain, which precises the shape of the miscibility gap in the ternary liquid. The miscibility gap was investigated theoretically by [1991But]. A comprehensive review may be found in [1979Cha, 1992Rag, 2002Rag]. Calphad assessments of the Cu-Fe-Si system have been carried out by [1997Oht, 1999Hin, 2003Mie] near the Cu-Fe border (< 40 mass% Si). Experimental works are listed in Table 1. Binary Systems After [Mas2], the liquidus and solidus phase boundaries in the Fe rich portion of the Fe-Si phase diagram were measured by [2005Mec] and thermodynamically reassessed by [1998Mie]. The DTA measurements done by [2005Mec] reveal a significant shift compared to [Mas2] in the solidus and liquidus curves in the range between 20 and 25 at.% Si (11.17 and 14.36 mass% Si), suggesting that the A2-B2 (and perhaps B2-D03) ordering transition occurs in large part through the two-phase temperature regions. Since these careful experimental data are identical to those used by [1998Mie] for his thermodynamic assessment, these data are retained in the present evaluation. Nevertheless, a more sophisticated thermodynamic treatment of the ordering transition is necessary for a better accuracy of the Fe-Si diagram in the range from 20 to 35 at.% Si (11.17 to 21.31 mass% Si). Cu-Fe is accepted from [2007Tur]. A metastable miscibility gap in the liquid phase exists. Cu-Si, based on the compilation of [1986Ole] is accepted from [2002Leb]. Solid Phases The binary solid compounds are listed in Table 2. Only one ternary metastable compound has been reported. The lattice parameters of Cu rich ternary single-phase alloys (up to 5 Fe, 5 Cu at.%) were reported by [1943And] for alloys quenched from various temperatures between 1025 and 800°C. Si decreases the Fe solubility in solid Cu [1934Han, 1943And]. Although [1933Rol] proposes an increase of Cu solubility in solid Fe, more recent studies [1997Oht] show also a decrease of the Cu solubility in solid Fe. The solubility of Cu in FeSi and FeSi2 phases is estimated less than 0.2 at.% [1997Yam]. Quasibinary Systems Cu3Si-FeSi section is found to be quasibinary with a eutectic point at 847°C and 2 mass% Fe [1953Vog]. Recent experimental data [2002Was] predict the existence of a liquid miscibility gap along this quasibinary section. Consequently the section has been modified taking into account this liquid miscibility gap. The temperature of 1400°C has been chosen since at 1450°C the liquid miscibility gap closes itself and at 1350°C the FeSi+L phase field was found by [2003Mie]. The critical temperature of this two-liquid domain was estimated to be 1632°C [1999Hin]. The not well know boundaries have been indicated as dashed lines in Fig. 1. The eutectic position has been moved to be in agreement with the accepted isothermal sections at 1100 and 900°C. The vertical section FeSi2-Cu3Si was constructed by [1975Abr] and claimed to be quasibinary. It presents a eutectic at 800°C towards 82 mol% Cu3Si. However, this section cannot be quasibinary because the transformation of βFeSi2 into αFeSi2 occurs following a peritectoid reaction at 982°C accompanied with a change in composition as shown in Table 2. Landolt-Börnstein New Series IV/11D4
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2
Cu–Fe–Si
Invariant Equilibria The first experimental study was reported by [1953Vog, 1979Cha] who detected four ternary U reactions in the Cu rich corner of the Cu-Fe-Si system. Three of these reactions are accepted in the present evaluation, their details are given in Table 3. The σ phase, wrongly labelled Fe3Si2 by [1953Vog] corresponds actually to the composition Fe5Si3. The temperatures of these ternary reactions are located between 900°C and 1100°C. Further experimental investigations are needed to determine precisely these reaction temperatures. The reaction (αFe)+FeSi ⇌ L+Fe5Si3 indicated by [1953Vog] has been rejected since no three-phase field (αFe) +FeSi+Fe5Si3 is indicated on the accepted isothermal sections below 916°C. More recently, from the isothermal sections constructed by [1999Hin], a ternary reaction of the U type may be supposed between 1450 and 1350°C: (γFe)+L1 ⇌ (αFe)+L2. No information is available regarding the compositions of the phases involved. Additional ternary reactions proposed by [1992Rag] are only speculative and the deduced reaction scheme for the Fe-FeSi-Cu3Si-Cu region could not be accepted. Liquidus, Solidus and Solvus Surfaces The first experimental results on the liquidus surface was reported by [1934Han] in the Cu rich corner showing the isothermal lines and the saturation curve between the (Cu) and (γFe) phases. Later [1953Vog] investigated the liquidus surface for the FeSi-Cu3Si-Cu region and suggested the existence of four ternary U reactions. After the review of [1979Cha], [1992Rag] reassessed critically the liquidus surface and speculated others ternary reactions without experimental findings. The dominant feature of this liquidus surface is the liquid miscibility gap that closes on itself. Two critical points have been measured at 1425°C (62 mass% Fe, 36.5 mass% Cu, 1.5 mass% Si) and 1110°C (18 mass% Fe, 70 mass% Cu, 12 mass% Si) [1953Vog]. From heat capacity measurements, [2002Was] deduced the liquid miscibility gap in the Cu rich corner at 1500°C and suggested a wider region of the two-liquid domain than the one measured by [1953Vog]. Liquid miscibility gap has been also calculated by [2002Wan], who extrapolated the two-liquid region up to 1727°C using thermodynamic assessment of experimental data obtained in the 800–1000°C temperature range. At 1800°C, the two-liquid region does not exist [2002Was]. [1999Hin] estimated the critical temperature of the two liquid region at 1632°C. Considering the disagreement about the miscibility gap and the location of the ternary reactions, the liquidus surface proposed by [1953Vog] could not be accepted. Further careful experimental data on the entire system are needed. Isothermal Sections The Fe rich corner of the Cu-Fe-Si system was first investigated by [1933Rol]. At room temperature, the solubility of Cu in Fe was found to increase with increasing Si content. The boundary was determined passing through about 1% Cu, 3% Si and 5% Cu, 16% Si (mass%). [1934Han] examined the Cu rich corner below 1100°C and reported that at temperatures between 1100 and 800°C, the solubility of Fe in Cu is decreased by the presence of Si. The equilibrium of (Cu) and (Fe) in the presence of Si has been measured by [1981Has] on one tie-line at 1000°C (xCu = 0.030, xSi = 0.057 for the α phase and xCu = 0.039, xSi = 0.046 for the γ phase), and at 1300°C (xCu = 0.041, xSi = 0.050 for the γ phase and xCu = 0.055, xSi = 0.042 for the liquid phase). A more thorough investigation of the (αFe)-(Cu) or (γFe)-(Cu) equilibria was carried out between 800 and 1000°C by [2002Wan]. At high temperature, a liquid miscibility gap is located near the Cu-Fe side. At 1450°C, it closes on itself and silicon stabilizes it with decreasing temperature. Large discrepancies are pointed out regarding the width of this miscibility gap. The maximum of the Si composition was estimated to be about 13 mass% by [1953Vog] whereas the computed data estimate this maximum around 30 mass% at 1450°C [1997Oht, 1999Hin, 2002Wan]. A wider miscibility gap at 1450°C was predicted by [2003Mie] with a maximum at 35 mass% Si. Recently, [2002Was] performed heat capacity measurements at 1800°C and extrapolated the liquid miscibility gap in the Fe rich corner (< 30 mass% Cu, < 50 at.% Si) at 1500°C. A wider liquid miscibility gap suggested in [2002Was] is in quite good agreement with the computed results of [2003Mie]. Consequently the work done by [2003Mie] is preferred. Nevertheless, further experimental studies on the liquid miscibility gap are needed. DOI: 10.1007/978-3-540-78644-3_3 # Springer 2008
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Cu–Fe–Si
3
The isothermal sections at 1450, 1350, 1250 and 1100°C are taken from the more recent calculations [2003Mie] and are depicted in Fig. 2 (1450°C), Fig. 3 (1350°C) and Fig. 4 (1100°C). The isothermal section at 1250°C is reported because of its complete similarity with the one calculated at 1350°C. Only the phase fields located along the Cu-Fe system up to 10 mass% Si were calculated by [2003Mie]. Consequently, some fields are shown by dashed lines at 1100°C. The isothermal sections calculated at 1000°C and 900°C by [2003Mie] are in disagreement with the binary system Fe-Si since the Fe5Si3 phase has not been taken into account in the thermodynamic assessment. The 900°C isothermal section calculated by [2002Wan] is preferred selected and reproduced in Fig. 5. Small modifications have been done regarding the maximum of the Fe composition of the liquid region in agreement with the quasibinary system FeSi-Cu3Si. Others isothermal sections at lower temperature have been established by [1934Han, 1943And, 1953Vog] in the Cu rich corner. The isothermal sections of [1934Han] do not include all the known Cu-Si intermediate phase and the phase rule is not obeyed. The isothermal sections presented by [1943And] in the Cu rich corner are in good agreement with the accepted binary systems. Nevertheless, some discrepancies are observed at 875°C compared to the one established at 900°C by [2002Wan]. The three-phase field L+(Cu)+Fe5Si3 is absent in [2002Wan]. Further experimental investigations are needed at low temperature. Consequently, these isothermal sections are not accepted in the present evaluation. Temperature – Composition Sections A number of vertical sections have been determined for this system: at 0.5 and 2 mass% Fe [1943And], at 20, 40 and 65 mass% Cu, 10 mass% Fe and 10 mass% Si, and along FeSi-Cu [1953Vog]. The vertical sections determined by [1943And] are in contradiction with the isothermal section at 900°C constructed by [2002Wan]. A two-phase field FeSi+(Cu) (which is absent in [1943And]) was experimentally measured by [2002Wan] at 900°C. The vertical sections proposed by [1953Vog] could not be accepted due to their discrepancies with the isothermal sections, in particular the absence of Fe2Si, and the deeper liquid miscibility gap compared to the more recent experimental work done by [2002Was]. [1975Abr] constructed a vertical section FeSi2-Cu3Si which cannot be quasibinary do to the composition change of the two FeSi2 phases in the binary system. A eutectic between αFeSi2 and η,Cu3Si was found at 82 mol% Cu3Si. At higher temperatures [1975Abr] stipulated the existence of a two-phase fields L+βFeSi2 above 960°C and L+αFeSi2 below 960°C. The accepted isothermal sections at 1100°C, 1350°C do not confirm their existence. Moreover, the liquid miscibility gap should exist along this vertical section. Consequently, this vertical section is not retained in this assessment. Further experimental data are needed. Thermodynamics The interaction coefficients of Cu and Si in liquid alloys, defined, in the iron rich alloys, as e(Si)Cu = ∂ ln fCu /∂ (mass% Si), where fCu = (mass% Cu in pure Fe)/(mass% Cu in alloy) at constant copper activity, has been evaluated as e(Si)Cu = + 0.027 at 1600°C by [1974Sig]. Si in liquid Fe increases the activity coefficient of Cu, thus decreases its solubility under an imposed copper activity. The interaction coefficient e(Cu)Si = ∂ ln fSi /∂ (mass% Cu) was calculated e(Cu)Si = + 0.014 at the same conditions. The excess Gibbs energy of a φ phase, in which φ represents any of the α, γ and liquid solutions were evaluated by [1997Oht, 1999Hin, 2003Mie] using the following model: Gxs(φ) = Gxs(Cu,Fe,φ) + Gxs(Cu,Mn,φ) + Gxs(Fe,Mn,φ) + Gxs(Cu,Fe,Mn,φ) The binary terms for a φ phase are given by a Redlich-Kister expansion: Gxs(i, j, φ) = xi xj Σk L(k, φ)(xi – xj) k The ternary term is given by Gxs(Cu,Fe,Mn,φ) = xCu xjFe xMn A(φ). The ternary interaction parameters A (φ) proposed by [1997Oht] and [1999Hin] are the same if φ = γ. For the liquid phase, [1997Oht] gives A(L) = 0. The ternary terms calculated by [1999Hin] are supported by its experimental determination of the two-liquid domain: A (α) = – 237300 + 100 T A (γ) = – 187300 + 100 T A (L) = 19000 – 22.5 T Landolt-Börnstein New Series IV/11D4
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Cu–Fe–Si
More elaborate expressions for these ternary terms have been proposed by [2003Mie]. The value accepted above for A(L) was later confirmed by [2002Was] who measured the heat content of the Cu-Fe-Si alloys at 1800°C in order to evaluate the energy needed to melt the metallic phase generated in the melting treatment process of domestic waste incineration residue. The heat content HT – H298 is shown at 1800°C in Fig. 6 and at 1500°C in Fig. 7. The values of HT – H298 on both figures are given in kJ·mol–1. An input energy of 50 to 65 kJ·mol–1 is necessary to melt the metallic phase generated by the incineration of domestic waste and keep it at 1500°C. Notes on Materials Properties and Applications Experimental works are gathered in Table 4. Cu is considered as the most promising alternative to Al based alloys for integration materials in silicon based integrated circuits. However, Cu reacts easily with Si to form the brittle compound Cu3Si even at temperatures as low as 200°C. The presence of Fe lowers the coefficient of thermal expansion of the Cu-Si matrix and forms the FeSi2 compound at the expanse of the brittle Cu3Si at the Cu-Si interface [2001Lee]. A small amount of Fe affects strongly the annealing behavior such as recovery, anneal hardening and recrystallization of cold worked Cu-Si alloys [1965Ued]. Cu-Fe alloys (1.5 to 2.5 mass% Fe) possess a higher tensile strength and a better conductibility than pure Cu. The presence of Si (1 mass% Si) causes an increase in the strength and an increase in the electrical resistivity [1983Str] because of the formation of Fe3Si precipitates. The presence of Si increases the evaporation rate of Cu [2003Ono] owing to the strong affinity between Fe and Si. At 1800°C, the activity coefficient of Cu at infinite dilution increases from γCu = 9.5 in pure liquid iron to a maximum γCu = 17.8 in Fe added with 20 at.% Si. A further increase of Si content decreases γCu because of the strong affinity between Si and Cu. As a consequence, its vapor pressure and its vaporization rate decrease with the Si content above 20 at.% Si. In pure Si at 1800°C, the Cu activity coefficient is only γCu = 0.16. Cu is considered as a tramp element in ferrous scraps which causes an adverse effect in the mechanical properties of steel products such as its hot workability. Several processes have been proposed for extracting Cu from the Fe-Cu recycling scraps. One is to use the miscibility gap which appears at addition of C or Si, which cause a separation between a Cu rich and a Fe rich phase [2004Wan]. The addition of less than 1 at.% B [2005Yam] makes possible an undercooling, thus, a phase separation in the metastable liquid domain. The vaporization processes are also considered as suitable for removal of Cu from iron alloys. Cu vaporizes at high temperatures (1800°C) and the presence of Si enhances the vaporization rate [2003Ono]. Miscellaneous The fcc metastable phase of an approximate composition Fe63Cu27Si10 has been obtained by bombarding iron rich Cu-Fe-Si multilayers made by depositing alternately pure elements in a high vacuum system [1992Liu] with Xe ions (1015 to 1016 Xe·cm–2 under 200 keV). Another phase reflecting sharp diffraction spots with 12-fold symmetry was also observed. The sharp incommensurate diffraction pattern bore a close relationship with the fcc metastable phase. βFeSi2 is known as an excellent high temperature thermoelectric material. Unfortunately, its preparation from the liquid goes through the eutectic transformation L ⇌ FeSi + αFeSi2 followed by the synthesis of βFeSi2 below 1255°C. Cu additions are quite effective to accelerate the βFeSi2 transformation rate [1997Yam]. The process is two orders shorter with small amount of Cu added. At 600°C, the time for complete βFeSi2 formation is 8 ·103 s [1999Yam]. Similar observations were made by [2001Nog] when Cu is added to the thermoelectric material Fe0.91Mn0.01Siy (y = 2, 2.5) synthesized by spark plasma sintering. Cu has also been shown to promote the eutectoid reaction αFeSi2 ⇌ βFeSi2 + (Si). The viscosity of Cu-Fe-Si liquid alloys has been measured in the composition range 0 to 5 mass% Cu and 10–20 mass% Si, which corresponds to the mean composition of alloys formed in the municipal wastes incinerators [2001Was]. The viscosity is in the range of 5 to 7 mPa·s, which is approximately twice that of pure Cu at the same temperature. Radioactive Cu and Fe ions implanted under 60 keV in Si single crystal at room temperature occupy lattice sites around 50 pm from substitutional positions [2006Wah], and are released by annealing at 300–600°C.
DOI: 10.1007/978-3-540-78644-3_3 # Springer 2008
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Cu–Fe–Si Table 1.
5
Investigations of the Cu-Fe-Si Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1933Rol]
Optical microscopy, solubility measurements
900°C, < 20 mass% Cu, < 20 mass% Si in Fe (Iron rich corner)
[1934Han]
Optical microscopy, solubility measurements
800–1100°C, < 8 mass% Fe, < 8 mass% Si (copper rich corner)
[1943And]
X-ray diffraction (XRD), chemical analysis
< 1025°C, Cu, <10 at.% Fe, < 10 at.% Si
[1953Vog]
Optical microscopy, thermal analysis
600–1500°C, Cu-Fe-FeSi-Cu3Si domain
[1975Abr]
XRD, thermal analysis, metallography
600–1200°C, FeSi2-Cu3Si quasibinary section
[1981Has]
Diffusion couple experiment, α-γ equilibrium
1000–1300°C, < 6 at.% Si, < 6 at.% Cu (Fe rich corner)
[1999Hin]
Optical microscopy, electron probe microanalysis
1250–1450°C, the whole diagram
[2002Wan]
SEM, electron dispersive X-ray analysis
800–1000°C, < 25 mass% Si, (αFe) + (Cu) equilibrium
[2002Was]
Drop calorimetry, heat content measurements
1500–1800°C, < 50 at.% Cu, < 25 at.% Si
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αFe) (Ferrite) < 912°C
cI2 Im 3m W
a = 286.65
pure Fe at 25°C [Mas2, V-C2] dissolves up to 10 at.% Si at 500°C [Mas2]. Dissolves 1.4 at.% Cu at 847°C [2007Tur] at 1390°C [V-C2] dissolves up to 19.5 at.% Si at 1280°C [Mas2]. Dissolves 5.6 at.% Cu at 1487°C [2007Tur]
(δFe) 1538–1394
a = 293.15
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, > 13 GPa [Mas2]
(Si) < 1414
cF8 Fd 3m C (diamond)
a = 543.06
at 25°C [Mas2]. 0 to 0.003 at.% Cu [2002Leb]. Practically no solubility for Fe [Mas2].
(γFe,Cu)
cF4 Fm 3m Cu
a = 364.67
at 915°C [V-C2, Mas2]. Dissolves 14.2 at.% Cu at 1425°C [2007Tur]. Dissolves up to 3.8 at.% Si at 1150°C [2005Mec]
(γFe) (Austenite) < 1394–912
(continued) Landolt-Börnstein New Series IV/11D4
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Cu–Fe–Si
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Cu) < 1084.62
Lattice Parameters [pm]
Comments/References
a = 361.46
pure Cu at 25°C [Mas2] melting point [2007Tur] dissolves up to 4.6 at.% Fe at 1096°C [2007Tur] and 11.25 at.% Si at 842°C [2002Leb] below critical temperature: γ - paramagnetic, Fe enriched
FeSi < 1410
cP8 P213 FeSi
a = 451.7 ± 0.5
at 300°C [V-C2, Mas2]
Fe2Si 1212–1040
cP2 Pm 3m CsCl or hP6 P 3m1 Fe2Si
a = 282
66 at.% Si [Mas2, V-C2]
Fe5Si3 1060–825
hP16 P63/mcm Mn5Si3
a = 675.9 ± 0.5 c = 472.0 ± 0.5
38.4 at.% Si [V-C2, Mas2]
βFeSi2(h) 1220–937
tP3 P4/mmm βFeSi2
a = 269.01 c = 513.4
69.5 to 73.5 at.% Si [Mas2, V-C2]
αFeSi2(r) < 982
oC48 Cmca αFeSi2
a = 986.3 ± 0.7 b = 779.1 ± 0.6 c = 783.3 ± 0.6
66.7 at.% Si [Mas2, V-C2]
α1, Fe3Si
cF16 Fm 3m BiF3
a = 565
10 to 30 at.% Si [Mas2, V-C2]
α2, Fe3Si
cP2 Pm 3m CsCl
-
10 to 22 at.% Si [Mas2]
κ, Cu7Si 842–552
hP2 P63/mmc Mg
β, Cu6Si 853–787
cI2 Im 3m W
δ, Cu5Si(h) 824–711
t**
a = 405.2 ± 0.2 c = 508.55 ± 0.3
11.05 to 14.5 at.% Si [2002Leb] at 12.75 at.% Si
a = 256.05 c = 418.46
14.2 to 16.2 at.% Si [2002Leb] at 14.9 at.% Si
a = 285.4
17.6 to 19.6 at.% Si [2002Leb] sample annealed at 700°C
a = 881.5 c = 790.3
(continued)
DOI: 10.1007/978-3-540-78644-3_3 # Springer 2008
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Cu–Fe–Si
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Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
ε, Cu15Si4 < 800
cI76 I 43d Cu15Si4
a = 961.5
21.2 at.% Si [V-C2]
η, Cu3Si(h2) 859–558
hR* R 3m or t**
23.4 to 24.9 at.% Si [2002Leb] a = 402.2 c = 243.7 a = 726.7 c = 789.2
[V-C2]
η’, Cu3Si(h1) 620–467
hR* R 3m
a = 699.98 c = 731.52
23.2 to 25.1 at.% Si [2002Leb]
η”, Cu3Si(r) < 570
o**
a = 767.6 b = 700 c = 2194
23.3 to 24.9 at.% Si [2002Leb]
Cu5Si
t**
a = 647 c = 873
at 17 at.% Si metastable [2002Leb]
* Fe63Cu27Si10 c**
a = 423
metastable [1992Liu]
Table 3.
Invariant Equilibria
Reaction
T [°C]
Type
Phase
Composition (mass%) Cu
Fe
Si
-
-
-
L1 + (γFe) ⇌ (αFe) + L2
1450–1350
U1
-
L + (γFe) ⇌ (Cu) + (αFe)
1068
U2
L (γFe) (Cu) (αFe)
93.5 7.5 96 2
1.5 92.0 0.5 96.5
5 0.5 3.5 1.5
L + (αFe) ⇌ Fe5Si3+ (Cu)
900–1068
U3
L (αFe) Fe5Si3 (Cu)
91.5 5 0 95
1 79 76.8 0.2
7.5 16 23.2 4.8
L + Fe5Si3 ⇌ FeSi + (Cu)
900–1068
U4
L Fe5Si3 FeSi (Cu)
91.4 0 0 95
1 76.8 66.5 0.1
7.6 23.2 33.5 4.9
L ⇌ FeSi + Cu3Si
847
e
L FeSi Cu3Si
84.5 0 87.16
2.0 66.54 0
13.5 33.46 12.84
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8 Table 4.
Cu–Fe–Si Investigations of the Cu-Fe-Si Materials Properties
Reference
Method/Experimental Technique
Type of Property/Conditions
[1930Cor]
Hardness and resistivity
< 900°C, < 2.5 mass% Fe, < 1 mass% Si
[1965Ued]
XRD, micrographic observation, hardness measurements
2.8 mass% Si, < 0.3 mass% Fe, cold rolled then annealed (< 800°C)
[1983Str]
Transmission Electron Microscopy (TEM), Mössbauer
< 950°C, < 2.5 mass% Fe, < 1 mass% Si, hardening
[1992Liu]
Energy dispersive spectroscopy, Scanning Electron Microscopy (SEM)
77 and 298 K, bombardment of Xe ions on thin layer
[1997Yam, 1999Yam]
Electron probe microanalysis, SEM, backscattered electrons
600–1050°C, FeSi2 + Cu (< 1 at.% Cu), synthesis of pure βFeSi2
[2001Cas]
Auger electron spectroscopy, low energy electron diffraction
Fe/Cu/Si thin layers (between 1 and 50 monolayers of Fe)
[2001Lee]
XRD, SEM, EPMA, DTA (microstructural analysis)
Cu + 50 vol% Si + 0 to 10 mass% Fe, compacted powders
[2001Nog]
XRD, electrical resistivity, seebeck coefficient
< 800°C, Fe0.91Mn0.09Si2.5 + Cu (< 20 mass% Cu)
[2001Was]
Viscosity measurements
1250–1400, < 3 mass% Cu, < 20 mass% Si
[2004Wan]
Thermal analysis, optical microscopy
1800°C, 50Cu-50Fe + 1 to 2 mass% Si, two liquids observations
[2005Yam]
XRD, SEM, EPMA, Differential Scanning Calorimetry (DSC)
1000–1800°C, 30 to 70 at.% Cu, 10 at.% Si, < 1 at.% B
[2006Wah]
XRD, activation analysis
Cu and Fe implanted in Si single crystal, then annealed (300–900°C)
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Cu–Fe–Si
Fig. 1. Cu-Fe-Si.
Landolt-Börnstein New Series IV/11D4
9
FeSi-Cu3Si quasibinary section
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Fig. 2. Cu-Fe-Si.
Cu–Fe–Si
Isothermal section at 1450°C
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Cu–Fe–Si
Fig. 3. Cu-Fe-Si.
Landolt-Börnstein New Series IV/11D4
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Isothermal section at 1350°C
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Fig. 4. Cu-Fe-Si.
Cu–Fe–Si
Isothermal section at 1100°C
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Cu–Fe–Si
Fig. 5. Cu-Fe-Si.
Landolt-Börnstein New Series IV/11D4
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Isothermal section at 900°C
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Fig. 6. Cu-Fe-Si.
Cu–Fe–Si
Heat content (H1773 – H298) in kJ·mol–1 of the liquid alloys at 1500°C
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Cu–Fe–Si
Fig. 7. Cu-Fe-Si.
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Heat content (H2073 – H298) in kJ·mol–1 of the liquid alloys at 1800°C
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16
Cu–Fe–Si
References [1927Cor] Corson, M.G., “Copper Hardened by New Method. What the Corson Alloys Are - Stronger Cable Wire Possible - Many Uses Suggested - Silicon-Aluminum and Silver-Silicon Alloys”, Trans. Amer. Inst. Min. Met. Eng., Inst. Metals Div., 421–424 (1927) (Experimental, Mechan. Prop., Phase Diagram, Phase Relations, 0) [1930Cor] Corson, M.G., “Systems of Copper Alloys with Phase and Variable Limits. Their Use for Copper Hardening” (in French), Rev. Metall., 27, 194–213 (1930) (Experimental, Mechan. Prop., Phase Relations, 6) [1933Rol] Roll, F., “Influence of Ni and Si on the Solubilities in the Fe-Cu System in the Solid State” (in German), Z. Anorg. Chem., 212, 61–64 (1933) (Phase Diagram, Phase Relations, Experimental, 15) [1934Han] Hanson, D., West, E.G., “The Constitution of Copper-Iron-Silicon Alloys”, J. Inst. Met., 54, 229–253 (1934) (Experimental, Phase Diagram, Phase Relations, 9) [1943And] Andersen, A.G.H., Kingsbury, A.W., “Phase Diagram of the Copper-Iron-Silicon System from 90 to 100 % Copper”, Trans. Am. Inst. Mining Metallurgy Pet. Engineering, 152, 38–47 (1943) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 21) [1949Jae] Jaenecke, E., “Cu-Fe-Si” (in German) in “Kurzgefasstes Handbuch aller Legierungen”, Winter Verlag Heidelberg, 660–661 (1949) (Phase Diagram, Phase Relations, Review, 2) [1953Vog] Vogel, R., Horstmann, D., “The Diagram of State Iron-Iron Silicide-Copper Silicide-Copper” (in German), Arch. Eisenhuettenwes., 24(9), 435–440 (1953) (Experimental, Phase Diagram, Phase Relations, Morphology, 6) [1965Ued] Ueda, J., Yamane, T., “A Note on the Effect of Iron on the Annealing Behavior of a Cold Rolled Copper-Silicon Alloy”, Trans. Jpn. Inst. Met., 6(3), 154–158 (1965) (Experimental, Mechan. Prop., Morphology, 7) [1974Sig] Sigworth, G.K., Elliott, J.F., “The Thermodynamics of Liquid Dilute Iron Alloys”, Met. Sci., 8, 298–310 (1974) (Review, Thermodyn., Calculation, 249) [1975Abr] Abrikosov, N.Kh., Petrova, L.I., “Preparation of Crystals of the Low-Temperature Modification of FeSi2”, Inorg. Mater., 11(2), 184–186 (1975), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 11(2), 223–225 (1975) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, 7) [1979Cha] Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Fe-Si” in “INCRA Monograph. Series 6. Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Systems”, Univ. Wisconsin-Milwaukee, U.S.A., 492–497 (1979) (Crys. Structure, Phase Diagram, Phase Relations, Review, Thermodyn., 6) [1981Has] Hasebe, M., Nishizawa, T., “Further Study on Phase Diagram of The Iron-Copper System”, Calphad, 5(2), 105–108 (1981) (Experimental, Phase Relations, Calculation, 4) [1983Str] Stroz, D., Panek, T., Morawiec, H., “Influence of the Addition of a Third Element to the Precipitation Process in Cu-Fe Alloys”, Mater. Sci. Engineering, 58(1), 43–53 (1983) (Experimental, Mechan. Prop., 20) [1986Ole] Olesinski, R.W., Abbaschian, G.J., “The Cu-Si (Copper-Silicon) System”, Bull. Alloy Phase Diagrams, 7(2), 170–178 (1986) (Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., Review, #, 230) [1991But] Butt, M.T.Z., Bodsworth, C., “Liquid Immiscibility in Ternary Metallic Systems”, Mater. Sci. Technol., 7(9), 795–802 (1991) (Theory, Phase Relations, Review, 39) [1992Liu] Liu, B.X., Shang, C.H., Liu, C.H., “Metastable Phases Formed in Fe-Cu-Si Multilayered Films by Ion Irradiation”, J. Mater. Res., 7(1), 85–88 (1992) (Crys. Structure, Experimental, 24) [1992Rag] Raghavan, V., “The Cu-Fe-Si (Copper-Iron-Silicon) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Metals, Calcutta, 6B, 759–767 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Review, 6) [1997Oht] Ohtani, H., Suda, H., Ishida, K., “Solid/Liquid Equilibria in Fe-Cu Based Ternary Systems”, ISIJ Int., 37(3), 207–216 (1997) (Calculation, Phase Relations, Review, Thermodyn., 47)
DOI: 10.1007/978-3-540-78644-3_3 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Cu–Fe–Si [1997Yam]
[1998Mie] [1999Hin]
[1999Yam]
[2001Cas]
[2001Lee]
[2001Nog]
[2001Was]
[2002Leb]
[2002Rag] [2002Wan]
[2002Was]
[2003Mie] [2003Ono]
[2004Wan]
[2005Mec]
[2005Yam]
Landolt-Börnstein New Series IV/11D4
17
Yamauchi, I., Suganuma, A., Okamoto, T., Ohnaka, I., “Effect of Copper Addition on the β-phase Formation Rate in FeSi2 Thermoelectric Materials”, J. Mater. Sci., 32(17), 4603– 4611 (1997) (Experimental, Kinetics, Phase Relations, Electr. Prop., 17) Miettinen, J., “Reassessed Thermodynamic Solution Phase Data for Ternary Fe-Si-C System”, Calphad, 22(2), 231–256 (1998) (Calculation, Assessment, Thermodyn., 36) Hino, M., Nagasaka, T., Washizu, T., “Phase Diagram of Fe-Cu-Si Ternary System Above 1523 K”, J. Phase Equilib., 20(3), 179–186 (1999) (Calculation, Experimental, Phase Diagrams, Phase Relations, 13) Yamauchi, I., Nagase, T., Ohnaka, I., “Temperature Dependence of β-Phase Transformation in Cu Added Fe2Si5 Thermoelectric Material”, J. Alloys Compd., 292, 181–190 (1999) (Experimental, Kinetics, Phase Relations, 10) Castrucci, P., Gunnella, R., Bernardini, R., Montecchiari, A., Carboni, R., De Crescenzi, M., “Epitaxy of Fe/Cu/Si(111) Ultrathin Films: an Auger Electron Diffraction Study”, Surf. Sci., 482(2), 916–921 (2001) (Experimental, Morphology, 18) Lee, Y.-F., Lee, S.-L., Huang, C.-H., Lee, C.K., “Effects of Fe Additive on Properties of Si Reinforced Copper Matrix Composites Fabricated by Vacuum Infiltration”, Powder Met., 44(4), 339–343 (2001) (Morphology, Electr. Prop., Phys. Prop., Experimental, 15) Nogi, K., Kita, T., “Influence of Copper Addition to α-Fe2Si5 on Thermoelectric Properties of Iron-Silicides Produced by Spark-Plasma Sintering”, Mater. Trans., 42(5), 862–869 (2001) (Morphology, Electr. Prop., Phys. Prop., Experimental, 14) Washizu, T., Nagasaka, T., Hino, M., “Viscosity of Liquid Fe-Cu-Si Alloy Formed in New Melting Process for Domestic Waste Incineration Residue”, Mat. Trans., 42(3), 471–477 (2001) (Experimental, Transport Phenomena, 28) Lebrun, N., Dobatkina, T., Kuznetsov, V., “Cu-Si (Copper-Silicon)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.12505.1.20, (2002) (Crys. Structure, Phase Diagram, Assessment, 23) Raghavan, V., “Cu-Fe-Si (Copper-Iron-Silicon)”, J. Phase Equilib., 23(3), 267–270 (2002) (Phase Diagram, Phase Relations, Review, 9) Wang, C.P., Liu, X.J., Ohnuma, I., Kainuma, R., Ishida, K., “Phase Equilibria in Fe-Cu-X (X: Co, Cr, Si, V) Ternary Systems”, J. Phase Equilib., 23(3), 236–245 (2002) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 38) Washizu, T., Nagasaka, T., Hino, M., “Heat Content of Liquid Fe-Cu-Si Alloys Formed in the Melting Treatment Process of Domestic Waste Incineration Residue”, Z. Metallkd., 93(4), 281–287 (2002) (Phase Relations, Thermodyn., Experimental, 22) Miettinen, J., “Thermodynamic Description of the Cu-Fe-Si System at the Cu-Fe Side”, Calphad, 27(4), 389–394 (2003) (Assessment, Phase Diagram, Phase Relations, Thermodyn., 26) Ono-Nakazato, H., Taguchi, K., Usui, T, “Estimation of the Evaporation Rate of Copper and Tin from Molten Iron-Silicon Alloy”, ISIJ Int., 43(7), 1105–1107 (2003) (Thermodyn., Kinetics, Calculation, 14) Wang, C.P., Liu, X.J., Takaku, Y., Ohnuma, I., Kainuma, R., Ishida, K., “Formation of Core-Type Macroscopic Morphologies in Cu-Fe Base Alloys With Liquid Miscibility Gap”, Metall. Mater. Trans. A, 35A(4), 1243–1253 (2004) (Calculation, Experimental, Morphology, Phase Relations, Thermodyn., 31) Meco, H., Napolitano, R.E., “Liquidus and Solidus Boundaries in the Vinicity of OrderDisorder Transitions in Fe-Si System”, Scripta Mater., 52, 221–226 (2005) (Experimental, Calculation, Phase Diagram, Phase Relations, Thermodyn., 30) Yamauchi, I., Irie, T., Sakaguchi, H., “Metastable Liquid Separation in Undercooled Fe-Cu and Fe-Cu-Si Melts Containing a Small B Concentration and Their Solidification Structure”, J. Alloys Compd., 403, 211–216 (2005) (Experimental, Morphology, Phase Relations, 10)
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DOI: 10.1007/978-3-540-78644-3_3 # Springer 2008
18 [2006Wah]
[2007Tur]
[Mas2] [V-C2]
Cu–Fe–Si Wahl, U., Correia, J.G., Rita, E., Araujo, J.P., Soares, J.C., “Fe and Cu in Si: Lattice Sites and Trapping at Implantation-Related Defects”, Nucl. Instrum. Methods Phys. Res./ B, 253, 167–171 (2006) (Experimental, Crys. Structure, Electronic Structure, 23) Turchanin, M., Agraval, P., “Cu-Fe (Copper-Iron)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.11107.1 (2007) (Phase Diagram, Crys. Structure, Thermodyn., Assessment, 31) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_3 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Cu–Fe– Ti
1
Copper – Iron – Titanium Tamara Velikanova, Mikhail Turchanin
Introduction Practical interest in the Cu-Fe-Ti system results from the hydrogen storage capacity of the alloys, the possibility to obtain amorphous materials by ultrarapid quenching, the high-strength Cu based alloys by precipitation hardening, and the promising application of diffusion bonding of titanium and iron parts using copper as an interlayer. A partial liquidus surface of the Cu-Fe-Ti system was first published by [1968Tak]. Later a number of experimental works [1970Kha, 1971Kha, 1971Zak, 1993Min] were devoted to the study of phase relationships in the Cu rich corner of the system. The results of [1968Tak, 1970Kha, 1971Kha, 1971Zak] were critically reviewed by [1979Dri, 1979Cha, 1992Rag]. Phase structures along the line TiCu1–xFex at 900°C were studied by [1989Wu] using Mössbauer spectroscopy. Experimental investigations of the TiCu-TiFe [1993Ali] and Ti2Cu-TiFe [1994Ali2] sections as well as the Ti-TiCu-TiFe [1994Ali1] part of the phase diagram were stimulated by the discovery of glass formation by ultrarapid quenching of liquid alloys. An isothermal section at 850°C of the whole Cu-Fe-Ti system was established by [1995Bee] using a special diffusion technique supplemented by conventional equilibration of alloys. Five ternary phases were reported to exist as equilibrium phases at 850°C. The results of [1993Ali, 1994Ali1, 1994Ali2, 1995Bee] were critically reviewed by [2002Rag]. Thermodynamic properties only for liquid alloys were investigated by [2006Abd]. The experimental investigations of phase relationships and thermodynamic properties of phases in the Cu-Fe-Ti system are summarized in Table 1. Binary Systems The binary systems Cu-Fe and Cu-Ti are accepted from [2007Tur] and [2002Ans] respectively. The Fe-Ti binary phase diagram is taken from [Mas2]. Solid Phases The crystal structure data of solid phases are listed in Table 2. Five ternary phases, denoted as τ1, τ2, τ3, τ4 and τ5, were reported to exist in the system at 850°C by [1995Bee]. They were found in the experiments which combined a method of diffusion couples and the traditional methods of investigation of equilibrated alloys. Two of the ternary phases with unknown crystal structure, τ4 and τ5, were reported to have no isostructural compounds in the related binary systems. The ternary τ1, τ2 and τ3 phases are reported to be based on the TiCu2, Ti2Cu3 and Ti3Cu4 binary phases, respectively. A continuous composition range of τ2 from Ti2Cu3 to Ti0.4Fe0.17Cu0.43 is not excluded (it is shown by a dashed line in Fig. 10) after [1995Bee]. The formation of the ternary phases τ2, τ3 and τ5 in the diffusion interfaces are confirmed in the work of [2005Kun] by the SEM-BSE method at the temperatures 850 and 900°C. However, at 950°C Cu-Fe-Ti and Cu-Ti based intermetallics were not found [2005Kun]. According to the results from the diffusion couple experiments supported by conventionally equilibrated alloys, TiFe dissolves up to 38 at.% Cu at 850°C [1995Bee]. This value is in excellent agreement with the value 37.5 at.% reported by [1989Wu] for alloys annealed 100 h at 900°C. [1994Ali1], however, deduced a much lower value of 25 at.% at 930°C from the study of alloys annealed for 50 h at 900°C. In view of this discrepancy, further investigations are needed to obtain a more precise value. The solubilities of Fe in Ti2Cu and TiCu were measured to be about 2 at.% Fe [1995Bee] and about 1 at.% Fe [1989Wu], respectively, via diffusion couple experiments. The homogeneity ranges of the ternary phases are given in Table 2. For all of them the variation of the Ti content is very small. For (Fe1–xCux)Ti [1994Ali1] reported a maximal Ti content of about 54 at.% Ti, whereas [1995Bee] has drawn it with less than 51 at.% Ti. The drawings of [1995Bee] agree much better with Schreinemakers’ rule than those of [1994Ali1]. Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_4 # Springer 2008
2
Cu–Fe– Ti
Cu decreases the solubility of Ti in (γFe) and Fe decreases the solubility of Ti in (Cu). Substitution of Cu by Fe and Ti atoms leads to an expansion of the cubic lattice of the Cu based solid solution as shown in Fig. 1 [1970Kha]. At the same time, substitution of the Fe atoms by Cu in TiFe leads to an expansion of its cubic lattice as it was found by [1989Wu, 1995Bee]. Figure 2 demonstrates the linear dependence of the lattice parameter of the CsCl type structure of TiFe solid solution as a function of the Fe content after [1989Wu]. Quasibinary Systems According to [1994Ali1], the TiFe-TiCu section is a quasibinary system with a eutectic composition of 5 at. % Fe and a eutectic temperature of 930°C. This statement is not compatible with the liquidus proposed by the same authors for the TiFe-Ti-TiCu region and the corresponding reaction scheme. There the abovementioned eutectic temperature is not a maximal one on the monovariant line of the equilibrium L ⇌ TiFe + TiCu, which is a prerequisite for a true quasibinary system. Because of the same reason, the Cu-TiFe2 section can not be quasibinary, as it was proposed by [1970Kha], notwithstanding on the reported constant solidus temperature for the alloys along this section. The corresponding vertical sections are given in the section “Temperature-Composition Sections” of this assessment. There was no tie line found by [1970Kha, 1971Zak] with a maximal temperature on the solidus surface (Cu) + TiFe2. The solidus temperature of this phase field decreases continuously from 1080°C at the L + (γFe) ⇌ (Cu) + TiFe2 equilibrium with increasing Ti content. Invariant Equilibria Invariant equilibria are studied systematically only for the Ti rich alloys in the TiFe-Ti-TiCu region by [1994Ali1]. The results [1994Ali1] are given in the reaction scheme in Fig. 3. Outside this region an invariant reaction L + (γFe) ⇌ (Cu) + TiFe2 at 1080°C is reported by [1970Kha]. Liquidus and Solidus Surfaces The liquidus and solidus surface projections for the Ti rich part of the system and TiFe+Ti+TiCu region, reported by [1994Ali1], are given in Figs. 4 and 5, respectively. The first liquidus isotherms of the region TiFe-Ti-TiCu were reported by [1968Tak], based on direct thermal analysis. A eutectic reaction L ⇌ (βTi) + TiFe + Ti2Cu at the Ti content of about 65 at.% and at 950°C was proposed. These data are close to those of [1994Ali1], but do not coincide. The supposed eutectic temperature is lower than that of the incongruent U1 process, but it is higher than the temperature of the quasibinary TiFe-TiCu eutectic after [1994Ali1]. Taking into account this circumstance the above mentioned ternary eutectic could be realized only if a quasibinary TiFe-Ti2Cu eutectic exists in the system above 950°C. Although the latter is not excluded, the data of such kind so far are not available. Isotherms of the solidus surface of the Cu rich region (95 to 100 mass% Cu) were reported by [1971Zak] indicating also the L + (βTi) + TiFe + Ti2Cu four-phase field. The Cu corner of the solidus surface is shown in Fig. 6, with the four-phase equilibrium L + (Cu) + Ti2Fe + TiFe tentatively added. The equilibria involving τ3 and τ2 are questionable because the temperature intervals of their stability are not known. The solidus temperatures were determined by [1971Zak] using metallographic investigation of alloys tempered in steps of 20°C and subsequently water quenched. Thus, the data precision can be estimated to be ±10°C. Isothermal Sections Isothermal sections at 900, 850 and 650°C for the TiFe-Ti-TiCu part of the system reported by [1994Ali1] are given in Figs. 7, 8, and 9, respectively. The alloys prepared by arc-melting were heat treated in the following process: at 900°C for 50 h and water quenched, 850°C for 100 h and water quenched, 650°C for 150 h and followed by air-cooling. Because the homogenization temperature is rather close to the solidus, one can suppose that equilibrium was achieved. No ternary compounds are found in the Ti corner of the system. The phase relations at 850°C in the entire composition range were measured by [1995Bee] using diffusion couples. Most of the critical conclusions were confirmed by conventional equilibration of alloys. Annealing times and phases found are well documented. The isothermal section at 850°C is given in Fig. 10 DOI: 10.1007/978-3-540-78644-3_4 # Springer 2008
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Cu–Fe– Ti
3
after [1995Bee]. Five ternary phases are found to exist at 850°C. A “ternary” τ1 phase, which is isostructural with the binary TiCu2, is treated by [1995Bee] as a solid solution of TiCu2 stabilized in the ternary system by some addition of Fe at a lower temperature than it is stable in the Cu-Ti binary system. The crystal structures of the phases τ2 and τ3 are closely related to Ti2Cu3 and Ti3Cu4, respectively. It is not excluded that the τ2 phase forms a continuous solid solution with the binary Ti2Cu3 phase, which is indicated in Fig. 10 by a thick dashed line. The τ4 and τ5 phases are without doubt true ternary compounds. The topology of the phase equilibria in the TiFe-Ti-TiCu region after [1995Bee] agrees qualitatively with that of [1994Ali1]. However, there is a noticeable difference concerning the compositions of the phases in the ternary equilibria, especially the solubility of Cu in TiFe after [1995Bee] is much larger and agrees very well with [1989Wu], whereas the Ti content of TiFe after [1994Ali1] reaches a maximum of about 55 at.%, but less than 51 at.% after [1995Bee]. Although the results of [1995Bee] are supported by those of [1989Wu], this question needs further experimental verification. In agreement with both [1994Ali1] and [1995Bee], the alloy Fe-8Cu-2Ti (mass%) compacted by sintering at 1350°C was found to be in the two-phase region of TiFe2+(αFe) [1999Pie]. Three isothermal sections at 900, 850 and 650°C in the Cu corner of the ternary system have been determined [1970Kha, 1971Kha, 1971Zak]. The isothermal sections at 900 and 650°C are given in Figs. 11, 12 with some corrections mainly according to the accepted binary phase diagrams. [1992Rag] in his review published a partial isothermal section from Cu-Fe up to 40 at.% Ti at 851°C based on the experimental data of [1970Kha]. Equilibrium between TiFe2 and Cu4Ti assumed in this section is in contradiction to [1995Bee] who assumed equilibria between TiFe2 and (Cu) as well as between TiCu and (Cu). Therefore the isothermal section proposed by [1992Rag] can not be accepted. Temperature – Composition Sections Five temperature-composition sections for the TiFe-Ti-TiCu part are published by [1993Ali, 1994Ali1, 1994Ali2]. Four of them are: at Fe:Cu = 3:1, Fe:Cu = 1:1, Fe:Cu = 1:3 [1994Ali1] and TiCu-TiFe [1993Ali]. The vertical section TiCu-TiFe is given in Fig. 13. It is treated by [1993Ali] as a quasibinary system of eutectic type (with eutectic temperature 930°C), which is not reasonable because its temperature is less than the temperature of the adjoining invariant equilibrium L + Ti2Fe ⇌ TiFe + TiCu of 950°C according to [1994Ali1]. In Fig. 13 this vertical section is given after data of [1993Ali] with minor correction made to account for the above remark: the three-phase field is added in the section. The vertical section Cu-TiFe2 is given in Fig. 14. According to [1970Kha] the section coincides with the line Cu-TiFe2, where the TiFe2 phase is of congruent composition. Therefore the section is considered by the authors to be a quasibinary system. However, no temperature maximum is shown on the solidus of the (Cu) + TiFe2 two-phase field in the vertical section at 95 mass% Cu by [1971Kha], which must exist in the case of a quasibinary L ⇌ (Cu) + TiFe2 equilibrium. The section Cu-TiFe2 given in Fig. 14 differs from that of [1970Kha] by adding the three-phase field L + (Cu) + TiFe2 according to the above remark. Thermodynamics Experimental information on the thermodynamic properties of ternary phases is available only for liquid alloys. In [2006Abd], the mixing enthalpy of liquid alloys was investigated at 1600°C and xTi < 0.6 by high temperature calorimetry. The investigation was carried out along three sections at xFe/xCu = 1/3, xFe/xCu = 1, and xFe/xCu = 3. The experiments show that the integral mixing enthalpy has negative values for alloys with xTi > 0.2. The maximum mixing enthalpy is related with the Cu-Fe binary system, the value is 10.8 kJ·mol–1 at xFe = 0.42 [2007Tur], its minimum value reaches –20.9 kJ·mol–1 at xTi = 0.47 in the binary Fe-Ti system [2006Abd]. The isotherm of integral mixing enthalpy at 1600°C according to [2006Abd] is shown in Fig. 15. The thermodynamic properties of the liquid phase were modeled by the ideal associated solution model, assuming a set of associates: TiCu, Ti2Cu, TiFe2, TiFe, and TiFeCu3. Notes on Materials Properties and Applications According to [1970Kha] the microhardness of alloys at 95 mass% Cu has a minimum at the composition of the crossing point with the Cu-Ti2Fe section. At the same composition the electrical conductivity shows a Landolt-Börnstein New Series IV/11D4
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Cu–Fe– Ti
maximum value. The microhardness of alloys along the Cu-TiFe2 section quenched from 650, 800, 850 and 900°C [1970Kha] is shown in Fig. 16. Mechanical properties of Cu alloys doped by Ti and Fe depending on the conditions of homogenization, quenching, preliminary deformation and final thermal treatment, are reported by [1983Hae]. The microhardness of alloys along the TiCu-TiFe section was studied by [1993Ali]. The microhardness of an annealed alloy at the eutectic composition Cu45Ti50Fe5 was found as 4000 MPa. An amorphous alloy of the same composition shows a microhardness value of 6100 MPa along with an ultimate strength of 2300 MPa. A study on electrical conductivity and Hall constant [1993Ali] demonstrates the change in electrical conductivity during the crystallization process of amorphous alloys. Two dilute alloys of the compositions Cu-0.8Fe-0.4Ti (at.%) and Cu-0.4Fe-0.4Ti (at.%) were studied in the temperature range from 500 to 950°C by electrical conductivity measurements. The electrical conductivities of the two alloys are much higher than theoretical values calculated assuming all the elements to be dissolved. The dependence of electrical conductivity on aging temperature was employed to evaluate the activation energy for the decomposition of the Cu matrix, which was found to be about 200 kJ·mol–1. Mechanical properties of an arc-melted alloy of the Ti-15Fe-15Cu (mass%) composition were studied by [2005Lou]. The alloy shows a good combination of compressive mechanical properties: Young’s modulus of 121 GPa, compressive fracture strength of 1.61 GPa, 0.2% yield strength of 1.53 GPa, and 8% ductility. An addition of B in small quantities (0.5 at.%) increases mechanical strength of the alloy up to 2.47 GPa, but decreases ductility. An addition of Cu improves ductility of the corresponding Fe-Ti binary alloy. The experimental investigations of the Cu-Fe-Ti materials properties are summarized in Table 3. Miscellaneous Amorphous alloys were obtained by rapid quenching of melt. Fully amorphous ribbons of composition Cu100–x–yFexTiy were formed for alloys with 0 ≤ x ≤ 10 and 40 ≤ y ≤ 60 (at.%) [1988Dun]. Amorphous alloys Cu100–xFexTi50 with 0 ≤ x ≤ 20 (at.%) were prepared by melt spinning at cooling rates of 106 K·s–1 [1993Ali]. The crystallization behavior of amorphous alloys was investigated in the above-mentioned works. The hydrogenation behavior of the amorphous Cu45Fe5Ti50 alloy was studied by [1988Dun]. The sintering process of Fe-8Cu-2Ti (mass%) compacts was studied by DTA [1999Pie]. Three different liquid phases formed and disappeared by interdiffusion during heating the compacts up to 1350°C. Diffusion bonding was carried out between commercially pure titanium and stainless steel using copper as interlayer in the temperature range of 850–950°C in vacuum [2005Kun]. The presence of different reaction layers in the diffusion zone was exhibited in the study. The occurrence of different intermetallic compounds such as Ti2Cu, TiCu, Ti2Cu3, Ti3Cu4, TiFe, TiFe2, τ2, τ3, and τ5 was confirmed by X-ray diffraction. A maximum bond strength of 318 MPa was obtained for the couple bonded at 900°C due to a better coalescence of mating surfaces. With rise of the joining temperature to 950°C, the bond strength decreased due to the formation of brittle Fe-Ti based intermetallics. At a lower joining temperature of 850°C, the bond strength was also lower due to incomplete coalescence of mating surfaces. Mössbauer spectra of different alloys of the ternary Cu-Fe-Ti system were studied by [1971Win, 1988Dun, 1989Wu].
Table 1.
Investigations of the Cu-Fe-Ti Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1968Tak]
Thermal analysis, measurements of microhardness
Partial liquidus surface
[1970Kha]
Optical microscopy, X-ray analysis, differential thermal analysis, measurements of microhardness, measurements of electrical conductivity, magnetic metallography
95 to 100 mass% Cu and 650 to 1100°C
(continued) DOI: 10.1007/978-3-540-78644-3_4 # Springer 2008
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Cu–Fe– Ti
5
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1971Kha]
magnetic metallography, magnetic susceptibility measurements
95 to 100 mass% Cu and 650°C
[1971Zak]
Optical microscopy, X-ray analysis, differential thermal analysis, measurements of microhardness, measurements of electroconductivity, magnetic metallography
95 to 100 mass% Cu and 650 to 1100°C
[1989Wu]
Mössbauer spectroscopy, X-ray analysis, EPMA, SEM
Phase structures of TiCu1–xFex alloys at x = 0.1–0.6
[1993Ali]
Optical microscopy, X-ray analysis, EPMA, differential thermal analysis, durometerical analysis, measurements of microhardness, measurements of electrical conductivity and Hall constant
TiCu-TiFe section
[1993Min]
Electrical conductivity measurements, TEM, SEM, STEM EDX microanalysis, field ion microscopy, atom probe microanalysis
Cu 0.8Fe-0.4Ti (at.%) and Cu-0.4Fe-0.4Ti (at.%) at 500 to 950°C
[1994Ali1]
Optical microscopy, X-ray analysis, differential thermal analysis
Ti-TiCu-TiFe portion at 650 to 900°C
[1994Ali2]
Optical microscopy, X-ray analysis, EPMA, thermal analysis, measurements of microhardness
Ti2Cu-TiFe section at 900 to 1300°C
[1995Bee]
Diffusion “sandwich” couples technique, optical microscopy, SEM, EPMA, X-ray analysis
Isothermal section at 850°C
[1999Pie]
Differential thermal analysis
Sintering Fe-8Cu-2Ti (mass%) at RT to 1350°C
[2005Kun]
Optical microscopy, X-ray analysis, SEM, mechanical properties testing
Diffusion bonding of Ti/Cu/Fe multilayers
[2005Lou]
X-ray analysis, SEM, mechanical properties measurements, compressive mechanical testing, density measurements
Ti-15Fe-15Cu (mass%) at RT
[2006Abd]
High temperature calorimetry
Liquid alloys at 1600°C and xTi < 0.6
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Cu), TixFeyCu1–x–y < 1095
cF4 Fm 3m Cu
Lattice Parameters [pm]
a = 361.46
Comments/References
at y = 0 and 885°C x = 0.08 [2002Ans] at x = 0 and 1095°C y = 0.05 [2007Tur] at x = 0 y = 0 and 25°C [V-C2, Mas2] (continued)
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Cu–Fe– Ti
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
γ, TixFe1–x–yCuy 1487–847 (γFe) 1394–912
cF4 Fm 3m Cu
α, TixFe1–x–yCuy < 1538
cI2 Im 3m W
(δFe) 1538–1394 (αFe) < 912 (βTi), Ti1–x–yFexCuy 1668–595
Lattice Parameters [pm]
a = 364.68
at x = 0 and 1417°C y = 0.135 [2007Tur] at y = 0 and 1150°C x = 0.008 [Mas2] at x = 0 y = 0 and 915°C [V-C2, Mas2]
a = 293.15
at x at x at y at x
a = 286.65
at x = 0 y = 0 and 25°C [V-C2, Mas2]
cI2 Im 3m W
a = 330.65 (αTi), Ti1–x–yFexCuy < 882
TiFe, Ti1–xFex–yCuy < 1317
hP2 P63/mmc Mg
a = 295.06 c = 468.35
cP2 Pm 3m CsCl a = 297.6 to 298.6
TiFe2, Ti1–xFex–yCuy < 1427
TiCu4, Ti1–xCux < 885
hP12 P63/mmc MgZn2 oP20 Pnma ZrAu2
Comments/References
a = 478.5 c = 779.9
= = = =
0 and 1487°C y = 0.058 0 and 847°C y = 0.016 [2007Tur] 0 and 1289°C x = 0.10 [Mas2] 0 y = 0 and 1390°C [V-C2, Mas2]
at x/y = 3 and 1040°C y + x = 0.18 [1994Ali1] at x/y = 1 and 990°C y + x = 0.17 [1994Ali1] at x/y = 1/3 and 1050°C y+x=0.14 [1994Ali1] at x = 0 and 1005°C y = 0.135 [2002Ans] at y = 0 and 1085°C x = 0.22 [Mas2] pure Ti(h) at T > 882°C [Mas2] at x = 0 and 790°C y = 0.016 [2002Ans] at y = 0 and 700°C x = 0.0005 [Mas2] pure Ti(l) at 25°C [Mas2] at 850°C y = 0 to 0.38 [1995Bee] at 930°C y = 0 to 0.25 [1993Ali] at y = 0 and 1085°C x = 0.475 to 0.503 [Mas2] at y = 0 and x = 0.47 to 0.50 [V-C2] at 850°C y = 0 to 0.02 [1995Bee] at y = 0 x = 0.646 to 0.724 [Mas2] at y = 0 and x = 0.717 [V-C2] x = 0.780 to 0.809 [2002Ans]
a = 452.5 b = 434.1 c = 1295.3
[V-C2]
(continued)
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Cu–Fe– Ti
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
TiCu2, Ti0.33FexCu0.67–x 890–870
oC12 Cmm2 VAu2
7
Lattice Parameters [pm]
Comments/References
[2002Ans]
a = 436.3 b = 797.7 c = 447.8
at x = 0 [V-C2]
τ1, Ti0.33FexCu0.67–x
oC12 Amm2
at 850°C x = 0.01 to 0.025 [1995Bee]
Ti2Cu3, Ti0.4FexCu0.6–x < 875
tP10
[2002Ans]
P4/nmm Ti2Cu3
a = 313 c = 1395
at x = 0 [V-C2]
τ2, Ti0.4FexCu0.6–x Ti3Cu4, Ti0.43FexCu0.57–x < 925
at 850°C x = 0.05 to 0.17 [1995Bee] tI14 I4/mmm Ti3Cu4
[2002Ans] a = 313.0 c = 1994
at x = 0 [V-C2]
τ3, Ti0.43FexCu0.57-x TiCu, Ti1–xFeyCux–y < 982
Ti2Cu, Ti0.67FexCu0.33–x < 1012
at 850°C x = 0.21 to 0.24 [1995Bee] tP4 P4/nmm TiCu
at 930°C y = 0 to 0.03 [1993Ali] at y = 0 x = 0.48 to 0.52 [2002Ans] at y = 0 [V-C2]
a = 310.8 to 311.8 c = 588.7 to 592.1
tI6 I4/mmm MoSi2
[2002Ans] at 850°C x = 0 to 0.02 [1995Bee] a = 295.3 c = 1073.4
at x = 0 [V-C2]
* τ4, Ti0.37FexCu0.63–x
-
-
at 850°C x = 0.05 to 0.07 [1995Bee]
* τ5, Ti0.45FexCu0.55–x
-
-
at 850°C x = 0.04 to 0.05 [1995Bee]
Table 3.
Investigations of the Cu-Fe-Ti Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1970Kha]
Measurements of microhardness and electrical conductivity
Microhardness, electrical conductivity
[1983Hae]
SEM, mechanical testing
Mechanical properties
[1988Dun]
Weight measurements
Hydrogen absorption of amorphous alloys (continued)
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Cu–Fe– Ti
Reference
Method/Experimental Technique
Type of Property
[1993Ali]
Tensile testing machine, measurements of electrical conductivity and Hall constant at heating
Ultimate tensile strength, microhardness, electrical resistively and Hall constant of amorphous alloys
[1993Min]
Electrical measurements with a Sigmatest apparatus
Electrical conductivity
[2005Kun]
Tensile testing machine
Bond strength, ultimate tensile strength, breaking strain
[2005Lou]
Mechanical properties testing Archimedian method
Young's modulus, compressive fracture strength, yield strength, ductility, density
Fig. 1. Cu-Fe-Ti. Lattice parameters of the (Cu) solid solution vs TiFe2 content in the alloys quenched from 900 and 650°C
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Fig. 2. Cu-Fe-Ti. Lattice parameters of the TiFe solid solution vs Fe content
Fig. 3. Cu-Fe-Ti. Partial reaction scheme Landolt-Börnstein New Series IV/11D4
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Cu–Fe– Ti
Fig. 4. Cu-Fe-Ti. Partial liquidus surface projection, Ti corner
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Fig. 5. Cu-Fe-Ti. Partial solidus surface projection, Ti corner
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Cu–Fe– Ti
Fig. 6. Cu-Fe-Ti. Partial solidus surface projection, Cu corner
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Fig. 7. Cu-Fe-Ti. Partial isothermal section at 900°C, Ti corner
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Cu–Fe– Ti
Fig. 8. Cu-Fe-Ti. Partial isothermal section at 850°C, Ti corner
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Fig. 9. Cu-Fe-Ti. Partial isothermal section at 650°C, Ti corner
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Fig. 10. Cu-Fe-Ti.
Cu–Fe– Ti
Isothermal section at 850°C
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Fig. 11. Cu-Fe-Ti.
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Partial isothermal section at 900°C, Cu corner
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Cu–Fe– Ti
Fig. 12. Cu-Fe-Ti. Partial isothermal section at 650°C, Cu corner
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Fig. 13. Cu-Fe-Ti. Temperature - composition section TiCu-TiFe
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Cu–Fe– Ti
Fig. 14. Cu-Fe-Ti. Partial temperature- composition section Cu-TiFe2
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Fig. 15. Cu-Fe-Ti. Integral mixing enthalpy of liquid alloys at 1600°C in kJ·mol–1
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Fig. 16. Cu-Fe-Ti. Microhardness of Cu-TiFe2 alloys, quenched from 650, 800, 850 and 900°C
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References [1968Tak] Takeuchi, Y., Watanabe, M., Yamabe, S., Wada, T., “Titanium- and Zirconium Eutectic Solders” (in German), Metall, 22, 8–15 (1968) (Phase Diagram, Experimental, 11) [1970Kha] Khan, M.G., Zakharov, A.M., Zakharov, M.V., “Study of the Ternary Cu-Fe-Ti Phase Diagram” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., 13(1), 104–109 (1970) (Phase Diagram, Experimental, 8) [1971Win] Window, B., “Mössbauer Studies of Iron in Copper Alloys”, J. Phys. F: Met. Phys., 1, 533–538 (1971) (Crys. Structure, Experimental, 20) [1971Kha] Khan, M.G., Zakharov, M.V., Zakharov, A.M., Zikeeva, V.N., “Identification of a Phase Based on Fe in Cu-Rich Alloys of the Cu-Fe-Ti Ternary System” (in Russian), Izv. Vyss. Uchebn. Zaved., Tsvetn. Metall., 14(4), 98–99 (1971) (Phase Diagram, Experimental, 2) [1971Zak] Zakharov, M.V., Zakharov, A.M., Khan, M.G., “The Cu Corner of the Ternary Cu-Fe-Ti System” (in Russian) in “Diagrammy Sostoyaniya Metallichskikh Sistem”, Mater. 4th Vses. Soveshch., 1969, Ageev, N.V., (Ed.), 139–141 (1971) (Phase Diagram, Experimental, 3) [1979Cha] Chang, Y.A., Neumann, J.P., Mikula, A., Goldberg, D., “Cu-Fe-Ti” in “INCRA Monograph. Series 6. Phase Diagrams and Thermodynamic Properties of Ternary Copper-Metall Systems”, Uni. Wisconsin-Milwaukee, USA, 501–503 (1979) (Phase Diagram, Review, 3) [1979Dri] Drits, M.E., Bochvar, N.R., Guzei, L.S., Lysova, E.V., Padezhnova, Rokhlin, L.L., Turkina, N.I., “Cu-Fe-Ti” (in Russian) in “Binary and Multicomponent Copper-Base System”, Nauka, Moscow, 119–120 (1979) (Phase Diagram, Review, 1) [1983Hae] Haeussler, D., Daut, H.H., “Electron-Microscopic Investigations in Dispersion Hardened Copper” (in German), Freib. Forsch. B, Metall. Werkstofftech., B234, 77–85 (1983) (Morphology, Mechan. Prop., Experimental, 3) [1988Dun] Dunlap, R.A., Stroink, G., Stadnik, Z.M., Dini, K., “Properties of As-Quenched and Hydrogenated Copper-Iron-Titanium Alloys”, Mater. Sci. Eng., 99, 543–546 (1988) (Morphology, Phys. Prop., Experimental, 10) [1989Wu] Wu, C., Li, J., “Phase Structure of the TiCu1–xFex System”, Metall. Trans. A, 20A, 981–985 (1989) (Crys. Structure, Experimental, 8) [1992Rag] Raghavan, V., “The Cu-Fe-Ti (Copper-Iron-Titanium) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Institut of Metals, Calcutta, 6B, 776–780 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Review, 7) [1993Ali] Alisova, S.P., Lutskaya, N.V., Budberg, P.B., Bychkova, E.I., “Phase Constitution of the TiCu-TiNi-TiCo (TiFe) Systems in the Equilibrium and Metastable States”, Russ. Metall. (Engl. Transl.), (3), 205–212 (1993), translated from Izv. Ros. Akad. Nauk, Metally, 3, 222–229 (1993) (Phase Diagram, Experimental, 16) [1993Min] Mineau, L., Hamar-Thibault, S.J., Allibert, C.H., “Precipitation in Cu-Rich Cu-Fe-Ti Ternary Alloys - a Continuous Process?”, Phys. Status Solidi A, 137, 87–100 (1993) (Phase Diagram, Crys. Structure, Experimental, 22) [1994Ali1] Alisova, S.P., Kovneristyi, Yu.K., Budberg, P.B., Lutskaya, N.V., “The Ti-TiFe-TiCu System”, Russ. Metall. (Engl. Transl.), (2), 122–125 (1994), translated from Izv. Ros. Akad. Nauk, (2), 157–161 (1994) (Phase Diagram, Experimental, 6) [1994Ali2] 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), translated from Izv. Ros. Akad. Nauk, (5), 170–172 (1993) (Phase Diagram, Experimental, 8) [1995Bee] van Beek, J.A., Kodentsov, A.A., van Loo, F.J.J., “Phase Equilibria in the Cu-Fe-Ti System at 1123 K”, J. Alloys Compd., 217, 97–103 (1995) (Phase Relations, Phase Diagram, Experimental, 16) [1999Pie] Pieczonka, T., Kaysser, W.A., Petzow, G., “Transient Liquid Phase Sintering of Fe-Cu-Ti Compacts”, J. Mater. Proces. Technology, 92, 21–24 (1999) (Phase Relations, Experimental, 28) [2002Rag] Raghavan, V., “Cu-Fe-Ti (Copper-Iron-Titanium)”, J. Phase Equilib., 23(2), 172–174 (2002) (Crys. Structure, Phase Relations, Review, 9)
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[2005Kun]
[2005Lou]
[2006Abd]
[2007Tur]
[Mas2] [V-C2]
Cu–Fe– Ti Ansara, I., Ivanchenko, V., “Cu-Ti (Copper-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) Kundu, S., Ghosh, M., Laik, A., Bhanumurthy, K., Kale, G.B., Chatterjee, S., “Diffusion Bonding of Commercially Pure Titanium to 304 Stainless Steel Using Copper Interlayer”, Mater. Sci. Eng. A, 407, 154–160 (2005) (Crys. Structure, Morphology, Experimental, 16) Louzguine-Luzgin, D.V., Louzguina-Luzgina, L.V., Inoue, A., “Investigation of the Structure and Properties of Hypereutectic Ti-based Bulk Alloys”, Mater. Res. Soc. Symp. Proc., 842, 133–138 (2005) (Crys. Structure, Mechan. Prop., Morphology, Phase Relations, Experimental, 16) Abdulov, A.R., Turchanin, M.A., Agraval, P.G., “Application of Ideal Associated Solution Model for Prediction of Glass Formation Composition Intervals in Ternary Liquid Alloys” (in Russian), Metallofiz. Noveishie Tekhnol., 28, 1247–1256 (2006) (Thermodyn., Experimental, Calculation) Turchanin, M., Agraval, P., “Cu-Fe (Copper-Iron)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart, Document ID: 20.11107.1 (2007) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Assessment, Phys. Prop., #, 36) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Fe–H–Mn
1
Iron – Hydrogen – Manganese Volodymyr Ivanchenko, Tetyana Pryadko
Introduction The solubility of hydrogen in molten Fe-Mn alloys was measured by [1961Mae] (1620°C and 101 kPa, to 8 mass% Mn), [1963Wei] (1592°C and 101 kPa, to 10 mass% Mn), [1970Fuk, 1970Kat] (at 1600°C to 3 mass% Mn) and [1975Bes] (at 1500, 1600 and 1700°C, to 12 mass% Mn), as listed in Table 1. Theories of hydrogen solubility were detailed in [1965Bur, 1967Bur] and the experimental data were reviewed by [1977Bes] and [1981Sch]. The hydrogen-induced phase transformation of the fcc γ phase was studied by [1978Pon, 1982Ant1, 1982Ant2, 1986Usu, 1993Aik1, 1993Aik2]. Hydrogen saturation increases with the temperature of the ε→γ transformation at high pressures [1978Pon]. It was shown that the hydrogen solubility is higher in the ε phase than in the γ phase regardless of the route of hydrogenation. [1993Aik1, 1993Aik2] studied the hcp ε phase in an Fe-21 at.% Mn alloy in order to compare the stability of the γ and ε phases in a two-phase alloy subjected to cathodic hydrogen charging in an aqueous solution. They showed that the hcp ε phase can always be stabilized more easily than the fcc γ phase, not only in a hydrogen-gas atmosphere at high pressures, as in the case studied by [1978Pon] and [1982Ant2], but also in an aquas solution under cathodic polarization. The effect of hydrogen on the magnetic properties and electroresistivity of an Fe-22.4 at.% Mn alloy was studied by [1982Ant1]. [1991Bol] proposed a theory for the calculation of the distribution coefficients of hydrogen between tetrahedral and octahedral interstices in Fe based alloys. The results of the calculations are in contradiction with the observation that hydrogen is not found in the tetrahedral interstices. The solubility and mobility of hydrogen in solid Fe-Mn alloys were studied experimentally by [1965Sch], and [1966Bur], as discussed by [1977Kar]. Investigations of phase relations in the Fe-H-Mn system, structures and thermodynamics are listed in Table 1. Binary Systems The assessed hydrogen-pressure-temperature Fe-H phase diagram (up to 10 GPa) is taken from [2003Fuk], who accepted the results of [1982Ant2] and [2002Ant]. The Fe-H system under a hydrogen pressure of 10 MPa is accepted from [2002Zin]. The phase relations in this diagram are similar to those of [1990San]. The accepted H-Mn phase diagram is taken from [2002Fuk] that has been established over a wide range of p(H2) (to 7.6 GPa), and T (to 1000°C). A new phase was found, ε’, which has a dhcp structure. This leads to the appearance of a third triple point γ-ε-ε’ at a hydrogen pressure of 3.5 GPa and temperature of 700°C. The curves below 3 GPa have been reproduced from [1996Ant], where the temperature of the ε→γ transition increases linearly with pressure. A fourth triple point, δ+γ+L, was included, but only qualitatively. The Fe-Mn phase diagram is accepted from [2004Wit]. Solid Phases The crystal structures of the unary and binary phases are presented in Table 2. No ternary phase of a crystal structure type different from those of the unary and binary phases was reported. In the Fe-H system within the studied T-p interval, there are 4 different phases, having the bcc, fcc, hcp and dhcp lattice, that are based on the α and γ phases of pure iron and the high pressure ε phase. Studies showed that hydrogen and deuterium occupy octahedral interstices in the close packed iron lattices of both the ε and ε’ phases. Deuterium randomly occupies octahedral interstices in the hcp iron lattice of the ε-FeD0.42. In this respect, the ε phase in the Fe-H(D) system is unlike the non-stoichiometric hcp phases TcH and MnD. The ε hydride, having a hcp iron lattice, is formed as an intermediate metastable phase in the course of the transformation of (αFe) to the stable ε’ hydride at high pressures and elevated temperatures [1998Ant, 2003Fuk, Landolt-Börnstein New Series IV/11D4
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2005Fuk]. The ε’ hydride, having a dhcp iron lattice, is thermodynamically stable at high hydrogen pressures [1991Bad, 1998Ant]. The composition of the ε’ phase is close to the stoichiometry FeH or FeD, and its metal lattice contains many stacking faults [1998Ant]. Hydrogen or deuterium atoms in this phase occupy all available octahedral interstitial sites in the metal lattice. The fcc γ phase is characterized by the formation of superabundant vacancies [2003Fuk, 2005Hir]. Both ε and ε’ hydrides are unstable at ambient conditions and rapidly decompose to (αFe) and molecular hydrogen. In the H-Mn system, there are 5 different phases: the primary solid solutions based on the α and β allotropic modifications of manganese containing only a few atomic percent of hydrogen, a hydride based on the γ fcc modification of manganese and the hydrides with hcp ε and dhcp ε’ metal lattices. The fcc manganese hydride MnHx with x ≈ 0.41 forms at high hydrogen pressures and can be considered as an interstitial solid solution of hydrogen in the γ modification of manganese. The dissolved hydrogen expands the crystal lattice of (γMn) at a rate of dVa/dx ≈ 1.85 Å3/H atom, which is typical of transition metal hydrides [1998Fed]. Under high pressures of hydrogen, manganese forms ε hydrides [1975Pon] with the H/Mn atomic ratio x ranging from 0.65 to 0.96 [1978Bel]. In the ε manganese hydrides MnHx with x ≥ 0.83, hydrogen atoms are randomly distributed over octahedral interstices in the hcp metal lattice (a deficient NiAs type structure) [1987Iro]. At x = 0.65, hydrogen forms a superstructure, presumably of the anti-CdI2 type [1987Iro]. The dhcp ε’ phase was found to exist [2002Fuk] in the region where the structure was previously thought to be hcp [1996Ant, 2002Ant]. The crystallographic data presented by [2002Fuk] indicate that in the course of the εv→⊊ε’⊊→⊊γ phase transitions, both the atomic volume and the spacing between close-packed planes are nearly conserved, and only their stacking sequence is changed: ε (ABAB…) → ε’ (ABAC...) → γ (ABCABC…). There is also a peculiar feature in the kinetics of the phase transitions. The transition between the ε hcp and ε’ dhcp phases is reversible and very fast: transitions in both directions were completed in less than 1 min in the course of stepwise changes of temperature across the phase boundary. In contrast, the ε → γ and ε’ → γ transitions were irreversible. Once the fcc structure has been formed, the reverse transition to the hcp and dhcp structures was sluggish and incomplete. Liquidus, Solidus and Solvus Surfaces The solubility of hydrogen in an Fe-Mn melt at 1620°C and 101 kPa was first measured by [1961Mae], who reported that alloying with Mn decreases the hydrogen solubility from 24 in pure Fe to 9 ppm (mass) at 8 mass% Mn. [1963Wei] found that alloying with Mn increased the hydrogen solubility at 1592°C and 101 kPa from 24.7 ppm for pure Fe to 25.7 ppm (mass) for the Fe-10 mass% Mn alloy. Later, [1970Fuk, 1970Kat] expressed the composition dependence of the hydrogen solubility as {H}alloy = {H}Fe/ 10(–0.0014 XMn) following studies at 1600°C and up to 3 mass% Mn using a sampling method. Their results are in good agreement with [1963Wei], but the results of [1963Wei, 1970Fuk, 1970Kat] are in disagreement with data of the later work of [1971Blo, 1975Bes]. [1971Blo] and [1975Bes] are consistent with each other and show a more drastic increase in the hydrogen solubility with Mn addition (from 25.6 for pure Fe to 30.5 ppm (mass) for the Fe-10 mass% Mn alloy [1975Bes]). The interaction parameters for Mn with hydrogen in liquid Fe, εMn H , presented in the review [1977Bes] (i.e. the slope of hydrogen solubility vs the Mn content), are: –0.30 calculated from data of [1963Wei], –0.73 from [1970Fuk, 1970Kat], –1.9 from [1975Bes], and –2.0 from [1971Blo]. In order to make a choice between the data of [1963Wei, 1970Fuk, 1970Kat] and [1971Blo, 1975Bes], the linear relationship between the logarithm of hydrogen solubility and the ratio of the metals can be derived (log {Halloy} = xFe·log {HFe} + xMn·log {HMn}, where {Halloy}, {HFe} and {HMn} are solubility of hydrogen in a Fe-Mn alloy and pure Fe and Mn). However, there are no reliable data for H solubility in Mn at 1600°C and 101 kPa. [1967Bur] extrapolated the data of [1937Sie] for the H-Mn system at 1350–1450 to 1600°C in order to describe the H solubility in the Fe-Mn melt from 0 to 100% Fe. However, there may be some error in the extrapolation leading to incorrect results. Burylev obtained an expression that was inconsistent with the earlier work [1965Bur] (the latter described well the experimental data of [1963Wei]), which turned out to be in better agreement with the later experimental data of [1971Blo, 1975Bes] rather than those of [1963Wei]. Thus, there would seem to be no way to chose unambiguously between the 1600°C liquidus isotherm of [1963Wei] and [1970Fuk, 1970Kat] or that of [1975Bes], which is close to [1971Blo]. DOI: 10.1007/978-3-540-78644-3_5 # Springer 2008
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Fe–H–Mn
3
In order to calculate the solvus isotherm at 1000°C, [1966Bur] proposed the equation logVHFe-Mn {ccm/100 g} = 0.73 + 0.78·xMn, which was derived using the H solubility in Fe and Mn in accordance with the later work of [1984Fro] ({H}Fe = 0.027 at.%, VHFe = 5.348 ccm/100g; and {H}Mn = 0.156 at.%, VHMn = 32 ccm/100g). [1965Sch] showed that alloying with Mn increases the solubility of hydrogen at 900 and 1000°C. The dependence of hydrogen solubility on the Mn content is not monotonous between 400 and 800°C, and in the temperature interval from 400 to 600°C and from 5 to 9 mass% Mn, the solubility of hydrogen is independent of the Mn content. The hydrogen saturation raises the temperature of the ε→γ transformation at high pressures [1978Pon]. The solubility of hydrogen in the ε phase, formed in the Fe-22 mass% Mn alloy at 280°C and 2 GPa hydrogen, is close to 0.1 H/Me, which is more than 3 times higher than in the γ phase formed in this alloy [1978Pon]. The hydrogen solubility was shown to be higher in the ε phase than in the γ phase irrespective of the route of hydrogenation. [1993Aik1, 1993Aik2] studied the hcp ε phase in an Fe-21Mn (at.%) alloy in order to compare the stability of the γ and ε phases in the two-phase alloy subjected to cathodic hydrogen charging in an aqueous solution and showed that the hcp ε phase can always be stabilized more easily than the fcc γ phase, not only in a hydrogen-gas atmosphere at high pressure, as in the case studied by [1978Pon] and [1982Ant2], but also in an aqueous solution under cathodic polarization. Isothermal Sections The isothermal section of the Fe-H-Mn system at 3 GPa and 600°C is given in Fig. 1. It is constructed using the data related to the solubility of hydrogen in the γ phase of pure iron as a function of temperature and hydrogen pressure [2005Hir, 2005Fuk], the p-T phase diagrams of the Fe-H [2003Fuk] and Mn-H [2002Fuk] systems and data relating to the solubility of hydrogen in the γ and ε phases in the Fe-22 mass% Mn alloy at 2 GPa and 350°C [1978Pon]. Because the section was constructed for pressures of 3 GPa, the experimentally derived data for the hydrogen solubility in the γ and ε phases of [1978Pon] pffiffiffiffiffiffiffiffi were increased by 3=2. Thermodynamics There are few thermodynamic data on the Fe-H-Mn system. The specific activity coefficient of hydrogen in liquid Fe-H-Mn alloys was evaluated by [1981Sch] as: log γHMn = (– 489.75·xMn + 138.73·xMn2/T) + (0.0695·xMn – 0.0243·xMn2), where xMn is the mole fraction of Mn in the melt. Notes on Materials Properties and Applications [1967Kod, 1968Kod] studied the effect of plastic deformation and annealing on the hydrogen permeability of ferromagnetic austenite (Fe-36 mass% Mn) over the temperature interval of 300 to 600°C and a hydrogen pressure of 101 kPa. It was shown that hydrogen permeability in the deformed alloy is more sensitive to point defects and pile-ups, and also to free dislocations. They excluded the role of short-range order, which varies considerably with plastic deformation and may have considerable influence on hydrogen mobility. [1982Ant1] reported that at 250°C and pH2 3 GPa, an isomorphic phase transition ε1 →⊊ε2 took place leading to an uneven rise in the hydrogen solubility, up to 0.4 H/Me. At this hydrogen content, the Fe77.6Mn22.4 becomes ferromagnetic. The high-temperature corrosion of Fe-Mn alloys in hydrogensulphide/hydrogen atmospheres was studied by [1984Sme]. A mechanism of hardening or softening of Fe-Mn alloys with hydrogen was studied by [1988Kim]. Literature data concerning investigations of the material properties of Fe-H-Mn alloys are listed in Table 3. Miscellaneous Alloying with Mn lowers the mobility of hydrogen in solid Fe-Mn alloys. The rate is at a maximum when the Mn content is higher than 8 mass% [1965Sch]. According to calculations performed by [2006Ela], Mn can form hydrides, which contributes to the impurity distribution between the liquid and gas phases. Landolt-Börnstein New Series IV/11D4
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Fe–H–Mn
[2006Ela] studied the removal of impurities from Fe by hydrogen plasma arc melting. It was shown that by using a gas mixture of Ar + 30 vol% H2, the Mn content could be reduced from 11.1 mass ppm to 0.31 mass ppm.
Table 1.
Investigations of the Fe-H-Mn Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1961Mae]
Saturation of liquid alloys by hydrogen in Sieverts’ type apparatus
1620°C, < 8 mass% Mn, liquid
[1963Wei]
Saturation of liquid alloys by hydrogen in Sieverts’ type apparatus
1592°C, < 10 mass% Mn, liquid
[1965Sch]
Effusion method and hot extraction method
400–1000°C, up to 14 mass% Mn
[1970Fuk] [1970Kat]
Sampling method
1600°C, < 3 mass% Mn, liquid
[1978Pon]
X-ray diffraction, electroresistivity at high temperatures and high pressures
Up to 250°C,in inert medium (silicone) pressure up to 2 GPa, hydrogen pressure up to 2.33 GPa, γ→ε and ε→γ transformations in the Fe-22 mass% Mn alloy
[1982Ant1]
X-ray diffraction, electroresistivity and spontaneous magnetization at high pressures of hydrogen
250°C, up to 7 GPa of hydrogen pressure, antiferromagnetic-ferromagnetic transition in Fe-22.4 at.% Mn - H alloy
[1993Aik1] [1993Aik2]
X-ray diffraction, the specimens were cathodically charged with hydrogen at 17°C, using a bath of 0.1 N H2SO4, poisoned with a small amount of As2O3, the current density was 500 A·cm–2
21.3 mass% Mn alloy annealed at 1000°C after that quenched in ace water, ε and γ phases saturated with hydrogen
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(βH) < – 259.34
hP2 P63/mmc Mg
a = 377.6 c = 616.2
[Mas2] triple point
(αH) < – 271.9
cF4 Fm 3m Cu
a = 533.8
[Mas2]
(δFe)(h2) 1538–1394
cI2 Im 3m W
a = 293.15
at 1390°C [V-C2, Mas2]
(continued)
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Fe–H–Mn
5
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αFe)(r) < 912
cI2 Im 3m W
a = 286.65
Pure Fe at 25°C [Mas2]
a a a a
Metastable in Fe-Mn binary system, quenched from the γ phase region 10.9 at.% Mn, [2000Mar] 12.2 at.% Mn, [2000Mar] 14.5 at.% Mn, [2000Mar] 16.9 at.% Mn, [2000Mar]
α’(Mn,Fe) = = = =
287.4 287.55 287.6 287.8
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, 13 GPa [Mas2]
(δMn)(h3) 1246–1138
cI2 Im 3m W
a = 308.0
Pure Mn, at 1138°C [Mas2]
γ, (γMn,γFe)
cF4 Fm 3m Cu
a = 357.6+0.122 xMn
0 ≤xMn ≤ 50 [2000Mar]
a = 386.0
Pure Mn, at 1100°C [Mas2]
a = 364.67
Pure Fe, at 915°C [V-C2, Mas2]
(γMn)(h2) 1138–1100 (γFe)(h1) 1394–912 γ’MnHx
tI2 P14/mmm
a = 267.0 c = 377.6
x = 0.41, at 27°C [1998Fed]
β, MnHx
cP20 P4132 βMn
a = 634.2
x = 0.08, at room temperature [1996Ant]
(βMn)(h1) 1100–727 α, MnHx (αMn)(r) < 727 ε, (Mn,Fe)Hx ε, (Mn,Fe) ε, MnHx
a = 631.52 Pure Mn, at 727°C [Mas2]
cI58 I 43/mmc αMn
a = 894.03
x = 0.073, 27°C [2002Ant]
a = 891.26
Pure Mn, at 25°C [Mas2]
hP2 P63/mmc Mg
Metastable in Fe-Mn binary system 18.5 at.% Mn [1993Oka]
a = 253.3 c = 408.8 a = 266.8 c = 433.0
x = 0.65 [1978Bel] (continued)
Landolt-Börnstein New Series IV/11D4
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Fe–H–Mn
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
a = 269.2 c = 435.5 a = 269.4 c = 435.7 a = 269.7 c = 436.8
x = 0.83, at –153°C [2002Ant] x = 0.86, at room temperature [1996Ant] x = 0.94 [1978Bel]
ε’, FeHx(I)
hP4 P63/mmc Nd
a = 267.9 c = 877.0
x = 1, at –183°C [2002Ant]
ε’, MnHx(I)
hP4 P63/mmc Nd
a = 271.3 c = 886.2
5.2 GPa, at 760°C [2002Fuk]
Table 3.
Investigations of the Fe-H-Mn Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1967Kod]
Hydrogen permeability using membrane technique, microhardness, volume resistivity, X-ray line broadening
Influence of plastic deformation on the hydrogen permeability
[1968Kod]
Hydrogen permeability using membrane technique, microhardness, X-ray line broadening
Influence of plastic deformation on the hydrogen permeability
[2006Ela]
GDMS was used for precise analysis of impurity content in specimens before and after plasma arc melting
Refining of iron
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Fe–H–Mn
Fig. 1. Fe-H-Mn.
Landolt-Börnstein New Series IV/11D4
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Isothermal section at 3 GPa and 600°C
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8 References [1937Sie] [1961Mae]
[1963Wei] [1965Bur] [1965Sch] [1966Bur] [1967Bur] [1967Kod]
[1968Kod]
[1970Fuk]
[1970Kat] [1971Blo] [1975Bes]
[1975Pon]
[1977Bes]
[1977Kar]
[1978Bel]
[1978Pon]
[1981Sch]
Fe–H–Mn
Sieverts, A., Moritz, H., “Manganese and Hydrogen” (in German), Z. Phys. Chem., Abt. A, 180, 249–263 (1937) (Phase Diagram, Phase Relations, Experimental, 5) Maekawa, S., Nakagawa, Y., “The Effect of Some Alloying Elements on the Solubility of Hydrogen in Liquid Iron” (in Japanese), Nippon Kinzoku Gakkai-shi, 25(9), 577–580 (1961) (Phase Relations, Thermodyn., Experimental, 6) Weinstein, M., Elliott, J.F., “Solubility of Hydrogen in Liquid Iron Alloys”, Trans. Met. Soc. AIME, 227, 382–393 (1963) (Phase Relations, Thermodyn., Calculation, Experimental, 27) Burylev, B.P., “Solubility of Hydrogen in Liquid Iron Alloys” (in Russian), Izv. Vyss. Uchebn. Zaved., Chern. Metall., 8(2), 17–22 (1965) (Phase Relations, Thermodyn., Calculation, 13) Schwarz, W., Zittter H., “Solubility and Diffusion of Hydrogen in Iron Alloys” (in German), Arch. Eisenhuettenwes., 36(5), 343–349 (1965) (Phase Relations, Experimental, 16) Burylev, B.P., “Solubility of Hydrogen in Solid Iron Alloys”, Russ. J. Phys. Chem., 40(4), 442–445 (1966) (Phase Relations, Calculation, 12) Burylev, B.P., “Solubility of Gases in Molten Mn-base Alloys” (in Russian), Izvest. Vyssh. Ucheb. Zaved., Chern. Met., (4), 5–11 (1967) (Phase Relations, Calculation, 17) Kodes, Ye.S., Geld, P.V., Goltsov, V.A., “Effect of Plastic Deformation and Annealing on the Hydrogen Permeability of Ferromanganese Austenite”, Phys. Met. Metallogr., 24, 144–149 (1967), translated from Fiz. Met. Metalloved., 24(3), 528–534 (1967) (Experimental, Electr. Prop., Mechan. Prop., Transport Phenomena, 10) Kodes, E.S., Goltsov, V.A., Geld, P.V., “Solubility and Rate of Permeation of H in Fe-Ni and Fe-Mn Austenitic Alloys” (in Russian), Tr. Uralsk. Politekhn. Inst., 167, 20–25 (1968) (Experimental, Mechan. Prop., Transport Phenomena, 10) Fukuda, S., Sugiyama, T., Furukawa, T., Kato, E., “Solubility of Hydrogen in Liquid Iron Alloys”, Rep. Casting Research Lab./Waseda Univ., (21), 35–46 (1970) (Phase Relations, Thermodyn., Experimental, 22) Kato, E., Fukuda, S., Sugiyama, T., Furukawa, T., “Solubility of Hydrogen in Liquid Iron Alloys” (in Japanese), Tetsu to Hagane, 56, 521–535 (1970) (Experimental, 28) Blossey, R.G., Pehlke, R.D., “Molubility of Hydrogen in Liquid Fe-Co-Ni Alloys”, Metallurg. Trans., 2(11), 3157–3161 (1971) (Phase Relations, Thermodyn., 27) Bester, H, “Kinetic of the Hydrogen Exchange Between the Gas Phase and Liquid Iron as well as his Binary Alloys”, Dr.-Ing. Diss. Techn. Hochsch. Aachen, (1975), as quoted by [1977Bes] Ponyatovskii, E.G., Belash, I.T., “Production and Physical-Chemical Properties of Manganese Hydride” (in Russian), Dokl. Acad. Nauk SSSR, 224 (3), 607–608 (1975) (Crys. Structure, Phase Relations, Experimental, 1) Bester, H., Lange, K.W., “H Solubility in Fe and in Liquid Fe-Mn, Fe-Cr, and Fe-Si Alloys” (in German), Stahl u. Eisen, 97, 1037–1039 (1977) (Phase Relations, Thermodyn., Experimental, 53) Karamysheva, G.A., Men, A.N., “Calculation of the Concentration-Dependence of Hydrogen Solubility in Binary Alloys Using the Cluster Component Method”, Russ. Metall. (Engl. Transl.), (2), 82–83 (1977), translated from Izv. Akad. Nauk SSSR, Met., (2), 95–96 (1977) (Calculation, Interface Phenomena, 5) Belash, I.T., Ponomarev, V.K., Tissen, V.G., Afonnikova, N.S., Shekhman, V.Sh., Ponyatovskii, E.G., “Ferromagnetism in Mn-H and Mn-D System” (in Russian), Fiz. Tverd. Tela, 20(2), 422–427 (1978) (Crys. Structure, Phase Relations, Experimental, 24) Ponyatovskiy, Y.G., Antonov, V.Y., Belash, I.T., “Influence of Hydrogen Pressure up to 23 kBars on Temperature of ε-γ Transformation in Fe-22 % Mn Alloy”, (in Russian), Fiz. Met. Metalloved., 46 (5), 1090–1092 (1978) (Phase Relations, Experimental, Electr. Prop., 7) Schuermann, E., Kaettlitz, W., “Equivalent Effect of the Alloying Elements on the Concentration- and Temperature-Dependent Hydrogen Solubility in Iron-Rich Ternary and Multicomponent Melts” (in German), Arch. Eisenhuettenwes., 52(8), 295–301 (1981) (Phase Relations, Thermodyn., Calculation, 20)
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Fe–H–Mn [1982Ant1]
[1982Ant2]
[1984Fro] [1984Sme]
[1986Usu]
[1987Iro]
[1988Kim]
[1990San] [1991Bad]
[1991Bol]
[1993Aik1]
[1993Aik2]
[1993Oka]
[1996Ant]
[1998Ant]
[1998Fed]
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Antonov, V.E., Belash, I.T., Georgieva, I.Ya., Degtyareva, V.F., Thiessen, V.G., Shalimova, A.V., “Magnetization of Solid Solutions of Hydrogen in HCP Fe+22.4 at.% Mn Alloy”, Sov. Phys. - Solid State, 24(4) 554–558 (1982), translated from Fiz. Tverd. Tela, 24(4), 975–981 (1982) (Crys. Structure, Phase Relations, Experimental, Magn. Prop., Electr. Prop., 23) Antonov, V.E., Belash, I.T., Ponyatovski, E.G., “T-p Phase Diagram of the Fe-H System at Temperatures to 450°C and Pressures to 6.7 GPa”, Scr. Metall., 16(2), 203–208 (1982) (Phase Relations, Phase Diagram, Electr. Prop., Magn. Prop., Experimental, 15) Fromm, E., Jehn, H., “Solubility Hydrogen in the Elements”, Bull. Alloy Phase Diagrams, 5(3), 323–326 (1984) (Phase Relations, Review, 2) Smeltzer, W.W., Jackson, P.R.S., “The High-Temperature Corrosion of Fe-Mn and Fe-Mn-Al Alloys in Hydrogen-Sulphide Hydrogen Atmospheres”, CIM Bull., 77(866), 52 (1984) (Abstract) Usui, M., Asano, S., “Hydrogen Induced Transformation in FCC Fe-Mn Alloys” (in Japanese), J. Jpn. Inst. Met., 50(11), 943–949 (1986) (Crys. Structure, Phase Relations, Experimental, 10) Irodova, A.V., Glazkov, V.P., Somenkov, V.R., Shilstein, S.Sh., Antonov, V.E., Ponyatovskii, E.G., “Neutron Diffraction Study of Antiferromagnetic Ordering in Manganese Hydrides”, Sov. Phys. Solid State, 29(8) 1562–1568 (1987) (Crys. Structure, Phase Relations, Experimental, Magn. Prop., 6) Kimura, T., Matsui, H., Kimura, H., “Effect of Hydrogen on the Mechanical Properties of Iron-Based Alloys” in “Strength of Metals and Alloys, (ICSMA8)”, Proceedings of the 8th International Conference, Pergamon, Oxford, Vol. 2, 541–546 (1988) (Morphology, Mechan. Prop., Experimental, 5) San-Martin, A., Manchester, F.D., “Fe-H (Iron-Hydrogen)”, Bull. Alloy Phase Diagr., 11(2), (1990) (Phase Diagram, Crys. Structure, Phase Relations, Assessment, #, 91) Badding, J.V., Hemley, R.J., Mao, H.K., “High-Pressure Chemistry of Hydrogen in Metals: In Situ Study of Iron Hydride”, Science, 253(5018), 421–424 (1991) (Crys. Structure, Morphology, Phase Relations, Experimental, 33) Bol’shov, L.A., “Hydrogen Localization in FCC Lattices of Alloys of Iron with Chromium, Manganese, Cobalt, and Nickel”, Russ. Metall.,(2), 165–169 (1991), translated from Izv. Akad. Nauk SSSR. Met., (2), 165–169 (1991) (Phase Relations, Thermodyn., Calculation, 6) Aikawa, T., Nishino, Y., Asano, S., “Preferential Absorption of Hydrogen in the ε Phase in Fe-Mn Alloys” (in Japanese), J. Jpn. Inst. Met., 57(4), 484–388 (Crys. Structure, Phase Relations, Experimental, 25) Aikawa, T., Nishino, Y., Asano, S., “Stabilization of the HCP Epsilon phase in an Fe-21 % Mn Alloy Subjected to Cathodic Hydrogen Charging”, Scr. Metall. Mater., 29(1) 135–137 (1993) (Crys. Structure, Phase Relations, Experimental, 16) Okamoto, H., “Fe-Mn (Iron-Manganese)” in “Phase Diagram of Binary Iron Alloys”, Okamoto, H. (Ed.), ASM International, Materials Park, OH, 9, 203–213 (1993) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, #, 91) Antonov, V.E., Antonova, T.E., Chirin, N.A., Ponyatovskii, E.G., Baier, M., Wagner, F.E., “T-p Phase Diagram of the Mn-H System at Pressures to 4.4 GPa and Temperatures to 1000°C”, Scr. Mater., 34(8), 1331–1336 (1996) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, #, 15) Antonov, V.E., Kornell, K., Fedotov, V.K., Kolesnikov, A.I., Ponyatovsky, E.G., Shiryaev, V.I., Wipf, H., “Neutron Diffraction Investigation of the dhcp and hcp Iron Hydrides and Deuterides”, J. Alloy Compd., 264, 214–222 (1998) (Crys. Structure, Phase Relations, Experimental, #, 24) Fedotov, V.K., Antonov, V.E., Kolesnikov, A.I., Beskrovnyi, A.I., Grosse, G., Wagner, F.E., “Neutron Diffraction Investigation of γ Manganese Hydride”, Solid State Commun., 107(12), 787–790 (1998) (Crys. Structure, Phase Relations, Experimental, #, 21)
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10 [2000Mar]
[2002Ant]
[2002Fuk]
[2002Zin]
[2003Fuk]
[2004Wit]
[2005Fuk]
[2005Hir]
[2006Ela]
[Mas2] [V-C2]
Fe–H–Mn Marinelli, P., Baruj, A., Fernandez Guillermet, A., Sade, M., “Lattice Parameters of Metastable Structures in Quenched Fe-Mn Alloys. Part I: Experimental Techniques, bcc and fcc Phases”, Z. Metallkd., 91(11), 957–962 (2000) (Crys. Structure, Experimental, Phase Relations, 38) Antonov, V.E., “Phase Transformations, Crystal and Magnetic Structures of High-Pressure Hydrides of d-Metals”, J. Alloys Compd., 330–332, 110–116 (2002) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, 46) Fukai, Y., Haraguchi, T., Shinomiya, H., Mori, K., “Constitution of the Mn-H System at High Hydrogen Pressure”, Scr. Mater., 46(9), 679–684 (2002) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, #, 25) Zinkevich, M., Mattern, N., Handstein, A., Gutfleisch, O., “Thermodynamics of Fe-Sm, Fe-H, and H-Sm Systems and its Application to the Hydrogen-DisproportionationDesorption-Recombination (HDDR) Process for the System Fe17Sm2-H2”, J. Alloys Compd., 339, 118–139 (2002) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 101) Fukai, Y., Mori, K., Shinomiya, H., “The Phase Diagram and Superabundant Vacancy Formation in Fe-H Alloys Under High Hydrogen Pressures”, J. Alloys Compd., 348,105–109 (2003) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, #, 42) Witusiewicz, W.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diffus., 25(4), 346–354 (2004) (Phase Diagram, Thermodyn., Assessment, 34) Fukai, Y., “The Structure and Phase Diagram of M-H Systems at High Chemical PotentialsHigh Pressure and Electrochemical Synthesis”, J. Alloys Compd., 404–406, 7–15 (2005) (Electrochemistry, Interface Phenomena, Phase Diagram, Review, Thermodyn., 40) Hiroi, Y., Fukai, Y., Mori, K., “The Phase Diagram and Superabundant Vacancy Formation in Fe-H Alloys Revisited”, J. Alloys Compd., 404–406, 252–255 (2005) (Crys. Structure, Phase Relations, Phase Diagram, Experimental, #, 23) Elanski, D., Lim, J.-W., Mimura, K., Isshiki, M., “Impurity Removal from Fe, Cr, Ti, and V Metals by Hydrogen Plasma arc Melting and Thermodynamic Estimation of Hydride and Sulfide Formation”, J. Alloys Compd., 421(1–2), 209–216 (2006) (Calculation, Phase Diagram, Phase Relations, 8) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Fe–H–Mo
1
Iron – Hydrogen – Molybdenum Kostyantyn Korniyenko
Introduction The constitution of the ternary Fe-H-Mo system is of great interest, primarily in relation to the use of Fe-Mo alloys for hydrogen permeation applications. On the whole, hydrogen has the unique property of being able to diffuse through certain metals at a rapid rate, and attempts have been made to exploit this property in the thermally regenerative cell system. The use of a permeable metal diaphragm as an electrode in fused salt systems alleviates some of the problems associated with porous metal gas electrodes, namely, concentration polarization and “flooding” of the electrode with either electrolyte or gas. Although iron and some iron alloys have proved to be compatible with a lithium hydride cell environment, their use as a hydrogen diffusion diaphragm-electrode limits the cell current density and therefore restricts the practical application of such a cell [1965Hei]. The dissolution of hydrogen in iron-molybdenum alloys was studied experimentally by [1934Sie] and has been reviewed by [1938Kub]. Subsequent studies of phase equilibria in the ternary Fe-H-Mo system have been concentrated mainly on the influence of molybdenum additions on the hydrogen solubility of liquid iron [1961Mae, 1964Gun, 1967Ban, 1967Som, 1970Fuk, 1970Kat, 1972Ngi, 1976Zhi]. A mathematical method of calculating the solubility of hydrogen in Fe-Mo alloys has been presented by [1966Bur]. A list of the investigations of phase relations, crystal structures and thermodynamics as well as the techniques employed are presented in Table 1. Binary Systems The Fe-H system under a hydrogen pressure of 0.1 MPa is accepted from [Mas2], which is based on the results of [1990San]. Later, this system was investigated carefully by [2003Fuk] up to hydrogen pressures of 10 GPa, temperatures up to 1500°C (Fig. 1). A drastic reduction in the melting point to 800°C owing to the dissolution of hydrogen was observed at a hydrogen pressure of 3 GPa. In comparison to the p-T phase diagram of pure iron, the α-γ-ε triple point of Fe at 8.4 GPa and 430°C was shifted to 5 GPa and 260°C. The hydrogen solubility in (γFe) and (αFe) under high pressure and its correlation with superabundant vacancy formation were reported by [2005Hir]. The Fe-Mo and H-Mo systems are accepted from [Mas2]. Solid Phases No ternary intermetallic phases have been found in the Fe-H-Mo ternary system. Crystallographic data relating to the known unary and binary phases and their compositions and temperature ranges of stability are listed in Table 2. Investigations of hydrogen solubility in solid Fe-Mo binary alloys were carried out by [1934Sie]. The data obtained were reproduced in the review of [1938Kub] in relation to the phase equilibria in the Fe-Mo binary system. The temperature dependence of this solubility for the alloys with different compositions is shown in Fig. 2; the solubility increases with increasing temperature. The most drastic increase in hydrogen solubility is observed for iron rich alloys. It should be noted that as the constitution of the Fe-Mo system was later revised, information about hydrogenation of Fe-Mo alloys is in need of amendment. A formula for calculating the solubility of hydrogen in solid iron-molybdenum alloys has been proposed by [1966Bur]. The calculated results were compared with the experimental data given by [1934Sie] for 1100 and 1000°C, and satisfactory agreement was observed. Isothermal Sections Knowledge of the phase relationships in the Fe-H-Mo system is limited since all investigations have been carried out on samples having a high concentration of iron. Isothermal sections of the Fe-H-Mo system have Landolt-Börnstein New Series IV/11D4
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Fe–H–Mo
not been constructed. Information in the literature relates to the solubility of hydrogen in liquid iron-based Mo-containing alloys [1961Mae, 1964Gun, 1967Ban, 1967Som, 1970Fuk, 1970Kat, 1972Ngi]. The published data are quite contradictory. [1964Gun, 1967Ban, 1967Som, 1970Kat, 1972Ngi] have reported a decrease in the solubility of hydrogen with increasing molybdenum content. However, [1961Mae, 1976Zhi] reported a contrary tendency for hydrogen solubility, it increasing with increasing molybdenum contents up to about 6 mass% (or 3.58 at.%) at 1600°C. [1981Sch] formulated a mathematical model on the basis of published experimental data, and used this to construct isotherms of hydrogen solubility at temperatures of 1605 and 1540°C. However, a number of the experimental data used were in contradiction with their original articles. Thermodynamics The effect of molybdenum on the activity of hydrogen in liquid iron alloys has been widely studied experimentally [1961Mae, 1964Gun, 1965Bag, 1967Ban, 1967Som, 1970Fuk, 1970Kat, 1972Ngi]. On the whole, the results obtained by applying both the Sieverts’ method [1964Gun, 1965Bag, 1967Ban, 1967Som, 1970Fuk, 1970Kat, 1972Ngi] and the samples method [1961Mae, 1970Fuk, 1970Kat] show that Mo additions increase the values of log fHMo. The interaction parameters for hydrogen in liquid Fe-Mo alloys obtained by various investigators are compiled in Table 3. The thermodynamic aspects of hydrogen solubility in liquid Fe-Mo alloys are discussed in relation to electronic structure [1963Wei, 1972Ngi] or based on an H– coordinate structure [1965Bag]. A mathematical description of the concentration- and temperature-dependent specific activity coefficients γHX in the iron rich Fe-H-X melts has been presented by [1981Sch], in particular, for the case of X = Mo. It was found that log γHMo at 1600°C decreases to about –0.04 on increasing the Mo content up to 20 at.%. The calculated enthalpy of dissolution ΔHHMo and entropy of dissolution ΔSHMo also decreased on increasing the molybdenum content up to 20 at.% - to about –1.5 kJ·mol–1 and –0.02 kJ·mol–1·K–1, respectively. Notes on Materials Properties and Applications The possibilities of using Fe-Mo alloys hydrogen absorbents are of great interest, and the conditions required to realize these possibilities are need in experimental verification. The main goal of the physical property investigations of Fe-H-Mo alloys reported in the literature [1967Kod, 1991Rie] was to establish the effects of the hydrogenation of steels on properties. It was shown by [1967Kod] that plastic deformation (in compression, Vickers hardness tester using a load of 20 kg) of a Fe based alloy with a Mo content of 0.42 mass% and a total content of C, Si, Mn, P, S and Cr of 0.544 mass% leads to a substantial increase in the activation energy and pre-exponential factor in the expression for the temperature dependence of hydrogen permeability. Also, the effect of deformation and annealing on the electrical resistance of alloys gradually heated to and cooled from 800°C was studied. These measurements were carried out using a standard potentiometric method on 1.0–1.5 mm diameter wire specimens deformed by drawing to 25–30% reduction. Two stages of recovery of the hydrogen permeability parameters were identified. The first stage corresponds to quite an intense recovery of hydrogen permeability and electrical resistivity parameters which terminates at 300–320°C. It was postulated that the recovery of hydrogen permeability under these conditions is associated mainly with the variation in the vacancy structure of the alloy. In the second stage, above 600°C, a marked reduction in hardness was observed (from about 1.59 GPa at 600°C to about 1.18 GPa at 700°C). In the author’s opinion, this change is associated with the variation in the dislocation structure whose influence on the diffusion of hydrogen is weaker than that of the vacancy structure. Mechanical tests were also carried out by [1991Rie] during the study of the effect Mo on hydrogen permeation, diffusion, solubility and its distribution on different binding states in the Fe-based alloys. The fracture behavior of the alloys affected by hydrogen in different binding states were tested under constant elongation rate conditions, and no specific effect of molybdenum was observed. It was concluded that crack initiation and fracture progress depend first of all on the external hydrogen activity and the grain size of the material, not on the total hydrogen content. Fine crystalline iron alloys of higher strength proved to be less sensitive to hydrogen damage than coarse grain weak structures.
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Miscellaneous Diffusion and permeation characteristics of hydrogen in Fe-Mo alloys are presented in [1952Cha, 1965Hei, 1967Kod, 1976Zhi, 1991Hag, 1991Rie, 1992Hag]. [1952Cha] studied iron-based alloys containing 0.98, 2.36 or 10.36 mass% Mo as well as C, Si, Mn, P, S and Ni with a total content up to 0.77 mass%. The specimens were heated to 1000°C and degassed for 1 h, after which they were soaked in pure hydrogen for 2 h before the system was pumped out. Permeation rates at various temperatures, both on heating and cooling, were then measured. It was concluded that molybdenum has no appreciable effect on the permeation rate of hydrogen in iron. So, activation energies for permeation were quite similar: 69.14, 71.23 and 64.11 kJ·mol–1 for alloys with 0.98, 2.36 and 10.36 mass% Mo, respectively. The hydrogen permeation of Fe-Mo alloys with molybdenum contents of 5.04 or 10 mass% (3.22 and 6.07 at.% Mo, respectively) was studied by [1965Hei] over a temperature range of 350–800°C and a hydrogen pressure of 101300 Pa. Activation energies were calculated from the experimental data for the alloys with 2.97 and 6.07 at.% Mo as 39.69 ± 0.59 kJ·mol–1 and 40.19 ± 0.80 kJ·mol–1, respectively. The application of iron diaphragm electrodes in lithium hydride cells under heavy current loads has resulted in current densities which were 1.5 times greater than those obtainable from data on the permeability of hydrogen through an iron diaphragm with a hydrogen pressure of 101300 Pa across the diaphragm. [1976Zhi] reported that a molybdenum content of up to 3.11 mass% (or about 1.83 at.%) in steels has little effect on the permeation coefficient of hydrogen. At the same time, with increasing Mo content, the activation energy of diffusion essentially increases. The diffusion coefficients of hydrogen in iron-based alloys with Mo contents from 0.51 to about 3.35 at.% annealed at 900°C for 3 h in vacuum and then slowly furnace cooled have been measured by [1991Hag, 1992Hag] using an electrochemical permeation method between 10 and 40°C. The temperature and alloy composition dependence of the diffusion coefficients have been discussed in relation to simple trapping models. It was concluded that the diffusivity of hydrogen is lowered by alloying, and this is caused by the interstitial hydrogen sites around the substitutional atoms acting as trapping sites. The diffusion behavior of hydrogen in the alloys is strongly influenced by the elastic interaction between hydrogen and the alloying atoms. [1991Rie] investigated the effect of Mo on the binding states of hydrogen in iron and the fracture behavior of iron alloys reporting that no specific effects of molybdenum were observed.
Table 1.
Investigations of the Fe-H-Mo Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1934Sie]
Sieverts’ method (apparatus consists of the reaction chamber, the gas-measuring system and the gas-purification equipment)
300–1100°C; 1.2–83.7 at.% Mo
[1961Mae]
Sampling method (Tamman furnace as the melting apparatus; sampler consisting of a quartz tube attached with a rubber bulb; tin fusion technique; streaming by pure hydrogen)
1600°C; < 6 mass% ( 3.58 at.%) Mo
[1964Gun]
Sieverts’ method
1550, 1600, 1650°C; < 10 mass% ( 6.07 at.%) Mo
[1965Bag]
Sieverts’ method
1600°C; ≤ 30 mass% ( 19.97 at.%) Mo
[1967Ban]
Sieverts’ method
1550–1670°C; ≤ 30 mass% ( 19.97 at.%) Mo
[1967Kod]
Induction melting, hydrogen permeability measurements, X-ray powder diffraction
250–750°C; 0.42 mass% Mo (continued)
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Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1967Som]
Sieverts’ method
1540 - 1605°C; ≤ 30 mass% ( 19.97 at.%) Mo
[1970Fuk]
Sieverts’ method, sampling method
1550–1600°C; ≤ 12 mass% ( 7.35 at.%) Mo
[1970Kat]
Sieverts’ method, sampling method
1550–1600°C; ≤ 12 mass% ( 7.35 at.%) Mo
[1972Ngi]
Sieverts’ method
1550–1600°C, 6664–13329 Pa pressure; ≤ 16 mass% ( 9.98 at.%) Mo
[1976Zhi]
Hydrogen stream technique (step annealing and cooling of the specimens)
300–800°C; 0.54, 1.27 and 1.83 at.% Mo
[1991Hag]
Electrochemical permeation technique
10–40°C; Fe with addition of 0.51 to 3.35 at.% Mo
[1991Rie]
Electrochemical permeation technique
Fe rich corner
[1992Hag]
Electrochemical permeation technique, X-ray diffraction
10–40°C; Fe with addition of 0.51 to 3.35 at.% Mo
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(δFe) (h2) 1538–1394
cI2 Im 3m W
a = 293.15
pure Fe, T = 1390°C [V-C2, Mas2] dissolves 0.0696 at.% H at 1536°C, 0.1MPa [1990San, Mas2] dissolves 1.6 at.% H at 1460°C, 1.8 GPa [2003Fuk] in the Fe-Mo binary system joins with the (αFe) [Mas2]
(γFe) (h1) (austenite) 1394–912
cF4
a = 364.67
pure Fe, T = 912°C [V-C2, Mas2]
(αFe) (r) (ferrite) < 912
Fm 3m Cu
cI2
dissolves 0.0466 at.% H at 1394°C, 0.1MPa [1990San, Mas2] dissolves 10.7 at.% H at 680°C, 2.5 GPa; 41.1 at.% H at 680°C, 6 GPa [2005Hir] dissolves 1.7 at.% Mo at 1140°C [Mas2] a = 286.65
pure Fe, T = 25°C [V-C2]
Im 3m W
dissolves 24.4 at.% Mo at 1449°C [Mas2] (continued)
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Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(εFe) ≥ 1.3·105 bar
hP2 P63/mmc Mg
a = 246.8 c = 396.0
pure Fe, T = 25°C, 13 GPa [Mas2] high-pressure modification
(βH) < –259.34
hP2 P63/mmc Mg
a = 377.6 c = 616.2
[V-C2, Mas2]
(αH) < –271.9
cF4 Fm 3m Cu
a = 533.8
[V-C2, Mas2]
(Mo) < 2623
cI2 Im 3m W
a = 314.70
T = 25°C [Mas2] dissolves 31.3 at.% Fe at 1611°C [Mas2]
σ, MoFe 1611–1235
tP30 P42/mnm σCrFe
a = 921.8 c = 481.3
μ, Mo6Fe7 < 1370
hR39 R 3m W6Fe7
a = 475.46 c = 2571.6
R, Mo2Fe3 1488–1200
hR159 R 3 Mo3Cr2Co5
λ, MoFe2 < 927
Table 3.
hP12 P63/mmc MgZn2
43.3 to 57.1 at.% Fe [Mas2] [V-C2] 56 to 61 at.% Fe [Mas2] [V-C2] 61.5 to 66.1 at.% Fe [Mas2]
a = 1091.0 c = 1935.4 a = 1095.6 c = 1935.3
Mo1.9Fe3.1, at 1250 to 1490°C [V-C2] annealed at 1250°C [V-C2] 66.7 at.% Fe [Mas2] C14 structure [V-C2]
a = 475.5 c = 776.7
Interaction Parameters for Hydrogen in Liquid Fe-Mo Alloys
Reference
Method/ Experimental Technique
εHMo
Temperature [°C]
Composition/ Range Studied
[1961Mae] [1964Gun]
Sampling method Sieverts’ method
– 0.013 + 0.0030 + 0.0029 + 0.0028 + 0.0029 (the average)
1600 1550 1600 1650 1550–1650
< 6 mass% ( 3.58 at.%) Mo < 10 mass% ( 6.07 at.%) Mo
(continued)
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Reference
Method/ Experimental Technique
εHMo
Temperature [°C]
Composition/ Range Studied
[1965Bag]
Sieverts’ method
+ 0.0022
1600
< 30 mass% ( 19.97 at.%) Mo
[1967Ban]
Sieverts’ method
+ 0.0014
1550–1670
< 30 mass% ( 19.97 at.%) Mo
[1967Som]
Sieverts’ method
+ 0.0009
1540–1605
< 30 mass% ( 19.97 at.%) Mo
[1970Fuk, 1970Kat]
Sieverts’ method Sampling method
+ 0.0048 + 0.0107
1550–1600 1600
< 5 mass% ( 2.97 at.%) Mo < 12 mass% ( 7.35 at.%) Mo
[1972Ngi]
Sieverts’ method
+ 0.0029
1550–1600 (6664–13329 Pa pressure)
< 10 mass% ( 6.07 at.%) Mo
Fig. 1. Fe-H-Mo.
The Fe-H phase diagram
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Fig. 2. Fe-H-Mo.
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Hydrogen solubility in the Fe-Mo alloys at different temperatures and atmospheric pressure
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References [1934Sie] Sieverts, A., Bruenig, K., “Absorptive Ability of the Iron-Molybdenum Alloys for Hydrogen and Nitrogen” (in German), Arch. Eisenhuettenwes., 7(11), 641–645 (1934) (Phase Diagram, Phase Relations, Experimental, *, 15) [1938Kub] Kubaschewki, O., “Absorption of Gases in the Metals” (in German), Z. Elektrochem., 44(2), 152–167 (1938) (Phase Relations, Review, Interface Phenomena, 187) [1952Cha] Chang, P.L., Bennett, W.D.G., “Diffusion of Hydrogen in Iron and Iron Alloys at Elevated Temperatures”, J. Iron Steel Inst., 170, 205–213 (1952) (Morphology, Experimental, Interface Phenomena, Kinetics, 19) [1961Mae] Maekawa, S., Nakagawa, Y., “The Effect of Some Alloying Elements on the Solubility of Hydrogen in Liquid Iron” (in Japanese), Nippon Kinzoku Gakkai-shi, 25(9), 577–580 (1961) (Phase Relations, Thermodyn., Experimental, *, 6) [1963Wei] Weinstein, M., Elliott, J.F., “Solubility of Hydrogen in Liquid Iron Alloys”, Trans. AIME, 227, 382–393 (1963) (Phase Relations, Thermodyn., Calculation, Experimental, Review, Theory, Electronic Structure, 27) [1964Gun] Gunji, K., Ono, K., Aoki, Y., “The Effect of Various Elements on the Solubility of Hydrogen in Liquid Pure Iron”, Trans. Nat. Res. Inst. Met. (Jpn.), 6(5), 209–213 (1964), translated from J. Jpn. Inst. Met., 28, 64–68 (1964) (Phase Relations, Thermodyn., Experimental, Review, *, 7) [1965Bag] Bagshaw, T., Engledow, D., Mitchell, A., “Solubility of Hydrogen in Some Liquid Iron-Based Alloys”, J. Iron Steel Inst., 203, 160–165 (1965) (Crys. Structure, Phase Relations, Thermodyn., Calculation, Experimental, Theory, 28) [1965Hei] Heinrich, R.R., Johnson, C.E., Crouthamel, C.E., “Hydrogen Permeation Studies. I. Armco Iron and Iron-Molybdenum Alloys”, J. Electrochem. Soc., 112(11), 1067–1070 (1965) (Morphology, Experimental, Interface Phenomena, 12) [1966Bur] Burylev, B.P., “Solubility of Hydrogen in Solid Iron Alloys”, Russ. J. Phys. Chem. (Engl. Transl.), 40(4), 442–445 (1966) (Phase Relations, Thermodyn., Calculation, 12) [1967Ban] Ban-ya, S., Fuwa, T., Ono, K., “Solubility of Hydrogen in Liquid Iron Alloys” (in Japanese), Tetsu to Hagane, 53(2), 101–116 (1967) (Phase Relations, Thermodyn., Experimental, Review, *, 18) [1967Kod] Kodes, E.S., Gel’d, P.V., “Recovery of Hydrogen Permeability of Ferritic Alloys of Iron with Molybdenum and Nickel”, Sov. Mater. Sci., 3(6), 478–482 (1967), translated from Fiz.-Khim. Mekh. Mater., 3(6), 656–662 (1967) (Crys. Structure, Experimental, Electr. Prop., Interface Phenomena, Mechan. Prop., 11) [1967Som] Someno, M., Nagasaki, K., Kadoi, K., “Solubility of Hydrogen in Liquid Iron and in Some Liquid Binary Iron Alloys” (in Japanese), J. Jpn. Inst. Met., 31, 729–734 (1967) (Phase Relations, Thermodyn., Experimental, Review, *, 27) [1970Fuk] Fukuda, S., Sugiyama, T., Furukawa, T., Kato, E., “Solubility of Hydrogen in Liquid Iron Alloys”, Rep. Casting Research Lab./Waseda Univ., (21), 35–46 (1970) (Phase Relations, Thermodyn., Calculation, Experimental, Review, *, 22) [1970Kat] Kato, E., Fukuda, S., Sugiyama, T., Furukawa, T., “Solubility of Hydrogen in Liquid Iron Alloys” (in Japanese), Tetsu to Hagane, 56, 521–535 (1970) (Phase Relations, Thermodyn., Experimental, *, 28) [1972Ngi] Ngia, N., Yavoyskiy, V.I., Kosterev, L.B., Afanas’yev, M.I., “Hydrogen Solubility in Binary Iron-Base Alloys”, Russ. Metall. (Engl. Transl.), (4), 11–15 (1972), translated from Izv. Akad. Nauk SSSR, Met., (4), 18–22 (1972) (Phase Relations, Thermodyn., Experimental, *, 19) [1976Zhi] Zhitenev, V.I., Ryabov, R.A., Geld, P.V., “Diffusion Coefficient and Solubility of Hydrogen in Solid Solutions of Molybdenum in Iron” (in Russian), Fiz. Met. Metalloved., 41, 650–652 (1976) (Phase Relations, Experimental, Interface Phenomena, 8) [1981Sch] Schuermann, E., Kaettlitz, W., “Equivalent Effect of the Alloying Elements on the Concentration- and Temperature-Dependent Hydrogen Solubility in Iron-Rich Ternary and Multicomponent Melts” (in German), Arch. Eisenhuettenwes., 52(8), 295–301 (1981) (Phase Relations, Thermodyn., Calculation, Review, *, 20)
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[1991Rie]
[1992Hag]
[2003Fuk]
[2005Hir]
[Mas2] [V-C2]
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San Martin, A., Manchester, F.D., “The Fe-H (Iron-Hydrogen) System”, Bull. Alloy Phase Diagrams, 11(2), 173–184 (1990) (Phase Diagram, Thermodyn., Review, #, 91) Hagi, H., “Effect of Substitutional Alloying Elements (Al, Si, V, Cr, Mn, Co, Ni, Mo) on Diffusion Coefficient of Hydrogen in α-Iron”, J. Jpn. Inst. Met., 55(12), 1283–1290 (1991) (Crys. Structure, Morphology, Experimental, Interface Phenomena, 23) cited from abstract Riecke, E., Johnen, B., “Effect of Mo, V, Nb, Ti, Zr and their Carbides on the Binding States of Hydrogen in Iron and the Fracture Behaviour of the Iron Alloys”, Werkstoffe und Korrosion, 42(12), 626–636 (1991) (Morphology, Phase Relations, Experimental, Review, Theory, Interface Phenomena, Mechan. Prop., 18) cited from abstract Hagi, H., “Effect of Substitutional Alloying Elements (Al, Si, V, Cr, Mn, Co, Ni, Mo) on Diffusion Coefficient of Hydrogen in α-Iron”, Mater. Trans. JIM, 33 (5), 472–479 (1992) (Crys. Structure, Morphology, Experimental, Interface Phenomena, 33) Fukai, Y., Mori, K., Shinomiya, H., “The Phase Diagram and Superabundant Vacancy Formation in Fe-H Alloys under High Hydrogen Pressures”, J. Alloys Compd., 348, 105–109 (2003) (Crys. Structure, Phase Diagram, Phase Relations, Experimental, #, 42) Hiroi, T., Fukai, Y., Mori, K., “The Phase Diagram and Superabundant Vacancy Formation in Fe-H Alloys Revisited”, J. Alloys Compd., 404–406, 252–255 (2005) (Crys. Structure, Phase Relations, Experimental, *, 23) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Fe–H–Ni
1
Iron – Hydrogen – Nickel Vasyl Tomashik, Larysa Shcherbak, Olga Fabrichnaya
Introduction The diffusion of hydrogen in Fe-Ni alloys is of great technological importance, especially for the design of controlled thermonuclear reactors [1973Dre]. Therefore, the main trends in the experimental and theoretical investigations of this system were the study of hydrogen dissolution and its diffusion and permeability in the Fe-Ni alloys. Hysteresis loops involving H absorption/desorption are notable in this ternary system [2002Ant, 2004Zag]. The first experimental results concerning the Fe-H-Ni ternary system were obtained more than 50 years ago [1952Kur], where the dissolution of hydrogen in Fe-Ni alloys, containing up to 20 mass% Ni, was investigated. The dissolution of hydrogen in such alloys at different temperatures and alloy compositions was investigated mainly using Sieverts’ method by [1957Izm, 1960Bus, 1961Mae, 1963Wei, 1964Gun, 1965Bag, 1965Bur, 1965Sch, 1966Bag, 1966Noz, 1966Sch1, 1966Sch2, 1968Dus, 1968Gol, 1969Lip, 1971Bec, 1971Blo, 1971Way, 1972Ngi, 1974Boo, 1974Pet, 1974Shv, 1975Dag, 1976Pet, 1976Sta, 1977Pet, 1977Pon, 1990Mat]. The experimental results of [1952Kur] indicated maximum hydrogen solubility at 4 mass% Ni: increasing the Ni content up to 20 mass% results in only a negligible decrease in the solubility. Later investigations showed that the hydrogen solubility in Fe-Ni alloys increases with increasing Ni content. An empirical formula for the level of hydrogen solubility in solid Fe-Ni alloys at 1200 and 1400°C was proposed by [1966Bur]. The temperature dependence of hydrogen solubility in Fe-Ni alloys within the temperature range from 1400 to 1800°C can be expressed by log10{H} = A/T + B [1976Pet, 1977Pet] over the whole composition range. It was shown than the hydrogen solubility depends essentially on the alloy composition and the level of disorder [1969Sim]. The equations for the hydrogen solubility and its diffusion coefficient in an FeNi3 alloy were obtained using the energies of the Fe(Ni)-H bond [1970Vyk]. The hydrogen solubility in the Fe-Ni alloys at 1150°C was calculated using the method of cluster components without correlation in the solution by [1977Kar]. The data obtained coincide with the experimental results of [1965Sch]. A statistical model of hydrogen solubility in Fe melts was proposed and the hydrogen solubility in Fe-Ni alloys at 1600°C and 0.1 MPa was calculated by [1984Fro1] which practically coincides with the data of [1966Sch1]. Diffusion regularities of hydrogen in the Fe-Ni alloys at different temperatures and pressures were studied by [1956Shc, 1973Dre, 1965Sch, 1967Dus, 1969Gol, 1969Sim, 1971Bec, 1972Dre, 1974Shv, 1979Sig]. The permeability of hydrogen in Fe-Ni alloys was investigated by [1967Fra, 1969Gol, 1969Sim, 1971Bec, 1974Shv, 1976Gol]. Between 0 and 40 mass% Ni the dominant factor in controlling the value of the permeability is the reduction of the mole fraction of the α phase in the alloy [1971Bec]. Some thermodynamic properties of the Fe-H-Ni alloys have been investigated both experimentally and theoretically by [1960Bus, 1961Mae, 1963Wei, 1964Gun, 1966Bag, 1966Noz, 1966Sch1, 1966Sch2, 1971Bec, 1971Blo, 1972Ngi, 1973Bar, 1974Boo, 1974Pet, 1974Shv, 1976Sta]. Details of the experimental studies are reported in Table 1. Binary Systems The Fe-Ni and Fe-H systems were taken from the MSIT Binary Evaluation Program [2007Kuz] and [Mas2], respectively. The solubility of H in solid and liquid Fe and Ni are accepted from [1984Fro2], corrected by [2006Tom]. These data may be used up to a H2 pressure of 50 MPa. [2002Ant] reported the stability of the FeH hydride phase ε’-FeH at pressures of 3.5–7 GPa. It must be pointed out that the H-Ni diagram given by [Mas2] for a hydrogen pressure of 50 MPa cannot be accepted, because it gives a eutectic composition of 0.036 at.% H at 1406°C whereas the actual composition would be 2.14 at.% H at the same temperature. The solubility of hydrogen in liquid Ni at ambient pressure has Landolt-Börnstein New Series IV/11D4
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also been reviewed by [1987Sch]. The best H-Ni phase diagram available to date is that given in [2002Shi] determined experimentally up to 800°C and a 5 GPa pressure of hydrogen. The H-Ni solid solution exhibits a miscibility gap, with a critical point, evaluated by [2004Fuk], of around 360°C, 1.4 GPa and 50 at.% H. Solid Phases Ternary compounds were not found in the Fe-H-Ni system at the standard temperature and pressure [2005Ide]. According to the data of [1972Bar, 1973Bar] the ternary solid solutions similar to the nickel hydrides are formed in this system at up to about 25 at.% Fe under high pressures (up to 3 GPa). [1976Miz] indicated that hydrides of Ni rich Fe-Ni alloys could be prepared by cathodic saturation. The decomposition of these hydrides at room temperature became faster with the increase of Fe content and sample preparation for an alloy above 10 at.% Fe was difficult. The increase of the lattice parameter due to the hydride formation became smaller and the hydrogen content decreases with the increase of the Fe concentration. The obtained hydrides were not microscopically homogeneous and coexist with non-hydrogenated phase. Crystallographic data of the unary and binary compounds are shown in the Table 2. The observed decrease in the lattice parameters of the Fe0.2Ni0.8Hx solid solutions suggests that the samples with x ≈ 0.8 and 0.63 were inhomogeneous [2002Ant]. The Fe0.2Ni0.8D0.63 sample appeared inhomogeneous throughout. The solubility of hydrogen in the Fe-Ni alloys at medium (400–1000°C) and low (70°C) temperatures according to the data of [1965Sch] and [1971Bec] is shown in Figs. 1 and 2, respectively. Solubility of hydrogen in the Fe-Ni alloys has been investigated both in solid states at 400–1000°C with the step of 100°C up to 12 mass% Ni by [1965Sch], at temperatures of 650, 600, 500, 400 and 300°C and compositions up to 11 at.% Ni by [1974Shv], at 70°C by [1971Bec]. The results of [1965Bag, 1966Bag] indicate an inflexion on the solubility curve at about 69 at.% Ni, coinciding with a minimum in the liquidus of the Ni-Fe binary system and a peak in the Curie temperature-composition relationship. Anomalous changes of the hydrogen solubility in the Fe-Ni alloys, containing 18.9 and 24.85 mass% Fe, were determined within the temperature range from 450 to 500°C [1968Gol]: the solubility increases for 20 and 65 % for the first and the second alloy respectively. The hydrogen solubility in the Fe-Ni alloys at 20°C is stopped at the adding more than 30 mass% Fe [1969Lip]. According to the data of [1990Zag] there is an order-disorder phase transition in the case of the alloy (24.85 mass% Fe + 75.15 mass% Ni) + H at 500°C (Kurnakov temperature). The appearance of long-range atomic order in this alloy should enhance the hydrogen solubility. [1990Mat] also noted that the ordering of this alloy leads to the increasing of hydrogen solubility. Sieverts’ law is obeyed at the hydrogen solubility in liquid Fe-Ni alloys up to 0.101 MPa pressures [1963Wei]. Isothermal Sections Solubility of hydrogen in the Fe-Ni alloys has been investigated in liquid alloys at 1400–1800°C by [1976Pet, 1977Pet], 1600°C by [1965Bag], [1971Blo] and [1974Pet], 1592°C up to 23 mass% Ni by [1963Wei]. The solubility of hydrogen in the Fe-Ni alloys at 1400–1700°C according to [1966Sch1] is presented in Fig. 3. Two fcc phases γ and β (hydride phase) with different hydrogen contents can be in equilibrium at room temperature in the H-Ni system under suitable conditions of electrolytic charging or under high hydrogen pressure [1982Sak]. In the Fe-H-Ni system the boundary between γ and β phases disappears for Fe contents above about 26 at.%. Phase diagram at 25°C from [1982Sak] is presented in Fig. 4. The γ ↾ β first order transition for the alloy containing 10 at.% Fe ends at the critical point with coordinates 250 < Tcr < 300°C and 1.6 MPa < pcr < 1.9 MPa [1976Pon1]. Critical temperature decreases linearly with Fe content increasing within the interval from 5 to 15 at.% Fe. Thermodynamics The Gibbs energies of formation of the Ni-Fe hydrides at room temperature forming at the pressure up to 3 GPa were calculated by [1973Bar]. The calculations were based on hydrogen pressure data of decomposition reaction. For the alloys containing from 1.20 to 36.5 at.% Fe the Gibbs energy of formation was calculated to be within the interval from 23.9 ± 0.2 to 50.2 ± 0.5 kJ·mol–1 H2 according to [1973Bar]. DOI: 10.1007/978-3-540-78644-3_7 # Springer 2008
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Fe–H–Ni
3
The heat of the hydrogen dissolution in the Fe-Ni alloys increases from 24.3 to 26.0 kJ·g-at–1 H in the concentration range 1.5–9 at.% Ni and then decreases to 23.4 kJ·g-at–1 H at 12 at.% Ni according to [1974Shv]. According to [1976Sta] the thermodynamic properties of dissolved hydrogen are not dependent upon the state of order of the Fe-Ni alloys and also invariant to the magnetic transformation. From the experimental data on solubility of hydrogen in Fe-Ni alloys at 1400, 1500–1550, 1600 and 1700°C Schenk [1966Sch1] determined first order interaction parameters εNiH in Fe solution and εFeH in Ni solution using Wagner’s formalism for concentrations of Fe up to 10 at.%, Ni up to 5 at.%. The parameters εNiH were calculated to be equal to –0.28, –0.23, –0.21 and –0.16 at the above mentioned temperatures, respectively [1966Sch1]. These data are in reasonable agreement with other determinations at 1600°C (εNiH = 0.000 [1963Wei, 1974Boo]; εNiH = –0.2 [1965Bag]; εNiH = –0.385 [1966Noz]; εNiH = –0.5 ± 0.3 [1971Blo]; εNiH = –0.534 [1964Gun]; εNiH = –0.12 [1960Bus]; εNiH = –0.606 up to 15 mass% Ni [1972Ngi]; εNiH = –0.16 [1974Pet]). The value εNiH = –4.6 obtained for concentrations up to 6 mass% Ni [1961Mae] seems to be erroneous. The first order interaction parameters εFeH in Ni solutions are equal 0.78 and 1.2 at 1400 and 1500–1700°C respectively [1966Sch1]. These data are in reasonable agreement with other measurements (εFeH = +1.17 up to 20 mass% Ni [1966Bag]; εFeH = +0.5 ± 0.3 [1971Blo]; εFeH = +0.91 at 1600°C [1974Pet] for dilute solutions). The value of εFeH = +0.009 [1960Bus] seems to be too low. The heat of hydrogen solution in Fe-Ni alloys is given in Fig. 5 [1971Bec]. The concentration dependence of the activity coefficients of hydrogen with respect to the third component in Fe-Ni alloys at 1400, 1550 and 1700°C are shown in Figs. 6 and 7 [1966Sch1]. The thermodynamic calculations of the hydrogen solubility in the Fe-Ni alloys at 1600°C [1965Bur] practically coincide with the experimental results of [1960Bus, 1963Wei, 1965Bag, 1966Bag]. Notes on Materials Properties and Applications [2006Ara] indicated that the Fe0.7Ni0.3 alloy is one of the best media to store hydrogen. The hydrogen is detrimental to the Fe-Ni alloys quality [1973Dre]. An influence of a plastic deformation with the next annealing on the hydrogen solubility in the Fe-Ni alloy containing 40 mass% Ni was not found [1968Kod]. The corrosion rate and hydrogen permeability, hence, susceptibility to hydrogen embrittlement, pass through a minimum at about 50 mass% Ni [1971Bec]. [1976Gol] noted that the hydrogen permeability and the diffusion rate decrease within the temperature interval from 250 to 850°C at the Fe addition to Ni. Anomalous behavior of thermal expansion and Young’s modulus of the Fe-Ni alloys containing 32.9 and 75.0 at.% Ni (Invar alloys) was faded with hydrogenation [1983Har]. This change corresponds to the disappearance of the characteristic weak ferromagnetic states of such alloys. According to [1976Miz], [1976Pon2] and [2002Ant] the hydrides of Ni rich Fe-Ni alloys were ferromagnetic at low temperatures and TC rose with the increase of Fe concentration. The value of dTC/dp increases at the hydrogen pressure increasing and become positive at p ≈ 1.5 MPa [1976Pon2]. The spontaneous magnetization of the Fe0.2Ni0.8Hx solid solutions at 124 K decreases approximately linearly from 1.1 μB per metal atom at x = 0 to 0.6 μB at x = 1 [1978Ant, 2002Ant]. The increasing of hydrogen concentration in the Fe-H-Ni alloys as its pressure increases leads to the TC increasing [1978Ant]. At the hydrogen pressure above 1.5 MPa TC of the Fe0.675Ni0.325-H begins to grow and reaches about 700 K at 5.0 MPa. There is a sharp analogy in the influence of the plastic deformation and cathodic hydrogenation on the internal friction parameters [1980Bar]. Electric resistance of Ni-Fe alloys at 25°C and pressures up to 3 GPa hydrogen pressure was measured “in situ” by [1973Bar]. The maximum of electric resistance was observed in alloys with concentrations of 16 at.% Fe at hydrogen pressure above 2 GPa. Further increase of hydrogen pressure had small influence on resistance of alloys. Miscellaneous Isothermal absorption relationships between H2 and D2 gas pressure and hydrogen or deuterium content of a Fe0.2Ni0.8 alloy is given in Fig. 8 [2002Ant]. The bcc Fe-Ni alloys containing 70–80 mass% Fe showed extreme instability when electrolytic charging was stopped [1971Way].
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Fe–H–Ni
The hydrogen solubility in Ni rich Fe-Ni alloys exhibits the same kind of solubility anomaly associated with the presence of grain boundaries which has been observed in pure Ni [1976Sta]. The hydrogen impurity presumably promotes ordering of the alloy by introducing stresses into the lattice [1968Dus]. The diffusion coefficient of hydrogen in the Fe-Ni alloys was found to remain constant up to 0.04 at.% Ni [1973Dre]. A drop by a factor 3 occurs between 0.04 and 0.15 at.% Ni and from 0.15 to 5 at.% Ni the diffusion coefficient decreases slightly within the bcc structure. In the two-phase bcc-fcc region a decrease of about 3 orders of magnitude was observed, while in the disordered fcc structure the diffusion coefficient was found to be nearly independent of the composition. The effect of long-range order was determined and it was shown that the diffusion coefficient exhibits an increase of a factor of about 4 at the composition Ni3Fe [1973Dre]. The concentration dependence of the diffusion coefficient of hydrogen in the Fe-Ni alloys at room temperature and 350, 650 and 1050°C was determined by [1965Sch]. [1980Bar] indicated that the diffusion coefficients of hydrogen at the degassing temperature for the Fe-Ni alloy, containing 3 mass% Ni, are equal 1.93·10–9 cm2·s–1 at 250 K and for the alloy, containing 1.5 mass% Ni, 2.98·10–9 cm2·s–1 at 288 K. According to the data of [1971Bec], the diffusion coefficient falls from ≈10–4 cm2·s–1 for pure Fe to ≈10–10·cm2·s–1for 40 mass% Ni in the alloy. Thereafter it rises slightly to that of pure Ni. The hysteresis which exists in the case of the nickel hydride disappeared in the ternary high-pressure Fe-H-Ni phase, containing 17 at.% Fe [1973Bar]. The hydrogen diffusion rate increases with increasing the Ni content and reaches a maximum at 6 mass% Ni [1956Shc]. The next increasing of Ni content leads to its decreasing with a minimum at 74 mass% Ni and the increasing of the diffusion rate at higher Ni content. It has been found that in the case of Fe-Ni specimens having the ability to form ordered structures there is no linear relationship between the logarithm of diffusion coefficient and the reciprocal of the absolute temperature [1967Dus]. Ordered solid solutions, and particularly those with a long-range order, manifest much higher activation energy values for the hydrogen diffusion than unordered ones. [1972Dre] indicated that the diffusion coefficient of hydrogen varies with the alloy composition even within a homogeneous structure. Generally, the diffusion of hydrogen occurs faster within bcc structures than fcc ones. A continuous decrease of the diffusion coefficient within the mixed bcc + fcc phase toward the fcc structure results from a change in crystal structure. The diffusivity increases by a factor of four from a completely disordered to ordered structure close to 75 at.% Ni [1972Dre]. The diffusion rate of the hydrogen in the alloys, inclined to an ordering, changes with the composition and temperature nonlinearly [1969Sim]. The activation energies of the hydrogen diffusion and permeability in FeNi3 alloy as well as the dissolution enthalpy are the same for the ordered and disordered samples [1969Gol]. Introduction of Fe into Ni reduces the hydrogen absorption [1971Bar]. The height of an internal friction peak due to hydrogen in the Fe1–xNix (x = 0.35–1.0) alloys did not show a monotonic change versus the increasing Ni content and depends strongly on the alloy composition [2005Ide]. There is a general decrease in the rate constant of hydrogen absorption as the amount of Fe in the alloy decreases except in the region between 65 and 75 at.% Fe, where an increase in the rate penetration is observed [1967Fra]. The peculiarities of hydrogen absorption by the alloy C-Fe-Ni, containing 0.38 mass% C and 23.29 mass% Ni, was investigated by [1975Dag]. It was shown that the absorbed hydrogen quantity for austenite is within the interval from 2 to 4 cm3/100 g as for the martensite its value could reach 9 cm3/100 g. A large hysteretic peak was observed in the internal friction spectrum of the single-phase Fe50Ni50 alloy corresponds to hydrogen degassing from the (γFe,Ni) solid solution [2004Zag]. A new low temperature relaxation peak with the activation enthalpy of 0.41 ± 0.04 eV is also observed in such spectrum of the hydrogen-charged Fe50Ni50 alloy. It is presumably related to the relaxation of H-H pairs or H-Fe complexes in the Ni rich regions of the ordered Fe-Ni solid solution. Degassing of hydrogen from these regions is probably responsible for a minor small additional hysteretic peak. The temperature of (αFe)/(γFe) transition for the Fe-Ni alloys containing 10, 20 and 30 at.% Ni is lower in the hydrogen atmosphere in comparison with inert medium [1977Pon]. This difference reaches about 45°C at 2.2 MPa and 30 at.% Ni. The hydrogen dissolution in the (γFe) phase leads to its stabilization. The hydrogen dissolution reduces the vacancy formation in the Fe-Ni alloys [1972Bru]. A significant feature of the alloy containing 74.94 at.% Ni is a hysteresis loop within the temperature range 510–570°C, i.e. near the order-disorder transition temperature, on the temperature dependence hydrogen solubility [1968Dus]. Lattice dynamics of the Fe0.7Fe0.3-H system was discussed by [2006Ara]. DOI: 10.1007/978-3-540-78644-3_7 # Springer 2008
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Fe–H–Ni Table 1.
5
Investigations of the Fe-H-Ni Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1952Kur]
Measuring of hydrogen absorbed
1590–1685°C/(Fe + up to 20 mass% Ni) + H2
[1956Shc]
Gas volumetric technique
200–600°C/(Fe-Ni) + H2 up to 10.1 MPa
[1957Izm]
Vacuum melting method
1650°C/(Fe + up to 11 at.% Ni) + H2
[1960Bus]
Sieverts’ method
1600°C/(Fe-Ni) + H2
[1961Mae]
Sieverts’ method
1600 and 1620°C/(Fe-Ni) + H2
[1963Wei]
Sieverts’ method
1592°C/(Fe + up to 23 mass% Ni) + H2
[1965Bag]
Sieverts’ method
1600°C/(Fe-Ni) + H2
[1965Sch]
Effusion method
350–1000°C/up to 12 mass% Ni
[1966Bag]
Sieverts’ method
1600°C/(Fe-Ni) + H2
[1966Noz]
Sieverts’ method
1488–1672°C/(Fe-Ni) + H2
[1966Sch1]
Sieverts’ method
900–1680°C/(Fe-Ni) + H2
[1967Dus]
Pressure measurements
400–800°C/(Fe-Ni) + H2
[1967Fra]
Chemical analysis
25°C/(Fe-Ni) + H2
[1968Dus]
Thermovolumetric method
300–800°C/(Fe-Ni) + H2
[1968Gol]
Pressure measurements
300–800°C/(Ni + 9.35, 18.9 and 24.85 mass% Fe) + H2
[1969Gol]
Pressure measurements
300–850°C/Ni3Fe
[1969Lip]
Cathodic saturation method
20°C/(Fe-Ni) + H2
[1969Sim]
Hot extraction method
300–800°C/Ni + 9.35, 18.9 and 24.85 mass% Fe + H2
[1971Bar]
XRD, mass-spectrometry, electric resistance measurement
25°C/(Fe-Ni) + H2 up to 2.5 GPa
[1971Bec]
Electrochemical method
27–90°C/(Fe-Ni) + H2
[1971Blo]
Sieverts’ method
1600°C/(Fe-Ni) + H2 at 0.101 MPa
[1971Way]
XRD
20°C/Fe-H-Ni
[1972Bar]
Pressure and electrical resistance measurements
25°C/(Fe-Ni alloys with 0.23, 0.50, 1.2, 5.5, 9.8 and 15.6 at.% Fe) + H2 up to 2.5 MPa
[1972Ngi]
Sieverts’ method
1550–1600°C/(Fe-Ni) + H2 at 6.67–13.33 kPa
[1973Bar]
Electric resistance measurements, XRD, mass-spectrometry
25°C/(Fe-Ni) + H2 up to 3 GPa
[1973Dre]
Electrochemical and gas volumetric technique, metallography
25 and 58°C/(Fe-Ni) + H2
[1974Boo]
Pressure measurements
1600°C/(Fe-Ni) + H2 at 0.101 MPa (continued)
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Fe–H–Ni
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1974Shv]
Diffusion and penetration methods
300–650°C/Fe + 1.5–12 at.% Ni + H2
[1976Gol]
Berrer’ method
250–850°C/(Fe-Ni at 3.27, 4.75 and 8.98 at.% Fe) + H2
[1976Miz]
XRD, Mössbauer spectroscopy
Room temperature/(Fe-Ni alloys up to 10 at.% Fe) + H2
[1976Pet, 1977Pet]
“Hot volume” method
1400–1800°C/(Fe-Ni) + H2
[1976Pon1, 1976Pon2]
Electroconductivity and magnetic permeability measurements
Up to 520°C/(Fe-Ni at 5, 10, 15, 32.5 and 36 at.% Fe) + H2 up to 2.3 MPa
[1976Sta]
Pressure measurements
300–1050°C/Fe-H-Ni
[1977Pon]
XRD, magnetic permeability measurements
Up to 650°C/(Fe-Ni at 10, 20, 25, 28 and 30 at.% Ni) + H2 up to 2.3 MPa
[1978Ant]
XRD, magnetization measurements
80–220 K/(Fe-Ni at 10, 20, 40, 60, 66.1 and 67.5 at.% Fe) + H2 up to 7.0 MPa
[1980Bar]
Internal friction measurements
80–673 K/(Fe + 1.5 or 3 mass% Ni) + H2
[1983Har]
Gas volumetric method
173–623 K/(Fe + 32.9 or 75.0 at.% Ni) + H2
[1982Sak]
XRD, cathodic saturation
25°C/Fe-H-Ni
[2002Ant]
XRD, neutron diffraction, Mössbauer spectroscopy
250 and 350°C/Fe0.2Ni0.8 + H2(D2) up to 7 GPa
[2004Zag]
Neutron scattering and internal friction measurements
50–500 K/Fe50Ni50 + H2
[2005Ide]
Gas-equilibrating method, XRD
1000°C/Fe1–xNix (x = 0.35–1.0) + H2
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(δFe) 1538–1394
cI2 Im 3m W
a = 293.15
pure Fe at 1390°C [V-C2, Mas2] dissolves 3.8 at.% Ni at 1517°C
γ, (γFe,Ni) < 1517 (γFe) 1394–912
cF4 Fm 3m Cu
a = 364.67
pure Fe at 915°C [V-C2, Mas2] dissolves 10 at.% H at 700°C and 6 GPa H2 pure Ni at 25°C [Mas2] below critical temperature: γ1 - paramagnetic, Fe enriched γ2 - ferromagnetic, Ni enriched
(Ni) < 1455
a = 352.40
(continued) DOI: 10.1007/978-3-540-78644-3_7 # Springer 2008
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7
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αFe) < 912
cI2 Im 3m W
a = 286.65 a = 291.6
at 25°C [Mas2] dissolves 4.6 at.% Ni at 495°C FeHx at 900–1200°C [2002Ant]
(βH) < –259.34
hP2 P63/mmc Mg
a = 377.6 c = 616.2
[Mas2] triple point
(αH) < –271.9
cF4 Fm 3m Cu
a = 533.8
[Mas2]
γ’, FeNi3 < 517
cP4 Pm 3m AuCu3
a = 355.50
[V-C2, Mas2]
γ’’, FeNi metastable
tP4 P4/mmm AuCu
a = 357.9
metastable [V-C2, Mas2] [V-C2] metastable ordering temperature 320°C at 51.2 at.% Ni [1984Ros]
β, (Ni1–xFex)H1–y NiH
cF8 Fm 3m NaCl
Ni2H
hP3 P3m1 Ni2H
a = 374.0 a = 266 c = 433
x < 0.25, y < 0.333 [V-C2] [V-C2]
metastable
implicit from [2004Fuk]
ε’-FeH(I)
hP4 P63/mmc
a = 269.6 c = 877.0
[2002Ant] 3.5–7 GPa, 25–400°C
ε-FeH0.42
hP2 P63/mmc Mg
a = 258.3 c = 417.6
[2002Ant] metastable
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Fig. 1. Fe-H-Ni.
Fe–H–Ni
Solubility of hydrogen in Fe-Ni alloys containing up to 11 mass% Ni at 400–1000°C
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Fe–H–Ni
Fig. 2. Fe-H-Ni.
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Solubility of hydrogen in Fe-Ni alloys at 70°C
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Fig. 3. Fe-H-Ni.
Fe–H–Ni
Solubility of hydrogen in Fe-Ni alloys at 1400–1700°C
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Fe–H–Ni
Fig. 4. Fe-H-Ni.
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Isothermal section at room temperature
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Fig. 5. Fe-H-Ni.
Fe–H–Ni
Heat of hydrogen dissolution in Fe-Ni alloys as a function of composition
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Fe–H–Ni
Fig. 6. Fe-H-Ni.
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Concentration dependence of the hydrogen coefficient of activity in Fe-Ni alloys at 1400°C
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Fig. 7. Fe-H-Ni. 1700°C
Fe–H–Ni
Concentration dependence of the hydrogen coefficient of activity in the Fe-Ni melts at 1550 and
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Fig. 8. Fe-H-Ni. Isothermal absorption relationships between H2 and D2 gas pressure and hydrogen or deuterium content of a Fe0.2Ni0.8 alloy
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DOI: 10.1007/978-3-540-78644-3_7 # Springer 2008
16 References [1952Kur]
[1956Shc]
[1957Izm] [1960Bus]
[1961Mae]
[1963Wei] [1964Gun]
[1965Bag] [1965Bur] [1965Sch] [1966Bag] [1966Bur] [1966Noz]
[1966Sch1]
[1966Sch2]
[1967Dus] [1967Fra]
[1968Dus]
[1968Gol]
[1968Kod]
Fe–H–Ni
Kurochkin, K.T., Yavoyskiy, V.I., Gel’d, P.V., “Influence of the Alloying Elements on the Hydrogen Solubility in Liquid Steel” (in Russian) in “Proizvodstvo Stali”, SverdlovskMoscow, 73–90 (1952) (Experimental, Phase Relations, 16) Scherbakova, A.A., “Hydrogen Diffusion through Iron and Binary Iron-Chromium and Iron-Nickel Alloys at High Pressures and Temperatures” (in Russian), Zh. Prikl. Khim., 29(6), 879–884 (1956) (Experimental, Interface Phenomena, 6) Izmanova, T.A., Klyachko, Yu.A., “Hydrogen in some Iron Alloys” (in Russian), Khim. Nauka i Promyshl., 2(4), 528–529 (1957) (Experimental, Phase Relations, 3) Busch, T., Dodd, R.A., “The Solubility of Hydrogen and Nitrogen in Liquid Alloys of Iron, Nickel, and Cobalt”, Trans. Met. Soc. AIME, 218, 488–490 (1960) (Experimental, Phase Relations, 16) Maekawa, S., Nakagawa, Y., “The Effect of some Alloying Elements on the Solubility of H in Liquid Fe” (in Japanese), Nippon Kinzoku Gakkai Shi, 25, 577–580 (1961) (Experimental, Phase Relations, Thermodyn., 6) Weinstein, M., Elliott, J.F., “Solubility of Hydrogen in Liquid Iron Alloys”, Trans. Met. Soc. AIME, 227, 382–393 (1963) (Experimental, Phase Relations, 27) Gunji, K., Ono, K., Aoki, Y., “The Effect of Various Elements on the Solubility of Hydrogen in Liquid Pure Iron”, Trans. Nat. Res. Inst. Met. (Jpn.), 6(5), 209–213 (1964) (Calculation, Phase Relations, 7) Bagshaw, T., Engledow, D., Mitchell, A., “Solubility of Hydrogen in Some Liquid IronBased Alloys”, J. Iron Steel Inst., 203, 160–165 (1965) (Experimental, Phase Relations, 28) Burylev, B.P., “Solubility of Hydrogen in Liquid Iron Alloys” (in Russian), Izv. Vyssh. Uchebn. Zaved., Chern. Met., (2), 17–22 (1965) (Calculation, Phase Relations, 13) Schwarz, W., Zittter, H., “Solubility and Diffusion of Hydrogen in Iron Alloys” (in German), Arch. Eisenhuettenwes., 36(5), 343–349 (1965) (Experimental, Phase Relations, #, *, 16) Bagshaw, T., Mitchell, A., “Solubility of Hydrogen in Some Liquid Alloys of Nickel”, J. Iron Steel Inst., 204, 87–90 (1966) (Experimental, Phase Relations, 18) Burylev, B.P., “Hydrogen Solubility in Solid Iron Alloys” (in Russian), Zh. Fiz. Khim., 40(4), 822–825 (1966) (Calculation, Phase Relations, 12) Nozaki, H., Ban-ya, S., Fuwa, T., Matoba, S., Ono, K., “Effect of Carbon, Silicon, Phosphorus and Nickel on the Solubility of Hydrogen in Liquid Iron” (in Japanese), Tetsu to Hagane, 52(13), 1823–1833 (1966) (Experimental, Phase Relations, Thermodyn., 17) Schenck, H., Lange, K.W., “Research on the Solubility of Hydrogen in Iron, Nickel, Cobalt, Copper and in their Binary Nickel Alloys” (in German), Arch. Eisenhuettenw., 37(9), 739–748 (1966) (Experimental, Phase Relations, Thermodyn., #, *, 54) Schenck, H., Lange, K.W., “An Application of the Hydrogen Solubility to the Physical and Thermodynamic Description of the Solvents” (in German), Z. Metallkd., 57(5), 378–384 (1966) (Experimental, Phase Relations, Thermodyn., 34) Dus, R., Smialowski, M., “Diffusion of Hydrogen in f.c.c. Alloys of Nickel with Iron” Acta Metall., 15, 1611 (1967) (Experimental, Phase Relations, 17) Franklin, T.C., Hudson, P.E., “Absorption of Electrolytic Hydrogen by Nickel and IronNickel Alloys”, J. Electrochem. Soc., 114(6), 568–572 (1967) (Experimental, Interface Phenomena, 21) Dus, R., “The Einstein Temperature of Hydrogen Atoms Dissolved in Ordered and Disordered Fe-Ni Alloys”, Bull. Acad. Polon. Sci., 16(7), 371–376 (1968) (Experimental, Phase Relations, 9) Gol’tsov, V.A., Gel’d, P.V., Simakov, Yu.P., Shteinberg, M.M., Vykhodets, V.B., “Effect of Ordering on the Solubility of Hydrogen in Ni-Fe Alloys” (in Russian) in “Diffuzionniye Procesy v Metalakh”, Naukova Dumka, Kiev, 92–94 (1968) (Experimental, Phase Relations, 3) Kodes, E.S., Gol’tsov, V.A., Gel’d, P.V., “Solubility and Rate of Hydrogen Permeation in Iron-Nickel and Iron-Manganese Austenitic Alloys”, Tr. Uralsk. Politekhn. Inst., (167), 20–25 (1968) (Experimental, Mechan. Prop., 10)
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Fe–H–Ni [1969Gol]
[1969Lip]
[1969Sim]
[1970Vyk]
[1971Bar]
[1971Bec]
[1971Blo] [1971Way] [1972Bar] [1972Bru]
[1972Dre]
[1972Ngi]
[1973Bar]
[1973Dre]
[1974Boo]
[1974Pet]
[1974Shv]
[1975Dag]
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17
Gol’tsov, V.A., Vykhodets, V.B., Gel’d, P.V., Simakov, Yu.P., “Effect of Ordering on the Diffusion Velocity and Penetration of Hydrogen into Ni3Fe Alloy” (in Russian), Fiz.-Khim. Mekh. Mater., 5(5), 597–601 (1969) (Experimental, Phase Relations, Mechan. Prop., 18) Lipets, T.V., Vert, Zh.L., Tverdovsky, I.P., “Solubility of Hydrogen in Alloys of Nickel with Iron and Cobalt” (in Russian), Zh. Fiz. Khim., 43(5), 1331–1333 (1969) (Experimental, Phase Relations, 14) Simakov, Yu.P., “The Diffusion and Solubility of Hydrogen in Nickel-Iron Alloys” (in Russian), Sbornik Nauchn. Tr. Permsk. Politekhn. Inst., (51), 87–95 (1969) (Experimental, Phase Relations, Interface Phenomena, 6) Vykhodets, V.P., Gol’tsov, V.A., Gel’d, P.V., “Diffusion and Solubility of Hydrogen in Ordering Alloys of the Cu3Au Type” (in Russian), Ukr. Fiz. Zhur., 15(1), 107–110 (1970) (Calculation, Phase Relations, 9) Baranowski, B., Filipek, S., “Electrical Resistance of Ni/Fe Alloys at High Pressures of Gaseous Hydrogen”, Roczn. Chem., 45(7–8), 1353–1354 (1971) (Experimental, Phase Relations, Electr. Prop., 5) Beck, W., Bockris, J.O’M., Genshaw, M.A., Subramanyan, P.K., “Diffusivity and Solubility of Hydrogen as a Function of Composition in Fe-Ni Alloys”, Metall. Trans., 2(3), 883–888 (1971) (Experimental, Phase Relations, #, *, 20) Blossey, R.G., Pehlke, R.D., “Solubility of Hydrogen in Liquid Fe-Co-Ni Alloys”, Metall. Trans., 2, 3157–3161 (1971) (1972) (Experimental, Phase Relations, Thermodyn., 26) Wayman, M.L., Smith, G.C., “Hydride Formation in Nickel-Iron Alloys”, J. Phys. Chem. Solids, 32(1), 103–108 (1971) (Experimental, Phase Relations, 21) Baranowski, B., “Thermodynamics of Metal/Hydrogen Systems at High Pressures”, Ber. Bunsen-Gesellschaft Phys. Chem., 76(8), 714–724 (1972) (Experimental, Phase Relations, 51) Brusetti, R., Chambron, W., “Vacancy Defects from Tempering in Iron-Nickel Alloy with 70 % Nickel. Influence of Hydrogen” (in French), J. Phys. Chem. Solids, 33(5), 993–1003 (1972) (Experimental, Interface Phenomena, 18) Dresler, W., Frohberg, M., “The Influence of the Structure of Binary Systems - Particularly Iron-Nickel - on the Diffusion Coefficient of Hydrogen”, Ber. Bunsen-Gesellschaft Phys. Chem., 76(8), 826 (1972) (Experimental, Interface Phenomena, 0) Ngia, N., Yavoiskiy, V.I., Kosterev, L.B., Afanas’yev, M.I., “Hydrogen Solubility in Binary Iron-Base Alloys”, Russ. Metall. (Engl. Transl.), (4), 11–15 (1977), translated from Izv. Akad. Nauk SSSR, Met., (4), 18–24 (1972) (Experimental, Phase Relations, Thermodyn., 19) Baranowski, B., Filipek, S., “Investigation on the Nickel-Iron-Hydrogen System within a Wide Range of Hydrogen Pressures. Part I. Electric Resistance at 25°C”, Roczn. Chem., 47(11), 2165–2177 (1973) (Experimental, Phase Relations, Electr. Prop., Thermodyn., 22) Dresler, W., Frohberg, M.G., “Diffusion Coefficient of Hydrogen in the System Iron-Nickel at 25 and 58°C”, J. Iron Steel Inst., 211, 298–302 (1973) (Experimental, Interface Phenomena, 13) Boorstein, W.M., Pehlke, R.D., “Measurement of Hydrogen Solubility in Liquid Iron Alloys Employing a Constant Volume Technique”, Metall. Trans., 5(2), 399–405 (1974) (Experimental, Phase Relations, 29) Petrushevsky, M.S., Gel’d, P.V., Abramycheva, I.E., “Short-Range Order and Thermodynamics of the Interaction of Hydrogen with Liquid Fe-Ni-Co Alloys” (in Russian), Tr. Uralsk. Politekhn. Inst., 231, 80–85 (1974) (Calculation, Phase Relations, Thermodyn., 8) Shvetsov, N.I., Levchenko, V.L., Ryabov, R.A., “Hydrogen Solubility in Iron-Nickel Alloys” (in Russian), Tr. Uralsk. Politekhn. Inst., 231, 140–141 (1974) (Experimental, Phase Relations, Thermodyn., 3) Dagbert, C., Galland, J., Azou, P., Bastien, P., “Temperature Influence on Hydrogen Behavior in Iron-Nickel Alloy” (in French), Compt Rend. Acad. Paris, B, 281(18), 437–441 (1975) (Experimental, Interface Phenomena, 6)
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DOI: 10.1007/978-3-540-78644-3_7 # Springer 2008
18 [1976Gol]
[1976Miz] [1976Pet]
[1976Pon1]
[1976Pon2]
[1976Sta] [1977Kar]
[1977Pet]
[1977Pon]
[1978Ant]
[1979Sig]
[1980Bar]
[1982Sak]
[1983Har]
[1984Fro1] [1984Fro2] [1984Ros] [1987Sch]
Fe–H–Ni Gol’tsov, V.A., Latyshev, V.V., “Effect of Magnetic Ordering on the Permeability, Diffusion and Solubility of Hydrogen in Nickel-Iron Alloys” (in Russian), Fiz.-Khim. Mekh. Mater., 12(5), 28–32 (1976) (Experimental, Phase Relations, Magn. Prop., 15) Mizutani, T., Shinjo, T., Takada, T., “Mössbauer Study of Ni-Fe Hydrides”, J. Phys. Soc. Jpn., 41(3), 794–797 (1976) (Experimental, Phase Relations, Magn. Prop., 13) Petrushevsky, M.S., Gel’d, P.V., Abramycheva, I.E., Kostina, T.K., “Investigation of the Hydrogen Solubility in the Liquid Fe-Co-Ni Alloys” (in Russian), Dokl. Akad. Nauk SSSR, 227(2), 337–340 (1976) (Experimental, Phase Relations, 7) Ponyatovskii, E.G., Antonov, V.E., Belash, I.T., “On the Critical Phenomena in the Ni-Fe-H System at High Hydrogen Pressures” (in Russian), Dokl. Akad. Nauk SSSR, 230, 649–651 (1976) (Experimental, Phase Relations, Electr. Prop., Magn. Prop., 8) Ponyatovskii, E.G., Antonov, V.E., Belash, I.T., “Effect of Hydrogen Pressure up to 20 kbar on the Curie Point of Fe-Ni Alloys in the Invar Region” (in Russian), Fiz. Tverd. Tela, 18(12), 3661–3665 (1976) (Experimental, Electr. Prop., Magn. Prop., 19) Stafford, S.W., McLellan, R.B., “The Thermodynamic Properties of the Fe-Ni-H Ternary System”, Acta Metall., 24, 553–558 (1976) (Experimental, Thermodyn., Phase Relations, 53) Karamysheva, G.A., Men’, A.N., “Calculation of the Concentration-Dependence of Hydrogen Solubility in Binary Alloys Using the Cluster Component Method”, Russ. Metall. (Engl. Transl.), (2), 82–83 (1977), translated from Izv. Akad. Nauk SSSR, Met., (2), 95–96 (1977) (Calculation, Phase Relations, 5) Petrushevsky, M.S., Gel’d, P.V., Abramycheva, I.E., Kostina, T.K., “Investigation and Calculation of the Hydrogen Solubility in the Liquid Fe-Co-Ni Alloys” (in Russian), Izv. Vyssh. Uchebn. Zaved., Chern. Met., (10), 5–9 (1977) (Experimental, Phase Relations, 7) Ponyatovskii, E.G., Antonov, V.E., Belash, I.T., “The Effect of Hydrogen Pressure up to 23 kbar on the Temperature of the αγ Transformation in the Iron-Nickel Alloys” (in Russian), Fiz. Met. Metalloved., 44(5), 1038–1043 (1977) (Experimental, Phase Relations, 13) Antonov, V.E., Belash, I.T., Degtyareva, V.F., Ponomarev, B.K., Ponyatovskii, E.G., Tissen, V.G., “Magnetic Properties of the Hydrogen Solutions in the Nickel-Iron Alloys” (in Russian), Fiz. Tverd. Tela, 20(9), 2680–2686 (1978) (Experimental, Magn. Prop., 13) Sigrist, P., Feichtinger, H.K., Marincek, B., “A New Method for the Determination of Diffusion Coefficients and Solubilities of Gases in Liquid Metals and Alloys” (in German) in “Gases in Metals (Proc. Conf.)”, Darmstadt, 1979, Deutsche Gesellschaft Metallkde, Oberursel, Germany, 15–22 (1979) (Calculation, Experimental, Interface Phenomena, 11) Barmin, N.I., Kodes, E.S., Nefedov, V.M., Ryabov, R.A., Geld, P.V., “Investigation of Hydrogen Behavior in Alloys of Iron with Nickel” (in Russian), Zh. Fiz. Khim., 54(11), 2895–2898, (1980) (Experimental, Mechan. Prop., 7) Sakamoto, Y., Yuwasa, K., Hirayama, K., “X-Ray Investigation of the Absorption of Hydrogen by Several Palladium and Nickel Solid Solution Alloys”, J. Less-Common Met., 88, 115–124 (1982) (Experimental, Crys. Structure, Phase Diagram, Phase Relations, #, 22) Harada, S., “Thermal-Expansion Coefficient and Youngs Modulus of Hydrogenated FCC Fe-Ni Invar-Alloys”, J. Phys. Soc. Jpn., 52(4), 1306–1310 (1983) (Experimental, Mechan. Prop., 14) Frohberg, M.G., Anik, S., “Prediction of Hydrogen Solubility in Binary Iron Melts” (in German), Arch. Eisenhuttenwes., 55(2), 45–48 (1984) (Calculation, Phase Relations, 20) Fromm, E., Jehn, H., “Hydrogen in Elements”, Bull. Alloys Phase Diagrams, 5(3), 323–326 (1984) (Review, Phase Relations, 3) Rossiter, P.L., Jago, R.A., “Towards a True Fe-Ni Phase Diagram”, Mater. Res. Soc. Symp. Proc., 21, 407–411 (1984) (Review, Phase Diagram, *, 26) Schuermann, E., Sittard, M., Voelker, R., “Equivalent Influence of Alloying Elements on the Nitrogen and Hydrogen Solubility in Liquid Ternary and Multicomponent Nickel-Base Alloys” (in German), Z. Metallkd., 78(6), 457–466 (1987) (Review, Phase Diagram, Phase Relations, Thermodyn., 53)
DOI: 10.1007/978-3-540-78644-3_7 # Springer 2008
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Fe–H–Ni [1990Mat]
[1990Zag]
[2002Ant]
[2002Shi] [2004Fuk]
[2004Zag]
[2005Ide]
[2006Ara] [2006Tom]
[2007Kuz]
[Mas2] [V-C2]
Landolt-Börnstein New Series IV/11D4
19
Matysina, Z.A, Milyan, M.I., “Connection of Some Phase Transformations with Alloy Phase Diagrams”, Russ. Metall. (Engl. Transl.), (2), 171–183 (1990), translated from Izv. Akad. Nauk SSSR, Met., (2), 175–188 (1990) (Review, Phase Relations, 48) Zaginaychenko, S.Yu., Matysina, Z.A., Milyan, M.I., “Solubility of Interstitial Impurities in Alloys”, Phys. Met. Metallogr., 70, 60–64 (1990), translated from Fiz. Met. Metalloved., (9), 63–67, (1990) (Calculation, Phase Relations, 20) Antonov, V.E., Baier, M., Dorner, B., Fedotov, V.K., Grosse, G., Kolesnikov, A.I., Ponyatovsky, E.G., Schneider, G., Wagner, F.E., “High-Pressure Hydrides of Iron and its Alloys”, J. Phys., Condens. Matter, 14, 6427–6445 (2002) (Experimental, Review, Phase Relations, #, 46) Shizuku, Y., Yamamoto, S., Fukai, Y., “Phase Diagram of the Ni-H System at High Hydrogen Pressures”, J. Alloys Compd., 336(1–2), 159–162 (2002) (Phase Diagram, Experimental, *, #, 19) Fukai, Y., Yamamoto, S., Harada, S., Kanazawa, M., “The Phase Diagram of the Ni-H System Revisited”, J. Alloys Compd., 372(1–2), L4–L5 (2004) (Phase Diagram, Experimental, #, 14) Zagorodniy, K., Shivanyuk, V., Toebbens, D., Danilkin, S., Gavriljuk, V., “Hydrogen-Caused Phase Transformations, Relaxation and Hysteretic Phenomena in FCC Alloys Fe55Cr25Ni20 and Fe50Ni50”, Scr. Mater., 50(12), 1467–1470 (2004) (Experimental, Phase Relations, 13) Ide, N., Naito, T., Asano, S., “Internal Friction Peak in FCC Fe-Cr-Ni Alloys HydrogenCharged by Gas-Equilibration Method”, Jpn. J. Appl. Phys., 44(11), 8088–8090 (2005) (Experimental, Morphology, Phase Relations, 11) Arasi, C.K., Balaguru, R.J.B., Raj, S.A.C., Lawrence, N., “Lattice Dynamics of the Ni0.3Fe0.7-H System”, Phys. Scri., 73(1), 11–16 (2006) (Experimental, Mechan. Prop., 14) Tomashik, V., Perrot, P., “Copper-Hydrogen-Nickel”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.16213.1.20, (2006) (Crys. Structure, Phase Diagram, Assessment, 45) Kuznetsov, V., “Fe-Ni (Iron-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published (2007) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 41) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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DOI: 10.1007/978-3-540-78644-3_7 # Springer 2008
Fe–H–O
1
Iron – Hydrogen – Oxygen Pierre Perrot
Introduction H2-H2O, as well as CO-CO2 atmospheres have been widely used to investigate the equilibria between iron and its oxides [1940Wib] because these atmospheres allow the determination of oxygen potentials with an accuracy which cannot be obtained with any other method such as electrochemistry or calorimetry. This happy result is easily explained by the affinity of oxygen towards iron, which is about the same as its affinity towards H2 or CO. The Fe-O phase diagram, widely accepted since [1945Dar, 1946Dar] has been confirmed by numerous and very precise experimental determinations, so, the works dealing with equilibria of iron oxides under H2-H2O atmospheres will not be taken into account in this report. For the same reason, the equilibria between iron, its oxides and hydroxides in aqueous solution will be presented from the only Pourbaix’ work [1963Pou]. On the other hand, the reaction of water vapor with liquid iron plays an important role in several steelmaking processes since the water content in the highly postcombusted gas may be as high as 30%. The main experimental results related to Fe-H2O reactions are gathered in Table 1. Binary Systems The Fe-H system under 0.1 MPa of hydrogen pressure is accepted from [1990San]. The Fe-H system has been carefully investigated by [2003Fuk] up to 10 GPa of hydrogen pressure and 1500°C. The hydrogen solubility in iron under low hydrogen pressures (up to 1 MPa) is given in [1980Fro]. The Fe-O system accepted by [Mas2], mainly from the fundamental work of [1945Dar, 1946Dar] has been carefully assessed by [1991Sun, 1995Kow]. The main compound of interest in the H-O system is H2O characterized by its oxidizing properties towards Fe. Mixtures H2-H2O are often used as a source of imposed oxygen potential. Solid Phases Numerous solid phases have been reported in the Fe-H-O system. They are often metastable and ill defined. The main solid phases, whose existence is widely recognized are presented in Table 2. Iron oxides Fe3O4 (magnetite) and Fe2O3 (hematite) are stable at room temperature whereas iron hydroxides Fe(OH)2 and Fe(OH)3 are metastable, lose water with formation of FeOOH (goethite) as an intermediate compound, Fe3O4 and Fe2O3 as final products. Gaseous Fe(OH)2 has been proposed as an intermediate species which forms during the formation of Fe3O4 by Fe oxidation under water vapor at 500°C [1973Sur]. Temperature – Composition Sections The phases in equilibrium under H2-H2O atmospheres are namely (αFe), (γFe), Fe1–xO, Fe3O4, L1 (liquid metal) and L2 (liquid oxide). Fe2O3 cannot be stable under such atmospheres, because it is reduced into Fe3O4 under pure water vapor. The stability domains of these phases are shown in Fig. 1, mainly from [1958Eds], and slightly modified to take into account the well accepted observation that the wuestite composition in equilibrium with iron between 1000°C and its melting point of 1380°C is Fe0.947O [1945Dar, 1970Gra]. Thermodynamics Information on the reactions between iron and water vapor may be found in [1980Fro, 1989Oet, 2003Per]. The solubility of water vapor in liquid iron may be described by the equilibrium: H2O (gas) ⇌ 2 {H}Fe + {O}Fe ΔrG° = 203130 + 1.72 T. In this expression, the reference states are the gaseous water under the standard pressure p° = 0.1 MPa and the elements (H or O) dissolved in liquid iron under the concentration c° = 1 mass%. It is in agreement with the expression proposed by [1989Rag] for the following Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
2
Fe–H–O
equilibrium: H2O (gas) ⇌ H2 (gas) + {O}Fe ΔrG° = – 134600 + 61.65 T. The reference states are the same (p° = 0.1 MPa for the gases and c° = 1 mass% for the solute). The oxygen solubility in solid iron in equilibrium with H2O-H2 atmospheres may be expressed by the following expressions [1976Kru]: In (γFe): log10 (cO / at.%) = log10(pH2O/pH2) + 0.26 – (3720/T ) (912–1394°C) In (δFe): log10 (cO / at.%) = log10(pH2O/pH2) + 1.17 – (4740/T ) (1394–1538°C) In liquid Fe: log10 (cO / at.%) = log10(pH2O/pH2) + 3.55 – (6750/T ) (1538–1700°C) In the three preceding expressions, cO represents the oxygen content of the iron, in at.%; it must be pointed out that these expressions may be used only for H2-H2O atmospheres corresponding to the iron domain shown in Fig. 1. The Fe-Fe3O4 equilibrium temperature depends on the total pressure [1984Fuk]. It increases by 400°C when the pressure goes from 0.1 MPa to 20 GPa. Notes on Materials Properties and Applications Water vapor acts as a decarburizing agent in steels [1994Nag] by reacting with iron droplets containing carbon according to the reaction H2O ⇌ 2 {H} + {O}. Dissolved oxygen reacts with C to form CO. This reaction is thermodynamically favorable and can cause a substantial hydrogen pick up in the steel. H2 has been proposed for the direct reduction of iron ores with the aim to reduce CO2 emissions due to the blast furnace-basic oxygen furnace route. A promising way is the hydrogen plasma smelting route [2005Pla]. αFeOOH (goethite) is used as a pigment [2005Kre] with color ranging from lemon yellow to dark brown. Miscellaneous The kinetics of oxidation of αFe with water vapor was investigated at 500°C by [1973Sur] and at 800–1000°C by [1970Gra]. The mechanism of Fe3O4 oxide layer growth is better described by a chemical transport model via a gaseous species such as Fe(OH)2 than by a diffusion model. The oxide layer is non protective and the kinetics of the layer growth tends to be linear with time. The potential-pH diagrams (Pourbaix’ diagram) of iron in aqueous solution [1963Pou] are shown in Figs. 2 and 3. The diagram shown in Fig. 2 is the stable diagram in which the solid phases involved are Fe, Fe3O4 and Fe2O3. Figure 3 gives the metastable diagram in which the solid phases involved are Fe, Fe(OH)2 and Fe(OH)3. In a neutral or basic medium, hydroxides appear first, then transform into oxy-hydroxides such as FeOOH, then in stable oxides Fe3O4 and Fe2O3. In both diagrams, the dashed lines delimit the stability domain of water under atmospheric pressure. Above the upper line, whose equation is E/ Volt = 1.23 – 0.06 pH, water is oxidized into gaseous oxygen. Below the lower line whose equation is E/Volt = – 0.06 pH, water is reduced into hydrogen. The stability domain of the main aqueous species (Fe++, Fe+++ in acidic medium and HFeO2– in basic medium) depends on its concentration. The values (0, –2, –4, –6) shown on the curves represent log10 (ci/mol·L–1), where ci represents the concentration of the considered species (Fe++, Fe+++ or HFeO2–). The potential at equilibrium Fe/Fe(OH)2 does not depend on the pressure whereas the potential at equilibrium H2/H+ decreases when the pressure increases. Below 74 MPa of hydrogen pressure, E(H2/H+) > E{Fe/Fe(OH)2} so that water corrodes iron, even in a reducing medium. Above 74 MPa, E(H2/H+) < E{Fe/Fe(OH)2} so that iron acts as a noble metal and is not more corroded by water in a reducing medium and at pH > 7.
Table 1.
Investigations of the Fe-H-O Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1958Eds]
Oxide equilibria under H2-H2O atmospheres
400–1600°C, Fe, FeO, Fe3O4 and liquid domains
[1965Bur]
Hydrogen solubility in liquid iron
1600°C (continued)
DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–H–O
3
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1970Gra]
Thermogravimetry
800–1000°C, Fe under H2-H2O atmospheres
[1973Sur]
Thermogravimetry
500°C, Fe under H2-H2O atmospheres
[1984Fuk]
XRD, micrography
< 1000°C, Fe under H2-H2O atmospheres, p < 20 GPa
[1987Ito]
XRD, ionic and electronic conductivity measurements
25°C. Preparation of HFe11O17,H2O par ion exchange
[1994Nag]
Thermogravimetry, chemical analysis, kinetics
1400–1600°C, H2O gas on liquid iron or steel, 10 to 30 kPa of H2O
[2005Kre]
XRD, FTIR (Fourier Transform Infrared, Mössbauer)
25°C, αFeOOH (Goethite)
[2005Pla]
MEB analysis
2.104 K, Ar-H2 plasma smelting of iron ores
[2005Was]
Anomalous X-ray scattering
αFeOOH
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(δFe) 1538–1394
cI2 Im 3m W
a = 293.15
at 1390°C [V-C2, Mas2]
a = 286.65
pure Fe at 25°C [Mas2, V-C2] (A2 structure)
(αFe) < 912 (γFe) 1394–912
cF4 Fm 3m Cu
a = 364.67
at 915°C [Mas2, V-C2]
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, 13 GPa [Mas2] Triple point α-γ-ε at 8.4 GPa, 430°C
Fe1–xO (Wuestite) 1422–569
cF8 Fm 3m NaCl
a = 431.0 a = 429.3
0.05 < x < 0.12 [1991Sun] x = 0.05 x = 0.12
Fe3O4 (I) < 580
oP56 Pbcm Fe3O4 I
a = 1186.8 b = 1185.1 c = 1675.2
[V-C2]
Fe3O4 (h) (Magnetite) 1596–580
cF56
a = 839.6
at 25°C
Fd 3m MgAl2O4
a = 854.5
at 1000°C [V-C2] inverse Spinel (continued)
Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
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Fe–H–O
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
αFe2O3 (Hematite) < 1451
hR30 R 3c αAl2O3 (Corundum)
a = 503.42 c = 1374.83
Melts at 1892°C under O2 pressure
βFe2O3
cI80 Ia 3 Mn2O3
a = 939.3
metastable phase [V-C2]
γFe2O3 (Maghemite)
cF56 Fd 3m MgAl2O4
a = 834
metastable phase [1989Rag]
αFeOOH (Goethite)
oP16 Pbnm
a = 460.8 b = 995.6 c = 302.1
[2005Kre] Loses 0.5 H2O at 136°C to give αFe2O3
βFeOOH (Akaganeite)
mI* C2/m
a = 1489.6 b = 303.4 c = 1060.0 β = 135.13°
βFeO1–x(OH)1+xClx [1991Pos]
γFeOOH (Lepidocrocite)
oC16 Cmcm
a = 307.2 b = 1251.6 c = 387.3
[2001Zhu] Loses 0.5 H2O by heating to give γFe2O3
Fe(OH)2
-
-
metastable
Fe(OH)3
-
-
metastable
HFe11O17,H2O
hP64 P63/mmc βAl2O3
a = 593.1 c = 2383
[1987Ito] decomposes at 400°C
DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–H–O
Fig. 1. Fe-H-O.
Landolt-Börnstein New Series IV/11D4
5
Stability domains of solid and liquid phases under H2-H2O atmospheres
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DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
6
Fig. 2. Fe-H-O.
Fe–H–O
Stable potential - pH diagram in aqueous solution at 25°C
DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–H–O
Fig. 3. Fe-H-O.
Landolt-Börnstein New Series IV/11D4
7
Metastable potential - pH diagram in aqueous solution at 25°C
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DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
8
Fe–H–O
References [1940Wib] Wiberg, M., “On the Reduction of Iron Oxides with Carbon Oxides, Water or Methane” (in Swedish), Jernkontorets Ann., 124(5), 179–212 (1940) (Experimental, Morphology, Phase Relations, Thermodyn., 38) [1945Dar] Darken, L.S., Gurry, G.W., “The System Iron-Oxygen. I- The Wuestite Field and Related Equilibria”, J. Am. Chem. Soc., 67, 1398–1412 (1945) (Experimental, Phase Diagram, Phase Relations, Thermodyn., *, 26) [1946Dar] Darken, L.S., Gurry, G.W., “The System Iron-Oxygen. II- Equilibria and Thermodynamics of Liquid Oxides and Other Phases”, J. Am. Chem. Soc., 68, 798–816 (1946) (Experimental, Phase Diagram, Phase Relations, Thermodyn., *, 24) [1958Eds] Edstroem, J.O., “Some Chemical Reactions Involved in Pelletizing and Iron Ore Reduction” (in Swedish), Jernkontorets Ann., 142, 401–466 (1958) (Crys. Structure, Experimental, Phase Relations, 40) [1963Pou] Pourbaix, M., “Atlas of Electrochemical Equilibria at 25°C” (in French), Gauthier-Villlars (Ed.), Paris, “Iron”, Sect. 12-1, 307–329 (1963) (Thermodyn., Phase Relations, Review, 18) [1965Bur] Burylev, B.P., “Solubility of Hydrogen in Liquid Iron Alloys” (in Russian), Izv. Vyss. Uchebn. Zaved., Chern. Metall., 8(2), 17–22 (1965) (Experimental, Phase Relations, Thermodyn., 13) [1970Gra] Grabke, H.J., Best, K.J., Gala, A., “Oxydation Reaction of Metal and Oxide under CO2-CO and H2O-H2 Mixtures” (in German), Werkstoffe und Korrosion, 21, 911–916 (1970) (Phase Relations, Thermodyn., Kinetics, 22) [1973Sur] Surmann, P.L., “The Oxidation of Iron at Controlled Oxygen Partial Pressures – I. Hydrogen/ Water Vapour”, Corros. Sci., 13(2), 113–124 (1973) (Experimental, Kinetics, Morphology, Thermodyn., 10) [1976Kru] Krueger, J., Kunze, H.D., Schuermann, E., “Gas and Carbon in Metals” (in German), Fromm, E., Gebhardt, E. (Eds.), Springer-Verlag, Berlin, 518–613 (1976) (Review, Phase Relations, Thermodyn., 241) [1980Fro] Fromm, E., Hoerz, G., “Hydrogen, Nitrogen, Oxygen, and Carbon in Metals”, Int. Mater. Rev., (5/6), 269–312 (1989) (Kinetics, Phase Diagram, Thermodyn., Review, 255) [1984Fuk] Fukai, Y., “The System Iron-Water-Hydrogen under High Pressure” (in Japanese), Nippon Kinzoku Gakkai Kaiho, 23(5), 369–372 (1984) (Experimental, Phase Relations, Thermodyn., 20) [1987Ito] Ito, S., Kubo, N., Nariki, S., Yoneda, N., “Ion Exchange in Alkali Layers of Potassium β-Ferrite ((1+x)K2O·11Fe2O3) Single Crystals”, J. Am. Ceram. Soc., 70(2), 874–879 (1987) (Crys. Structure, Experimental, 29) [1989Oet] Oeters, F., “Metallurgy of the Steel Production” (in German), Springer-Verlag, 503 pp. (1989) (Phase Diagrams, Phase Relations, Thermodyn., Review, 723) [1989Rag] Raghavan, V., “The Fe-H-O (Iron-Hydrogen-Oxygen) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Metals, Calcutta, 5, 140 (1989) (Phase Relations, Review, Thermodyn., 5) [1990San] San Martin, A., Manchester, F.D., “The Fe-H (Iron-Hydrogen) System”, Bull. Alloys Phase Diagrams, 11(2), 173–184 (1990) (Phase Diagram, Phase Relations, Review, Thermodyn., 86) [1991Pos] Post, E., Buckwald, V.F., “Crystal Structure Refinement of Akaganeite”, Amer. Mineralogist, 76, 272–277 (1991) (Crys. Structure, Experimental, 17) [1991Sun] Sundman, B., “An Assessment of the Fe-O System”, J. Phase Equilib., 12(1), 127–140 (1991) (Phase Diagram, Phase Relations, Thermodyn., Assessment, 53) [1994Nag] Nagasaka, T., Fruehan, R.J., “Kinetics of the Reaction of H2O Gas with Liquid Iron”, Metall. Mater. Trans., B, 25B(2), 245–253 (1994) (Kinetics, Experimental, 21) [1995Kow] Kowalski, M., Spencer, P.J., “Thermodynamic Reevaluation of the Cr-O, Fe-O and Ni-O Systems: Remodelling the Liquid, BCC and FCC Phases”, Calphad, 19(3), 229–243 (1995) (Assessment, Phase Diagram, Phase Relations, Thermodyn., Review, 47) [2001Zhu] Zhukhlistov, A.P., “Crystal Structure of Lepidocrocite FeO(OH) from the ElectronDiffractometry Data”, Crystallography Rept., 46(5), 730–733 (2001), translated from Kristallografiya, 46(5), 805–808 (2001) (Crys. Structure, Experimental, 14) DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–H–O [2003Fuk]
[2003Per]
[2005Kre]
[2005Pla] [2005Was]
[Mas2] [V-C2]
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9
Fukai, Y., Mori, K., Shinomiya, H., “The Phase Diagram and Superabundant Vacancy Formation in Fe-H Alloys under High Hydrogen Pressures”, J. Alloys Compd., 348, 105–109 (2003) (Phase Diagram, Phase Relations, Experimental, 42) Perrot, P., Foct, J., “Gases other than Hydrogen in Iron and Steels” (in French) Techniques de l’Ingenieur, Ser. Metalllurgie, M-4275, 1–23 (2003) (Phase Relations, Thermodyn., Review, 127) Krehula, S., Music, S., Popovic, S., “Influence of Ni-dopant on the Properties of Synthetic Goethite”, J. Alloys Compd., 403, 368–375 (2005) (Crys. Structure, Experimental, Magn. Prop., Optical Prop., 30) Plaul, J.F., Krieger, W., Baeck, E., “Reduction of Fine Ores in Argon-Hydrogen Plasma”, Steel Res. Int., 76(8), 548–554 (2005) (Experimental, Kinetics, Phase Relations, 14) Waseda, Y., Suzuki, S., Saito, M., “Atomic-scale Structure in α-FeOOH Particles Estimated from the Anomalous X-ray Scattering Data with Reverse Monte Carlo Simulation”, J. Alloys Compd., 401, 24–28 (2005) (Calculation, Crys. Structure, Experimental, 16) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
MSIT®
DOI: 10.1007/978-3-540-78644-3_8 # Springer 2008
Fe–H–P
1
Iron – Hydrogen – Phosphorus Pierre Perrot
Introduction Hydrogen and phosphorus are known to play an important role in steelmaking because of their deleterious effect on the mechanical behavior of the steels. Hydrogen causes embrittlement and phosphorus worsens the situation because of the existence of a low melting (977°C) ternary eutectic in the C-Fe-P system. Most of the investigations, gathered in Table 1, have been directed towards the influence of P on the hydrogen solubility in Fe. Binary Systems The Fe-H system under 0.1 MPa of hydrogen pressure is accepted from [1990San]. The Fe-P diagram is mainly accepted from [2002Per]. The main hydrogen phosphide is the gaseous PH3, an endothermal compound which has never been reported in steelmaking. Solid Phases Solid phases are shown in Table 2. No ternary hydrides are known. Isothermal Sections Phosphorus reduces slightly the hydrogen solubility in liquid iron [1963Wei, 1966Noz, 1981Sch, 1988Rag] in agreement with the calculations of [1993Din]. Figure 1 shows the hydrogen solubility in liquid Fe-P alloys measured at four temperatures by [1981Sch]. Thermodynamics The Wagner first order interaction parameter eH(P) has been experimentally determined at 1592°C by [1963Wei], between 1450 and 1620°C in liquid (Fe,P) alloys by [1966Noz]. It is defined by eH(P) = {∂ log10 fP/∂ (% P)} where (% P) represents the phosphorus content of the alloy expressed in mass% and fP the activity coefficient of P in the alloy defined by fP = (% P in pure iron)/(% P in the alloy). It is calculated from the solubility measurements at constant temperature and hydrogen pressure and has been measured at eH(P) = 0.01 in the temperature range 1450–1620°C for less than 6 mass% P in the alloy. The interaction coefficient may also be defined by ɛH(P) = {∂ ln γP/∂ xP}, where xP is the mole fraction of P in the liquid alloy and γP = (xP in pure iron)/(xP in the alloy). The careful investigation of [1981Sch] leads to ɛH(P) = – 0.4495 + (3166/T ) between 1420 and 1610°C for less than 1 at.% P in good agreement with the preceding relation. This experimental result contradicts the calculation of [1999Din], based on the Miedema’s model of the heat of formation of binary alloys, which proposes ɛH(P) = 15.7. However, it must be pointed out that the Miedema’s model can hardly be extrapolated to non metallic elements. Notes on Materials Properties and Applications Fe-P alloys sintered in H2 or N2 form soft magnetic materials whose use is promising in micromotors [2001Chi].
Landolt-Börnstein New Series IV/11D4
MSIT®
DOI: 10.1007/978-3-540-78644-3_9 # Springer 2008
2 Table 1.
Fe–H–P Investigations of the Fe-H-P Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1963Wei]
Hydrogen solubility measurements, Sievert’s method
1592°C, < 1 mass% P, < 0.1 MPa of H2 pressure
[1966Noz]
Hydrogen solubility measurements, Sievert’s method
1450–1670°C, < 6 mass% P, < 0.1 MPa of H2 pressure
[1981Sch]
Hydrogen solubility measurements, Sievert’s method
1420–1610°C, xP < 0.20, < 0.1 MPa of H2 pressure
[2001Chi]
Magnetization curves, hardness measurements
Fe-P (< 1 mass% P) powders sintered. Soft magnetic material preparation
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(P) (red) < 417
c*66
a = 1131
sublimation at 1 bar. Stable form of P. Triple point at 576°C, > 36.3 bar; triple point at 589.6°C at 1 atm [Mas2, V-C2]
(P) (white) < 44.14
c** P (white)
a = 718
common form of P [Mas2, V-C2]
(P) (black)
oC8 Cmca P (black)
a = 331.36 b = 1047.8 c = 437.63
at 25°C [Mas2, V-C2]
(δFe) 1538–1394
cI2 Im 3m W
a = 293.15
pure Fe at 1390°C [V-C2, Mas2] dissolves up to 4.52 at.% P at 1048°C at 1394°C pure Fe at 25°C [Mas2, V-C2]
(αFe) < 912°C
a = 286.65
(γFe) 1394–912
cF4 Fm 3m Cu
a = 364.67
pure Fe at 915°C [Mas2, V-C2] dissolves up to 0.5 at.% P at 1150°C
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, 13 GPa [Mas2] triple point α-γ-ε at 8.4 GPa, 430°C
Fe3P < 1166
tI32 I 4 Ni3P
a = 910.8 c = 445.5
at 25°C [Mas2, V-C2]
a = 917.4 c = 453.0
at 678°C [1990Oka] (continued)
DOI: 10.1007/978-3-540-78644-3_9 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–H–P
3
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
Fe2P < 1370
hP9 P 62m Fe2P
a = 586.4 c = 346.0
33.3–34 at.% P [V-C2]
Fe2P (HP)
oP12 Pnma Co2Si
a = 577.5 b = 357.1 c = 664.1
at 800°C under 8 GPa [V-C2]
FeP < 1370
oP8 Pna21 AsCo or oP8 Pnma MnP
a = 519.3 b = 579.2 c = 309.9
[2002Per]
a = 520.8 b = 316.0 c = 581.2
[V-C2]
FeP2
oP6 Pnnm FeS2 (marcasite)
a = 497.29 ± 0.07 b = 565.68 ± 0.08 c = 272.30 ± 0.04
[V-C2]
FeP4
mP30 P21/c FeP4
a = 461.9 ± 0.1 b = 1367.0 ± 0.2 c = 700.2 ± 0.1 β = 101.48°
[V-C2]
Fe4P (HP)
oC20 C2221
a = 500.5 ± 0.1 b = 1021.3 ± 0.3 c = 553.0 ± 0.1
at 1100°C under 6 GPa [V-C2]
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DOI: 10.1007/978-3-540-78644-3_9 # Springer 2008
4
Fig. 1. Fe-H-P.
Fe–H–P
Hydrogen solubility in Fe-P liquid alloys under 1 atm H2
DOI: 10.1007/978-3-540-78644-3_9 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–H–P
5
References [1963Wei] Weinstein, M., Elliott, J.F., “Solubility of Hydrogen in Liquid Iron Alloys”, Trans. Met. Soc. AIME, 227, 382–393 (1963) (Calculation, Experimental, Thermodyn., 27) [1966Noz] Nozaki, H., Ban-ya, S., Fuwa, T., Matoba, S., Ono, K., “Effect of Carbon, Silicon, Phosphorus and Nickel on the Solubility of Hydrogen in Liquid Iron” (in Japanese), Tetsu to Hagane, 52(13), 1823–1833 (1966) (Experimental, Phase Relations, Thermodyn., 17) [1981Sch] Schuermann, E., Kaettlitz, W., “Equivalent Effect of the Alloying Elements on the Concentration- and Temperature-Dependent Hydrogen Solubility in Iron-Rich Ternary and Multicomponent Melts” (in German), Arch. Eisenhuettenwes., 52(8), 295–301 (1981) (Calculation, Experimental, Phase Relations, Thermodyn., 20) [1988Rag] Raghavan, V., “The Fe-H-P (Iron-Hydrogen-Phosphorus) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Metals, Calcutta, Vol. 3, 80 (1988) (Review, Phase Relations, Thermodyn., 3) [1990Oka] Okamoto, H., “The Fe-P (Iron-Phosphorus) System”, Bull. Alloy Phase Diagrams, 11(4), 404–412 (1990) (Phase Diagram, Phase Relations, Review, 89) [1990San] San Martin, A., Manchester, F.D., “The Fe-H (Iron-Hydrogen) System”, Bull. Alloys Phase Diagrams, 11(2), 173–184 (1990) (Phase Diagram, Phase Relations, Review, Thermodyn., 86) [1993Din] Ding, X., Wang, W., Han, Q., “Thermodynamic Calculation of Fe-P-H System Melt”, Acta Metall. Sin. (China), 29(12), B527–B532 (1993) (Calculation, Theory, Thermodyn., 7) [1999Din] Ding, X., Fan, P., Wang, W., “Thermodynamic Calculation for Alloy Systems”, Metall. Trans B, 30B(2), 271–277 (1999) (Thermodyn., Theory, 18) [2001Chi] Chistyakov, V.K., Dorogina, G.A., Korobka, O.B., “Influence of a Protective Medium on the Magnetic and Mechanical Properties of Materials of the Iron - Phosphorus System”, Powder Metall. Met. Ceram., 40(9–10), 526–532 (2001) (Experimental, Magn. Prop., Mechan. Prop., 19) [2002Per] Perrot, P., Batista, S., Xing, X., “Fe-P (Iron-Phosphorus)”, Diagrams as Published in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.16107.1.20, (2002) (Phase Diagram, Phase Relations, Crys. Structure, 23) [Mas2] Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) [V-C2] Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
Landolt-Börnstein New Series IV/11D4
MSIT®
DOI: 10.1007/978-3-540-78644-3_9 # Springer 2008
Fe–H–Si
1
Iron – Hydrogen – Silicon Pierre Perrot
Introduction Fe-Si alloys, in presence of H do not form hydrides, but interstitial solid solutions which play an important role in steelmaking because of undesirable effects in the embrittlement of steel products. On the other hand, alloying Si in iron increases the activation energy of H diffusion. Most of the investigations have been directed towards the H solubility measurements in Fe-Si melts. The main experimental results are gathered in Table 1. Binary Systems The Fe-Si system, accepted from [Mas2], has been later assessed by [1999Liu]. The Fe-H system under 0.1 MPa of hydrogen pressure is accepted from [1990San]. The hydrogen solubility in liquid iron may be expressed by [1984Fro]: log10xH = 0.5 log10 (pH2/Pa) – 4.44 – (1720/T) The Fe-H system has been carefully investigated by [2003Fuk] up to 10 GPa of hydrogen pressure and 1500°C. The hydrogen solubility in liquid Si under 0.1 MPa of hydrogen pressure is 0.087 at.% at 1414°C according to [Mas2], value which does not agree the experimental value of 0.020 at.% at 1600°C reported by [1981Sch] and accepted by [1984Fro]. Solid Phases The solid phases are presented in Table 2. No ternary compounds are known. Isothermal Sections A number of investigations [1977Bes] on the solubility of H in liquid Fe-Si alloys are reported. [1970Fuk], using a Sievert’s method, reported first order Wagner interaction parameters to be positive (eHSi = 0.03 at 1600°C), indicating that Si increases the activity coefficient of H in the melt and decreases its solubility at a given H2 pressure. The solubility of hydrogen in liquid Si decreases also with the iron content. At 1600°C, the minimum of solubility is observed at 0.02 at.% H in a liquid melt containing 60 at.% Si [1971Gel, 1971Pet, 1973Gel], value justified by a theoretical model proposed by [1984Fro]. The minimum of hydrogen solubility corresponds to a maximum for the enthalpy of hydrogen dissolution in liquid alloys as shown in Table 3 [1971Kos, 1977Bes]. The H solubility in Fe-Si liquid alloys measured at 1500, 1600 and 1700°C under 0.1 MPa of hydrogen pressure is shown in Fig. 1 for xSi < 0.25 [1981Sch]. The hydrogen solubility calculated at 1600°C by a statistical model at 1600°C in the whole composition range [1984Fro] is shown in Fig. 2. The solubility of H in solid alloys, investigated by [1971Rya] between 300 and 700°C, decreases with the silicon content of the alloy. Thermodynamics The Wagner first order interaction parameter, defined by eH(Si) = {∂ log10 fSi/∂ (% Si)}, where (% Si) represents the silicon content of the alloy expressed in mass% with fSi = (% Si in pure iron)/(% Si in the alloy), has been early experimentally determined at eH(Si) = 0.030 ± 0.08 [1970Fuk, 1971Kos, 1972Ngi]. A more precise value eH(Si) = 0.0286 at 1600°C was measured by [1974Boo], which corresponds to ɛH(Si) = 3.81 [1977Bes]. The interaction coefficient ɛH(Si) is defined by: ɛH(Si) = {∂ ln γSi/∂ xSi}, where xSi is the mole fraction of Si in the liquid alloy and γSi = (xSi in pure iron)/(xSi in the alloy). The careful investigation of [1981Sch] leads to ɛH(Si) = – 0.4467 + (3498/T) between 1500 and 1700°C for less than 1 at.% Si in agreement with the preceding relation and in contradiction with the probably too high values ɛH(Si) = 4.20 proposed by [1973Gel] at 1600°C and ɛH(Si) = 8.0 calculated by [1999Din] on the basis of Miedema’s model for the estimation of the heat of formation of binary alloys. Landolt-Börnstein New Series IV/11D4
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2
Fe–H–Si
Notes on Materials Properties and Applications Hydrogen plays an important role in environmental embrittlement of Fe-Si alloys. These alloys may be hardened of softened by hydrogen depending on the temperature and compositions [1988Kim] whereas high purity iron is always softened by hydrogen. Miscellaneous Hydrogen diffusion in metallic glasses gives an insight on the glass structure. So, H diffusion in Fe-Si alloys presents a deviation from Arrhenius’ law [1999Eli] which is explained by the temperature dependence of the short range order in Fe-Si alloy. The refining effect of Hydrogen Plasma Arc Melting (HPAM) was investigated for impurities in Fe [2006Ela]. Under 5.3 kPa of H, the Si content of Fe has been decreased from 3.5 to 0.9 ppm whereas, under 100 kPa (1 bar) of H, the Si content decreases from 3.5 to 0.6 ppm. The removal degree of impurities seems lower under H than under Ar plasma.
Table 1.
Investigations of the Fe-H-Si Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1970Fuk]
Hydrogen solubility in liquid alloys, Sievert’s method
1550–1600°C, < 2.5 mass% Si
[1971Gel]
Hydrogen solubility in liquid alloys
1500–1700°C, 0 to 100 % Si, 0.1 MPa of H2 pressure
[1971Kos]
Hydrogen solubility in liquid alloys, hot volume method
1415–1610°C, 0 to 100 % Si, 0.1 MPa of H2 pressure
[1971Pet]
Hydrogen solubility in liquid alloys
1415–1600°C, 0 to 100 % Si, 0.1 MPa of H2 pressure
[1971Rya]
Hydrogen absorption isotherms
300–700°C, < 6 mass% Si
[1972Ngi]
Hydrogen solubility in liquid alloys, Sievert’s method
1550–1700°C, < 10 mass% Si, 0.1 MPa of H2 pressure
[1974Boo]
Hydrogen solubility in liquid alloys, constant volume method
1560–1800°C, < 12 mass% Si, 0.1 MPa of H2 pressure
[1981Sch]
Hydrogen solubility in liquid alloys, Sievert’s method
1500–1700°C, < 25 mol% H2, 0.1 MPa of H2 pressure
[1988Kim]
Hardness measurements
< 5 mass% Si
[2006Ela]
Hydrogen plasma arc melting, Si impurities removal in Fe
2000–2300°C, 5.3 and 100 kPa of H pressure
DOI: 10.1007/978-3-540-78644-3_10 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–H–Si Table 2.
3
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αFe) (Ferrite) < 912°C
cI2 Im 3m W
a = 286.65
pure Fe at 25°C [Mas2, V-C2] dissolves up to 10 at.% Si at 500°C [Mas2]. at 1390°C [V-C2, Mas2] dissolves up to 19.5 at.% Si at 1280°C [Mas2].
(δFe) 1538–1394
a = 293.15
(γFe) (Austenite) < 1394–912
cF4 Fm 3m Cu
a = 364.67
at 915°C [V-C2, Mas2]
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, > 13 GPa [Mas2]
(Si) < 1414
cF8 Fd 3m C (diamond)
a = 543.06
at 25°C [Mas2]. Practically no Fe solubility [Mas2].
FeSi < 1410
cP8 P213 FeSi
a = 451.7 ± 0.5
at 300°C [V-C2, Mas2]
Fe2Si 1212–1040
cP2 Pm 3m CsCl or hP6 P 3m1 Fe2Si
a = 282
66 at.% Si [Mas2, V-C2]
Fe5Si3 1060–825
hP16 P63/mcm Mn5Si3
a = 675.9 ± 0.5 c = 472.0 ± 0.5
38.4 at.% Si [V-C2, Mas2]
βFeSi2(h) 1220–937
tP3 P4/mmm βFeSi2
a = 269.01 c = 513.4
69.5 to 73.5 at.% Si [Mas2, V-C2]
αFeSi2(r) < 982
oC48 Cmca αFeSi2
a = 986.3 ± 0.7 b = 779.1 ± 0.6 c = 783.3 ± 0.6
66.7 at.% Si [Mas2, V-C2]
α1, Fe3Si
cF16 Fm 3m BiF3
a = 565
10 to 30 at.% Si [Mas2, V-C2]
α2, Fe3Si
cP2 Pm 3m CsCl
-
10 to 22 at.% Si [Mas2]
Landolt-Börnstein New Series IV/11D4
a = 405.2 ± 0.2 c = 508.55 ± 0.3
MSIT®
DOI: 10.1007/978-3-540-78644-3_10 # Springer 2008
4 Table 3.
Fe–H–Si Thermodynamic Data of Reaction or Transformation
Reaction or Transformation
Temperature [°C]
Quantity, per mole of atoms [J, mol, K]
Comments
H ⇌ {H} in pure Fe
1600
ΔH = – 187000 ΔS = – 0.6
[1971Kos]
H ⇌ {H} in Fe65Si35
1600
ΔH = – 119000 ΔS = – 6.2
[1971Kos]
H ⇌ {H} in Fe50Si50
1600
ΔH = – 65500 ΔS = – 29
[1971Kos]
H ⇌ {H} in Fe33Si67
1600
ΔH = – 56500 ΔS = – 27
[1971Kos] ΔH is maximum
H ⇌ {H} in pure Si
1600
ΔH = – 106500 ΔS = – 35.5
[1971Kos]
Fig. 1. Fe-H-Si. pressure
Hydrogen solubility in Fe-Si liquid alloys at 1500, 1600 and 1700°C under 0.1 MPa of hydrogen
DOI: 10.1007/978-3-540-78644-3_10 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–H–Si
Fig. 2. Fe-H-Si.
Landolt-Börnstein New Series IV/11D4
5
Hydrogen solubility in Fe-Si liquid alloys at 1500, 1600 and 1700°C in the whole composition range
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DOI: 10.1007/978-3-540-78644-3_10 # Springer 2008
6
Fe–H–Si
References [1970Fuk] Fukuda, S., Sugiyama, T., Furukawa, T., Kato, E., “Solubility of Hydrogen in Liquid Iron Alloys”, Rep. Casting Research Lab./Waseda Univ., (21), 35–46 (1970) (Calculation, Experimental, Phase Relations, Thermodyn., 22) [1971Gel] Geld, P.V., Kostina, T.K., Baum, B.A., Kurochkin, K.T., “Thermodynamic Characteristics of the Dissolution of Hydrogen in Binary Alloys of Fe, Co, Ni, and Cr with Si”, Russ. J. Phys. Chem., 45(8), 1134–1136 (1971), translated from Zh. Fiz. Khim., 45(8), 1999–2003 (1971) (Experimental, Thermodyn., 2) [1971Kos] Kostina, T.K., Baum, B.A., Kurochkin, K.T., Geld, P.G., “Solubility of Hydrogen in Liquid of Iron-Silicon Alloys”, Russ. J. Phys. Chem., 45(4), 453–455 (1971), translated from Zh. Fiz. Khim., 45(4), 813–816 (1971) (Experimental, Phase Relations, Thermodyn., 13) [1971Pet] Petrushevskiy, M.S., Geld, P.V., Baum, B.A., Kostina, T.K., “Calculation of the Solubility of Hydrogen in Fe-Si and Mn-Si Melts”, Russ. Metall., (2), 38–42 (1971), translated from Izv. Akad. Nauk SSSR, Met., (2), 59–64 (1971) (Thermodyn., Experimental, Phase Relations, 13) [1971Rya] Ryabov, R.A., Saliy, V.I., Tesler, M.L., Geld, P.V., “Diffusion and Solubility of Hydrogen in Solid Solutions of the Fe-Si System” (in Russian) in “Metody Opred. Gazov v Metal. Splavakh”, Moscow, 77–79 (1971) (Experimental, Phase Relations, 6) [1972Ngi] Ngia, N., Yavoyskiy, V.I., Kosterev, L.B., Afanas’yev, M.I., “Hydrogen Solubility in Binary Iron-Base Alloys”, Russ. Metall., (4), 11–15 (1972), translated from Izv. Akad. Nauk SSSR, Met., (4), 18–22 (1972) (Experimental, Phase Relations, Thermodyn., 19) [1973Gel] Geld, P.V., Baum, B.A., Petrushevsky, M.S., Kostina, T.K., Kurochkin, K.T., “Solubility of Hydrogen in Melts of 3d-transition Metals with Silicon” (in Russian) in “Vzaimodeistvie Gazov s Metal.”, Trudi 3 Sov. Yaponskogo Simp. Fiz.-Khim. Osn. Met. Protsessov, Nauka, Moscow, 173–181 (1973) (Calculation, Review, Phase Relations, 10) [1974Boo] Boorstein, W.M., Pehlke, R.D., “Measurement of Hydrogen Solubility in Liquid Iron Alloys Employing a Constant Volume Technique”, Metall. Trans., 5(2), 399–405 (1974) (Calculation, Experimental, Phase Relations, 29) [1977Bes] Bester, H., Lange, K.W., “Hydrogen Solubility in Iron and in Liquid Fe-Mn, Fe-Cr, and Fe-Si Alloys” (in German), Stahl Eisen, 97, 1037–1039 (1977) (Thermodyn., Phase Relations, Review, 53) [1981Sch] Schuermann, E., Kaettlitz, W., “Equivalent Effect of the Alloying Elements on the Concentration- and Temperature-Dependent Hydrogen Solubility in Iron-Rich Ternary and Multicomponent Melts” (in German), Arch. Eisenhuettenwes., 52(8), 295–301 (1981) (Calculation, Experimental, Phase Relations, 20) [1984Fro] Frohberg, M.G., Anik, S., “The Prediction of Hydrogen Solubility in Binary Iron Melts” (in German), Arch. Eisenhuettenwes., 55(2), 45–48 (1984) (Calculation, Phase Relations, Thermodyn., 20) [1988Kim] Kimura, T., Matsui, H., Kimura, H., “Effect of Hydrogen in the Mechanical Properties of Iron Based Alloys” in “Strength of Metals and Alloys”, Proc 8th Int. Conf. ISCMA, Pergamon Press, Oxford, 541–546 (1988) (Experimental, Mechan. Prop., 5) [1990San] San Martin, A., Manchester, F.D., “The Fe-H (Iron-Hydrogen) System”, Bull. Alloys Phase Diagrams, 11(2), 173–184 (1990) (Phase Diagram, Phase Relations, Review, Thermodyn., 86) [1999Din] Ding, X., Fan, P., Wang, W., “Thermodynamic Calculation for Alloy System”, Metall. Mater. Trans B, 30B(2), 271–277 (1999) (Thermodyn., Theory, 17) [1999Eli] Eliaz, N., Fuks, D., Eliezer, D., “A New Model for the Diffusion Behaviour in Metallic Glasses”, Acta Mater., 47(10), 2981–2989 (1999) (Theory, Transport Phenomena, Interface Phenomena, 47) [1999Liu] Liu, Z.-K., Chang, Y.A., “Thermodynamic Assessment of the Al-Fe-Si System”, Met. Mat. Trans. A, 30A(4), 1081–1095 (1999) (Thermodyn., Assessment, Phase Diagrams, Phase Relations, 56) [2003Fuk] Fukai, Y., Mori, K., Shinomiya, H., “The Phase Diagram and Superabundant Vacancy Formation in Fe-H Alloys under High Hydrogen Pressures”, J. Alloy Compd., 348, 105–109 (2003) (Phase Diagram, Phase Relations, Experimental, 42) DOI: 10.1007/978-3-540-78644-3_10 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–H–Si [2006Ela]
[Mas2] [V-C2]
Landolt-Börnstein New Series IV/11D4
7
Elanski, D., Lim, J.-W., Mimura, K., Isshiki, M., “Impurity Removal from Fe, Cr, Ti, and V Metals by Hydrogen Plasma Arc Melting and Thermodynamic Estimation of Hydride and Sulfide Formation”, J. Alloys Compd., 421, 209–216 (2006) (Thermodyn., Experimental, Calculation, Phase Relations, 8) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
MSIT®
DOI: 10.1007/978-3-540-78644-3_10 # Springer 2008
Fe–H–V
1
Iron – Hydrogen – Vanadium Pierre Perrot
Introduction Vanadium metal inside a bcc structure is a promising hydrogen storage material with a high capacity because it can absorb hydrogen up 50 at.% under 0.1 MPa of hydrogen pressure at temperature not higher than 400°C [2002Yuk]. Unfortunately, hydrogen is hardly released and iron makes easier the hydrogen release. Most investigations have been directed towards the solubility of H in Fe-V melts. The main experimental results are gathered in Table 1. Binary Systems The Fe-H diagram under 0.1 MPa H2 is accepted from [1990San]. The Fe-H system has been carefully investigated by [2003Fuk] up to 10 GPa of hydrogen pressure and 1500°C. The hydrogen solubility in vanadium under pressures lower than 0.1 MPa was measured by [1938Kub] and an equilibrium diagram under 0.1 MPa H2 and up to 425°C was proposed by [1982Smi] and reproduced by [Mas2], diagram accepted in the present report. An extension of the H-V diagram up to the vanadium melting point was presented by [1992Oka], but this diagram cannot be accepted, because too much thermodynamically improbable features are involved. A H-V phase under 5 GPa of H2 was proposed by [2005Fuk]. The Fe-V system, is accepted from the assessment of [1984Smi] reproduced in [Mas2]. Solid Phases The solid phases are presented in Table 2. Vanadium may dissolve up to 50 at.% H above 170°C under 0.1 MPa of hydrogen pressure. Below 175°C, vanadium hydrides, namely β1, β2, γ and δ are formed. No ternary hydrides are known. Isothermal Sections The hydrogen solubility in α(Fe,V) alloys investigated by [1954Kri] from 411 to 900°C, was found to decrease with the increase of temperature and vanadium content of the alloy. The same trend was observed α alloys by [1974Egu]. The existence of a minimum in the hydrogen solubility on Fe-V alloy is thus probable. In the liquid state [1962Yak, 1981Sch], the hydrogen solubility in (Fe,V) alloys increases with increasing the V content. The hydrogen solubility in Fe-V alloys (< 25 at.% V) between 1550 and 1650°C is shown in Fig. 1. Thermodynamics The Wagner first order interaction parameter, defined by eH(V) = {∂ log10 fV/∂ (% V)} where (% V) represents the vanadium content of the alloy expressed in mass% with fV = (% V in pure iron)/(% V in the alloy), has been early experimentally determined at eH(V) = – 0.0074 ± 0.0001 between 1500 and 1600°C [1964Gun] and as eH(V) = 0.0087 – (28.9/T ) between 1550 and 1700°C [1967Ban]. An expression of the interaction coefficient εH(V) defined by: εH(V) = {∂ ln γV/∂ xV} where xV is the mole fraction of V in the liquid alloy and γV = (xV in pure iron)/(xV in the alloy) has been proposed by [1981Sch]: εH(V) = 0.2419 – (1704/T ) between 1550 and 1650°C for less than 20 at.% V in agreement with the preceding relation. In solid (V,Fe) alloys (< 1 at.% Fe) εH(Fe) = 4.8 at 927°C [1999Bal], which means that Fe decreases the H solubility in V, in agreement with the observation of [1974Egu].
Landolt-Börnstein New Series IV/11D4
MSIT®
DOI: 10.1007/978-3-540-78644-3_11 # Springer 2008
2
Fe–H–V
Notes on Materials Properties and Applications Hydrogen plays an important role in environmental embrittlement of Fe-V alloys. These alloys may be hardened or softened by hydrogen depending on the temperature and compositions [1988Kim] whereas high purity iron is always softened by hydrogen. Vanadium absorbs easily hydrogen [2002Yuk]. At 70°C, the isobaric plateau forms under 0.1 Pa of hydrogen pressure, so that hydrogen is not easily released. The presence of 3 at.% Fe raises the plateau up to 1 Pa of hydrogen pressure. Miscellaneous Hydrogen pick-up lowers the superconducting transition of Fe-V alloys [1979Ohl, 1981Ohl] and increases the saturation magnetization [2002Uzd]. In non magnetic Fe-V alloys (< 20 at.% Fe), absorption of H generates local magnetic moments and a superparamagnetic contribution. H follows V in Fe-V alloys and the H distribution in rapidly quenched alloys is correlated with Fe segregation [2001Hom], a result which is in agreement with the decreasing of H solubility when the iron content of the alloy increases.
Table 1.
Investigations of the Fe-H-V Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1954Kri]
Hydrogen solubility measurements in solid alloys
400–900°C, < 50 at.% V, 0.1 MPa of H2 pressure
[1962Yak]
Hydrogen solubility measurements in liquid alloys
1560°C, < 10 mass% V, 0.1 MPa of H2 pressure
[1964Gun]
Hydrogen solubility in liquid alloys by the hot volume method
1500–1650°C, < 35 mass% V, 0.1 MPa of H2
[1967Ban]
Hydrogen solubility measurements in liquid alloys
1548–1672°C, < 60 mass% V, 0.1 MPa of H2
[1974Egu]
Hydrogen solubility measurements in solid Fe-V alloys
600–1200°C, < 10 at.% Fe, 0.1 MPa of H2 pressure
[1979Ohl, 1981Ohl]
Cp measurements, Cu-tabletting method
5–15 K, V2H, VH, VH2, V0.9Fe0.1Hn (n < 1)
[1981Sch]
Hydrogen solubility measurements in liquid alloys
1550–1650°C, < 25 at.% V, 0.1 MPa of H2 pressure
[1988Kim]
Hardness measurements
< 5 mass% V
[1992Ost]
Mössbauer at 4.2 K
Magnetic behavior
[1999Bal]
Hydrogen solubility measurements in solid Fe-V alloys
927°C, < 1 at.% Fe, 0.1 MPa of H2 pressure
[2001Hom]
Tritium radioluminography, X-ray dispersive micrography
V0.95Fe0.05Hn, H distribution in alloys
[2002Uzd]
Faraday’ balance
Saturation magnetization measures
[2002Yuk]
Coulometric titration, hydrogen absorption measurements
V0.97Fe0.03Hn
DOI: 10.1007/978-3-540-78644-3_11 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–H–V Table 2.
3
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
α, (αδFe,V) (δFe) 1538–1394 (αFe) < 912 (V) < 1910
cI2 Im 3m W
a = 293.15
1394°C [Mas2]
a = 286.65
at 25°C [Mas2]
a = 302.40
pure V at 25°C [Mas2] (A2 structure) dissolves up to 50 at.% H above 220°C [1982Smi] > 180, limit of H solubility into V under 0.1 MPa H2 [1982Smi]
a = 303.7 (γFe) 1394–912
cF4 Fm 3m Cu
a = 364.67
at 915°C [Mas2] dissolves up to 1.3 at.% V at 1180°C
σ, FeV < 1252
tP30 P42/mnm σCrFe
a = 886.5 to 901.5 c = 460.5 to 464.2
29.6 to 60.1 at.% V [1984Smi]
β2, V2H (h) < 205
tF6 (?)
a ≈ 447 c ≈ 300
31 to 50 at.% H [1982Smi]
a = 600.12 ± 0.03 c = 661.88 ± 0.09
at 28°C [V-C2]
mC6 C2/m AuTe2 or mC6 Cm V 2H
a ≈ 447 b ≈ 300 c ≈ 447 β ≈ 92° a = 445.66 ± 5 b = 300.22 ± 2 c = 447.60 ± 5 β = 95.61 ± 1°
31 to 45 at.% H [1982Smi] No superconductivity [1981Ohl]
δ, V3H2 < – 70
mC10
a ≈ 452 b ≈ 302 c ≈ 692 β ≈ 77°
40 at.% H [1982Smi]. Structure closely related to β1, V2H. Probably ordered β1, V2H.
VH2 < – 15
cF12 Fm 3m CaF2
a = 427.1
< – 15°C under 0.1 MPa H2 [1982Smi] Superconductor [1981Ohl]
or tI24 I41/amd V 2H β1, V2H (r) < 175
Landolt-Börnstein New Series IV/11D4
MSIT®
at 28°C [V-C2]
DOI: 10.1007/978-3-540-78644-3_11 # Springer 2008
4
Fe–H–V
Fig. 1. Fe-H-V. Hydrogen solubility in liquid Fe-V alloys under 0.1 MPa H2
DOI: 10.1007/978-3-540-78644-3_11 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–H–V References [1938Kub] [1954Kri]
[1962Yak]
[1964Gun]
[1967Ban] [1974Egu]
[1979Ohl]
[1981Ohl]
[1981Sch]
[1982Smi]
[1984Smi] [1988Kim]
[1990San]
[1992Oka] [1992Ost]
[1999Bal]
[2001Hom]
[2002Uzd]
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5
Kubaschewki, O., “Absorption of Gases into the Metals” (in German), Z. Elektrochem., 44(2), 152–167 (1938) (Phase Relations, Review, 187) Krivoglaz, M.A., “Solubility in Homogeneous Alloys” (in Russian), Zh. Tekh. Fiz., 24(6), 1077–1089 (1954) (Crys. Structure, Experimental, Phase Relations, Experimental, Thermodyn., 11) Yakushev, A.M., Yavoiskii, V.I., “Effect of V and B on the Solubility of H in Liquid Fe” (in Russian), Izv. Vyssh. Ucheb. Zaved., Chern. Met., (1), 52–56 (1962) (Experimental, Phase Relations, 5) Gunji, K., Ono, K., Aoki, Y., “The Effect of Various Elements on the Solubility of Hydrogen in Liquid Pure Iron”, Trans. Natl. Res. Inst. Metals, 6(5), 209–213 (1964), translated from J. Jpn. Inst. Met., 28, 64–68 (1964) (Experimental, Phase Relations, Thermodyn., 7) Ban-ya, S., Fuwa, T., Ono, K., “Solubility of Hydrogen in Liquid Iron Alloys” (in Japanese), Tetsu to Hagane, 53(2), 101–116 (1967) (Experimental, Phase Relations, Thermodyn., 18) Eguchi, T., Morozumi, S., “Influence of Alloying Elements on the Solubility of Hydrogen in Vanadium” (in Japanese), J. Jpn. Inst. Met., 38, 1025–1030 (1974) (Experimental, Phase Relations, Thermodyn., 24) Ohlendorf, D., Wicke, E., Obermann, A., “Low Temperature Specific Heat and Magnetic Properties of Fe-V and V-Fe-H Alloys”, J. Phys. Chem. Solids, 40(11), 849–856 (1979) (Experimental, Thermodyn., Magn. Prop., 25) Ohlendorf, D., Severin, H.G., Wicke, E., “Heat Capacity Measurements of Pulverized Metal Hydrides between 1.5 and 15 K by Means of the Cu-Tabletting Method”, Z. Phys. Chem., 128(2), 137–146 (1981) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 9) Schuermann, E., Kaettlitz, W., “Equivalent Effect of the Alloying Elements on the Concentration- and Temperature-Dependent Hydrogen Solubility in Iron-Rich Ternary and Multicomponent Melts” (in German), Arch. Eisenhuettenwes., 52(8), 295–301 (1981) (Experimental, Phase Relations, Thermodyn., 20) Smith, J.F., Peterson, D.T., “The H-V (Hydrogen-Vanadium) System”, Bull. Alloys Phase Diagrams, 3(1), 55–60 (1982) (Phase Diagram, Phase Relations, Crys. Structure, Thermodyn., Review, #, 49) Smith, J.F., “The Fe-V (Iron-Vanadium) System”, Bull. Alloys Phase Diagrams., 5(2), 184–194 (1984) (Phase Diagrams, Phase Relations, Crys. Structure, Thermodyn., Review, #, 99) Kimura, T., Matsui, H., Kimura, H., “Effect of H in the Mechanical Properties of Iron Based Alloys” in “Strength of Metals and Alloys”, Proc. 8th Int. Conf. ISCMA, Pergamon Press, Oxford, 541–546 (1988) (Experimental, Mechan. Prop., 5) San Martin, A., Manchester, F.D., “The Fe-H (Iron-Hydrogen) System”, Bull. Alloys Phase Diagrams, 11(2), 173–184 (1990) (Phase Diagram, Phase Relations, Review, Thermodyn., 86) Okamoto, H., “H-V (Hydrogen-Vanadium)”, J. Phase Equilib., 13(6), 679 (1992) (Phase Diagram, Phase Relations, Review, 6) Ostrasz, A., Szuskiewicz, M., “Study of the Influence of Hydrogen Absorption on the Magnetic Properties of V-Fe-H Alloys by Low Temperature Mössbauer Spectrometry” Hyperfine Interactions, 68, 369–372 (1992) (Experimental, Magn. Prop., Electronic Structure, 6) Balabaeva, R.F., Vasil’eva, I.A., Sukhushina, I.S., Alekseev, I.V., “Thermodynamic Properties and Stability of Nb, Ti, Zr, Hf Intermetallics with Small Content of Transition Metals and Hydrogen (Nitrogen)” (in Russian), Zh. Fiz. Khim., 73(8), 1345–1347 (1999) (Calculation, Phase Relations, Thermodyn., 6) Homa, H., Satoh, H., Misawa, T., Ohnishi, T., “Observation of H Distribution in V-5 at.% Fe Alloys by Tritium Radioluminography”, J. Jpn. Inst. Met., 65(3), 163–166 (2001) (Crys. Structure, Experimental, 11) Uzdin, V., Labergerie, D., Westerhoff, K., Zabel, H., Hjovarsson, B., “Evolution of Atomic Magnetic Moments in Fe-V Multilayers with Hydrogen Loading”, J. Magn. Magn. Mat., 240, 481–484 (2002) (Experimental, Magn. Prop., 15) MSIT®
DOI: 10.1007/978-3-540-78644-3_11 # Springer 2008
6 [2002Yuk]
[2003Fuk]
[2005Fuk]
[Mas2] [V-C2]
Fe–H–V Yukawa, H., Teshima, H., Yamashita, D., Ito, S., Yamaguchi, S., Morinaga, M., “Alloying Effects on the Hydriding Properties of V at Low Hydrogen Pressures”, J. Alloys Compd., 337, 264–268 (2002) (Phase Relations, Thermodyn., Experimental, 7) Fukai, Y., Mori, K., Shinomiya, H., “The Phase Diagram and Superabundant Vacancy Formation in Fe-H Alloys under High Hydrogen Pressures”, J. Alloy Compd., 348, 105–109 (2003) (Phase Diagram, Phase Relations, Experimental, 42) Fukai, Y., “The Structure and Phase Diagrams of M-H Systems at High Chemical Potential, High Pressure and Electrochemical Synthesis”, J. Alloys Compd., 404/406, 7–15 (2005) (Phase Diagrams, Phase Relations, Thermodyn., Review, 40) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_11 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–La–Si
1
Iron – Lanthanum – Silicon Gabriele Cacciamani
Introduction Phase equilibria in the Fe-La-Si system are not well known: one of the boundary binary phase diagrams (La-Si) is not reported in literature and, in the ternary, only one isothermal section [1972Bod] and one sub-solidus section [1996Lia] have been experimentally investigated. A summary of the experimental works carried out on crystal structures, phase equilibria and thermodynamics is reported in Table 1. In the last ten years, attention of researchers has been mainly attracted by the peculiar properties of the La (Fe1–xSix)13 ternary phase which shows large magnetic moment, high magnetocaloric effect and giant isotropic magnetostriction. Binary Systems The Fe-La and Fe-Si binary systems are accepted from [Mas2]. La-Si phase diagram is not known, but six La-Si intermediate phases are known in literature. Solid Phases Crystal structure data of the Fe-La-Si phases are summarized in Table 2. Six ternary phases have been identified in the system. Four of them (τ1, τ2, τ5 and τ6) show appreciable solution ranges at constant La composition. The relation between τ5 and τ6 is still not clear. In previous studies one phase was found with NaZn13 cubic structure [1968Kry] and/or a pseudo-cubic variant with tetragonal symmetry [1969Bod]. More recently a tetragonal structure has been refined and two separate composition ranges have been determined by [1994Tan] but temperature ranges of stability of the two phases and the transition mechanism between them have not been clearly identified La(FexSi1–x)2 has been cited by [1985Hul] while discussing structural properties of BaAl4 derivative Mooser-Pearson phases. Isothermal Sections The Fe-La-Si isothermal section at 600°C is reported in Fig. 1, based on the experimental results by [1972Bod] (also reported by [1992Rag]) modified in order to be consistent with the accepted binary systems. The sub-solidus section obtained by [1996Lia] is shown in Fig. 2. It was determined by analyzing samples homogenized at 900°C and then cooled down into the furnace. According to [2001Rag] it may be supposed that the equilibria reported are those stable at about room temperature. Thermodynamics The enthalpy of mixing of ternary Fe-La-Si alloys has been measured at 1645°C by calorimetry at Fe/Si equiatomic ratio, between 0 and 45 at.% La by [1983Esi] and at Fe/Si 1/3 ratio by [1982Esi]. The minimum enthalpy of mixing (–22.7 kJ·mol–1) was found at Fe/Si equiatomic ratio and 30 at.% La. Notes on Materials Properties and Applications La(FexSi1–x)2 is a Pauli paramagnet [1974Nar] with a temperature-independent susceptibility of about 3·10–6 emu·g–1. LaFe2Si2 is also a Pauli paramagnet. Low temperature magnetic susceptibility has been measured by [1975Fel] while effective magnetic moment has been measured by [2000Ger]. Low temperature specific heat has been measured by [2003Svo, 2005Svo].
Landolt-Börnstein New Series IV/11D4
MSIT®
DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
2
Fe–La–Si
The by far most studied Fe-La-Si phase is La(Fe1–xSix)13 (τ6). It shows important magnetocaloric effect (MCE), a Curie temperature of about 242 K at x = 0.185 [2000Zha] and it is magnetically soft enough to be a magnetic refrigerant. The magnetic entropy change was studied by several authors: according to [2000Zha], at x = 0.185, TC = 242 K and ΔS = –3.2 J·(kg·K)–1 upon a 2 T magnetic field; according to [2001Hu], at x = 0.123, TC = 208 K and ΔS reaches –19.2 J·(kg·K)–1 under a field of 5 T, which exceeds that of most other materials with a reversible magnetic transition in the same temperature range. Magnetic entropy change at various compositions has been also investigated by [2002Wen]. Thermal expansion coefficient has been measured at x = 0.1 in the 85 to 413 K temperature range by [2003Cha1]: they found a variation from 3.6·10–6 to 1.2·10–5 K–1 between 143 and 233 K and no structural transformation in the same temperature range. In [2003Cha2] a theoretical discussion of the stability of the R(TX)13 NaZn13 type phases is presented while in [2006Fuj2] the entropy change mechanism is discussed. The interplay between structural and magnetic properties has been studied by [2003Wan]. Giant magnetovolume (MVE) and MCE effects in hydrogenated La(Fe1–xSix)13 have been investigated by [2002Fuj1, 2002Fuj2, 2002Fuj3, 2003Fuj, 2004Fuj2, 2005Fuj, 2005Man, 2006Fuk1, 2006Fuk2, 2006Med]. MVE has also been investigated by X-ray diffraction in high magnetic field by [2004Fuj1]. Mössbauer spectroscopy was used by [2004Ham] in order to investigate the site preference of Fe/Si substitution and explain its effect on magnetic properties (Si substitution decreases magnetization while rising the Curie temperature). Structure and magnetic properties in melt-spun ribbons have been investigated by [2005Yan]. The effect of Fe partial substitution by Co has been considered by [2003Liu, 2005Ily] while partial substitution of La by Ce has been considered by [2006Fuj1].
Table 1.
Investigations of the Fe-La-Si Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1967Ram]
XRD
La2FeSi3
[1968Kry]
XRD
cubic La(Fe1–xSix)13
[1969Bod]
XRD
pseudo-cubic La(Fe1–xSix)13
[1969May]
XRD, Mössbauer spectroscopy
La(FexSi1–x)2
[1970Bod]
XRD
LaFeSi
[1972Bod]
XRD, LOM
isothermal section at 600°C
[1978Ros]
XRD
LaFe2Si2
[1982Esi]
Calorimetry
1640°C, liquid at Fe/Si 3/1, up to 45 at.% La
[1983Esi]
Calorimetry
1645°C, liquid at Fe/Si 1/1, up to 45 at.% La
[1994Tan]
XRD
tetragonal La(Fe1–xSix)13
[1996Lia]
XRD
sub-solidus section
[1997Mor]
XRD, DTA
La(FexSi1–x)2
[1998Wel]
XRD
LaFeSi
DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–La–Si Table 2.
3
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
(δFe) 1538–1394
cI2 Im 3m W
a = 293.15
(γFe) 1394–912
cF4 Fm 3m Cu
a = 364.67 a = 357.3
(αFe) < 912
cI2 Im 3m W
Comments/References
connected to αFe in Fe-Si, dissolves up to 19.5 at.% Si at 1390°C [Mas2] dissolves up to 3.8 at.% Si [Mas2] at 915°C [Mas2]
a = 286.65 a = 286.0
connected to δFe in Fe-Si, dissolves up to 19.5 at.% Si at 1390°C [Mas2] at 10 at.% Si [V-C2]
(β’La)
cF4 Fm 3m Cu
a = 517
at 25°C, 2.0 GPa [Mas2]
(γLa) 918–865
cI2 Im 3m W
a = 426
[Mas2]
(βLa) 865–310
cF4 Fm 3m Cu
a = 530.3
[Mas2]
(αLa) < 310
hP4 P63/mmc αLa
a = 377.40 c = 1217.1
at 25°C [Mas2]
(δSi)
hP4 P63/mmc αLa
a = 380 c = 628
at 25°C, 16 GPa → 1 atm [Mas2]
(γSi)
cI16 Im 3m γSi
a = 663.6
at 25°C, 16 GPa [Mas2]
(βSi)
tI4 I41/amd βSn
a = 468.6 c = 258.5
at 25°C, 9.5 GPa [Mas2]
(αSi) < 1414
cF8 Fd 3m C (diamond)
a = 543.06
at 25°C [Mas2]
α2, Fe-Si ≲ 1280
cP2 Pm 3m CsCl
α1, Fe-Si ≲ 1240
cF16 Fm 3m BiF3
10 to 22 at.% Si
a = 565
10 to 30 at.% Si [V-C2] (continued)
Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
4
Fe–La–Si
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
Fe2Si ∼1212–∼1040
hP6 P 3m1
a = 405.2 c = 508.5
[V-C2]
Fe5Si3 1060–825
hP16 P63/mcm Mn5Si3
a = 675.6 c = 471.8
[V-C2]
FeSi < 1410
cP8 P213 FeSi
a = 451.7
[V-C2]
βFeSi2 (h) 1220–937
tP3 P4/mmm βFeSi2
a = 268.4 c = 512.8
69.5 to 73 at.% Si [V-C2]
αFeSi2 (r) < 982
oC48 Cmca αFeSi2
a = 986.3 b = 779.1 c = 783.3
[V-C2]
LaSi2 (h)
tI12 I41/amd Th2Si
a = 431.8 c = 1384
[V-C2]
LaSi2 (r)
oI12 Imma Gd2Si3
a = 427 b = 419 c = 1394
[V-C2]
LaSi
oP8 Pnma FeB
a = 840.4 ± 0.5
[V-C2]
La5Si4
tP36 P41212 Zr5Si4
a = 804 c = 1543
[V-C2]
La3Si2
tP10 P4/mbm U3Si2
a = 788.5 c = 443.4
[V-C2]
La5Si3
tI32 I4/mcm Cr5B3
a = 794.9 c = 1407
at La rich boundary
a = 795.0 c = 1404
at Si rich boundary [V-C2]
a = 406.2 ± 0.3 c = 7117.9 ± 0.5
[1970Bod]
a = 409.8 c = 713.3
(refinement) [1998Wel]
* τ1, LaFeSi
tP6 P4/nmm Cu2Sb CeFeSi
b = 401.0 ± 0.3 c = 605.9 ± 0.3
(continued)
DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–La–Si
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
* τ2, La(FexSi1–x)2 (also denoted as La5Fe2Si8 or La2FeSi3)
hP3 P6/mmm AlB2
Lattice Parameters [pm]
a = 406.9 c = 410.1 a = 409.7 c = 433.1 a = 409.2 c = 437.0
5
Comments/References
x = 0.15–0.25 at x = 0.25 [1967Ram] at x = 0.15 [1967Ram] at x = 0.2 [1969May]
* τ3, LaFeSi2
oC16 Cmcm CeNiSi2
a = 418.4 ± 0.3 b = 1724 ± 1 c = 405.8 ± 0.3
[V-C2]
* τ4, LaFe2Si2 < 1547
tI10 I4/mmm BaAl4 (ThCu2Si2) CeGa2Al2
a = 405.3 c = 1015.3
[1978Ros]
a = 404.9 c = 1014.7
[1997Mor]
tI56 I4/mcm
a = 799.4 to 795.4 c = 1166.3 to 1172.0
at x = 0.25–0.38 [1996Lia]
a = 793.2 c = 1167.7
at x = 0.31, after annealing at 900°C (refinement) [1994Tan]
a = 1148.7 to 1143.9 a = 1148.62 a = 1154.36
at x = 0.11–0.20 [1996Lia]
* τ5, La(Fe1–xSix)13(h?)
LaFe9Si4 (NaZn13 derivative) * τ6, La(Fe1–xSix)13(r)
cF112 Fm 3c NaZn13
a = 1147.86 a = 1152.89 a = 1143.4
Landolt-Börnstein New Series IV/11D4
MSIT®
at x = 0.19, T = 298 K (refinement) at x = 0.19, T = 85 K (refinement) sharp decrease of the lattice parameter at about 190 K [2003Cha2] at x = 0.11 and T = 193 K at x = 0.11 and T = 185 K (refinement) [2003Wan] at x = 0.19 [2004Ham]
DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
6
Fig. 1. Fe-La-Si.
Fe–La–Si
Isothermal section at 600°C
DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–La–Si
Fig. 2. Fe-La-Si.
Landolt-Börnstein New Series IV/11D4
7
Sub-solidus section
MSIT®
DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
8 References [1967Ram] [1968Kry]
[1969Bod]
[1969May] [1970Bod]
[1972Bod]
[1974Nar] [1975Fel]
[1978Ros]
[1982Esi]
[1983Esi]
[1985Hul] [1992Rag]
[1994Tan] [1996Lia]
[1997Mor]
[1998Wel]
Fe–La–Si
Raman, A., “Structural Data of Some Rare-Earth Alloys.”, Naturwissenschaften, 54, 560 (1967) (Crys. Structure, Experimental, 0) Krypyakevych, P.I., Zarechniuk, O.S., Gladyshevskiy, E.I., Bodak, O.I., “Ternary Compounds of the NaZn13-Type” (in German), Z. Anorg. Chem., 358, 90–96 (1968) (Crys. Structure, Experimental, Phase Relations, 26) Bodak, O.I., Gladyshevskiy, E.I., “Ternary Compounds of the NaZn13 and Related Types in the (La, Ce, Pr, Nd, Sm, Eu, Gd)-Ni-Si and (La, Ce)-(Fe, Co)-Si Systems”, Dop. Akad. Nauk Ukrain. RSR, Ser. A, Fiz-Mat. Tekh. Nauki, (12), 1125–1129 (1969) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 3) Mayer, I., Tassa, M., “Rare-Earth-Iron(Cobalt, Nickel)-Silicon Compounds”, J. LessCommon Met., 19, 173–177 (1969) (Crys. Structure, Experimental, 9) Bodak, O.I., Gladyshevskiy, E.I., Krypyakevych, P.I., “Crystal Structure of Ce-Fe Silicide and Similar Compounds” (in Russian), Zhur. Strukt. Kim., 11(2), 305–310 (1970) (Crys. Structure, Electronic Structure, Experimental, 20) Bodak, O.I., Gladyshevskiy, E.I., “The Lanthanum-Iron-Silicon System”, Visn. Lviv. Derzh. Univ., Ser. Khim., 14, 29–34 (1972) (Phase Diagram, Phase Relations, Crys. Structure, Experimental, 12) Narasimhan, K.S.V.L., Steinfink, H., “Magnetic Investigations on AlB2 Type Structures”, J. Solid State Chem., 10, 137–141 (1974) (Crys. Structure, Experimental, Magn. Prop., 10) Felner, I., Mayer, I., Grill, A., Schieber, M., “Magnetic Ordering in Rare-earth Fe Silicides and Germanides of the RFe2X2 Type”, Solid State Commun., 16, 1005–1009 (1975) (Crys. Structure, Experimental, Magn. Prop., 17) Rossi, D., Marazza, R., Ferro, R., “Lattice Parameters of Some ThCu2Si2-Type Phases in Ternary Alloys of Rare Earths with Cobalt (or Iron) and Silicon (or Germanium)”, J. Less-Common Met., 58(2), 203–207 (1978) (Crys. Structure, Experimental, 10) Esin, Yu.O., Ermakov, A.F., Valishev, M.G., Gel`d, P.V., Levin, E.S., “Enthalpies of Formation of the Liquid Alloys of the System Fe3Si with Yttrium and Lanthanum” (in Russian), Zh. Fiz. Khim., 56(10), 2619–2620 (1982) (Experimental, Thermodyn., 5) Esin, Yu.O., Ermakov, A.F., Valishev, M.G., Geld, P.V., Levin, E.S., “Enthalpy of Formation of Liquid Alloys of Iron Monosilicide with Yttrium and Lanthanum”, Russ. Metall., (3), 49–50 (1983) (Experimental, Thermodyn., 5) Hullinger, F., “Mooser-Pearson Phases Among the BaAl4-Type Derivations”, Helv. Phys. Acta, 58, 216–225 (1985) (Crys. Structure, Review, 47) Raghavan, V., “The Fe-La-Si (Iron-Lanthanum-Silicon) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Metals, Calcutta, 6B, 936–939 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Review, 10) Tang, W.-H., Liang, X.-K., Chen, X.-L., Rao, G.-H., “Structure of LaFe9Si4 Intermetallic Compound”, J. Appl. Phys., 76(7), 4095–4098 (1994) (Crys. Structure, Experimental, 13) Liang, J., Rao, G., Tang, W., Zhao, Y., Guo, Y., Yan, X., Zhang, Y., Cheng, X., Xie, S., Chen, X., Yu, Y., Tian, J., Liang, J., “Phase Relations of R-T-M Ternary System, Crystal Structure and Magnetic Properties of R(T, M) Compounds (R = Light Rare Earth; T = Co, Fe; M = Al, Si) (I)”, Prog. Nat. Sci., 6(6), 641–650 (1996) (Experimental, Phase Diagram, Phase Relations, 6) Morozkin, A.V., Seropegin, Yu.D., Gribanov, A.V., Barakatova, J.M., “Analysis of the Melting Temperatures of RT2 Compounds (MgCu2 structure) (R = Rare Earth, T = Mn, Fe, Co, Ni, Ru, Rh, Pd, Os, Ir, Pt) and RT2X2 Compounds (R = La, Ce, Sm, Er; T = Mn, Fe, Co, Ni, Cu, Ru, Rh, Pd, Pt; X = Si, Ge)”, J. Alloys Compd., 256, 175–191 (1997) (Crys. Structure, Experimental, Thermodyn., 74) Welter, R., Ijjaali, I., Venturini, G., Malaman, B., “X-Ray Single Crystal Refinement on Some CeFeSi Type RTX Compounds (R = RE Elements; T = Mn, Fe, Co, Ru; X = Si, Ge). Evolution of the Chemical Bonds”, J. Alloys Compd., 265, 196–200 (1998) (Crys. Structure, Experimental, 16)
DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–La–Si [2000Ger]
[2000Zha]
[2001Hu]
[2001Rag] [2002Fuj1]
[2002Fuj2]
[2002Fuj3]
[2002Wen]
[2003Cha1]
[2003Cha2]
[2003Fuj]
[2003Liu]
[2003Svo]
[2003Wan]
[2004Fuj1]
[2004Fuj2]
Landolt-Börnstein New Series IV/11D4
9
Gerasimov, E.G., Gavilko, V.S., Zaikov, N.K., Podgornykh, S.M., Korolev, A.V., Pirogov, A.N., “Magnetic State of Fe Atoms in La(Fe1–xVx)2Si2 (0 ≤ x ≤ 0.5) Compounds”, Phys. Met. Metallogr., 90(3), 231–238 (2000) (Crys. Structure, Experimental, Magn. Prop., 7) Zhang, X.X., Wen, G.H., Wang, F.W., Wang, W.H., Yu, C.H., Wu, G.H., “Magnetic Entropy Change in Fe-Based Compound LaFe10.6Si2.4”, Appl. Phys. Lett., 77(19), 3072–3074 (2000) (Crys. Structure, Experimental, Magn. Prop., 17) Hu, F.-X., Shen, B.-G., Sun, J.-R., Cheng, Z.-H., Rao, G.-H., Zhang, X.-X., “Influence of Negative Lattice Expansion and Metamagnetic Transition on Magnetic Entropy Change in the Compound LaFe11.4Si1.6”, Appl. Phys. Lett., 78(23), 3675–3677 (2001) (Crys. Structure, Experimental, Magn. Prop., 22) Raghavan, V., “Fe-La-Si (Iron - Lanthanum - Silicon)”, J. Phase Equilib., 22(2), 158–159 (2001) (Crys. Structure, Phase Relations, Review, 4) Fujita, A., Fujieda, S., Fukamichi, K., “Giant Magnetovolume and Magnetocaloric Effects in Itinerant-Electron Metamagnetic La(FexSi1–x)13 Compounds” (in Japanese), Materia Japan, 41, 269–275 (2002) (Phys. Prop., 47) Fujita, A., Fujieda, S., Fukamichi, K., Mitamura, H., Goto, T., “Itinerant-Electron Metamagnetic Transition and Large Magnetovolume Effects in La(FexSi1–x)13 Compounds”, Phys. Rev. B: Condens. Matter, 65(1), 014410–1–6 (2002) (Experimental, Magn. Prop., Phase Relations, 31) Fujieda, S., Fujita, A., Fukamichi, K., “Large Magnetocaloric Effect in La(FexSi1–x)13 Itinerant-Electron Metamagnetic Compounds”, Appl. Phys. Lett., 81(7), 1276–1278 (2002) (Experimental, Magn. Prop., Thermodyn., 17) Wen, G.H., Zheng, R.K., Zhang, X.X., Wang, W.H., Chen, J.L., Wu, G.H., “Magnetic Entropy Change in LaFe13–xSix Intermetallic Compounds”, J. Appl. Phys., 91(10), 8537–8539 (2002) (Crys. Structure, Experimental, Phys. Prop., 14) Chang, H., Liang, J.-K., Shen, B.-G., Yang, L.-T., Wang, F., Chen, N.-X., Rao, G.-H., “Thermal Expansion Coefficients of LaFe11.7Si1.3”, J. Phys. D: Appl. Phys., 36(2), 160–163 (2003) (Crys. Structure, Experimental, Magn. Prop., Phys. Prop., 11) Chang, H., Chen, N., Liang, J., Rao, G., “Theoretical Study of Phase Forming of NaZn13 Type Rare-Earth Intermetallics”, J. Phys.: Condens. Matter, 15, 109–120 (2003) (Calculation, Crys. Structure, Phase Relations, 28) Fujita, A., Fujieda, S., Hasegawa, Y., Fukamichi, K., “Itinerant-Electron Metamagnetic Transition and Large Magnetocaloric Effects in La(FexSi1–x)13 Compounds and Their Hydrides”, Phys. Rev. B: Condens. Matter, 67(10), 104416–1–12 (2003) (Experimental, Magn. Prop., Phase Relations, Thermodyn., 38) Liu, X.B., Altounian, Z., “Effect of Co Content on Magnetic Entropy Change and Structure of La(Fe1–xCox)11.4Si1.6”, J. Magn. Magn. Mater., 264, 209–213 (2003) (Crys. Structure, Experimental, Magn. Prop., Thermodyn., 12) Svoboda, P., Vejpravova, J., Honda, F., Santava, E., Schneeweiss, O., Komatsubara, T., “The Analisis of the Specific Heat of RFe2Si2 Compounds”, Physica B, 328(1–2), 139–141 (2003) (Crys. Structure, Experimental, Thermodyn., 6) Wang, F., Wang, G.-J., Hu, F., Kurbakov, A., Shen, B., Cheng, Zh., “Strong Interplay Between Structure and Magnetism in the Giant Magmetocaloric Intermetallic Compound LaFe11.4Si1.6: a Neutron Diffraction Study”, J. Phys.: Condens. Matter, 15(30), 5269–5278 (2003) (Crys. Structure, Experimental, Phys. Prop., 24) Fujita, A., Fukamichi, K., Koyama, K., Watanabe, K., “X-ray Diffraction Study in High Magnetic Fields of Magnetovolume Effect in Itinerant-Electron Metamagnetic La (Fe0.88Si0.12)13 Compound”, J. Appl. Phys., 95(11), 6687–6689 (2004) (Crys. Structure, Experimental, Magn. Prop., Phase Relations, Phys. Prop., 20) Fujieda, S., Hasegawa, Y., Fujita, A., Fukamichi, K., “Thermal Transport Properties of Magnetic Refrigerants La(FexSi1–x)13 and Their Hydrides, and Gd5Si2Ge2 and MnAs”, J. Appl. Phys., 95(5), 2429–2431 (2004) (Crys. Structure, Experimental, Phase Relations, Transport Phenomena, 20)
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10 [2004Ham]
[2005Fuj]
[2005Ily]
[2005Man]
[2005Svo]
[2005Yan]
[2006Fuj1]
[2006Fuj2]
[2006Fuk1]
[2006Fuk2]
[2006Med]
[Mas2] [V-C2]
Fe–La–Si Hamdeh, H.H., Al-Ghanem, H., Hikal, W.M., Taher, S.M., Ho, J.C., Anh, D.T.K., Thuy, N.P., Duc, N.H., Thang, P.D., “Mössbauer Spectroscopic Evaluation of Chemical and Electronic Distributions in La(Fe0.81Si0.19)13”, J. Magn. Magn. Mater., 269(3), 404–409 (2004) (Crys. Structure, Experimental, Magn. Prop., 17) Fujita, A., Fukamichi, K., “Control of Large Magnetocaloric Effects in Metamagnetic La (FexSi1–x)13 Compounds by Hydrogenation”, J. Alloys Compd., 404–406, 554–558 (2005) (Crys. Structure, Experimental, Interface Phenomena, Magn. Prop., 30) Ilyn, M., Tishin, A.M., Hu, F.X., Gao, J., Sun, J.R., Shen, B.G., “Magnetocaloric Properties of the LaFe11.7Si1.3 and LaFe11.2Co0.7Si1.1 Systems”, J. Magn. Magn. Mater., 290–291(1), 712–714 (2005) (Experimental, Magn. Prop., Phase Relations, Thermodyn., 12) Mandal, K., Gutfleisch, O., Yan, A., Handstein, A., Mueller, K.-H., “Effect of Reactive Milling in Hydrogen on the Magnetic and Magnetocaloric Properties of LaFe11.57Si1.43”, J. Magn. Magn. Mater., 290–291(1), 673–675 (2005) (Experimental, Magn. Prop., Phase Relations, Thermodyn., 8) Svoboda, P., Mihalik, M., Mihalik, M., Vejpravova, J., Rusz, J., “Specific Heat Analysis of Heavy REFe2Si2 Compounds”, J. Magn. Magn. Mater., 290–291(1), 609–611 (2005) (Crys. Structure, Experimental, Magn. Prop., Thermodyn., 6) Yan, A., Mueller, K.-H., Gutfleisch, O., “Structure and Magnetic Entropy Change of MeltSpun LaFe11.57Si1.43 Ribbons”, J. Appl. Phys., 97(3), 036102–1–3 (2005) (Crys. Structure, Experimental, Magn. Prop., Thermodyn., 17) Fujieda, S., Fujita, A., Fukamichi, K., Hirano, N., Nagaya, S., “Large Magnetocaloric Effects Enhanced by Partial Substitution of Ce for La in La(Fe0.88Si0.12)13 Compound”, J. Alloys Compd., 408–412, 1165–1168 (2006) (Crys. Structure, Experimental, Magn. Prop., 18) Fujita, A., Fukamichi, K., “Enhancement of Isothermal Entropy Change Due to Spin Fluctuations in Itinerant-Electron Metamagnetic La(Fe0.88Si0.12)13 Compound”, J. Alloys Compd., 408–412, 62–65 (2006) (Experimental, Magn. Prop., Thermodyn., 21) Fukamichi, K., Fujita, A., Fujieda, S., “Large Magnetocaloric Effects and Thermal Transport Properties of La(FeSi)13 and Their Hydrides”, J. Alloys Compd., 408–412, 307–312 (2006) (Crys. Structure, Electr.Prop., Experimental, Magn. Prop., Thermodyn., 30) Fukamichi, K., Fujita, A., Fujieda, S., “Application of Large Magnetocaloric Effects in Itinerant-Electron Metamagnets to Cooling Systems”, Mat. Sci. Forum, 512, 137–144 (2006) (Experimental, Magn. Prop., Thermodyn., 20) de Medeiros, L.G. Jr., de Oliveira, N.A., “Magnetocaloric Effect in La(FexSi1–x)13 Doped with Hydrogen and Under External Pressure”, J. Alloys Compd., 424(1–2), 41–45 (2006) (Calculation, Crys. Structure, Magn. Prop., Thermodyn., 18) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_12 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mg–O
1
Iron – Magnesium – Oxygen Nathalie Lebrun, Olga Fabrichnaya
Introduction Knowledge of the phase equilibria in the Fe-Mg-O system is important in the fields of ceramics, metallurgy and petrology. For example, magnesiowüstite (Fe,Mg)O is a constituent of the lower mantle of the Earth. In geophysics, phase equilibria between magnesiowüstite and other minerals and their elastic properties at high pressures are important for the calculation of the seismic profile of the Earth. Interest in this ternary system also arises from the electrical and magnetic properties of magnesioferrite MgFe2O4 (spinel structure). In addition, the melting behavior of compositions in the MgO-FeO-Fe2O3 region provide a basis for evaluating the performance of magnesite bricks subjected to iron-oxide attack in steel making furnaces. The Fe-Mg-O ternary system has been studied extensively, in particular in the FeO-MgO-Fe2O3 region. [1931Rob] first attempted to determine the phase relationships in FeO-MgO-Fe2O3 from dissociation measurements in air at one atmosphere and below 1000°C. The temperature was measured using an optical pyrometer and the compositions were determined by chemical analysis. [1954Ric] investigated the reversible dissociation of ferric oxide in the FeO-MgO-Fe2O3 region as a function of temperature in air between 800 and 1500°C. Thermogravimetry was used to determine the loss of oxygen and the composition was checked by Xray and microscopic analysis. [1955Woo] also investigated the phase relations in this ternary subsystem in air from 700 up to 1650°C using the same experimental procedure as [1954Ric]. [1960Pal] studied the dissociation as a function of pressure (1 to 0.01 bar) and temperature (1000 to 1300°C) in the ferrite region. [1962Phi] equilibrated samples in sealed tubes, which were then quenched and the resulting phase assemblages studied by X-ray analysis and microscopic examinations. Isothermal sections between 1400 to 1800°C were constructed from the results that were then used to draw the liquidus surface. However, the oxygen partial pressure in these experiments was not known. [1965Kat] used thermogravimetry at 1160°C and various oxygen pressures and also constructed an isothermal section at 1160°C. [1965Ole] determined the isothermal section for 1400°C from oxidation measurements made at various pressures from 1 to 10–13 bar. Using X-ray analysis and microscopic investigation, they studied the concentration dependence of the lattice parameter of magnesioferrite solid solutions. [1967Spe] also used thermogravimetry to determine phase equilibria at 1300°C as function of the fugacity of oxygen. [1968Rei] studied the equilibria between the solid phases and oxygen by thermogravimetry in air and temperatures up to 1000°C. [1970Ber] established the phase relationships at 850°C under various oxygen pressures using chemical analysis and X-ray diffraction. [1980Shi] established a partial isothermal section at 1000°C by measuring the emf of galvanic cells incorporating a solid electrolyte. [1961Phi] used a quenching method to determine the phase relationships in MgO-Fe2O3 in air at temperatures between 1600 and 1800°C. This method consisted of annealing the samples in air in the furnace until equilibrium was established between the gaseous and solid phases. The samples were then quenched to room temperature. Compositions were deduced from X-ray analysis and microscopic examination. The MgO rich region of MgO-Fe2O3 system was studied by high-temperature XRD by [2003Hon], in air and temperatures between 25 and 1400°C. [1935Bow] first reported measurement of the refractive index of FeO-MgO solid solutions revealing a continuous solid solution, named magnesiowüstite. [1961Sch1] measured the solidus curve of this solid solution in air between 1520 and 1750°C. [1953Sha] determined the activities of MgO and FeO in magnesiowüstite solid solutions at 600, 900 and 1060°C and measured their lattice parameters using X-ray diffraction. A continuous variation in the lattice parameter was observed in agreement with [1970Aub, 1970Ber]. [1997Byg] used EDX to study cation diffusivities in the MgO-FeO system over the temperature range 1200–1330°C in order to understand the interdiffusion of iron and magnesium in the oxide lattice. [1961Sch2] established oxidation curves for different mixtures of Fe2O3 and MgO at 800, 900 and 1000°C from which the activities of MgO and FeO were deduced along with an isothermal section of the Fe-Mg-O system at 900°C. [1957Flo] investigated the effect of MgO on the oxidation state of iron oxide at 1400–1500°C. [1962Hah] determined the activities of FeO in magnesiowüstite solid solutions between Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
2
Fe–Mg–O
1100 and 1300°C. The oxide samples were equilibrated with pure metallic iron in atmospheres of known oxygen partial pressures, and the compositions were checked by X-ray analysis. [1962Hah] calculated the activity of MgO. [1965Tre1] calculated the activities of Fe3O4 and MgFe2O4 and γFe2O3 by a statistical method. [1965Tre2] calculated the activity of Fe3O4 by consideration of phase equilibria involving magnesiowüstite, accepting the activity of FeO in magnesiowüstite from the experimental data of [1963Gor]. [1970Aub, 1970Ber, 1972Bal] determined the activities of FeO at 850°C from oxidation data obtained from chemical analysis. [1983Dav] calculated the activities of MgO and FeO using regular and subregular thermodynamic models for the mixing of the two components. [1987Sre] studied the activity-composition relationships of the magnesiowüstite solid solution in equilibrium with metallic iron in the temperature range 777–1127°C. [1970Ber, 1980Shi] estimated the activities of Fe3O4 and MgFe2O4 in the spinel solid solution from oxidation measurements. Using thermoelectric coefficient measurements, [1984Tre] determined the cation distributions from 600 to 1300°C and calculated the activities of Fe3O4 and MgFe2O4. [2002Kat] determined activities in spinel solid solutions by an emf method at 887–1117°C. [1984Dob] calculated the activities in the Fe-Mg-O system from oxygen activity data. [1966Ole] evaluated the thermodynamic properties of solid solutions with the spinel structure, MgxFe3–xO4 (x ≤ 1), from phase equilibrium data. [1974Sah] studied the reduction of magnesiowüstite over the temperature range 1050–1350°C using CO2/H2 mixtures. [1977Tri] established a schematic O2 partial pressure-composition phase diagram for the Fe3O4-MgFe2O4 system at 1000°C consistent with [1975Zal]. [1978Lyk] determined the fugacity from FeO to Fe0.504Mg0.496O1.13 at 800, 900, 1000 and 1100°C. [1988Lue] studied the oxidation kinetics of Mg0.99Fe0.01O in air over the temperature range 873 to 1176°C. [1988Rom] studied the reversible decomposition of MgFe2O4 in an (Fe,Mg)O solid solution during cooling and heating in an inert atmosphere and in air. [1967Guz] studied the system (MgO)0.07MgFe2On (3.06 ≤ n ≤ 4.00) by the emf method using a solid electrolyte and calculated the thermodynamics of the formation of magnesium ferrite from its component oxides at 900–1100°C under standard conditions. [1958Bry] measured the equilibrium between Fe3+ and Fe2+ ions in wüstite containing various amounts of magnesium oxide at various partial pressures (10–6.71 to 10–3.1 bar) and 1400°C. Using both calorimetric and volumetric methods, [1967Alc] measured the variation in the Fe3+/Fe2+ ratio in magnesiowüstites containing 0–20 mol% of iron oxide, which had been brought to equilibrium with known oxygen pressures. The temperature range for the measurements was 900–1200°C. [1965Rig] and [1969Bla] used microscopy to study the diffusion of FeO in MgO. [1963Gor] investigated the dependence on composition of the dissociation pressure of oxygen in spinel solid solutions MgxFe3–xO4 (x ≤ 0.5) at temperatures between 800 and 1100°C and determined the free energy of formation of MgxFe3–xO4 (x ≤ 0.5) solid solutions at 860, 960 and 1060°C. [1978Cor] studied the redox behavior of Fe2+ in solid solutions with MgO by ESR and X-ray diffraction techniques. [1972Rez] measured the electrical conductivity of MgFe2O4 at 27°C. [1974Kra] measured the magnetization of a single crystal of magnesiowüstite, which was prepared by annealing MgO single crystals in magnesiowüstite powder containing (Fe)/(Fe+Mg) = 0.22 at 1400°C for 15 days in a CO/CO2 atmosphere (p(O2) = 10–9 bar). A Foner-type vibrating sample magnetometer was used to measure the magnetization. The Fe-Mg-O ternary system was reviewed by [1964Lev, 1981Rot, 1989Rag1, 2000Leb]. Using the sublattice model, subsolidus phase equilibria in the FeO-MgO-Fe2O3 system were calculated by [1998Fab]. The solidus and liquidus of the FeO-MgO system were calculated by [2000Fab] using the same model. [1993Wu] preformed a critical assessment and optimization of the thermodynamic properties and phase diagram of the FeO-MgO system using the modified quasichemical model for the liquid phase and simple polynomial excess Gibbs energy terms for magnesiowüstite. The modified quasichemical model for the liquid phase was also applied by [2004Jun1]. [2002Dec, 2004Jun1] were able to describe phase relations in the Fe-Mg-O system by modeling magnesiowüstite MgO-FeO-FeO1.5 using a simple polynomial expansion while the magnesioferrite (spinel) phase was modeled using the compound energy formalism. The thermodynamic assessment by [2004Jun1] reproduces all of the available thermodynamic and phase equilibrium data within their uncertainty limits. The Fe-Mg-O system is of interest for geological applications, and therefore, elastic and thermal properties and high temperature transformations have been studied at high temperature and pressures [2001And, 2001Dub, 2002Jac, 2002Zha, 2003Lin, 2004Kon, 2004Lev, 2005Wes, 2005Jac, 2006Kan1, 2006Kan2]. Table 1 lists the experimental and theoretical studies of the Fe-Mg-O system occurring after 1998. DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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Binary Systems The Fe-Mg binary system is accepted from [Mas2]. The phase diagram is consistent with the version presented by [1982Kub] except for some small details; the γFe/δFe transition temperature ([1982Kub] at ∼ 1392°C, [Mas2] at ∼ 1356°C) and monotectic temperature ([1982Kub] at ∼ 1515°C, [Mas2] at ∼1526°C). Extensive miscibility gaps exist both in the liquid and solid states. The estimated solubility of liquid Mg in liquid Fe under 25 bar of argon is given by [1982Kub] as log(at.% Mg) = – 11.250/T(K) + 5.31. The solid solubility of Mg in solid Fe is negligible. The solubility of Fe in liquid Mg increases from 0.0026 at.% at the eutectic temperature of 650°C to ∼ 1.6 at.% at 1600°C. The Fe-O phase diagram is accepted from [Mas2]. The thermodynamic parameters of the Fe-O system have been assessed by [1991Sun] and [2002Dec]. Both thermodynamic assessments are in very good agreement with [Mas2]. The assessment of [2002Dec] is used in the work of [2004Jun1] reproducing most of experimental data in the Fe-Mg-O ternary system. The stoichiometric composition FeO is outside the range of wüstite. In the Fe rich part, wüstite coexists with (αFe) from 570 to 912°C, with (γFe) from 912 to 1371°C, and with the liquid from 1371 to 1424°C. In the oxygen rich part, it coexists with Fe3O4 between 570 and 1429°C, which are the temperature limits of its stable range. At high pressures, the Fe rich boundary shifts towards the composition of stoichiometric FeO and reaches a maximum value of x = 0.98 at 10 GPa, while with a further increase in pressure, x slightly decreases to a value of 0.96 at 20 GPa. At higher pressures, the wüstite composition remains practically constant [1984McC]. The Mg-O phase diagram was accepted from [1993Hal]. Periclase MgO is practically stoichiometric, and a very small amount of oxygen can be dissolved in solid and liquid Mg. Solid Phases FexO wüstite has a defect structure in which the ratio Fe/O deviates from simple stoichiometry [1965Kat]. For this reason, FeO is indicated in quotation marks throughout this assessment. The wüstite and MgO periclase form a continuous series of solid solutions (magnesiowüstite) [1931Rob, 1935Bow, 1945Jay, 1946Pet]. [1970Aub, 1970Ber, 1980Sim] measured the lattice parameter of magnesiowüstite at various compositions of MgO. The lattice parameter of “FeO” varies from 428.7 pm at its upper limit of oxidation to 431.0 pm at its lower limit. The results are consistent with [1953Sha]. The eutectoid decomposition temperature of magnesiowüstite is lowered with increasing Mg content [1968Rei, 1971Gur]. The calculated temperature of dissociation of magnesiowüstite with respect to composition as presented by [2004Jun1] is shown in Fig. 1. The non-stoichiometry in magnesiowüstite decreases with increasing Mg substitution [1970Ber, 1971Gur, 1987Sre]. [1970Ber] and [1977Tri] determined the solubility range of magnesiowüstite in the ternary system at 850°C and 1000°C in terms of the ratio O/Fe and Mg/(Mg+Fe). Figures 2a and 2b show the variation of Fe+3/O ratio and log(pO2) with respect to Fe/(Mg+Fe) ratio for magnesiowüstite in equilibrium with iron as calculated by [2004Jun1]. The results are in agreement with experimental data [1987Sre, 1965Kat, 1980Sim] and [1953Sha, 1973Maj 1974Sah, 1987Sre, 1991Wis]. Figures 3a and 3b show the nonstoichiometry of magnesiowüstite by the variation in Fe+3 content with respect to log(pO2) for different temperatures and Mg/(Mg+Fe) ratios as presented by [2004Jun1]. The calculations are in a good agreement with the experimental results of [1958Bry], [1965Kat], [1967Alc], [1967Spe], [1975Val1], [1970Lyk] and [1987Sre]. [1980Sim] measured the lattice parameters of magnesiowüstite and concluded that deviations from Vegard’s law could be due to the presence of Fe+3 cations. All the solid phases in the MgO-FeO-Fe2O3 ternary system contain cations that are smaller than the oxygen ion. The ionic arrangement in these oxides generally consists of hexagonal or cubic close-packed oxygen lattices in which the cations fill the interstitial tetrahedral or octahedral holes [1968Rei]. MgFe2O4 and Fe3O4 form a continuous series of solid solutions of the spinel structure [1931Rob, 1943Whi, 1954Ric, 1955Woo, 1977Tri, 1980Shi]. At room temperature, both end members, MgFe2O4 and Fe3O4, have inverse spinel structures [1932Bar, 1983One]. The octahedral sites are occupied by Mg+2/Fe+2 and Fe3+ ions, whereas the tetrahedral sites are occupied mainly by Fe3+ ions [1968Rei]. With increasing temperature, the degree of inversion decreases for both end-members. Cationic distributions between tetrahedral and octahedral cites in spinel were investigated by [1984Tre] and [1989Nel]. Both sets of experimental results are in reasonable agreement with the calculation of [2004Jun1] for 1000°C, presented in Fig. 4. The spinel in the phase diagram is, however, not limited by the line of MgFe2O4-Fe3O4. When the Landolt-Börnstein New Series IV/11D4
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temperature is raised, the spinel phase expands, being enriched with oxygen [1961Sch2, 1963Gor]. At the Fe3O4 end, the spinel field has some extension to the oxygen rich side. This width decreases with increasing Mg content, becoming negligible at the composition MgFe2O4 [1955Woo, 1966Ole]. The MgFe2O4 spinel is practically stoichiometric [1955Woo, 1960Pal, 1968Rei, 1966Sch, 1998Kan]. The Fe2O3 hematite is considered to be a stoichiometric phase. Table 2 lists crystallographic data for the solid phases. Quasibinary Systems The solidus in the “FeO”-MgO system has been studied extensively [1931Rob, 1935Bow, 1954Ric, 1955Woo, 1961Sch1]. According to [2004Jun1], the reported liquidus data are very scattered, possibly owing to the difficulty in quenching very fluid slags. The phase diagram of the “FeO”-MgO system was calculated by [1978Kau], [1993Wu], [2000Fab] and [2004Jun1] using different models for the solid and liquid solutions. The phase diagram calculated by [2004Jun1] is presented in Fig. 5. Based on the experimental data of [1931Rob, 1954Ric, 1955Woo], the Fe2O3-MgO “quasibinary” phase diagram has been constructed by [1961Phi]. However, it is not really quasibinary; the investigation was carried out under a constant oxygen partial pressure of 0.21 atm. The bulk oxygen content in the solid phases increases with Fe/Mg ratio. Therefore, the phase diagram of the Fe2O3-MgO system will be presented in a chapter “Temperature – Composition Section”. The liquidus of the Fe3O4-MgO system was determined experimentally by [1932War] for compositions up to 60 mass% MgO. No change in the slope of liquidus line was reported by [1932War]. Therefore, the position of the peritectic formation of spinel is impossible to deduce from these data. The complete phase relationships in the Fe3O4-MgO quasibinary system are available only through the thermodynamic calculations of [1978Kau]. The phase diagram of the Fe3O4-MgO section is not presented here because the calculations of [1978Kau] are found to be in contradiction with experimental data for the rest of the Fe-Mg-O ternary system (see Temperature – Composition Section). Therefore, these calculations are not recommended in the present evaluation. Liquidus, Solidus and Solvus Surfaces [1962Phi] determined a hypothetical liquidus surface for the MgO-FeO-Fe2O3 system based only on the data taken from the isothermal sections containing liquidus isotherms at 1400, 1500, 1600, 1700 and 1800°C. Because of the uncertainties in this liquidus surface, it has not been reproduced in this assessment. No four phase invariant reaction has been found. The univariant liquidus line starting from the peritectic reaction at 1711°C (in air) on the Fe2O3-MgO vertical section ends at the peritectic reaction of 1424°C of the Fe-O binary system. A three-phase equilibrium Lm + Lo + magnesiowüstite (where Lo-oxide liquid. Lm - metallic liquid) starting at the critical point continues down to the solidification temperature of iron. Another three-phase equilibrium, Fe2O3 + spinel + Lo, starting presumably from the MgO-Fe2O3 edge, ends at the eutectic reaction at high oxygen pressure between Fe3O4 and Fe2O3. Isothermal Sections The calculated isothermal section for the Fe-Mg-O system at 1200°C and 1 bar total pressure is presented in Fig. 6, taken from the work of [2004Jun1]. Numerous isothermal sections have been established for the partial ternary system FeO-MgO-Fe2O3 at a variety of temperatures and oxygen partial pressures: 850°C [1970Ber, 1977Tri], 1000°C [1960Pal, 1977Tri, 1980Shi], 1100°C [1960Pal], 1160°C [1965Kat, 1967Spe, 1974Sah, 1998Fab], 1200°C [1960Pal], 1300°C [1960Pal, 1967Spe], 1400°C [1962Phi, 1965Ole, 1967Spe], 1500°C [1962Phi], 1600°C [1962Phi], 1700°C [1962Phi] and 1800°C [1962Phi]. Others have determined phase relationships at temperatures up to 1400°C [1968Rei], up to 1650°C [1955Woo] and up to 1500°C [1954Ric] from dissociation curves. The calculated isothermal sections of the MgO-FeO-Fe2O3 system along with oxygen isobars at 1160, 1300 and 1400°C [2004Jun1] are presented in Figs. 7 to 9. The calculated isobars in the magnesiowüstite phase field at 1300 and 1400°C are in good agreement with the experimental data of [1967Spe] and [1958Bry] DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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respectively. However, the slope of the two isobars (log p(O2) = –7 and –6) in the magnesiowüstite + spinel two-phase region determined by [1967Spe] is steeper than that calculated by [2004Jun1]. Figures 10 to 13 show phase diagrams calculated over the temperature range of 1500 to 1800°C. Experimentally determined phase assemblages given by [1962Phi] are in agreement with the calculations of [2004Jun1] at all studied temperatures. [1988Sch] equilibrated molten MgO-FeO-Fe2O3 slags and gas mixtures of known oxygen potentials in MgO crucibles at 1600°C. Also, the calculations of oxygen partial pressure by [2004Jun1] are in agreement with the experimental data, while the calculated MgO content in the liquid is less than obtained experimentally. Temperature – Composition Sections The calculated phase diagram of the Fe2O3-MgO system given by [2004Jun1] shown in Fig. 14 is in agreement with the experimental data of [1931Rob, 1954Ric, 1955Woo, 1961Phi, 1968Rei, 2003Hon]. However, no experimental data on phase relations are available for temperatures above 1800°C. The only intermediate compound encountered is the spinel phase MgFe2O4, which melts at 1711°C (in a good agreement with experimental value of 1713 ± 5°C given by [1961Phi]) through a peritectic reaction between liquid and magnesiowüstite. At this isobaric invariant, the liquid, magnesioferrite and the magnesiowüstite phases contain approximately 68, 60, and 38 mol% iron oxide (calculated as Fe2O3), respectively. Based on their dissociation data for temperatures up to 1500°C, [1954Ric] found, that the extension of the magnesiowüstite homogeneity range towards the spinel phase boundary is very restricted up to 1500°C. This is in contradiction with [1931Rob] and [1961Phi]. The solid solubility of MgFe2O4 in MgO reported by [1931Rob] was confirmed by [1955Woo]. The phase diagram of the Fe2O3-MgO system calculated by [1978Kau] is different from the experimental diagram determined by [1961Phi]. The dissociation curves for MgO-Fe2O3 mixtures in air for increasing temperatures are shown in Fig. 15, as calculated by [2004Jun1]. The calculated results are in good agreement with the experimental data of [1954Ric, 1955Woo, 1967Wil]. Mixtures containing less than 50% of MgO are reduced to spinel with increasing temperature, while mixtures with a MgO content higher than 50% are reduced to magnesiowüstite. The phase diagram of the MgO-FeO-Fe2O3 system under air is presented in Fig. 16, taken from [2004Jun1]. This diagram is another representation of that for the MgO-Fe2O3-O2 system under air. Potential Diagrams Phase diagrams showing oxygen partial pressure against metal ratio calculated by [2004Jun1] are presented in Figs. 17 to 21. The oxygen partial pressures along the phase boundaries were determined by gas equilibration [1955Woo, 1960Pal, 1962Hah, 1965Kat, 1967Spe] or by an emf technique [1973Maj, 1987Sre], except for the study by [2000Kan] where it was determined by measurement of electrical conductivity and thermoelectric power. All experimental data are reproduced in the calculation of [2004Jun1] within experimental error limits. Thermodynamics Review of the thermodynamic measurements for the end members of the magnesioferrite and magnesiowüstite solid solutions and Fe2O3 can be found in [1973Bar, 1979Sek, 1991Sun, 2004Fab, 2004Jun1]. The enthalpy of formation of MgFe2O4 at 970 K (697°C) was measured by solution calorimetry [1968Nav] and found to be consistent with values derived from the gas equilibration data of [1970Ber] measured at 850°C and the emf data of [1965Tre3] but different from that resulting from the emf study of [1967Guz] and data derived from reduction isotherms and phase analysis [1977Tri]. The activities of “FeO” and MgO in “FeO”-MgO solid solutions have been studied extensively by the gas equilibration method and the emf technique [1953Sha, 1961Sch2, 1962Hah, 1970Aub, 1970Ber, 1972Bal, 1973Abb, 1973Maj, 1977Tri, 1983Dav, 1984Dob, 1987Sre]. Deviation from ideality is observed for the “FeO”-MgO solid solution owing to the attractive interaction of Mg2+ and Fe2+ in the lattice [1977Tri]. The results of [1970Aub, 1970Ber, 1972Bal] are in good agreement with those of [1953Sha, 1961Sch2, 1962Hah, 1987Sre]. [1973Abb] and [1973Maj] derived activity values for the magnesiowüstite solid solution from emf measurements performed at 800 and 1100°C. Positive deviations from ideal behavior, Landolt-Börnstein New Series IV/11D4
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decreasing with temperature, were observed. The observed deviations from ideality at 1100°C are lower than those obtained by [1962Hah]. The latest data for the activity of FeO, MgO and Fe2/3O from [1987Sre] are presented in Fig. 22. These data are in good agreement with other investigations, and the presence of Fe+3 is taken into account. Solid solutions between MgFe2O4 and Fe3O4 are not ideal, although the deviation from ideality is comparatively small [1966Ole, 1980Shi, 1984Dob]. This small positive departure from ideal behavior was indicated by [1963Gor] who calculated the Gibbs energy of formation of magnesioferrite - magnetite solid solutions of MgxFe3–xO4 with 0 < x ≤ 0.5. [1980Shi] determined the activities of Fe3O4 and MgFe2O4 via a galvanic cell. The results were confirmed by [1984Tre] who calculated the activity of Fe3O4 at 1000°C from a consideration of cation distribution. However, these results disagree with those of [1965Tre2], who found an almost ideal solution since the activities were considered to be independent if the phase was in equilibrium with metallic iron or with the spinel phase. From the studies undertaken by [1970Ber, 1977Tri], larger positive deviations from ideality were found, most probably owing to uncertainties in the position of the homogeneity boundaries of solid solutions of MgO in wüstite. The emf measurements of [2002Kat] indicated positive deviation from ideal behavior in spinel solid solutions at 1000°C. The values obtained lie approximately between the results of [1980Shi] and [1977Tri]. The data of [2002Kat] are presented in Fig. 23. [1978Lyk] used the regular ionic solution model to describe the thermodynamic properties of magnesiowüstite. The equilibrium partial pressure of oxygen and the homogeneity ranges of magnesiowüstite were calculated and compared with experimental data [1970Lyk]. [1983Dav] parameterized the deviations from ideal mixing for magnesiowüstite solid solutions using regular and subregular thermodynamic mixing models. [1980Gor] calculated the interaction parameter between Mg and O in liquid iron from the Mg solubility in liquid iron and constants for the deoxidation equilibrium of Mg in liquid iron. [1997Ito] studied experimentally the deoxidation equilibrium at 1600 and 1750°C, and the relationship between the concentration of magnesium and oxygen was obtained. The experimental data are in good agreement with calculations made using a thermodynamic model incorporating first and second order interaction parameters. A modified Wagner formalism was used by [2004Jun2] to model deoxodation equilibria in steel for 15 metals including Mg. The Calphad method was used by [1993Wu] to model the liquidus and solidus lines in the MgO-FeO system using a Redlich-Kister solution model for the solid phase and a quasichemical model for the liquid. [1998Fab] applied the Calphad method in the derivation of thermodynamic parameters for the solid phases and to calculate phase diagrams for the Fe-Mg-O system. However, the calculated phase boundaries deviated from the experimental data of [1967Spe] and [1958Bry]. Liquidus and solidus lines of the FeO-MgO system were calculated by [2000Fab]. The Calphad-type assessment of [2002Dec] incorporated the compound energy formalism for the spinel phase, simple polynomial expansions for magnesiowüstite and the modified quasichemical model for liquid phase. The modeling was found to reproduce most of the available experimental data within uncertainty limits. The same models were used by [2004Jun1]. In the work of [2004Jun1], the thermodynamic parameters of the Fe-Mg-O system were reassessed and a number of phase diagrams were produced showing good agreement with available experimental data. The calculated activity for magnesiowüstite and magnesioferrite were not presented in the work of [2004Jun1]. Therefore, activity data are recommended from other sources, which are in good agreement with other available experimental data. Notes on Materials Properties and Applications The phases observed in the Fe-Mg-O system are constituent minerals of the Earth. Magnesioferrite is observed at relatively low pressures in the Earth’s crust, while magnesiowüstite is a constituent of the low mantle where pressures are above 26 GPa. At these pressures, the (Mg,Fe)2SiO4 phase with the spinel structure decomposes to a mixture of the (Mg,Fe)SiO3 phase with the perovskite structure and magnesiowüstite. In order to interpret seismological data, the elastic properties, thermal expansion and seismic wave velocities of magnesiowüstite have been studied experimentally. A review of these data is presented in [1989Duf, 1992Ita, 1993Sax, 2004Fab]. Recently, volumetric measurements of magnesiowüstite at high temperatures and pressures have been performed using energy dispersive synchrotron XRD by [2002Zha, 2005Jac, 2005Wes] to derive thermal and elastic properties. [2002Jac] used ultrasonic interferometry at high pressure and temperature to determine the elastic constants of magnesiowüstite. [2006Tom] deformed DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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magnesiowüstite with a composition of xMg/(xFe+xMg) = 0.1, 0.25 and 0.4 at room temperature and pressures up to 37, 16 and 18 GPa. Elastic constants were calculated and elastic anisotropy was observed. [2004Lev] studied the behavior of MgFe2O4 spinel by synchrotron XRD at high pressures and temperatures and derived elastic properties and thermal expansion. [1965Ole] observed an increase in the electrical resistance of spinel solid solutions with reciprocal temperature. The electrical conductivity at 27°C is 4.6·10–2 Ω–1·cm–1and it increases with time [1972Rez]. This has been attributed to a re-ordering of the cations. The electronic conduction is explained by the electronic exchange of Fe3+ and Fe2+ on octahedral sites. The magnetic susceptibility of magnesiowüstite was studied between 120 and 293 K and for total Fe contents in the range of 0–6.6 mol% FeO by [1975Val2], and at 78–460 K and up to 20 mol% of FeO by [1980Bra]. [1974Kra] measured the magnetization of a magnesiowüstite single crystal. [1972Rez] measured the electrical conductivity of MgFe2O4 at 27°C. [1999Har] measured the magnetic susceptibility of MgFe2O4. It was shown that the Curie temperature of MgFe2O4 is very sensitive to Fe+3 and Mg+2 distributions between octahedral and tetrahedral sites within the spinel structure. The magnetic properties of MgFe2O4 were also measured by [1999Cro]. [1964Gro] studied strengthening of MgO by Fe. It was demonstrated that this effect is due to the presence of Fe+3 in contrast to Fe+2. [1974Sin1, 1974Sin2] determined edge and screw dislocation velocities for Fe+3 doped MgO (90 and 150 ppm of Fe+3) as a function of temperature and stress, and estimated the activation energy for edge and screw dislocations. [1988Wol] measured the creep resistance of a magnesiowüstite single crystal at 1300–1500°C, 20–70 MPa applied stress and oxygen partial pressure in the range of 10–4-102 Pa. The thermal conductivity of MgO containing 300, 3000 and 7500 ppm Fe was measured by [1971Mor]. It was shown that the thermal conductivity is reduced with Fe content and has a minimum at 80 K for highly doped samples (1 mass% Fe). [2000Kan] measured the electrical conductivity and thermoelectric power of the Fe-Mg-O system as a function of oxygen partial pressure at 1100–1300°C and Mg/(Mg+Fe) ratio in the range 0–0.4. The compositional dependence of electrical conductivity of magnesiowüstite was studied by [1971Gur] at 800–1000°C. Experimental studies of properties are summarized in Table 3. Miscellaneous [1954Ric] investigated the influence of temperature and composition on the rate of dissolution of ferric oxide in mixtures of “FeO”, MgO and Fe2O3. It was found that the oxygen loss is completely reversible; the rates of oxidation and dissolution are relatively rapid. [1988Lue] determined the kinetics of the internal oxidation of (Mg,Fe)O solid solutions and found that a two-phase region of (Mg,Fe)O and MgFe2O4 grows inward from the original surface under oxidizing conditions. The diffusion of “FeO” into MgO occurs by the counterdiffusion of Mg2+ with Fe2+ and Fe3+ ions through a relatively rigid oxygen lattice [1965Rig]. The diffusion coefficient increases exponentially with cation vacancy concentration [1969Bla]. [1997Byg] calculated the interdiffusion coefficients at different temperatures (1200, 1250, 1300 and 1330°C). Results are in good agreement with those of [1969Bla] but are lower than values obtained by [1965Rig]. This difference has been attributed to the difference in the oxygen partial pressures which would in turn affect the vacancy concentrations in the oxide samples. Assuming that the diffusion of oxygen is negligible, the interdiffusion can be considered to occur essentially by cationic transport. [2003Yam] measured interdiffusivities in magnesiowüstite at high temperatures and pressures. Diffusion profiles were measured across the interface between MgO and (Fe0.5Mg0.5)O samples by microprobe analysis. [1954Hen] studied the decomposition kinetics of Fe2O3 in the presence of MgO at 783–1020°C. [1978Cor] studied the effect of electronic structure on the redox behavior of the guest specie Fe2+ in the magnesium oxide solid solution and showed that Fe2+ is oxidized to Fe3+ in the temperature range 500–700°C. The formation of MgFe2O4 is eventually observed. [1980Vla] used the electron paramagnetic resonance method to study the mechanism of the solid state reaction taking place between MgO and Fe2O3. [1990Bec] used Mössbauer spectroscopy to investigate the kinetics of the oxidation of magnesiowüstite to spinel. [1999Har] studied the kinetics of non-convergent cation ordering in MgFe2O4 by measuring the Curie Landolt-Börnstein New Series IV/11D4
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temperature of synthetic samples as a function of isothermal annealing time. [1999Cro] prepared spinel MgFe2O4 by self-propagating high-temperature synthesis from iron, iron(III) oxide and Mg oxide. The reaction was carried out in air, under an oxygen atmosphere and in the presence of an external magnetic field of 1.1 T. The inversion parameter of spinel was influenced by the external field. Samples were characterized by XRD, EDXA, SEM, FTIP, Mössbauer spectroscopy and vibrating sample magnetometry. [2001Dub] found that magnesiowüstite solid solutions dissociate to Fe rich and Mg rich components at T > 1000 K (727°C) and P > 80 GPa. [2001Boi] found that under lithospheric conditions, periclase could dissolve a limited amount of FeO. Mineral and synthetic samples were studied by microprobe analysis, XRD and SEM. [2003Lin] studied magnesiowüstite at 86 GPa and 1000 K (727°C) using a laser-heated diamond anvil cell combined with synchrotron XRD. It was shown that the addition of Mg to FeO stabilizes the cubic structure at very high pressures that correspond to those at the boundary between the lower mantle and the Earth’s core. [2005Jac] studied wüstite and magnesiowüstite under lower mantle pressures by synchrotron XRD. It was found that wüstite Fe0.93O transformed to a rhombohedral structure at 23 GPa while (Mg0.73Fe0.27)O remained cubic at pressures of 51 GPa. The elastic properties of both phases were determined by [2005Jac]. [2006Lin] studied the electronic transition of Fe in magnesiowüstite by synchrotron Mössbauer and XRD spectroscopy under high pressure. Contrary to [2003Lin, 2004Kon, 2005Jac], where the rhombohedral phase was found only in Fe rich magnesiowüstite, [2006Kan1] discovered the transformation of cubic (Mg0.8Fe0.2)O to the rhombohedral phase at 35 GPa and room temperature, using XRD, Mössbauer and XANES spectroscopy in situ experiments. [2006Zha] presented shock-compression data for magnesiowüstite for pressures up to 130 GPa. The volume change for (Mg0.6Fe0.4)O is consistent with a cubic to NiAs structural transition. It is possible that contradictions between the different experimental results relating to the transformation of magnesiowüstite can be explained by the fact that this transformation may occur under non-hydrostatic conditions. [2006Kan2] presented high-pressure Mössbauer and X-ray emission spectra for (Mg0.8Fe0.2)O. [2001And, 2004Lev] studied the crystal structure of MgFe2O4 at high pressure, and it was indicated that spinel MgFe2O4 transforms to the orthorhombic CaMn2O4 structure at pressures between 5 and 17.7 GPa and temperatures above 1800 K (1527°C). At lower temperatures, the spinel decomposes to an oxide mixture. [2005Ant] investigated the unit cell parameters of MgFe2O4 by synchrotron radiation at 25–982°C and normal pressure. A linear dependence was observed below and above 581°C. The discontinuity in the cell parameters at 581°C coincides with maximum cation ordering. An exothermic peak was also observed at 550°C by TG-DSC measurements. [1998Ber] studied the defect structure of divalent Mg-doped αFe2O3 by Rietveld structure refinement of XRD data. It was shown that Mg+2 occupied vacant interstitial sites and adjacent octahedral Fe+3 sites in the corundum related structure. This result was confirmed by neutron powder diffraction studies [2000Ber]. [2000Ber] also measured the magnetic properties of Mg-doped αFe2O3 and Mössbauer spectra. [1997Fon] measured optical and magneto-optical polar Kerr spectra for Fe3O4 doped by Mg+2. [2000Vas] performed a thermodynamic analysis of molecular beam epitaxial growth of MgO(s) and MgO/Fe(001) layers from flow balance at the growth surface, taking into account condensation, evaporation and growth flows. Predominance diagrams were proposed where the domains for the formation of the different phases were predicted. The geometric structure of MgO deposited on Fe(001) under ultra high vacuum by electron evaporation was studied by surface XRD by [2001Mey]. An interface layer of FeO between the substrate and the growing MgO was indicated. [2004Shi] studied the contact angle between liquid Fe and a MgO substrate at 1600°C and oxygen partial pressures in the range 5.2·10–4-1·10–9 atm. [2002Mer] studied the electron density of MgO and (Mg0.963Fe0.037)O by a maximum entropy method and multipole refinements. Single crystal XRD and Mössbauer spectroscopy were used in the investigations. [2005Gat] studied the three-dimensional structure of nanocrystalline MgFe2O4 by synchrotron XRD. The catalytic activity of iron oxide and magnesium oxide solid solutions for N2O decomposition was studied by [1975Sch, 1975Val2, 1977Cor].
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Fe–Mg–O Table 1.
9
Investigations of the Fe-Mg-O Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1999Cro]
XRD, EDXA, SEM, FTIP, Mössbauer spectroscopy
MgFe2O4, self-propagating synthesis at 950–1050°C
[2000Fab]
Calphad
Solidus/liquidus in FeO-MgO system
[2000Kan]
Electric conductivity and thermoelectric power measurements
1100–1300°C, p(O2) = 0.21–10–16 atm x (Mg)/(x(Mg)+x(Fe)) = 0–0.4
[2001And]
In situ XRD in diamond anvil cell
MgFe2O4, T < 1727°C, p < 60 GPa
[2001Boi]
Microprobe analysis, XRD, SEM-WDS
Magnesiowüstite with FeO content of 5, 23 and 8.22 mol%
[2001Dub]
Diamond anvil cell, XRD
T > 727°C, p = 40–80 GPa (Mg0.8Fe0.2)O and (Mg0.5Fe0.5)O
[2002Dec]
Calphad
Phase diagrams in the whole composition range at 800–1800°C and p(O2) = 10–20 1 atm
[2002Kat]
emf
887–1117°C, MgFe2O4-Fe3O4
[2003Hon]
High-temperature XRD
25–1400°C, air, MgO 100–80 mol%
[2003Yam]
Microprobe
Interdiffusivities in magnesiowüstite
[2004Jun1]
Calphad
Phase diagrams in the whole composition range at 800–1800°C and p(O2) = 10-20 - 1 atm
[2004Jun2]
emf
887–1117°C, MgFe2O4-Fe3O4
[2004Kon]
In situ XRD in diamond anvil cell, synchrotron radiation
(MgxFe1–x)O, x = 0–0.2, T = 0–1800°C, p < 140 GPa
[2004Lev]
Synchrotron XRD in diamond anvil cell
MgFe2O4, 25–2227°C, p < 60 GPa
[2005Jac]
Synchrotron XRD in diamond anvil cell
(Mg0.73Fe0.27)O, p < 51 GPa
[2005Wes]
Synchrotron XRD in multi-anvil press
(Mg0.64Fe0.36)O, T = 1900°C, p = 26.7 GPa
[2006Kan1]
Synchrotron XRD, X-ray absorption spectroscopy, Mössbauer spectroscopy in diamond anvil cell
(Mg0.8Fe0.2)O, p < 60 GPa, 25°C
[2006Kan2]
Mössbauer spectroscopy in diamond anvil cell
(Mg0.8Fe0.2)O, p < 100 GPa, 25°C
[2006Lin]
Synchrotron Mössbauer, X-ray emission spectroscopy in diamond anvil cell
(Mg0.75Fe0.25)O, p < 100 GPa
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10
Fe–Mg–O
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(δFe) (h2) 1538–1356
cI2 Im 3m W
a = 293.15
at 1390°C [Mas2]
(γFe) (h1) 1394–912
cF4 Fm 3m Cu
a = 364.67
at 915°C [Mas2]
(αFe) (r) < 912
cI2 Im 3m W
a = 286.65
at 25°C [Mas2]
(Mg) < 650
hP2 P63/mmc Mg
a = 320.94 c = 521.01
at 25°C [V-C2]
αFe2O3 (hematite) < 1457
hR30 R 3c Al2O3
a = 503.42 c = 1374.83
∼60 at.% O [V-C2]
γFe2O3
cF56 Fd 3m MgAl2O4
a = 834
low temperature or metastable [1989Rag2]
(Fe1–x–yMgx)O magnesiowüstite (MW) MgO (periclase) < 2827 ± 30 Fe1–yO (wüstite) 1424–570
cF8 Fm 3m NaCl
(Mg1–xFex)Fe2O4+y spinel (magnesioferrite) Fe3O4 (magnetite) < 1596
cF56 Fd 3m MgAl2O4
MgFe2O4 < 1713 ± 5
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
0≤x≤1
a = 421.12
[V-C2, Mas2, 1993Hal]
a = 435.35
at 1000°C [V-C2]
a = 841.1 a = 838.65 a = 839.46
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57.1 to 58.0 at.% O at 200°C [V-C2] [1961Phi] at 1100°C and 53.1 mol% Fe2O3 [1960Pal] at 1100°C and 49.0 mol% Fe2O3 [1960Pal]
Landolt-Börnstein New Series IV/11D4
Fe–Mg–O Table 3.
11
Investigations of the Fe-Mg-O Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1964Gro]
Microhardness measurement
Strengthening of MgO by Fe+3 doping
[1965Ole]
Electrical conductivity measurements
Electrical resistance for (MgO)0.092-MgFe2O4 and MgFe2O4
[1971Gur]
Conductometry
Electroconductivity of magnesiowüstite 800–1160°C and Mg/(Mg+Fe) = 0.05–0.5
[1972Rez]
Bridge-method
Electroconductivity of MgFe2O4 spinel at 27°C
[1974Kra]
Foner-type magnetometer
Magnetization of (Fe0.157Mg0.843)O at 77 K
[1974Sin1]
Yield stress measurements
Edge and screw dislocation velocities in MgO doped with 90 ppm Fe+3
[1974Sin2]
Yield stress measurements
Edge and screw dislocation velocities in MgO doped with 150 ppm Fe+3
[1975Val2]
Magnetic measurements
Magnetic susceptibility (Fe0.066Mg0.934)O at 120–293 K
[1980Bra]
Faraday method
Magnetic susceptibility of magnesiowüstite, up to 20 mol% FeO at 78–460 K
[1988Wol]
Uniaxial compression
Creep of (Mg,Fe)O single crystal, 70–11900 ppm Fe, at 1300–1500°C, 20–70 MPa stress, p(O2)=10–4-10–2 Pa
[1999Cro]
Magnetometer
Magnetic properties for MgFe2O4
[1999Har]
Magnetic property study
Magnetic susceptibility of MgFe2O4 at 200–400 K
[2000Kan]
DC four-probe and pulse techniques
Electric conductivity and thermal power
[2001And]
In situ XRD in diamond anvil cell
Elastic constants for MgFe2O4 spinel and high-pressure phase of CaMn2O4 structure
[2002Jac]
Ultrasonic interferometry
Elastic constants of magnesiowüstite
[2002Zha]
Energy-dispersive synchrotron XRD in large-volume high-pressure apparatus
Elastic constants (Mg0.6Fe0.4)O
[2004Lev]
Synchrotron XRD in diamond anvil cell
Elastic constants and thermal expansion of MgFe2O4
[2005Jac]
Synchrotron XRD in diamond anvil cell
Elastic properties (Mg0.73Fe0.27)O
[2005Wes]
Synchrotron XRD in multianvil press
Unit cell volume, elastic constants, thermal expansion for (Mg0.64Fe0.36)O
[2006Lon]
Deformation experiments
Mechanical properties of (Mg,Fe)O
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Fig. 1. Fe-Mg-O.
Fe–Mg–O
Calculated temperature of dissociation of magnesiowüstite into spinel and Fe
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Landolt-Börnstein New Series IV/11D4
Fe–Mg–O
Fig. 2a. Fe-Mg-O.
Landolt-Börnstein New Series IV/11D4
13
Calculated Fe+3/O vs Fe/(Fe + Mg) ratio in magnesiowüstite in equilibrium with metallic iron
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14
Fig. 2b. Fe-Mg-O.
Fe–Mg–O
Calculated oxygen partial pressure for the equilibrium between magnesiowüstite and metallic iron
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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Fe–Mg–O
15
Fig. 3a. Fe-Mg-O. Calculated nonstoichiometry of magnesiowüstite as a function of oxygen partial pressure at various temperatures and Mg/(Mg + Fe) ratios
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Fe–Mg–O
Fig. 3b. Fe-Mg-O. Calculated nonstoichiometry of magnesiowüstite as a function of oxygen partial pressure at various temperatures and Mg/(Mg + Fe) ratios
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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Fe–Mg–O
17
Fig. 4. Fe-Mg-O. Calculated cation distribution in the Fe3O4-MgFe2O4 spinel solid solution at 1000°C. (O)–octahedral sites, (T) – tetrahedral sites
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Fig. 5. Fe-Mg-O.
Fe–Mg–O
Calculated phase diagram of the MgO-FeO system in equilibrium with metallic Fe
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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Fe–Mg–O
Fig. 6. Fe-Mg-O.
Landolt-Börnstein New Series IV/11D4
19
Calculated isothermal section at 1200°C and ptot=1 bar
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Fig. 7. Fe-Mg-O.
Fe–Mg–O
Calculated FeO-MgO-Fe2O3 isothermal section at 1160°C with log(pO2) isobars
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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Fe–Mg–O
Fig. 8. Fe-Mg-O.
Landolt-Börnstein New Series IV/11D4
21
Calculated FeO-MgO-Fe2O3 isothermal section at 1300°C with log(pO2) isobars
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22
Fig. 9. Fe-Mg-O.
Fe–Mg–O
Calculated FeO-MgO-Fe2O3 isothermal section at 1400°C with log(pO2) isobars
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Fe–Mg–O
23
Fig. 10. Fe-Mg-O. Calculated FeO-MgO-Fe2O3 isothermal section at 1500°C
Landolt-Börnstein New Series IV/11D4
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Fig. 11. Fe-Mg-O.
Fe–Mg–O
Calculated FeO-MgO-Fe2O3 isothermal section at 1600°C
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Landolt-Börnstein New Series IV/11D4
Fe–Mg–O
Fig. 12. Fe-Mg-O.
Landolt-Börnstein New Series IV/11D4
25
Calculated FeO-MgO-Fe2O3 isotherm at 1700°C
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Fig. 13. Fe-Mg-O.
Fe–Mg–O
Calculated FeO-MgO-Fe2O3 isotherm at 1800°C
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Fe–Mg–O
Fig. 14. Fe-Mg-O.
Landolt-Börnstein New Series IV/11D4
27
Calculated phase diagram of the MgO-Fe2O3-O2 system in air
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Fe–Mg–O
Fig. 15. Fe-Mg-O. Calculated dissociation curves of MgO-Fe2O3 initial mixtures in air with increasing temperature
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Fe–Mg–O
Fig. 16. Fe-Mg-O.
Landolt-Börnstein New Series IV/11D4
29
Calculated dissociation curves of MgO-FeO-Fe2O3 system in air
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Fig. 17. Fe-Mg-O.
Fe–Mg–O
Calculated phase diagram at 1000°C log(pO2) vs Mg/(Mg+Fe) ratio
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Fe–Mg–O
Fig. 18. Fe-Mg-O.
Landolt-Börnstein New Series IV/11D4
31
Calculated phase diagram at 1100°C log(pO2) vs Mg/(Mg+Fe) ratio
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Fig. 19. Fe-Mg-O.
Fe–Mg–O
Calculated phase diagram at 1200°C log(pO2) vs Mg/(Mg+Fe) ratio
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Fe–Mg–O
33
Fig. 20. Fe-Mg-O. Calculated phase diagram at 1300°C log(pO2) vs Mg/(Mg+Fe) ratio
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Fig. 21. Fe-Mg-O.
Fe–Mg–O
Calculated phase diagram at 1500°C log(pO2) vs Mg/(Mg+Fe) ratio
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mg–O
Fig. 22. Fe-Mg-O. 1400°C
Landolt-Börnstein New Series IV/11D4
35
Activity-composition relations for FeO, Fe2/3O and MgO in equilibrium with metallic iron at
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Fig. 23. Fe-Mg-O.
Fe–Mg–O
Activity-composition in the Fe3O4-MgFe2O4 system in equilibrium with Fe2O3 at 1000°C
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Landolt-Börnstein New Series IV/11D4
Fe–Mg–O References [1931Rob]
[1932Bar] [1932War]
[1935Bow] [1943Whi] [1945Jay]
[1946Pet]
[1953Sha]
[1954Hen]
[1954Ric]
[1955Woo]
[1957Flo] [1958Bry]
[1960Pal]
[1961Phi]
[1961Sch1]
[1961Sch2]
[1962Hah]
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37
Roberts, H.S., Merwin, H.E., “The System MgO-FeO-Fe2O3 in Air at One Atmosphere”, Amer. J. Sci., 21(122), 145–157 (1931) (Experimental, Phase Diagram, Phase Relations, *, #, 15) Barth, T.F.W., Posnijak, E., “Spinel Structures: with and without Variate Atom Equipoints”, Z. Kristallogr., 82, 325–341 (1932) (Experimental, Structure, 10) Wartenberg, H.V., Prophet, E., “Melting Diagrams of Maximumfireproof Oxides. V. Systems with MgO” (in German), Z. Anorg. Chem., 208, 369–379 (1932) (Experimental, Phase Relations, 21) Bowen, N.L., Schairer, J.F., “The System MgO-FeO-SiO2” (in German), Amer. J. Sci., 29(170), 151–217 (1935) (Experimental, Phase Diagram, Phase Relations, *, #, 47) White, J., “The Physical Chemistry of Open-hearth Slags”, J. Iron Steel Inst., 148, 579–625 (1943) (Phase Diagram, Phase Relations, Review, 195) Jay, A.H., Andrews, K.W., “Note on Oxyde Systems Pertaining to Steel-Making Furnace Slags. FeO-MnO, FeO-MgO, CaO-MnO, MgO-MnO”, J. Iron Steel Inst., London, 152, 15–18 (1945) (Crys. Structure, Experimental, 7) Pettersson, H., “Investigations of the Solubility in the Solid State in the Binary Systems of the Oxides CaO, MnO, MgO and FeO”, Jernkontorets Ann., 130(12), 653–663 (1946) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 19) Shashkina, A.V., Gerasimov, Y.L., “Equilibrium of FeO-MgO Solid Solution with Hydrogen and Activities of Solution Components” (Russian), Zh. Fiz. Khim., 27, 399–410 (1953) (Experimental, Thermodyn., 6) Henrich, G., “Over the Influence of the Gas Atmosphere and the Compacting Pressure on Reaction in the Solid State” (in German), Z. Elektrochem., 58(3), 183–196 (1954) (Experimental, 40) Richards, R.G., White, J., “Phase Relationships of Iron-Oxide-Containing Spinels. II Relations in the Systems Fe-Cr-O, Fe-Mg-O, Fe-Al-Cr-O, Fe-Al-Cr-Mg-O”, Trans. Brit. Ceram. Soc., 53(7), 422–459 (1954) (Experimental, Phase Diagram, Phase Relations, *, #, 11) Woodhouse, D., White, J., “Phase Relationships of Iron Oxide-containing Spinels: III. Further Investigations on the System Fe-Mg-O and Fe-Mg-Cr-O, and General Relationships in the System Fe-Mg-Cr-Al-O”, Trans. Brit. Ceram. Soc., 54(6), 333–366 (1955) (Experimental, Phase Diagram, Phase Relations, *, #, 10) Flood, H., Hill, D.G., “The Redox Equilibrium in Iron Oxide Spinels and Related Systems”, Z. Elektrochem., 61(1), 18–24 (1957) (Experimental, 16) Brynestad, J., Flood, H., “The Redox Equilibrium in Wüstite and Solid Solutions of Wüstite and Magnesium Oxide”, Z. Elektrochem., 62(9), 953–958 (1958) (Experimental, Thermodyn., 14) Paladino, A.E., “Phase Equilibria in the Ferrite Region of the System FeO-MgO-Fe2O3”, J. Am. Ceram. Soc., 43(4), 183–191 (1960) (Experimental, Phase Diagram, Phase Relations, 9) Philips, B.L., Somiya, S., Muan, A., “Melting Relations of Magnesium Oxide-Iron Oxide Mixtures in Air”, J. Am. Ceram. Soc., 44(4), 167–169 (1961) (Experimental, Phase Relations, *, #, 6) Von Schenck, H., Pfaff, W., “The System FeO-MgO and Distribution Equilibria with Liquid Iron from 1520 to 1750°C”, Arch. Eisenhüttenw., 32(11), 741–751 (1961) (Experimental, Phase Relations, Thermodyn., *, #, 49) Von Schmahl, N.G., Frisch, B., Stock, G., “Equilibrium investigations at Magnesiowüstite and Magnesioferrite” (in German), Arch. Eisenhttenw., 32, 297–302 (1961) (Experimental, Thermodyn., 17) Hahn, W.C., Muan, A., “Activity Measurements in Oxide Solid Solutions: The System FeO-MgO in the Temperature Interval 1100 to 1300°C”, Trans. AIME, 224, 416–420 (1962) (Experimental, Thermodyn., 5)
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38 [1962Phi]
[1963Gor] [1964Gro] [1964Lev] [1965Kat] [1965Ole]
[1965Rig] [1965Tre1]
[1965Tre2]
[1965Tre3]
[1966Ole]
[1966Sch] [1967Alc] [1967Guz]
[1967Spe]
[1967Wil]
[1968Nav] [1968Rei] [1969Bla] [1970Aub]
Fe–Mg–O Philips, B., Muan, A., “Phase Equilibria in the System MgO-FeO-Fe2O3 in the Temperature Range 1400 to 1800°C”, J. Am. Ceram. Soc., 45(2), 588–591 (1962) (Experimental, Phase Relations, Phase Diagram, *, #, 18) Gordeev, I.V., Tret'yakov, Y.D., “Thermodynamics of Solid Solutions of Magnesium Ferrite with Magnetite”, Russ. J. Inorg. Chem., 8(8), 943–947 (1963) (Experimental, Thermodyn., 21) Groves, G.W., Fine, M.E., “Solid Solution + Precipitation Hardening in Mg-Fe-O Alloys”, J. Appl. Phys., 35(12), 3587–3593 (1964) (Experimental, Mechan. Prop., 14) Levin, E.M., Robbins, C.R., McMurdie, H.F., Phase Diagrams for Ceramists, Am. Ceram. Soc., Columbus, Ohio, 54–56 (1964) (Phase Diagram, Phase Relations, Review, 4) Katsura, T., Kimura, S., “Equilibria in the System FeO-Fe2O3-MgO at 1160°C”, Bull. Chem. Soc. Jap., 38(10), 1664–1670 (1965) (Experimental, 16) Oleinikov, N.N., Saksonov, Y.G., Tret"yakov, Y.D., “A Study of Phase Equilibria in the MgO-FeO-Fe2O3 System at 1400°C”, Inorg. Mater., 1(2), 227–233 (1965), translated from Zh. Neorg. Mater., 1(2), 246–253 (1965) (Experimental, Phase Relations, 21) Rigby, E.B., Cutler, I.B., “Interdiffusion Studies of the System FexO-MgO”, J. Am. Ceram. Soc., 48, 95–99 (1965) (Experimental, Interface Phenomena, Phase Relations, 13) Tret"yakov, Yu.D., Oleinikov, N.N., “Activity of the Components of Solid Solution with a Spinel Structure in the Iron-Magnesium-Oxygen System”, Inorg. Mater. (Engl. Trans.), 1(2), 234–236 (1965) (Thermodyn., Experimental, 8) Tret"yakov, Yu.D., “Methods for Estimating the Thermodynamic Properties of Ferrites and Solid Solutions of Ferrites with Magnetite”, Inorg. Mater., 1, 221–226 (1965) translated from Izv. Ak. Nauk SSSR, Neorg. Mater., 1, 240–245 (1965) (Experimental, Thermodyn., 16) Tret"yakov, Yu.D., Schmalzried, H., “Thermodynamics of Spinel Phase (Chromite, Ferrite, Aluminate)” Ber. Bunsen-Ges. Phys. Chem., 69, 396–402 (1965) (Experimental, Thermodyn., 50) Oleinikov, N.N., Tretyakov, Yu.D., “Thermodynamic Properties of Solid Solutions with Spinel Structure in the System Iron-Magnesium-Oxygen”, Inorg. Mater., 2, 131–135 (1966) translated from Izv. Akad. Nauk. SSSR, Neorg. Mater., 2(1), 155–160 (1966) (Experimental, Thermodyn., 7) Schmalzried, H., Tretjakow, J.D., “Irregularity in Ferrite”, Ber. Bunsenges Phys. Chem., 70, 180–188 (1966) as quoted in [2001Jun1] Alcock, C.B., Iyengar, G.N., “A Study of the Oxidation-Reduction Equilibria of Dilute Magnesiowüstite”, Proc. Brit. Ceram. Soc., 87, 219–229 (1967) (Experimental, Thermodyn., 10) Guzei, A.S., Lavrent'ev, V.I., Bulgakova, T.I., “The Study of the Thermodynamic Properties of Magnesium Ferrite by the emf Method Using a Solid Electrolyte”, Inorg. Mater., 3, 767–772 (1967) translated from Izv. Ak. Nauk SSSR, Neorg. Mater., 3(5), 860–867 (1967) (Experimental, Thermodyn., 19) Speidel, D.H., “Phase Equilibria in the System MgO-FeO-Fe2O3: The 1300°C Isothermal Section and Extrapolations to Other Temperatures”, J. Am. Ceram. Soc., 50(5), 243–248 (1967) (Experimental, Phase Diagram, Phase Relations, *, #, 12) Willshee, J.C., White, J., “An Investigation of Equilibrium Relationships in the System MgO-FeO-Fe2O3 up to 1750°C in Air”, Trans. Brit. Ceram. Soc., 66, 541–555 (1967) as quoted in [2001Jun1] Navrotsky, A., Kleppa, O.J., “Thermodynamics of Formation of Simple Spinels”, J. Inorg. Nucl. Chem., 30, 479–498 (1968) (Experimental, Thermodyn., 30) Reijen, P., “Phase Equilibria in the System MgO-FeO-Fe2O3”, Phillips Res. Rep., 23, 151–188 (1968) (Experimental, Phase Relations, Thermodyn., 34) Blank, S.L., Pask, J.A., “Diffusion of Iron and Nickel in Magnesium Oxide Single Crystals”, J. Am. Ceram. Soc., 52, 669–675 (1969) (Experimental, 33) Aubry, J., Berthet, A., Duchene, R., Etienne, H., Evrard, O., Jeannot, F., Gleitzer, C., Offroy, C., Perrot, P., “Stabilization of Iron Monoxide by Formation of Solid Solution” (in French), Ann. Chim., 5, 299–308 (1970) (Experimental, Thermodyn., 45)
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mg–O [1970Ber] [1970Lyk] [1971Gur]
[1971Mor]
[1972Bal]
[1972Rez]
[1973Abb] [1973Bar]
[1973Maj] [1974Kra] [1974Sah] [1974Sin1] [1974Sin2]
[1975Sch]
[1975Val1]
[1975Val2]
[1975Zal]
[1977Cor]
[1977Tri] [1978Cor]
Landolt-Börnstein New Series IV/11D4
39
Berthet, A., Perrot, P., “Equilibria in the Fe-Mg-O System at 850°C”, Mem. Sci. Rev. Met., 67(11), 747–753 (1970) (Experimental, Phase Relations, Thermodyn., 19) Lykasov, A.A., Kozheurov, V.A., “Equilibrium Oxygen Pressures over Magnesiowustite” Izv. Vyssh. Ucheb. Zaved., Chern. Metall., 13, 5–8 (1970) (Experimental, Thermodyn.) Gurevich, S.Yu., Lykasov, A.A., Mikhaylov, G.G., “Investigation of the Boundaries of the Wustite Region in the Fe-Mg-O System” (in Russian), Izv. Vyss. Uchebn. Zaved., Chern. Metall., 14(8), 28–31 (1971) (Electr. Prop., Experimental, Phase Diagram, Phase Relations, 12) Morton, I.P., Lewis, M.F., “Effect of Iron Impurities on Thermal Conductivity of Magnesium Oxide Single Crystals Below Room Temperature”, Phys. Rev. B, 3(2), 552–559 (1971) (Calculation, Experimental, 34) Balesdent, D., Evrard, O., Gleitzer, C., “Extension of Darken's Method to Heterogeneous Systems for the Calculation of the Activities” (in French), Rev. Chim. Miner., 9(1), 233–144 (1972) (Experimental, Thermodyn., 5) Rezlescu, N., Cuciureanu, E., Ioan, C., Luca, E., “Time Variation of the Electrical Conductivity in Spinel Ferrites”, Phys. Status Solidi A, 11(1), 351–9 (1972) (Experimental, Electr. Prop., 9) Abbattista, F., Borroni Grassi, G., Maja, M., “Activity of the Wuestite in (Fe,Mg)Ox Solid Solutions” (in Italian), Metall. Ital., (9), 485–488 (1973) (Experimental, 20) Barin, I., Knacke, O., Thermochemical Properties of Inorganic Substances, Springer-Verlag, Berlin, Heidelberg, New York, Verlag Stahleisen m.b.H., Düsseldorf, 292–295, 441, 462 (1973) (Thermodyn., Review, 7) Maja, M., Abbattista, F., “Thermodynamics Characteristic of the Magnesia-Wuestite System” (in Italian), Metall. Ital., 65(10), 565–570 (1973) (Experimental, Thermodyn., 18) Krawitz, A., Cohen, J.B., “Phase Relations at Low Temperatures in the Fe-Mg-O System”, J. Am. Ceram. Soc., 57(4), 186–189 (1974) (Experimental, Phase Relations, 22) Saha, P., Biggar, G.M., “An Investigation of Magnesiowüstite as a Calibrant of Oxygen Fugacy”, Indian J. Earth Sci., 1(2), 131–140 (1974) (Experimental, 7) Singh, R.N., Coble, R.L., “Dynamic Dislocation Behavior in Iron-doped Magnesium-Oxide Crystals”, J. Appl. Phys., 45(3), 990–995 (1974) (Experimental, Thermodyn., 11) Singh, R.N., Coble, R.L., “Dynamic Dislocation Behavior in Iron-Doped Magnesium-Oxide Crystals Containing Dislocation Dipoles”, J. Appl. Phys., 45(12), 5129–5135 (1974) (Experimental, Mechan. Prop., 16) Schiavello, M., Valigi, M., Pepe, F., “Structure and Catalytic Activity of Iron-Oxide and Magnesium-Oxide Solid-Solutions. 2. Catalytic Activity for N2O Decomposition”, J. Chem. Soc., Faraday Trans. 1, 71, 1642–1648 (1975) (Experimental, Thermodyn., 19) Valet, P.-M., Pluschkell, W., Engell, H.-J., “Equilibria Between MgO-FeO-Fe2O2 Solid Solutions and Oxygen” (in German), Arch. Eisenhuettenwes., 46(6), 383–388 (1975) (Calculation, Experimental, Phase Relations, 40) Valigi, M., Pepe, F., Schiavello, M., “Structure and Catalytic Activity of Iron-Oxide and Magnesium-Oxide Solid-Solutions. Structural and Magnetic Investigations”, J. Chem. Soc., Faraday Trans. 1, 71, 1631–1641 (1975) (Experimental, Crys. Structure, Magn. Prop., 27) Zalazinskii, A.G., Balakirev, V.F., Barkhatov, V.P., Chufarov, G.I., “Oxygen PressureComposition Diagram in the Mg-Fe-O System” (in Russian), Zh. Fiz. Khim., 49(6), 1551–1552 (1975) (Experimental, Thermodyn., 8) Cordischi, D., Pepe, F., Schiavello, M., Valigi, M., “Structure and Catalytic Activity of IronOxide and Magnesium-Oxide Solid-Solutions. 3. ESR Characterization”, J. Chem. Soc., Faraday Trans. 1, 73, 62–70 (1977) (Experimental, Thermodyn., 10) Trinel-Dufour, M.C., Perrot, P., “Thermodynamic Study of the Solid Solutions of the System Fe-Mg-O” (in French), Ann. Chim., 2, 309–318 (1977) (Experimental, Thermodyn., 21) Cordischi, D., Gazzoli, D., Valigi, M., “Redox Behavior of Mn, Fe and Cr in Solid Solution in Magnesium Oxide by ESR and X-Ray Diffraction Techniques”, J. Solid State Chem., 24, 371–379 (1978) (Experimental, 17)
MSIT®
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
40 [1978Kau] [1978Lyk]
[1979Sek] [1980Bra]
[1980Gor]
[1980Shi]
[1980Sim]
[1981Rot] [1980Vla]
[1982Kub] [1983Dav]
[1983One]
[1984Dob]
[1984McC]
[1984Tre]
[1987Sre]
[1988Lue] [1988Rom]
Fe–Mg–O Kaufman, L., Nesor, H., “Calculation of Quasibinary and Quasiternary Oxide Systems - I”, Calphad, 2, 35–53 (1978) (Thermodyn., Theory, *, #, 30) Lykasov, A.A., Tarvid, L.S., “Use of the Theory of Ionic Regular Solutions for Describing the Thermodynamic Properties of Magnesiowüstite”, Izv. Vyssh. Ucheb. Zaved., Chern. Metall., 10, 9–12 (1978) (Experimental, Thermodyn., 9) Sekine, K., Yamaguchi, T., “Thermodynamics of Ferrite Systems” (in Japanese), Yogyo Kyokai Shi, 87(1009), 443–452 (1979) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 60) Brach, B.Ya., Bobrysheva, N.P., Belozerskii, G.N., “Nature of Interactions between iron Atoms in Solid Solutions of the Systems FeO - MgO, FeO(1+x) - MgO, Fe3O4 - MgAl2O4, α-Fe2O3 - Al2O3”, Inorg. Mater., 16(1), 65–68 (1980), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 16(1), 84–87 (1980) (Experimental, Magn. Prop., 8) Gorobets, A.P., “Investigation Into the Thermodynamics of the Deoxidation of Iron by a Magnesium Solution” (in Russian), Metall. Koksokhim, 69, 34–37 (1980) (Thermodyn., Experimental, 8) Shishkov, V.I., Lykasov, A.A., Il"ina, A.F., “Activity of the Components of Iron-Magnesium Spinel”, Russ. J. Phys. Chem., 54(3), 440–441 (1980) translated from Zh. Fiz. Khim., 54(3), 766–767 (1980) (Experimental, Phase Relations, Thermodyn., *, #, 3) Simons, B., “Composition-Lattice Parameter Relationship of the Magnesiowuestite Solid Solution Series”, Carnegie Inst. Washington, Yearbook, 79, 376–380 (1980) (Crys. Structure, Experimental, 13) Roth, R.S., Negas, T., Cook, L.P., Phase Diagrams for Ceramists, Am. Ceram. Soc., Columbus, Ohio, 34–34 (1981) (Phase Diagram, Phase Relations, Review, 7) Vlasova, M.V., Kakazei, N.G., Ristich, M.M., “Radiospectroscopic Investigation of Solid Phase Interaction in the System MgO-Fe2O3”, Inorg. Mater., 17(5), 598–602 (1981) (Experimental, 16) Kubaschewski, O., “Iron-Magnesium” in “Iron-Binary Phase Diagram”, Springer-Verlag, Berlin, 59–60 (1982) (Review, Phase Diagram, Phase Relations, 7) Davies, P.K., Navrotsky, A., “Quantitative Correlations of Deviations from Ideality in Binary and Pseudobinary Solid Solutions”, J. Solid State Chem., 46, 1–22 (1983) (Thermodyn., Theory, 96) O’Neill, H.S.C., Navrotsky, A., “Simple Spinels: Crystallographic Parameters, Cation Radii, Lattice Energies, and Cation Distribution”, Am. Mineral., 68, 181–194 (1983) (Crys. Structure, Review, Theory, Thermodyn., 49) Dobrovinskiy, P.Y., Mesnyankina, C.L., Men", A.N., “The Registration of Mutual Solubility of Phases in Thermodynamic of Phase Equilibria”, Zh. Fiz. Khim., 58, 2326–2328 (1984) (Experimental, Thermodyn., 4) Mc Cammon, C.A., Liu, L., “The Effect of Pressure and Temperature on Nonstoichiometric Wustite FexO: The Iron rich Phase Boundary”, Phys. Chem. Miner., 10, 106–113 (1984) (Experimental, 48) Trestman-Matts, A., Dorris, S.E., Mason, T.O., “Thermoelectric Determination of Cation Distributions in Fe3O4-MgFe2O4”, J. Am. Ceram. Soc., 67, 69–74 (1984) (Experimental, Thermodyn., 39) Srecec, I., Ender, A., Woermann, E., Gans, W., Jacobsson, E., Eriksson, G., Rosen, E., “Activity-Composition Relations of the Magnesiowustite Solid Solution Series in Equilibrium with Metallic Iron in the Temperature Range 1050–1400 K”, Phys. Chem. Miner., 14, 492–498 (1987) (Experimental, Thermodyn., 15) Luecke, W., Kohlstedt, D.L., “Kinetics of the Internal Oxidation of (Mg, Fe) Solid Solution”, J. Am. Ceram. Soc., 71(1), 189–195 (1988) (Experimental, Theory, 21) Romanovskii, L.B., Terekhin, V.A., “Transformation of Magnesium Iron Oxide (MgFe2O4) into Magnesium Oxide-Iron Oxide (MgO-FeO) Solid Solution on Cooling and Heating” (in Russian), Ukr. Khim. Zh. (Russ. Ed.), 54(1), 107–109 (1988) (Experimental, Phase Relations, 8)
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mg–O [1988Sch]
[1988Wol] [1989Duf] [1989Nel]
[1989Rag] [1989Rag2] [1990Bec]
[1991Sun] [1991Wis] [1992Ita] [1993Hal] [1993Sax]
[1993Wu]
[1997Byg] [1997Fon]
[1997Ito] [1998Ber]
[1998Fab]
[1998Kan] [1999Cro]
Landolt-Börnstein New Series IV/11D4
41
Schurmann, E., Kolm, I., “Effect of the Iron(III) Oxide Contents in Steelmaking Slag on the Magnesia Solubility at Different Oxygen Partial Pressure on the Gas-Slag Equilibrium”, Steel Res., 59, 185–191 (1988) as quated in [2004Jun1] Wolfenstine, J., Kohlstedt, D.L., “Creep of (Mg, Fe)O Single Crystals”, J. Mater. Sci., 23(10), 3550–3557 (1988) (Experimental, Mechan. Prop., 42) Duffy, T.S., Anderson, D.L., “Seismic Velocities in Mantle Minerals and the Mineralogy of Upper Mantle”, J. Geophys. Res . 94, 1989, 1895–1912 (Review, Phys. Prop., 89) Nell, J., Wood, B.J., Mason, T.O., “High Temperature Cation Distribution Fe3O4-MgAl2O4MgFe2O4-FeAl2O4 Spinel from Thermopower and Conductivity Measurements”, Am. Mineral., 74, 339–351 (1989) as quoted in [2004Jun1] Raghavan, V., “The Fe-Mg-O System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Met., Calcutta, 5, 170–180 (1989) (Review, Phase Relations, Phase Diagram, 19) Raghavan, V., “The Fe-O System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. of Met., Calcutta, 5, 5–8 (1989) (Phase Diagram, Crys. Structure, Review, 3) Becker, K.D., “Spectroscopic in Situ Investigations of Solids at High Temperatures: A Mössbauer and Optical Spectroscopy Study of Diffusion and of the Kinetics of Solid State Reactions”, Solid State Ionics, Diffusion & Reactions, 39(1–2), 27–35 (1990) (Experimental, Kinetics, 29) Sundman, B., “An Assessment of the Fe-O System”, J. Phase Equilib., 12(1), 127–140 (1991) (Thermodyn., Assessment, *, #, 53) Wiser, N.M., Wood, B.J., “Experimental Determination of Activities in Fe-Mg Olivine at 1400 K”, Contrib. Mineral. Petrol., 108, 146–153 (1991) (Experimental, Thermodyn., 27) Ita, J., Stixrude, L., “Petrology, Elasticity and Composition of the Mantle Transition Zone”, J. Geophys. Res. – Solid Earth, 97B, 6849–6866 (1992) (Review, Phys. Prop., 178) Hallstedt, B., “The Magnesium-Oxygen System”, Calphad, 17(3), 281–286 (1993) (Phase Relations, Thermodyn., Theory, *, #, 8) Saxena, S.K., Chatterjee, N., Fei, Y., Chen, G., Thermodynamic Data on Oxides and Silicates: an Assessed Data Set Based on Thermochemistry and High Pressure Phase Equilibrium, Springer Verlag, Berlin, New-York, Heidelberg, pp. 428 (1993) (Review, Thermodyn.) Wu, P., Eriksson, G., Pelton, A.D., “Critical Evaluation and Optimization of the Thermodynamic Properties and Phase Diagrams of the CaO-FeO, CaO-MgO, CaO-MnO, FeO-MgO, FeO-MnO, and MgO-MnO Systems”, J. Am. Ceram. Soc., 76(8), 2065–2075 (1993) (Phase Diagram, Phase Relations, Thermodyn., Review, *, #, 73) Bygden, J., Jacobsen, A., Sichen, D., Seethraman, S., “Interdiffusion Studies in the System MgO-FeO”, Z. Metallkd., 88(5), 433 (1997) (Experimental, 17) Fontijn, W.F.J., van der Zaag, P.J., Devillers, M.A.C., Brabers, V.A.M., Metselaar, R., “Optical and Magneto-Optical Polar Kerr Spectra of Fe3O4 and Mg2+- or Al3+-Substituted Fe3O4”, Phys. Rev. B, 56(9), 5432–5442 (1997) (Experimental, 42) Itoh, H., Hino, M., Ban-ya, S., “Deoxidation Equilibrium of Magnesium in Liquid Iron” (in Japanese), Tetsu to Hagane (J. ISIJ), 83(10), 623–628 (1997) (Experimental, 13) Berry, F.J., Bohorquez, A., Greaves, C., McManus, J., Moore, E.A., Mortimer, M., “Structural Characterization of Divalent Magnetism-Doped α-Fe2O3”, J. Solid State Chem., 140, 428–430 (1998) (Crys. Structure, Experimental) Fabrichnaya, O., “The Assessment of Thermodynamic Parameters for Solid Phase in the Fe-Mg-O and Fe-Mg-Si-O Systems”, Calphad, 22(1), 85–125 (1998) (Thermodyn., Assessment, 61) Kang, S.-H., Yoo, H.-J., “Composition (x) Dependence of Nonstoichiometry (δ) in Ferrite Spinel (MgxFe1–x)3–δO4”, J. Solid State Chem., 139, 128–134 (1998) (Experimental, 23) Cross, W.B., Affleck, L., Kuznetsov, M.V., Parkin, I.P., Pankhurst, Q.A., “Self-Propagating High-Temperature Synthesis of Ferrites MFe2O4 (M = Mg, Ba, Co, Ni, Cu, Zn); Reactions in an External Magnetic Field”, J. Mater. Chem., 9, 2545–2552 (1999) (Crys. Structure, Experimental, Magn. Prop., 26)
MSIT®
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
42 [1999Har]
[2000Ber]
[2000Fab]
[2000Kan] [2000Leb]
[2000Vas]
[2001And]
[2001Boi]
[2001Dub]
[2001Mey]
[2002Dec]
[2002Jac]
[2002Kat]
[2002Mer]
[2002Zha] [2003Hon]
Fe–Mg–O Harrison, R.J., Putnis, A., “Determination of the Mechanism of Cation Ordering in Magnesioferrite (MgFe2O4) from the Time- and Temperature-Dependence of Magnetic Susceptibility”, Phys. Chem. Miner., 26, 322–332 (1999) (Crys. Structure, Experimental, Kinetics, Magn. Prop., 39) Berry, F.J., Greaves, C., Helgason, O., McManus, J., Palmer, H.M., Williams, R.T., “Structural and Magnetic Properties of Sn-, Ti-, and Mg-substituted α-Fe2O3: a Study by Neutron Diffraction and Mössbauer Spectroscopy”, J. Solid State Chem., 151, 157–162 (2000) (Crys. Structure, Experimental, Magn. Prop., 16) Fabrichnaya, O.B., “Thermodynamic Modelling of Melting in the System FeO-MgO-SiO2O2 at Pressure of 1 bar”, Calphad, 24(2), 113–131 (2000) (Thermodyn., Phase Relations, Phase Diagram, Calculation, 24) Kang, S-H., Chang, S-H., Yoo, H-I., “Phase Stability in the System Mg-Fe-O”, J. Solid State Chem., 149, 33–40 (2000) (Crys. Structure, Experimental, Phase Relations, 29) Lebrun, N., “Iron-Magnesium-Oxygen”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.10323.1.20, (2000) (Crys. Structure, Phase Diagram, Assessment, 53) Vassent, J.L., Marty, A., Gilles, B., Chatillon, C., “Thermodynamic Analysis of Molecular Beam Epitaxy of MgO(s). II. Epitaxial Growth of MgO Layers on Fe(001) Substrates”, J. Cryst. Growth, 219(4), 444–450 (2000) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 13) Andrault, D., Bolfan-Casanova, N., “High-Pressure Phase Transformation in the MgFe2O4 and Fe2O3-MgSiO3 Systems”, Phys. Chem. Miner., 28, 211–217 (2001) (Crys. Structure, Experimental, 26) Boiocchi, M., Caucia, F., Merli, M., Prella, D., Ungaretti, L., “Crystal-Chemical Reason for the Immiscibility of Periclase and Wuestite under Lithospheric P, T Conditions”, Eur. J. Mineral, 13(5), 871–881 (2001) (Crys. Structure, Experimental, Morphology, Phase Relations, 36) Dubrovinsky, L., Dubrovinskaia, N., Annersten, H., Halenius, E., Harryson, H., “Stability of (Mg0.5Fe0.5)O and (Mg0.8Fe0.2)O Magnesiowüstite in the Lower Mantle”, Eur. J. Mineral, 13(5), 857–861 (2001) (Crys. Structure, Experimental, Phase Relations, 22) Meyerheim, H.L., Popescu, R., Krischner, J., Jedrecy, N., Sauvage-Simkin, M., Heinrich, B., Pinchaux, R., “Geometrical and Compositional Structure at Metal-Oxide Interfaces: MgO on Fe(001)”, Phys. Rev. Lett., 87(7), 076102–1-4 (2001) (Crys. Structure, Experimental, Magn. Prop., 25) Decterov, S.A., Jung, I.-H., Pelton, A.D., “Thermodynamic Modeling of the FeO-Fe2O3MgO-SiO2 System”, J. Am. Ceram. Soc., 85(12), 2903–2910 (2002) (Calculation, Phase Diagram, Phase Relations, 68) Jacobsen, S.D., Spetzler, H.A., Reichmann, H.J., Smyth, J.R., Mackwell, S.J., Angel, R.J., Bassett, W.A., “Gigahertz Ultrasonic Interferometry at High P and T: New Tools for Obtaining a Thermodynamic Equation of State”, J. Phys. Condens. Matter., 14(44), 11525–11530 (2002) (Experimental, Thermodyn., 11) Katayama, I., Iseda, A., “Thermodynamic Study of Spinel-Type Solid Solutions of the Fe3O4-MgFe2O4 Coexisting with Fe2O3 by Emf Method”, Scand. J. Metall., 31(6), 374–378 (2002) (Experimental, Thermodyn., 12) Merli, M., Pavese, A., Ranzini, M., “Study of the Electron Density in MgO, (Mg0.963Fe0.037) O and Cu2O by the Maximum Entropy Method and Miltipole Refinements: Comparison between Methods”, Phys. Chem. Miner., 29(7), 455–464 (2002) (Crys. Structure, Electronic Structure, Experimental, 38) Zhang, J., Kostak, P., “Thermal Equation of State of Magnesiowüstite”, Phys. Earth Planet. Interiors, 129, 301–311 (2002) (Crys. Structure, Experimental, Magn. Prop., 31) Honda, T., Kaneko, S., “Phase Relations in the MgO-Rich Region of the System MgOFe2O3”, J. Ceram. Soc. Jpn., 111(11), 841–847 (2003) (Electronic Structure, Phase Diagram, Phase Relations, 19)
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mg–O [2003Lin]
[2003Yam] [2004Fab]
[2004Jun1]
[2004Jun2]
[2004Kon]
[2004Lev]
[2004Shi]
[2005Ant]
[2005Gat]
[2005Jac]
[2005Wes]
[2006Kan1]
[2006Kan2]
[2006Lin]
[2006Lon]
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43
Lin, J.F., Heinz, D.L., Mao, H.K., Hemley, R.J., Devine, J.M., Li, J., Shen, G.Y., “Stability of Magnesiowüstite in Earths Lower Mantle”, Proc. Nat. Acad. Sci. USA, 100(8), 4405–4408 (2003) (Experimental, Thermodyn., 35) Yamazaki, D., Irifune, T., “Fe-Mg Interdiffusion in Magnesiowüstite up to 35 GPa”, Earth Planet. Sci. Letters, 216(3), 301–311 (2003) (Experimental, Phys. Prop., 35) Fabrichnaya, O., Saxena, S.K., Richet, P., Westrum, E.F., Thermodynamic Data, Model and Phase Diagrams in Multicomponent Oxide Systems, Springer Verlag, Berlin, Heidelberg, pp. 198 (2004) (Review, Phase Diagram, Phase Relations, Thermodyn., 479) Jung, I.-H., Decterov, S.A., Pelton, A.D., “Critical Thermodynamic Evaluation and Optimization of the Fe-Mg-O System”, J. Phys. Chem. Solids, 65(10), 1683–1695 (2004) (Assessment, Calculation, Phase Diagram, Phase Relations, Thermodyn., 81) Jung, I.-H., Decterov, S.A., Pelton, A.D., “A Thermodynamic Model for Deoxidation Equilibria in Steel”, Metall. Mater. Trans. B, 35b(3), 493–507 (2004) (Calculation, Theory, Thermodyn., 100) Kondo, T., Ohtani, E., Hirao, N., Yagi, T., Kikegawa, T., “Phase Transitions of (Mg, Fe)O at Megabar Pressures”, Phys. Earth Planet. Interiors, 143–144, 201–213 (2004) (Crys. Structure, Experimental, Optical Prop., Phase Relations, 30) Levy, D., Diella, V., Dapiaggi, M., Sani, A., Gemmi, M., Pavese, A., “Equation of State, Structural Behaviour and Phase Diagram of Synthetic MgFe2O4, as a Function of Pressure and Temperature”, Phys. Chem. Miner., 31(2), 122–129 (2004) (Crys. Structure, Experimental, Phase Relations, 38) Shibata, H., Jiang, X., Valdez, M., Cramb, A.W., “The Contact Angle Between Liquid Iron and Single Crystal Magnesium Oxide Substrate at 1873 K”, Metall. Mater. Trans. B, 35B(1), 179–181 (2004) (Interface Phenomena, Experimental, 13) Antao, S.M., Hassan, I., Parise, J.B., “Cation Ordering in Magnesioferrite, MgFe2O4, to 982°C Using in Situ Synchrotron X-Ray Powder Diffraction”, Am. Mineral., 90(1), 219–228 (2005) (Calculation, Crys. Structure, Experimental, Thermodyn., 38) Gateshki, M., Petkov, V., Pradhan, S.K., Vogt, T., “Structure of Nanocrystalline MgFe2O4 from X-ray Diffraction, Rietveld and Atomic Pair Distribution Function Analysis”, J. Appl. Crystallogr., 38(5), 772–779 (2005) (Crys. Structure, Experimental, Magn. Prop., 46) Jacobsen, S.D., Lin, J.-F., Angel, R.J., Shen, G.Y., Prakapenka, V.B., Dera, P., Mao, H.-K., Hemley, R.J., “Single-Crystal Synchrotron X-ray Diffraction Study of Wüstite and Magnesiowüstite at Lower-Mantle Pressures”, J. Synchrotron Radiation, 12, 577–583 (2005) (Crys. Structure, Experimental, Kinetics, 38) Van Westrenen, W., Li, J., Fei, Y., Frank, M.R., Hellwig, H., Komabayashi, T., Mibe, K., Minarik, W.G., Van Orman, J.A., Watson, H.C., Funakoshi, K.-I., Schmidt, M.W., “Thermoelastic Properties of (Mg0.64Fe0.36)O Ferropericlase Based on in Situ X-Ray Diffraction to 26.7 GPa and 2173K”, Phys. Earth Planet. Interiors, 151(1–2), 163–176 (2005) (Crys. Structure, Experimental, Mechan. Prop., Thermodyn., 42) Kantor, I., Dubrovinsky, L., McCammon, C., Kantor, A., Pascarelli, S., Aquilanti, G., Crichton, W., Mattesini, M., Ahuja, R., Almeida, J., Urusov, V., “Pressure-Induced Phase Transition in Mg0.8Fe0.2O Ferropericlase”, Phys. Chem. Miner., 33, 35–44 (2006) (Calculation, Experimental, Phase Relations, Kinetics, 53) Kantor, I.Y., Dubrovinsky, L.S., McCammon, C.A., “Spin Crossover in (Mg, Fe)O: a Mössbauer Effect Study with an Alternative Interpretation of X-ray Emission Spectroscopy Data”, Phys. Rev. B, 73(B), 100101 (2006) (Experimental, Crys. Structure, Electronic Structure, Kinetics, 14) Lin, J.-F., Gavriliuk, A.G., Struzhkin, V.V., Jacobsen, S.D., Sturhahn, W., Hu, M.Y., Chow, P., Yoo, C.S., “Pressure-Induced Electronic Spin Transition of Iron in Magnesiowüstite-(Mg,Fe) O”, Phys. Rev. B, 73, 113107 (2006) (Experimental, Phys. Prop., 28) Long, M.D., Xiao, X.H., Jiang, Z.T., Evans, B., Karato, S., “Lattice Preferred Orientation in Deformed Polycrystalline (Mg, Fe)O and Implications for Seismic Anisotropy in D”, Phys.
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[2006Tom]
[2006Zha]
[Mas2] [V-C2]
Fe–Mg–O Earth Planet. Interiors, 156(1–2), 75–88 (2006) (Crys. Structure, Experimental, Mechan. Prop., 36) Tommaseo, C.E., Devine, J., Merkel, S., Speziale, S., Wenk, H.-R., “Texture Development and Elastic Stresses in Magnesiowüstite at High Pressure”, Phys. Chem. Miner., 33, 84–97 (2006) (Crys. Structure, Experimental, Kinetics, Mechan. Prop., 54) Zhang, L., Gong, Z.-Z., “Shock Compression and Phase Transitions of Magnesiowüstite (Mg, Fe)O up to Earths Lowermost Mantle Conditions”, Chin. Phys. Lett., 23(11), 3049–3051 (2006) (Experimental, Mechan. Prop., 21) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_13 # Springer 2008
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Fe–Mg–S
1
Iron – Magnesium – Sulfur Ludmyla Tretyachenko
Introduction Investigation of the phase relations in the Fe-Mg-S system was started because some compositions of the MgS-FeS system were found to be constituents of meteorites (niningerite [1971Ski] as well as keilite [2002Shi, 2007Kei]). Niningerite is unknown in terrestrial rocks. The phase relations in the MgS-FeS section were mostly studied [1971Ski, 1973Kur, 1984McC, 1984Osb, 2006And]. Electronic structure of MgS-FeS compositions has been investigated by [2000Far, 2001Ste, 2002Far, 2002Kra, 2004Kra]. The equilibria between Mg and S in liquid Fe have been studied by [1992Han, 1995Zho, 1997Han]. Desulfurization of molten iron with magnesium has been investigated by [1978Nas, 1979Iro, 2007Yan]. A mechanism of resulfurization of liquid Fe with MgS has been studied by [2005Yan, 2007Yan]. Phase equilibria in the MgS-FeS section at high pressure were studied experimentally and calculated from thermodynamic data by [1984McC]. Samples along the MgS-FeS section were prepared of mixtures of pure elements [1973Kur, 1984McC, 2000Far, 2004Kra] or MgS and FeS preliminary prepared [1971Ski, 1984Osb, 2006And] by reactions during heating in sealed evacuated silica glass tubes at appropriate temperatures mostly followed by quenching in cold water. Earlier works [1971Ski] and [1973Kur] on the phase reactions in the Fe-Mg-S system (FeS-MgS section) were reviewed by [1988Rag]. A more complete review was presented by [2000Tre]. Experimental investigations on the phase relations and the structure of phases in the Fe-Mg-S system are listed in Table 1. Binary Systems The Fe-Mg phase diagram is given in [1982Kub1]. The Fe-S phase diagram is taken from [Mas2], where it is shown according to [1982Kub2]. The Mg-S phase diagram is not determined. The MgS compound melts congruently at 1927°C [1989And]. Solid Phases Ternary compounds are not found. Data on the binary compounds and solid solutions on their base are given in Table 2. Extended ternary solid solutions with the NaCl type (B1) crystal structure were found only for MgS [1971Ski, 1973Kur, 1984McC, 1984Osb, 2000Far, 2006And]. The solid solutions (Mg1–xFex)S extend up to 75 mol% FeS at ∼1250°C [2006And]. Concentration dependence of the (Mg1–xFex)S lattice parameters was determined by [1971Ski, 1973Kur, 1984McC, 1984Osb] and found to be linear. The equation a = 520.4 – 0.1299x was proposed by [1984McC] for the concentration dependence of the lattice parameter of (Mg1–xFex)S. [2006And] revealed a positive deviation from the Vegards’ law. The MgS solubility in FeS was found to be below 0.1 mol% up to 1000°C [1971Ski]. The maximum solubility was determined to be 2 mol% MgS at ∼1200°C [2006And]. Quasibinary Systems Taking into account that the congruent melting of FeS takes place at 52 at.% S (FeS1.08) and the results of [2006And] obtained for the FeS-MgS and FeS1.03-MgS sections, the FeS1.08-MgS section can be considered as the quasibinary one of the peritectic type. The peritectic reaction L + MgS ⇌ γFeS takes place at ∼1200°C, the peritectic point is located at 2 ± 0.5 mol% MgS. The solubility of MgS in γFeS does not exceed ∼1 mol% but FeS dissolves in MgS up to ∼75 mol%. Landolt-Börnstein New Series IV/11D4
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Fe–Mg–S
Invariant Equilibria The presence of Fe in as-cast MgS-FeS samples rich in FeS evidenced the existence of a three-phase region MgS + γFeS + (γFe) below the solidus surface [2006And]. Taking into account a eutectic melting at 1070°C in similar compositions reported by [1971Ski], the invariant equilibrium MgS + L ⇌ γFeS + (γFe) at ∼1070°C can be assumed to take place in the ternary system. Temperature – Composition Sections The section FeS-MgS is shown in Fig. 1 according to [2006And] with some corrections taking into account that the congruent melting point of FeS corresponds to 52 at.% S rather than 50 at.% S. The peritectic reaction MgS + L ⇌ γFeS was found to take place in this section at 1197°C. Eutectic type grains of iron were observed in as-cast samples containing ∼3–4 mol% FeS. The samples annealed at 497 and 897°C for 1300 and 740 h respectively did not contain the Fe grains. A eutectic melting at 1070°C in samples containing Fe was observed also by [1971Ski]. The Fe phase grains were not revealed in FeS1.03-MgS samples [2006And]. The peritectic temperature in this section was determined as 1200°C. The solvus line determined by [1971Ski, 1984McC] is consistent with that shown by [2006And]. [1973Kur] determined the solubility of FeS in MgS to be 50 mol% at 900°C, that is lower than the values found in other works. [1984McC] assumed that this may be caused by too slow quenching from a high temperature. A dependence of the MgS solid solutions composition on the quenching rate was studied by [1971Ski]. The FeS phase in samples quenched from a high temperature and examined by means of XRD at room temperature was shown to be the low temperature modification of this sulfide, αFeS (troilite) [1971Ski] in spite it is well established that in pure FeS the troilite αFeS is not stable above 138°C. Polymorphic transitions of FeS were revealed in the MgS-FeS section by means of DTA at 312°C (γFeS/ βFeS) and 132°C (βFeS/αFeS) [2006And]. Thermodynamics Thermodynamic data were derived from the FeS-MgS solid solutions to calculate the phase diagram at high pressure [1984McC]. The values of ΔG0 (0, 1000°C) and ΔV (0, 25°C) have been determined to be –4.1 kJ·mol–1 and –1.4 cm3·mol–1 respectively. The equilibrium between magnesium and sulfur in liquid iron at 1600°C has been studied by [1992Han, 1995Zho, 1997Han] using the vapor-liquid equilibrium method. The solubility product for MgS has been measured to be K′MgS = 2.36·10–4 and the equilibrium constant has been obtained to be KMgS = 2.0·10–4 for the reaction MgS(s) = {Mg} + {S}. Notes on Materials Properties and Applications The microhardness of the Fe and MgS solid solutions in the samples annealed at 897°C for 740 h measured under a load of 0.4 N by [2006And] is shown in Fig. 2. Magnesium can be used as a desulfurizing agent for molten iron [1978Nas, 1979Iro, 1997Han, 2005Yan, 2007Yan]. Desulfurization kinetics was studied by [1979Iro, 2005Yan, 2007Yan]. [2007Yan] studied a mechanism of resulfurization and proposed methods of its prevention. Miscellaneous The FeS-MgS phase diagram was determined at high pressures using samples with compositions close to the phase boundary between the single-phase MgS (NaCl type) region and the two-phase MgS + FeS (NiAs type) region [1984McC]. The single-phase MgS sample (50 mol% FeS) and the two-phase MgS + FeS one (80 mol% FeS) were used. Experiments were carried out at 1000°C at high pressures up to 6.0 GPa. The obtained samples have been studied by means of XRD. The phase diagram of FeS-MgS determined from the high pressure experiments is shown in Fig. 3. The most notable feature of the phase diagram is a considerable displacement of the MgS phase boundary towards MgS at high pressure and a pressure insensitivity of the composition of the FeS phase. DOI: 10.1007/978-3-540-78644-3_14 # Springer 2008
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Fe–Mg–S
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The composition of the MgS/(MgS + FeS) phase boundary in the FeS-MgS system at high pressure and 1000°C has been calculated from the thermodynamic data derived from the solid solution. In the calculation FeS (NiAs type) was assumed to be a pure phase stable at 1000°C from 0.1 to 6.0 GPa. Nonstoichiometry of FeS was neglected. A good agreement was found between the experimental results and the calculated curve. So, the analysis, which predicts exsolution of FeS (NiAs type) from MgS (NaCl type) at high pressure, was shown to be valid. Mössbauer spectra of (Mg1–xFex)S at 78 and 300 K have been obtained by [1973Kur]. The Mössbauer spectra obtained for these solid solutions at 4.2 and 298 K, as well as for this phase quenched from high pressure, have been discussed by [1984McC]. Mössbauer spectra of (Mg,Fe)S solid solutions within the composition range 8.4 to 50.0 mol% FeS quenched from 1000°C have been obtained at 300 K and discussed by [1984Osb]. X-ray absorption near-edge structure (XANES) spectra were used to analyze an evolution of local electronic structure of (Mg1–xFex)S solid solutions [2000Far, 2001Ste, 2002Far, 2002Kra, 2004Kra]. The XANES spectra were obtained experimentally and calculated theoretically. It was shown that in the (Mg1–xFex)S solid solutions FeS changes its structure from NiAs type (B8) to NaCl type (B1) under the influence of the MgS matrix at x below 0.68 (the samples were annealed at 1000°C). The S K-edge XANES spectrum of the hypothetical cubic B1 type FeS phase was derived from the spectrum of Mg0.375Fe0.625S by [2000Far]. The lattice parameters of the hypothetical B1 type FeS was evaluated by means of extrapolation to be a = 508.4 pm for the linear concentration relationship and a = 504.5 pm for the second-order polynomial [2006And]. Table 1.
Investigations of the Fe-Mg-S Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/ Phase Range Studied
[1971Ski]
XRD, light microscopy, EMPA
600–1000°C; MgS-FeS
[1973Kur]
XRD, Mössbauer spectroscopy (78 and 300 K)
Quenching from 900°C, 0 to 70 mol% FeS
[1984Osb]
XRD, Mössbauer spectroscopy (78 and 300 K)
Quenching from 1000°C, 8.39 to 50 mol% FeS
[1984McC]
XRD, Mössbauer spectroscopy (4.2 and 298 K), experiments at high pressures (up to 6 GPa), calculation
1000°C, (Mg1–xFex)S, 0 ≤ x ≤ 0.67
[1992Han]
Vapor - molten metal equilibrium method
1600°C, MgS in molten Fe
[1995Zho]
Vapor - liquid equilibrium method
1600°C, MgS in molten Fe
[1997Han]
Vapor - liquid equilibrium method
1600°C, MgS in molten Fe
[2000Far]
XRD, X-ray absorption spectroscopy
1000°C (annealed and quenched), (Mg1–xFex)S, 0 ≤ x ≤ 1
[2001Ste]
X-ray absorption spectroscopy, calculation
(Mg1–xFex)S, 0 ≤ x ≤ 0.68
[2002Far]
X-ray absorption spectroscopy
(Mg1–xFex)S
[2002Kra]
X-ray absorption spectroscopy, calculation
(Mg1–xFex)S
[2004Kra]
X-ray absorption spectroscopy
(Mg1–xFex)S
[2006And]
Microscopy, XRD, microhardness measurement, DTA, melting temperature determination by means of visual polythermal analysis (VPTA) (incipient melting)
up to 2227°C, MgS-FeS, MgS-FeS1.03
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4 Table 2.
Fe–Mg–S Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(δFe) 1538–1394
cI12 Im 3m W
a = 293.15
at 1390°C [Mas2]
(γFe) 1394–912
cF4 Fm 3m Cu
a = 364.67
at 915°C [Mas2]
(αFe) < 912
cI2 Im 3m W
a = 286.65
at 25°C [Mas2]
(Mg) < 650
hP2 P63/mmc Mg
a = 320.94 c = 521.07
[Mas2]
(βS) 115.22–95.5
mP64 P21/c βS
a = 1102 b = 1096 c = 1090 β = 96.7
at 103°C [Mas2, V-C2]
(αS) < 95.5
oF128 Fddd αS
a = 1046.4 b = 1286.60 c = 2448.60
at 25°C [Mas2, V-C2]
γFeS (h1) 1188–315
hP4 P63/mmc NiAs
βFeS (h2) 315–∼60
αFeS (r) < 138
hP24 P 62c FeS
hP6 P63/mmc FeS
a = 344.36 c = 587.59 a = 345.59 to 344.40 c = 577.89 to 587.00 a = 344.7 ± 0.5 c = 588 ± 1
a = 596.3 c = 1175.4 a = 586.1 c = 1157.7 a = 599.8 c = 1171 a = 597 c = 1173 a = 345.59 ± 0.05 c = 577.89 ± 0.05
50 to 55 at.%, congruent melting at 52 at.% S [Mas2] [V-C2] for FexS, x = 0.94 to 1 (given as NiS type) [V-C2] at 6 GPa, 1000°C in two-phase (80FeS-20MgS) (mol%) composition [1984McC] troilite [Mas2, 1982Kub2] at 21°C, p = 0.1 MPa [V-C2] at 21°C, p = 3.33 GPa [V-C2] at 120°C [V-C2] [2006And] in Fe0.94S [V-C2]
(continued) DOI: 10.1007/978-3-540-78644-3_14 # Springer 2008
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Fe–Mg–S
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
FeS (I)
oP8 Pnma MnP
FeS (mackinawite)
tP4 P4/nmm PbO
βFeS2 (h) 743–∼425
cP12 Pa3 FeS2
αFeS2 (r) ≲ 425
oP6 Pnnm FeS2
(Mg1–xFex)S
cF8 Fm 3m NaCl
MgS (niningerite) < 1927
Lattice Parameters [pm]
a = 582.5 b = 346.8 c = 693.5 a = 565 b = 331.6 c = 663.1 a = 571.6 b = 334.7 c = 669.4
Comments/References
50 to ∼52 at.% S [Mas2] at 50 at.% S [Mas2] at 190°C [V-C2]
at 21°C, p = 6.35 GPa [V-C2]
at 21°C, p = 4.15 GPa [V-C2]
a = 368 c = 504
at 17°C [V-C2]
a = 541.8
pyrite [V-C2, Mas2, 1982Kub2]
a = 443.6 b = 541.6 c = 338.1
a = 520 a = 520.20 to 511.16 a = 520.3±0.2 to 515 a = 519.5 to 513.5 a = 519.6 a = 514.6 a = 512.7
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marcasite [V-C2]
[1971Ski, 1973Kur, 1984Osb, 1984McC, 2000Far, 2006And] 0 ≤ x ≤ 0.75 at 1200°C [2006And] [V-C2, 1989And] 0 ≤ x ≤ 0.67 [1971Ski] 0 ≤ x ≤ 0.6 [1973Kur] 0.0839 ≤ x ≤ 0.5, quenched from 1000°C [1984Osb] at x = 0 [2006And] at x = 0.35 [2006And] at x = 0.58 [2006And]
DOI: 10.1007/978-3-540-78644-3_14 # Springer 2008
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Fig. 1. Fe-Mg-S.
Fe–Mg–S
The FeS-MgS section
DOI: 10.1007/978-3-540-78644-3_14 # Springer 2008
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Fe–Mg–S
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Fig. 2. Fe-Mg-S. Concentration dependence of microhardness of the phases in the FeS-MgS samples annealed at 897°C: a - FeS; b – MgS
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Fig. 3. Fe-Mg-S.
Fe–Mg–S
Phase diagram of the FeS-MgS determined from high-pressure experiments at 1000°C
DOI: 10.1007/978-3-540-78644-3_14 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mg–S References [1971Ski]
[1973Kur]
[1978Nas] [1979Iro] [1982Kub1] [1982Kub2] [1984McC]
[1984Osb]
[1988Rag]
[1989And] [1992Han]
[1995Zho]
[1997Han] [2000Far]
[2000Tre]
[2001Ste]
[2002Far]
[2002Kra]
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Skinner, B.J., Luce, F.D., “Solid Solutions of the Type (Ca, Mg, Mn, Fe)S and Their Use as Geothermometers for the Enstatite Chondrites”, Am. Mineral., 56, 1269–1296 (1971) (Crys. Structure, Experimental, 19) Kurash, V.V., Kulikov, G.A., Makarov, E.S., “Investigation of the Isomorphism of Mg and Fe in the Niningerite MgS - Troilite FeS System” (in Russian), Geokhimiya, (8), 1266–1267 (1973) (Crys. Structure, Experimental, Phase Relations, 2) Nashiwa, H., Ueda, S., Yamaguchi, S., Nagao, N., “Desulfurization of Hot Metal With Magnesium”, Proc. Annu. Meet.-Int. Magnesium Assoc., (35), 19–23 (1978) (Experimental, 4) Irons, G.A., “Desulphurization Kinetics of Molten Iron by Magnesium Vapour”, Diss. Abstr. Int., 39B, 4535 (1979) (Experimental, Kinetics, 0) Kubaschewski, O., “Iron – Magnesium” in “Iron Binary Phase Diagrams”, Springer Verlag, Berlin, Verlag Stahleisen, Duesseldorf, 59–60 (1982) (Phase Diagram, Review, #, 7) Kubaschewski, O., “Iron – Sulfur” in “Iron Binary Phase Diagrams”, Springer Verlag, Berlin, Verlag Stahleisen, Duesseldorf, 364–365 (Phase Diagram, Review, #, 20) McCammon, C.A., Jackson, I., Ringwood, A.E., Cashion, J.D., “The Binary Systems FeSMgS and FeS-MnS: Mössbauer Spectroscopy of the B1 Solid Solutions and High Pressure Phase Equilibria”, Phys. Chem. Miner., 11, 182–193 (1984) (Calculation, Crys. Structure, Electronic Structure, Experimental, Phase Diagram, #, 39) Osborne, M.D., Fleet, M.E., “Mössbauer Investigation of Niningerite Solid Solutions (Mg, Fe)S”, Phys. Chem. Miner., 10, 245–249 (1984) (Crys. Structure, Electronic Structure, Experimental, Phase Relations, 23) Raghavan, V., “The Fe-Mg-S (Iron-Magnesium-Sulphur) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Met., Calcutta, 2, 153 (1988) (Crys. Structure, Review, Phase Relations, 3) Andreev, O.V., Kertman, A.V., Kislovskaya, T.M., “The MgS-La2S3 System” (in Russian), Zh. Neorgan. Khim., 34, 2913–2915 (1989) (Crys. Structure, Experimental, Phase Diagram, 11) Han, Q., Zhang, X., Wang, P., Chen, D., Wang, C., “Thermodynamic Behavior of Alkaline Earth Elements in Molten Iron and Nickel” (in Chinese), Beijing Keji Daxue Xuebao (Univ. Sci. Technol. Bejing), 14, 85–93 (1992) (Experimental, Review, Thermodyn., 37) Zhou, D., Han, Q., “Investigation of Equilibrium Between Magnesium and Sulfur in Liquid Iron” (in Chinese), Beijing Keji Daxue Xuebao (Univ. Sci. Technol. Bejing), 17, 576–579 (1995) (Experimental, Thermodyn., 8) Han, Q., Zhou, D., Xiang, C., “Determination of Dissolved Sulfur and Mg-S, Mg-O Equilibria in Molten Iron”, Steel Res., 68, 9–14 (1997) (Calculation, Experimental, Thermodyn., 13) Farrell, S.P., Fleet, M.E., “Evolution of Local Electronic Structure in Cubic Mg1–xFexS by S K-edge XANES Spectroscopy”, Solid State Commun., 113, 69–72 (2000) (Electronic Structure, Experimental, Phase Relations, 21) Tretyachenko, L., Bondar, A., Tretyachenko, C., “Iron – Magnesium – Sulfur” in “Ternary Alloys”, Ed. Effenberg, G., Weinheim (FRG); VCH., Vol. 17, 558–563 (2000); MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G., (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.24243.1.20 (1992) (Assessment, Crys. Structure, Phase Diagrams, 9) Stekhin, I.E., Soldatov, A.V., Farrel, S.P., Fleet, M.E., “Electronic Structure Investigation of Mg1–xFexS Solid Solution: X-ray Absorption Study” (Abstract), J. Synchrotron Rad., 8, 238–239 (2001) (Calculation, Electronic Structure, Experimental, 8) Farrell, S.P., Fleet, M.E., Stekhin, I.E., Kravtsova, A., Soldatov, A.V., Liu, X., “Evolution of Local Electronic Structure in Alabandite and Niningerite Solid Solutions {(Mn,Fe)S, (Mg, Mn)S, (Mg,Fe)S} Using Sulfur K- and L-Edge XANES Spectroscopy”, Am. Mineral., 87, 1321–1332 (2002) (Electronic Structure, Experimental, 59) Kravtsova, A.N., Stekhin, I.E., Soldatov, A.V., Liu, X., Fleet, M.E., “B1 Phase of FeS in Mg1–xFexS Solid Solution: X-ray Absorption Study”, Phys. Status Solidi B, 234, R4–R5 (2002) (Calculation, Electronic Structure, Experimental, Phase Relations, 4) MSIT®
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[2004Kra]
[2005Yan]
[2006And]
[2007Kei]
[2007Yan]
[Mas2] [V-C2]
Fe–Mg–S Shimizu, M., Yoshida, H., Mandarino, J.A., “The New Mineral Species Keilite, (Fe,Mg)S, the Iron-Dominant Analogue of Niningerite”, Can. Mineral., 40, 1687–1692 (2002) (Review, 18) Kravtsova, A.N., Stekhin, I.E., Soldatov, A.V., Liu, X., Fleet, M.E., “Electronic Structure of MS (M = Ca, Mg, Fe, Mn): X-Ray Absorption Analysis”, Phys. Rev. B, 69, 134109–1-12 (2004) (Calculation, Electronic Structure, Experimental, Phase Relations, 62) Yang, J., Kuwabara, M., Teshigawara, T., Sano, M., “Mechanism of Resulfurization in Magnesium Desulfurization Process of Molten Iron”, ISIJ Int., 45, 1607–6015 (2005) (Experimental, Kinetics, 25) Andreev, O.V., Solov’eva, A.V., Burkhanova, T.M., “MgS-FeS Phase Diagram”, Russ. J. Inorg. Chem. (Engl. Transl.), 51, 1826–1828 (2006), translated from Zh. Neorg. Khimii, 51, 1938–1941 (2006) (Experimental, Phase Diagram, #, 9) Keil, K., “Occurrence and Origin of Keilite, (Fe > 0.5, Mg < 0.5)S, in Enstatite Chondrite Impact-Melt Rocks and Impact-Melt Breccias”, Chemie der Erde (Geochemistry), 67, 37–54 (2007) (Experimental, Morphology, Phase Relations, 52) Yang, Y., Kuwabara, M., Sakai, T., Uchida, N., Liu, Z.Z., Sano, M, “Simultaneous Desulfurization and Deoxidation of Molten Steel with in Situ Produced Magnesium Vapor”, ISIJ Int., 47, 418–426 (2007) (Experimental, Abstract, 26) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Iron – Magnesium – Silicon Nathalie Lebrun, Christian Baetzner, Athanasios Stamou, James Robinson, Lazar Rokhlin
Introduction The Fe-Mg-Si system alloys were studied in a number of the works, Table 1. A great part of them was devoted to the Fe-Mg-Si phase diagram. [1954Zwi] presented an equilibrium liquidus surface obtained from thermal analysis, metallography and Debye-Scherrer patterns in the silicon rich composition range. The part of the diagram which was not investigated experimentally, was presented schematically [1954Zwi]. He stipulated that the primary crystallization of the field FeSi phase occupies the largest part of the liquidus surface. [1970Gui, 1971Gui] determined at 1454.5°C the extent of the miscibility gap between iron rich and magnesium rich liquids. The magnesium solubility was measured at lower silicon levels (5 to 17 mass% Si) in the iron rich region and the solubility of iron in the magnesium rich liquid. Measurements have been performed with a new technique for vapor pressure measurements. A series of melts was made from high purity materials. Preliminary experiments have been done to determine the time needed to establish the equilibrium and led to a good reproducibility after five minutes of measurements. The profile of the miscibility gap was constructed from a standard polynomial regression program. The tie-lines have also been reported. [1967Tka] measured the dependence of the solubility of iron and silicon in liquid magnesium at 800°C on the annealing time. [1967Rod, 1990Nor] determined liquidus curves in the magnesium rich corner. The solubility of iron in magnesium has been examined at 800°C by measuring the solubility rate constant and the diffusion coefficient of iron in magnesium [1967Rod]. [1990Nor] determined the solubility at 670, 710 and 750°C. For this last study [1990Nor], alloys were prepared from 15 kg melts to which the required amount of Si was added at 780°C. The concentration of Si has been reduced to 10 ppm by cooling the melt to 670°C and decanting metal from the top of the melt bath before alloying. The melt was then saturated with iron by adding 50 g FeCl3. Samples were taken at 760, 720 and 680°C before and after 1.5 h holding periods. At each sampling, molten alloy taken from the top of the melt was cast in three ways: in a standard disc mould, in a OKI mould for remelting and spun on a water-cooled brass wheel to form 0.3–1.0 mm thick strips. The OKI samples were reheated to the original sampling temperature and centrifuged during cooling and solidification. The standard discs were examined by atomic absorption spectroscopy. The melt-spun and the centrifuged samples were analyzed by electron microprobe. The solubility of magnesium in the iron rich liquid has been measured at 1600°C with a micro-oven (molybdenum heater) in an atmosphere of purified argon (2–5 atmospheres) [1984Age]. The iron alloys of mass 8 g have been melted from special purity iron powder and magnesium of purity not less than 99.9 mass%. A theoretical estimate has been made which accounted for the experimental data. [1996Li] determined experimentally effect of Si on the solubility of Mg vapor in liquid Fe at 1600°C. In the experiments iron, magnesium and silicon taken in the respective quantities were placed into crucibles which were sealed in a molybdenum reaction chamber and heated in the furnace at 1600°C for 4 h. The chamber was then taken out from the furnace and immediately quenched into ice water. The Fe ingots obtained were cleaned and analyzed for Mg and Si. By their contents effect of Si on solubility of Mg in liquid Fe was established. The iron of high purity (99.94435%) was used in the experiments. The Mg vapor partial pressure was estimated to be pMg = 0.41 bar. [2000Pie] established the phase equilibria in the Fe-Mg-Si system at 727°C and constructed the isothermal section of the phase diagram at this temperature. In this investigation the studied Fe-Mg-Si alloys were prepared by annealing the powder mixtures of Fe, Mg and Si under Ar at the atmosphere pressure at 727°C and cooled then down to room temperature. The phase equilibria in the system were determined by characterization of the phases by optical microscopy, X-ray diffraction, scanning electron microscopy and electron probe microanalysis. [2003Loe] established the phase relations in the Fe-Mg-Si system at 800 and 1000°C by identifying phases in some ternary alloys equilibrated at these temperatures and by the results of the diffusion couple experiments. Landolt-Börnstein New Series IV/11D4
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[1992Rag] presented a review on the Fe-Mg-Si phase diagram which contained the reaction scheme, the partial liquidus surface projection and the isothermal section at 1455°C. This review was supplemented by the isothermal section at 727°C [2000Pie] in the next review [2002Rag]. Binary Systems The three binary systems Fe-Mg, Fe-Si and Mg-Si are accepted from [Mas2]. Mg-Si and Fe-Mg are based on the compilations of [1984Nay] and [1985Nay], respectively. Solid Phases Based on the available data no ternary phases exist in this system. The data on the crystallography of the stable binary solid phases are summarized in Table 2. According to [2003Loe], the binary phases FeSi and FeSi2 dissolve at 800°C up to 0.03 at.% Mg and 0.05 at.% Mg, respectively. Solubility of Fe in Mg2Si at 800°C amounts to 0.3 at.%. The binary phase Fe5Si3 dissolves up to ∼0.5 and ∼1.1 at.% Mg at 800 and 1000°C, respectively. Quasibinary Systems The FeSi-Mg2Si quasibinary system was reported by [1992Rag] based on [1954Zwi]. For the respective three-phase invariant reaction the eutectic type was assumed by [1992Rag] with the eutectic composition close to Mg2Si and eutectic temperature of ∼1110°C. The eutectic type of this three-phase invariant reaction contradicts the accepted binary system Mg-Si [Mas2] with the melting point of Mg2Si being 1085°C. Therefore, the more correct type of the invariant reaction is peritectic, L + FeSi ⇌ Mg2Si, as it was assumed in the review [2000Leb] with the peritectic point being close to Mg2Si and, respectively, peritectic temperature being some above 1085°C. There is not enough experimental data to construct reliably the FeSi-Mg2Si section and it may turn out to be quasibinary only in a limited concentration and temperature range because of a possible crossing with the miscibility gap (L1+L2) extended from the binary Fe-Mg into the ternary system. Invariant Equilibria The invariant equilibria deduced from the liquidus projection [1954Zwi] are presented in the reaction scheme given in Fig. 1. The temperatures of the two ternary reactions determined from the liquidus projection [1954Zwi] were estimated in accordance with the binary Mg-Si system as 1000°C for the transition reaction U (L + FeSi ⇌ Mg2Si + βFeSi2) and < 945.6°C for the eutectic reaction E (L ⇌ (Si) + βFeSi2 + Mg2Si). In addition, the quasibinary peritectic reaction p1: L + FeSi ⇌ Mg2Si is suggested to be slightly above 1085°C based on [1954Zwi, 1992Rag] and the binary Mg-Si phase diagram [Mas2]. Characteristics of the invariant equilibria in the ternary system are described in Table 3. Liquidus, Solidus and Solvus Surfaces A partial liquidus surface represented in Fig. 2 is based on the work of [1954Zwi]. The author detected three regions of primary crystallization: FeSi, (Si) and the intermediate phase βFeSi2. Primary crystallization field of the FeSi phase occupies the largest part of the liquidus surface. The isotherms of the liquidus surface reported in Fig. 2 in the silicon rich region were constructed from [1954Zwi] and from the binary phase diagrams given by [Mas2]. The liquidus isotherm at 1454.4°C deduced from [1971Gui], [1970Gui] is also shown in Fig. 2. Because of the unknown impact of the miscibility gap [1971Gui] on the ternary system, the schematic isothermal liquidus proposed by [1954Zwi] has been rejected up to 50 mass% Si. Four binary eutectic reactions were not mentioned by [1954Zwi]: e4 at 1203°C (l ⇌ Fe2Si + FeSi) for the system Fe-Si, e8 at 637.6°C (l ⇌ (Mg) + Mg2Si) for the system Mg-Si, e7 at 649°C (l ⇌ (Mg) + αFe) and e1 at 1526°C (l1 ⇌ l2 + δFe) for the system Mg-Fe. These reactions certainly lead to the formation of one or several ternary reactions unknown to all of the authors. Because of the lack of information, no double saturation lines have been drawn from the eutectic reactions e1, e2, e5, e7 and e8 (compare question mark in Fig. 2). DOI: 10.1007/978-3-540-78644-3_15 # Springer 2008
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[1984Age, 1985Age] have determined the solubility of magnesium in binary alloys of iron with silicon in the iron rich liquid below 9.45 mass% Si at 1600°C. The results showed that the concentration of silicon used has no effect on the solubility of magnesium. However, [1996Li] revealed some solubility decrease of Mg vapor in liquid Fe at 1600°C when Si was added. The isotherm of the Mg solubility in liquid Fe containing Si at 1600°C after the results of the [1996Li] experiments is shown in Fig. 3. No solubility of iron and silicon in liquid magnesium at 800°C depending on the annealing time [1967Tka] are given here because no equilibrium state was reached for iron at the end of the measurements. [1967Rod, 1990Nor] measured liquidus curves in the magnesium rich corner. Liquidus curves were obtained at 670, 710 and 750°C from [1990Nor] and at 800°C from [1967Rod]. The solubility of iron in magnesium at 800°C [1967Rod] is not in agreement with the temperature of the liquidus curve for the Fe-Mg binary. There is a discrepancy between the value of about 0.017 at.% Fe at 800°C [1967Rod] and the value of only 4.10–3 at.% at the other temperatures [1990Nor]. At first sight, the experimental results from [1967Rod] seem to be less reliable. Therefore, the data [1990Nor] were opted as preferable and reproduced in Fig. 4. Isothermal Sections The miscibility curve deduced from [1971Gui] was reported in Figs. 2 and 5 at 1454.4°C. The tie lines in Fig. 5, illustrate the respective compositions of equilibrated liquid phases (L1 and L2). The line slopes indicate that the critical point of the miscibility gap lies near the summit of the curve. The part of the curve up to 30 mass% Si was drawn as a dashed line because of missing experimental results. The imprecise determination of the critical point position has been attributed to the nearly identical densities of the two liquids which were indiscernible with the experimental technique used in the study of [1971Gui]. In Fig. 6 the isothermal section of the phase diagram at 727°C is shown after [2002Rag] who used the work [2000Pie] to construct the section. In Fig. 7 the isothermal section at 800°C is shown after [2003Loe]. For a better view the Fe rich part of the section is shown enlarged in Fig. 8 [2003Loe]. In Fig. 9 the Fe rich part of the isothermal section at 1000°C is shown [2003Loe]. Thermodynamics On the basis of a quasiregular solution [1984Age, 1985Age] calculated the thermodynamic interaction parameters which describe the effect of silicon on the solubility of magnesium metal in liquid iron. From emf measurements of liquid Fe-Mg-Si alloys at 1350°C the activities of the components were calculated by the Darken-Wagner method [1966Lep]. The melts investigated exhibit both positive and negative deviations from Raoult’s law. Thermodynamic properties of the melts were approximated by regular solutions [1966Lep]. The effects of silicon on the solubility of Mg in liquid iron at 1600°C were studied by a vapor-molten Fe equilibration method and the interaction coefficient εMgSi = –10.90 was calculated in terms of the Margules-Wagner equation [1996Li]. Notes on Materials Properties and Applications The Fe-Mg-Si phase diagram is important for industrial production of primary magnesium and magnesium alloys because Fe and Si are the main impurities in them. Increase of impurities of Fe and Si decreases corrosion resistance of Mg base commercial alloys and decreases their plasticity. One of the main origins of the Fe and Si impurities in Mg base commercial alloys is the primary Mg. Another important origin of the Fe impurity in Mg base commercial alloys and also in primary Mg is the mild steel of crucibles where the Mg products are commonly melted during preparation [1990Nor, 2000Pie, 2002Pie]. The Fe-Mg-Si phase diagram is of great importance also for treatment of cast iron for spheroidization of graphite inclusions. The treatment is performed by addition of Mg into liquid cast iron, and because of high inclination of Mg to burning and high Mg vapor pressure it is reasonable to use alloys of Mg with high contents of Si close to Mg2Si compound and Fe. These alloys have high melting temperature [1954Zwi, 1975Dan, 1993Sol, 1996Li, 2003Kob]. Another importance of the Fe-Mg-Si phase diagram is related to technology of siliconizing surface of mild steels. This technology is used for hardening surface of steels and can be performed by dip of the steel pieces into Mg-Si melts [1993Sai, 1997Sai1, 1997Sai2, 2002Pie]. Landolt-Börnstein New Series IV/11D4
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Miscellaneous The crystalline Fe-montmorillonite (Na0.3Fe1.73+Mg0.3Si4O10(OH)2·H2O) was reported to be synthesized hydrothermally for 24 h at 100 and 200°C using the Fe-Mg-Si system. In [1998Sug] the thermoelectric material Mg2Si-FeSi2 was prepared by mechanical alloying and spark plasma sintering. Table 1.
Investigations of the Fe-Mg-Si Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1954Zwi]
DTA, optical microscopy, X-ray diffraction
Below liquidus surface (<1500°C)/0–50 at.% Fe, 0–100 at.% Mg, 0–50 at.% Si/Mg-FeSi-Si
[1966Lep]
emf measurements
1399–1350°C/ 7.3–69 at.%Fe, 2.9–58.8 at.% Mg, 11.6–69.7 at.% Si/LFe, LMg, (αFe,δFe), FeSi
[1967Rod]
Chemical analysis
700–800°C/Mg containing 0.002–0.59 mass% Si, 0.018–0.16 mass% Fe/liquid Mg
[1967Tka]
Dissolution of Si and Fe in liquid Mg/ chemical analysis of the samples
750–800°C/Mg containing 0.002–0.06 mass% Si, 0, 11 mass% Fe/liquid Mg
[1970Gui]
Sampling from the melts/chemical analysis, vapor pressure measurements
1450–1510°C, up to 15 atm/0–95%Fe, 0–95%Mg, 5–34%Si/LMg, LFe
[1971Gui]
Sampling from the melts/chemical analysis, vapor pressure measurements
1450–1510°C, up to 15 atm/0–95 mass% Fe, 0–95 mass% Mg, 5–34 mass% Si/LMg, LFe
[1984Age]
Saturation of Fe-Si melts by the Mg vapor/ determination of Mg contents by atomic absorption spectroscopy
1600°C, 2–5 atm over vapor pressure of Mg/ up to 9.45 mass% Si, up to 1.1 mass% Mg, remainder Fe/liquid Fe
[1985Age]
Calculation of the Mg activity coefficient and interaction parameter in Fe-Si alloys/theory of quasi-regular solutions
1600°C/up to ∼10 mass% Si, up to solubility of Mg in liquid Fe, remainder Fe/liquid Fe
[1990Nor]
Sampling from the Mg rich melts/chemical analysis by atomic absorption spectroscopy
680–750°C/0–0.027 at.% Fe, 0–0.75%Si, remainder Mg, liquid Mg
[1996Li]
Saturation of liquid Fe by Mg vapor/ chemical analysis by atomic absorption spectroscopy after quenching of melts
1600°C, Mg vapor pressure 0.41 bar/0.0184–0.0321 at.% Mg, (0.6–4.6)·10–4 at.% Si, remainder Fe/liquid Fe
[2000Pie]
Heating for equilibrium of the Fe-Mg-Si powder mixtures and growth of reaction zones of diffusion couples/X-ray powder diffraction, optical metallography, scanning electron microscopy, electron probe microanalysis
727°C/0–100 at.% Fe, 0–100 at.% Mg, 0–100 at.% Si/Fe, liquid Mg, Si, Mg2Si, α1, α2, FeSi, βFeSi2
[2003Loe]
Investigation of the alloys prepared by melting and diffusion couples/chemical analysis, electron dispersive microanalysis, scanning electron microscope, X-ray analysis, optical metallography
800°C, 1000°C/0–100 at.% Fe, 0–100 at.% Mg, 0–100 at.% Si/Mg2Si, α1, FeSi, Fe5Si3, FeSi2
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Fe–Mg–Si Table 2.
5
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(Mg) < 650
hP2 P63/mmc Mg
a = 320.94 c = 521.07
at 25°C [Mas2]
(Si) < 1414
cF8 Fd 3m C(diamond)
a = 534.06 a = 543.088
at 25°C [Mas2] at 20°C, 99.999 at.% purity [V-C2]
(αδFe)
cI2 Im 3m W
(δFe)(h2) < 1538 (αFe)(r) < 912
[Mas2] a = 293.15
at 1390°C [V-C2, Mas2]
a = 286.65
at 25°C [V-C2, Mas2]
(γFe)(h1) 1394–912
cF4 Fm 3m Cu
a = 364.67
at 915°C [Mas2, V-C2]
α2 ≲ 1280
cP2 Pm 3m CsCl
a = ∼281
Ordered (αδFe), ∼10 to 22 at.% Si [Mas2, V-C2]
α1 ≲ 1250
cF16 Fm 3m BiF3
a = ∼560.8
Ordered (αδFe), ∼10 to 30 at.% Si [Mas2]
FeSi < 1410
cP8 P213 FeSi
a = 451.7 ± 0.5
at 300°C [V-C2]
Fe2Si(h) 1212–1040
hP6 P 3m1 Fe2Si
a = 405.2 ± 0.2 c = 508.55 ± 0.03
[Mas2, V-C2]
Fe5Si3 1060–825
hP16 P63/mcm Mn5Si3
a = 675.6 c = 471.8
[V-C2]
βFeSi2(h) 1220–937
tP3 P4/mmm βFeSi2
a = 269.01 c = 513.4
69.5 to 73.5 at.% Si [Mas2, V-C2]
αFeSi2 (r) < 982
oC48 Cmca αFeSi2
a = 986.3 ± 0.7 b = 779.1 ± 0.6 c = 783.3 ± 0.6
[V-C2]
Mg2Si < 1085
cF12 Fm 3m CaF2
a = 634.7 ± 0.4
[V-C2]
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6 Table 3.
Fe–Mg–Si Invariant Equilibria
Reaction
T [°C]
Type
Phase
Composition (at.%) Fe
Mg
Si
L + FeSi ⇌ Mg2Si
> 1085
p1
L FeSi Mg2Si
∼0 ∼50 ∼0.3
∼66.7 ∼0.03 ∼66.7
∼33.3 ∼50 ∼33.3
L + FeSi ⇌ Mg2Si + βFeSi2
ca. 1000
U
L FeSi Mg2Si βFeSi2
∼3 ∼50 ∼0.3 ∼33.3
∼42 ∼0.03 ∼66.7 ∼0.05
∼55 ∼50 ∼33.3 ∼66.7
L ⇌ (Si) + βFeSi2 + Mg2Si
< 945.6
E
L (Si) βFeSi2 Mg2Si
∼2 ∼0 ∼33.3 ∼0.3
∼43 ∼0 ∼0.05 ∼66.7
∼55 ∼100 ∼66.7 ∼33.3
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Fig. 1. Fe-Mg-Si.
Partial reaction scheme
Fe–Mg–Si
Landolt-Börnstein New Series IV/11D4
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Fig. 2. Fe-Mg-Si.
Fe–Mg–Si
Tentative partial liquidus surface
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Fig. 3. Fe-Mg-Si.
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The effect of Si on the Mg solubility in liquid iron at 1600°C
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Fig. 4. Fe-Mg-Si.
Fe–Mg–Si
Joint solubility of Fe and Si in liquid Mg at 670, 710 and 750°C
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Fig. 5. Fe-Mg-Si.
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11
Isothermal section at 1454°C
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Fig. 6. Fe-Mg-Si.
Fe–Mg–Si
Isothermal section at 727°C
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Fig. 7. Fe-Mg-Si.
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Isothermal section at 800°C
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Fig. 8. Fe-Mg-Si.
Fe–Mg–Si
Fe rich part of the isothermal section at 800°C
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Fig. 9. Fe-Mg-Si.
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Fe rich part of the isothermal section at 1000°C
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Fe–Mg–Si
References [1954Zwi] Zwicker, V., “Reactions of Some Magnesium - Alloys with Cast Iron Melts” (in German), Z. Metallkd., 45, 31–35 (1954) (Phase Diagram, Experimental, *, #, 13) [1966Lep] Lepinskikh, B.M., Sryvalin, I.T., Tikhomorov, A.A., “Thermodynamic Properties of IronSilicon-Magnesium Alloys”, Russ. J. Phys. Chem., 40(7), 840–842 (1966) (Thermodyn., 9) [1967Rod] Rodyakin, V.V., Andreyev, A.E., Tkalich, V.S., “On the Solubility of Iron in Magnesium”, Russ. Metall., 23–24 (1967) abridged translated from Izv. Akad. Nauk. SSSR, Met., (3), 63–64 (1967) (Experimental, 1) [1967Tka] Tkalich, V.S., Andreev, A.E., Rodyakin, V.V., “The Effect of the Si on the Solubility of Fe in Mg” (in Russian), Vses. Nauchn. Issled. i Proekt. Inst. Titana, (1), 190–195 (1967) (Experimental, 8) [1970Gui] Guichelaar, P.J., Trojan, P.K., McCluhan, T., Flin, R.A., “New Techniques for the Vapour Pressure Measurement Applied to the Fe-Si-Mg System”, J. Metals, 22(12), 19A (1970) (Experimental, Abstract, 0) [1971Gui] Guichelaar, P.J., Trojan, P.K., Mccluhan, T., Flinn, R.A., “A New Techniques for Vapor Pressure Measurement Applied to the Fe-Si-Mg System”, Metal. Trans., 2(12), 3305–3313 (1971) (Experimental, Phase Diagram, *, #, 16) [1975Dan] Dang-Moon, W., Soo-Young, K., “Study on the Relation Between Graphite Spheroidization and Amount of Additional Fe-Si-Mg in Spheroidal Graphite Cast Iron” (in Korean), J. Korean Inst. Met., 13(4), 407–414 (1975) (Morphology, Experimental, 12) [1984Age] Ageev, Yu.A., Archugov, S.A., “Solubility of Mg in Some Molten Binary Iron-Base Alloys”, Russ. Metall., 72–74 (1984), translated from Izv. Akad. Nauk. SSSR, Met., (3), 78–80 (1984) (Experimental, Thermodyn., Theory, 6) [1984Nay] Nayeb-Hashemi, A.A., Clark, J.B., “The Magnesium-Silicon System”, Bull. Alloy Phase Diagrams, 5(6), 584–592 (1984) (Phase Diagram, Thermodyn., Review, 49) [1985Age] Ageev, Yu.A., Archugov, S.A., “Solubility of Alkaline Earth Metals in Liquid Iron and its Alloys”, Russ. J. Phys. Chem., 59(4), 488–490 (1985), translated from Zh. Fiz. Khim., 59, 838–841 (1985) (Experimental, Theory, Thermodyn., 13) [1985Nay] Nayeb-Hashemi, A.A., Clark, J.B., “The Iron-Magnesium System”, Bull. Alloy Phase Diagrams, 6(3), 235–238 (1985) (Phase Diagram, Thermodyn., Review, 27) [1990Nor] Normann, H.H., Thoresen, H., Tibballs, J.E., Simensen, C.J., “Impurities and Mg - Base Phase Diagrams” in “Advanced Aluminium and Magnesium Alloys”, Proc. Conf. Light Metals, Amsterdam, 823–828 (1990) (Experimental, *, 7) [1992Rag] Raghavan, V., “The Fe-Mg-Si (Iron-Magnesium-Silicon) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Institut of Metals, Calcutta, 6B, 942–947 (1992) (Crys. Structure, Phase Diagram, Phase Relations, Review, 4) [1993Sai] Saikawa, H., Matsumoto, M., Minegishi, T., Morozumi, S., “Mineralization of the Surface of Iron Alloys by Siliconizing in a Molten Magnesium Bath Containing Silicon Compounds”, J. Mater. Sci. Lett., 12, 157–159 (1993) (Mechan. Prop., Morphology, Experimental, 0) [1993Sol] Solntsev, K.A., Shifrin, V.D., Miroshnichenko, O.N., Serkhovets, S.I., “A Study of HighSilicon Structural Constituents of Magnesium Iron Cast in Metal Molds”, Russ. Metall., (4), 89–92 (1993) (Morphology, Experimental, 5) [1996Li] Li, X., Song, B., Han, Q., “Thermodynamic Properties of Liquid Fe-Mg-Al and Fe-Mg-Si Dilute Ternary Solutions”, J. Phase Equilib., 17(1), 21–23 (1996) (Experimental, Thermodyn., 11) [1997Sai1] Saikawa, H., Ono, T., Minegishi, T., Morozumi, S., “Siliconizing of Iron by Molten Mg-Si Alloy Bath”, Tetsu to Hagane (Iron Steel Inst. Jpn.), 83(7), 437–441 (1997) (Experimental, Phase Relations, 13) [1997Sai2] Saikawa, H., Ono, T., Minegishi, T., Morozumi, S., “Siliconizing of Iron and Steel by Molten Magnesium Bath with Silicon Carbide Powder”, Tetsu to Hagane (Iron Steel Inst. Jpn.), 83(7), 442–447 (1997) (Experimental, Kinetics, Morphology, Phase Relations, 7) [1998Sug] Sugiyama, A., Kobayashi, K., Ozaki, K., Nishio, T., Matsumoto, A., “Preparation of Functionally Graded Mg2Si-FeSi2 Thermoelectric Material by Mechanical Alloying Pulsed Current DOI: 10.1007/978-3-540-78644-3_15 # Springer 2008
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[1999Nag]
[2000Leb]
[2000Pie]
[2002Pie]
[2002Rag] [2003Kob]
[2003Loe]
[Mas2] [V-C2]
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Sintering Process” (in Japanese), J. Jpn. Inst. Met., 62(11), 1082–1087 (1998) (Morphology, Experimental, 3) Nagase, T., Iwasaki, T., Ebina, T., Hayashi, H., Onodera, Y., Dutta, N.C., “Hydrothermal Synthesis of Fe-Montmorillonite in Si-Fe-Mg System”, Chem. Lett. (4), 303–304 (1999) (Morphology, Crys. Structure, Experimental, 9) Lebrun, N., Baetzner, C., Stamou, A., Robinson, J., “Iron-Magnesium-Silicon”, MSIT Ternary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 10.15873.2.20, (2000) (Crys. Structure, Phase Diagram, Phase Equilibria, Assessment, 11) Pierre, D., Peronnet, M., Bosselet, F., Viala, J.C., Bouix, J., “Solid-Liquid Phase Equilibria at 727°C in the Ternary System Fe-Mg-Si”, J. Phase Equilib., 21(1), 78–86 (2000) (Experimental, Kinetics, Phase Relations, 21) Pierre, D., Peronnet, M., Bosselet, F., Viala, J.C., Bouix, J., “Chemical Interaction Between Mild Steel and Liquid Mg-Si Alloys”, Mater. Sci. Eng. B, 94, 186–195 (2002) (Morphology, Phase Diagram, Phase Relations, Calculation, Experimental, Kinetics, 26) Raghavan, V., “Fe-Mg-Si (Iron-Magnesium-Silicon)”, J. Phase Equilib., 23(2), 175–176 (2002) (Phase Relations, Review, 6) Kobayashi T., Maruyama T., “Thermal Decomposition Behaviour of Expandable Pattern Including Blended Metal or Alloy Powder in Evaporative Pattern Casting Process of Cast Iron”, Mater. Trans., 44(11), 2396–2403, (2003) (Experimental, Kinetics, 13) Loeffler, F., “Investigations in the Ternary Systems Fe-Si-Mg and Fe-Si-Ti: Phase Equilibria and Mechanical Behaviour of Selected Alloys” (in German), Fortschritt-Berichte VDI. Reihe 5: Grund- und Werkstoffe/Kunststoffe, 5(676), 1–106 (2003) (Phase Relations, Phase Diagram, Experimental, Interface Phenomena, 95) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Fe–Mn–N
1
Iron – Manganese – Nitrogen Lazar Rokhlin, Evgeniya Lysova
Introduction The Fe-Mn-N system has raised much interest because Mn is a key element in the preparation of High Nitrogen Steels (HNS). Indeed, the presence of Mn in Fe increases the nitrogen solubility in the steel and the increase is not limited by the manganese nitride precipitation. Different experimental works were reviewed by [1984Rag, 1987Rag, 1993Rag] and thermodynamically assessed by [1993Qiu2]. In the detailed review [1987Rag] the partial isothermal sections of the Fe-Mn-N phase diagram at 500 and 550°C, the solubility of N in the liquid Fe-Mn alloys, the solubility of N in Fe-Mn solid austenite and Fe-Mn ferrite were presented. The isothermal sections were reproduced from the experimental work of [1975Raw, 1984Bur]. The N solubility in liquid Fe-Mn was assessed by [1993Qiu2] from the experimental determinations of [1960Peh, 1961Bee, 1961Dod, 1974Gri, 1982Ish, 1984Lee, 1986Wad, 2001Kim]; the N solubility in the solid Fe-Mn phases: austenite and ferrite, was assessed from [1963Mor, 1963Sch1, 1963Sch2, 1975Gra, 1983Ko] and [1962Enr, 1973Pip, 1975Raw, 1980Bra], respectively. In the next review, [1993Rag] updated [1987Rag] using the experimental work of [1991Sat] on the N solubility in the Fe-Mn alloys. Moreover, a new isothermal section of the Fe-Mn-N phase diagram at 700°C was presented. A thermodynamic assessment of the ternary system was presented by [1993Qiu2] which reproduces with a rather agreement the experimental data about the solubilities and which takes into account the nitrogen pressures. The main experimental investigations of the Fe-Mn-N phase diagram are presented in Table 1. Binary Systems The Fe-Mn system has been recently reassessed and the Calphad description was updated by [2004Wit]. This new reevaluation of the phase equilibria leads to consistently better fits to the available experimental data. Consequently, the Fe-Mn system is accepted from [2004Wit]. There is however, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have been written as L ⇌ (δMn) + (γMn,γFe). The Fe-N phase diagram in the solid state is accepted from the review of [1987Wri]. The Calphad assessment carried out by [1991Fri] and justified by the model proposed by [1994Fer] gives an insight on the phase equilibria under high nitrogen pressures. The Mn-N system at high temperatures is accepted from the reevaluation of [1993Qiu1], which takes into account the nitrogen pressure. The nitrogen solubility in pure Mn may be expressed by [1967Bur]: log10 (mass% N) = – 1.457 + (3010/T ) + 0.5 log10 (pN2/bar) (1246–1600°C). Solid Phases No ternary compound was reported in the Fe-Mn-N system. Other solid phases are given in Table 2. The nitrides Mn4N and Fe4N have the same cP5 cubic structure with parameters close enough to suggest the formation of a continuous solid solution Mn4(1–x)Fe4xN. Actually, a complete solid solution can not be observed because of the mismatch in the stability domains of both nitrides. Mn4N reacts with Fe4N by giving Fe and Mn nitrides with higher nitrogen content. In [1959Juz] the lattice parameter measurements were extended up to x = 0.56. At higher Fe content dissolved in Mn4N, the lattice parameter remains constant. However, the Curie point of Mn4N continues to decrease with addition of Fe up to the highest values x = 0.72 used in the experiments. The solubility of Mn in Fe4N was reported by [2004Pad].
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Liquidus Surface Using thermodynamic calculation, [1993Qiu2] presents the partial liquidus surface of the Fe-Mn-N phase diagram near the Fe corner for different values of the N2 pressure. Projection of the liquidus surface with the respective isothermal and isobaric curves is shown in Fig. 1. Isothermal Sections The partial isothermal section of the Fe-Mn-N phase diagram at 500°C and 0.1 MPa of N2 pressure is presented in Fig. 2 after [1987Rag], mainly from experimental data [1984Bur]. The homogeneity ranges of the binary Mn-N phases are taken from [Mas2]. Figure 3 shows the solubility of nitrogen in Fe-Mn ferrite at 550°C, experimentally determined by [1975Raw] and reproduced by [1987Rag]. Figures 4 and 5 show the nitrogen solubility at various temperatures in the Fe-Mn austenite and at 1550°C in the liquid phase, respectively [1987Rag]. The solubility curves for austenite and liquid at 1550°C are well reproduced by thermodynamic calculations of [1993Qiu2]. The nitrogen solubility in liquid alloys deviates from the Sievert’s law above 0.1 MPa of N2 pressure [1963Sch2, 1990Siw, 1991Sat]. However, it may be shown that Sievert’s law is better obeyed with nitrogen fugacities rather than with nitrogen pressures. Temperature – Composition Sections The temperature-composition section Mn4N-Fe4N, calculated under 0.1 MPa of N2 pressure [1993Qiu2] is presented in Fig. 6. Thermodynamics The interaction coefficients between Mn and N in liquid iron at 1600°C, (< 20 mass% Mn) were first evaluated as eN(Mn) = {∂ log10 fN/∂ (mass% Mn)} = – 0.020 at 1600°C by [1960Mae, 1960Peh, 1982Ish] and as –0.023 at 1400–1650°C by [1961Bee, 1991Sat], with fN = (mass% N in pure Fe)/(mass% N in alloy). A temperature dependent expression, usable around 1600°C was proposed by [1977Lin]: in the iron rich liquid alloys: eN(Mn) = 0.022 – 73/T = – 0.016 at 1600°C Around 1600°C, eN(Mn)< 0: Mn increases the nitrogen solubility in iron whereas eN(Fe) > 0 [1961Dod]: Fe decreases the nitrogen solubility in manganese. A statistical model of the nitrogen-metal interaction in iron based liquid alloys has been developed by [1994Tan]. The interaction coefficient between Mn and N in γ austenite has been investigated by [1963Mor], which propose, eN(Mn) = –0.034, –0.037 and –0.041 at 1250, 1150 and 1050°C, respectively. By using mole fractions instead of mass%, the interaction parameter in the γ austenite may be expressed by: εN(Mn) = {∂ lnγN/∂ xMn} = (8.2 – 2100/T ) ± 0.5 [1975Gra] in the temperature range 900–1100°C with γN = (xN in pure Fe)/(xN in alloy). This last expression agrees with the evaluation of [1975Kik] εN(Mn) = – 9.0 at 1100°C and with the experimental measurements of [1983Ko]: εN(Mn) = – 10.34 at 1000°C. Mn up to 0.75 mass% seems to have little effect on the nitrogen solubility in (αFe) [1962Enr], an observation which is confirmed by the low value of the interaction parameter εN(Mn) = + 0.21 proposed in the temperature range 450–550°C by [1975Gra]. Notes on Materials Properties and Applications The alloys (Mn,Fe)4N were reported to be ferrimagnetic [1959Juz]. However, a ferromagnetic order was observed for MnFe3N and calculations carried out by [2004Pad] showed that a high spin ferromagnetic state is the stable phase at low temperatures. Aiming to investigate the precipitation phenomenon in the Fe-Mn-N alloys the change of internal friction in them with increasing temperature and strength of coercivity field were measured [1962Enr, 1966Nac1, 1966Nac2, 2002Bli].
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[1974Ste] investigated the internal friction in the Fe-Mn-N alloys containing 0–2 mass% Mn and 0.0007– 0.0011 mass% N in limits from –53 to 252°C at ∼1 Hz using the torsion pendulum method. The obtained results give a procedure for determining the nitrogen content in the Fe-Mn-N alloys. [1975Tsu] investigated the strain aging of Fe-Mn-N polycrystalline specimens and showed, that strain aging takes place in two stages corresponding to different change of the yield point with ageing time. Authors explained the yield point behavior during the strain aging of the alloys by an increase of the average distance between pinning points in screw dislocations. [1976Sal] studied the low frequency friction spectrum in Fe-Mn-N alloys and stress relaxation at constant strain in limits of the –4 to 150°C temperature range. The obtained results are connected with motion of single nitrogen atoms around single manganese atoms, and two types of motion of single nitrogen atoms around a manganese atom pair. [1998Gav] examined influence of 0.2 mass% N on vibration damping and mechanical properties of Fe13.7Mn (mass%) alloy. The damping measurements were conducted at the torsion deformation amplitude of (0.2–8.0)·10–4 at frequency of 1–2 Hz. Presence of nitrogen resulted in increase of damping and strength properties of the alloy, especially the ultimate strength. The favorable nitrogen effects are attributed to the strengthening of austenite and higher mobility of the austenite/martensite. [1999Ari] has investigated the effect of 0.26 mass% N on the shape memory behavior of a Fe-Mn based shape memory alloy. Addition of nitrogen increased the strength of the alloy and suppressed the shape memory effect by approximately to 80%. [2002Lee] has examined the effects of 0.015–0.05 mass% N on the damping capacity and austenite/martensite transformation in the Fe-16Mn (mass%) alloy. Increase in Ms (Martensite start temperature) with N content was pointed out and explained by decrease in stacking fault energy (SFE) of austenite (γ). The damping capacity of the alloys decreases when increasing nitrogen content. [2006Che] studied the electronic structure and magnetism of Fe4–xMnxN compounds by a periodic quantum-mechanical calculation based on density functional theory. The results showed that ferromagnetic ordered phase is stable when Fe is substituted by Mn on cube corner sites, whereas the antiferromagnetic phase is energetically favored when Mn substitutes for Fe on face-centered sites. Details of the properties investigations are presented in Table 3. Miscellaneous [1967Cho1] measured the rates of N absorption into liquid iron containing 0–4.3 mass% Mn at 1600°C and [1979Cho] investigated the kinetics of nitrogen desorption in liquid Fe-Mn alloys with 1.1–6.8 mass% Mn at 1600°C and 0.1 mm Hg pressure. [1968Roy] investigated the high temperature strengthening mechanism in a dilute Fe-∼1.5Mn-0.01N (mass%) alloy by using the Mössbauer effect. The Mössbauer spectrum observed was interpreted as evidence of the Fe-Mn-N chains formation in the Fe base solid solution. In [1973Kun] effect of Mn on the diffusion coefficient of nitrogen in Fe-Mn melted alloys at 1550–1700°C was investigated. Actually increase of Mn contents up to about 15 mass% did not result in any change of the N diffusion coefficient. However, [1981Ers] reported about decrease of the nitrogen mobility in melted Fe-Mn alloys at 1600°C with increasing Mn contents up to 8 mass% Mn. Prolonged mechanical milling of a mixture of Fe and Mn nitrides [1989Roc] leads to oversaturated solid solutions and metastable phases. [1978Sos] established the linear increase of the γ solid solution lattice parameter in the Fe-18Mn (mass%) alloys with increasing N contents up to about 2.3 mass%. [1982Gar] presented the detailed diffraction data used for identification of a nitrogen rich orthorhombic phase which is considered as a M3N type nitride formed during decomposition of the Fe-Mn-N alloys with low N content. In [1992Vol] structure and breakdown of the Fe-Mn alloys prepared by hot extrusion of powders were studied. The alloys contained 4 to 32 mass% Mn and 0.0002–0.01915 mass% N, because the powders were obtained by atomizing in the nitrogen atmosphere. Distribution of the N atoms and other gas impurities on the grain boundaries of the alloys and dependence of it on temperature were established. The change of the breakdown mechanism of the alloys with increasing temperature is attributed to the change of the impurity distribution. Landolt-Börnstein New Series IV/11D4
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In [1999Sum] neutron-spectroscopy investigation of the Fe-Mn-N alloys containing 17–35 mass% Mn was studied aiming to estimate interaction between Mn and N atoms in Fe-Mn solid solution. Investigation showed weak interaction between each of the atom kinds as compared with Cr-N interaction. [2000Gou] provided a realistic diffusion - precipitation model, which describes the volume diffusion of nitrogen in ferrite and the simultaneous precipitation of fine scale alloying element nitrides in the diffusion zone. Theoretical calculations were in good agreement with experimental measurements. Table 1.
Investigations of the Fe-Mn-N Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1960Mae]
Nitrogen solubility in liquid alloys, sampling method
1500–∼1700°C, < 4 mass% Mn < 0.1 MPa of N2 pressure
[1960Peh]
Nitrogen solubility in liquid alloys, Sievert’s method
1606°C, < 6 mass% N, 0.1 MPa of N2 pressure
[1961Dod]
Nitrogen solubility in liquid alloys, Sievert’s method
1550°C, 0–100 mass% Mn, 0.1 MPa of N2 pressure
[1961Bee]
Nitrogen solubility in liquid alloys, sampling method
1400–1650°C, 0–100 mass% Mn, 0.1 MPa of N2 pressure
[1962Enr]
Nitrogen solubility in α alloys, internal friction measurements
200–350°C, 0.15 to 0.75 mass% Mn, 0.1 MPa of N2 pressure
[1963Mor]
Nitrogen solubility in γ alloys, chemical analysis
1050–1250°C, < 6 mass% Mn, H2-N2 atmospheres (5% H2)
[1963Sch1]
Nitrogen solubility in γ alloys, sampling method
700–1200°C, < 9.5 mass% Mn, 0.1 MPa of N2 pressure
[1963Sch2]
Nitrogen solubility in γ alloys, sampling method
1000°C, 0–100 mass% Mn, 0.1 to 6.4 MPa of N2 pressure
[1966Bai]
XRD, electron diffraction, nitride precipitation investigation
450–650°C, 1.6 mass% Mn, (0.040 to 0.063) mass% N
[1966Nac2]
Nitrogen solubility in α alloys, internal friction measurements
585°C, 0.45 to 2.01 mass% Mn
[1968Nem]
Nitrogen solubility in liquid alloys, sampling method
1550°C, 26 to 76 mass% Mn, 0.1 MPa N2
[1974Gri]
Nitrogen solubility in liquid alloys, sampling method
1550–2100°C, 3.6 to 59.2 mass% Mn, 25 to 101 kPa N2
[1975Gra]
Nitrogen solubility in γ alloys, sampling method
900–1100°C, 3.4 to 45 mass% Mn, 1 to 95 kPa N2
[1975Raw]
Fe4N solubility in α alloys, sampling method
450–550°C, < 4 mass% Mn, H2-NH3 atmospheres
[1978Sos]
XRD, lattice parameters measurements
25°C, 18 mass% Mn, < 2.3 mass% N
[1982Gar]
XRD, identification of the nitrogen rich orthorhombic phase
25°C, 4.2 mass% Mn, 0.5 to 0.9 mass% N (continued)
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Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1982Ish]
Nitrogen solubility in liquid alloys, sampling method
1540–1680°C, < 9 mass% Mn, 0.1 MPa N2
[1983Ko]
Nitrogen solubility in γ alloys, sampling method
867–1265°C, < 15 at.% Mn, 0.104 MPa N2
[1984Lee]
Nitrogen solubility in liquid alloys, levitation melting method
1489–1998°C, < 15 mass% Mn, 0.1 MPa of N2 pressure
[1985Bas]
Nitrogen solubility in γ under high pressures, XRD, TEM
1000–1070°C, 7 and 18 mass% Mn, < 120 MPa N2
[1986Wad]
Nitrogen solubility in liquid alloys, levitation melting method
1489–1998°C, 3.8–15.8 mass% Mn, 0.1 MPa of N2 pressure
[1990Siw]
Nitrogen solubility in liquid alloys, sampling method
1600°C, 18 mass% Mn, 0.125 to 1.0 MPa of N2pressure
[1991Sat]
Nitrogen solubility in liquid alloys under high pressure
1600°C, 18–28 mass% Mn, < 10 MPa of N2 pressure
[1994Kun]
Nitrogen solubility in γ alloys, sampling method
1000–1200°C, 6.26 to 18.21 mass% Mn, 98.7 kPa N2 + 2.6 kPa H2
[1995Nec]
Nitrogen solubility in γ alloys under high pressure
900–1200°C, < 18 mass% Mn, 3 to 348 MPa of N2 pressure
[2001Kim]
Nitrogen solubility in liquid alloys, sampling method
1350–1550°C, 0.1 MPa of N2 pressure
[2003Gou1, 2003Gou2]
TEM, metallography, electron diffraction
570°C, 1.62 mass% Mn nitrided in bath salt; a new θ′ phase
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group Prototype
(δFe) 1538–1394
cI2 Im 3m W
γ, (γFe,γMn) < 1473
cF4 Fm 3m Cu
(γFe) 1394–912 (γMn) 1138–1100
Lattice Parameters [pm]
a = 293.15
Comments/References
dissolves up to 10.5 at.% Mn, pure Fe at 1390°C [V-C2, Mas2] continuous solid solution, dissolves up to 11.3 at.% N [Mas2]
a = 364.67
pure Fe at 915°C [Mas2, V-C2]
a = 386.0
pure Mn [Mas2] (continued)
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Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group Prototype
(αFe) < 912
cI2 Im 3m W
(δMn) 1246–1138
cI2 Im 3m W
(βMn) 1155–727 (αMn) < 727 γ′, Mn4(1–x)Fe4xN γ′, Mn4N < 890
cP20 P4132 βMn cI58 I 43m αMn
Mn2N
hP12 P6322 Mn12N5 hP3 P63/mmc -
Comments/References
a = 286.65
dissolves 3 at.% Mn and 0.4 at.% N, pure Fe at 25°C [Mas2, V-C]
a = 308.0
dissolves up to 8.85 at.% Fe and ∼2 at.% N, pure Mn [Mas2]
a = 631.2
dissolves up to 34 at.% Fe and ∼4 at.% N, pure Mn [Mas2]
a = 891.26
dissolves 29.7 at.% Fe and ∼0.5 at.% N, pure Mn at 25°C[Mas2]
cP5 Pm 3m Fe4N
γ′, Fe4N < 680 ζ Mn12N5
Lattice Parameters [pm]
solid solution stable at both ends (x < 0.3 and x∼1 at 900°C [1984Bur] a = 383.4 to 387.7
17.4 to 21.6 at.% N [Mas2], [1987Rag]
a = 387.6 a = 378.7
at 20.45 at.% N 19.3 to 20 at.% N [Mas2]
-
∼ 29 at.% N [Mas2]
a = 281.6 c = 453.4
32.6 at.% N [Mas2] at 32.6 at.% N [V-C2]
Mn2N
oP12 Pbcn PbO2
η, Mn6N4 ≲ 710
tI* I4/mmm -
θ, Mn6N5 ≲ 580
hP* P6/mmm
θ′
tP* P42/mmc -
∼33 at.% N [Mas2]
a = 421.1 c = 409.5 a = 420.5 c = 404.2 a = 421.1 c = 414.5 a = 287.6 c = 575.2
from ∼38 to 41 at.% N [Mas2] [1987Rag] at 39.1 at.% N [V-C2] from ∼ 45.8 to 49 at.% N [Mas2] [2003Gou1] Metastable phase [2003Gou1]
(continued)
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Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group Prototype
εFe3N < 580
hP3 P63 22 Fe3N
Fe2(1 –x)Mn2x N < 500
oP12 Pbcn Fe2N
7
Lattice Parameters [pm]
Comments/References
from 15.8 to 33.2 at.% N [1987Rag] εFe5.23N
a = 259.7 c = 434.1 a = 278.7 c = 444.8
εFe2N x < 0.30 [1984Bur]
a = 551.2 b = 482.0 c = 441.6
at 25°C for x = 0 [1987Wri]
α′, Fe16N2
tI* I4/mmm -
a = 572.0 c = 629.2
Ordered fcc structure, metastable [1987Rag]
α′, Fe-N (Martensite)
cI2 Im 3m W bct -
-
from 0 to 2.5 at.% N [Mas2]
a = 284.4 c = 310.0
from 2.7 to 9.5 at.% N [Mas2]
Table 3.
Investigations of the Fe-Mn-N Materials Properties
Reference
Method/Experimental Technique
Type of Property/Conditions
[1959Juz]
XRD, Curie temperature, saturation magnetization
< 400°C, magnetic properties of the solid solution (Fe1–xMnx)4N
[1962Enr]
Internal friction measurements
Kinetics of nitrogen precipitation from (αFe).
[1966Bai]
Normal ageing tests, tensile properties and hardness
450–600°C, 1.0 to 1.6 mass% Mn, 0.008 to 0.063 mass% N.
[1966Nac1]
Damping behavior
< 200°C, < 1.5 mass% Mn, < 0.07 mass% N
[1966Nac2]
Damping behavior, Mn precipitation
585°C, < 2 mass% Mn, < 0.08 mass% N
[1967Cho2]
Creep resistance under stress change
450°C, 0.24 to 0.50 mass% Mn, 0.034 to 0.046 mass% N
[1974Ste]
Nitrogen determination by internal friction measurements
– 52 to + 252°C, < 2 mass% Mn, < 0.001 mass% N
[1975Tsu]
Electrical resistivity and recovery of yield point
Ageing of polycrystalline Fe-Mn-N alloys
[1976Sal]
Internal friction, stress relaxation at constant strain
– 40 to + 150°C, 3.6 mass% Mn (continued)
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Fe–Mn–N
Reference
Method/Experimental Technique
Type of Property/Conditions
[1981Ers]
Nitrogen diffusion in liquid alloy, volumetric method
1600°C, < 7 mass% Mn, diffusion coefficient
[1989Roc]
XRD, SEM, Mössbauer
Mechanical alloying Fe4N + Mn mixtures, annealing at 190°C
[1998Gav]
XRD, internal friction, nitrogen influence on damping behavior,
< 497°C, 13.7 mass% Mn, 0.2 mass% N, mechanical properties
[1999Ari]
Shape memory effect, thermodynamic modeling
0.26 mass% N, effect of nitrogen on shape memory effect
[2002Bli]
Amplitude - dependent internal friction technique
Dependence of internal friction on strain amplitude
[2002Lee]
XRD, dilatometry, damping behavior
16 mass% Mn, < 0.05 mass% N, γ → ε martensitic transformation
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Fig. 1. Fe-Mn-N.
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The calculated liquidus projection together with isobaric lines (dashed curves)
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Fig. 2. Fe-Mn-N.
Fe–Mn–N
Partial isothermal section at 500°C
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Fig. 3. Fe-Mn-N.
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Solubility of N in Fe-Mn ferrite at 550°C
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Fig. 4. Fe-Mn-N.
Fe–Mn–N
Solubility of N in Fe-Mn austenite at 1000 and 1200°C
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Fig. 5. Fe-Mn-N.
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Solubility of N in liquid Fe-Mn alloys under 0.1 MPa of N2 pressure
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Fig. 6. Fe-Mn-N.
Fe–Mn–N
The temperature - composition section Mn4N - Fe4N
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Fe–Mn–N References [1959Juz]
[1960Mae]
[1960Peh]
[1961Bee] [1961Dod] [1962Enr] [1963Mor]
[1963Sch1]
[1963Sch2]
[1966Bai] [1966Nac1]
[1966Nac2]
[1967Bur] [1967Cho1]
[1967Cho2] [1968Nem]
[1968Roy]
[1973Kun]
[1973Pip]
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Juza, R., Deneke, K., Puff, H., “Ferromagnetism of Solid Solutions of Mn4N with Cr, Fe, and Ni” (in German), Z. Elektrochem., 63, 551–557 (1959) (Experimental, Thermodyn., Magn. Prop., 11) Maekawa, S., Nakagawa, Y., “Solubility of Nitrogen in Liquid Iron and Iron Alloys. I. Solubility of Nitrogen in Liquid Iron and Effect of Carbon, Silicon and Manganese on the Solubility” (in Japanese), Tetsu to Hagane, 46, 748–753 (1960) (Experimental, Phase Relations, Thermodyn., 14) Pehlke, R.D., Elliott, J.F., “Solubility of Nitrogen in Liquid Iron Alloys. I. Thermodynamics”, Trans. AIME, 218, 1088–1101 (1960) (Experimental, Thermodyn., Phase Relations, 32) Beer, S.Z., “Solubility of Nitrogen in Molten Iron-Manganese Alloys”, Trans. AIME, 221, 2–8 (1961) (Electrical Prop., Experimental, Phase Relations, 14) Dodd, R.A., Gokcen, N.A., “Solubility of Nitrogen in Liquid-Iron Manganese Alloys”, Trans. AIME, 221, 233–236 (1961) (Experimental, Phase Relations, Theory, Thermodyn., 9) Enrietto, J.F., “The Solubility and Precipitation of Nitrides in α-Iron Containing Manganese”, Trans. AIME, 224, 43–48 (1962) (Experimental, Phase Relations, Morphology, 12) Mori, T., Shimmyo, K., Ichise, E., Morooka, A., “Effects of Alloying Elements on the Activity of Nitrogen in Austenite. I. Effects of Chromium and Manganese”, Mem. Fac. Eng. Kyoto Univ., 25, 164–174 (1963) (Experimental, Calculation, 11) Schenck, H., Frohberg, M.G., Reinders, F., “Contribution to the Study of the Solubility of N in Fe Alloys over the Temperature Range from 700 to 1200°C” (in German), Stahl Eisen, 83(2), 93–99 (1963) (Experimental, Phase Relations, 26) Schenck, H., Frohberg, M.G., Kranz, W., “The Solubility of N in Fe-Mn Alloys at 1000°C and Pressures up to 65 Atmospheres” (in German), Arch. Eisenhuettenwes., 34, 825–830 (1963) (Experimental, Theory, Thermodyn., 33) Baird, J.D., “Precipitation of Nitrides in Iron-Manganese-Nitrogen Alloys”, J. Iron Steel Inst., 204(11), 1122–1130 (1966) (Experimental, Phase Relations, 18) Nacken, M., Kuhlmann, U., “Effect of Manganese Content on the Damping Behaviour of Nitrogen-Containing Iron Alloys” (in German), Arch. Eisenhuettenwes., 37, 235–243 (1966) (Kinetics, Thermodyn., 20) Nacken, M., Kuhlmann, U., “Effect of Manganese Content on the Precipitation Behaviour of Nitrogen-Containing Ferrous Alloys” (in German), Arch. Eisenhuettenwes., 37, 331–339 (1966) (Kinetics, Thermodyn., 23) Burylev, B.P., “Solubility of Gases in Molten Mn-base Alloys” (in Russian), Izv. Vyssh. Ucheb. Zaved. Chern. Met., (4), 5 (1967) (Theory, Calculation, 7) Choh, T., Inouye, M., “On the Rate of Absorption of Nitrogen in Liquid Iron and Iron Alloys, Containing Carbon, Silicon, Manganese and Chromium” (in Japanese), Tetsu to Hagane, 53(12), 1393–1406 (1967) (Experimental, Kinetics, 26) Choubey, R., Nijhawan, B.R., “Effects of Stress Changes on Creep Resistance of Fe-Mn-N Alloys”, J. Iron Steel Inst, 205(5), 554–555 (1967) (Mechan. Prop., Experimental, 6) Nemchenko, V.P., Malkin, I.P., Popel, S.I., “To the Thermodynamic of the Solution of the Nitrogen in the Melts of Fe-Cr, Fe-Mn and Fe-Cr-Mn” (in Russian), Izv. Vyss. Uchebn. Zaved., Chern. Metall., (12), 5–8 (1968) (Calculation, Experimental, Thermodyn., 9) Roy, R.B., Solly, B., Wappling, R., “An Investigation of High Temperature Strengthening Mechanism in a Dilute Fe-Mn-N Alloy by Means of Mössbauer Effect” (in German), Z. Metallkd., 59(7), 563–566 (1968) (Experimental, Mechan. Prop., 7) Kunze, H.D., “Effect of the Elements Chromium, Manganese, Cobalt, Nickel, Molybdenum and Tungsten on the Diffusion of Nitrogen in Liquid Iron Alloys” (in German), Arch. Eisenhuettenwes., 44(2), 71–80 (1973) (Experimental, Kinetics, Thermodyn., 50) Pipkin, N.J., Roberts, W., Speirs, D.L., Grieveson, P., Jack, K.H., “Effect of Substitutional Alloying Elements on the Activity Coefficients and Behaviour of Interstitial Solutes in Iron” in “Chemical Metallurgy of Iron and Steel”, Proc. Int. Symp. Metallurgical Chemistry. MSIT®
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16
[1974Gri]
[1974Ste] [1975Gra]
[1975Kik] [1975Raw]
[1975Tsu] [1976Sal]
[1977Lin] [1978Sos]
[1979Cho]
[1980Bra] [1981Ers]
[1982Gar] [1982Ish] [1983Ko] [1984Bur]
[1984Lee] [1984Rag] [1985Bas]
Fe–Mn–N Applications in Ferrous Metallurgy, Iron and Steel Inst., London, 351–352 (1973) (Morphology, Thermodyn., 7) Grigorenko, G.M., Pomarin, Yu.M., Lakomsky, V.I., Sherevera, A.V., “The Influence of Temperature on the Solubility of Nitrogen in Iron-Manganese Melts”, Russ. Metall., (6), 9–13 (1974), translated from Izv. Akad. Nauk SSSR, Met., (6), 11–15 (1974) (Kinetics, Phase Relations, Thermodyn., 20) Stephenson, E.T., “Determination of Total Nitrogen Solution Fe-Mn-N Alloys by Internal Friction”, Metall. Trans., 5(9), 1987–1992 (1974) (Experimental, Calculation, 24) Grabke, H.J., Iyer, S.K., Srinivasan, S.R., “The Solubility of Nitrogen in Austenitic IronManganese and Iron-Chromium Alloys”, Z. Metallkd., 66, 286–292 (1975) (Experimental, Phase Relations, Thermodyn.,13) Kikuchi, M., Tanaka, R., “Activity of Nitrogen in Austenitic Steels” (in Japanese), Tetsu to Hagane, 61(13), 2892–2903 (1975) (Phase Relations, Thermodyn., 46) Rawlings, R., Hatherley, P.G., “Iron-Manganese-Nitrogen Ferrite: The Activity of Nitrogen and the Solubility of Manganese Nitrides”, Met. Sci., 9, 97–103 (1975) (Crys. Structure, Phase Relations, Thermodyn., 34) Tsunoyama, K., Suzuki, H., “Strain Aging of Fe-Mn-N Alloy” (in Japanese), J. Jpn. Inst. Met., 39(8), 836–840 (1975) (Experimental, Mechan. Prop., 21) Salzbrenner, R.J., Carpenter, C.H., “Anelastic Relaxations in the Fe-Mn-N System”, Physica Status Solidi A, 36(1), 117–126 (1976) (Experimental, Phase Relations, Mechan. Prop., Calculation, 26) Lin, R.Y., Chang, Y.A., “Activity-Coefficient of Nitrogen in Binary-Liquid Metal-Alloys”, Metall. Trans. B, 8B(2), 293–300 (1977) (Calculation, Thermodyn., Review, 72) Soshnikov, V.I., Omelchenko, A.V., Georgieva, I.V., Bashchenko, A.P., “Nitrogen Influence on Lattice - Parameter of Iron - Manganese Alloy Austenite” (in Russian), Fiz. Met. Metalloved., 45(5), 1120 (1978) (Crys. Structure, Experimental, 7) Chon, T., Moritani, T., Inouye, M., “Kinetics of Nitrogen Desorption of Liquid - Iron, Liquid Fe-Mn and Fe-Cu Alloys under Reduced Pressures”, Trans. Iron Steel Inst. Jpn., 19(4), 221–230 (1979) (Experimental, Kinetics, 48) Brandes, E.A., Flint, R.F., “Mn-Fe-N” in “Manganese Phase Diagrams”, Manganese Centre, Paris, 113 (1980) (Phase Diagram, Review, 2) Ershov, G.S., Kasatkin, A.A., “Influence of Alloying Elements on the Diffusion of Nitrogen in Liquid Iron”, Russ. Metall., (3), 24–27 (1981) (Experimental, Transport Phenomena, Thermodyn., 13) Garwood, R.D., “Cementite - Type (D011) Orthorhombic Phase in Tempered Fe-Mn-N Alloys”, Met. Sci., 16(12), 587–590 (1982) (Crys. Structure, Experimental, 12) Ishii, F., Banya, S., Fuwa, T., “Solubility of Nitrogen in Liquid Iron Alloys”, Tetsu to Hagane, 68(10), 1551–1559 (1982) (Experimental, Phase Relations, Thermodyn., 43) Ko, C., McLellan, R.B., “Thermodynamics of Ternary Nitrogen Austenites”, Acta Metall., 31(11), 1821–1827 (1983) (Experimental, Phase Relations, Thermodyn., 26) Burdese, A., Firrao, D., Rolando, P., Rosso, M., “About the Iron-Manganese-Nitrogen System” (in Italian), Chim. Ind., 66(7–8), 456–460 (1984) (Experimental, Phase Relations, Thermodyn., 44) Lee, S-W., Young H-Y., “Effect of Mn on the Nitrogen Solubility in Iron Melt” (in Korean), J. Korean Inst. Met., 22(9), 789–793 (1984) (Experimental, Thermodyn., 18) Raghavan, V., “The Fe-Mn-N (Iron-Manganese-Nitrogen) System”, Trans. Indian Inst. Met., 37(3), A-F (1984) (Phase Diagram, Phase Relations, Review, 30) Bashchenko, A.P., Izotov, V.I., Omelchenko, A.V., Soshnikov, V.I., Shcherbedinskii G.V., “Alloying Iron and its Alloys with Nitrogen Under High Pressure”, Russ. Metall., (4), 169–174 (1985), translated from Izv. Akad. Nauk SSSR, Met., (4), 173–178 (1985) (Crys. Structure, Morphology, Thermodyn., 13)
DOI: 10.1007/978-3-540-78644-3_16 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–N [1986Wad] [1987Rag]
[1987Wri]
[1989Roc]
[1990Siw]
[1991Fri] [1991Sat]
[1992Vol]
[1993Qiu1] [1993Qiu2] [1993Rag] [1994Fer]
[1994Kun] [1994Tan]
[1995Nec]
[1998Gav]
[1999Ari] [1999Sum]
[2000Gou]
[2001Kim]
Landolt-Börnstein New Series IV/11D4
17
Wada, H., Lee, S.W., Pehlke, R.D., “Nitrogen Solubility in Liquid Fe and Fe-Mn Alloys”, Met. Trans. B, 17B(1), 238–239 (1986) (Experimental, Phase Relations, Thermodyn., 17) Raghavan, V., “The Fe-Mn-N (Iron-Manganese-Nitrogen) System” in “Phase Diagrams of Ternary Iron Alloys”, Ind. Inst. Metals, Delhi, 1, 183–189 (1987) (Crys. Structure, Phase Diagram, Phase Relations, Review, 26) Wriedt, H.A., Gokcen, N.A., Nafziger, R.H., “The Fe-N (Iron-Nitrogen) System”, Bull. Alloy Phase Diagams, 8(4), 355–377 (1987) (Crys. Structure, Phase Diagrams, Thermodyn., Review, *, #, 126) Rochegude, P., Foct, J., “Behaviour of Nitrogen During Mechanical Alloying at Solid Solutions and Iron Nitrides”, Compt. Rend. Acad. Sci. Paris, Ser. 2, 309, 1545–1549 (1989) (Experimental, Phase Relations, 9) Siwka, J., Tochowicz, S., “Nitrogen Solubility in Liquid Iron Alloys Containing High Amounts of Chromium and Manganese in Hyperbaric Conditions”, Arch. Metall., 35(3), 449–461 (1990) (Experimental, Phase Relations, 26) Frisk, K., “A Thermodynamic Evaluation of the Fe-Ni-N System”, Z. Metallkd., 82, 59–66 (1991) (Phase Diagram, Phase Relations, Assessment, Thermodyn., 36) Satir-Kolorz, A.H., Feichtinger, H.K., “On the Solubility of Nitrogen in Liquid Iron and Steel Alloys Using Elevated Pressure”, Z. Metallkd., 82(9), 689–697 (1991) (Experimental, Calculation, Phase Relations, 24) Volynova, T.F., Sidorova, I.B., Emelyanova, I.Z., “State of the Boundaries and Features of the Fracture of Powder - Metallurgical Iron - Manganese Alloys”, Metal Sci. Heat Treat., 34(11–12), 693–699 (1992) (Experimental, Mechan. Prop., 4) Qiu, C., Fernandez-Guillermet, A., “A Thermodynamic Evaluation of the Mn-N System”, Z. Metallkd., 84(1), 11–22 (1993) (Phase Diagram, Phase Relations, Assessment, Thermodyn., 54) Qiu, C., “A Thermodynamic Evaluation of the Fe-Mn-N System”, Metall. Trans. A, 24A, 629–645 (1993) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 49) Raghavan, V., “Fe-Mn-N (Iron-Manganese-Nitrogen)”, J. Phase Equilib., 14(5), 629 (1993) (Phase Diagram, Phase Relations, Review, 8) Fernandez-Guillermet, A., Du, H., “Thermodynamic Analysis of the Fe-N System using the Compound-Energy Model with Prediction of the Vibrational Entropy”, Z. Metallkd., 85(3), 154–163 (1994) (Phase Diagrams, Phase Relations, Theory, Assessment, 75) Kunze, J., Rothe, I., “Solubility of Nitrogen in Austenitic FeCrMn Alloys”, Steel Research, 65(8), 331–337 (1994) (Calculation, Experimental, Thermodyn., 25) Tanaka, T., Gokcen, N.A., Iida, T., Morita, Z.-I., “Thermodynamic Relationship Between the Enthalpy Interaction Parameter and the Entropy Interaction Parameter in Liquid IronNitrogen Based Ternary Alloys”, Z. Metallkd., 85, 696–700 (1994) (Theory, Thermodyn., 17) Nechaev, Yu.S., Omelchenko, A.V., “Solubility of Molecular Nitrogen in Austenite”, Russ. J. Phys. Chem., 69(9), 1408–1413 (1995), translated from Zh. Fiz. Khim., 69(9), 1556–1561 (1995) (Experimental, Phase Relations, 23) Gavriljuk, V.G., Yakovenko, P.G., Ullakko, K., “Influence of Nitrogen on Vibration Damping and Mechanical Properties of Fe-Mn Alloys”, Scr. Mater., 38(6), 931–935 (1998) (Experimental, Mechan. Prop., 7) Ariapour, A., Yakubsov, I., Perovic, D.D., “Effect of Nitrogen on Shape Memory Effect of a Fe-Mn-based Alloy”, Mater. Sci. Eng. A, 262(1–2), 39–49 (Experimental, Mechan. Prop., 24) Sumin, V.V., Chimid, G., Rashev, Ts., Saryivanov, L., “The Neutron-Spectroscopy Proof of the Strong Cr-N Interactions in Nitrogen Stainless Steels”, Mater. Sci. Forum, 318–320, 31–40 (1999) (Crys. Structure, Electronic Structure, Experimental, 11) Goune, M., Belmonte, T., Fiorani, G.M., Chomer, S., Michel, H., “Modelling of Diffusion Precipitation in Nitrided Alloyed Iron”, Thin Solid Films, 377, 543–549 (2000) (Interface Phenomena, Calculation, Phase Relations, 27) Kim, E.J., You, B.D., Pak, J.J., “Nitrogen Solubility in Liquid Manganese and Ferromanganese Alloys”, Metall. Mater. Trans. B, 32B(4), 659–668 (2001) (Experimental, Thermodyn., 22)
MSIT®
DOI: 10.1007/978-3-540-78644-3_16 # Springer 2008
18 [2002Bli]
[2002Lee]
[2003Gou1]
[2003Gou2]
[2004Pad] [2004Wit]
[2006Che] [2007Wit] [Mas2] [V-C2]
Fe–Mn–N Bliznuk, V.V., Glavatska, N.I., Soederberg, O., Lindroos, V.K., “Effect of Nitrogen on Damping, Mechanical and Corrosive Properties of Fe-Mn Alloys”, Mater. Sci. Eng. A, 338, 213–218 (2002) (Experimental, Interface Phenomena, Mechan. Prop., Phase Relations, 10) Lee, Y.K., “Effects of Nitrogen on γ -ɛ Martensitic Transformation and Damping Capacity of Fe-16%Mn-X%N Alloys”, J. Mater. Sci. Lett., 21(15), 1149–1151 (2002) (Experimental, Mechan. Prop., 15) Goune, M., Belmonte, T., Redjaimia, A., Weisbecker, P., Fiorani, J.M., Mochel, H., “Thermodynamic and Structure Studies on Nitrided Fe-1.62%Mn and Fe-0.56%V Alloys”, Mater. Sci. Eng. A, 351, 23–30 (2003) (Calculation, Crys. Structure, Experimental, Phase Relations, 22) Goune, M., Redjaimia, A., Belmonte, T., Michel, H., “Identification and Characterization of a Novel Mn-N Nitride Formed in Fe-Mn-N Alloy”, J. Appl. Crystallogr., 36(1), 103–108 (2003) (Crys. Structure, Experimental, Mechan. Prop., Phase Relations, 26) Paduani, C., “Electronic Structure and Magnetic Properties of MnFe3N”, J. Appl. Phys., 96(3), 1503–1506 (2004) (Calculation, Electronic Structure, Magn. Prop., 6) Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diff., 25(4), 346–354 (2004) (Experimental, Phase Diagram, Calculation, Thermodyn., #, 34) Chen, L., “Electronic Structure and Magnetism of Fe4–xMnxN Compounds”, J. Appl. Phys., 100(11), 13717 (2006) (Electronic Structure, Magn. Prop., Calculation, 20) Witusiewicz, V.T., private communication to MSI, (2007) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson′s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_16 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni
1
Iron – Manganese – Nickel Nathalie Lebrun, Pierre Perrot
Introduction Mn and Ni are important constituents of steels, so, the Fe-Mn-Ni ternary system raised much interest. [1909Jae, 1913Par] gave a first insight to the general shape of the liquidus surfaces which has largely been confirmed in more recent works [1985Koc, 1994Rag] and reproduced in the reviews of [1983Riv, 1988Ray]. There is a remarkable consistency between the data of [1913Par] and those related to binary systems determined half a century later. A more recent determination of the liquidus was carried out by [1997Sch]. A stability domain for the phases α, γ and ε is shown by [1980Bra], but the conditions of pressure and quenching are not given, so that the diagram cannot be used. No Calphad assessment exists for the whole systems, but partial data have been gathered [1998Mie, 2000Mie]. Main experimental data are reported in Table 1. Binary Systems Using new experimental thermodynamic data, the Fe-Mn system has been recently reassessed and the Calphad description was updated by [2004Wit]. This new reevaluation of the phase equilibria leads to consistently better fits to the available experimental data. There is however, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have been written as L ⇌ (δMn) + (γMn,γFe). Consequently, the Fe-Mn system is accepted from [2004Wit]. The Fe-Ni and Mn-Ni phase diagrams are taken from [2007Kuz] and [2007Wat], respectively. In a thermodynamic assessment of the Mn-Ni binary carried out by [2005Guo], the two intermetallic phases η’MnNi and η”MnNi are considered as a unique phase. Solid Phases All the data on unary, binary and ternary phases are indicated in Table 2. The γ’(MnxFe1–x)Ni3 solid solution, ordered below 525°C with a L12 structure (AuCu3 type) presents a maximum in the Young’s modulus and in the order-disorder transition temperature, [1966Ibr, 1967Yeg] together with a minimum in the electrical conductivity around x ∼ 0.5 [1966Ibr] and an anomaly in the thermal capacity and in the saturation magnetization [1966Ibr], which is an indication of a short-range ordering of Fe and Mn around x ∼ 0.5 in addition to the long-range order of the L12 (AuCu3) type which superimposes in the whole composition range. Liquidus, Solidus and Solvus Surfaces The liquidus surface has been experimentally determined by [1913Par, 1985Koc], critically reinterpreted by [1997Sch] and calculated, together with the solidus surface, by [1985Koc], reproduced in [1994Rag]. But this last calculation does not match well with the accepted binary systems and can not be taken into account. Figure 1 presents the liquidus projection from [1997Sch], slightly modified to take into account the accepted binary systems. The peritectic monovariant lines have also been added. Figure 2 represents the liquidus line in the Fe rich corner of the diagram, calculated by [1998Mie] and showing the primary crystallization fields of the α and γ solutions. Isothermal Sections The α-γ equilibria in the Fe rich corner were calculated by [1989Har] at 850, 750, 650 and 550°C. Some of these partial isotherms are shown in Figs. 3a and 3b. The same equilibria, calculated at 700°C by [1995Tan] Landolt-Börnstein New Series IV/11D4
MSIT®
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
2
Fe–Mn–Ni
using a central atom model differ from the preceding one by the shape of the (α+γ) domain which is shrunk along the Mn/Ni = 1 line, which does not agree with the experimental observations [2002Mun]. The shape of the (α+γ)/γ border at 550°C (Fig. 3b) is very similar to the shape of the same border calculated at 450°C by [2002Mun] and confirmed by microscopic observations, which increases the degree of confidence which may be attributed to the calculations of [1989Har]. Figure 4 shows the whole Fe-Mn-Ni diagram calculated at 450°C by [2004Mun]. This diagram has been widely modified to take into account the non stoichiometry of the intermetallic compounds and the non existence of the Mn3Ni and MnNi2 phases at 450°C. Thermodynamics Mn activities have been measured by a Knudsen effusion cell [1964Smi] at 1059°C in the whole concentration range. Fe and Ni activities were calculated by integration of the Gibbs-Duhem equation through the ternary system. In the γ field, Mn-Ni alloys present a strong negative departure from ideality, departure which decreases as Fe replaces Ni in the solid solution. Expressions of the Gibbs energy for the liquid and γ phases were presented by [1987Dan] using a subregular model with ternary interaction parameters, but the calculated liquidus surface agree poorly with the experimental one. The existence of a metastable miscibility gap island in the α(Fe,Mn,Ni) alloy was postulated [1970Suz] to explain the age hardening observed in martensitic alloys before their decomposition into a biphased (α + γ) stable state. The miscibility gap was deduced by combination of the energy parameters related to the three binary systems, but was not confirmed by further studies. Careful investigation on the Fe-12Ni-6Mn (mass%) alloy [1981Ray] explain age hardening by the precipitation of the η’MnNi in the matrix, a tetragonal compound whose ratio c/a lies close to unity. Notes on Materials Properties and Applications Main experimental works are reported in Table 3. Mn is known to decrease the Curie temperature of Fe-Ni alloys [1969Col] and increase the atomic magnetic moment of alloys in which the atomic ratio Ni:Fe = 4:1 up to a maximum observed at 5.0 at.% Mn in the alloy. However, Mn in Ni3Fe alloy increases the Curie temperature up to a maximum of 2.5 at.% Mn in the alloy [1991Gan]. Such an ordering explains the difficulty of cold rolling Ni3Fe alloys containing more than 0.5 at.% Mn. The atomic magnetic moment μB = 1.074 for Ni80Fe20 increases up to 1.107 for Ni76Fe19Mn5, decreases down to 1.033 for Ni72Fe19Mn10. For atomic ratios Ni:Fe = 3:2 and 2:3, Mn decreases the atomic magnetic moment. Mn was also shown to lower the degree of long range atomic order in Ni3Fe [1987Bab1]. Neutron scattering on Ni3Fe1–xMnx shows that Fe and Mn are randomly arranged on the site. When x rises, the order-disorder transition is replaced by a second order phase transition [1987Bab2]. The magnetic phase diagram has been investigated along the 50 at.% Fe section by [1974Ett, 1985Der], at 15 and 20 at.% Mn by [1981Men], at 72 at.% Fe by [1990Dub], at a ratio Fe/Ni = 65/45 and < 30 at.% Mn by [1992Wul, 1995Hoe] and in the whole composition range by [1976Men]. The magnetic moment in the Fe-Mn-Ni alloys is calculated by [1990Jep] and a statistical theory of the magnetic ordering in Fe-Mn-Ni alloys was developed by [1990Her]. The experimental magnetic diagram [1976Men, 1990Jep, 1995Hoe] is represented in Fig. 5. Both magnetic orderings (ferro + antiferro) exist at low temperature around the transition. The heat capacity was shown to present a maximum near the concentration for which the alloys looses its long range ferromagnetic behavior [1975Lev], but on both sides of this concentration, the heat capacity variations are weak. A spin glass state has been observed [1983Der, 1984Der, 1985Der, 1986Der] below 100 K in the concentration range corresponding to the ferromagnetic-antiferromagnetic transition. Invar behavior is characterized by a minimum of the linear expansion coefficient observed around 35 at.% Ni [1977Zak]. The Invar anomaly is connected to the ferromagnetic properties of the alloy, a minimum in the expansion coefficient corresponding to a maximum in the magnetic moment. The Invar effect disappears when the Mn content of the alloys exceeds 10 at.%. Characteristics of the antiferromagnetic alloys (Mn: Fe = 2:3) are well described by electron concentration [1971Col]. Linear spontaneous magnetostriction and Neel temperature decrease with the Ni content.
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni
3
Miscellaneous [2006Li] investigated the microstructural evolution of 36 mass% Ni austenitic steel under severe plastic deformation using the accumulative roll-bonding process at 500°C. During the mechanical deformation, a grain subdivision process occurs leading to ultrafine grain formation. The structure formed is found finer and straighter than that formed in ferrite steel deformed under similar conditions. Fe-Mn-Ni alloys are reported to show a remarkable age hardening behavior [2002Mun, 2006Ned] attributed to the precipitation of ηMnNi particles. However, they suffer grain boundary embrittlement depending on the ageing process. The possible causes of that effect are discussed by [2004Wil] which concludes that Mn is a major embrittlement cause of its own right due to its affinity towards N, S and P which segregate towards the grain boundaries before the austenitization, a point of view defended earlier by [1983Ray]. Fe-Mn-Ni presents a martensitic transformation which may be athermal (or burst type) or isothermal [2000Rag]. An athermal transformation starts abruptly at a well defined Ms (martensite start) temperature. The amount of martensite formed at Ms varies from a few percent to 50%. An athermal transformation is observed in alloys containing between 20 and 30 at.% Ni. In an isothermal transformation, the Ms point is not well defined and the transformation occurs in a narrow temperature range or after some finite incubation time during isothermal holding. Fe-Mn-Ni alloys are typical for presenting an isothermal martensitic transformation. In Fe containing 23 at.% Ni, the nature of the transformation is very sensible to the Mn content between 3 and 4 at.% Mn [1998Pal, 1999Cha, 2000Cha]. A hydrostatic pressure or a magnetic field [2000Kak] may also change the nature of the transformation. For Fe72Ni24Mn4 (in at.%) alloy, which exhibits an isothermal martensitic transformation under the room temperature, a static magnetic field (up to 3.2 MA·m–1) lowers the nose temperature of the TTT curve together with the incubation time whereas a hydrostatic pressure (up to 0.5 GPa) increases the nose temperature of the TTT curve together with the incubation time [2001Kak]. There is a strong correlation between the magnetic behavior and the nature of the austenite-martensite transformation [2002Ayd]. The Fe alloy with 10 mass% Ni and 10 mass% Mn, austenitic in the quenched state from 1100°C, presents a Ms point at –20°C [1967Bou]. By heating, the martensite-austenite transformation is observed between 450 (As point) and 640°C. The antiferromagnetic alloy Fe72Mn22Ni6 presents a Ms point at 47°C [1988Dub1] and a Neel temperature at 85°C [1988Dub2]. Whereas the austenitic phase is paramagnetic, the martensitic phase may be ferro- or antiferromagnetic [2005Akt]. Alloys with 3 to 7 mass% Ni and 11 to 15 mass% Mn present high damping capacity under frequencies as low as 0.2 kHz, a shape memory behavior and superelasticity [1995Wat] which is attributed to the austenitemartensite transformation. The best property is observed for the 5Ni-15Mn (in mass%) alloy. (Fe,Mn,Ni) alloys have been electrodeposited from a mixture of sulfates or chloride [1980Sri]. The structure of the layer is α or γ and the composition depends strongly on the temperature (10–60°C) and on the pH (2-8) imposed. A deposit with a minimum internal stress can be obtained using a high bath temperature, a low cathode current density, a low pH and a rough cathode surface.
Table 1.
Investigations of the Fe-Mn-Ni Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1913Par]
Thermal analysis
1100–1450°C, liquidus and solidus surfaces (whole diagram)
[1954Kur]
Thermal analysis, dilatometry, microscopic investigations
< 800°C, phase equilibria in the solid state (whole diagram)
[1964Smi]
Knudsen effusion cell, Mn vapor pressure measurements
1059°C, the whole system, Mn activity measured, Fe and Ni activities calculated
[1975Lev]
Magnetization, heat capacity measurements
< 0°C, Fe / Mn = 86 / 14, 15 to 40 at.% Ni (invar alloys) (continued)
Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
4
Fe–Mn–Ni
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1983Der]
Magnetic transitions, heat capacity measurements
5–90 K, 50 at.% Fe
[1984Der]
Magnetic transitions, heat capacity measurements
80–320 K, 20 at.% Mn, spin glass state
[1985Koc]
Thermal analysis
1000–1500°C, liquidus and solidus surfaces (whole diagram)
[1986Der]
Magnetic transitions, heat capacity measurements
5–20 K, 15 at.% Mn, 45 to 65 at.% Fe, spin glass state
[1986Gom]
Neutron diffraction, long range atomic order determination
15 to 75 at.% Ni, 10 to 70 at.% Fe
[1987Bab1, 1987Bab2]
Neutron diffraction, long range atomic order determination
Ni3Fe1–xMnx (x < 0.3), order-disorder transition
[1987Dan]
Calorimetry (temperature not given)
Integral mixing enthalpy of the liquid in the whole composition range
[1988Dub1]
Diffuse neutron scattering on single crystal
Fe72Mn22Ni6 (in at.%), martensitic transformation
[1988Gom]
Neutron diffraction
400°C, phase diagram
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
α, (αδFe,δMn)
cI2 Im 3m W
(αFe) (Ferrite) < 912
Lattice Parameters [pm]
Comments/References
a = 286.65
(δFe) 1538–1394
a = 293.15
(δMn) 1246–1138
a = 308.13
pure Fe at 25°C [Mas2, V-C2] dissolves 5 at.% Mn at 527°C [2004Wit], dissolves up to 4.6 at.% Ni at 495°C [2007Kuz] at 1390°C [V-C2, Mas2] dissolves 10 at.% Mn at 1474.5°C [2004Wit], dissolves up to 3.8 at.% Ni at 1517°C [2007Kuz] pure Mn at 1137°C [2007Wat] dissolves up to 12.2 at.% Fe at 1236.6°C [2004Wit]. Dissolves 6 at.% Ni [2007Wat] (continued)
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
α’ (Martensite)
tI4 I4/mmm -
(εFe)
hP2 P63/mmc Mg
γ, (γFe,γMn,Ni)
cF4 Fm 3m Cu
(γFe) (Austenite) 1394–912 (γMn) 1138–1100 (Ni) < 1455 (Fe0.65Ni0.35)1–xMnx
Lattice Parameters [pm]
5
Comments/References
metastable phase obtained by quenching from austenite a = 246.8 c = 396.0
at 25°C, 13 GPa [Mas2] triple point α-γ-ε at 8.4GPa, 430°C
a = 364.67
at 915°C [V-C2, Mas2]. Complete solubility with (γMn) [2004Wit] and Ni [2007Kuz] pure Mn at 1097°C [2007Wat]. Complete solubility with (γFe) [2004Wit] and (Ni) [2007Wat]. [Mas2]. Complete solubility with (γFe) [2007Kuz] and (γMn) [2007Wat] x < 0.12, invar alloy [1986Huc]
a = 386.26 a = 352.40 a = 358.3 ± 0.5
(βMn) 1100–727
cP20 P4132 βMn
a = 631.52
[Mas2]. Dissolves up to 32.9 at.% Fe [2004Wit]. Dissolves up to 18 at.% Ni [2007Wat].
(αMn) < 727
cI58 I 43m (αMn)
a = 891.26
pure Mn at 25°C [Mas2]. Dissolves up to 35.2 at.% Fe [2004Wit]. Dissolves up to 9 at.% Ni [2007Wat].
φ, Mn3Ni < 430
tF*
a = 369.8 c = 369.2
75 at.% Mn [2007Wat] antiferromagnetic
ε, Mn2Ni 720–560
-
-
31.5–35.5 at.% Ni [2007Wat]
η, MnNi 911–675
cP2 Pm 3m CsCl
a = 374.5 ± 0.5 a = 297.7
45–52 at.% Ni [1991Gok] at 50 at.% Ni, 750°C [2007Wat]
η’, MnNi 775–620
tP4 P4/mmm AuCu
a = 373.1 c = 363.2
47–55.5 at.% Ni [2007Wat] at 50 at.% Ni [2007Wat] L10 structure
η”, MnNi < 480
t** -
γ’, MnNi3
cP4
-
∼ 46–54 at.% Ni [2007Wat] Antiferromagnetic [1991Gok] L12 structure (continued)
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6
Fe–Mn–Ni
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
< 520
Pm 3m AuCu3
Lattice Parameters [pm]
Comments/References
71–85 at.% Ni [2007Wat] at 75 at.% Ni [2007Wat]
a = 358.9
γ’Fe0.5Mn0.5Ni3 < 525
a = 356.2
γ’, FeNi3 < 517
a = 354.7
Fe-Mn ordering [1966Ibr]
63 to 85 at.% Ni [2007Kuz]
ζ,MnNi2 710-580
t** -
-
64–68.5 at.% Ni [2007Wat]
ζ’, MnNi2 < 440
t** -
-
∼ 66.7–73 at.% Ni [2007Wat]
γ”, FeNi
tP4 P4/mmm AuCu
a = 357.9 c = 357.9
L10 structure [V-C2]. Metastable ordering temperature 320°C at 51.2 at.% Ni [2007Kuz]
Table 3.
Investigations of the Fe-Mn-Ni Materials Properties
Reference
Method/Experimental Technique
Conditions/Type of Property
[1950Kra]
Dilatometry, hardness, tensile strength, TTT curves
< 950°C, < 17 mass% Mn, < 5 mass% Ni, α/γ transformations
[1957Rav]
Dilatometry, electrical conductivity
< 900°C, MnxFe1–x Ni3 (x < 0.03), orderdisorder transitions
[1966Ibr]
XRD, dilatometry, Young’s modulus, tensile strength, electrical conductivity, saturation magnetization
< 800°C, MnxFe1–x Ni3 (0 < x < 1), orderdisorder transitions
[1967Bou]
XRD, dilatometry, hardness measurements
< 650°C, Fe + 10 mass% Mn + 10 mass% Ni, phase transitions
[1967Yeg]
XRD, dilatometry, Young’s modulus
< 750°C, MnxFe1–x Ni3 (0 < x < 1), orderdisorder transitions, ordering kinetics
[1968Min]
Electrical resistance
< 1000°C, MnxFe1–x Ni3 (0 < x < 1), diffusion, kinetics of ordering
[1969Col]
Curie temperature, saturation magnetization
77–840 K, the whole concentration range
[1969Pan]
Neutron diffraction, resistivity, Yield point, magnetization
700°C, Ni3Mn0.95Fe0.05, kinetics of ordering (continued)
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni
7
Reference
Method/Experimental Technique
Conditions/Type of Property
[1970Gom]
Neutron diffraction, saturation magnetization
(Ni3Fe)1–xMnx (0.01 < x < 0.05), Mn influence on ordering
[1971Col]
Expansion, magnetostriction, Neel temperature
< 300°C, Mn / Fe = 2 / 3, < 25 at.% Ni, invar anomaly
[1974Ett]
Mössbauer, antiferromagnetic-ferromagnetic transitions
< 500°C, 50 at.% Fe, 10 to 41 at.% Ni, magnetic phase diagram
[1976Men]
Antiferromagnetic-ferromagnetic transitions
< 800°C, the whole composition range, magnetic phase diagram
[1977Jo]
Nuclear magnetic resonance (NMR)
Magnetic state of Mn diluted in (Fe,Ni) alloys
[1977Sch]
Inelastic neutron scattering, time of flight method
25 and 225°C, Fe60Mn30.5Ni9.5 (in at.%) alloy
[1977Zak]
Dilatometry
4.2–300 K, 35 at.% Ni, < 20 at.% Mn Invar behavior
[1979Che]
Resistivity, saturation magnetization
4.2–90 K, 8.6 mass% Mn + 1.1 mass% Fe annealed 1 h at 900°C
[1980Bar]
Mössbauer, magnetic anisotropy
500–625°C, Fe85Ni10Mn5 (in at.%)
[1980Geo]
Dilatometry, strain gauge, electrical resistivity, martensitic transformation
Fe70+xNi30–2xMnx (x < 3) quenched from 1100–1150°C, then annealed 2 to 30 min at 200-650°C
[1981Men]
Antiferromagnetic-ferromagnetic transitions
< 800°C, 15 and 20 at.% Mn, magnetic phase diagram
[1981Ray]
Scanning and transition electron microscopy (SEM and TEM), neutron diffraction, hardness measurements
Fe-12Ni-6Mn (in mass%), age hardening, mechanism investigation
[1983Zol]
Magnetic susceptibility, magnetostriction (100-300 kOe)
77–900 K, Fe70+xNi30–2xMnx (x < 3)
[1984Kaj]
Martensitic transformation, optical microscopy at 143 K
Fe-23Ni-3.9Mn (in mass%), austenetized at 1250°C, water quenched then observed at 143 K
[1985Der]
Magnetic susceptibility, magnetostriction, elastic modulus, dilatometry
4.2–1000 K, 50 at.% Fe, magnetic phase diagram, spin glass state
[1986Huc]
Mössbauer
(Fe0.65Ni0.35)1–xMnx (x < 0.12) invar behavior
[1987Sin]
Optical microscopy, TEM, martensitic transformation
500–600°C, 19.8 mass% Ni, 5.2 mass% Mn
[1988Dub2]
Neutron scattering on single crystal, magnetic susceptibility
Fe72Mn22Ni6 (in at.%), magnetic ordering near the Neel temperature
[1990Dub]
Neutron scattering on single crystal, magnetization
4.2–800 K, 72 at.% Fe, magnetic ordering, martensitic transitions (continued)
Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
8
Fe–Mn–Ni
Reference
Method/Experimental Technique
Conditions/Type of Property
[1991Gan]
DTA, XRD, electrical conductivity, Curie temperatures
< 700°C, 20 to 26 at.% Fe, < 5 at.% Mn, magnetic ordering
[1991Wat]
SEM, tensile strength
3 to 7 at.% Ni, 11 to 15 a % Mn, damping properties
[1992Wul]
Magnetization, Curie temperature
4–280 K, Fe / Ni = 65 / 35, < 30 at.% Mn, magnetic diagram
[1995Hoe]
Electrical resistivity, Mössbauer
4–280 K, Fe / Ni = 65 / 35, < 30 at.% Mn, spin glass state
[1995Wat]
Dilatometry, hardness, stress-strain diagram
3 to 7 mass% Ni, 11 to 15 mass% Mn, damping effect
[1996Heo]
SEM, electrical resistivity on age hardened alloy
400–550°C, 8 mass% Mn, 7 mass% Ni, Mn segregation
[1998Pal]
XRD (Rietveld refinement), microhardness
77 and 143 K, 23 at.% Ni, 3.8 at.% Mn, martensitic transformation
[1999Cha]
XRD (Rietveld refinement), SEM, microhardness
140 to 250 K, 23 at.% Ni, 3.3 at.% Mn, martensitic transformation
[2000Cha]
XRD (Rietveld refinement), SEM, microhardness
77 to 200 K, 23 at.% Ni, 3.6 at.% Mn, martensitic transformation
[2000Kak, 2001Kak]
SEM, electrical resistivity, hydrostatic pressure (< 1 GPa)
77 to 210 K, Fe72Ni24Mn4 (in at.%), martensitic transformation
[2002Ayd]
SEM, Mössbauer at 77 K
30 mass% Ni, 0.2 mass% Mn, martensitic transformation
[2002Mun, 2004Mun]
TEM, HR-TEM (high spatial resolution TEM), selected area electron diffraction (SAED), X-Ray energy dispersive spectroscopy (EDS)
8 mass% Mn, 7 mass% Ni annealed at 1200°C, then aged 0.5 to 48 h at 450°C, ageing behavior, precipitation of austenite particles at grain boundaries
[2002Tsu]
Magnetization curves, magnetic torque curves
Ni-Fe/Mn82Ni18 thin layers on MgO substrate, exchange anisotropy of ferroantiferromagnetic bilayers
[2005Akt]
TEM, Mössbauer, magnetic properties of the martensite
Fe-25Ni-5Mn and Fe-34.5Ni-10Mn (in mass%), austenitized at 1100°C then quenched at 77 K
[2005Bel]
Stress-strain curbes, electron backscattered diffraction (EBSD)
Ni-30 Fe-0.02 Mn-0.01 C (in mass%), hot rolled at 1000–1200°C
[2006Gor]
Magnetization curves, magneto-optical indicator film (MOIF) technique
Fe50Mn50/Ni81Fe19 thin layers, exchange anisotropy of ferro-antiferromagnetic bilayers
[2006Li]
FEG-SEM (field emission gun SEM), EBSD
36 mass% Ni, accumulative roll bonding deformation at 500°C
[2006Ned]
FEG-TEM (field emission gun TEM), HREDS, SAED
Fe-10Ni-7Mn (mass%), aged at 480°C
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni
Fig. 1. Fe-Mn-Ni.
Landolt-Börnstein New Series IV/11D4
9
Liquidus surface projection
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Fig. 2. Fe-Mn-Ni.
Fe–Mn–Ni
The Fe rich corner of the liquidus surface
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni
Fig. 3a. Fe-Mn-Ni.
Landolt-Börnstein New Series IV/11D4
11
The α/γ equilibrium in the Fe rich corner 650°C
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12
Fig. 3b. Fe-Mn-Ni.
Fe–Mn–Ni
The α/γ equilibrium in the Fe rich corner 550°C
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni
Fig. 4. Fe-Mn-Ni.
Landolt-Börnstein New Series IV/11D4
13
Calculated isothermal section at 450°C
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Fig. 5. Fe-Mn-Ni.
Fe–Mn–Ni
The magnetic phase diagram
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni References [1909Jae] [1913Par]
[1950Kra]
[1954Kur]
[1957Rav] [1964Smi]
[1966Ibr]
[1967Bou]
[1967Yeg]
[1968Min]
[1969Col] [1969Pan]
[1970Gom]
[1970Suz] [1971Col] [1974Ett] [1975Lev]
[1976Men]
Landolt-Börnstein New Series IV/11D4
15
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, 22) Parravano, N., “The Ternary Alloys of Iron - Nickel - Manganese, Nickel - Manganese Copper, Iron - Manganese - Copper Systems” (in German), Z. Metallkd., 4, 171–203 (1913), reproduced in Int. Z. Metallography, 4, 187–202 (1913) (Experimental, Phase Diagram, Phase Relations, 14) Kramer, I.R., Toleman, S.L., Haswell, W.T., “Iron-Manganese and Iron-Manganese-Nickel Alloys”, Trans. Amer. Soc. Met., 42, 1260–1294 (1950) (Experimental, Phase Relations, Magn. Prop., Kinetics, 4) Kurnakov, N.N., Troneva, M.Ya, “The Iron-Manganese-Nickel System Investigation” (in Russian), Izv. Sekt. Fiz.-Khim. Anal., 24, 132–147 (1954) (Experimental, Magn. Prop., Morphology, Phase Diagram, Phase Relations, 17) Ravdel, M.P., Selissky, Ya.P., “Transitions in Ternary Ni3Fe Solid Solutions” (in Russian), Dokl. Akad. Nauk SSSR, 115(2), 319–321 (1957) (Experimental, Morphology, 5) Smith, J.H., McCabe, C.L., Paxton, H.W., “Manganese Vapor Pressures in Equilibrium with Manganese-Iron-Nickel Solid Solutions”, J. Phys. Chem., 68(6), 1345–1354 (1964) (Experimental, Phase Relations, Thermodyn., 37) Ibragimov, E.A., Selissky, Ya.P., Sorokin, M.N., “Investigation of the Effects of Ordering on Ternary Ni-Fe-Mn Solid Solutions”, Russ. Metall., (5), 82–87 (1966), translated from Izv. Akad. Nauk SSSR, Met., (5), 152–157 (1966) (Crys. Structure, Electr. Prop., Magn. Prop., Experimental, 3) Bourgeot, J., Manenc, J., “Structure of an Fe-Mn-Ni Alloy” (in French), Mem. Scient. Revue Metall, 64(11), 1006–1007 (1967) (Experimental, Phase Relations, Crys. Structure, Magn. Prop., 6) Yeganyan, I.L., Selisskiy, Ya.P., “Atomic Ordering Effects in Binary Solid Solutions (Ni3Fe, Mn)”, Phys. Met. Metallogr., 23(2), 195–197 (1967), translated from Fiz. Met. Metalloved., 23(2), 369–371 (1967) (Crys. Structure, Magn. Prop., Experimental, 5) Minaeva, G.G., Panin, V.E., Fadin, V.P., Pyabyshkina, F.A., Zhukova, V.M., “The Effect on the Ordering Kinetics of Alloying Ni3Mn Alloy with Co and Fe” (in Russian) in “Diffuzionnye Protsessy v Metallakh”, Naukova Dumka, Kiev, 84–88 (1968) (Experimental, Kinetics, Electr. Prop., 12) Colling, D.A., “Intrinsic Magnetization of Fe-Ni-Mn Alloys”, J. Appl. Phys., 40(3), 1379–1381 (1969) (Experimental, Magn. Prop., 6) Panin, V.E., Prushinsky, V.V., Fadin, V.P., Novak, L.I., “Effect of Alloying a Ni3Mn Alloy with a Third 3d-Transition Element on Ordering Kinetics”, Sov. Phys., Doklady, 14(11), 1119–1122 (1970), translated from Dokl. Akad. Nauk SSSR, 189(1), 84–87 (1969) (Crys. Structure, Experimental, Kinetics, 15) Goman’kov, V.I., Puzey, I.M., Mal’tsev, E.I., “Effect of Alloying Elements on the Superstructure of Ni3Fe” (in Russian), Dokl. Akad. Nauk SSSR, 194(2), 309–311 (1970) (Crys. Structure, Experimental, Magn. Prop., 6) Suzuki, T., “Metastable Miscibility Gap Island in Fe-Ni-Mn Ternary Martensitic Alloys”, Trans. Jpn. Inst. Met., 11(4), 257–263 (1970) (Calculation, Thermodyn., Magn. Prop., 21) Colling, D.A., Mathur, M.P., “Invar Behavior in Antiferromagnetic Fe-Mn-(Ni) Alloys”, J. Appl. Phys., 42(13), 5699–5703 (1971) (Experimental, Magn. Prop., 27) Ettwig, H.H., Pepperhoff, W., “On Magnetism of γ-Fe-Ni-Mn Alloys”, Phys. Status Solidi A, 23(1), 105–111 (1974) (Experimental, Magn. Prop., Crys. Structure, 18) Levesque, B., Caudron, R., Costa, P., “Low Temperature Specific Heat of Ni1–x(Mn0.14Fe0.86)x Alloys”, Proceedings of the International Conference on Low Temperature Physics, 3, 246–249 (1975) (Magn. Prop., Thermodyn., Experimental, 6) Menshikov, A.Z., Kazantsev, V.A., Kuzmin, N.N., “Magnetic State of Fe-Ni-Mn Alloys” (in Russian), Zh. Eksp. Teor. Fiz., 71, 648–656 (1976) (Experimental, Magn. Prop., 26)
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16 [1977Jo]
[1977Sch] [1977Zak]
[1979Che]
[1980Bar]
[1980Bra] [1980Geo]
[1980Sri] [1981Men]
[1981Ray] [1983Der]
[1983Ray] [1983Riv]
[1983Zol]
[1984Der]
[1984Kaj] [1985Der]
Fe–Mn–Ni Jo, T., “Magnetic State of Mn Atom in Ferromagnetic Ni-Fe-Mn and Ni-Mn Alloys: Local Environment Effect”, Physica B/C, 86–88(pt.2), 747–748 (1977) (Experimental, Magn. Prop., 4) Schreiber, J., Matz, W., “Antiferromagnons in Fe-Mn-Ni Alloys”, Physica B/C, 86–88(pt.1), 272–274 (1977) (Experimental, Magn. Prop., 5) Zakharov, A.I., Basargin, O.V., Menshikov, A.Z., Kalinin, V.M., “Temperature-Coefficient of Linear Expansion of Fe-Ni-Mn Alloys”, Phys. Met. Metallogr., 43(5), 182–185 (1977), translated from Fiz. Met. Metallov., 43(5), 1103–1105 (1977) (Experimental, Magn. Prop., 12) Cherenkov, V.A., Rogel`berg, I.L., Beylin, V.M., “Low-Temperature Resistivity Anomalies in Nickel-Manganese, Nickel-Manganese-Cobalt (or Iron) Alloys”, Phys. Met. Metallogr., 47(4), 168–170 (1979), translated from Fiz. Met. Metalloved., 47(4), 865–867 (1979) (Experimental, Electr. Prop., Magn. Prop., 12) Barton, J., Kansy, J., Panek, T.J., “Influence of Mn and Ni on the Magnetic-Properties of FeNi-Mn Alloys”, Phys. Status Solidi A, 62(2), 407–412 (1980) (Experimental, Magn. Prop., Morphology, 12) Brandes, E.A., Flint, R.F., “Mn-Fe-Ni” in “Manganese Phase Diagrams”, Manganese Centre, Paris, France, 114 (1980) (Phase Diagram, Phase Relations, Review, 3) Georgieva, I.Ya., Matyushenko, L.A., “Effect of Heat Treatment on the Kinetics of the TwoStage Martensitic Transformation in Fe-Ni-Mn and Fe-Ni-Mo Alloys”, Met. Sci. Heat Treat., 22(5-6), 311–314 (1980), translated from Metallov. Term. Obrab. Metal., SSSR, 22(5), 3–5 (1980) (Experimental, Kinetics, Phase Relations, 7) Srivastava, S.C., “Electrodeposition of Ternary Alloys: Developments in 1972 - 1978”, Surf. Technol., 10, 237–257 (1980) (Electrochemistry, Magn. Prop., Review, 121) Menshikov, A.Z., Burlet, P., Chamberod, A., Tholence, J.L., “Magnetic Phase Diagram of γ-FeNiMn Alloys with a Mixed Exchange Interaction”, Solid State Commun., 39(10), 1093–1095 (1981) (Experimental, Magn. Prop., Phase Relations, 14) Ray, R.K., “Structure and Properties of an Fe-Ni-Mn Alloy”, Z. Metallkd., 72(3), 203–210 (1981) (Crys. Structure, Morphology, Magn. Prop., Experimental, 16) Deryabin, A.V., Rimlyand, V.I., Larionov, A.P., “Low-temperature Specific Heat of Fe-Ni-Cr and Fe-Ni-Mn Alloys”, Sov. Phys., Solid-State, 25(7), 1109–1111 (1983), translated from Fiz. Tverd. Tela, 25(7), 1921–1925 (1983) (Experimental, Thermodyn., 23) Ray, R.K., “Strenghtening Mechanisms in a Fe-11 Mn-6.5 Ni Alloy”, Arch. Eisenhuettenw., 54(12), 513–518 (1983) (Experimental, Magn. Prop., 15) Rivlin, V.G., Raynor, G.V., “Phase Equilibriums in Iron Ternary Alloys 9: Critical Review of Constitution of Ternary Systems Fe-Mn-X (X = Ti, V, Cr, Co, Ni, Cu)”, Int. Met. Rev., 28(1), 23–64 (1983) (Phase Diagram, Phase Relations, Review, 68) Zolotarevskiy, I.V., Snezhnoy, V.L., Sheyko, L.M., “Magnetostriction of the Austenite of Iron-Nickel-Manganese Alloys and the Martensitic Transformation Induced by a Strong Magnetic Field”, Phys. Met. Metallogr., 55(3), 115–121 (1983), translated from Fiz. Met. Metallov., 55(3), 548–553 (1983) (Experimental, Magn. Prop., Phase Relations, 20) Deryabin, A.V., Menshikov, A.Z., Larionov, A.P., Kazantsev, V.K., “Heat Capacity of FexNi80–xMn20 Alloys”, Phys. Met. Metallogr., 57(3), 182–185 (1984), translated from Fiz. Met. Metalloved., 57(3), 614–616 (1984) (Experimental, Magn. Prop., Thermodyn., 12) Kajuwara, S., “Continuous Observation of Isothermal Martensite Formation in Fe-Ni-Mn Alloys”, Acta Metall., 32(3), 407–413 (1984) (Experimental, Morphology, Thermodyn., 16) Deryabin, A.V., Tkov, A.V., Kazantsev, V.K., Shvetsov, B.N., “Magnetic Phase Diagram, Thermal and Magnetoelastic Properties of Alloys of System Fe50Ni50–xMnx”, Phys. Met. Metallogr. 59(3), 52–56 (1985), translated from Fiz. Met. Metalloved., 59(3), 476–480 (1985) (Experimental, Magn. Prop., Phase Relations, 11)
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Fe–Mn–Ni [1985Koc]
[1986Der]
[1986Gom]
[1986Huc] [1987Bab1]
[1987Bab2]
[1987Dan]
[1987Sin]
[1988Dub1]
[1988Dub2]
[1988Gom]
[1988Ray] [1989Har]
[1990Dub]
[1990Her]
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17
Kocherzhinskii, Yu.A., Kulik, O.G., Turkevich, V.Z., “Melting Diagram of the Fe-Mn-Ni System”, Russ. Metall., (4), 204–207 (1985), translated from Izv. Akad. Nauk SSSR, Met., (4), 210–213 (1985) (Experimental, Calculation, Phase Diagram, Phase Relations, 4) Deryabin, A.V., Kourov, N.I., Larionov, A.P., Menshikov, A.Z., “On the Nature of the Low Temperature Heat Capacity Maximum in the Range of the Ferro-Antiferromagnetic Transition in Iron-Nickel-Manganese Alloys”, Phys. Met. Metallogr., 62(1), 83–88 (1986), translated from Fiz. Met. Metalloved., 62(1), 95–100 (1986) (Experimental, Kinetics, Thermodyn., 27) Gomankov, V.I., Kleynerman,V.I., Tret’yakov, B.N., Zaytsev, A.I., “The Coexistence of Ni3(FeMn) and Ni(FeMn) Superlattices in the Iron-Nickel-Manganese System”, Phys. Met. Metallogr., 61(6), 93–96 (1986), translated from Fiz. Met. Metalloved., 61(6), 1136–1139 (1986) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Review, Theory, Thermodyn., 7) Huck, B., Saurenbach, F., Hesse, J., “Investigation of Spin Structure in FCC Fe-Ni-Mn”, Hyperfine Interact., 28(1–4), 479–482 (1986) (Electronic Structure, Experimental, 11) Babayev, Z.M., Matysina, Z.A., Mekhrabov, A.O., “Ordering Temperatures and Order Parameters of Alloy Ni3Fe with Impurity Mn or Cr”, Phys. Met. Metallogr., 64(1), 191–194 (1987), translated from Fiz. Met. Metalloved., 64(1), 202–205, (1987) (Experimental, Theory, Phase Relations, 11) Babayev, Z.M., Valiyev, E.Z., Menshikov, A.Z., Mekhrabov, A.O., “Atomic Ordering in Ni3(Fe1–xMnx) Alloys”, Phys. Met. Metallogr., 64(4), 124–127 (1987), translated from Fiz. Met. Metalloved., 64(4), 762–766 (1987) (Crys. Structure, Experimental, 9) Danilenko, V.M., Turkevich, V.Z., “Thermodynamic Properties of Alloys, and Constitution Equilibrium Diagram of the Fe-Ni-Mn System”, Russ. Metall. (4) 210–212 (1987), translated from Izv. Akad. Nauk SSSR, Met., (4), 209–211 (1987) (Phase Diagram, Phase Relations, Thermodyn., Experimental, 7) Singh, J., Wayman, C.M., “Formation of Secondary Austenite During the α’-γ Transformation in an Fe-Ni-Mn Alloy”, Mater. Sci. Eng., 93, 227–233 (1987) (Experimental, Morphology, Phase Relations, 16) Dubinin, S.F., Mikhaylov, Yu.N., Sidorov, S.K., “Diffuse Scattering of Neutrons in the PreMartensitic Range of Alloys Fe72Mn22Ni6”, Phys. Met. Metallogr., 65(4), 188–190 (1988), translated from Fiz. Met. Metalloved., 65(4), 823–824, (1987) (Experimental, Phase Relations., 1) Dubinin, S.F., Mikhaylov, Yu.N., Skorobogatov, V.P., Sidorov, S.K., “Critical Effects in the Antiferromagnetic Crystal Fe72Mn22Ni6 with Unstable Lattice”, Phys. Met. Metallogr., 66(2), 67–74 (1988), translated from Fiz. Metal. Metalloved., 66(2), 283–291 (1988) (Experimental, Phase Relations, 11) Gomankov, V.I., Zaitzev, A.I., Kleinerman, V.I., “Diagram of Structural States of Alloys of the System Ni-Fe-Mn”, Russ. Metall., (2) 196–200 (1988), translated from Izv. Akad. Nauk SSSR, Met., (2), 204–208 (1988) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 7) Raynor, G.V., Rivlin, V.G., “Fe-Mn-Ni” in “Phase Equilibria in Iron Ternary Alloys”, Inst. Metals, London, 361–362 (1988) (Phase Diagram, Phase Relations, Review, 5) Hari Kumar, K.C., Raghavan, V., “The bcc-fcc Equilibrium in Ternary Iron Alloys - III”, J. Alloy Phase Diagrams, 5(3), 201–220 (1989) (Calculation, Phase Diagram, Phase Relations, 28) Dubinin, S.F., Mikhaylov, Yu.N., Skorobogatov, V.P., “Magnetic and Structural Transformations in Fe72MncNi28–c Alloys”, Phys. Met. Metallogr., 69(6), 64–68 (1990), translated from Fiz. Met. Metalloved., 69(6), 69–73 (1990) (Experimental, Magn. Prop., Phase Diagram, Phase Relations, 12) Herman, F., Jepsen, O., “Electron and Magnetic Structure of the Ternary fcc Mn-Fe-Ni System. II. Disordered Alloys and Ferromagnetic-Antiferromagnetic Interfaces”, Phys. Rev. B,
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[1990Jep]
[1991Gan]
[1991Gok]
[1991Wat]
[1992Wul]
[1994Rag] [1995Hoe]
[1995Tan]
[1995Wat]
[1996Heo]
[1997Sch] [1998Mie]
[1998Pal]
[1999Cha]
[2000Cha]
[2000Kak]
Fe–Mn–Ni 41(10), 6811–6819 (1990) (Electronic Structure, Experimental, Magn. Prop., Phase Relations, 17) Jepsen, O., Herman, F., “Electronic and Magnetic-Structure of the Ternary fcc Mn-Fe-Ni System. 1. Ordered Compounds”, Phys. Rev. B, 41(10), 6801–6810 (1990) (Electronic Structure, Experimental, Magn. Prop., 39) Ganguli, B., Meherotra, V., Gupta, K.P., “Measurement of Order-Disorder Transition Temperature of Some Ni-Fe-Mn Alloys”, Trans. Indian Inst. Met., 44(2), 187–189 (1991) (Crys. Structure, Electric Prop., Magn. Prop., Experimental, 3) Gokcen, N.A., “Mn-Ni (Manganese-Nickel)” in “Phase Diagrams of Binary Nickel Alloys”, Nash, P. (Ed.), ASM Int., Materials Park, OH, 199–206 (1991) (Crys. Structure, Phase Diagram, Phase Relations, Review, 47) Watanabe, S., Sato, S., Nakagami, I., Nagashima, S., “Damping Behavior of Fe-Ni-Mn Alloys” (in Japanese), Tetsu to Hagane, 77(2), 306–313 (1991) (Experimental, Morphology, Magn. Prop., 17) Wulfes, A., Bottger, C.H., Hesse, J., Sievert, J., Ahlers, H., “Magnetic Phase Diagram of the Reentrant Spin Glass System (Fe0.65Ni0.35)1–xMnx”, J. Magn. Magn. Mater., 104–107(III), 2069–2071 (1992) (Abstract, Magn. Prop., Phase Diagram, Phase Relations, 11) Raghavan, V., “Fe-Mn-Ni (Iron-Manganese-Nickel)”, J. Phase Equilib., 15(6), 617–618 (1994) (Phase Diagram, Phase Relations, Review, 14) Hoefer, A., Fricke, M., Boettger, C., Hesse, J., “Resistivity of (Fe0.65Ni0.35)1–xMnx versus Ferro- and Antiferromagnetic Re-Entrant Spin Glass Transitions”, Phys. Solid State A, 148, 551–564 (1995) (Electr. Prop., Magn. Prop., Phase Relations, Experimental, 22) Tanaka, T., Aaronson, H.I., Enomoto, M., “Calculations of α/γ Phase Boundaries in Fe-CX1-X2 Systems from the Central Atoms Model”, Metall. Mater. Trans. A, 26A(3), 535–545 (1995) (Calculation, Phase Diagrams, Phase Relations, Thermodyn., 27) Watanabe, S., Sato, S., Miura, K., “An Unique Feature of Mechanical Property of IronNickel-Manganese Alloy”, J. Mat. Proc. Tech., 53(1–2), 467–476 (1995) (Experimental, Magn. Prop., 14) Heo, N.H., “Theory of Nonequilibrium Segregation in an Fe-Mn-Ni Ternary Alloy and a Ductile-Brittle-Ductile Transition”, Acta Mater., 44(7), 3015–3023 (1996) (Experimental, Magn. Prop., Thermodyn., 25) Schuermann, E., Djurdjevic, M., Nedeljkovic, L., “Liquid Equilibria in the Fe-Ni-Mn System”, Steel Res., 68(12), 512–515 (1997) (Assessment, Phase Diagram, 7) Miettinen, J., “Approximate Thermodynamic Solution Phase Data for Steels”, Calphad, 22(2), 275–300 (1998) (Assessment, Calculation, Phase Diagram, Phase Relations, Thermodyn., 83) Pal, H., Chanda, A., De, M., “Characterisation of Microstructure of Isothermal Martensite in Fe-23Ni-3.8Mn by Rietvield Method”, J. Alloys Compd., 278, 209–215 (1998) (Crys. Structure, Phase Relations, 25) Chanda, A., Pal, H., De, M., Kajiwara, S., Kikuchi, T., “Characterisation of Microstructure of Isothermal Martensite Formed at Low Temperature in Powder of Fe-23Ni-3.3Mn Alloy by Rietveld Method”, Mater. Sci. Eng. A, 265A, 110–116 (1999) (Experimental, Crys. Structure, Morphology, Magn. Prop., Phase Relations, 15) Chanda, A., De, M., Kajiwara, S., “Characterisation of Microstructure of Isothermally Transformed Martensite in Low Temperature in Powder and Bulk State of Fe-23Ni-3.6Mn (mass%) Alloy by Rietveld`s Method”, Jpn. J. Appl. Phys., 39(2A), 539–544 (2000) (Experimental, Crys. Structure, Morphology, Magn. Prop., Phase Relations, 25) Kakeshita, T., Saburi, T., Shimizu, K., “Kinetics of Martensitic Transformations in Some Ferrous and Non-Ferrous Alloys”, Philos. Mag. B, 80B(2), 171–181 (2000) (Experimental, Kinetics, Phase Relations, 36)
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ni [2000Mie]
[2000Rag]
[2001Kak]
[2002Ayd]
[2002Mun]
[2002Tsu]
[2004Mun]
[2004Wil]
[2004Wit]
[2005Akt]
[2005Bel]
[2005Guo] [2006Gor]
[2006Li] [2006Ned]
[2007Kuz]
Landolt-Börnstein New Series IV/11D4
19
Miettinen, J., Howe, A.A., “Estimation of Liquidus Temperatures for Steels Using Thermodynamic Approach”, Ironmaking Steelmaking, 27(3), 212–227 (2000) (Calculation, Phase Relations, Thermodyn., 53) Raghavan, V., “The Martensitic Transformation in Fe-Ni and Fe-Ni-X Alloys at Subzero Temperatures”, Trans. Indian Inst. Met., 53(6), 597–603 (2000) (Review, Phase Relations, Kinetics, 44) Kakeshita, T., Katsuyama, J., Fukuda, T., Saburi, T., “Time-Dependent Nature of Displacive Transformations in Fe-Ni and Fe-Ni-Mn Alloys Under Magnetic Field and Hydrostatic Pressure”, Mater. Sci. Eng. A, 312A(1–2), 219–226 (2001) (Experimental, Kinetics, Phase Relations, 12) Aydin, A., Guler, E., Aktas, H., Gungunes, H., “Mössbauer Studies on Athermal Martensite Formation in an Fe-Ni-Mn Alloy”, Bull. Mater. Sci., 25(5), 359–360 (2002) (Electronic Structure, Experimental, 10) Mun, S.-H., Watanabe, M., Li, X., Oh, K.H., Williams, D.B., Lee, H.-C., “Precipitation of Austenite Particles at Grain Boundaries During Aging of Fe-Mn-Ni Steel”, Metall. Mater. Trans. A, 33A(4), 1057–1067 (2002) (Experimental, Morphology, Phase Diagram, 34) Tsunoda, M., Konoto, M., Takahashi, M., “Magnetic Anisotropy of Antiferromagnet in Exchange Coupled Ni-Fe/Mn-Ni Epitaxial Bilayers”, Phys. Status Solidi A, 189A(2), 449–457 (2002) (Experimental, Magn. Prop., 12) Mun, S.-H., Heo, Y.U., Watanabe, M., Williams, D.B., Lee, H.-C., “Discussion of “Precipitation of Austenite Particles at Grain Boundaries During Aging of Fe-Mn-Ni Steel”. Authors’ Reply”, Metall. Mater. Trans. A, 35A(1), 355–356 (2004) (Experimental, Morphology, Phase Diagram, Phase Relations, 13) Wilson, E.A., “Discussion of “Precipitation of Austenite Particles at Grain Boundaries During Aging of Fe-Mn-Ni Steel””, Metall. Mater. Trans. A, 35A(1), 352–355 (2004) (Morphology, 18) Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diff., 25(4), 346–354 (2004) (Experimental, Phase Diagram, Phase Relations, Calculation, Thermodyn., #, 34) Akturk, S., Durlu, T.N., “The Magnetic Properties of Martensitic Phase in Fe-Ni-Mn Alloys: Mössbauer Spectroscopy Observation”, J. Mater. Sci., 40(22), 6023–6026 (2005) (Electronic Structure, Experimental, Magn. Prop., Morphology, 13) Beladi, H., Hodgson, P.D., Barnett, M.R., “Mapping the Hot Deformation Microstructure of Ni-30Fe Alloy”, ISIJ Int., 45(12), 1893 (2005) (Experimental, Mechan. Prop., Magn. Prop., Morphology, 16) Guo, C., Du, Z., “Thermodynamic Optimization of the Mn-Ni System”, Intermetallics, 13(5), 525–534 (2005) (Phase Diagram, Phase Relations, Assessment, Thermodyn., 42) Gornakov, V.S., Nikitenko, V.I., Shapiro, A.J., Shull, R.D., Yangund, F.Y., Chien, C.L., “Switching of Domains and Domain Walls in Fe50Mn50/Ni81Fe19 Bilayers with non-180° Ferromagnetic Domains”, Phys. Met. Metallogr., 101(Sup.1), S51–S55 (2006) (Experimental, Magn. Prop., Morphology, 10) Li, B.L., Tsuji, N., Minamino, Y., “Microstructural Evolution in 36%Ni Austenitic Steel during the ARB Process”, Mat. Sci. Forum, 512, 73–78 (2006) (Experimental, Morphology, 10) Nedjad, S.H., Ahmadabadi, M.N., Mahmudi, R., Furuhara, T., Maki, T., “Analytical Transmission Electron Microscopy Study of Grain Boundary Precipitates in an Fe-Ni-Mn Maraging Alloy”, Mater. Sci. Eng. A, 438A, 288–291 (2006) (Electronic Structure, Experimental, 7) Kuznetsov, V., “Fe-Ni (Iron-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenbrg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published (2007) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, 41)
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DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
20 [2007Wat]
[2007Wit] [Mas2] [V-C2]
Fe–Mn–Ni Watson, A., Wagner, Z., E. Lysova, E. Rokhlin, L., “Mn-Ni (Manganese-Nickel)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenbrg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published (2007) (Phase Diagram, Phase Relations, Crys. Structure, 13) Witusiewicz, V.T., private communication to MSI, (2007) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_17 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–O
1
Iron – Manganese – Oxygen Pierre Perrot
Introduction The Fe-Mn-O system has raised very soon a huge interest because manganese is a useful element generally found in iron ores. Mn has a beneficial effect on the plasticity of steels, decreases the stacking fault energy, increases the nitrogen solubility and gives a shape memory effect. Mn being easily oxidized gives non metallic inclusions in iron [1981Mik] and their conditions of formation have ben well explored. [1929Sch] showed that MnO has for effect to stabilize iron oxides, thus making harder their reduction by CO. The first FeO-MnO tentative diagram, found in [1930Ben] presents a miscibility gap in the solid state which is considered as doubtful by the author and has not been confirmed since, although [1943Whi, 1964Lev] reproduce it in their reviews. [1937Wen] presents a Fe-FeO-MnO phase diagram accepted today, phase diagram without any miscibility gap, a fact which has been definitively proved by X-Ray diffraction [1946Pet]. Later, this system was investigated with respect to melting, refining and corrosion of stainless steels. The main experimental investigations on phase equilibria and thermodynamics are gathered in Table 1. An assessment of the reactions occurring in the liquid phases may be found in [1993Tur]. No Calphad assessment were carried out. Binary Systems Using new experimental thermodynamic data, the Fe-Mn system has been recently reassessed and the Calphad description was updated by [2004Wit]. This new reevaluation of the phase equilibria leads to consistently better fits to the available experimental data. There is however, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have been written as L ⇌ (δMn) + (γMn,γFe). Consequently, the Fe-Mn system is accepted from [2004Wit]. The Fe-O system is accepted from [Mas2], which reproduces the fundamental work of [1945Dar, 1946Dar]. Two careful assessment may be found in [1991Sun, 1995Kow]. They reproduce well the experimental equilibria and differ by the description of the (Fe,O) liquid solution. The Mn-O binary system is accepted from [2007Fen]. Solid Phases The solid phases are presented in Table 2. Wüstite Fe1–xO gives a complete solid solution “manganowüstite” (Fe1– yMny)1–xO whose x, the departure from stoichiometry decreases when y, the Mn content, increases [1967Sch1]. For y > 0.5, the manganowüstites in equilibrium with iron are stoichiometric [1970Aub]. For pure Fe1–xO, 0.05 < x < 0.12, whereas for pure Mn1–xO, 0 < x < 0.05. The lattice parameter of manganowüstite in equilibrium with a metallic phase (reduced side of the solid solution) and with a spinel phase (oxidized side of the solid solution), measured by [1968Oki] is shown in Fig. 1. Magnetite βFe3O4 (inverse spinel) gives also a complete spinel solid solution with βMn3O4 (normal spinel). In the iron rich part, whose composition can be represented by β(Fe1–yMny)Fe2O4 (0 < y < 1), the structure is of the type “disordered spinel” because Mn2+ ions which have a preference for tetrahedral sites [1969Baf] replace the Fe2+ ions in the octahedral sites of the spinel structure. However, [1985Tay] presents wrongly MnFe2O4 as inverse. A model of the cations distribution amongst the sites of the spinel structure, based on site preference energy, may be found in [1979Pel, 1983ONe, 2003Mue]. These iron rich spinels may be easily obtained by mixing Fe2O3 and MnO2 at 950°C under a (CO2 + 1% CO) atmosphere whereas the mixing under pure CO2 leads to the bixbyite (Fe,Mn)2O3 solid solution. The crystal parameter of the spinel solution β(Fe1–yMny)Fe2O4 (0 < y < 1) increases linearly with the Mn content [1984Pun, 1995All, 2005Zin] according to the Vegard’s law. Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_18 # Springer 2008
2
Fe–Mn–O
The non stoichiometry of the Mn ferrite MnFe2O4 has been investigated by [1983Tsu] at 1150°C and by [1998Bon] between 400 and 900°C under various oxygen pressures. Below 3 Pa O2, the parameter remains constant at 861.2 pm, corresponding to the stoichiometric spinel. Between 3 and 300 Pa O2, the crystal parameter decreases down to 850.5 pm corresponding to the hyperstoichiometric, metal defective spinel. Above 300 Pa O2, the hematite phase forms and the parameter of the remaining spinel increases, which corresponds to an increase of the Mn content of the spinel. Below 550°C, and at high oxygen partial pressure, MnFe2O4 oxidizes into MnFe2O4+δ metastable. A value of δ = 0.5 corresponds to the γ(Mn0.5Fe)2O3 metastable phase [1998Bon]. As a consequence, it is hard to define an upper limit of stability for MnFe2O4 at low temperatures. In the manganese rich part of the solid solution, whose composition may be represented by Mn(Fe1–yMny)2O4 (0 < y < 1), the spinel solid solution is direct [1986Maz], with the Mn2+ ions in the tetrahedral sites and the Mn3+ and Fe3+ ions in the octahedral sites. The Mn rich part of the solid solution is tetragonally deformed [1969Wic, 1971Bra, 1973Hol, 1995Bat] as shown in Fig. 2. The origin of the tetrahedral distortion of the lattice for the Mn rich spinels is the effort of Mn3+ ions to remove their orbitally degenerated ground state (3d4) [2000Pol]. This is the cooperative Jahn-Teller effect on Mn3+ ions in octahedral sites which provokes a lowering of the symmetry of their surroundings. The cubic-tetragonal transition varies from 200°C when y = 0.5 to 1171°C when y = 1, which corresponds to Mn3O4. At 940°C, the cubic spinel Mn2.54Fe0.46O4 is in equilibrium with the tetragonal spinel Mn2.73Fe0.27O4 [1973Hol]. The Mn rich spinel are stable under oxygen pressures higher than those generated by pure CO2. The spinel solid solutions may be metal defective (Fe1–yMny)3(1–x)O4 under higher oxygen pressures [1969Wic]. The departure from stoichiometry of these spinels increases with temperature and with iron content. The lattice parameter of some spinel (0.27 < y < 0.32), investigated by [1974Tan], increases with the oxygen pressure when x < 0, presents a maximum which corresponds to the stoichiometric composition (x = 0), then decreases, which corresponds to x > 0 (the ratio O/Metal in the spinel is then higher than 4/3). The valences of the elements (Fe2+, Fe3+, Mn2+, Mn3+) in the spinel phase have been discussed in [1967Sch2]. αFe2O3 (stable under the form of hematite) and αMn2O3 (stable under the form of bixbyite) present limited solubility, although Mn2O3 may dissolve Fe2O3 up to 50 mol% at 800°C and 70 mol% at 1020°C. γMn2O3 decomposes slowly at 1000°C to give Mn3O4 through αMn2O3 as an intermediate phase. The solid solutions γ(Fe1–yMny)2O3 (y > 0.45) prepared by co-precipitation in aqueous medium [1990Pat] decomposes at lower temperatures to give MnFe2O4 and Mn2O3 which transforms into Mn3O4 at 1000°C. Quasibinary Systems The αFe2O3-βMn2O3 quasibinary system under 0.21 kPa of oxygen pressure (air atmosphere), established by [1969Wic] and reproduced in [1980Bra] is shown in Fig. 3. At higher temperatures, αFe2O3 and βMn2O3 are not stable. Under air atmosphere, αFe2O3 loses oxygen and gives βFe3O4 at 1380°C and βMn2O3 transforms into αMn3O4 at 977°C. Above 1380°C, the solid phase in equilibrium with the liquid, is always the β(Fe,Mn)3O4 spinel solid solution [1958Hoo]. It is not clear whether or not the liquid phase is in the same oxidation state as the solid. At lower temperatures (< 877°C), Mn and Fe are in the state of oxidation +3 and the only stable solid solutions are the manganese rich (Mn,Fe)2O3 and iron rich (Fe,Mn)2O3 solid solutions. Between 877 and 1380°C, the (Fe,Mn)2O3 solid solutions may be in equilibrium with the (Fe,Mn)3O4 spinel solutions [1962Mua]. Metastable equilibria may easily been observed [1996Gol] by sintering pellets of Fe2O3-Mn2O3 mixtures under 10 MPa followed by quenching, annealing or slow cooling. The FeO-MnO quasibinary diagram in equilibrium with the (Fe,Mn) metallic phase, experimentally determined by [1977Oet] and calculated by [1993Wu] is presented in Fig. 4, and the metallic phase in equilibrium with the solid manganowüstite (MW) or with the liquid oxide (L2) is shown. The manganowüstite (Fe1–yMny)1–xO in equilibrium with the metallic phase is characterized by x = 0.05 for y = 0. When the Mn content increases, x decreases and becomes 0 for y ∼ 0.5. For y > 0.5 the manganowüstite is stoichiometric (x = 0).
DOI: 10.1007/978-3-540-78644-3_18 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–O
3
Liquidus Surface [1989Rag] presents a tentative liquidus projection which is not reproduced in this report because the tie lines suggested are not credible. Actually, experimental data are scarce and the position of the invariant points are unknown. The two-liquid domain (Fe+FeO) originated along the Fe-O side continues through the diagram up to the Mn-O side, but the tie lines (Fe,Mn)-(Fe,Mn)O have never been determined. As Mn is more easily oxidized than Fe, it is probable that Mn goes preferentially in the liquid oxide phase and Fe in the metallic phase. Isothermal Sections The oxygen solubility in liquid iron first investigated by [1950Hil] was shown to decrease with the manganese content. However, it is not clear wether the oxygen solubility in the Fe-Mn alloys present a minimum [1973Buz, 1976Jan]. [1981Ave] carried out oxygen solubility measurements between 1600 and 1700°C under H2-H2O atmospheres and claimed the existence of a minimum in the oxygen solubility towards 5 mass% Mn, but without a convincing experimental evidence. More credible are the results presented by [1987She] which, through oxygen solubility measurements towards the Mn rich end of the alloy, show the presence of a minimum of solubility at 0.015 mass% O in a Fe + 9.6 mass% Mn at 1600°C under 0.1 MPa of oxygen pressure. The oxygen solubility shown in Fig. 5 from [2000Tak] represent the oxygen content of the liquid alloys in equilibrium with the (Fe,Mn)O liquid slag in the iron rich corner or with the manganowüstite solid solution in the manganese rich corner. Calculations carried out by [1998Lee, 2004Jun] did not take into account the existence of a minimum. The isothermal section at 1600°C, mainly from [1977Oet] is represented in Fig. 6 and shows the isobaric curve 0.1 MPa O2 proposed by [1950Gur] and reproduced by [1964Lev]. The tie lines in the two-phase domain manganowüstite (Fe,Mn)O - spinel (Fe,Mn)3O4, were investigated at 1300°C [1967Sch1], 1200°C [1993Sub, 1994Sub], 1150°C [1967Sch1], 1000°C [1965Tre, 1975Duq, 2005Zin], 987°C [1968Oki] and 700°C [1984Pun]. Figure 7 gives isothermal section at 1300°C [1967Sch1, 1989Rag]. Figure 8 taken from [1967Sch1], gives the position of each tie line together with the oxygen pressure at equilibrium at 1300 and 1150°C. The potential diagram taking into account the equilibria between the solid solutions (metal, manganowüstite, spinel, hematite and bixbyite) at 900°C is shown in Fig. 9 from [1979Pel]. Oxygen pressure at equilibrium has been measured by equilibration with a CO-CO2 [1967Sch1] or with a CO2-H2 [1968Oki, 1984Pun] atmosphere. Figure 10 represents the manganowüstite field at 1150°C together with the isobaric curves [1967Sch1]. Under higher oxygen pressures, manganowüstite oxidizes into spinel. Under lower oxygen pressures, manganowüstite is reduced into a metallic phase constituted mainly from iron. Thermodynamics The thermodynamic quantities of formation of MnFe2O4 from its oxides [1979Sek] are given in Table 3 together with the enthalpy of dissolution of O2 in the spinel solid solution [1974Tan] and the Gibbs energy of dissolution of MnO in liquid iron [1991Mik, 2003Per]. Heat capacities of MnFe2O4, Mn1.5Fe1.5O4 and Mn2FeO4 which belong to the solid solution (Fe1–yMny)3O4, measured between 200 and 740 K by [1981Nai], are presented in Table 4. The ferrimagnetic-paramagnetic transition, observed around 550 K (180 ± 10°C) for these three compounds is characterized by a λ shape accident in the heat capacity vs temperature curves and the transition entropy was evaluated at 55 ± 5 J·K–1·mol–1 for each compound, with a slight tendency to decrease with the Mn content of the spinel solid solution. The activities of FeO in the manganowüstite solid solutions in equilibrium with Fe has been evaluated from the oxygen pressures at equilibrium Fe/(Fe1–yMny)1–xO, the oxygen pressures being generated by CO-CO2 atmospheres [1929Sch, 1971Ono, 1990Fra], H2-H2O atmospheres [1970Aub, 1970Bul, 1991Ban], H2-CO2 atmospheres [1968Oki] or measured from the emf of a solid electrolyte cell [1965Tre]. Authors agree to propose a slight positive departure towards ideality, which is confirmed by the calculations of [1983Dav] based on the volume mismatch between FeO and MnO. The FeO-MnO solid solution is described by a regular model with an energy parameter α = 5250 J·mol–1 at 1150°C. For calculating the FeO-MnO diagram in equilibrium with a (Fe,Mn) alloy, [1993Wu] supposes the liquid ideal (α = 0) and proposes α = 6126 J·mol–1 Landolt-Börnstein New Series IV/11D4
MSIT®
DOI: 10.1007/978-3-540-78644-3_18 # Springer 2008
4
Fe–Mn–O
between 1380 and 1850°C, in good agreement with the experimental diagram. A thorough investigation [1972Bal] showed that, for a given ratio Fe/Mn of the solid solution, the FeO activity decreases with the departure from stoichiometry. The liquid (Fe,Mn)O solution in equilibrium with iron behaves quasi ideally at 1500°C [1991Ban]. The manganowüstites in equilibrium with a spinel phase behaves ideally. Activities of FeO and MnO inside the manganowüstite field have been calculated by [1988Lyk, 1990Fra] from the oxygen potential vs composition experimental curves. Surprisingly, [1997Pet] calculates, for the manganowüstite, spinel and hematite solid solutions, a strong negative departure from ideality. Such a result may be explained, as pointed out by the authors [1996Gol, 1997Pet], by a non equilibrium state and an ill controlled oxygen pressure. A thermodynamic model has also been proposed [1993Tyu] for the spinel solid solution. The activities of Fe3O4 in the spinel solid solution have been also evaluated from the oxygen pressures at equilibrium spinel/manganowüstite under various CO-CO2 gaseous atmosphere [1967Sch1, 1971Ono, 1990Fra], H2-CO2 atmospheres [1983Ter] or from oxygen pressure at the spinel/corundum equilibrium measured directly [1970Roe, 1971Ono, 1988Yoo] provided that oxygen pressures lie above 1 kPa. For instance, the upper domain of stability of the spinel (Mn0.482Fe0.518)Fe2O4, that is the oxygen pressure at equilibrium spinel-bixbyite, measured by [1988Yoo] in the temperature range 1000–1200°C is given by: log10 (pO2/bar) = (11.7 ± 0.6) – (1.87 ± 0.08) / T The spinel solid solutions present a large positive departure from ideal behavior [1979Sek, 1990Fra]. The interaction coefficients of Mn and O in liquid iron at 1600°C, (< 2 mass% Mn) have been evaluated by [1966Sch, 1974Sig, 1991Mik] as: eO(Mn) = {∂ log10 fO / ∂ (mass% Mn)} = – 0.021 at 1600°C, with fO = {(mass% O in pure Fe) / (mass% O in alloy)}, as eO(Mn) = – 0.040 ± 0.005 by [1967Buz], as eO(Mn) = – 0.054 by [1981Mik] and as eO(Mn) ∼ 0 by [1972Mat, 1976Jan] at the same temperature. Actually, Mn in liquid Fe increases slightly the oxygen solubility and the value accepted by [2003Per] is eO(Mn) = – 0.025. [1981Ave] pointed out the difficulties arising in the measurements at high temperatures because of the easy volatilization of Mn but presents interaction parameters in good agreement with the preceding ones. [2000Tak, 2002Hin] showed that the value eO(Mn) = – 0.037, although less precise at infinite dilution, may be used up to 20 mass% Mn. On the other hand, [1970Nis] used the oxygen distribution between solid and liquid alloy to calculate eO(Mn) = – 1.83 in the (δFe). Mn increases strongly the oxygen solubility in (δFe). Notes on Materials Properties and Applications The main experimental works are summarized in Table 5. Mn may be used as deoxidizer in steelmaking [1972Mat]. Although it is not so strong than Al or Si, it has many beneficial effects on the plasticity of steel and it represents, with Ni, the main constituent of High Nitrogen Steels. Mn increases the plasticity of steels but decreases the machinability [2004Umi] because the formation of oxide films with high melting temperatures reduces the lubrification effects. The spinel solid solution (Fe1–yMny)Fe2O4 may be used as soft magnetic material [1993Alc] and used in the fabrication of components such as high density magnetic core or read/write heads for high speed tapes or magnetic recording [2005Pad]. MnFe2O4 may be prepared by solid state reaction between MnCO3 and Fe2O3 [2001Der] at 900–1100°C. Li2O may be added under the form of LiNO3 as dopant for stimulating the ferrite formation. The activation energy for the ferrite formation goes down from 174 kJ·mol–1 without catalyst to 117 kJ·mol–1 with the addition of 3.5 mol% Li2O. Mechanical alloying of αFe2O3 (hematite) and βMn2O3 (bixbyite) forms βFeMnO3, a ferrimagnetic insulator with a magnetic moment of 2.8 μB per formula unit and a Curie temperature around 40 K [2000Sei]. The size of the crystals lie between 10 and 20 nm. Mechanical alloying of MnO + Fe2O3 mixtures forms MnFe2O4, but a prolonged milling time may induce a reduction due to metallic iron originating from ball debris, thus transforming MnFe2O4 into manganowüstite [2005Pad]. Mechanical alloying of (1–y) Fe2O3 + y Mn2O3 (y > 0.3) forms only bixbyite like crystals [2004Med] whereas the mixture (1–y)Fe + y Mn2O3 gives a spinel phase for y > 0.6 and a manganowüstite phase for y < 0.6. Nanosized particles may also be obtained [2001Mur] from the displacement reaction: 2 FeCl3 + 4 MnO → MnFe2O4 + 3 MnCl2
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5
activated by high energy ball milling. The spontaneous magnetization of the nanopowder obtained is lower than that of the bulk and decreases with the nanoparticle size. Below 50 K, the coercitivity increases sharply with the decreasing temperature. Colloid nanocrystals (Fe,Mn)3O4 have been prepared [2002Sar] by chemical precipitation of FeCl2+FeCl3+MnCl2 in basic aqueous medium (pH > 12). The MnFe2O4 nanocrystals synthesized by radiofrequency plasma torch [2006Swa] exhibit a cuboctahedral morphology having only (111) and (110) surfaces. A new phase MnFe2O4–δ (δ unknown), disordered spinel characterized by a ratio c/a < 1 has been prepared by reacting MnO2 + Fe2O3 at 1200°C [2001Hu]. Its Curie temperature is TC ∼ 400°C (the Curie temperature of the normal cubic spinel is ∼ 280–297°C) and it possesses a negative magnetoresistance. Miscellaneous The electrical conductivity of MnFe2O4 sintered at 1250°C, measured at 100–200°C [1972Rez] was shown to decrease with time, a behavior attributed to a redistribution of metallic ions on the tetrahedral and octahedral sites. The Curie point lies at about 300°C, and seems to shift towards 350°C through the oxidation process of MnFe2O4 [1986Maz]. The Fe3+ of the maghemite γFe2O3 may be substituted by Mn3+ or Mn2+. The substitution by Mn2+ introduces vacancies [1984Gil] which are ordered if the substitution rate is low. Metal Ceramic microstructures may be obtained by partial reduction of (Fe1–yMny)O manganowüstites [1995Sub]. The presence of refractory oxides (Al2O3, Cr2O3, ZrO2, alkaline earth oxides) as dopants enhances the rate of the metal precipitation during the reduction. The precipitation of the metallic phase, a rich iron alloy, occurs mainly in grain boundaries if Al2O3 or Cr2O3 are present as dopants. Table 1.
Investigations of the Fe-Mn-O Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1929Sch]
Reduction isotherm under CO-CO2 atmospheres
600–800°C, Fe2O3-MnO mixtures
[1930Ben]
Thermal analysis, metallography
1300–1650°C, Mn-O and FeO-MnO phase diagrams
[1937Wen]
Thermal analysis, metallography
1300–1650°C, Fe-FeO-MnO phase diagram
[1946Pet]
X-ray diffraction
FeO-MnO solid solution
[1950Gur]
Isobaric curves determination by weight loss measurements
1600°C, Fe2O3-FeO-MnO, < 0.1 MPa of oxygen pressure
[1950Hil]
Oxygen solubility measurements, chemical analysis
1550–1650°C, < 2 mass% Mn, < 0.2 mass% O
[1958Hoo]
Thermal analysis, X-ray diffraction
< 1600°C, Fe3O4-Mn3O4 phase diagram
[1962Mua]
Thermal analysis, X-ray diffraction
800–1585°C, Fe-Mn-O phase diagram under 0.21 kPa O2
[1965Tre]
Powder sintering, X-ray diffraction, EMF measurements
1000°C, from 10 nPa to 10 mPa of oxygen pressure
[1967Sch1]
XRD, chemical analysis and thermogravimetry
1000–1300°C, CO-CO2 atmospheres, spinel + manganowüstite field (continued)
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6
Fe–Mn–O
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1968Oki]
XRD, chemical analysis, Neel temperatures measurements
987°C, equilibrium Fe-(Fe,Mn)O-(Fe,Mn) Fe2O4 under CO2-H2
[1969Wic]
XRD, chemical analysis
< 1400°C, Fe3O4-Mn3O4 solid solutions under air atmosphere
[1970Aub]
XRD, chemical analysis
< 1000°C, (Fe,Mn)1–xO solid solutions under H2-H2O atmospheres
[1970Bul]
XRD, chemical analysis
969–1000°C, (Fe1–yMny)O (0.2 < y < 0.5) under H2-H2O atmospheres
[1970Nis]
Oxygen distribution between solid and liquid alloys
1534–1600°C, < 1.6 mass% Mn, < 0.2 mass% O
[1970Roe]
XRD, thermogravimetry
1000–1400°C, spinel-corundum equilibrium, 1 to100 kPa O2
[1971Bra]
XRD, crystal parameters, thermal expansion, conductivity
< 1400°C, (Fe1–yMny)3O4, cubic-tetragonal transition
[1971Ono]
Thermogravimetry
800–1300°C, Fe-(Fe,Mn)O-(Fe,Mn)3O4-(Fe, Mn)2O3 equilibria
[1973Hol]
XRD, cubic-tetragonal equilibrium
940°C, (Fe1–yMny)3(1–x)O4,
[1974Tan]
XRD, crystal parameters vs oxygen pressure
1200–1350°C, (Fe1–yMny)3(1–x)O4, 0.27 < y < 0.32, 1 Pa to 100 kPa O2
[1975Duq]
XRD, thermogravimetry, electrical conductivity
1000°C, 0 to 100 mass% Mn, 10–18 < pO2 / bar < 1
[1976Jan]
EMF from ZrO2(CaO) coated with ThO2(Y2O3) cell
1600°C, < 20 mass% Mn, < 0.2 mass% O
[1977Oet]
Thermal analysis, equilibrium measurements
1370–1875°C, CO-CO2 atmospheres, equilibrium alloy-manganowüstite
[1979Kop]
XRD, EMF and solubility measurements
1600°C, < 10 mass% Mn, < 0.2mass% O
[1981Ave]
Oxygen analysis in alloys under H2-H2O (< 40 % H2O)
1600–1700°C, < 9 mass% Mn, < 0.2 mass% O,
[1981Mik]
Chemical analysis, solubility measurements
1525–1600°C, < 10 mass% Mn, < 0.2 mass% O, 0.1 MPa O2
[1981Nai]
Calorimetry, Curie temperature measurements
200–740 K, (Fe1–yMny)3O4 (y = 0.33, 0.50, 0.67)
[1983Ter]
Thermogravimetry, EMF with stabilized ZrO2 as solid electrolyte
880–1000°C, (Fe1–yMny)3O4 under CO2-H2 atmospheres
[1983Tsu]
Thermogravimetry, electrical conductivity
950–1120°C, Fe/Mn = 2, O2 pressure between 0.1 Pa and 0.1 MPa
[1984Pun]
Thermogravimetry
700°C, (Fe,Mn)3O4-(Fe,Mn)2O3 tie lines under CO2-H2 atmospheres (continued)
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7
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1987She]
Chemical analysis, solubility measurements
1600°C, < 87 mass% Mn, < 0.2 mass% O, 0.1 MPa O2
[1990Fra]
Thermogravimetry
1000°C, (Fe1–yMny)1–xO under CO-CO2 atmospheres
[1991Ban]
Chemical analysis, metal-slag equilibria
1450–1500°C, < 20 mass% MnO, H2-H2O atmospheres
[1992Sub, 1993Sub]
Thermogravimetry, electrical conductivity
1000–1200°C, (Fe1–yMny)1–xO under COCO2 or N2-O2 atmospheres
[1995All]
XRD, lattice parameters measurements
25°C, β(Fe1–xMnx)Fe2O4 spinel solid solution (0 < x < 1)
[1995Bat]
XRD, electrical conductivity, Mössbauer
(Fe1–yMny)3O4 prepared by coprecipitation (0.35 < y < 1)
[1996Gol, 1997Pet]
High temperature XRD
600–1400°C, sintered Fe2O3-Mn2O3 samples, then annealed or quenched
[1998Bon]
XRD, Mössbauer, saturation magnetization, coulometry
400–900°C, (MnFe2)1–xO4, 10-20 < pO2 / bar < 0.21
[2000Tak, 2002Hin]
ERD, IR spectroscopy analysis, solubility measurements
1550–1650°C, < 24 mass% Mn, < 0.2 mass% O
[2005Zin]
XRD, chemical analysis
1000°C, manganowüstite-spinel equilibrium under H2-H2O
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(δFe) 1538 – 1394
cI2 Im 3m W
a = 293.15
at 1390°C [V-C2, Mas2]. Dissolves 10 at.% Mn at 1474.5°C [2004Wit]
(δMn) 1246 – 1138
cI2 Im 3m W
a = 308.0
[Mas2]. Dissolves 12.2 at.% Fe at 1236.6°C [2004Wit]
(αFe) (Ferrite) < 912
cI2 Im 3m W
a = 286.65
pure Fe at 25°C [Mas2, V-C2]. Dissolves 5 at.% Mn at 527°C [2004Wit]
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, > 13 GPa [Mas2]
(continued)
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Fe–Mn–O
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
γ(Fe,Mn)
cF4 Fm 3m Cu
(γFe) (Austenite) 1394–912 (γMn) 1138–1087
Lattice Parameters [pm]
Comments/References
a = 364.67
at 915°C [V-C2, Mas2]
a = 386.0
[Mas2]
(βMn) 1087–707
cP20 P4132 βMn
a = 631.52
[Mas2]
(αMn) < 707
cI58 I 43m (αMn)
a = 891.26
pure Mn at 25°C [Mas2]
(Fe,Mn)O (Manganowüstite)
cF8 Fm 3m NaCl
Fe1–xO (Wüstite) 1422–569 MnO < 1850 Fe3O4 (r) < 580
oP56 Pbcm Fe3O4 I
β(Fe,Mn)3O4
cF56 Fd 3m MgAl2O4 Inverse spinel
βFe3O4 (h) (Magnetite) 1597–580 β(Fe1–xMnx)Fe2O4
βMnFe2O4 βMn3O4 1680–1171
Disordered spinel
Normal spinel Normal spinel
actually (Fe,Mn)1-xO with x < 0.12 a = 431.0 a = 429.3 a = 444.5
x = 0.05 x = 0.12 [2007Fen]. Actually Mn1–xO (x < 0.1 at 1200°C [1993Sub]
a = 1186.8 b = 1185.1 c = 1675.2
[V-C2]
spinel solid solution
a = 854.5 a = 839.6
at 1000°C [2007Fen] at 25°C [1995All]
a a a a a a
x = 0.2 [1995All] x = 0.4 [1995All] x = 0.6 [1995All] x = 0.8 [1995All] x = 1 [1995All] at 1200°C [2007Fen]
= = = = = =
841.9 844.6 847.0 849.5 850.2 862.6
(continued)
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Fe–Mn–O
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
αMn(Fe1–yMny)2O4
tI28 I41/amd αMn3O4 (Hausmannite)
αMn3O4 (r) < 1171 MnFe2O4–δ
-
(Fe,Mn)2O3
hR30 R 3c αAl2O3 (Corundum)
αFe2O3 (Hematite) < 1451 αMn2O3 < 2330 β(Fe,Mn)2O3 βFe2O3
γFe2O3 (Maghemite) γMn2O3
Comments/References
distorted spinel with c/a > 1 (y > 0.475) [1995Bat] a = c = 854.1 a = 824.3 c = 918.4 a = 815.5 c = 948.4
y = 0.475 [1995Bat] y = 0.710 [1995Bat]
a = 919.03 c = 893.96
distorted spinel with c/a < 1 [2001Hu]
y = 1 [1995Bat]
sesquioxide solid solution
a = 503.42 ± 0.03 c = 1374.83 ± 0.04
dissolves 10 mol% Mn2O3 at 1000°C metastable phase
cI80 Ia 3 βMn2O3 (Bixbyite)
βFeMnO3 βMn2O3 (Bixbyite) γ(Fe,Mn)2O3
Lattice Parameters [pm]
9
cF56 Fd3m MgAl2O4
a = 939.3 ± 0.2
[V-C2]
a = 940.7 a = 940.8
[2000Sei] stable dissolves up to 70 mol% Fe2O3 at 1000°C [1969Wic] metastable phase [1989Rag]
a = 834
[1989Rag] maghemite tetrahedrally distorted [1990Pat]
αMnO2 (Pyrolusite)
tP6 P42/mnm TiO2
a = 439.83 c = 287.30
[1989Rag]. Many polymorphs are known
βMnO2 (Cryptomelane)
tI24 I4/m -
a = 981.5 b = 284.7
[1989Rag] diaspore like structure
γMnO2 (Ramsdellite)
oP12 Pnma -
a = 927.0 b = 286.6 c = 453.3
[1989Rag] disordered diaspore like structure
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10 Table 3.
Fe–Mn–O Thermodynamic Data of Reaction or Transformation
Reaction or Transformation
Temperature [°C]
Quantity, per mol of atoms [J, mol, K]
Comments
MnO + Fe2O3 ⇌ MnFe2O4
25 25 25–900
ΔrH° = – 16 300 ΔrS° = 5.02 ΔrCp = – 30.0 + 0.0269 T + 1.56·106 T –2
[1979Sek]
1/2 O2 ⇌ {O} dissolved in (Fe1–yMny)3O4
1200–1350
ΔrH° = – 314000 (for y = 0.27) ΔrH° = – 331000 (for y = 0.30) ΔrH° = – 308000 (for y = 0.33)
[1974Tan]
MnO ⇌ {Mn} + {O} dissolved in liquid Fe
Table 4.
1500–1800
[2005Zin]
ΔrG° = 245250 – 108.74 T Standard state: 1 mass% in liquid Fe
[2003Per]
Thermodynamic Properties of Single Phases
Phase
Temperature Range [°C]
Property, per mole of atoms [J, mol, K]
Comments
(1/7)(MnFe2O4)
27 467
Cp = 23.49 Cp = 27.61
[1981Nai]
(1/7)(Mn1.5Fe1.5O4)
27 467
Cp = 23.29 Cp = 27.56
[1981Nai]
(1/7)(Mn2FeO4)
27 467
Cp = 23.57 Cp = 27.50
[1981Nai]
Table 5.
Investigations of the Fe-Mn-O Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1953Gal]
Magnetic moment measurement
25°C, Fe3O4-Mn2FeO4 solid solution
[1967Mar]
Magnetic hysteresis curves
< 400°C, Mn0.2Fe2.8O4
[1967Sch2]
Chemical analysis
Determination of Mn2+ and Mn3+ in manganese ferrites
[1969Baf]
XRD, neutron diffraction
Distribution of Mn2+ and Mn3+ ions in the spinel structure
[1972Rez]
Electrical conductivity vs time
MnFe2O4 sintered at 1250°C
[1978Sid]
XRD, SEM (Scanning Electron Microscopy)
Segregation of Mn in (Fe,Mn)Fe2O4 (< 1 at.% Mn)
[1984Gil]
IR spectrometry
Order-disorder transition in γFe2O3 substituted by MnO
[1985Tay]
XRD, dilatometry
–182–480°C (91–753 K), MnFe2O4, thermal expansion (continued)
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Reference
Method/Experimental Technique
Type of Property
[1986Maz]
XRD, electrical conductivity, Curie point
25–400°C, MnFe2O4, influence of the preparation conditions
[1988Yoo]
Conductivity, thermoelectric power measurements
1000–1200°C, (Mn0.482Fe0.518)Fe2O4, 10-14 < pO2 / bar < 1
[1990Pat]
XRD, IR spectroscopy
< 1000°C, γ(Fe1–yMny)2O3 (y > 0.45), preparation and decomposition
[1995Sub]
SEM, Chemical Analysis
Preparation of metal (Fe) ceramic (Fe1–yMny)O structures
[2000Sei]
XRD, SEM, magnetic moment, Curie temperature, Mössbauer
MnFeO3 prepared by mechanical alloying
[2001Der]
XRD, DTA
< 1100°C, MnCO3 + Fe2O3 (+LiNO3 as dopant), synthesis of MnFe2O4
[2001Hu]
XRD, thermal magnetic analysis, magnetoresistance
1200°C, MnO2 + Fe2O3 → new tetragonal ferrite MnFe2O4–δ
[2001Mur]
XRD, TEM, spontaneous magnetization, coercitivity
MnFe2O4 from displacement reaction activated by high energy ball milling
[2002Sar]
XRD, EDX, TEM, electron diffraction, magnetization
Colloids (Fe,Mn)3O4 precipitated in aqueous medium (pH > 12)
[2004Med]
XRD, Mössbauer, magnetization
(Fe1–yMny)2O3 (0.3 < y < 0.9) and {(1-y) Fe+yMn2O3} ball milled
[2004Umi]
Metallography, hardness
< 1.5 mass% Mn, machinability investigations
[2005Pad]
XRD, X-ray fluorescence analysis
MnO + Fe2O3 ball milled, preparation of MnFe2O4 nanocrystalline
[2006Swa]
High Resolution TEM,
MnFe2O4 nanocrystals synthesized by radiofrequency plasma torch
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Fig. 1. Fe-Mn-O.
Fe–Mn–O
Crystal parameter of the manganowüstite (Fe1–yMny)1–xO
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Fig. 2. Fe-Mn-O.
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Crystal parameters of the solid solutions α(Fe,Mn)3O4 (Hausmannite) and β(Fe,Mn)3O4 (Spinel)
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Fig. 3. Fe-Mn-O.
Fe–Mn–O
The Fe2O3-Mn2O3 quasibinary system under air atmosphere
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Fig. 4. Fe-Mn-O. The FeO-MnO quasibinary system in equilibrium with a metallic phase. L1: liquid metal, L2: liquid oxide, MW: manganowüstite
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Fig. 5. Fe-Mn-O. mass%)
Fe–Mn–O
Oxygen solubility in liquid (Fe,Mn) alloys under 0.1 MPa of oxygen pressure (solubility in
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Fig. 6. Fe-Mn-O.
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Isothermal section at 1600°C showing, in the L2 domain, the isobaric curve 0.1 MPa O2
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Fig. 7. Fe-Mn-O.
Fe–Mn–O
Isothermal section at 1300°C
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Fig. 8. Fe-Mn-O.
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Oxygen pressure for the equilibrium Spinel (Sp) – Manganowüstite (MW) at 1300 and 1150°
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Fig. 9. Fe-Mn-O.
Fe–Mn–O
Oxygen pressure for the equilibrium Spinel (Sp) – Manganowüstite (MW) at 987 and 700°C
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Fig. 10. Fe-Mn-O.
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The manganowüstite (MW) field at 1150°C with isobaric curves given in log10 (pO2 / bar)
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22 References [1929Sch]
[1930Ben] [1937Wen] [1943Whi] [1945Dar]
[1946Dar]
[1946Pet]
[1950Gur]
[1950Hil]
[1953Gal]
[1958Hoo] [1962Mua] [1964Lev] [1965Tre]
[1966Sch]
[1967Buz]
[1967Mar] [1967Sch1]
[1967Sch2]
Fe–Mn–O
Schenck, R., Franz, H., Willeke, H., “Investigations of Equilibrium of the Reduction, Oxidation and Carbonization Processes of Iron” (in German), Z. Anorg. Allg. Chem., 184, 1–38 (1929) (Experimental, Phase Relations, 7) Benedicks, C., Loefquist, H., “The Binary System Mn-O”, Non-Metallic Inclusions in Iron and Steel., 72–73 (1930) (Experimental, Morphology, Phase Diagram, Phase Relations, 2) Wentrup, H., “The Phase Diagram of Sulfide Inclusions in Steel” (in German), Techn. Mitt. Krupp, 5, 131–152 (1937) (Morphology, Phase Diagram, Experimental, Phase Relations, 27) White, J., “The Physical Chemistry of Open-Hearth Slags”, J. Iron Steel Inst., London, 148, 579–694 (1943) (Phase Diagram, Phase Relations, Review, 195) Darken, L.S., Gurry, G.W., “The System Iron-Oxygen. I. The Wuestite Field and Related Equilibria”, J. Am. Chem. Soc., 67, 1398–1412 (1945) (Experimental, Phase Diagram, Thermodyn., *, 26) Darken, L.S., Gurry, G.W., “The System Iron-Oxygen. II. Equilibria and Thermodynamics of Liquid Oxides and Other Phases”, J. Am. Chem. Soc., 68, 798–816 (1946) (Experimental, Phase Diagram, Thermodyn., *, 24) Pettersson, H., “Investigations of the Solubility in the Solid State in the Binary Systems of the Oxides CaO, MnO, MgO and FeO” (in Swedish), Jernkontorets Ann., 130(12), 653–663 (1946) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 19) Gurry, R.W., Darken, L.S., “The Composition of CaO-FeO-Fe2O3 and MnO-FeO-Fe2O3 Melts at Several Oxygen Pressures in the Vicinity of 1600°C”, J. Am. Chem. Soc., 72, 3906–3910 (1950) (Experimental, Phase Relations, Phase Diagrams, 10) Hilty, D.C., Crafts, W., “Solubility of Oxygen in Liquid Iron Containing Silicon and Manganese”, Trans. Am. Inst. Min. Metall. Eng., 188, 425–436 (1950) (Phase Relations, Experimental, 18) Gal’perin, F., “Magnetic Moments and the Crystal Structure of Ferromagnetic Metals and Alloys” (in Russian), Dokl. Akad. Nauk SSSR, 88, 643–644 (1953) (Crys. Structure, Magn. Prop., Experimental, 11) Van Hook, H.J., Keith, M.L., “The System Fe3O4-Mn3O4”, Am. Mineral., 43, 69–83 (1958) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 20) Muan, A., Somiya, S., “The System Iron Oxide - Manganese Oxide in Air”, Amer. J. Sci., 260, 230–240 (1962) (Experimental, Phase Diagram, Phase Relations, 16) Levin, E.M., Robbins, C.R., McMurdie, H.F., Phase Diagrams for Ceramists, Am. Ceram. Soc., Columbus, Ohio, 1, Fig. 70–74 (1964) (Phase Diagram, Phase Relations, Review, 2) Tret`yakov, Yu.D., Saksonov, Yu.G., Gordeev, I.V., “Phase Diagram of the System Fe3O4Mn3O4-MnO-FeO at 1000°C and Thermodynamic Properties of the Coexisting Phases”, Inorg. Mater., 1(3), 382–386 (1965) (Crys. Structure, Thermodyn., Experimental, Phase Diagram, 10) Schenck, H., Steinmetz, E., “Activity, Standard Condition and Coefficient of Activity” (in German), Stahleisen-Sonderberichte, Duesseldorf: Verlag Stahleisen, (7), 1–36 (1966) (Thermodyn., Review, 161) Buzek, Z., Schindlerova, V., Macoszek, M., “The Influence of Cr, Mn, V, Si, Ti, Al, Zr, Ce and Ca on the Activity and Solubility of Oxygen in Liquid Iron” (in Czech), Sb. Ved. Pr. Vys. Sk. Banske Ostrave, Rada Hutn., 13(2–3), 175–193 (1967) (Thermodyn., Review, Phase Relations, 26) Marais, A., Porte, M., “Special Barkhausen Jumps in Ferrites” (in French), Compt. Rend. Acad. Sci. Paris, Ser. A B, 265B(1), 12–15 (1967) (Magn. Prop., Experimental, 7) Schwerdtfeger, K., Muan, A., “Equilibria in the System Fe-Mn-O Involving (Fe, Mn)O and (Fe, Mn)3O4 Solid Solutions”, Trans. Met. Soc. AIME, 239, 1114–1119 (1967) (Phase Diagram, Phase Relations, Thermodyn., Experimental, 44) Schwerdtfeger, K., “Measurement of Oxygen Activity in Iron, Iron-Silicon, Manganese, and Iron-Manganese Melts Using Solid Electrolyte Galvanic Cells”, Trans. Nat. Res. Inst. Met. (Jpn.), 239, 1267–1281 (1967) (Electrochemistry, Experimental, Thermodyn., 30)
DOI: 10.1007/978-3-540-78644-3_18 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–O [1967Sch2]
[1968Oki]
[1969Baf]
[1969Wic]
[1970Aub]
[1970Bul]
[1970Nis]
[1970Roe]
[1971Bra]
[1971Ono]
[1972Bal]
[1972Mat]
[1972Rez]
[1973Buz]
[1973Hol]
[1974Sig] [1974Tan]
Landolt-Börnstein New Series IV/11D4
23
Schnitt, G., Funke, A., “Remarks to the Chemical Analysis of Two Valence Iron and Three Valence Manganese in Manganese Containing Ferrites” (in German), Z. Anal. Chem., 230 (5), 338–344 (1967) (Electronic Structure, Experimental, Phase Relations, 6) Okinaka, H., Kosuge, K., Kachi, S., “Phase Diagram of (Fe, Mn)O at 1260K”, J. Japan. Soc. Powder Met., 8, 295–301 (1968) (Phase Diagram, Thermodyn., Experimental, Phase Relations, 21) Baffier, N., Huber, M., “Determination of Distribution in a Four-cation Mixed Spinel by X-ray and Neutron Diffraction” (in French), Compt. Rend. Acad. Sci. Paris, Ser. C, 268(17), 1521–1524 (1969) (Crys. Structure, Experimental, 6) Wickhalm, D.G., “The Chemical Composition of Spinels in the System Fe3O4-Mn3O4”, J. Inorg. Nucl. Chem., 31, 313–320 (1969) (Crys. Structure, Phase Diagrams, Experimental, 12) Aubry, J., Berthet, A., Duchene, R., Etienne, H., Evrard, O., Jeannot, F., Gleitzer, C., Offroy, C., Perrot, P., “Stabilization of Iron Monoxide by Formation of Solid Solutions” (in French), Ann. Chim., 5, 299–308 (1970) (Experimental, Phase Diagram, Phase Relations, 45) Bulgakova, T.I., Rozanov, A.G., “Phase Equilibria in the Ferrite Region of the Mn-Fe-O System”, Russ. J. Phys. Chem., 44(3), 385–388 (1970), translated from Zh. Fiz. Khim., 44(3), 693–699, (1970) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 15) Nishikawa, K., Kusano, A., Ito, K., Sano, K., “The Effect of Alloying Elements on the Solubility of Oxygen in δ -Iron”, Trans. Iron Steel Inst. Jpn., 10, 83–88 (1970), translated from Tetsu to Hagane, 55, 1193–1198 (1969) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 17) Roethe, A., Roethe, K.-P., Jerschkewitz, H.-G., “Equilibrium Measurements in the Me3O4 Phase in the Fe-Mn-O Ternary System” (in German), Z. Anorg. Chem., 378(1), 1–13 (1970) (Experimental, Phase Relations, Phase Diagrams, Thermodyn., 21) Brabers, V.A.M., “Cation Migration, Cation Valencies and the Cubic-Tetragonal Transition in MnxFe3-xO4”, J. Phys. Chem. Solids, 32, 2181–2191 (1971) (Crys. Structure, Experimental, 40) Ono, K., Ueda, T., Ozaki, T., Ueda, Y., Yamaguchi, A., Moriyama, J., “Thermodynamic Study of the Fe-Mn-O System” (in Japanese), Nippon Kinzoku Gakkai-Si, 35(8), 757–763 (1971) (Experimental, Phase Diagrams, Phase Relations, Thermodyn., 27) Balesdent, D., Evrard, O., Gleitzer, C., “Extension of the Darken’s Method to Heterogeneous Systems for the Calculation of the Activities” (in French), Rev. Chim. Miner., 9(1), 233–244 (1972) (Thermodyn., Calculation, Review, 5) Mathew, P.M., Kapoor, M.L., Frohberg, M.G., “Thermodynamics of the Deoxidation of Fe by Mn” (in German), Arch. Eisenhuettenwes., 43(5), 389–396 (1972) (Calculation, Phase Relations, Thermodyn., 26) Rezlescu, N., Cuciureanu, E., Ioan, C., Luca, E., “Time Variation of the Electrical Conductivity in Spinel Ferrites”, Phys. Status Solidi A, 11(1), 351–359 (1972) (Electr. Prop., Kinetics, Experimental, 9) Buzek, Z., “Effect of Alloying Elements on the Solubility and Activity of Oxygen and Sulphur in Liquid Iron at 1600°C”, Int. Symp. Metall. Chem. - Appl. Ferrous Metall., Sheffield, July 1971, Iron and Steel Inst., London, 173–177 (1973) (Thermodyn., Phase Relations, Review, 8) Holba, P., Khilla, M.A., Krupicka, S., “On the Miscibility Gap of Spinels MnxFe3–xO4+y”, J. Phys. Chem. Solids, 34, 387–395 (1973) (Crys. Structure, Phase Diagram, Experimental, 18) Sigworth, G.K., Elliott, J.F., “The Thermodynamics of Liquid Dilute Iron Alloys”, Met. Sci., 8, 298–310 (1974) (Review, Thermodyn., 249) Tanaka, T., “Lattice Constant and Nonstoichiometric of Mn-Fe Ferrites”, Jpn. J. Appl. Phys., 13(8), 1235–1237 (1974) (Crys. Structure, Experimental, Thermodyn., 17)
MSIT®
DOI: 10.1007/978-3-540-78644-3_18 # Springer 2008
24 [1975Duq]
[1976Jan]
[1977Oet]
[1978Sid]
[1979Kop]
[1979Pel]
[1979Sek]
[1980Bra] [1981Ave]
[1981Mik]
[1981Nai] [1983Dav]
[1983ONe]
[1983Ter]
[1983Tsu]
[1984Gil]
[1984Pun]
Fe–Mn–O Duquesnoy, A., Couzin, J., Gode, P., “Isothermal Representation of Ternary Phase Diagrams A-B-O. Study of the System Mn-Fe-O” (in French), Compt. Rend. Acad. Sci. Paris, 281C, 107–109 (1975) (Experimental, Phase Diagram, Phase Relations, 6) Janke, D., Fischer, W.A., “Equilibrium of Chromium and Manganese with Oxygen in Iron Melts at 1600°C” (in German), Arch. Eisenhuettenwes., 47(3), 147–151 (1976) (Experimental, Thermodyn., Phase Relations, 23) Oeters, F., Koch, K., Scheel, R., Noelle, U., “Investigation of the Equilibria in the System FeMn-O” (in German), Arch. Eisenhuettenwes., 48(9), 475–480 (1977) (Experimental, Phase Diagram, Phase Relations, 11) Sidhu, P.S., Gilkes, R.J., Posner, A.M., “The Synthesis and Some Properties of Co, Ni, Zn, Cu, Mn and Cd Substituted Magnetites”, J. Inorg. Nucl. Chem., 40, 429–435 (1978) (Crys. Structure, Morphology, Experimental, 39) Kopitsa, N.M., Vladimirov, L.P., “Interaction Between Manganese and Oxygen in Iron Melts with Different Carbon Contents”, Russ. J. Phys. Chem., 53(3), 337–339 (1979) translated from Zh. Fiz. Khim., 5(3), 596–599 (1979) (Experimental, Phase Relations, Thermodyn., 11) Pelton, A.D., Schmalzried, H., Sticher, J., “Thermodynamics of Mn3O4-Co3O4, Fe3O4Mn3O4, and Fe3O4-Co3O4 Spinels by Phase Diagram Analysis”, Ber. Bunsen-Ges. Phys. Chem., 83, 241–252 (1979) (Phase Diagram, Calculation, Thermodyn., Phase Relations, 29) Sekine, K., Yamaguchi, T., “Thermodynamics of Ferrite Systems” (in Japanese), Yogyo Kyokai Shi, 87(1009), 443–452 (1979) (Phase Diagram, Phase Relations, Thermodyn., Review, 60) Brandes, E.A., Flint, R.F., “Mn-Fe-O” in “Manganese Phase Diagrams”, Manganese Centre, Paris, France, 115 (1980) (Phase Diagram, Phase Relations, Review, 4) Averbukh, S.M., Smirnov, L.A., “Equilibrium between Manganese and Oxygen in Molten Iron”, Steel USSR, (3), 123–126 (1981), translated from Izv. Vys. Ucheb. Zav., Chern. Metall., (3), 18–23 (1981) (Experimental, Phase Relations, Thermodyn, 15) Mikhailov, G.G., Tankleskaya, N.M., “A Procedure for Calculating the Diagram of Phase Equilibrium in the Fe-Mn-O System during Solidification of Melts”, Steel USSR, (6), 313–315 (1981), translated from Izv. Vys. Ucheb. Zav., Chern. Metall., (6), 5–9 (1981) (Phase Relations, Review, 8) Naito, K., Inaba, H., Yagi, H., “Heat Capacity Measurements of MnxFe3–xO4”, J. Solid State Chem., 36, 28–35 (1981) (Experimental, Magn. Prop., Thermodyn., 45) Davies, P.K., Navrotsky, A., “Quantitive Correlations of Deviations from Ideality in Binary and Pseudobinary Solid Solutions”, J. Solid State Chem., 46, 1–22 (1983) (Review, Theory, 96) O’Neill, H.S.C., Navrotsky, A., “Simple Spinels: Crystallographic Parameters, Cation Radii, Lattice Energies, and Cation Distribution”, Am. Mineral., 68, 181–194 (1983) (Crys. Structure, Review, Theory, Thermodyn., 49) Terayama, K., Ikeda, M., Taniguchi, M., “Phase Equilibria in the Mn-Fe-O System in CO2-H2 Mixtures”, Trans. Jpn. Inst. Met., 24(7), 514–517 (1983) (Experimental, Phase Relations, Thermodyn., 14) Tsuji, T., Asakura, Y., Yamashita, T., Naito, K., “Phase Equilibria of the Mn-Fe-O System (Fe/Mn = 2)”, J. Solid State Chem., 50(3), 273–280 (1983) (Experimental, Phase Relations, Thermodyn., 11) Gillot, B., “Study by Infrared Absorption Spectra of the Factors Influencing the OrderDisorder Transformation in γ Vacancy Ferrites Obtained from the Oxidation of Ferrous Spinels” (in French), Mater. Chem. Phys., 10(4), 375–384 (1984) (Crys. Structure, Optical Prop., Experimental, 17) Punge-Witteler, B., “Phase Equilibria between Spinel and Wustite in the Iron-ManganeseOxygen System at 700°C at 1 bar” (in German), Z. Phys. Chem. (Munich), 142(2), 239–248 (1984) (Experimental, Crys. Structure, Phase Diagrams, Phase Relations, Thermodyn., 22)
DOI: 10.1007/978-3-540-78644-3_18 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–O [1985Tay] [1986Maz]
[1987She]
[1988Lyk]
[1988Yoo]
[1989Rag]
[1990Fra]
[1990Pat]
[1991Ban]
[1991Mik]
[1991Sun] [1992Sub]
[1993Alc]
[1993Sub]
[1993Tur]
[1993Tyu]
[1993Wu]
Landolt-Börnstein New Series IV/11D4
25
Taylor, D., “Thermal Expansion Data. VI. Complex Oxides, AB2O4, the Spinels”, Br. Ceram. Trans. J., 84(4), 121–127 (1985) (Calculation, Crys. Structure, Review, 99) Mazen, S.A., Sabran, B.A., “Thermal Effect on Formation and Conduction Mechanism of MnFe2O4”, Thermochim. Acta, 105, 1–8 (1986) (Crys. Structure, Magn. Prop., Experimental, 13) Shevtsov, V.E., Merker, E.E., Luzgin, V.P., “Thermodynamics of Oxygen Solutions in IronManganese Melts” (in Russian), Izv. Vys. Ucheb. Zav., Chern. Metall., (7), 1–3 (1987) (Experimental, Phase Relations, Thermodyn., 8) Lykasov, A.A., Pavlovaskaya, M.S., “Thermodynamic Properties of Manganowustite”, Russ. J. Phys. Chem., 62(5), 602–605 (1988), translated from Zh. Fiz. Khim., 62(5), 1188–1193 (1988), (Experimental, Theory, Thermodyn., 6) Yoo, H.-I., Tuller, H.L., “In situ Phase Equilibria Determination of a Manganese Ferrite by Electrical Means”, J. Mater, Res., 3(3), 552–556 (1988) (Experimental, Phase Relations, Electr. Prop., Thermodyn., 16) Raghavan, V., “The Fe-Mn-O System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Metals, Calcutta, 5, 181–192 (1989) (Phase Diagram, Phase Relations, Crys. Structure, Review, 34) Franke, P., Dieckmann, R., “Termodynamics of Iron Manganese Mixed Oxides at High Temperatures”, J. Phys. Chem. Solids, 51(1), 49–57 (1990) (Phase Diagrams, Phase Relations, Thermodyn. Experimental, 51) Pattanayak, J., Sitakara Rao, V., Maiti, H.S., “Effect of Iron Concentration on the Thermal behaviour of γ-Mn2O3”, Thermochim. Acta, 160, 233–242 (1990) (Crys. Structure, Experimental, 15) Ban-ya, S., Hino, M., Yuge, N., Kikuchi, I., “Activity Measurement of the Constituents in FeO-MnO Slag Equilibrated with Iron” (in Japanese), Tetsu to Hagane, 77(9), 1419–1425 (1991) (Experimental, Thermodyn., 25) Mikhailov, G.G., Lopatko, N.N., “Phase Equilibria in Iron-Based Melts on Interaction of Oxygen with Manganese, Vanadium, and Chromium”, Russ. Metall., (2), 7–11 (1991), translated from Izv. Akad. Nauk SSSR, Met., (2), 11–15, (1991) (Phase Relations, Thermodyn., Review, 10) Sundman, B., “An Assessment of the Fe-O System”, J. Phase Equilib., 12(2), 127–140 (1991) (Phase Diagram, Thermodyn., Assessment, 53) Subramanian, R., Dieckmann, R., “Limits of the Thermodynamic Stability of Cobalt-IronManganese Mixed Oxides at 1200°C”, J. Am. Ceram. Soc., 75(1), 382–391 (1992) (Phase Relations, Thermodyn., 40) Alcock, C.B., “Thermodynamic and Transport Properties of Electroceramic Oxide Systems”, J. Alloys Compd., 197, 217–227 (1993) (Phys. Prop., Magn. Prop., Transport Phenomena, Review, 28) Subramanian, R., Dieckmann, R., “Nonstoichiometry and Thermodynamics of the Solid Solution (Fe, Mn)1–ΔO at 1200°C”, J. Phys. Chem. Solids, 54(9), 991–1000 (1993) (Experimental, Phase Relations, Thermodyn., 27) Turkdogan, E.T., “Synchronising Free Energy Data on Reactions in Liquid Fe-Mn-O and Solid Mn-S-O”, Ironmaking and Steelmaking, 20(6), 469–475 (1993) (Assessment, Thermodyn., 35) Tyurin, A.G., “Thermodynamics of Molecular and Ionic Solutions”, Russ. Metall., (2), 39–47 (1993), translated from Izv. Ross. Akad. Nauk, Met., (2), 48–56 (1993) (Theory, Thermodyn., 35) Wu, P., Eriksson, G., Pelton, A.D., “Critical Evaluation and Optimization of the Thermodynamic Properties and Phase Diagrams of the CaO-FeO, CaO-MgO, CaO-MnO, FeO-MgO, FeO-MnO, and MgO-MnO Systems”, J. Am. Ceram. Soc., 76(8), 2065–2075 (1993) (Phase Diagram, Phase Relations, Thermodyn., Assessment, 73)
MSIT®
DOI: 10.1007/978-3-540-78644-3_18 # Springer 2008
26 [1994Sub]
[1995All]
[1995Bat]
[1995Kow]
[1995Sub]
[1996Gol]
[1997Pet]
[1998Bon]
[1998Lee]
[2000Pol] [2000Sei]
[2000Tak]
[2001Der]
[2001Hu]
[2001Mur]
[2002Hin]
Fe–Mn–O Subramanian, R., Dieckmann, R., Eriksson, G., Pelton, A., “Model Calculations of Phase Stabilities of Oxide Solid Solutions in the Co-Fe-Mn-O System at 1200°C”, J. Phys. Chem. Solids, 55(5), 391–404 (1994) (Calculation, Thermodyn., Phase Diagram, Phase Relations, 56) Allen, G.C., Hallam, K.R., Jutson, J.A., “X-ray Diffraction Studies of Solid Solutions of Cr-, Mn-, Fe-, Co-, and Ni-containing Transition Metal Spinel Oxides”, Powder Diffr., 10(3), 214–220 (1995) (Crys. Structure, Experimental, 8) Battault, T., Legros, R., Rousset, A., “Structural and Electrical Properties of Iron Manganite Spinels in Relation with Cationic Distribution”, J. Eur. Ceram. Soc., 15, 1141–1147 (1995) (Crys. Structure, Electrical Prop., Electronic Structure, Experimental, 25) Kowalski, M., Spencer, P.J., “Thermodynamic Revaluation of the Cr-O, Fe-O and Ni-O Systems: Remodelling the Liquid, BCC and FCC Phases”, Calphad, 19(3), 229–243 (1995) (Assessment, Phase Diagram, Phase Relations, Thermodyn., Review, 47) Subramanian, R., Üstündag, E., Sass, S.L., Dieckmann, R., “Metal-Ceramic Microstructure Control in Partial Reduction Reactions in the Model System Fe-Mn-O by Doping”, Mater. Sci. Eng. A, A195, 51–63 (1995) (Experimental, Morphology, Phase Relations, Phase Diagram, Thermodyn., 33) Golikov, Yu.V., Petrova, S.A., Zakharov, R.G., Antonov, A.V., Balakirev, V.F., “Equilibrium and Unstable States of the Fe-Mn-O System in Air”, Russ. J. Inorg. Chem., 41(9), 1510–1513 (1996), translated from, Dokl. Ross. Akad. Nauk, 344(6), 770–772 (1995) (Experimental, Phase Relations, 19) Petrova, S.A., Golikov, Yu.V., Zakharov, R.G., Balakirev, V.F., “Heterogeneous Equilibria in the Fe-Mn-O System in Air”, Russ. J. Phys. Chem., 71(2), 184–186 (1997), translated from Zh. Fiz. Khim., 71(2), 240–242 (1997) (Experimental, Phase Relations, Thermodyn., 10) Bonsdorf, G., Schafer, K., Teske, K., Langbein, H., Ullmann, H., “Stability Region and Oxygen Stoichiometry of Manganese Ferrite”, Solid State Ionics, 110, 73–82 (1998) (Electronic Structure, Experimental, Phase Diagram, Phase Relations, 19) Lee, B.-J., “Thermodynamic Calculation for Stability of Oxides in Steel Systems” (in Korean), J. Korean Inst. Met., 36(2), 217–224 (1998) (Calculation, Thermodyn., Phase Relations, 57) Pollert, E., “Influence of Mn3+ Ions on Ordering in Magnetic Oxides”, Int. J. Inorg. Mater., 2, 661–670 (2000) (Crys. Structure, Theory, Review, 18) Seifu, D., Kebede, A., Oliver, F.W., Hoffman, E., Hammond, E., Wynter, C., Aning, A., Takacs, L., Siu, I.-L., Walker, J.C., Tessema, G., Seehra, M.S., “Evidence of Ferrimagnetic Ordering in FeMnO3 Produced by Mechanical Alloying”, J. Magn. Magn. Mater., 212, 178–182 (2000) (Crys. Structure, Experimental, Magn. Prop., Morphology, Nano, 13) Takahashi, K., Hino, M., “Equilibrium between Dissolved Mn and O in Molten High-Manganese Steel”, High Temp. Mater. Proc., 19(1), 1–10 (2000) (Experimental, Phase Relations, Thermodyn., 28) Deraz, N.A.M., El-Shobaky, G.A., “Solid-Solid Interaction Between Ferric Oxide and Manganese Carbonate as Influenced by Lithium Oxide Doping”, Thermochim. Acta, 375, 137–145 (2001) (Crys. Structure, Experimental, Transport Phenomena, 37) Hu, J., Qin, H., Wang, Y., Zhang, S., Wang, Z., “Magnetic Properties and Magneto-Transport in MnFe2O4–δ with Tetragonal Structure”, J. Mater. Sci. Lett., 20, 1531–1532 (2001) (Crys. Structure, Experimental, Magn. Prop., 5) Muroi, M., Street, R., McCormick, P.G., Amighian, J., “Magnetic Properties of Ultrafine MnFe2O4 Powders Prepared by Mechanochemical Processing”, Phys. Rev. B, 63(18), 184414_1–7 (2001) (Crys. Structure, Experimental, Magn. Prop., 28) Hino, M., Takahashi, K., Itoh, T., Kikuchi, I., Nagasaka, T., “Equilibrium Between Liquid Iron and FeO-containing Solid Solution”, Scand. J. Metall., 31(3), 169–177 (2002) (Experimental, Phase Diagram, Phase Relations, Thermodyn., 64)
DOI: 10.1007/978-3-540-78644-3_18 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–O [2002Sar]
[2003Mue] [2003Per] [2004Jun]
[2004Med]
[2004Umi]
[2004Wit]
[2005Pad]
[2005Zin]
[2006Swa]
[2007Fen]
[2007Wit] [Mas2] [V-C2]
Landolt-Börnstein New Series IV/11D4
27
Saravanan, P., Alam, S., Kandpal, L.D., Mathur, G.N., “Effect of Substitution of Mn Ion on Magnetic Properties of Fe3O4 Nanocrystallites”, J. Mater. Sci. Lett., 21(14), 1135–1137 (2002) (Crys. Structure, Experimental, Morphology, Magn. Prop., 15) Mueller-Buschbaum, H., “The Crystal Chemistry of AM2O4 Oxometallates”, J. Alloys Compd., 349(1–2), 49–104 (2003) (Crys. Structure, Electronic Structure, Review, 476) Perrot, P., Foct, J., “Gases other than Hydrogen in Iron and Steels” (in French), Techniques de l’Ingenieur, M4275, 1–23 (2003) (Thermodyn., Kinetics, Phase Relations, Review, 127) Jung, I.-H., Decterov, S.A., Pelton, A.D., “A Thermodynamic Model for Deoxidation Equilibria in Steel”, Metall. Mater. Trans. B, 35B(3), 493–507 (2004) (Theory, Thermodyn., Review, 100) de Medeiros, S.N., Luciano, A., Cotica, L.F., Santos, I.A., Paesano, A.Jr., da Cunha, J.B.M., “Structural and Magnetic Characterization of the Ball-Milled αFe2O3-Mn2O3 and αFeMn2O3 Systems”, J. Magn. Magn. Mater., 281, 227–233 (2004) (Crys. Structure, Electronic Structure, Magn. Prop., Experimental, 15) Umino, M., Sera, T., Ikenaga, Y., Okada, Y., Murakami, R., Tubakino, S., “Role of Alloying Elements on Machinability of Plastic-molding Steels”, Z. Metallkd., 95(2), 109–114 (2004) (Mechan. Prop., Morphology, Experimental, 13) Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diff., 25(4), 346–354 (2004) (Experimental, Phase Diagram, Calculation, Thermodyn., #, 34) Padella, F., Alvani, C., La Barbera, A., Ennas, G., Liberatore, R., Varsano, F., “Mechanosynthesis and Process Characterization of Nanostructured Manganese Ferrite”, Mater. Chem. Phys., 90(1), 172–177 (2005) (Morphology, Nano, Experimental, Phase Relations, 27) Zinovik, E.V., Zinovik, M.A., “Hydrogen Reduction of Cu-Mn-Fe-O Spinel Solid Solutions”, Inorg. Mater., 41(3), 272–278 (2005), translated from Neorg. Mater., 41(3), 332–338, (2005) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Thermodyn., 14) Swaminathan, R., Willard, M.A., McHenry, M.E., “Experimental Observations and Nucleation and Growth Theory of Polyhedral Magnetic Ferrite Nanoparticles Synthesized using an RF Plasma Torch”, Acta Mater., 54(3), 807–816 (2006) (Experimental, Morphology, Nano, Phase Relations, Theory, 37) Fenstad, J., Perrot, P., “Mn-O (Manganese-Oxygen)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services, GmbH, Stuttgart; to be published, (2007) (Crys. Structure, Thermodyn., Phase Diagram, Phase Relations, Review, #, 87) Witusiewicz, V.T., private communication to MSI, (2007) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Iron – Manganese – Phosphorus Kostyantyn Korniyenko
Introduction Phase relationships in the Fe-Mn-P system are of great interest, in the first instance because of the importance of manganese as an alloying element in steels. Another important aspect is related to the use of phosphorus as a glass-formation promoting element in the production of iron-based amorphous alloys. Recently, considerable research effort has been focused on nanocrystalline materials produced by controlled crystallization of amorphous alloys, and Fe-Mn-P amorphous alloys are candidate materials for this application. However, current knowledge of the constitution of the Fe-Mn-P system is lacking. Experimental studies of phase equilibria were undertaken long ago [1952Vog]. The reaction scheme, liquidus surface projection and isothermal section at 20°C of the partial Mn-Mn2P-Fe2P-Fe system as well as a series of vertical sections were proposed. More recently, information on the solubility of manganese in iron at different temperatures was presented in [1965Kan2], and phase contents of as-cast alloys in [1993Vav, 1997Vav, 1999Vav]. In parallel with the studies of the ternary system, the constitution of the Fe-Mn binary system was revised and new crystal structures for the ternary solid solutions have been determined. Therefore, a reinvestigation of the Fe-Mn-P system using modern physico-chemical analysis techniques is necessary. Publications devoted to experimental study of the phase relations, crystal structures and thermodynamics are listed in Table 1 together with the techniques used. The first studies of crystal structure of ternary alloys were focused on the phases present in alloys of compositions along the Mn3P-Fe3P and Mn2P-Fe2P sections, and were carried out by [1948Now]. The crystal structures of the unary, binary and ternary phases of the Fe-Mn-P system are given also in [1952Vog, 1969Fru, 1969Rog, 1973Mae, 1973Nag, 1973Suz, 1977Got, 1981Maa, 1986Fje, 2000Bro, 2003Bal, 2003Liu]. Data on thermodynamic properties, mainly on the activity of phosphorus in liquid iron with manganese additions, were obtained experimentally by [1969Sch, 1972Sha, 1980Kun, 1983Yam, 1984Ban, 1995Zai, 1996Zai, 1998Zai]. Reviews of the literature deal with information on phase equilibria in the Fe-Mn-P system [1965Kan1, 1965Kan2, 1980Bra, 1988Rag, 1999Vav], crystal structures [1948Now, 1988Rag, 1999Vav] and thermodynamics [1974Sig, 1984Ban]. A database of thermodynamic model parameters for liquid Fe-based alloys containing Mn and P was proposed by [1995Bou] and a model of mass action concentrations for the Fe-Mn-P melt was presented by [2000Zha]. Binary Systems The Fe-P boundary binary system is accepted from [2002Per] and the Fe-Mn system is accepted from the recent assessment by [2004Wit]. There is however, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have been written as L ⇌ (δMn) + (γMn,γFe). The Mn-P boundary system is fairly well determined and is accepted from [Mas2]. However, there is some uncertainty relating to the reactions involving (δMn), (γMn) and (βMn) with liquid at virtually pure Mn. According to [Mas2], there are two invariant reactions that seem to take place at the transformation temperatures between pure (δMn) and (γMn), and between pure (γMn) and (βMn), 1138 and 1100°C, respectively. According to the article from where these data are taken [1950Ber], these temperatures are 1143 and 1072°C. From the data points given on the phase diagram presented by [1950Ber], it would seem that these relate again to the transformation temperatures of pure Mn, and the differences between these values and the accepted values being to due a reassessment of pure Mn. Therefore, it is not possible to be able to determine the nature of these binary invariants from these data, and they are thus classed as degenerate. Solid Phases Crystallographic data for the known unary, binary and ternary phases of the Fe-Mn-P system are listed in Table 2. Information concerning the phase relationships is incomplete; only one report of the joint solubility Landolt-Börnstein New Series IV/11D4
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of two of the components in the third has been presented (related to the (αFe)-based solid solution) [1965Kan2]. However, the isomorphous Mn2P- and Fe2P-based phases (Fe2P-type structure) have been found to form a continuous series of solid solutions (M2P) at high temperatures [1948Now, 1952Vog]. At lower temperatures in the (Mn1–yFey)2P composition range, a ternary phase (labeled as τ1) with an orthorhombic structure appears at compositions of 0.38 ≤ x ≤ 0.69 in alloys annealed at 850–900°C [1969Fru, 1969Rog]. The lattice parameters presented for this phase in Table 2 have been estimated from the figures given in the original publications. In the same work, it was also reported that the orthorhombic τ1 and hexagonal M2P phases are separated by very narrow two-phase regions. The stoichiometry of the τ1 phase centers around MnFeP (x = 0.5) and has the Co2Si structure prototype. The τ1 phase was also obtained by [1973Nag] but in somewhat different temperature and composition ranges (Table 2). Later, the crystal structures of the M2P and the τ1 phases were investigated in more detail by [1973Mae, 1973Suz, 1981Maa, 1991Sjo]. In particular, the first band structure calculations of the MnFeP compound indicated that the crystallographic position, rather than the type of metallic atom that occupies each position, determines the electronic and magnetic properties. The Mn3P and Fe3P phases are isomorphous (Ni3P structure type) but they have not been reported as forming a continuous series of solid solutions. At the same time, the solubility of the third component in both phases is quite large. A number of diverse solubility values have been presented [1977Got, 2000Bro, 2003Liu] but they probably correspond to different temperatures. Lattice parameters of the Mn3P- and Fe3P-based solid solutions are presented in Table 2 and have been estimated from figures given in [1977Got]. Included in the same work is evidence of another ternary phase that occurs at a composition lying between the Mn3P- and Fe3P-based solid solutions. However, according to [1988Rag], an ordered orthorhombic phase lying in this section cannot be considered as a true ternary phase. But in the present evaluation this phase is accepted (denoted as τ2, (Mn1–yFey)3P, (Table 2)) because its structure differs from that of other unary and binary phases participating in the phase equilibria in the Fe-Mn-P system. More work on the temperature and composition ranges and crystal structures are needed for confirmation. The crystal structures of the FeP and MnP phases belong to the same MnP type structure type. The crystal structure of the FeP phase was studied by [1986Fje] with additions of manganese of only up to 10 at.%, but the existence of a continuous series of solid solutions between these two phases is possible. Further study of the complete MnP-FeP section is necessary. [1972Sel] determined that in pure FeP, the z-value of the position of the P atoms is 0.296, being considerably shifted from the value of 0.25 demanded by the mirror plane in the Pnma space group. Thus the mirror plane disappears and the space group transforms to the subgroup Pna21 with the same symmetry elements, except for this mirror plane. Mössbauer effect measurements and physico-chemical analysis revealed that annealing amorphous Fe-Mn-P alloys leads to the formation of a nanocrystalline structure [2003Bal]. However, information about phase equilibria in the Fe-Mn-P ternary system involving the stable binary Mn3P2 phase is absent. According to the data of [1965Kan2], the solubility of phosphorus in (αFe) decreases from 4.4 at.% (2.5 mass%) in the absence of manganese, to 3.0 at.% (1.6 mass%) at 4.0 at.% (3.8 mass%) Mn. With temperature decreasing from 1000 to 800°C, the solubility of P in (αFe) containing 1 at.% Mn decreases from 4.0 to 2.0 at.%. Quasibinary Systems [1948Now] reported the quasibinary character of the Mn2P-Fe2P section from their X-ray diffraction studies. The phase diagram of the Mn2P-Fe2P system was constructed by [1952Vog] on the basis of thermal and chemical analyses coupled with metallographic examination. The section is shown in Fig. 1 with amendment of the melting temperature of the Fe2P (1350°C in [1952Vog]) by 1370°C [Mas2]. Corresponding corrections to the liquidus curves were also carried out. Approximate phase boundaries of the τ1 phase were plotted using the data of [1969Fru, 1969Rog], which had been obtained from the study of alloys annealed at 850–900°C. The quasibinary character of the Mn2P-Fe2P section enables triangulation of the ternary Fe-Mn-P system along this line and to consider the phase equilibria in the Mn-Mn2P-Fe2P-Fe partial system independently of the more P rich part of the ternary system.
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Invariant Equilibria Table 3 presents the temperatures, types of reactions and the available compositions of the phases taking part in the invariant equilibria in the partial Mn-Mn2P-Fe2P-Fe system. The information, together with the partial reaction scheme shown in Fig. 2, are based mainly on the data of [1952Vog] who presented information on the invariant reactions and liquidus surface. Corrections have been made in accordance with the accepted boundary binary systems. As a consequence, because the accepted Fe-Mn system [Mas2] shows (γMn) and (γFe) forming a continuous series of solid solutions in contrast to the earlier version of the phase diagram accepted by [1952Vog], the number of four-phase invariant equilibria involving the liquid phase has decreased by one. The invariant reaction involving the liquid, γ, M2P and Fe3P phases given by [1952Vog] as of transition U type is in error. Reanalysis of their data indicates that this reaction should really be of the decomposition E type. In the review of [1988Rag], two additional solid state transition reactions were proposed to account for the ordering reaction in the Mn2P-Fe2P section, but this has not been confirmed experimentally. Liquidus, Solidus and Solvus Surfaces Figure 3 shows the liquidus surface projection of the partial Mn-Mn2P-Fe2P-Fe system. It is based principally on the data of [1952Vog] (reproduced similarly in the review of [1980Bra]) with some modifications similar to those were carried out by [1988Rag]. Therefore, in comparison with the original source material, the peritectic reaction L + (γFe) ⇌ (γMn) does not take place as (γFe) and (γMn) form a continuous series of solid solutions, and therefore a common field of primary crystallization exists on the liquidus surface. Other changes were made in accordance with the boundary binary systems and invariant equilibria (see “Binary Systems” and “Invariant Equilibria” sections). The crystallization path of rapidly quenched alloys of compositions Mn5.04Fe86.5P8.46, Mn4.80Fe83.2P12.0 and Mn9.8Fe81.6P8.6 (in at.%) was studied by [1993Vav], and the obtained results are in a good agreement with the accepted liquidus surface. Information on the liquid-solid equilibria in the Fe-Mn-P system is incomplete, especially with regard to the solidus and solvus surfaces and the liquidus surface in the P rich range of the diagram. Further experimental investigations of these aspects are necessary. Isothermal Sections An isothermal section of the Mn-Mn2P-Fe2P-Fe partial system at a temperature of 20°C was presented by [1952Vog], based on the results of their own experimental investigations. [1988Rag] attempted to revise this section for 25°C using later data regarding phase relations at room temperature of the boundary binary and ternary systems. A nominal width of about 1 at.% was assumed by [1988Rag] for the two-phase regions between the hexagonal (Mn2P and Fe2P) and the orthorhombic (τ1) phases on the Fe2P-Mn2P section. According to the variant of the Fe-Mn system accepted by [1952Vog], at 20°C the (γFe) solid solution region extends from about 29 to 61 at.% Mn, but according to the Fe-Mn phase diagram accepted by [1988Rag], this region was replaced by (γFe,γMn). The compilation of the data of [1952Vog] and [1988Rag] with corrections according to the accepted boundary binary systems is presented in Fig. 4. From the data of [2004Wit] concerning the Fe-Mn system it follows that at 25°C a wide two-phase region α + (αMn) stretches from the iron apex up to about 80 at.% Mn, at the same time the γ phase does not exist at this temperature in the binary system. Because in the ternary Fe-Mn-P system three-phase fields with the participation of the γ phase exist [1952Vog], we suppose that the γ phase is stabilized in the ternary system by phosphorus. Therefore the γ single-phase field and the attendant three-phase field α + γ + (αMn) are plotted in Fig. 4. [1965Kan1] studied phosphides isolated electrolytically from ternary Fe-P-M alloys containing 2.5 mass% P in order to clarify the phase relationships between the phosphide phases and the iron phase in these systems and also to determine the phosphide-forming tendency of alloying elements in steel. It was observed that manganese has a weak tendency to combine with phosphorus in steel, because no phosphide phases, except the Fe3P phase were found in the ternary Fe-Mn-P alloys. It was reported that at 800°C, on increasing the manganese concentration in steel from 1.80 to 4.87 mass%, its concentration in the Fe3P phase correspondingly increases from 5.7 to 13.3 mass%. Landolt-Börnstein New Series IV/11D4
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Phase contents of a number of different Fe-Mn-P alloys annealed at different temperatures were presented in [1997Vav, 1998Vav, 1999Vav]. In particular, in rapidly quenched and annealed (at 500°C for 10 minutes) alloys of the compositions (in at.%) Mn4.7Fe77.1P18.2 and Mn9.3Fe72.5P18.2, the presence of the α and (MnxFe1–x)3P phases, and the α, (MnxFe1–x)3P and M2P phases, respectively, was established. In [1998Vav], the same information was reproduced but a different value of annealing temperature (770 K or 497°C) was reported. Rapid quenching and annealing (400°C for 10 min) of these two alloys resulted in the presence of the α and (MnxFe1–x)3P phases as well as the paramagnetic γ phase [1997Vav]. Temperature – Composition Sections A series of temperature-composition sections has been constructed by [1952Vog] based on the results of their experimental investigations. The temperature-composition section at 95 mass% Fe and phosphorus contents up to 4 mass% (Mn5.08Fe94.92-Mn0.98Fe92.03P6.99 in at.%) is presented in Fig. 5, with corrections in accordance with the accepted Fe-Mn phase diagram (existence of the α + γ two-phase region at the room temperature). The temperature-composition section at 90 mass% Fe proposed by [1952Vog] is incorrect because the phosphorus content is in error (up to 21.5 mass%). Moreover, at the Fe-Mn system side it contradicts the binary system accepted by [1952Vog]. Also the temperature-composition sections at 50 and 10 mass% Fe and at 1 mass% P do not comply with the accepted Fe-Mn binary system. The temperature-composition sections at 12 and 6 mass% P (Mn80.52P19.48-Fe80.27P19.73 and Mn89.83P10.17-Fe89.68P10.32 in at.%, respectively) do not agree with the accepted invariant reaction U2. Thermodynamics The effect of manganese on the activity coefficient of phosphorus in liquid iron alloys at different temperatures was investigated by [1969Sch, 1983Yam, 1984Ban] and also by [1972Sha, 1980Kun] (cited in [1995Zai]). It was concluded by [1969Sch] that the addition of Mn (up to 16 mol%) to liquid Fe-P alloys at 1550°C does not influence on the activity of phosphorus and therefore the interaction parameter εPMn = 0. Knudsen cell - mass spectrometry measurements carried out at 1600°C by [1983Yam] demonstrated that the vapor pressure of manganese in liquid iron was high. However, during the experiment, the inside wall of the ion source became severely contaminated by Mn vapor making accurate determinations impossible. Later, the vapor pressure of phosphorus in liquid Fe-Mn-P alloys with manganese contents of up to 19.3 mass% was measured by [1984Ban] by a transportation method at the lower temperature of 1400°C. The results were treated using the interstitial solution model of J. Chipman. The effect of manganese on the activity coefficient of phosphorus in liquid iron was determined by assuming the manganese dissolves substitutionally. A value of εPMn = – 7.17 ± 1.16 at 1400°C was obtained for the interaction parameter. The set of experimental data obtained by [1984Ban] was optimized in [1995Bou], and satisfactory agreement between the calculated and experimental data was found. [1993Din] calculated the activity interaction parameter between Mn and P in liquid Fe at 1600°C. The value derived of εPMn = – 5.03 differs from the experimental value cited in [1988The] (– 7.2). [1972Sha] measured the activity of phosphorus in Fe-Mn-P melts containing 0.8 to 30.7 at.% P and 1.7 to 2.0 at.% Mn using the emf technique. The electrolyte used was a synthetic slag containing 41.18 at.% CaO, 39.87 at.% SiO2, 9.20 at.% Al2O3, 2.64 at.% B2O3 and 1.27 at.% P2O5. A melt containing 30.7 at.% P and 1.8 at.% Mn was used as the reference electrode. It was assumed that, in view of the low concentration of iron and manganese ions in the slag, the potential was controlled by phosphorus transfer between the alloy and the electrolyte. Based on consideration of Henry’s law, [1972Sha] concluded that in calculations for low-phosphorus alloys (0.8 at.% of P) the mole fraction of phosphorus can be used instead of its activity. In the determination of the activity of Mn at 1500°C, [1980Kun] used the same technique and an electrolyte with a similar composition: 40.85 at.% CaO, 43.26 at.% SiO2, 8.34 at.% Al2O3, 2.31 at.% B2O3 and 1.46 at.% MnO. Knudsen-cell mass-spectrometry was used by [1995Zai, 1996Zai, 1998Zai] to investigate the thermodynamic properties of liquid alloys in the temperature range from 1029 to 1428°C (1302 to 1701 K). The concentrations of iron, manganese and phosphorus of the alloys studied were in the ranges from 20.5 to 73.8, from 22 to 69.5 and from 15.2 to 36.5 at.%, respectively. The thermodynamic properties of the melts DOI: 10.1007/978-3-540-78644-3_19 # Springer 2008
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were approximated by application of the associated-solution model. For adequate representation of the experimental data it was necessary to take into account binary associates formed by iron and manganese atoms with phosphorus, the ternary associative complex FeMnP and excess terms describing the interactions between the monomer species. The thermodynamic characteristics of the binary associates were found to agree with the values established earlier from the experimental data for binary systems within the limits of experimental error. The obtained molar Gibbs free energies and enthalpies of formation from the liquid components at 1377°C are presented in Figs. 6 and 7, respectively, in the form of isovalue lines. These lines are asymmetric, and the extrema are displaced from the phosphorus corner towards the binary Fe-Mn side and from the Fe-P side towards the Mn-P side. The entropy is negative for most of the concentration triangle. It becomes positive only in the regions close to the corners of the triangle. [1974Sig] collected and assessed literature data on the thermodynamic behavior of dilute solute elements (including manganese and phosphorus) in liquid iron. A non-equilibrium thermodynamic model that describes the effect of solute grain boundary segregation on grain boundary cohesion was proposed by [1984Lee] for iron based ternary systems, in particular, for the Fe-Mn-P system. This model directly and simply predicts the effect of alloying elements on impurity-induced grain boundary embrittlement. According to the model, Mn enhances impurity-induced grain boundary embrittlement in a Fe ternary system. The equilibrium segregation of phosphorus at austenite grain boundaries in a Mn10Fe89P1 (mass%, in at.% – Mn10.07Fe88.15P1.78) alloy was modeled by [1989Paj]. The segregation was described using the Langmuir-McLean equation, and a value for the free energy of phosphorus segregation ΔGoP was estimated as 47 kJ·mol–1. The activation energy of crystallization of Mn4.7Fe77.1P18.2 and Mn9.8Fe72.5P18.2 (in at.%) amorphous alloys was estimated by [1999Vav, 2000Vav] using the Kissinger method, in terms of the variation in crystallization temperature (Tc) as a function of heating rate in the range 3.75 to 40°C·min–1. For these two alloys, the values of the activation energy obtained were 230 ± 6 and 137 ± 6.5 kJ·mol–1, respectively. Based on the phase diagrams and the coexistence theory of the structure of metallic melts involving compound formation, models for mass action concentrations for the Fe-Mn-P melt have been formulated in [2000Zha]. The calculated mass action concentrations agree well with the corresponding measured activities, and this in turn shows that the deduced model can reflect the structural characteristics of the melt concerned and that there is no phosphorus saturation of the melt. Notes on Materials Properties and Applications Fe-Mn-P and related alloys (complicated manganese-containing steels and ferromanganese alloys) find wide practical applications in various fields of modern technology. Another important aspect of their use is related to the application of phosphorus as a glass forming element for the production of iron based amorphous alloys. However, the preparation of Fe-P master alloys involves a number of practical difficulties, and the addition of some transition metals, in particular manganese, can help in the preparation of alloys suitable for amorphization [1997Vav]. Material property investigations are listed in Table 4. On the whole, magnetic and mechanical properties have been studied in some detail. It has been observed that very small substitutions of manganese in Fe2P induce metamagnetism [1969Fru, 1969Rog]. The τ1 phase was reported to demonstrate antiferromagnetic properties with a Néel temperature of 320 K [1973Nag] or 340 K [1973Suz]. It was shown by [1993Vav, 1997Vav, 1998Vav, 1999Vav] that the embrittlement of iron based Fe-Mn-P alloys is governed by the amount of α phase precipitating during crystallization. The first band structure calculations of the MnFeP compound were carried out by [1991Sjo]. It was found that the crystallographic position, rather than the type of metallic atoms that occupy these positions, determines the electronic and magnetic properties. Miscellaneous The lattice and grain boundary tracer diffusion coefficients D*PL and P*P of phosphorus in the Mn0.38Fe99.45P0.17 and Mn2.11Fe97.78P0.11 (at.%) alloys were measured by [1983Mat]. There is a break in each of the D*PL – 1/T and P*P – 1/T lines at temperatures near and below the Curie temperature, respectively. It was considered that these anomalous breaks were caused by magnetic transformations taking place in the bulk and at grain boundaries. It was shown that alloying with manganese increased D*PL and P*P. Landolt-Börnstein New Series IV/11D4
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[1984Lee] studied the effect of manganese on phosphorus grain boundary segregation. Their experimental results showed that after aging for 300 hours, maximum P grain boundary segregation occurs at the aging temperature of 550°C. It was concluded that manganese possesses a much lower grain boundary segregation enrichment ratio than P. Equilibrium segregation of phosphorus at the grain boundaries of austenite was studied by [1989Paj] in a Mn10Fe89P1 (mass%, in at.% - Mn10.07Fe88.15P1.78) alloy equilibrated at temperatures between 750 and 1200°C before rapid quenching. It was shown that manganese increased the level of phosphorus segregation. It was suggested that the level of phosphorus segregation is enhanced by manganese because manganese intensifies the additional segregation of phosphorus during quenching. It was established by [1993Vav, 1997Vav, 1998Vav, 1999Vav] that alloys of the Fe-Mn-P system have a tendency to amorphization at concentrations between 15.1 and 21.2 at.% P and between 4.7 and 9.3 at.% Mn at cooling rates of 105 to 106 K/sec. The viscosity of Mn4.7Fe77.1P18.2 and Mn9.3Fe72.5P18.2 (at.%) alloys after melt homogenization for 1 h at 1150–1200°C and for an additional 0.5 h at the set temperature was measured by [2000Vav].
Table 1.
Investigations of the Fe-Mn-P Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1948Now]
X-ray diffraction
Mn3P-Fe3P and Mn2P-Fe2P sections
[1952Vog]
Melting in Tamman furnace (sintered corundum crucibles), thermal analysis, optical microscopy, chemical analysis
≤ 1500°C, Mn-Mn2P-Fe2P-Fe partial system
[1965Kan1]
Electrolytical isolation, X-ray diffraction, chemical analysis
1000°C, 800°C, 2.5 mass% P in the alloys
[1965Kan2]
X-ray diffraction, chemical analysis
1100–700°C, ≤ 4 at.% P
[1969Fru]
X-ray diffraction (Seeman-Bohlin camera), Mössbauer spectroscopy
Mn2P-Fe2P section
[1969Rog]
X-ray diffraction (Seeman-Bohlin camera), Mössbauer spectroscopy
Mn2P-Fe2P section
[1969Sch]
Equilibrium P (vapor) – molten Fe with Mn additions
1550°C
[1972Sha]
Emf measurements in a liquid electrolyte
1550°C, 1.7 to 2.0 at.% Mn, 0.8 to 30.7 at.% P
[1973Mae]
X-ray powder diffraction, Mössbauer spectroscopy
900°C, MnFeP
[1973Nag]
X-ray powder diffraction (Debye-Scherrer technique)
Mn2P-Fe2P section
[1973Suz]
X-ray powder diffraction, neutron powder diffraction, Mössbauer spectroscopy
MnFeP
[1977Got]
X-ray powder diffraction
Mn3P-Fe3P section
[1980Kun]
Emf measurements in a liquid electrolyte
1500°C, 0.7 to 10.8 at.% Mn, 1.1 to 32.5 at.% P
[1981Maa]
X-ray powder diffraction, Mössbauer spectroscopy
Mn0.7Fe1.3P (continued)
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Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1983Yam]
Melting, Knudsen cell-mass spectrometry
1600°C
[1984Ban]
Transportation method
εPMn in liquid phase, 1400°C, ≤ 19.3 mass% Mn
[1986Fje]
X-ray powder diffraction, neutron diffraction
50 at.% P, 40 to 50 at.% Fe
[1993Vav]
Arc melting, quenching, X-ray diffraction, differential thermal analysis
8.46 to 21.2 at.% P, 4.60 to 9.80 at.% Mn
[1995Zai]
Knudsen effusion and cell-mass spectrometry, X-ray diffraction, chemical analysis
1428–1029°C, 20.5 to 73.8 at.% Fe, 2.2 to 69.5 at.% Mn, 15.2 to 36.5 at.% P
[1996Zai]
Arc melting, Knudsen effusion and cell-mass spectrometry
1428–1029°C, 20.5 to 73.8 at.% Fe, 2.2 to 69.5 at.% Mn, 15.2 to 36.5 at.% P
[1997Vav]
Arc melting, X-ray diffraction, thermal analysis, TEM, electron diffraction, Mössbauer spectroscopy
≤ 10 at.% Mn
[1998Vav]
Mössbauer spectroscopy, differential thermal analysis, X-ray diffraction
4.7 and 9.3 at.% Mn, 18.2 at.% P
[1998Zai]
Knudsen effusion and cell-mass spectrometry, X-ray diffraction, chemical analysis
1428–1029°C, 20.5 to 73.8 at.% Fe, 2.2 to 69.5 at.% Mn, 15.2 to 36.5 at.% P
[1999Vav]
Mössbauer spectroscopy, differential thermal analysis
4.7 and 9.3 at.% Mn, 18.2 at.% P
[2000Bro]
X-ray diffraction, neutron diffraction, Mössbauer spectroscopy
25 at.% P, 0 to 18.75 at.% Mn
[2003Bal]
Mössbauer spectroscopy, physico- chemical analysis techniques
Amorphous Fe-Mn-P alloys
[2003Liu]
X-ray powder diffraction, neutron powder diffraction, Mössbauer spectroscopy, chemical analysis
880°C, 25 at.% P
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
α, (αδFe)
cI2 Im 3m W
Lattice Parameters [pm]
Comments/References
dissolves 10 at.% Mn at 1473°C, 3 at.% Mn at ∼400°C [Mas2], dissolves 4.55 at.% P at 1048°C, 4 at.% P at 1000°C [2002Per] dissolves 3 at.% P at 4 at.% Mn, 1000°C [1965Kan2] (continued)
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Fe–Mn–P
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(δFe) (h2) 1538–1394 (αFe) (r) (ferrite) < 912 γ, (Mn1–xFex)
Lattice Parameters [pm]
Comments/References
a = 293.15
pure Fe at 1390°C [V-C2, Mas2]
a = 286.65
pure Fe at 25°C [Mas2]
a = 386
0 ≤ x ≤ 1, dissolves 0.56 at.% P at x = 1 and 1150°C [Mas2] pure Mn, at > 1100°C [Mas2]
a = 364.67
pure Fe at 915°C
cF4 Fm 3m Cu
(γMn) 1138–1100 (γFe) (h1) (austenite) 1394–912
[Mas2]
(δMn) 1246–1138
cI2 Im 3m W
a = 308
pure Mn, at > 1138°C [Mas2] dissolves 9 at.% Fe at 1235°C [Mas2]
(βMn) 1100–727
cP20 P4132 βMn
a = 631.52
pure Mn, at > 727°C [Mas2] dissolves 29.88 at.% Fe at 700°C [Mas2]
(αMn) < 727
cI58 I 43m αMn
a = 891.26
pure Mn, at 25°C [Mas2] dissolves 29.7 at.% Fe at 700°C [Mas2]
(P) (red) < 417
c*66
a = 1131
sublimation at 1 bar. Stable form of P. Triple point at 576°C, > 36.3 bar; triple point at 589.6°C at 1 atm [Mas2, V-C2]
(P) (white) < 44.14
c** P (white)
a = 718
common form of P [Mas2, V-C2]
(P) (black)
oC8 Cmca P (black)
a = 331.36 b = 1047.8 c = 437.63
T = 25°C [Mas2, V-C2]
(continued)
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Fe–Mn–P
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(MnxFe1–x)3P
tI32 I 4 Ni3P
Lattice Parameters [pm]
MnxFe1–xP
a = 913.10 c = 451.40 a = 910.46 c = 446.34 a = 910.8 c = 445.5 a = 910.81 c = 446.06 oP8 Pnma MnP
FeP < ∼ 1370
FeP2
FeP4
oP6 Pnnm FeS2 (marcasite) mP30 P21/c FeP4
Comments/References
0 ≤ x ≤ 0.3 [1977Got] a = 911.0 to 911.7 c = 446.77 to 450.0 a = 911.16 to 911.44 c = 446.77 to 449.70
Fe3P < 1166
9
a = 520.20 b = 310.82 c = 580.58 a = 519.81 b = 309.48 c = 580.10 a = 519.69 b = 309.41 c = 580.02 a = 520.92 b = 311.35 c = 581.73 a = 520.30 b = 310.60 c = 581.11 a = 519.1 b = 309.9 c = 579.2 a = 519.3 b = 579.3 c = 309.9 a = 497.29 b = 565.68 c = 272.30 a = 461.9 b = 1367.0 c = 700.2 β = 101.48°
at x = 0.1 to 0.3, after quenching from ∼ 1300°C [1977Got] at x = 0.05 to 0.25, in the alloys annealed at 950°C and slowly cooled [2000Bro] x = 0.15, 247°C [2000Bro] x = 0.15, 10 K [2000Bro] 25 at.% P at x = 0 [2002Per] after annealing at 950°C [2000Bro] 0 ≤ x ≤ 0.2 [1986Fje] x = 0.1, 20°C [1986Fje]
x = 0.1, 100 K [1986Fje]
x = 0.1, 10 K [1986Fje]
x = 0.2, 20°C [1986Fje]
x = 0.2, 10 K [1986Fje]
at x = 0 [P]
at x = 0 [2002Per], space group Pna21 66 to 67 at.% P [2002Per] at 66.7 at.% P [2002Per] 80 at.% P [2002Per] [2002Per]
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10
Fe–Mn–P
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
(Mn1–xFex)3P
tI32 I 4 Ni3P
Mn3P < 1105
Lattice Parameters [pm]
a = 918.0 to 915.4 c = 457.2 to 454.8 a = 917.12 to 913.91 c = 456.45 to 455.20 a = 918.1 c = 456.8 a = 918.26 c = 456.85
M2P, (Mn1–xFex)2P
Mn2P < 1327 a = 607.4 c = 345.4 a = 608.1 c = 346.0 a = 604.54 to 598.0 c = 346.2 to 347.5 Fe2P < 1370 a = 586.4 c = 346.0 a = 587.5 to 589.4 c = 346.3 to 346.9 -
MnP < 1147
oP8 Pnma MnP
0 ≤ x ≤ 0.33 [2003Liu] at x = 0.1 to 0.3, after quenching from ∼ 1300°C [1977Got] at x = 0.1 to 0.33, annealing at 880°C [2003Liu] 25 at.% P [V-C2]
[2003Liu] 0 ≤ x ≤ 1 at high temperature [1948Now, 1952Vog]
hP9 P 62m Fe2P
Mn3P2 1090 - 1002
Comments/References
-
0 ≤ x ≤ 0.38 after annealing at 850-900°C [1969Fru, 1969Rog] at x = 0 [P] at x = 0 after annealing at 850-900°C [1969Fru, 1969Rog] at x = 0.1 to 0.38 after annealing at 850-900°C [1969Fru, 1969Rog] 0.69 ≤ x ≤ 1 after an nealing at 850-900°C [1969Fru, 1969Rog] at x = 1 [2002Per] at x = 0.9 after annealing at 850-900°C [1969Fru, 1969Rog] 40 at.% P [Mas2]
a = 525 b = 317 c = 591.7
50 at.% P [Mas2] [1988Rag]
(continued)
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Fe–Mn–P
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
* τ1, (Mn1–yFey)2P
oP12 Pnma Co2Si
11
Lattice Parameters [pm]
Comments/References
0.38 ≤ y ≤ 0.69 (25.3 to 46 at.% Fe, 33.3 at.% P), annealing at 850–900°C [1969Fru, 1969Rog] 21.7 to 45 at.% Fe, 33.3 at.% P, heating at 1000∼1100°C for 48 h and slow cooling [1973Nag]
MnFeP
* τ2, (Mn1–yFey)3P
a = 595.3 b = 356.7 c = 673.9
at x = 0.5 (33.3 at.% Fe), annealing at 850-900°C [1969Fru, 1969Rog]
a = 596.6 to 590.0 b = 357.8 to 356.0 c = 676.5 to 670.0
at x = 0.38 to 0.69 (25.3 to 46 at.% Fe, 33.3 at.% P), annealing at 850-900°C [1969Fru, 1969Rog] 0.3 ≤ y ≤ 0.7, quenched from ∼ 1300°C [1977Got]. Existence needs further confirmation
oI*
0.4 ≤ x ≤ 0.6 [1977Got]
a = 918 to 914.5 b = 858.3 to 855 c = 486.5 to 483
Table 3.
Invariant Equilibria
Reaction
T [°C]
Type
Phase
Composition (at.%) Fe
Mn
P
L + α ⇌ γ + Fe3P
1025
U1
L
∼ 75.1
∼ 8.0
∼ 16.9
L + Fe3P ⇌ γ + M2P
1015
U2
L
∼ 65.0
∼ 18.2
∼ 16.8
L + (βMn) ⇌ γ + Mn3P
958
U3
L
∼ 35.0
∼ 50.0
∼ 15.0
L ⇌ M2P + Mn3P + γ
955
E?
L
∼ 40.4
∼ 44.2
∼ 15.4
Table 4.
Investigations of the Fe-Mn-P Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1969Fru]
Magnetic measurements (thermomagnetic balance)
Curie temperatures, magnetic moment, magnetization
[1969Rog]
Magnetic measurements (thermomagnetic balance)
Curie temperatures, magnetic moment, magnetization (continued)
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Fe–Mn–P
Reference
Method/Experimental Technique
Type of Property
[1973Nag]
Magnetic measurements (Weiss-Forrer type magnetometer, Bozorth pendulum type magnetometer)
Magnetic susceptibility, Néel and Curie temperatures, magnetic moment, magnetization
[1973Suz]
Magnetic measurements (magnetometer with a differential transformer, pendulum type magnetometer)
Magnetic susceptibility, Néel and Curie temperatures, magnetic moment
[1977Got]
Magnetic measurements (sample vibrating magnetometer; magnetic balance)
Magnetization, saturation magnetic moment, magnetic transition points
[1979Iwa]
Magnetic measurements
Magnetization, magnetic susceptibility
[1984Lee]
Bend and tension tests (Instron universal testing machine), SEM, Auger spectroscopy
Yield stress, grain boundary segregation, embrittlement and strength, hardness, tensile fracture stress, elongation
[1986Fje]
Magnetic measurements
Magnetic susceptibility, Néel temperatures, magnetic moments
[1990Bac]
Magnetic measurements, neutron diffraction
Magnetic susceptibility, magnetic ordering temperature
[1993Vav]
Electrical resistivity measurements, mechanical properties tests
Electrical resistivity, microhardness, plasticity to bending
[1997Vav]
Magnetic measurements
Effective magnetic field
[1997Zho]
Mechanical tests
Effect of Mn additions on the P embrittlement of the Fe grain boundary
[1998Vav]
Mechanical properties tests, electrical resistivity measurements
Brittle temperature, electrical resistivity
[1999Vav]
Mechanical properties tests, magnetic measurements
Microhardness, plasticity, effective magnetic field, chemical shift
[2000Bro]
Magnetization measurements (vibrating sample magnetometer (VSM); quantum design MPMS 5.5T SQUID magnetometer)
Magnetization, Curie temperatures, magnetic moments
[2000Vav]
Rapid quenching, X-ray diffraction, thermal analysis, Mössbauer spectroscopy; viscosity measurements (rotational oscillations of a quartz beaker filled with the melt
Kinematic viscosity
[2001Vav]
Mechanical properties tests
Structural hardening
[2003Liu]
Magnetization measurements (quantum design MPMS 5.5T SQUID magnetometer)
Magnetization, Néel temperatures, intensity of magnetic reflections, density of states
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Fe–Mn–P
13
Fig. 1. Fe-Mn-P. The Mn2P-Fe2P phase diagram
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Fe–Mn–P
Fig. 2. Fe-Mn-P.
Reaction scheme
14
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Fe–Mn–P
15
Fig. 3. Fe-Mn-P. Partial liquidus surface projection
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Fe–Mn–P
Fig. 4. Fe-Mn-P. Partial isothermal section at 25°C
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Fe–Mn–P
17
Fig. 5. Fe-Mn-P. Temperature-composition section at 95 mass% Fe, 0 to 4 mass% P (Mn5.08Fe94.92Mn0.98Fe92.03P6.99 in at.%)
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Fe–Mn–P
Fig. 6. Fe-Mn-P. Lines of constant molar Gibbs free energy of formation from melts at 1377°C
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Fe–Mn–P
19
Fig. 7. Fe-Mn-P. Lines of constant molar enthalpy of formation from melts at 1377°C
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20 References [1948Now] [1950Ber] [1952Vog]
[1955Sul] [1965Kan1]
[1965Kan2]
[1969Fru]
[1969Rog]
[1969Sch]
[1972Sel] [1972Sha]
[1973Mae] [1973Nag] [1973Suz] [1974Sig] [1977Got]
[1979Iwa]
[1980Bra] [1980Kun]
[1981Maa]
Fe–Mn–P
Nowotny, H., Henglein, E., “Study of Ternary Alloys with Phosphorus” (in German), Monatsh. Chem., 79, 385–393 (1948) (Crys. Structure, Phase Relations, Experimental, Review, *, 18) Berak, J., Heumann, T., “The Phosphorus-Manganese System” (in German), Z. Metallkd., 41, 19–23 (1950) (Phase Diagram, Phase Relations, Experimental) as quoted by [1955Sul] Vogel, R., Berak, J., “The Iron-Phosphorus-Manganese System” (in German), Arch. Eisenhuettenwes., 23(5–6), 217–223 (1952) (Crys. Structure, Morphology, Phase Diagram, Phase Relations, Experimental, *, 14) Sully, A.H., Manganese, Butterworth Scientific Publications, London, 295pp. (Crys. Structure, Phase Diagram, Phase Relations, Review, Phys. Prop.) Kaneko, H., Nishizawa, T., Tamaki, K., “Phosphide-Phases in Ternary Alloys of Iron, Phosphorus and Other Elements” (in Japanese), Nippon Kinzoku Gakkai-shi, 29(2), 159–165 (1965) (Morphology, Phase Diagram, Phase Relations, Experimental, Review, *, 24) Kaneko, H., Nishizawa, T., Tamaki, K., Tanifuji, A., “Solubility of Phosphorus in α- and γ-Iron” (in Japanese), Nippon Kinzoku Gakkai-shi, 29(2), 166–170 (1965) (Phase Relations, Experimental, Review, *, 20) Fruchart, R., Roger, A., Senateur, J.P., “Crystallographic and Magnetic Properties of Solid Solutions of the Phosphides M2P, M = Cr, Mn, Fe, Co, and Ni”, J. Appl. Phys., 40(3), 1250–1257 (1969) (Crys. Structure, Experimental, Magn. Prop., 45) Roger, A., Senateur, J.-P., Fruchart, R., “Crystallographic and Magnetic Properties of Solid Solutions Among the Phosphides Ni2P - Co2P - Fe2P - Mn2P and Cr2P” (in French), Ann. Chim. (Paris), 4(2), 79–91 (1969) (Crys. Structure, Experimental, Magn. Prop., 44) Schenck, H., Steinmetz, E., Gitizad, H., “Activity of Phosphorus in the Molten Iron and its Control by Nickel, Manganese and Chromium” (in German), Arch. Eisenhuettenwes., 40(8), 597–602 (1969) (Thermodyn., Experimental, 24) Selte, K., Kjekshus, A., “Structural and Magnetic Properties of FeP”, Acta Chem. Scand., Ser. A, 26(3), 1276–1277 (1972) (Crys. Structure, Experimental, Magn. Prop., 17) Shabdenov, B.A., Kunaev, A.M., Ilieva, A.A., “Thermodynamic Properties of Phosphorus in Liquid Metallic Alloys” (in Russian), Khim. Khim. Tekhnol. (Alma-Ata), 13, 228–233 (1972) (Thermodyn., Experimental) as quoted by [1995Zai] Maeda, Y., Takashima, Y., “Mössbauer Studies of FeNiP and Related Compounds”, J. Inorg. Nucl. Chem., 35(6), 1963–1969 (1973) (Crys. Structure, Experimental, Electronic Structure, 12) Nagase, S., Watanabe, H., Shinohara, T., “Magnetic Properties of the System Fe2P-Mn2P”, J. Phys. Soc. Jpn., 34(4), 908–910 (1973) (Crys. Structure, Experimental, Magn. Prop., 11) Suzuki, T., Yamaguchi, Y., Yamamoto, H., Watanabe, H., “Magnetic Structure of FeMnP”, J. Phys. Soc. Jpn., 34(4), 911–916 (1973) (Crys. Structure, Experimental, Magn. Prop., 12) Sigworth, G.K., Elliott, J.F., “The Thermodynamics of Liquid Dilute Iron Alloys”, Met. Sci., 8, 298–310 (1974) (Thermodyn., Review, 249) Goto, M., Tange, H., Tokunaga, T., Fujii, H., Okamoto, T., “Magnetic Properties of the (Fe1–xMx)3P Compounds”, Japan. J. Appl. Phys., 16(12), 2175–2179 (1977) (Crys. Structure, Experimental, Magn. Prop., 16) Iwata, N., Fujii, H., Okamoto, T., “Magnetic Properties of (Mn,Cr)P and (Mn,Fe)P Compounds”, J. Phys. Soc. Jpn., 46(3), 778–783 (1979) (Morphology, Experimental, Theory, Magn. Prop., 11) cited from abstract Brandes, E.A., Flint, R.F., “Mn-Fe-P” in “Manganese Phase Diagrams”, Manganese Centre, Paris, France, 116 (1980) (Phase Diagram, Phase Relations, Review, *, 1) Kunaev, A.M., Shabdenov, B.A., Beisembaev, B.B., Omarov, N.G., “A Study of the Thermodynamic Behavior of Manganese in Iron-Phosphorus-Manganese Melts” (in Russian), Vestn. Akad. Nauk Kaz. SSR, (7), 43–46 (1980) (Thermodyn., Experimental) as quoted by [1995Zai] Maaref, S., Madar, R., “Crystal Chemistry of M12P7 Phases in Relation with the M2P Phosphides”, J. Solid State Chem., 40, 131–135 (1981) (Crys. Structure, Experimental, Theory, 11)
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Fe–Mn–P [1983Mat]
[1983Yam]
[1984Ban]
[1984Lee]
[1986Fje]
[1988Rag]
[1988The]
[1989Paj]
[1990Bac]
[1991Sjo]
[1993Din]
[1993Vav]
[1995Bou]
[1995Zai]
[1996Zai]
[1997Vav]
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21
Matsuyama, T., Hosokawa, H., Suto, H., “Tracer Diffusion of P in Iron and Iron Alloys”, Trans. Jpn. Inst. Met., 24(8), 589–594 (1983) (Morphology, Experimental, Interface Phenomena, Kinetics, Magn. Prop., 14) Yamada, K., Kato, E., “Effect of Dilute Concentrations of Si, Al, Ti, V, Cr, Co, Ni, Nb and Mo on the Activity Coefficient of P in Liquid Iron”, Trans. Iron Steel Inst. Jpn., 23(1), 51–55 (1983) (Thermodyn., Calculation, Experimental, 16) Ban-Ya, S., Maruyama, N., Kawase, Y., “Effects of Ti, V, Cr, Mn, Co, Ni, Cu, Nb, Mo and W on the Activity of Phosphorus in Liquid Iron” (in Japanese), Tetsu to Hagane, 70(1), 65–72 (1984) (Thermodyn., Calculation, Experimenta1, Review, 21) Lee, D.Y., Barrera, E.V., Stark, J.P., Marcus, H.L., “The Influence of Alloying Elements on Impurity Induced Grain Boundary Embrittlement”, Metallurg. Trans. A. (Phys. Met. and Mater. Sci.), 15A(7), 1415–1430 (1984) (Morphology, Thermodyn., Calculation, Experimental, Interface Phenomena, Mechan. Prop., 41) Fjellvaeg, H., Kjekshus, A., Andresen, A.F., “Structural and Magnetic Properties of Mn1–tFetP (0.80 ≤ t ≤ 1.00)”, Acta Chem. Scand., Ser. A, A40, 227–229 (1986) (Crys. Structure, Experimental, Magn. Prop., 15) Raghavan, V., “The Fe-Mn-P (Iron-Manganese-Phosphorus) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Metals, Calcutta, 3, 91–99 (1988) (Crys. Structure, Phase Diagram, Phase Relations, Assessment, Review, #, 11) “The 19th Committee on Steelmaking, The Japan Society for the Promotion of Science”, Steelmaking Data Sourcebook, Gordon and Bresch Science Publishers, 280 (1988) (Morphology, Thermodyn., Review) as quoted by [1993Din] Paju, M., Grabke, H.J., “Segregation of Phosphorus in Austenite in Fe-P and Fe-10Mn-P Alloys”, Mater. Sci. Technol., 5(2), 148–154 (1989) (Morphology, Thermodyn., Experimental, Interface Phenomena, 33) cited from abstract Bacmann, M., Fruchart, D., Chenevier, B., Fruchart, R., Puertolas, J.A., Rillo, C., “Magnetic Phase Diagram of the (Fe1–xMnx)2P System”, J. Magn. Magn. Mater., 83(1–3), 313–314 (1990) (Morphology, Abstract, Phase Relations, Experimental, Magn. Prop., 13) Sjostrom, J., Jarlborg, T., “Band Structures and Magnetic Properties in the Iron Phosphide Compounds, FeMnP, FeCrP and FeVP”, J. Magn. Magn. Mater., 98(1–2), 85–91 (1991) (Crys. Structure, Calculation, Electronic Structure, Magn. Prop., 15) cited from abstract Ding, X., Wang, W., Han, Q., “Thermodynamic Calculation of Fe-P-j System Melt”, Acta Metall. Sin. (China), 29(12), B527–B532 (1993) (Thermodyn., Calculation, Phase Relations, Theory, 7) Vavilova, V.V., Kovneristyi, YuK., Levintov, M.B., “Phase Equilibrium and Tendency to Amorphization in Alloys of the System Fe-P-Mn”, Inorg. Mater. (Engl. Trans.), 29(6), 883–887 (1993), translated from Neorg. Mater., 29(6), 770–773 (Morphology, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., *, 10) Bouchard, D., Bale, C.W., “Simultaneous Optimization of Thermochemical Data for Liquid Iron Alloys Containing C, N, Ti, Si, Mn, S, and P”, Metall. Mater. Trans. B, 26B, 467–484 (1995) (Phase Relations, Thermodyn., Calculation, Review, Theory, 85) Zaitsev, A.I., Dobrokhotova, ZhV., Litvina, A.D., Mogutnov, B.M., “Thermodynamic Properties of Iron-Manganese-Phosphorus Melts”, Inorg. Mater. (Engl. Trans.), 31(9), 1064–1072 (1995), translated from Neorg. Mater., 31(9), 1164–1173 (1995) (Phase Relations, Thermodyn., Experimental, 36) Zaitsev, A.I., Litvina, A.D., Shelkova, N.E., Mogutnov, B.M., “Thermodynamic Properties of {xFe+yMn+(1–x–y)Si} (l) and {xFe+yMn+(1–x–y)P} (l)”, J. Chem. Thermodyn., 28, 285–306 (1996) (Thermodyn., Experimental, 66) Vavilova, V.V., Kovneristyi, Y.K., “Preparation and Thermal Stability of Fe-P-M (M = Mo, V, Nb, Mn, Si) Amorphous Alloys”, Inorg. Mater. (Engl. Trans.), 33(3), 275–281 (1997), translated from Neorg. Mater., 33(3), 333–339 (1997) (Phase Diagram, Phase Relations, Thermodyn., Experimental, Electr. Prop., Magn. Prop., 15)
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22 [1997Zho]
[1998Vav]
[1998Zai]
[1999Vav]
[2000Bro]
[2000Vav]
[2000Zha]
[2001Vav]
[2002Per]
[2003Bal]
[2003Liu]
[2004Wit]
[2007Wit] [Mas2] [V-C2]
Fe–Mn–P Zhong, L., Wu, R., Freeman, A.J., Olson, G.B., “Effects of Mn Additions on the P Embrittlement of the Fe Grain Boundary”, Phys. Review B. (Condens. Mater.), 55(17), 11133–11137 (1997) (Morphology, Calculation, Electronic Structure, Mechan. Prop., 14) cited from abstract Vavilova, V.V., Kovneristyi, YuK., Palii, N.A., “Correlation between the Annealing-induced Embrittlement of Fe-P-M (M = V, Mn, Si) Amorphous Alloys and the Content of α -Fe Precipitates”, Inorg. Mater. (Engl. Trans.), 34(6), 566–570 (1998), translated from Neorg. Mater., 34(6), 692–696 (1998) (Crys. Structure, Morphology, Phase Relations, Experimental, Electr. Prop., Mechan. Prop., 17) Zaitsev, A.I., Litvina, A.D., Shelkova, N.E., Mogutnov, B.M., “Association in Ternary Metallic Melts Fe-Mn-Si, Fe-Cr-P and Fe-Mn-P”, Thermochim. Acta, 314, 307–315 (1998) (Phase Relations, Thermodyn., Calculation, Experimental, 31) Vavilova, V.V., Baldokhin, Y.V., “Mössbauer Study of Rapidly Quenched Fe-P-E Alloys (E = V, Nb, Mo, Mn, Si)”, Russ. Metall. (Engl. Transl.), (1), 122–132 (1999) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Experimental, Review, Magn. Prop., Mechan. Prop., 20) Broddefalk, A., James, P., Liu, H.-P., Kalska, B., Anderson. Y., Granberg, P., Nordblad, P., Haeggstroem, L., Eriksson, O., “Structural and Magnetic Properties of (Fe1–xMnx)3P (x < 0.25)”, Phys. Rev. B, 61(1), 413–421 (2000) (Crys. Structure, Calculation, Experimental, Magn. Prop., 25) Vavilova, V.V., Baldokhin, Y.V., Kovneristyi, Y.K., Matveev, V.M., “Fe-P-M (M = Si, Mn, V) Alloys: Viscosity in the Liquid State and Tendency to Amorphization”, Inorg. Mater. (Engl. Trans.), 36(7), 703–708 (2000), translated from Neorg. Mater., 36(7), 845–851 (2000) (Morphology, Thermodyn., Calculation, Experimental, Kinetics, Phys. Prop., 13) Zhang, J., Zhu, R., “Thermodynamic Properties of Mn-P and Fe-Mn-P Melts”, J. Univ. Sci. Technol. Beijing, 7(1), 10–13 (2000) (Phase Diagram, Phase Relations, Thermodyn., Calculation, 6) cited from abstract Vavilova, V.V., Kovneristyi, Y.K., Palii, N.A., “Physico-Chemical Properties of Amorphous and Crystalline Fe-P-Mn and Fe-P-Mn-V Alloys”, Inorg. Mater., 37(2), 166–169 (2001) (Morphology, Experimental, Mechan. Prop., 5) cited from abstract Perrot, P., Batista, S., Xing, X., “Fe-P (Iron-Phosphorus)”, MSIT Binary Evaluation Program, in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; Document ID: 20.16107.1.20, (2002) (Crys. Structure, Phase Diagram, Phase Relations, Thermodyn., Assessment, Phys. Prop., #, 23) Baldokhin, Y.V., Vavilova, V.V., Kolotyrkin, P.Y., Kovneristyi, Y.K., Palii, N.A., Solomatin, A.S., “Nanoscale Crystallization in Amorphous Fe-P-Mn Alloys”, Inorg. Mater., 39(6), 562–567 (2003) (Crys. Structure, Morphology, Experimental, 8) cited from abstract Liu, H.-P., Andersson, Y., James, P., Satula, D., Kalska, B., Haeggstroem, L., Eriksson, O., Broddefalk, A., Nordblad, P., “The Antiferromagnetism of (Fe1–xMnx)3P, x ≥ 0.67, Compounds”, J. Magn. Magn. Mater., 256, 117–128 (2003) (Crys. Structure, Phase Relations, Calculation, Experimental, Electronic Structure, Magn. Prop., 26) Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diff., 25(4), 346–354, (2004) (Phase Diagram, Phase Relations, Thermodyn., Assessment, Calculation, 34) Witusiewicz, V.T., private communication to MSI, (2007) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_19 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
1
Iron – Manganese – Sulfur Vasyl Tomashik, Hans-Leo Lukas
Introduction Numerous studies of the Fe-Mn-S ternary system have been done since 1913 [1913Roe] because of its importance in understanding desulfurization of molten steels and the formation of sulfides in the solidification process [1976Hil2, 1980Bra, 1988Rag, 2000Oht]. The first experimental results concerning this system were included in the review of [1949Jae]. A critical assessment of the Fe-Mn-S system has been published by [1988Rag], which includes the literature up to 1983. The literature up to 2000 is discussed in [2004Rag]. The present assessment takes into account all available data. Most of the studies of this ternary system have been completed in the Fe-FeS-MnS-Mn part of the composition triangle. The FeS-MnS system is described by a phase diagram of the eutectic type with the eutectic temperature reported between 1110 and 1181°C and eutectic composition of 6 and 10 mol% MnS [1913Roe, 1928Shi, 1964Cha, 1971Ski, 1976Man, 1978Fis2, 1980Ito]. The existence of a ternary compound Fe3Mn2S5, reported by [1913Roe] with a melting temperature of 1365°C, was not confirmed by the later investigations. This composition corresponds to the FexMn1–xS solid solutions which crystallize in the cubic structure of MnS extending at 700°C up to the composition Fe0.6Mn0.4S [1971Bru]. The phase diagram of the quasibinary FeS-MnS system at 1000°C and high pressures was determined by [1984McC]. The first experimental results on the mutual solubility of FeS and MnS were obtained by [1928Shi, 1964Cha, 1966Kie, 1969Sal]. The data of [1928Shi] were included in the review of [1964Kul]. Later the mutual solubility was investigated by [1971Ski]. Much work has been done in studying the liquidus of the Fe-FeS-MnS-Mn part of the Fe-Mn-S system [1933Vog, 1934Mey, 1937Vog, 1937Wen, 1971Tur, 1972Kor, 1975Big, 1975Ito2, 1975Sch, 1975Tur, 1976Man, 1980Ito, 1983Uto]. The phase relations in the Fe-Mn-S system at 1330 and 1615°C were investigated by [1971San] and included in the review of [1980Bra]. Phase diagrams calculated from thermochemical data are given by [1976Hil1, 1978Uhr, 1976Hil2, 1981Nis, 1998Mie, 1998Val, 2000Oht]. Isothermal sections of the Fe-Mn-S system at 1610, 1600, 1330, 1300, 1200, 1100, 950, 800°C were constructed by [1963Kan, 1971Nak, 1971Ski, 1973Kir, 1975Sch, 1978Fis2]. Some polythermal sections of the Fe-Mn-S system were constructed by [1933Vog], [1937Vog] and [1975Ito2]. The solid solubility of S in (γFe) at 1000, 1200 and 1335°C in the presence of 0.37, 0.40, 1.07, 1.16 and 1.30 mass% Mn was determined by [1955Tur1, 1955Tur2]. The presence of Fe in Mn melts leads to decreasing of the S solubility [1973Das]. Chemical compositions of sulfides in Fe-Mn-S ternary alloys containing between 0.01 and 1.0 mass% S and 0.01 to 1.5 mass% Mn within the temperature range 800 to 1500 were calculated by [1976Efi]. The equilibrium in the interaction FeS + Mn ⇌ MnS + Fe at 1600°C is shifted to the formation of MnS [1934Mey]. According to [1968Ska] FexMn1–xS solid solutions are formed during cooling Fe-Mn-S alloys containing 2.60 or 1.88 mass% Mn and 0.051 or 0.78 mass% S, respectively. [1971Nak] reports in the equilibrium (γFe)-FeS-MnS at 1300°C the compositions of the phases to be Fe0.988Mn0.012S, Fe0.642Mn0.358S and (γFe), containing 0.15 mass% Mn and 0.009 mass% S. Thermodynamic properties of solid and liquid alloys in the Fe-Mn-S ternary system were investigated by [1952She, 1955Tur1, 1966Sch, 1967Bro, 1969Ban, 1969Sch, 1970Nis, 1973Buz, 1973Das, 1974Sig, 1977Fis, 1995Bou, 1998Val]. Table 1 lists the numerous experimental works on the phase equilibria, crystal structures and thermodynamics of the Fe-Mn-S system. Binary Systems The binary systems are accepted as calculated from the dataset of [1998Mie]. The Fe-Mn dataset is adopted from [1989Hua], Fe-S and Mn-S were assessed in the work of [1998Mie]. Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
2
Fe–Mn–S
A newer calculation of the Fe-Mn system by [2004Wit] differs only within the accuracy of the experimental knowledge. The binary systems given by [Mas2] agree fairly well with those of [1998Mie]. Solid Phases Crystallographic data of solid phase are given in Table 2. The solid solubility of S in (γFe) at 1000, 1200 and 1335°C was found to be 0.013, 0.031 and 0.046 mass%, respectively [1955Tur1]. In the presence of 0.37, 1.07 and 1.30 mass% Mn the S solubility is 0.0018, 0.0007 and 0.0006 mass%, respectively, at 1200°C (0.0019 and 0.007 mass% S at 0.40 and 1.16 mass% Mn respectively [1955Tur2]) and increases up to 0.0058, 0.0031 and 0.0018 mass% at 1335°C [1955Tur1]. The effect of Mn on the S solubility in Fe at 1600°C is shown in Fig. 1 [1973Buz]. The solubility of MnS in FeS and FeS in MnS at 600, 700, 800, 900 and 1000°C is equal to 5.0, 6.0, 6.6, 7.0, 7.4 and 53.6, 58.8, 64.6-65.0, 68.0-68.8, 73.6-73.8 mol%, respectively [1971Ski]. The mutual solubilities of FeS and MnS significantly depend on the S vapor pressure [1980Nis]. The Fe content in rambergite (FexMn1–xS) varies between 0.1 and 6 at.% for different samples [2001Eri]. FeS and other binary sulphides are in equilibrium with the Fe phase in the Fe-Mn-S system [1963Kan]. Quasibinary Systems By comparing the locations of the solvus at the MnS side of the quasibinary FeS-MnS phase diagram of [1928Shi] with own results, [1971Ski] concluded that the temperatures reported by [1928Shi] may be overestimated. [1976Man] reported the eutectic temperature to be 1110°C and the eutectic composition to be about 7.5 mol% MnS. Although [1928Shi] stated that FeS in their studies was stoichiometric, [1976Man] found some metal-deficiency in FeS. More reliable eutectic temperatures and eutectic compositions in the system Fe1–xS-MnS are 1110°C and about 10 mol% MnS respectively, proposed in the review of [1988Rag] (Fig. 2, dashed lines). A calculation using the dataset of [1998Mie] gives a higher eutectic temperature of 1181°C but the eutectic compositions are nearly the same (Fig. 2, solid lines). This may be an artefact of modelling the phases P and Q to be restricted to strictly stoichiometric 50 at.% S. The lattice parameters of the FexMn1–xS solid solutions based on MnS can be expressed by the equation a = 522.4 - 0.1653 mol% FeS (pm) [1984McC]. The phase diagram of the FeS-MnS system determined at 1000°C in high-pressure experiments is illustrated in Fig. 3 [1984McC]. The maximum concentration of MnS in FeS decreases from 7.4 mol% at 0.1 MPa to 5.4 mol% at 3.5 GPa, and remains relatively constant to at least 7 GPa. The limiting solubility of FeS in MnS appears to change dramatically between 6 and 7 GPa. It is likely that the discontinuity is caused by a high-pressure phase transition. [1988Rag] indicated the Fe-MnS section to be also quasibinary, but precise experimental data on this section are not available. The quasibinary character can be only an approximation as the monotectic maximum L1 ⇌ Fe + Q and the eutectic maximum L2 ⇌ Fe + Q may be somewhat off this plane and there is no reason, why both these three-phase maxima should be exactly in the same plane. In calculations using the dataset of [1998Mie] the monotectic appears in the plane going from Fe1.04Mn48.96S50 to Fe90.12Mn9.88 and the eutectic in the plane from Fe0.57Mn49.43S50 to Fe99.0Mn1.0. Invariant Equilibria A reaction scheme for the Fe-FeS-MnS-Mn part of the Fe-Mn-S system is given in Fig. 4, calculated from the dataset of [1998Mie]. The calculated phase compositions are given in Table 3. The deviation of the sulfur content from stoichiometry in the P phase is neglected in the dataset of [1998Mie], thus the composition of sulfur in P and Q appears always as exactly 50 at.% in Table 3. The three different miscibility gaps of liquid were first time mentioned by [1976Hil1], their temperatures and phase compositions must be taken as tentative as not enough experimental data are available. The parts below U4 agree with the reaction scheme assessed in the review of [1988Rag], except U4 itself, where [1988Rag] gave the reaction: (δFe) + L ⇌ Q + γ, whereas the calculation yields: (δFe) + Q ⇌ L + γ. [1988Rag] as a consequence postulated a minimum of the three-phase field (δFe) + Q + γ between U2 and U4. The three-phase equilibria along the sequence p2-U5-U2-U4-e4 containing the two phases γ and (δFe) describe mainly the transformation DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
3
(δFe) ⇌ γ, the third phases participate only with very small amounts in this transformation. Both modifications of Fe are very dilute with respect to S and Mn, and the S content in both phases decreases monotonically along this sequence and the Mn content increases monotonically. Due to Henry’s law thus also the temperature must change monotonically and a minimum is very unlikely. The congruent formation of (βMn) from γ in the Fe-Mn binary system results in a pair of γ + (βMn) + Q phase fields at ∼ 1150°C. The order-disorder transitions in Fe1–xS at lower temperatures and the transitions in (S) are not considered separately in the reaction scheme. Regarding the sulfur rich part of the system [1988Rag] indicated that there are literature data suggesting a transition reaction P + L ⇌ Q + FeS2 just below the Fe-S peritectic temperature of 743°C, by which pyrite comes into two-phase equilibrium with the Q phase. [1988Rag] assumed the maxima e1 and e3 to be exactly on the line Fe-MnS. Thermodynamically there is no reason for this assumption and the calculation gives deviations, for e1 much more than for e3. Liquidus, Solidus and Solvus Surfaces A liquidus surface of the Fe-FeS-MnS-Mn part of the Fe-Mn-S system was constructed by [1988Rag] using the experimental results of [1933Vog, 1934Mey, 1937Vog, 1937Wen, 1971Tur, 1972Kor, 1975Big, 1975Ito2, 1975Sch, 1975Tur, 1976Man, 1980Ito, 1983Uto]. The main features, especially those proved by experimental data agree with the liquidus surface calculated from the dataset of [1998Mie] shown in Fig. 5a. The isotherms are drawn in 100°C steps up to 2000°C. Above this temperature in the Mn corner appears the gas phase (at 0.1 MPa pressure). Parts of this liquidus projection are shown enlarged in rectangular coordinates in Figs. 5b to 5d. A liquid miscibility gap that originates on the Mn-MnS side dominates this region. The eutectic through, L + γ + Q, starting near the Mn corner at e5 runs around the liquidus boundary of the miscibility gap. In the calculated liquidus surface it is interrupted by reaction U3 and ends at U6. Contrary to the calculation and to some of the earlier studies including that of [1975Sch], [1988Rag] assumed that the miscibility gap does not reach and interrupt this eutectic through at any point. There is a temperature maximum e3 at 1510°C on the Fe-MnS quasibinary section, however not exactly on the join FeMnS. A number of early reports give the invariant reaction U6 as a ternary eutectic. However, the results of [1976Man] and [1980Ito] confirm the earlier finding of [1937Vog] that this reaction is transition [1988Rag]. According to thermochemical calculations there are three miscibility gaps in the Fe-FeS-MnS-Mn subsystem [1976Hil1]. These authors calculated the locus of the critical points of the miscibility gaps. This calculation is not directly implemented in Thermo-Calc. Figure 5 does not show this locus, but the univariant lines of three-phase equilibria containing two or three different liquids, thus the results of [1976Hil1] and Fig. 5 are difficult to compare. The calculated liquidus surface near the Fe corner is shown in Fig. 6, it agrees well with a diagram given by [1975Big]. Isothermal Sections The calculated isothermal section at 1600°C is given in Fig. 7. The diagram reported by [1975Sch] agrees well at the metal rich side but shows a significantly lower S concentration in the sulfidic melt. The scatter of the experimental data however is even larger and thus this problem cannot be solved. Furthermore [1975Sch] shows only a single critical point at 4.5 at.% Mn and 33.5 at.% S instead of the two calculated ones. The calculated isothermal sections at 1500, 1400, 1300, 1200, 1100 and 1000°C are given in Figs. 8 to 13. In the two-phase fields, where the direction of tie-lines is an important information, some calculated tie lines are drawn. This is the field L + Q in Figs. 7 and 8 and γ + Q in Figs. 10 to 13. In Fig. 14 the composition of FexMn1–xS (Q) in equilibrium with γ alloys is shown in more detail at 1100, 1200 and 1300°C from [1988Rag] due to [1978Fis1]. Figure 15 shows an isothermal section at 800°C according to [1971Ski] and [1988Rag], showing the homogeneity range of Fe1–xMnxS1+y (P). Temperature – Composition Sections Five polythermal sections of the Fe-Mn-S system from the S corner to the Fe-Mn binary system and one section from Mn to the Fe-S binary system were constructed by [1933Vog]. [1937Vog] constructed 4 polythermal sections: 97Fe3Mn (mass%) - S, 94Fe6Mn (mass%) - S, 90Fe10Mn (mass%) - S and 80Fe20Mn Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
4
Fe–Mn–S
(mass%) - S. Some vertical sections were calculated by [2000Oht] and it was shown that the obtained phase boundaries agree well with the experimental data points from [1937Vog], although the configuration of the phase fields is different from that drawn by [1937Vog]. An Fe1–xS-MnS section “in equilibrium with excess Fe” has been investigated by [1975Ito1, 1975Ito2, 1976Man, 1978Fis2, 1980Ito]. Figure 16 gives this FeS-MnS figure calculated from the dataset of [1998Mie]. Figure 16 is not really a T-c section but a projection of the univariant lines of the three-phase fields containing iron (α or γ) and any two of the three phases L, P or Q, projected from the S corner onto the Fe-Mn plane. The invariant horizontal at 1005°C corresponds to the four-phase equilibrium U6. Thermodynamics The activity coefficient of sulfur in Fe-Mn-S alloys decreases with increasing Mn content. The magnitude of this effect varies with temperature and Mn content of the alloy [1952She, 1955Tur1, 1955Tur2, 1966Sch, 1969Ban, 1969Sch, 1970Nis, 1973Buz]. The first order interaction coefficients are eMnS = –0.025 and εMnS = –5.7 at 1600°C and up to 5 mass% Mn [1952She] and εSMn = –0.048 at 1600°C [1974Sig] (eMnS = –0.017 and εMnS = –3.86 at 1600°C and up to 3.5 mass% Mn [1966Sch, 1969Sch]; eMnS = –0.026 and interaction parameter θMnS is equal –5.87 at 1550°C and up to 8 at.% Mn [1969Ban, 1974Sig]; eMnS = –0.021 and εMnS = –4.76 at 1600°C and at 0.5-20 mass% S [1973Buz]). [1998Val] notes that eMnS = –0.024 and eSS = 0.0087 at 1487°C. The temperature dependence of eMnS at 1200–1600°C can be expressed by the equation eMnS = –215/T + 0.097 [1955Tur1]. The calculated values of eSMn reproduce the experimental data satisfactorily and can be expressed at the temperatures 1230–1930°C by the equation eSMn = –420683.7/T + 218.5 [1995Bou]. [1970Nis] reports a very strong interaction between Mn and S in δFe around 1500°C: eMnS = –0.58, valid up to 0.15 mass% Mn. The second order interaction coefficient rMnS is equal zero [1974Sig]. According to the data of [1952She] up to and possibly beyond 5 mass% Mn the activity coefficient f MnS is independent on the S content. The first order interaction coefficient eFeS is equal to 0.008 between 1300 and 1500°C and up to 22 mass% Fe [1973Das]. [1977Fis] equilibrated (Fe,Mn)-sulfide solid solutions with a gas phase of H2/H2S mixtures and measured the ratio p(H2S)/p(H2), which is proportional to the S-activity. Thermochemical data on the Fe-Mn-S system are summarized on a single sulfur potential diagram by [1967Bro], which allows to predict the solubility of S in liquid and solid Fe-Mn alloys. Thermodynamic datasets to calculate phase diagrams of the Fe-FeS-MnS-Mn partial system were assessed by [1976Hil2, 1998Mie, 2000Oht]. Notes on Materials Properties and Applications Mn has a favorable influence on the hot workability of steel [1955Tur1, 1960Mch, 1971Tur, 1976Hil2]. Without Mn hot-shortness may occur in steel containing more than 0.01 mass% S. Presence of Mn appreciably reduces the S solubility. The amount of sulfide inclusions formed during cooling from 1200–1300°C is almost negligible as MnS inclusions, which may also contain some ferrous sulfide are already solid before solidification starts. There is a limiting Mn content for each S concentration, above which no hot-short cracks appear in the steel [1960Mch]. Mn does not influence the distribution of sulfide inclusions but influences their dimensions and chemical compositions. The sulfide inclusions become smaller and contain more MnS with increasing Mn content [1960Mch]. The hot-short range varies with the type of steel, particularly with the Mn/S ratio [1971Tur]. Cold and hot fragility of Fe-Mn-S alloys is related to the formation of sulfides depleted by Mn which have a low melting temperature [1973Wyj]. Outside the critical range of Mn content these sulfides are replaced by intragranular sulfides enriched in Mn and such alloys are ductile in the cold state and forgeable in the hot state. Electric resistance of the FexMn1–xS solid solutions decreases with increasing x up to x = 0.6 [1971Bru]. In the field of the NiAs type hexagonal structure at 90–100 at.% Fe the solid solutions have a metallic conductivity like pure FeS. According to [1978Shi] the viscosity of FeS-MnS melts at 1300–1400°C slightly increases with increasing MnS content. The coefficients of self-diffusion of the ions in such melts were estimated using experimental data. It was shown that most probably the particle responsible for the viscous flow is the S–2 ion. DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
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Fe–Mn–S
5
MnS is not hardened measurably by dissolved FeS [1964Cha]. Excess FeS will, however, increase the hardness of the sulfide by the formation of a two-phase microstructure. Miscellaneous Sulfides of the MnS type are common non-metallic inclusions in steel [1966Kie]. The change in physical properties of the sulfide phase due to solid solution can be rather large. The change in composition of the steel matrix around MnS inclusions due to diffusion phenomena of solute atoms may also influence the deformation properties of the steel. [1969Sal] indicates that MnS inclusions in Fe-Mn-S alloys occur either as isolated inclusions or in groups, particularly in the interdendritic interstices. According to [2004Liu] MnS is nucleated more easily in γFe rather than in αFe and MnS precipitation in γFe can be suppressed by rapid cooling. The particles of MnS are usually larger than 100 nm. [1975Tur] indicates that aside from the precipitation of minute MnS inclusions along the interdendritic regions, there are larger particles of MnS with diameters of approximately 20 μm which contain one or two pieces of Fe particles. According to [1937Wen] decreasing of the ratio Mn/S in Fe-Mn-S alloys containing 0.31–2.0 mass% Mn and 0.27–2.03 mass% S, leads to increasing of the number of inclusions. Sulfides are formed in the alloys by S transfer from a synthetic slag to a surrounding Fe-Mn matrix at 1100 to 1450°C [1971His]. The Mn concentration in the matrix decreases rapidly. The region of precipitation of sulfide particles widens with increasing of time and temperature and with decreasing of initial Mn matrix concentration. Sulfide precipitates coagulate to larger particles during repeated solution and precipitation. As a result, there appears a zone without precipitates in the neighborhood of the slag/Me interface. The shapes and locations of the precipitates are affected by the oxygen concentration of the alloy. Si in the alloy prevents somewhat the S transfer and the precipitation of sulfides [1971His]. [1971Kor] notes that during oxidation of Fe containing 0.34 mass% Mn and 150 ppm sulfur depletion of Mn and accumulation of S occurs in the metal near the scale/metal interface. These changes of the Mn and S concentrations lead to the formation of liquid oxysulfide near the surface at 900°C. The control of morphology of MnS inclusions plays an important role in the field of ferrous metallurgy [2000Oht]. In S-lean melts the morphology of secondary MnS inclusions formed during solidification after primary crystallization of the Fe phase can be classified under the following three categories [1995Oik]: a) globular or droplet shaped MnS resulting from a monotectic reaction; b) rod-like MnS formed by an eutectic reaction, or c) a fish-bone type MnS formed as a result of an irregular eutectic reaction. It was found that the corrosion kinetics of the Fe-Mn alloys in pure S vapor (101.3 kPa) at 700–1000°C can be divided into four groups, depending on the alloy concentration [1980Nis]. The first group, up to 11 mass % Mn, has a constant rate. The second group, up to ∼ 63 mass% Mn has an exponentially decreasing rate of corrosion with increasing Mn content. In the third group, ∼ 63-80 mass% Mn alloys, the corrosion rate does not follow the parabolic law and the corrosion rates become constant above 80 mass% Mn. According to [1987Pap] the sulfidation kinetics are parabolic in most cases for the alloys containing from 20.3 to 99.95 mass% Mn, but two different diffusion processes affect this kinetics. In the sulfidation at 700 and 800°C the diffusion of Mn atoms in the alloy is the rate-determining step and an interlocked scale/alloy interface is observed. At 900°C there is a transition to rate-control by the diffusion of Mn in the sulfide scale connected with transition from rugged to a planar scale/alloy interface. Steady-state parabolic kinetics were always observed after an initial period of the sulfidation of the Fe + 27 mass% Mn alloy at the temperatures 700–900°C in flowing H2S/H2/N2 atmospheres corresponding to equilibrium S pressures of 8 Pa [1990Sou]. The deviation of the S-content from stoichiometry of MnS based solid solutions FexMn1–xS is very narrow (about 10–3) but in FeS based solid solutions it increases approximately by 2 orders [1973Bru, 1980Nis]. The results of a qualitative analysis of sulfide inclusions in Fe-Mn-S alloys indicate that even at Mn:S ratios up to 5:1 such inclusions are the MnS based FexMn1–xS solid solutions containing from 28 to 52 mass% Mn, from 3.5 to 25 mass% Fe and from 25 to 32 mass% S [1971Ska]. It was found by Mössbauer measurements on Fe0.975Mn0.025S that Mn impurities profoundly affect both, the crystallographic and the spin rotation transitions of FeS [2005Nam].
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DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
6 Table 1.
Fe–Mn–S Investigations of the Fe-Mn-S Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1913Roe]
DTA, metallography
Up to 1620°C/FeS-MnS
[1928Shi]
DTA, metallography, chemical analysis
Up to 1610°C/FeS-MnS
[1933Vog]
DTA, metallography
Up to 1500°C/Fe-FeS-MnS-Mn
[1934Mey]
Chemical analysis
1600°C/FeS + Mn
[1933Vog]
DTA, metallography, chemical analysis
Up to 1600°C/Fe-FeS-MnS-Mn
[1937Wen]
Metallography
Up to 1600°C/Fe-Mn-S at 0.31-2.0 mass% Mn and 0.27–2.03 mass% S
[1952She]
Equilibration with H2S/H2 gas and chemical analysis
Up to 1600°C/(Fe-Mn) + H2S/H2 up to 8.2 mass% Mn
[1955Tur1]
Equilibration with H2S/H2 gas and chemical analysis
1200–1600°C/(Fe-Mn) + H2S/H2 up to 1.3 mass% Mn
[1955Tur2]
Equilibration with H2S/H2 gas and chemical analysis
1200–1600°C/(Fe-Mn) + H2S/H2 at 0.4 or 1.16 mass% Mn
[1960Mch]
Metallography, chemical analysis
1520 and 1600°C/Fe-Mn-S at 0.20-0.30 mass% S and up to 3 mass% Mn
[1963Kan]
XRD, chemical analysis
Fe-Mn-S
[1964Cha]
Microhardness measurements
Up to 1600°C/FeS-MnS
[1966Kie]
XRD, EPMA, microhardness measurements
1150°C/FeS-MnS
[1966Sch, 1969Sch]
Chemical analysis
1600°C/Fe-Mn-S up to 3.5 mass% Mn
[1968Ska]
EPMA
Fe + 2.60 or 1.88 mass% Mn + 0.051 or 0.78 mass% S
[1969Ban]
Chemical analysis
1550°C/Fe-Mn-S
[1969Sal]
Chemical analysis
Fe-Mn-S
[1970Nis]
Chemical analysis
1600°C/Fe-Mn-S up to 0.15 mass% Mn
[1971Kor]
EPMA
900–1300°C/Fe-Mn-S
[1971Nak]
EPMA, metallography
1300°C/γFe-FeS-MnS up to 9.3 mass% Mn
[1971San]
EPMA, levitation melting
1330 and 1615°C/Fe-Mn-S
[1971Ska]
EPMA, TEM
(Fe + 1.88 or 2.6 mass% Mn)-S
[1971Ski]
XRD, EPMA
600–1000°C/FeS-MnS (continued)
DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
7
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1973Bru]
TGA, XRD
600–900°C/FexMn1–xS
[1973Buz]
XRD, EPMA
1600°C/Fe-Mn-S
[1973Das]
Chemical analysis
1300°C/Fe-Mn-S up to 22 mass% Fe
[1973Kir]
EPMA, metallography
1190–1380°C/Fe-Mn-S
[1973Wyj]
EPMA, metallography, SEM
Fe + 0.05-0.50 mass% Mn + 0.25–0.30 mass% S
[1975Big]
DTA, SEM
Up to 1650°C/Fe-Mn-S
[1976Man]
XRD, EPMA, DTA, metallography
Up to 1160°C/Fe-FeS-MnS-Mn
[1977Fis]
XRD, EPMA, metallography, vapor pressure measurements
1100–1400°C/Fe-Mn-S
[1978Fis1]
Chemical analysis, metallography
1100–1300°C/(Fe-Mn) + H2S/H2 up to 0.39 mass% Mn
[1978Fis2]
EPMA, chemical analysis, metallography
1100–1300°C/Fe-Mn-S
[1980Ito]
X-ray microanalysis, chemical analysis
670–1500°C/Fe-Mn-S
[1980Nis]
XRD, EPMA, TGA, metallography
670–1500°C/Fe-Mn-S
[1983Uto]
Metallography, chemical analysis, X-ray microanalysis
1600°C/Fe-Mn-S
[1984McC]
XRD,EPMA, Mössbauer spectroscopy
Up to 1000°C/FeS-MnS
[1987Pap]
TGA, SEM, chemical analysis, X-ray EDA, metallography
700, 800 and 900°C/(Fe + 27 mass% Mn) + H2S/ H2/N2
[1990Sou]
TGA, XRD, chemical analysis, X-ray EDA
700, 800 and 900°C/(Fe-Mn) + H2S/H2 at 20.3, 29.9, 40.6, 50 and 99.95 mass% Mn
[1995Oik]
SEM, metallography
Fe + 1 mass% Mn + 0.3 mass% S and Fe + 2.5 mass% Mn + (1–1.7) mass% S
[2000Kim]
XRD, Mössbauer spectroscopy
77–600 K/Fe0.975Mn0.025S
[2001Eri]
XRD, EPMA
Room temperature/Fe0.05Mn0.95S
[2004Liu]
TEM, EDS
Up to 800°C/Fe-Mn-S
[2005Nam]
XRD, Mössbauer spectroscopy
77–600 K/Fe0.975Mn0.025S
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DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
8 Table 2.
Fe–Mn–S Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(δFe)1538–1394
cI2 Im 3m W
a = 293.15
dissolves up to 10 at.% Mn pure Fe at 1390°C [V-C2, Mas2]
γ, (γFe,γMn) (γFe) 1394–912 (γMn) 1138–1100
cF4 Fm 3m Cu
a = 364.67
continuous solid solution at 915°C [Mas2]
a = 386.0
[Mas2]
α, (αFe) < 912
cI2 Im 3m W
a = 286.65
dissolves up to 3 at.% Mn, up to 0.033 at.% S pure Fe at 25°C [Mas2].
(δMn) 1246–1138
cI2 Im 3m W
a = 308.0
dissolves up to 9 at.% Fe pure Mn [Mas2].
(βMn) 1100–727
cP20 P4132 βMn
a = 631.52
dissolves up to 35 at.% Fe pure Mn [Mas2]
(αMn) < 727
cI58 I 43m αMn
a = 891.26
dissolves up to 35 at.% Fe pure Mn at 25°C [Mas2]
(βS) 115.22–95.5
mP64 P21/c βS
a = 1102 b = 1096 c = 1090 β = 96.7°
[Mas2]
(αS) < 95.5
oF128 Fddd αS
a = 1046.4 b = 1286.60 c = 2448.60
pure S at 25°C [Mas2]
P, γFeS
hP4 P63/mmc NiAs
a = 344.36 ± 0.05 c = 587.59 ± 0.05
[V-C2, Mas2]
βFeS 315–∼60
hP24 P 62c FeS, troillite
Fe0.975Mn0.025S
a = 596.3 ± 0.1 c = 1175.4 ± 0.1 a = 586.1 c = 1157.7 ± 0.1 a = 599.8 ± 1.1 c = 1171 ± 1 a = 596.7 ± 0.2 c = 1172 ± 5
mineral troillite at 21°C [V-C2, Mas2] at 21°C and 3.33 GPa [V-C2, Mas2] at 120°C [V-C2, Mas2] [2000Kim, 2005Nam] (continued)
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
9
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
αFeS < 138
hP6 P63/mmc FeS
a = 345.59 ± 0.05 c = 577.89 ± 0.05
[V-C2, Mas2]
FeS(I)
oP8 Pnma MnP
a = 582.5 ± 0.2 b = 346.8 ± 0.1 c = 693.5 ± 0.6 a = 571.6 ± 0.9 b = 334.7 ± 0.3 c = 669.4 ± 0.9 a = 565 ± 1 b = 331.6 ± 0.3 c = 663.1 ± 0.8
at 190°C [V-C2]
at 21°C and 4.15 GPa [V-C2]
at 21°C and 6.35 GPa [V-C2]
FeS (mackinawite)
tP4 P4/nmm PbO
a = 376.8 c = 503.9
mineral mackinawite [V-C2]
αFeS2 ≲ 444.5
oP6 Pnnm FeS2 (marcasite)
a = 444.1 b = 542.5 c = 338.7 a = 446.4 b = 544 c = 339 a = 443 ± 2 b = 550 ± 4 c = 342 ± 4 a = 446 b = 538 c = 335 a = 446 b = 540 c = 336
mineral marcasite [V-C2, Mas2, 1982Kub] at 327°C [V-C2]
at p = 16.8 GPa [V-C2]
at p = 29.1 GPa [V-C2]
at p = 33.4 GPa [V-C2]
βFeS2 < 743
cP12 Pa 3 FeS2 (pyrite)
Fe2S3
tP80 P43212
a = 1053 c = 1001
[V-C2]
Fe3S4
cF56 Fd 3m Al2MgO4
a = 987.6 ± 0.2
mineral greigite [V-C2]
Fe3S4 (smythite)
hR21 R 3m Fe3S4
a = 347 ± 2 c = 3450 ± 20
mineral smythite [V-C2]
a= a= a= a=
541.79 ± 0.11 534.8 ± 0.2 529.3 ± 0.2 525.5 ± 0.2
mineral pyrite [Mas2, 1982Kub] [V-C2, Mas2] at 1.57 GPa [V-C2] at 2.87 GPa [V-C2] at 3.85 GPa [V-C2]
(continued)
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10
Fe–Mn–S
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
Fe7S8
hP45 P3121 Fe7S8
a = 686.52 ± 0.06 c = 1704.6 ± 0.2
metastable phase, mineral pyrrhotite-3C [V-C2]
Q, MnS < 1655 Fe0.65Mn0.35S
cF8 Fm 3m NaCl
a = 522.26
mineral alabandite [V-C2], dissolves up to 65 mol% FeS [1966Kie] [1966Kie] [1984McC]
MnS
hP4 P63mc ZnS
a = 398.8 c = 643.3 a = 398.2 ± 0.2 c = 644.5 ± 0.3
[V-C2] mineral rambergite [2001Eri]
MnS
cF8 F 43m ZnS
a = 559
high temperature phase [V-C2]
MnS2
cP12 Pa 3 FeS2
a= a= a= a=
mineral mauerite, at 20°C [V-C2] at 110°C [V-C2] at 200°C [V-C2] at 252°C [V-C2]
Fe0.05Mn0.95S
a = 513.0 a = 511.66
610.16 611.09 611.80 612.24
a = 607 a = 598 a = 587
Table 3.
at p = 2.13 GPa [V-C2] at p = 7.35 GPa [V-C2] at p = 12.3 GPa [V-C2]
Invariant Equilibria
Reaction or Critical Tie Line
Type
Temperature [°C]
Phase
Composition (at.%) Fe
Mn
S
L1, L3 - L2
c1
1829
L1, L3 L2
70.1 9.7
7.6 40.7
22.3 49.6
L2 ⇌ L1 + Q
e1
1672
L2 L1 Q
1.1 88.9 0.6
49.2 9.8 49.4
49.7 1.3 50.0
L1 + L2 ⇌ L3 + Q
U1
1548
L1 L2 L3 Q
83 14 58 9
1 36 7 41
16 50 35 50
L1 ⇌ (δFe) + Q
e3
1510
L1 (δFe) Q
96.5 98.14 1.0
2.6 1.83 49.0
0.9 0.03 50.0 (continued)
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
Reaction or Critical Tie Line
Type
11
Temperature [°C]
Phase
Composition (at.%) Fe
Mn
S
L1 + (δFe) ⇌ γ + Q
U2
1472
L1 (δFe) γ Q
87.4 90.6 90.5 0.2
12.3 9.4 9.5 49.8
0.3 0.007 0.004 50.0
L2, L3 − Q
c2
1437
L2, L3 Q
33.4 15.5
17.9 34.5
48.7 50.0
L1 + Q ⇌ L3 + (δFe)
U3
1369
L1 Q L3 (δFe)
84.4 12.4 62.9 99.75
0.3 37.6 2.8 0.05
15.3 50.0 34.3 0.20
(δFe) + Q ⇌ L3 + γ
U4
1363
(δFe) Q L3 γ
99.75 12.6 62.2 99.87
0.05 37.4 2.9 0.05
0.20 50.0 34.9 0.08
L3 + (δFe) ⇌ L1 + γ
U5
1362
L3 (δFe) L1 γ
68.6 99.77 81.8 99.88
1.2 0.03 0.3 0.03
30.2 0.20 17.9 0.09
L1, L3 − γ
c3
1357
L1, L3 γ
76.31 99.89
0.33 0.02
23.36 0.09
L1 ⇌ (δMn) + γ + Q
E1
1233
L1 (δMn) γ Q
12.4 12.25 13.15 0.0005
87.5 87.69 86.80 49.9995
0.1 0.06 0.05 50.0
L3 ⇌ P + Q, Gas
D1/e6
1181
L3 P Q Gas
45.80 46.24 39.15 0.0000
4.22 3.76 10.85 0.0000
49.98 50.0 50.0 100.0
L3 + Q ⇌ γ + P
U6
1005
L3 Q γ P
55.4 37.7 99.98 46.7
0.8 12.3 0.00 3.3
43.8 50.0 0.02 50.0
γ+P⇌α+Q
U7
916
γ P α Q
99.99 47.0 99.96 36.5
0.0000 3.0 0.0000 13.5
0.01 50.0 0.04 50.0
(βMn) ⇌ (αMn) + γ, Q
D2
700
(βMn) (αMn) γ Q
31.9 31.0 39.4 0.0
68.1 69.0 60.6 50.0
0.000 0.000 0.000 50.0 (continued)
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12
Reaction or Critical Tie Line γ ⇌ α + (αMn), Q
Fe–Mn–S
Type
D3
Temperature [°C]
Phase
248
γ α (αMn) Q
Composition (at.%) Fe
Mn
56.4 96.7 34.1 0.0
43.6 3.3 65.9 50.0
S 0.000 0.000 0.000 50.0
Fig. 1. Fe-Mn-S. Boundary of the homogeneity range of liquid Fe at 1600°C
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Fe–Mn–S
13
Fig. 2. Fe-Mn-S. Quasibinary system FeS-MnS. Full lines calculated from the dataset of [1998Mie], dashed lines as assessed by [1988Rag]
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14
Fe–Mn–S
Fig. 3. Fe-Mn-S. Composition - pressure diagram of the quasibinary system MnS-FeS at 1000°C
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Landolt-Börnstein New Series IV/11D4
15
Fig. 4. Fe-Mn-S.
Reaction scheme, calculated from the dataset of [1998Mie]
Fe–Mn–S
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16
Fig. 5a. Fe-Mn-S.
Fe–Mn–S
Calculated liquidus surface of the partial system Fe-Fes-MnS-Mn with isotherms up to 2000°C
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MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
Fig. 5b. Fe-Mn-S.
Landolt-Börnstein New Series IV/11D4
17
Enlarged part of the liquid surface between Fe and FeS from 0 to 10 at.% Mn
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18
Fig. 5c. Fe-Mn-S.
Fe–Mn–S
Enlarged part of liquidus surface along the Fe-Mn edge
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
Fig. 5d. Fe-Mn-S.
Landolt-Börnstein New Series IV/11D4
19
Enlarged part of liquidus surface between 40 and 50 at.% S
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20
Fe–Mn–S
Fig. 6. Fe-Mn-S. Calculated liquidus surface of the Fe corner
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21
Fig. 7. Fe-Mn-S. Calculated isothermal section at 1600°C
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22
Fe–Mn–S
Fig. 8. Fe-Mn-S. Calculated isothermal section at 1500°C
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
23
Fig. 9. Fe-Mn-S. Calculated isothermal section at 1400°C
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24
Fe–Mn–S
Fig. 10. Fe-Mn-S. Calculated isothermal section at 1300°C
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–S
Fig. 11. Fe-Mn-S.
Landolt-Börnstein New Series IV/11D4
25
Calculated isothermal section at 1200°C
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26
Fig. 12. Fe-Mn-S.
Fe–Mn–S
Calculated isothermal section at 1100°C
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Fe–Mn–S
Fig. 13. Fe-Mn-S.
Landolt-Börnstein New Series IV/11D4
27
Calculated isothermal section at 1000°C
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28
Fig. 14. Fe-Mn-S.
Fe–Mn–S
Compositions of coexistent phases γ and Q at 1100, 1200 and 1300°C
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Fe–Mn–S
Fig. 15. Fe-Mn-S.
Landolt-Börnstein New Series IV/11D4
29
Isothermal section at 800°C
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30
Fig. 16. Fe-Mn-S.
Fe–Mn–S
Phase diagram of the sulfides P and Q in equilibrium with γ
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–S References [1913Roe]
[1928Shi]
[1933Vog]
[1934Mey]
[1937Vog]
[1937Wen] [1949Jae]
[1952She] [1955Tur1]
[1955Tur2]
[1960Mch]
[1963Kan]
[1964Cha] [1964Kul] [1966Kie]
[1966Sch]
[1967Bro] [1968Ska]
[1969Ban]
Landolt-Börnstein New Series IV/11D4
31
Roehl, G., “About Sulfide Inclusions in Iron and Steel: A Contribution to the Theory of Desulphurization of Irons” (in German), Stahl Eisen, 33(14), 565–567 (1913) (Experimental, Phase Diagram, Phase Relations, 3) Shibata, Z., “The Equibrium Diagram of the Iron Sulphide-Manganese Sulphide System”, Tech. Rep. Tohoku Imp. Univ., 7, 279–289 (1928) (Experimental, Phase Diagram, Phase Relations, 10) Vogel, R., Baur, H., “The Iron-Iron Sulphide-Manganese Sulfide-Manganese System” (in German), Arch. Eisenhuettenwes., 6(11), 495–500 (1933) (Experimental, Phase Relations, Phase Diagram, #, *, 6) Meyer, O., Schulte, F., “The Equilibrium FeS + Mn = MnS + Fe at High Temperatures” (in German), Arch. Eisenhuettenwes., 8, 187–195 (1934) (Experimental, Phase Relations, Phase Diagram, 21) Vogel, R., Hotop, W., “The Phase Diagram of the Iron-Iron Sulphide-Manganese SulhpideManganese System” (in German), Arch. Eisenhuettenwes., 11(1), 41–54 (1937) (Experimental, Phase Diagram, Phase Relations, #, *, 13) Wentrup, H., “Formation of Sulphide Inclusions in Steel” (in German), Techn. Mitt. Krupp, 5, 131–152 (1937) (Review, Morphology, 27) Jänecke, E., “S-Fe-Mn” (in German) in “Kurzgefasstes Handbuch Aller Legierungen”, Winter Verlag, Heidelberg, 722–724 (1949) (Review, Crys. Structure, Phase Diagram, Phase Relations, 5) Sherman, C.W., Chipman, J., “Activity of Sulphur in Liquid Iron and Steel”, J. Metals, 4, 597–602 (1952) (Experimental, Phase Relations, Thermodyn., 12) Turkdogan, E.T., Ignatowicz, S., Pearson, J., “The Solubility of Sulphur in Iron and IronMangannese Alloys”, J. Iron Steel Inst., 180, 349–354 (1955) (Experimental, Phase Relations, Thermodyn., 10) Turkdogan, E.T., Ignatowicz, S., Pearson, J., “Solid Solubility of Sulfur in Iron and IronManganese Alloys” (in French), Rev. Metall., 52(9), 725–730 (1955) (Experimental, Phase Relations, 7) Mchedlishvili, V.A., Lyubimova, A.M., Samarin, A.M., “The Manganese Function at the Elimination of the Unfavourable Sulfur Effect on the Steel Quality” (in Russian), Gos. Nauchno-Tekhn. Izd. Literatury po Chern. i Tsvet. Metall, Moscow, 56pp. (1960) (Experimental, Phase Relations, Morphology, 37) Kaneko, H., Nishizawa, T., Tamaki, K., “Study on Phase Diagrams of Sulphides in Steels” (in Japanese), Nippon Kinzoku Gakkai Shi, 27(7), 312–319 (1963) (Experimental, Phase Relations, 23) Chao, H.C., Van Vlack, L.H., Oberin, F., Thomassen, L., “Hardness of Inclusion Sulfides”, Trans. Am. Soc. Met., 57, 885–891 (1964) (Experimental, Phase Relations, Phase Diagram, 5) Kullerud, G., “Review and Evaluation of Recent Research on Geologically Significant SulfideType Systems”, Fortschritte Mineral., 41(2), 221–270 (1964) (Review, Phase Diagram, 109) Kiessling, R., Westman, C., “Sulphide Inclusions and Synthetic Sulphides of the (Mn, Me)S Type”, J. Iron Steel Inst., 204, 377–379 (1966) (Experimental, Phase Relations, Crys. Structure, Morphology, 19) Schindlerova, V., Buzek, Z., “Effect of Aluminium, Titanium, Manganese, Zirconium and Cerium on the Solubility and Activity of Sulfur in Molten Iron at 1600°C” (in Czech), Hutnicke Listy, 21(3), 169–175 (1966) (Experimental, Thermodyn, 14) Brown, J.R., “Sulphur Potentials in Iron and in Iron-Manganese Alloys”, J. Iron Steel Inst., 205, 154–157 (1967) (Calculation, Thermodyn., 22) Skala, J., Riman, R., “The Influence of Certain Elements on the Chemical Composition of Iron Sulphides” (in Czech), Sb. Ved. Pr. Vys. Sk. Banske Ostrave, Rada Hutn., 14(3), 115–122 (1968) (Experimental, Phase Relations, Morphology, 0) Ban-ya, Sh., Chipman, J., “Sulfur in Liquid Iron Alloys: II. Effects of Alloying Elements”, Trans. Met. Soc. AIME, 245(1), 133–143 (1969) (Experimental, Thermodyn., 18) MSIT®
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32 [1969Sal] [1969Sch]
[1970Nis]
[1971Bru]
[1971His]
[1971Kor]
[1971Nak]
[1971San]
[1971Ska]
[1971Ski]
[1971Tur]
[1972Kor]
[1973Bru]
[1973Buz]
[1973Das]
[1973Kir]
[1973Wyj]
Fe–Mn–S Salter, W.J.M., Pickering, F.B., “Composition and Constitution of Nonmetallic Inclusions in 1 % C-Cr Steels”, J. Iron Steel Inst., 207(7), 992–1002 (1969) (Experimental, Phase Relations, 26) Schindlerova, V., Buzek, Z., “The Effect of Aluminium, Titanium, Manganese, Zirconium and Cerium on the Solubility and Activity of Sulfur in Molten Iron at 1600°C” (in German), Freiberger Forschungshefte, 117B, 43–58 (1969) (Experimental, Phase Relations, Thermodyn., 14) Nishikawa, K., Ito, K., Sano, K., “The Effect of the Alloying Elements on the Solubility of Sulfur in δ -Iron” (in Japanese), Tetsu to Hagane, 56, 1467–1476 (1970) (Experimental, Thermodyn., 16) Brusq Le, H., Delmaire, J.-P., Marion, F., “About Electric Properties of the FeS, CoS and NiS Sulfides at High Temperatures. Comparison with the MnS Sulfide” (in French), C. R. Acad. Sci. Paris, C273(2), 139–142 (1971) (Experimental, Electr. Prop., 7) Hisada, H., Ito, K., Sano, K., “On the Solid State Reactions Between Mn-Fe Alloy and Synthetic Sulfide Inclusion”, Trans. Iron Steel Inst. Jpn., 11(2), 142–143 (1971) (Experimental, Phase Relations, Morphology, 0) Kor, G.J.W., Turkdogan, E.T., “Sulfides and Oxides in Fe-Mn Alloys: Part II. Precipitation of Solid Sulfides and Liquid Oxysulfides During Oxidation of Iron Alloys”, Metall. Trans., 2(6), 1571–1578 (1971) (Experimental, Phase Relations, Morphology, 18) Nakao, H., Kirkaldy, J.S., “Thermodynamics and Kinetic Reactions in the γ-Fe-FeS-MnS System at 1300°C” (in Russian), in “Fiz. Khim. Osn. Proizvod. Stali”, Moscow, Nauka, 32–40 (1971) (Experimental, Phase Relations, 23) Sano, N., Iwata, M., Hosoda, H., Matsushita Y., “Phase Relations of Fe-Mn-S System at 1330 and 1615°C”, (in Japanese), Tetsu to Hagane, 57, 1984–1989 (1971) (Experimental, Phase Diagram, 7) Skala, I., Riman, R., “The Elements Influence on the Composition of the Sulfide Inclusions in Iron” (in Russian) in “Fiz.-Khim. Osnovy Proizvodstva Stali”, Moscow, Nauka, 216–220 (1971) (Experimental, Phase Relations, Morphology, 0) Skinner, B.J., Luce, F.D., “Solid Solutions of the Type (Ca, Mg, Mn, Fe)S and their Use as Geothermometers for the Enstatite Chondrites”, Am. Mineral., 56(7–8), 1269–1296 (1971) (Experimental, Phase Relations, 19) Turkdogan, E.T., Kor, G.J.W., “Sulphides and Oxides in Fe-Mn Alloys. Part I. Phase Relations in the Fe-Mn-S-O System”, Metall. Trans., 2(6), 1561–1570 (1971) (Review, Experimental, Phase Relations, 10) Kor, G.J.W., Turkdogan, E.T., “Sulphides and Oxides in Fe-Mn Alloys. Part III. Formation of Oxysulphides During Freezing of Steel”, Metall. Trans., 3(5), 1269–1278 (1972) (Calculation, Phase Relations, 19) Brusq Le, H., Delmaire, J. P., “Determination of the Regions of Stability of the Cubic and Hexagonal Solid Solutions (Mn,Fe)S as a Function of Sulfur Pressure and Temperature” (in French), Compt. Rend. Acad. Sci. Paris, Ser. C, 276(7), 603–605 (1973) (Experimental, Phase Relations, 6) Buzek, Z., “Effect of Alloying Elements on the Solubility and Activity of Oxygen and Sulphur in Liquid Iron at 1600°C”, Int. Symp. Metall. Chem. - Appl. Ferrous Metall., Sheffield, July 1971, Iron and Steel Inst, London, 173–177 (1973) (Experimental, Review, Crys. Structure, *, 8) Dashevsky, V.Ya., Kashin, V.I., “Solubility and Activity of Sulfur in Manganese and its Alloys” (in Russian), Izv. Akad. Nauk SSSR, Met., (5), 85–88 (1973) (Experimental, Thermodyn., Phase Relations, 2) Kirkaldy, J.S., Nakao, H., Clark, I.S.R., Smith, P.N., “An Investigation of the Phase Constitution of Fe-Mn-S in the Vicinity of 1300°C. Paper I”, Canad. Metall. Quart., 12(1), 45–51 (1973) (Experimental, Phase Diagram, Phase Relations, 14) Wyjadlowski, J., Boos, J.-Y., Poyet, P., Goux, C., “Influence of Manganese on the Nature of Sulphides and the Cold and Hot Brittleness of Iron-Sulfur-Manganese Alloys with High Purity” (in French), C. R. Acad. Sci. Paris, C, 276(5), 383–386 (1973) (Experimental, Interface Phenomena, Mechan. Prop., 4)
DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mn–S [1974Sig] [1975Big] [1975Ito1]
[1975Ito2] [1975Sch]
[1975Tur] [1976Efi]
[1976Hil1]
[1976Hil2]
[1976Man]
[1977Fis]
[1978Fis1]
[1978Fis2]
[1978Shi]
[1978Uhr]
[1980Bra] [1980Ito]
[1980Nis]
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Sigworth, G.K., Elliott, J.F., “The Thermodynamics of Liquid Dilute Iron Alloys”, Met. Sci., 8, 298–310 (1974) (Review, Thermodyn., 249) Bigelow, L.K. Flemings, M.C., “Sulphide Inclusions in Steel”, Metall. Trans. B, 6, 275–283 (1975) (Experimental, Phase Diagram, Phase Relations, 24) Ito, I., Kamazawa, T., Matsubara, I., “The Fe-Mn-S Phase Diagram. I. Three-Phase Regions” (in Japanese), Tetsu to Hagane, 61, 503 (1975) (Experimental, Phase Relations, Phase Diagram, 0) Ito, I., Kamazawa, T., “The Fe-Mn-S Phase Diagram. II. The eutectic reaction L⇌ Fe + FeS + Q” (in Japanese), Tetsu to Hagane, 61, 504 (1975) (Experimental, Phase Diagram, Phase Relations, 0) Schuermann, E., Stroesser, H.J., “A Study of the Melting Equilibria in the Partial Diagram Fe-FeS-Mn-MnS” (in German), Arch. Eisenhuettenwes., 46, 761–766 (1975) (Experimental, Phase Diagram, Phase Relations, 20) Turkdogan, E.T., “Discussion of “Sulphide Inclusions in Steel” ”, Metall. Trans. B, 6, 663–664 (1975) (Experimental, Phase Diagram, Phase Relations, Morphology, 3) Efimov, V.A., Nakonechnyi, N.F., Lesnikova, E.V., Ishchuk, N.Ya., “The Chemical Composition of Sulphides in the Fe-Mn-S System” (in Russian), Izv. Akad. Nauk SSSR, Met., (1), 201–205 (1976) (Calculation, Phase Relations, 5) Hillert, M., Staffansson, L.-I., “On the Interaction between Three Liquid Miscibility Gaps in a Reciprocal System”, Acta Metall., 24, 1079–1082 (1976) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 5) Hillert, M., Staffansson, L.-I., “A Thermodynamic Analysis of the Phase Equilibria in the Fe-Mn-S System”, Metall. Trans. B, 7, 203–211 (1976) (Theory, Phase Relations, Phase Diagram, Thermodyn., 23) Mann, G.S., Van Vlack, L.H., “FeS-MnS Phase Relationships in the Presence of Excess Iron”, Metall. Trans. B, 7, 469–475 (1976) (Experimental, Phase Relations, Phase Diagram, 23) Fischer, M. Schwerdtfeger, K., “Thermodynamics of the System Fe-Mn-S. I. Activities in Iron Sulfide-Manganese Sulfide Solid Solutions in the Temperature Range 1100 to 1400°C”, Metall. Trans., 8, 467–470 (1977) (Experimental, Thermoyn., Phase Relations, 14) Fischer, M., Schwerdtfeger, K., “Thermodynamics of the System Fe-Mn-S: Part II. Solubility of Sulfur and Manganese in γ -Iron Coexisting with Sulfide in the Temperature Range 1100 to 1300°C”, Metall. Trans. B, 9, 631–634 (1978) (Experimental, Phase Relations, 12) Fischer, M., Schwerdtfeger, K., “Thermodynamics of the System Fe-Mn-S: Part III. Equilibria Inlvolving Solid and Liquid Phases in the Temperature Range 1100 to 1300°C”, Metall. Trans. B, 9, 635–641 (1978) (Experimental, Thermodyn., Phase Relations, Phase Diagram, 15) Shibanova, L.N., Lepinskikh, B.M., Vostryakov, A.A., “Transport Characteristics of Melts FeS-MnS, FeS-MnO and FeS-Cr2O3” (in Russian), Izv. Akad. Nauk SSSR, Neorg. Mater., 14(6), 1036–1039 (1978) (Experimental, Transport Phenomena, 8) Uhrenius, B., “A Compendium of Ternary Iron-base Phase Diagrams”, Hardenability Concepts with Applications to Steel, (Eds.): Doane, D.V., Kirkaldy, J.S., Proc. Symp. Held Sheraton-Chicago Hotel, Oct. 24–26, 1977, Metall. Soc. AIME Heat Treat. Com./Amer. Soc Met. Activ. Phase Trans., 28–81 (1978) (Phase Diagram, Phase Relations, Thermodyn., Calculation, 53) Brandes, E.A., Flint, R.F., “Mn-Fe-S” in “Manganese Phase Diagrams”, Manganese Centre, Paris, France, 117–117 (1980) (Review, Crys. Structure, Phase Relations, Phase Diagram, 3) Ito, Y., Yonezawa, N., Matsubara, K., “The Composition of Eutectic Conjugation in the Fe-Mn-S System”, Trans. Iron Steel Inst. Jpn., 20, 19–25 (1980), translated from Tetsu to Hagane, 65(3), 1979, 391–398 (Experimental, Phase Relations, Phase Diagram, 21) Nishida, K., Narita, T., Tani, T., Sasaki, G., “High-Temperature Sulfidation of Fe-Mn Alloys”, Oxid. Met., 14(1), 65–83 (1980) (Experimental, Phase Diagram, Phase Relations, Interface Phenomena, 16)
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DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
34 [1981Nis]
[1982Kub] [1983Uto]
[1984McC]
[1987Pap]
[1988Rag]
[1990Sou] [1995Bou]
[1995Oik]
[1998Hua] [1998Mie]
[1998Val]
[2000Kim]
[2000Oht]
[2001Eri] [2004Liu]
[2004Rag] [2004Wit]
Fe–Mn–S Nishizawa, T., Hasebe, M., “Computer Calculation of Phase Diagrams of Iron Alloys” (in Japanese), Tetsu to Hagane, 67(14), 2086–2097 (1981) (Review, Phase Relations, Phase Diagram, Thermodyn., 110) Kubaschewski, O., “Iron-Sulfur” in “Iron - Binary Phase Diagrams”, Springer-Verlag, New York, 125–128 (1982) (Review, Phase Diagram, #, 20) Utochkin, Yu.I., Mendelev, V.A., Grigoryan, V.A., Vestfal, S.V., “Experimental Study of Lamination in Metallurgical System” (in Russian), Izv. Vyss. Uchebn. Zaved., Chern. Metall., (7), 5–8 (1983) (Experimental, Phase Relations, 5) McCammon, C.A., Jackson, I., Ringwood, A.E., Cashion, J.D., “The Binary Systems FeSMgS and FeS-MnS: Mössbauer Spectroscopy of the B1 Solid Solutions and High-Pressure Phase Equilibria”, Phys. Chem. Miner., 11(4), 182–193 (1984) (Experimental, Electronic Structure, Crys. Structure, Phase Relations, Phase Diagram, #, *, 39) Papaiacovou, P., Schmidt, H.P., Erhart, H., Grabke, H.J., “High Temperature Sulfidation of Fe-Mn Alloys and Mn in H2-H2S”, Werkst. Korros., 38(9), 498–506, (1987) (Experimental, Interface Phenomena, 20) Raghavan, V., “Fe-Mn-S (Iron-Maganese-Sulphur) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Met., Calcutta, 2, 154–173 (1988) (Review, Phase Relations, Phase Diagram, #, *, 47) Southwell, G., Young, D.J., “Sulfidation Behavior of a Binary Fe-Mn Alloy”, Oxid. Met., 34 (3–4), 161–172, (1990) (Experimental, Interface Phenomena, 16) Bouchard, D., Bale, C.W., “Simultaneous Optimization of Thermochemical Data for Liquid Iron Alloys Containing C, N, Ti, Si, Mn, S, and P”, Metall. Trans., B, 26, 467–484 (1995) (Calculation, Thermodyn., 85) Oikawa, K., Ohtani, H., Ishida, K., Nishizawa, T., “The Control of the Morphology of MnS Insclusions in Steel during Solidification”, ISIJ Int., 35(4), 402–408 (1995) (Experimental, Phase Relations, 20) Huang, W., “An Assessment of the Fe-Mn System”, Calphad, 13, 243–252 (1989) (Assessment, Calculation, Thermodyn., 35) Miettinen, J., Hallstedt, B., “Thermodynamic Assessment of the Fe-FeS-MnS-Mn System”, Calphad, 22(2), 257–273 (1998) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 29) Valeev, R.G., Tanklevskaya, N.M., Mikhaylov, G.G., “Phase Equilibria in the Fe-Mn-S System at the Temperatures of the δ-Iron Existence” (in Russian), Izv. Vyssh. Uchebn. Zaved. Chern. Metall., (3), 72–73 (1998) (Calculation, Phase Relations, Phase Diagram, Thermodyn., 4) Kim, E.C., “Crystallographic and Magnetic Properties of Iron Sulfides Doped with 3d Transition Metals”, J. Mater. Sci. Letter., 19, 693–694 (2000) (Experimental, Crys. Structure, Magn. Prop., 8) Ohtani, H., Oikawa, K., Ishida, K., “Optimization of the Fe-Rich Fe-Mn-S Ternary Phase Diagram”, High Temp. Mater. Proces., 19(3–4), 197–210, (2000) (Calculation, Phase Diagram, Phase Relations, Thermodyn., #, *, 28) Eriksson, L., Kalinowski, M.P., “Mn1–xFexS, x = 0.05, an Example of an Anti-Wurtzite Structure”, Acta Crystallogr., Sect. E, E57(11), i92–i93 (2001) (Experimental, Crys. Structure, 10) Liu, Z., Kobayashi, Y., Nagai, K., “Crystallography and Precipitation Kinetics of Copper Sulfide in Strip Casting Low Carbon Steel”, ISIJ Int., 44(9), 1560–1567 (2004) (Experimental, Crys. Structure, Morphology, 37) Raghavan, V., “Fe-Mn-S (Iron-Manganese-Sulfur)”, J. Phase Equilib. Diffus., 25(4), 371–372 (2004) (Review, Phase Diagram, Phase Relations, #, *, 11) Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diff., 25(4), 346–354 (2004) (Experimental, Phase Diagram, Calculation, Thermodyn., #, 34)
DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
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Fe–Mn–S [2005Nam]
[Mas2] [V-C2]
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35
Nam, H.D., Kim, E.C., Han, J.S., “Mössbauer Study of Iron Sulfides Doped with 3d-transition Metals”, Solid State Commun., 135(5), 327–329 (2005) (Experimental, Crys. Structure, Electronic Structure, Phase Relations, 8) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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DOI: 10.1007/978-3-540-78644-3_20 # Springer 2008
Fe–Mn–Si
1
Iron – Manganese – Silicon Nathalie Lebrun
Introduction Being driven by the potential in industrial application as shape memory material, the Fe-Mn-Si based alloys have attracted intensive studies and considerable achievements were made in past decades. For example, Fe-Mn-Si alloys have recently brought out a promising potential to be used as tighteners or pipe couplings at a large scale [2005Wan]. Addition of Cr or Ni to the Fe-Mn-Si alloys improves their Shape Memory Effect (SME) and corrosion resistance [2005Lin]. A martensite transformation can occur if the alloy is cooled down to below the Ms (martensite start) temperature of the martensitic transformation (γFe) → (εFe), or, if a stress is applied even at temperature above Ms. On the other hand, a reverse martensitic transformation (εFe) → (γFe) can take place if the alloy is heated above the As (austenite start) temperature and presents non-thermoelastic property. The iron disilicide αFeSi2 is also a promising material for photovoltaics, optoelectronics and thermoelectrics. It is a semiconductor that presents a p-type conduction with different doping additives as for example Mn [2005Jia]. Fe-Mn-Si has superior features such as good formability, excellent workability, low electrical resistivity, chemical stability, high oxidation resistance, low cost of raw material and nontoxicity. It is also a thermoelectric good candidate for use up to 927°C. But compared with the traditional Ni-Ti and Cu based shape memory alloys, the recoverable strain and recovery stress of Fe-Mn-Si alloys are quite low which limit their practical applications [2000Dai]. That is why several works on thermomechanical treatment, change of composition, doping process were done in order to improve the shape memory properties of Fe-Mn-Si alloys. A comprehensive review on SME regarding the Fe-Mn-Si system is done by [2000Hsu]. [1988Ray] and [1994Rag] present a complete review of the Fe-Mn-Si ternary system. All experimental and thermodynamic optimization works are reported in Table 1. Binary Systems Using new experimental thermodynamic data, the Fe-Mn system has been recently reassessed and the Calphad description was updated by [2004Wit]. This new reevaluation of the phase equilibria leads to consistently better fits to the available experimental data. There is however, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have been written as L ⇌ (δMn) + (γMn,γFe). Consequently, the Fe-Mn system is accepted from [2004Wit]. After [Mas2], the liquidus and solidus phase boundaries in the Fe rich portion of the Fe-Si phase diagram was measured by [2005Mec] and thermodynamically reassessed by [1998Mie]. The DTA measurements done by [2005Mec] reveal a significant shift compared to [Mas2] in the solidus and liquidus curves in the range between 20 and 25 at.% Si, suggesting that the A2-B2 (and perhaps B2-D03) ordering transition occurs in large part through the two-phase temperature regions. Since these careful experimental data are identical to those used by [1998Mie] for its thermodynamic assessment, these data are retained in this report. Nevertheless, a more sophisticated thermodynamic treatment of the ordering transition is necessary for a better accuracy of the Fe-Si diagram in the range from 20 to 35 at.% Si. No new phase diagram has been available in the literature regarding the Mn-Si binary system and it has been accepted from [Mas2]. Solid Phases All the data on unary, binary and ternary phases are indicated in Table 2. Several ternary compounds have been reported in the literature. [1937Vog] and [1973Ish] detected respectively the Mn5Fe12Si11 and Mn3Fe5Si2 compounds which are not retained in this assessment since later works did not confirm their existence. Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
2
Fe–Mn–Si
[1964Gup] found a ternary phase Mn77Fe4Si19, called τ, at 1000°C with a non stoichiometric composition. The crystal structure of τ has been precisely characterized by [1977Sho] and the corresponding parameters and others characteristics are reported in Table 2. The unit cell is presented as a superstructure containing 2 · 110 atoms with the c axis repeat equal to (873.4 ± 0.1) · 2 pm. A complete solid solution is observed between FeSi and MnSi [1956Gla1, 1957Tav, 1958Aro, 1962Gla, 1968Set], Mn5Si3 and Fe5Si3 [1956Gla1, 1958Aro, 1962Gla, 1968Set] and α2Fe3Si and αMn3Si [1956Gla2, 1962Gla]. The lattice parameters of the cubic compound Mn1–xFexSi (0 ≤ x ≤ 1) vary linearly [1968Sid, 1958Aro, 1996Spo]. Using density-functional theory, [2003Luo] noticed a transition from an insulator for x = 1 to a metallic state with long range magnetic order at the Mn rich side. The lattice parameters of the compound (Mn1–xFex)5Si3 (0 ≤ x ≤ 1) first display an approximately linear increase as Mn is initially substituted by Fe [1958Aro, 1972Joh, 1973Bin, 1976Sud] with ∼2.2 pm/Mn atom for a and 1.7 pm/Mn atom for c at room temperature [2004Can]. In the Fe rich region, [1973Bin] reported a helical spin structure for Mn2Fe3Si3 with a spiral period of 3000 pm along the c axis and a collinear c axis ferromagnetic alignment for MnFe4Si3. This spiral structure also exists for the Mn1.5Fe3.5Si3 below –213°C with an estimated turn angle of 24° between successive planes along the axis of the spiral [2004Can]. In the (Mn1–xFex)5Si3 series, Mn and Fe atoms occupy preferably the 6(g) (FeII or MnII) and/or the 4(d) (FeI or MnI) sites respectively with small amount of disorder [1972Joh, 1973Bin]. In the MnFe4Si3 phase, splitting is observed regarding some magnetic peaks. [1970Nar] concluded that essentially all the first iron going to be substituted into Mn5Si3 go preferentially into the single crystallographic 4(d) sites with a slightly distorted octahedral array of silicon. For all the others compounds Mn3Fe2Si3, Mn2Fe3Si3 and MnFe4Si3, the 4(d) and 6(g) sites are occupied with an increasing occupation probability of the 6(g) with increasing Fe content [1970Nar]. The analysis of spin echo spectra of Mn in Fe3Si shows that Mn atoms prefer the Fe site (FeI) with eight first-neighbor Fe atoms [1974Bab, 1974Bur]. This result was recently confirmed by [2002AlN] using a theoretical Mössbauer study of the order-disorder phenomena in the (Mn1–xFex)3Si (x = 0, 0.967, 0.917, 0.833) leading to a split into a number of sextets [1975Pic, 1977Ver, 2002AlN] separated by about 40-45 kOe (3.2 to 3.6 MA·m–1) of the low magnetic field component [2002AlN]. As Mn concentration increases, the probability of the existence of Fe atoms decreases indicating that the Mn atoms are substituted on the Fe sites [1977Ver]. The corresponding structure is a L12 structure [1974Yoo, 1977Yoo] instead of D03 structure found by [1956Gla1, 1956Gla2]. Higher Mn doping induces additional atomic disorder in the form of a partial occupation of FeII sites (surrounded by four Fe atoms) and SiII (surrounded by four Si atoms) sites. This ordering effect leads to a break on the concentration curve of the lattice parameter when half of the atoms is substituted in sublattice FeII and SiII by Mn atoms [1975Ver, 1977Yoo]. The solubility limit for 3d transition metal is found to be x = 0 in (Mn1–xFex)3Si alloys [1977Nic]. In the (Mn1–xFex)3Si series, the MnFe2Si compound was particularly investigated. [1956Gla1, 1956Gla2] first measured the lattice parameter of this intermetallic compound and found a D03 type structure. [1976Zie] confirmed the value of the lattice parameter and concluded that the structure is L12 type together with some preferential D03 Fe-Mn disorder. The solubility ranges of the other binary phases in to the ternary system are rather limited. The (βMn) phase exhibits a large homogeneity range with a maximum estimated to be 43.9 at.% Fe and 16.5 at.% Si at 1000°C [1966Bar]. The homogeneity range of (αFe) is found to be maximum at 10-12 at.% Mn and 10 at.% Si at 1000°C [1968Set]. [1962Wit] and [1962Gla] suggest that the βFeSi2 phase exhibits a solubility range up to 25 mol% MnSi2 (i.e. 8.3 at.% Mn, 25 at.% Fe and 66.7 at.% Si) and 10 mol% MnSi2 (i.e. 3.33 at.% Mn, 30 at.% Fe and 66.67 at.% Si), respectively. [1962Wit] measured the lattice parameters going from 268.5 to 269.0 pm for a and 512.3 to 511.8 pm for c. Unfortunately no temperature is indicated. The lattice parameters of the binary βFeSi2 phase are slightly lower than that accepted in this assessment and reported in Table 2. Finally, the maximum of the solubility of this phase inside the ternary is found at about 1120°C at 6.8 at.% Mn, 27.2 at.% Fe and 66.0 at.% Si [1974Abr]. The βFeSi2 structure presents two crystallographically non equivalent sites for Fe and Si (FeI, FeII, SiI, SiII). Using first principles pseudopotential calculations based on Generalized Gradient Approximation (GGA) density function theory, [2002Tan] calculated the total
DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Fe–Mn–Si
3
energy of substitution of Mn dopant at the FeI and FeII sites and concluded that the dopant Mn prefers the FeI site. The MnSi1.75–x presents a solubility range with an upper limit of 26.1 at.% Mn, 11.2 at.% Fe and 62.7 at.% Si at 1000°C [1968Fli]. This result confirms the assumption of [1962Wit] who also suggested the existence of this binary phase inside the ternary. [1962Gla] found a limit of 31.67 at.% Mn, 1.67 at.% Fe and 66.66 at.% Si with no indication of the temperature. Finally, the maximum of the solubility of this phase inside the ternary is found by [1974Abr] at about 1120°C and 26.9 at.% Mn, 9.0 at.% Fe and 64.1 at.% Si. The lattice parameters vary from 552.5 to 548.3 pm for a and 6555 to 9517.2 pm for c at 1000°C [1962Wit]. The lattice parameter c of the binary MnSi1.75–x phase is higher than that accepted in this assessment and reported in Table 2. ν, Mn4Si and R, Mn6Si present smaller solubilities which extend at 1000°C up to 1.5 at.% Fe [1964Gup] and 5 at.% [1968Set] Fe, respectively. Mn5Si2, considered as a metastable phase in the Mn-Si binary system, could be stabilized by Fe substitution. The (Mn,Fe)5Si2 is absent for the Fe content higher than 34 mass% at 750°C [1964Sen]. [1999Guo] presented the description of crystallographic approach of (γFe) → (εFe) martensitic transformation in Fe-Mn-Si based alloys. A set of crystallographic formulas was deduced by introducing shuffle and shear probability into Wechsler-Lieberman-Read theory. It was concluded that Bain distortion does not occur and the normal vector of simple shear corresponds to {111} of the (γFe) structure. The amplitude of the shuffle direction on (0001)hcp plane is very small and within the unit cell. Quasibinary Systems Three quasibinary sections were proposed by [1937Vog]: FeSi - MnSi, Mn5Fe12Si11 - Mn5Si3 and FeSi Mn5Fe12Si11. A continuous solid solution is observed between FeSi and MnSi [1937Vog, 1956Gla1, 1957Tav]. A minimum of the liquidus curve is measured at 1260°C and 50 mass% MnSi, i.e. 50.00 mass% Mn, 16.26 mass% Fe and 33.74 mass% Si (37.88 at.% Mn, 12.12 at.% Fe and 50.00 at.% Si). A solid-solid transformation detected by [1937Vog] between 1150 and 1230°C may be explained by the crystallization of the Mn5Si3 phase involved in a eutectic reaction L ⇌ Mn5Si3 + MnSi reported in the binary Mn-Si system. This is certainly due to a shift of the composition of the alloys. Consequently this solid-solid transformation is not retained. The MnSi-FeSi quasibinary section is shown in Fig. 1. The quasibinary sections Mn5Fe12Si11-Mn5Si3 and FeSi-Mn5Fe12Si11 are not retained in the present evaluation since the existence of the ternary compound Mn5Fe12Si11 is doubtful. Invariant Equilibria Four ternary reactions of the U type were determined by [1937Vog]: L + (γMn) ⇌ (βMn) + (γFe) at 1120°C, L + (βMn) ⇌ Mn3Si + (γFe) at 1020°C, L + Mn3Si ⇌ Mn5Si3 + (γFe) at 1010°C and L + FeSi ⇌ (αFe) + (γFe) at 1160°C. These ternary reactions measured by [1937Vog] are not in agreement regarding the phases involved in various binary reactions. Later, [1938Jae] reviewed the ternary system based on the experimental data of [1937Vog]. The ternary reaction found at 1010°C by [1937Vog] becomes a eutectic ternary reaction in the review of [1938Jae]: L ⇌ (γFe) + Mn3Si + Mn5Si3. This leads to a new invariant reaction at 1060°C [1949Jae]: L + (δFe) ⇌ (γFe) + Mn5Si3. The reaction measured at 1160°C then becomes in the work of [1938Jae]: L + FeSi ⇌ (δFe) + Mn5Si3 with a reaction temperature at 1030°C [1949Jae]. [1938Jae] deduced these new ternary reactions from more recent binaries than those used by [1937Vog] but they are not in agreement with the binary phase diagrams accepted in this assessment. Later [1988Ray] assessed the ternary system and proposed other reactions in good agreement with the accepted binary systems. However, [1988Ray] suggested a large existence of the Mn9Si2 which is in fact the ν (Mn4Si) compound. A smaller primary crystallization field has been suggested in this assessment based on experimental data. The ternary reactions of [1988Ray] involving this compound have been replaced and a new ternary reaction L + ν + (βMn) ⇌ (Mn,Fe)3Si is proposed. On the other hand, the ternary reactions proposed by [1988Ray] regarding the βFeSi2 and the MnSi1.75–x could not been retained since they are not based on experimental data. Compositions and temperatures of all the ternary reactions are indicated in
Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Fe–Mn–Si
Table 3. Compositions data of the ternary reactions U2, U3 and U5 have been taken from the experimental work of [1937Vog]. The ternary reaction L ⇌ (γFe) + Mn3Si + Mn5Si3 proposed by [1938Jae] could not be retained in this assessment since the (Mn,Fe)3Si phase has a continuous solid solution whatever the temperature. This reaction has been replaced by the existence of a minimum involving the (Mn,Fe)3Si and the (γMn,γFe) phases. The reaction scheme is reported in Fig. 2. Liquidus, Solidus and Solvus Surfaces A partial liquidus surface has been presented by [1937Vog] in the Mn-Fe-MnFe-FeSi portion of the ternary system. This liquidus presents four ternary U type reactions and two maxima L ⇌ (αFe) + (γFe) and L ⇌ MnSi + Mn5Si3. Large disagreement is observed with the accepted binary Fe-Si system. In [1938Jae] version, one ternary reaction is eliminated as well as the two maxima. [1938Jae] proposed a new liquidus surface based on the work of [1937Vog] and new version of binary systems which differ from those retained in this assessment. Nevertheless large discrepancies are also noticed with the binaries accepted in this assessment. Later [1988Ray] assessed the liquidus surface which has been widely modified to be in agreement with the binaries. Nevertheless, the homogeneity range of the ν phase suggested by [1988Ray] is in contradiction with the small homogeneity range of this phase observed inside the ternary at 1000°C, i.e. 60°C lower than the in congruent melting temperature of this phase. Consequently the region of crystallization for this compound is taken narrower as indicated in Fig. 3. The Si rich corner is not complete since the ternary reactions involving βFeSi2 and MnSi1.75–x are unknown (question marks on the liquidus surface drawing). The isothermal lines are drawn from the measured data of [1937Vog]. Slight modifications were done regarding these curves to be in agreement with the accepted binaries. The partial liquidus surface is depicted in Fig. 3. Isothermal Sections An experimental partial isothermal section at 400°C was constructed by [1937Vog] in the Mn-Fe-MnFeFeSi part of the ternary system. Inconsistencies are observed with the binary Mn-Si and Fe-Si systems. The ν (Mn4Si) and R (Mn6Si) phases are missing on the isothermal section. Later [1938Jae] proposed a new version of this isothermal section along the Fe-Si binary system without the ternary compound suggested by [1937Vog] and using more recent versions of the binary systems. A Fe3Si2 phase was mentioned by these authors. This phase is certainly the η (Fe5Si3) which does not exist at 400°C according to the accepted Fe-Si binary system. Consequently these two versions of the isothermal section at 400°C could not be retained in this assessment. The Fe rich region of the isothermal section at 700°C has been studied by [1973Ish] showing the (γFe) (αFe) tie lines in the iron corner with Mn and Si varying up to 30 at.%. Using a Central Atoms Model, [1995Tan] calculated the (γFe)/(αFe) phase boundaries in the Fe rich corner (1 to 5 at.% Si and 1 to 12 at.% Mn) showing a good agreement with the experimental data. These results are in agreement with those thermodynamically assessed by [1993For]. Slight change has been done taking into account the homogeneity range of the (Mn,Fe)3Si phase in agreement with the accepted binaries. Moreover the η (Fe5Si3) phase has not been taken into account to be in agreement with the accepted Fe-Si binary system. Consequently, a new three-phase equilibrium has been added involving the Mn5Si3, (Mn,Fe)Si and (Mn,Fe)3Si phases. The maximum limits of the homogeneity ranges of MnSi1.75–x and αFeSi2 are estimated from the measurement at 1000°C assuming that it does not really change with variation of temperature of 300°C. Results of this compilation are presented in Fig. 4. The uncertain phase boundaries are indicated as dashed lines. The partial isothermal section at 900°C in the Fe rich corner has been experimentally investigated with Mn and Si content up to 30 mass% (30.34 at.% Mn and 46.01 at.% Si). [1973Ish] measured the tie lines (αFe) (γFe) and (αFe) - (βMn). Later [1991Lan] investigated the same region and included the (Mn,Fe)3Si phase, also suggested by [1968Set] at 1000°C. Since this last result is in agreement with the binary system Fe-Si, the phase relations proposed by [1991Lan] are retained in this assessment, whereas the results of [1973Ish] are rejected. The data of [1991Lan] in the Fe rich corner are compiled with the calculated isotherm done by [1993For] in the Mn rich part and are shown in Fig. 5. Slight modification has been done regarding the homogeneity range of the (Mn,Fe)3Si, (αFe), (βMn), (γFe) and ν phases to be in agreement with the accepted binaries and are shown as dashed lines on the drawing. The maximum limit of the homogeneity DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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ranges of MnSi1.75–x and αFeSi2 are estimated from the isotherm at 1000°C. The uncertain phase boundaries are indicated as dashed lines. From the composition curves of the lattice parameters along the FeSi-MnSi and Fe5Si3-Mn5Si3 sections at 950°C, the composition of the two phases (Mn1–xFex)Si and (Mn1–xFex)5Si3 in equilibrium which each other was obtained by [1958Aro] and the tie lines in the two-phase region were estimated. [1964Gup] investigated the partial isothermal section at 1000°C in the range Mn30Fe70-Mn10Fe90 and up to 20 at.% Si. The homogeneity range of the ternary phase τ was measured. At the high Fe content, τ is found to be in equilibrium with the (βMn) and Mn3Si phases, whereas at lower Fe content an additional phase ν is found in equilibrium with τ. The ternary phase τ was not found by [1966Bar] who experimentally investigated the (βMn)/(Mn,Fe)3Si phase boundaries along the Mn-Fe binary system with increasing Si up to 30 at.%. Using the measured data of [1964Gup], [1966Bar] redrawn the isothermal section showing a large homogeneity range of the (βMn) phase. Later [1968Set] proposed a complete isothermal section where the solubility of MnSi1.75–x, first suggested by [1968Fli], was taken into account. However, only schematic phase relations have been proposed in the Fe rich corner involving (αFe), (γFe) and Fe3Si phases. More recently, [1973Ish] determined the tie lines between these phases in the Fe rich ternary region with Mn and Si being up to 30 at.%. Unfortunately, [1973Ish] did not consider the (Mn,Fe3)Si phase which exists along the binary edge as a consequence of an order-disorder transition. Consequently, these last results could not be retained in this assessment. Recently, [1993For] thermodynamically assessed the isothermal section using the experimental data of [1966Bar]. An agreement is observed between experimental and calculated data. However the results on the Fe rich corner are not retained here since the βFe2Si phase does not exist at 1000°C in the accepted binary Fe-Si system and the (Mn,Fe)3Si phase was considered as stoichiometric. Finally, the isothermal section suggested by [1968Set] in the Mn rich corner is retained taking into account the phase boundaries (βMn)/(Mn,Fe)3Si experimentally determined by [1966Bar]. Moreover, the homogeneity ranges of the ν phase and the ternary phase τ are added from the good quality experimental work done by [1964Gup]. In agreement with the Mn-Si binary system, the homogeneity range of the Mn3Si phase is slightly modified on the drawing compared to the one proposed by [1968Set]. The solubility ranges of the βFeSi2 and MnSi1.75–x are deduced from the vertical section measured by [1974Abr]. The compiled isotherm is reproduced in Fig. 6. Temperature – Composition Sections Six isopleths have been constructed by [1937Vog] in the Fe-Mn-MnSi-FeSi region. Large discrepancies are observed with the accepted binary systems. [1937Vog] reported the Fe3Si2 phase which is in fact the η phase which does not exist below 825°C. Moreover the βFe2Si, η, ν, R, and (δMn) phases were not considered. This leads to large discrepancies on the solid-solid transformations involving these phases. Consequently, the isopleths of [1937Vog] can not be retained in this assessment. Using the liquidus data of [1937Vog], [1993For] thermodynamically reassessed these isopleths. These calculated results are in agreement with the accepted isothermal section at 1000°C for the two isopleths at xMn/xFe = 7.76 and 2.14 with Si varying from 0 to 35 mass% (0 to about 51 at.%). On the contrary, large discrepancies are noticed between the accepted isothermal section at 1000°C and the isopleths calculated by [1993For] at xMn/xFe = 0.837, 0.439, 0.176 and 0.0873 and 2.14 with Si varying from 0 to 35 mass%. The large homogeneity range of the (Mn,Fe)Si3 phase was not considered in the calculation and the (αFe) and the (βMn) phases have larger homogeneity ranges in [1993For] compared to the experimental data of [1966Bar, 1968Set]. Consequently, only the two first isopleths of [1993For] are accepted in this assessment and are depicted in Figs. 7 and 8. A better agreement is observed between the experimental liquidus data of [1937Vog] and the calculation for the Mn rich part of the isopleths of [1993For], except the isopleths at xMn/xFe = 0.176 and 0.0873. The fit of the liquidus curve is less satisfactory at high Si content. The phase boundary (βMn)/(βMn) + (Mn,Fe) Si3 is changed in agreement with the measured values of [1966Bar] at 1000°C and the homogeneity range of (Mn,Fe)3Si phase is considered in agreement with the accepted binary Mn-Si phase diagram. [1938Jae] assessed the Fe-Mn-Si ternary system and suggested an isopleth along Fe5Si3-Mn5Si3 over the temperature range 1000–1400°C which is not quasibinary since the Fe5Si3 phase only crystallizes between 825 and 1060°C. A continuous solid solution has been suggested between Fe5Si3 and Mn5Si3 which is reliable since the two phases have the same crystal structure and close lattice parameter values. This feature has Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Fe–Mn–Si
been confirmed later in the review of [1956Gla1]. The ternary compound Mn5Fe12Si11 found by [1937Vog] was not taken into account and [1938Jae] suggested an equilibrium between liquid, FeSi and Mn5Si3 at the same composition. Unfortunately the isopleth proposed by [1938Jae] is not retained since large discrepancies (more than 50°C) are observed regarding the melting temperature of the L + (Mn,Fe)Si region at the Fe5Si3 composition compared to the one accepted in the binary Fe-Si. Moreover the variation of the composition along this section is not clear. Indeed the composition Mn5Fe12Si11 mentioned by [1938Jae] along the vertical section does not correspond to the Mn composition on the axis. From tie lines determinations in the temperature range 700–1000°C between (αFe), (γFe) and Fe3Si, [1973Ish] constructed the isopleth at xMn/xSi = 1 which is not in agreement with the accepted isothermal sections. Consequently this isopleth is rejected. The system MnSi1.75–x-FeSi2 is of eutectic type [1974Abr]. The eutectic corresponds to the composition 21.2 mass% Mn, 14.1 mass% Fe and 64.7 mass% Si (13.12 at.% Mn, 8.58 at.% Fe and 78.30 at.% Si) and occurs at 1125°C. Ambiguities arise from a more detailed interpretation of this isopleth. In the text, the results are discussed in terms of mol% FeSi2 whereas the axis composition of the diagram is expressed as mass% FeSi2. Large discrepancies are observed at the binary edge with the accepted Mn-Si and Fe-Si phase diagrams. In the diagram proposed by [1974Abr], the MnSi phase is absent at the MnSi1.72 composition and the αFeSi2 does not appear as a single phase. Finally, this vertical section is not retained in this assessment. Thermodynamics [1998Zai] used an ideal associated solution model for the description of the activities of the three components of the Fe-Mn-Si liquid alloys in the temperature range 1162–1536°C and the composition range xMn ≤ 60 at.%. The experimental activities can be described by this model if the ternary complex MnFeSi is considered. The calculated ln(γMn) values deduced from this model are in good agreement with the experimental data obtained in the same work but are higher than the values measured by [1978Gee] and only in qualitative agreement with the measured data of [1979Ari]. [1996Zai1] pointed out that a possible reason for these discrepancies could be found in the experimental method used by [1978Gee] where the gas carrier was not saturated with Mn vapor and the oxidation of Si to SiO with oxygen contained in the carrier gas. Consequently the calculation of [1997Li2] and [1993For] using the unified interaction parameter model (UIPM) and the substitutional solution model respectively, are not retained in this assessment since these works are based on the experimental data of [1978Gee]. The activity data measured by [1996Zai1, 1996Zai2, 1998Zai] seem to be more reliable and less scattered. Consequently, these more recent experimental data are retained in this assessment. The corresponding silicon, manganese and iron isoactivity lines are depicted in Figs. 9, 10 and 11, respectively. The isoactivity lines of silicon are close to the secants connecting the Fe-Si and Mn-Si binary sides. There is a tendency for silicon to decrease the activity coefficient of Mn as pointed out by [1987Bal, 1993For, 1995Bou]. The corresponding first-order interaction parameter in the liquid was measured to be εMn(Si)(liq) = 4.9 by [1987Bal] where εMn(Si)(liq) = ∂ln γMn/∂ xSi and γMn = (xMn in pure Fe)/(xMn in the solution). The first-order parameters of Mn and Si were measured in the (δFe) phase: εMn(Si)(δ) = εSi(Mn)(δ) = – 13±2 at 1527°C [1981Fuj]. [1984Bon] investigated experimentally the mixing enthalpy of liquid Fe-Mn-Si alloys (xFe/xSi = 8/2, 7/3, 4/6 and 3.5/6.5) at 1687°C showing large negative values due to strong attractive interactions Fe-Si and Mn-Si of the donor-acceptor type. Later [1988Sud, 1989Bon] completed by the measurement of the mixing enthalpy of liquid alloys at 1572°C along the Mn-FeSi and Fe-MnSi sections. [1980Bob] measured the mixing enthalpy in liquid alloys at 1527°C along the Mn5Si3-Fe5Si3 line showing a minimum enthalpy of –2.6 kJ·mol–1 at 45 mol% Mn5Si3. Some results are reported in Table 4. Using an associated-solution model which assumes the existence of some binary and one ternary complexes (MnFeSi) in the melts, the molar Gibbs energy and enthalpies of formation of Fe-Mn-Si liquid alloys have been calculated by [1996Zai1, 1996Zai2, 1998Zai]. The optimal value of the molar energy of formation for the ternary complex MnFeSi found by [1998Zai] is reported in Table 5. The enthalpy of formation data are in agreement with the values calculated by [1958Bon, 1960Bon] at 1327°C and [1984Bon] at 1687°C for three cross sections Fe70Si30 - Mn28.9Fe49.7Si21.4, Fe40Si60 - Mn33.7Fe26.5Si39.8 and Mn64.9Fe35.1 Mn29.4Fe24.8Si45.8. For the cross section Fe80Si20 - Mn25.5Fe59.6Si14.9 agreement is seen only for alloys DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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at high manganese content. [1988Sud] found higher heat effect of mixing than in [1996Zai1, 1996Zai2] and [1984Bon]. These discrepancies could be explained by the non-controlled oxidation of silicon in samples studied by [1988Sud]. Moreover, the results of [1958Bon] are in only qualitative agreement with the data of [1996Zai1, 1996Zai2] since the thermodynamic properties were predicted from the binary subsystems. In conclusion, the data of [1996Zai1, 1996Zai2] are retained in this assessment and their corresponding lines with the same Gibbs energy are shown in Fig. 12. Using a thermodynamic model introducing magnetic contributions from binaries, [1997Li1] evaluated the Gibbs energy of the (γFe) and (εFe) phases for the Fe rich alloys with 30 mass% Mn and 6 mass% Si (28.65 at.% Mn and 11.21 at.% Si). A decrease of these energies is observed in the temperature range 627°C to –73°C with G(γFe) < G(εFe). For the Fe alloy with 28.2 mass% Mn and 5.49 mass% Si (27.07 at.% Mn and 10.31 at.% Si), the measured critical driving force ΔGC for the reverse (α’Fe) → (γFe) transformation is estimated to be 68 J·mol–1 by [1997Li1] at As = 168°C (441 K), temperature calculated by [1997And]. This measured value of ΔGC is not in agreement with the calculated critical driving force of [1997Li1] for the Fe alloys with 30 mass% Mn and 6 mass% Si estimated to be ΔGC = 142 J·mol–1 at As = 157°C (430 K), temperature calculated by [1998Cot1]. A better agreement is observed between the experimental and calculated data of [1997And] using calorimetry technique and Thermocalc software with the thermodynamic description of [1993For]. Despite the fact that the surface of the alloy is probably oxidized, the measured values of [1997And] are preferred here. [1980Smi] measured the heat capacity for (Mn1–xFex)3Si (0.47 ≤ x ≤ 0.8) at temperature lower than 187°C. The magnetic contribution to the heat capacity was deduced and has been estimated to be 4.3 and 4.8 J·(mol ·K)–1 for x = 0.75 and 0.58, respectively. These values are quite reasonably in agreement with those calculated from a molecular field approximation by the same authors. The entropy associated with magnetic ordering SF and reordering SR were derived. For example, the values regarding the MnFe2Si compound is found to be SR = 3.2 J·(mol·K)–1 and SF = 3.4 J·(mol ·K)–1 with a uncertainty of 10–20 %. SF decreased significantly with the increase of the concentration x whereas SR is more or less constant. [1980Smi] concluded that the re-ordering transition at TR plays a more significant role in the establishment of the magnetic state than a simple realignment of the magnetic moment as suggested by [1977Yoo]. The total magnetic entropy Sm is considerably lower than the 3Rln2 = 17.3 J·(mol·K)–1 for the 3 N magnetic spins with S = ½ [1980Smi, 1988Mil]. Notes on Materials Properties and Applications There exist three deformation stages in the Fe-Mn-Si alloys when an increasing stress is applied: (a) elastic deformation of (γFe); (b) formation of a stress induced martensite; (c) plastic deformation of (γFe). The martensitic reaction (γFe) → (εFe) is classified as a semi-thermoelastic transformation since the dislocation can move reversibly upon heating and the thermal hysteresis is as high as 100°C. The reverse martensite transformation (εFe) → (γFe) can also be assisted by intensive strain and occurs at a temperature far below the normal As temperature [1996Zha]. Moreover the nucleation process of the (εFe) phase does not strongly depend on soft modes [1999Hsu]. The parent austenite (γFe) phase transforms into martensite phase (εFe) by the formation and overlap of a stacking fault mechanism with low stacking fault energy based on the reversible motion of Schockley partial dislocation [2000Jin1, 2000Hsu, 2001Wan1, 2001Wan2]. The martensite start temperature does not markedly vary with the grain size of the (γFe) phase and there does not show a considerable lowering of the elastic modulus during the (γFe) → (εFe) transformation [2000Hsu]. Tension-compression asymmetry was observed in Fe-Mn-Si alloys strained by bending. [2005Sta] concluded that formation and reversion of stress-induced martensite occurs more readily in tension than in compression. Moreover, irreversible plastic deformation is more pronounced in the compression region, contributing to reduced shape memory. Tensile stress should be better for shape memory than compressive stress. The shape memory effect (SME) is associated with this stress-induced martensitic transformation by a predeformation at a temperature between Ms and Md [2005Wan] with increase of Si content in the alloy. The SME could improve by a decrease of the critical stress for inducing martensite, an increase of the amount of martensite and a strengthening of austenite [2000Dai, 2000Hsu]. The shape recovery rate in Fe-Mn-Si alloys is proportional to the volume of reversible martensite before recovery which can be estimated as the area covering an internal friction peak. Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Fe–Mn–Si
The good SME strongly depends on the composition of the alloys. Fe-Mn-Si alloys, which contain 28–34 mass% Mn and 4–6.5 mass% Si (27.26-32.30 at.% Mn and 7.62-12.08 at.% Si), exhibit a nearly perfect SME [2005Wan]. [2000Jin1] concluded that stacking fault energy, related to the critical driving force ΔGC, increases with addition of Mn (low SME) and decreases with the addition of Si (high SME) linearly as followed ΔGC = 215.23 + 7.923 {mass% Mn} – 46.38 {mass% Si} (in J·mol–1) leading to a linear evolution of the martensite start temperature as follows: Ms = 284.7 – 7.857 {mass% Mn} + 46.0 {mass% Si} (in K). Finally, the stacking fault energy of the (γFe) phase can be expressed as 180.54 + 7.923 {mass% Mn} – 46.38 {mass% Si} (in J·mol–1) [2000Jin2]. These three equations are only available within the composition range 24 < Mn < 32 and 4 < Si < 6 (in mass%). The linear evolution of Ms suggested by [2000Jin1] is not in agreement with the findings of [1998Cot1, 1999Cot] which stipulated that small additions of Si (less than 5 mass%) increase both Ms and Md temperatures, but further Si additions start to progressively decrease them at sufficiently high Mn contents [1990Sad]. These last results are qualitatively in fairly agreement with those reported by [1993For]. However various differences do exist: the lines for equal Ms and As are essentially parallel in the work done by [1998Cot2], while some of the Ms lines calculated by [1993For] significantly approach each other at high Mn contents. These isolines calculated by [1998Cot2] would be preferred since a complete thermodynamic description of the martensitic transformation and its reverse transformation upon heating has been done considering the magnetic contribution which was not taken into account in the work done by [1993For]. Moreover, [1993For] considered the Resistance to Start Transformation Energy ΔGC (critical driving force) as independent of composition and equal to 50 J·mol–1. The experimental values of Ms and As measured by [1996Zha] for a Fe alloy with 29.44 mass% Mn and 6.2 mass% Si (28.07 at.% Mn and 11.56 at.% Si) and [1995Ron] for a Fe alloy with 29.9 mass% Mn and 6.0 mass% Si (28.56 at.% Mn and 11.21 at.% Si) are in general agreement with the calculated values of [1998Cot2] but 30°C lower than the calculated temperatures. Ms and As were estimated to be about 27°C and 127°C [1996Zha], and 34°C and 129°C [1995Ron] respectively instead of the respective calculated values 57°C and 157°C [1998Cot2]. The SME is also influenced by the machining of the materials. For example, Wire Electro-Discharge Machining (WEDM) surface of Fe alloy with 30 mass% Mn and 6 mass% Si present slight degradation of shape recovery due to the depression of the re-cast layer [2005Lin]. Deformation of the Fe alloy with 30.7 mass% Mn and 6.54 mass% Si (29.17 at.% Mn and 12.16 at.% Si) produced a complete SME in a wide temperature range (–269°C to 127°C) [1986Sat]. [1996Mat] compared the SME on hot rolled sheets for Fe alloy with 28 mass% Mn and 6 mass% Si (26.75 at.% Mn and 11.21 at.% Si) between mid-thickness and near surface layers. It was concluded that a maximum of shape recovery strain is obtained for specimen with a tensile axis parallel to the rolling direction (RD) or the transverse direction of the RD in the plane of the surface layer, i.e. axis preferentially oriented to <110>. The room temperature prestrained materials lead to lower SME due to the increase of the gap between the reverse starting and finishing martensite temperature As and Af. As the strain is increasing, the Af is greatly increased whereas the As keeps constant [1995Ron]. For example the shape recovery ratio decreases from 53 to 35% with increase strain from 2.4 to 6.2%. [1995Ron] suggested that a larger driving force ΔGC is required to reverse the stress-induced (εFe) martensite. Consequently the deformed specimen must be heated to higher temperature in order to obtain a better SME. The SME properties can also be improved by the repeated deformation and the number of heating cycles leading to an increase of the martensite start temperature. This training effect has been studied by [1993Wat] who found the major recovery of the SME at a temperature lower by 30°C after 13 deformation-heating cycles on Fe alloy with 31 mass% Mn and 6 mass% Si (29.60 at.% Mn and 11.21 at.% Si). The homogenization of the internal structure is improved during the cycles and is achieved by introduction of the (εFe) phase distributed in fine structure which favors the SME. The addition of Al, Nb or Ti at about 1 mass% favors the ferrite formation and significantly reduces the austenite stability resulting in duplex microstructure in austenite and ferrite. Austenite stabilizing additions of Ni or Mn are then necessary. Nevertheless, too many alloying additions are detrimental to the shape memory behavior. This decrease of the shape memory with addition of Ni and/or Mn is the results of a suppression in the temperature Ms [2007Sta]. On the contrary, carbon, nitrogen and rare earth elements were found to improve shape memory effect in Fe-Mn-Si alloys [2000Dai, 2000Hsu, 2005Wan]. The stacking fault energy of the (γFe) phase is then reduced and thus favors the SME.
DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Several works show that the transition antiferromagnetic → paramagnetic of the (γFe) phase characterized by the Neel temperature TN(γFe) strongly affected the SME and the martensitic transformation [2000Wu, 2000Hsu]. In fact, the amount of (εFe) is related to Ms and (Ms–TN) [1996Qin]. For example in the Fe alloy with 30.30 mass% Mn and 6.10 mass% Si (28.91 at.% Mn and 11.39 at.% Si), the TN(γFe), equal to 27°C [1999Jin, 2000Wu], is slightly below the martensitic temperature Ms (Ms – TN(γFe) = 30°C) and the martensitic transformation is then distinctly suppressed [2000Wu]. These values are slightly different from those reported by [2000Hsu] (TN(γFe) = 20°C) and [1999Che] (TN(γFe) = 6°C) with Ms – TN(γFe) = 50°C and a stacking fault probability Psf estimated to be 2.73·10–3 [2000Hsu]. [1995Ron] found TN(γFe) = 24°C with Ms – TN(γFe) = 10°C for a Fe alloy with 29.9 mass% Mn and 6.0 mass% Si. A systematic investigation of the martensite transformation temperature and their influence on the SME has been carried out by [1996Qin] on Fe alloy with 27.77–32.01 mass% Mn and 4.10–6.10 mass% Si (27.01–30.54 at.% Mn and 7.80–11.38 at.% Si). It was concluded that a decrease of the Ms– TN(γFe) quantity with increasing Mn and Si content leads to a higher quality of the SME: Ms– TN(γFe) = 52°C with a SME of 35% for Fe with 27.77 mass% Mn and 4.10 mass% Si to Ms– TN(γFe) = 23°C with a SME of 82% for Fe alloys with 32.01 mass% Mn and 6.10 mass% Si. Finally, SME increases rapidly when Ms–TN(γFe) becomes positive [1986Sat]. Upon further cooling below TN(γFe), the martensitic transition could still occur, but the amount of the (εFe) phase formed is rather small and leads to a decrease of the area of the internal friction peak. With (Ms–TN) being smaller, the martensite transformation is suppressed, the pre-existing (α’Fe) martensite phase becomes smaller and the SME is good [1996Qin]. Contrary to the assumptions done by [2000Hsu, 2000Wu], [1999Che] concluded that, even when Ms is lowered below TN, martensite can also be induced to a great extent from the paramagnetic or the antiferromagnetic parent phase at different temperatures by application of stress. The alloy composition evidently affects the TN(γFe) temperature while the Neel temperature for the (εFe) phase TN(εFe) is little affected [1999Jin]. Addition of Si to Fe-Mn alloys lowers TN(γFe) [1986Mur], increases the susceptibility [1996Qin, 1998Qin, 2000Wu] and leads to a higher and sharper peak of the magnetic susceptibility at TN [1996Qin]. [2000Wu] found that when Mn content is 30 mass% and Si content is in the range 0–6 mass% the TN(γFe) temperature reduces about 25°C with increasing 1 mass% Si. These results are not in agreement with the empirical equations of [1999Jin] deduced from thermodynamic consideration of the antiferromagnetic (γFe) → (α’Fe) for Fe with 24–32 mass% Mn and 0–6 mass% Si (24.30– 30.55 at.% Mn and 0–11.21 at.% Si). It was also pointed out that when Si content is about 6 mass% (11.21 at.%) and Mn content is in the range of 25–35 mass% (23.90–33.40 at.%), the increase of the TN(γFe) is about 7°C with increasing 1 mass% Mn content [2000Wu]. TN of Mn rich (MnxFe1–x)Si alloys (x > 0.9) decreases rapidly as x is decreased [1987Gel]. The addition of rare earth elements to Fe-Mn-Si also gives lower TN(γFe) values [1999Hsu, 2000Wu]. In conclusion, the occurrence of the transformation (γFe) → (α’Fe) in Fe-Mn-Si based alloys is determined by the chemical driving force, the magnetic structure of the parent phase (γFe) and the external stress applied [1999Che]. An increase of the electrical resistance is observed during the martensitic transformation in Fe alloy with 30 mass% Mn and 6 mass% Si (28.65 at.% Mn and 11.21 at.% Si) whereas a decrease is noticed during the reverse martensitic transformation [1996Zha]. These results are in agreement with those of [1996Qin] for Fe alloys ranging from 27.77 mass% Mn and 4.10 mass% Si (27.01 at.% Mn and 7.80 at.% Si) to 32.01 mass% Mn and 6.10 mass% Si (30.54 at.% Mn and 11.38 at.% Si). This behavior was also observed in [1993Wat] showing a higher resistivity for powder alloys (31 mass% Mn and 6 mass% Si or 29.60 at.% Mn and 11.21 at.% Si) after 13 deformation-heating cycles. The hardness is not homogenous in the wire electro-discharge machining (WEDM) surface of the Fe alloys with 30 mass% Mn and 6 mass% Si (28.65 at.% Mn and 11.21 at.% Si). The hardness near the outer surface can reach 550 Vickers (5396 MPa), whereas for depth more than 50 μm the value is about 225 Vickers (2207 MPa) [2005Lin]. On bulk samples, the mechanical strength is improved after annealing for iron alloys containing about 20–30 mass% Mn (18.45–27.63 at.% Mn) and up to 10 mass% Si (18.02 at.% Si). The best mechanical performance was optimum for a Fe alloy with 25 mass% Mn and 2 mass% Si (24.82 at.% Mn and 3.88 at.% Si) after a plastic deformation [1976Zvi]. After deformation at 1000°C by rolling of 20%, this alloy gives a yield point of 85 kg·mm–2 (833.6 MPa) and an ultimate strength of 105 kg·mm–2 (1029.7 MPa). Landolt-Börnstein New Series IV/11D4
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Fe–Mn–Si
The magnetic phase diagram and magnetocaloric properties of the system Fe-Mn-Si have been investigated [2004Can, 2002Son]. MnFe2Si: This cubic compound has a saturation magnetic moment of 2.33 μB [1983Bus]. MnFe4Si3: This compound orders ferromagnetically [1970Nar] just below 27°C and the moments are aligned at approximately 40° to the c axis with higher values (1.4 μB and 2.5 μB for the 4d and 6g sites) than the site moments for Fe5Si3 [2004Can]. From the fact that the total magnetic moment of this compound is higher than in Fe5Si3, [2002Son] concluded that the Mn moments are coupled parallel to the Fe moments. A considerable magnetic-entropy change is observed around room temperature, the maximum value being 4 J·kg–1·K–1 for a field change of 5 T. This relatively high value has been attributed by [2002Son] to the large value of the Mn moment when forced into a ferromagnetic arrangement. Mn1.5Fe3.5Si3: [2004Can] noticed that its magnetic structure is similar to that of MnFe4Si3 at –123°C with lower moments for 4d and 6g sites (0.9 μB and 1.9 μB). The Curie temperature is –8°C. The transition from the spiral to ferromagnetic phase is accompanied by a small increase in the cell volume suggesting that magnetoelastic effects may play a role. Mn2Fe3Si3: Magnetic order occurs below –196°C with however a higher Curie point (–23°C) than the previous components and the presence of a low temperature transition at –223°C [1973Bin]. Both of these findings are in contrast with neutron diffraction studies done by [2004Can]. This more recent investigation leads to the conclusion that the magnetic structure of Mn2Fe3Si3 seems to be an incommensurate antiferromagnet since a continuous increase of the intensity of the magnetic peak is observed between –265°C and –123°C. The magnetic entropy changes are positive, but the change in magnetic order with changing magnetic field (0 to 5 T) is rather low [2002Son]. Mn3Fe2Si3: Two transitions have been observed with the first corresponding to a paramagnetic to antiferromagnetic ordering at –148°C followed by a successive transition below –228°C [2002Son, 2004Can]. Between these two transitions, the antiferromagnetic phase [1970Nar] has a hexagonal cell [1973Bin]. The high temperature phase corresponds to a commensurate antiferromagnet with an ordering wave vector {1/2, 0, 0}. The lower transition is probably based on a spin reorientation of the moments from the hexagonal axis to a low temperature planar configuration [2002Son]. The Curie temperature is located at 57°C [1973Bin]. At –196°C the saturation moment was measured to be 0.3 μB [1973Bin]. Mn4FeSi3: At temperatures between –203°C and –178°C, its magnetic structure is of the same type as that of the low temperature phase of Mn3Fe2Si3 below –223°C with a saturation moment at –196°C to be 1.4 μB [1973Bin]. No field-induced magnetic phase transition has been observed [2002Son]. The Curie temperature of this antiferromagnetic phase [1970Nar] is located at 53°C [1973Bin]. The atomic magnetic moments of the (Mn1–xFex)Si (0 ≤ x ≤ 1) were calculated by [2003Luo] varying from –0.13μB (Mn0.5Fe0.5Si) to 0.55μB (Mn0.75Fe0.25Si) and from 0μB (Mn0.125Fe0.875Si) to 1.44μB (Mn0.75Fe0.25Si) for the Fe and Mn ions respectively. The calculated magnetic moment of the Mn ion is significantly less than 3μB expected for the Mn4+ configuration while Fe ions have a positive or a negative magnetic moment. Moreover it was found that for x < 0.875, the alloys start to show ferromagnetic order. These results are in agreement with the experiments of [1996Spo]. Indeed, [1996Spo] measured the magnetic transition temperature at about –244°C for MnSi and observed its decreasing down to –251°C if 2 at.% of Mn are substituted by Fe. The effect of Mn substitution on magnetic and thermal properties of (MnxFe1–x)3Si have been studied by [1998Muk] (0.4 ≤ x ≤ 0.6), [1977Yoo] (0 ≤ x ≤ 0.75) and [1976Sid] (0 ≤ x ≤ 0.67). Addition of Mn leads to an increase of the Debye temperature and an increase of the electron-magnetic contribution to the thermal expansion and the entropy [1988Mil]. On the contrary, the dilution of Fe atoms in Fe3Si by Mn atoms reduces the Curie temperature [1974Yoo, 1976Zie, 1988Mil]. Otherwise the magnon contribution to thermal expansion decreases with addition of Mn. [1976Lut, 1976Sid] demonstrated that series of transitions exist in these alloys leading to: 1) a ferromagnetic solution for alloys from Fe3Si to Mn0.99Fe2.01Si; 2) a ferromagnetic solution for alloys from Mn0.99Fe2.01Si to Mn2.01Fe0.99Si; 3) an antiferromagnetic solution for alloys from Mn2.01Fe0.99Si to Mn3Si. In the ferromagnetic solution, positive exchange interactions of Fe-Mn and Fe-Fe exist. At high Mn content the antiferromagnetic order is established because of the strong negative interaction of Mn-Mn type. In the intermediate solutions, a complex ferromagnetic structure is observed with a competition between the positive and negative interactions as described above. Various limits of this complex intermediate state were suggested Mn24.75Fe50.25Si25.00-Mn50.25Fe24.75Si25.00 [1976Sid], DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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11
Mn15Fe60Si25-Mn43.75Fe31.25Si25.00 [1998Muk], Mn18.75Fe56.25Si25.00-Mn43.75Fe31.25Si25.00 [1977Yoo] and Mn15Fe60Si25.00-Mn50.25Fe24.75Si25.00 [1976Lut]. The magnetic reordering temperature TR is very sensitive to the sample condition (ageing, heat treatment, magnetic field) [1991Mil]. [1976Nic] showed that the Mn atoms which substituted the Fe atoms on the FeI sites have a magnetic moment which increased from 0 to 2.2 μB. This effect produced a decrease in the average moment of the other Fe sites from 1.35 μB at x = 0.87 to 0 μB at x = 0.75. [1991Mil] found that a long range ordering can be suppressed in Mn0.8Fe2.2Si compound in applied magnetic field of about 0.4 T. Moreover low fields (1 to 2 mT) strongly influence the bulk magnetization. The (MnxFe1–x)5Si3 phase also presents a magnetic transition. Two transformations occur at low temperature characterized by the rearrangement of the antiferromagnetic ordering at TR and the transition to a paramagnetic state at TN [1976Sud]. A rapid decrease of the Curie temperature TC is observed for the Fe5Si3 rich alloys. [1976Sud] noticed that replacement of Mn atoms by Fe atoms weakly affects the temperatures TR for x > 0 [1977Yoo, 1980Smi] and TN. Agreement had been observed regarding the measured data of TR from [1977Yoo] and [1980Smi]. From magnetic hyperfine fields, [1972Joh] suggested that the relative magnitudes of the moments for FeI and FeII sites are the same for the entire series. Changes on magnetic properties of Fe alloy with 6.5 mass% Si and Mn addition less than 0.20 mass% (6.59 at.% Mn and 0.4 at.% Si), are fairly small [1976Nar, 1978Nar], except in the case of the maximum permeability (at 0.15 mass% Mn and annealing at 900°C for 20 h). The non-doped semiconducting material αFeSi2 usually presents p-type conduction over a wider range of Si/Fe ratios, but n-type conduction has sometimes been reported in the stoichiometric or Si excess conditions. [2005Jia] observed an evolution of the conduction to n-type during the eutectoid decomposition of a Mn-added αFe2Si5 alloy. This change is caused by the abundant Fe vacancies acting as donors in the supersaturated αFeSi2(Si) solution, which is a metastable state of the αFeSi2 phase which appeared in the eutectoid reaction βFeSi2 ⇌ αFeSi2 + Si and can be eliminated by further annealing. A small amount of excess silicon in αFeSi2 based thermoelectric materials Mn0.08Fe0.92Six (1.9 ≤ x ≤ 2.5) increases the thermoelectric sensitivity with a maximum Seebeck coefficient of 315 μV·K–1 for x = 2.0 at 150°C [2003Che]. However a large amount of excess silicon deteriorates the thermoelectric properties due to the high thermal conductivity of silicon. Compared to the undoped αFeSi2, the electrical conductivity of p-type Mn0.08Fe0.92Six (1.9 ≤ x ≤ 2.5) is greatly enhanced. This electrical conductivity increases with increasing pressureless sintering temperature, i.e. density, whereas the Seebeck coefficient becomes lower. The optimum thermoelectric properties of the Mn0.1Fe0.9Si2 alloy were obtained for a sintering temperature at about 400–600°C. The best Seebeck coefficient S of 380 μV·K–1 at 400–600°C was obtained for coarse powders Mn0.1Fe0.9Si2 sintered at 1150°C for 2 h and annealed at 800°C [1999Cho, 2001Cho]. The highest power factor P = S2 / ρ of 1.7·10–3 W·mK–2 was obtained at 500°C [1999Cho]. The p-type conduction of the Mn doped αFeSi2 (αMnxFe1–xSi2 with x = 0.01, 0.03, 0.06 and 0.1) is revealed by an exponential growth of the Hall coefficient with decreasing temperature down to about –103°C to –153°C with a well pronounced maximum which is a characteristic sign of the presence of a conduction in a band formed by shallow impurity levels. This feature is similar to that observed in the p-type αFeSi2. However the obtained values for the activation energies of the shallow ε1 and deep ε2 acceptor levels are higher in αMnxFe1–xSi2 (x = 0.01, 0.03, 0.06 and 0.1) (ε1 = 90 to 135 meV, ε2 = 219 to 299 meV) than those determined for the undoped αFeSi2 (ε1 = 70 meV, ε2 = 120 meV) [2000Aru]. The electrical resistivities of melt-spun and levitation-melted Mn doped Fe2Si5 alloys decrease with increasing temperature since an increasing number of carriers are thermally excited. The electrical conductivity of the melt-spun sample is lower than the one of the levitation-melted counterparts due to enhanced carrier scattering and hence reduced carrier mobility. The n-type conduction of the Mn doped βFeSi2, called Fe2Si5 by [2003Zhu], is revealed by a negative polarity of the Hall coefficient and a higher electrical conductivity than the non doped βFeSi2. The hardness and the electrical resistivity has been measured along the quasibinary system MnSi-FeSi showing a maximum for the two properties: at 50 mass% FeSi (17 at.% Mn and 66 at.% Si) for the electrical resistivity and about 35 mass% FeSi (22 at.% Mn and 66 at.% Si) for the hardness [1957Tav]. It was found that carbon impurity has a light influence on the position of the maximum of the electrical resistance but its strongly affects the composition curve of the hardness with a maximum at 55 mass% FeSi (15 at.% Mn and 66 at.% Si).
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12
Fe–Mn–Si
After appropriate oxydation, the alloys with 22.7–23.2 mass% Mn, 67.3 mass% Fe and 10-9.5 mass% Si (20.93–21.48 at.% Mn and 18.03–17.21 at.% Si) have very good corrosion resistance in sulfuric acid at room temperature. Corrosion resistance can also be enhanced by addition of Cr to Fe-Mn-Si alloys [1990Ots]. All experimental data regarding works on material properties are reported in Table 6.
Table 1.
Investigations of the Fe-Mn-Si Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1937Vog]
Thermal analysis and optical microscopy
400–1500°C, alloys in the Mn-Fe-MnSi-FeSi ternary area
[1956Gla2]
X-ray diffraction
400°C, Mn3–xFexSi (0 ≤ x ≤ 1)
[1958Aro]
X-ray diffraction
950°C, alloys along FeSi-MnSi and Mn5Si3Fe5Si3
[1962Gla]
X-ray diffraction and microstructure method
MnSi2-FeSi2, MnSi-FeSi, Mn5Si3-Fe5Si3, Mn3SiFe3Si
[1964Gup]
X-ray diffraction
1000°C, Mn30Fe70-Mn10Fe90 (at.%) and up to 20 at.% Si
[1964Sen]
X-ray diffraction
750°C, (Mn,Fe)5Si2 with 34 mass% Fe
[1966Bar]
X-ray diffraction
1000°C, Mn-Fe with Si up to 30 at.%.
[1968Sid]
X-ray diffraction
Annealed for 72 h between 890–1050°C, MnSiFeSi
[1968Fli]
X-ray diffraction
Annealing at 1000°C for 50 h, 53Mn47Si35.5Mn16.4Fe-46.1Si (mass%)
[1968Set]
X-ray diffraction
1000°C, Fe-Mn-Si
[1970Nar]
X-ray diffraction, Mössbauer spectroscopy
–268.8 to 25°C, Mn5–xFexSi3 with 1 ≤ x ≤ 5
[1972Joh]
X-ray diffraction, Mössbauer spectroscopy
Annealed at 950 and then at 750°C, Mn5–xFexSi3 with 0 ≤ x ≤ 5
[1973Bin]
Neutron diffraction
Annealed 3 days at 950°C / Mn5–iFexSi3 with x = 0, 1, 2, 3 and 4
[1973Ish]
X-ray diffraction and thermodynamic calculation
700–1000°C, Fe-Fe1–x(Mn,Si)x with x ≤ 30 at.%.
[1974Abr]
Thermal and microstructural analysis
900–1250°C, MnSi1.72-FeSi2
[1974Bab]
Neutron and X-ray diffraction
Annealed for 50 h at 950°C, (MnxFe1–x)3Si with 0≤x≤1
[1974Bur]
Spin-echo spectroscopy
Annealed for 2 h at 600°C, Mn0.005Fe0.745Si0.25 and Mn0.01Fe0.76Si0.23
[1974Yoo]
Neutron and X-ray diffraction
MnxFe3–xSi with 0 ≤ x ≤ 2 (continued)
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13
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1975Pic]
Neutron diffraction
Annealed for 1 h at 600°C, MnxFe3–xSi with 0 ≤ x ≤ 0.1
[1975Ver]
X-ray diffraction
900°C, (MnxFe1–x)3Si with 0 ≤ x ≤ 1
[1976Sud]
X-ray diffraction
Annealed for 100 h at 947°C, (MnxFe1–x)5Si3 with 0 ≤ x ≤ 1
[1976Zie]
X-ray and neutron powder diffraction measurements
MnFe2Si
[1977Ver]
Mössbauer and neutron diffraction measurements
–195°C and 27°C, (MnxFe1–x)3Si with x = 0, 0.05, 0.10, 0.15, 0.20 and 0.25
[1977Sho]
X-ray diffraction
Mn77Fe4Si19 (at.%)
[1977Yoo]
Neutron and X-ray diffraction
MnxFe3–xSi with 0.75 ≤ x ≤ 3
[1978Gee]
Vapor transport method
1427°C, Fe-Mn-Si
[1980Bob]
High-temperature calorimeter
1527°C, Mn5Si3-Fe5Si3
[1980Smi]
Calorimetry
–269 to 187°C, MnxFe3–xSi with 0.6 ≤ x ≤ 1.6
[1979Ari]
Vapor transport method
1560°C, Fe-Mn-Si with 0.1 < xMn < 20 and 1 < xSi < 20 (at.%)
[1981Fuj]
Electric resistance furnace, EPMA
1515–1534°C, Fe-Mn0-2.62Si0-1.990
[1984Bat]
High-temperature calorimeter
1926, 1953 and 1994°C, Fe-Mn0.0192-0.0258Si0.0040-0.0195
[1984Bon]
High-temperature calorimeter
1687°C, Fe-Mn-Si alloys with xFe:xSi =0.8:0.2, 0.7:0.3, 0.4:0.6 and 0.35:0.65
[1987Bal]
Evaporation measurements
1500°C, Fe-Mn alloys with 3.14 mass% Si
[1988Mil]
Microprobe analysis, X-ray diffraction, capacitance dilatometer
–73 to 27°C, MnFe2Si
[1988Sud]
Calorimetry
1572°C, Mn-FeSi
[1989Bon]
Thermodynamic measurements
1572°C, Mn-FeSi and Fe-MnSi
[1991Lan]
Electron microprobe analysis
900°C, Fe alloys with Mn and Si up to 35 and 12 mass%
[1992For, 1993For]
Thermodynamic models
527–1527°C, Fe-Mn-Si with xMn/xFe = 7.76, 2.14, 0.837, 0.439, 0.176 and 0.0873 with Si from 0 to 35 (in mass%)
[1996Zai1, 1996Zai2, 1998Zai]
Knudsen effusion technique, mass spectrometry analysis of evaporated species
1162–1536°C, liquid Fe-Mn-Si alloys with x(Fe) ≥ 0, x(Si) ≤ 1 and 0 ≤ x(Mn) ≤ 0.6 (at.%)
[1997And]
X-ray, DSC and dilatometer measurements
–73–227°C, Fe-(28.0-28.3)Mn(1.99-7.01)Si (mass%)
[1997Li1]
Substitutional solution model
627 to –73°C, Fe-30Mn6Si (mass%) (continued)
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14
Fe–Mn–Si
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1997Li2]
Unified interaction parameter model
1400°C, Fe-Mn-Si
[2004Can]
Neutron diffraction
Annealed at 952°C for 5 d and water quenched, Mn5–xFexSi3 (x = 4, 3.5, 3, 2, 1)
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
α, (αFe,δFe)
cI2 Im 3m W
(αFe) (Ferrite) < 912
(δFe) 1538–1394
(εFe)
Comments/References
a = 286.65
pure Fe at 25°C [Mas2, V-C2]. Dissolves 5 at.% Mn at 527°C [2004Wit]. Dissolves up to 10 at.% Si at 500°C [Mas2]. pure Fe at 1390°C [V-C2, Mas2] Dissolves 10 at.% Mn at 1474.5°C [2004Wit]. Dissolves up to 19.5 at.% Si at 1280°C [Mas2].
a = 293.15
hP2 P63/mmc Mg
(δMn) 1246–1138
cI2 Im 3m W
γ, (γFe,γMn)
cF4 Fm 3m Cu
(γFe) 1394–912
Lattice Parameters [pm]
(γMn) 1138–1087
metastable martensite obtained by quenching γ(Fe,Mn) alloys [1990Sad] a = 254.34 a = 412.12
Fe-30 mass% Mn-6 mass% Si, (aγ = 360.05) [1999Guo]
a = 308.0
pure Mn [Mas2]. Dissolves up to 12.2 at.% Fe at 1236.6°C [2004Wit]. Dissolves 2 at.% Si [Mas2]
a = 364.67
pure Fe at 915°C [V-C2, Mas2]. Complete solubility with (γMn). Dissolves up to 3.8 at.% Si at 1150°C [2004Wit] pure Mn [Mas2] Complete solubility with (γFe). Dissolves up to 2.8 at.% Si [Mas2]
a = 386.0
(βMn) 1087–707
cP20 P4132 βMn
a = 631.52
pure Mn [Mas2] Dissolves up to 32.9 at.% Fe [2004Wit]. Dissolves up to 16.7 at.% Si [Mas2].
(αMn) < 707
cI58 I 43m (αMn)
a = 891.26
pure Mn at 25°C [Mas2]. Dissolves up to 35.2 at.% Fe [2004Wit]. Dissolves up to 6 at.% Si [Mas2]. (continued)
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15
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αSi) < 1414
cF8 Fd 3m C (diamond)
a = 543.06
at 25°C [Mas2] Practically no solubility for Fe and Mn [Mas2].
(δSi)
hP4 P63/mmc αLa
a = 380 c = 628
at 25°C, 16 GPa [Mas2]
(γSi)
cI16 Im 3m γSi
a = 663.6
at 25°C, 16 GPa [Mas2]
(βSi)
tI4 I41/amd βSn
a = 468.6 c = 258.5
at 25°C, 9.5 GPa [Mas2]
(Mn1–xFex)Si
cP8 P213 FeSi
FeSi < 1410
a = 451.7 ± 0.5
0≤x≤1 x = 1 and at 25°C [1996Spo] x = 0 and at 25°C [1996Spo] 33.36 to 34.46 at.% Si [Mas2] at 300°C [V-C2, Mas2]
MnSi < 1276
a = 455.8
49.5 to 50.2 at.% Si [Mas2] at 10 K [V-C2]
(Mn1–xFex)5Si3 η, Fe5Si3 1060–825 Mn5Si3 < 1300
β, Fe2Si 1212–1040
hP16 P63/mcm Mn5Si3
cP2 Pm 3m CsCl or hP6 P 3m1 Fe2Si
a = 448.8 a = 456.0
0≤x≤1 a = 675.9 ± 0.5 c = 472.0 ± 0.5 a = 691.0 c = 481.4 a = 688.40 ± 0.06 c = 478.76 ± 0.08 a = 686.02 ± 0.06 c = 475.98 ± 0.09 a = 683.22 ± 0.09 c = 474.26 ± 0.11 a = 681.78 ± 0.04 c = 473.60 ± 0.05 a = 680.04 ± 0.05 c = 472.98 ± 0.08
[V-C2]
a = 281
[V-C2] 19.93 to 21.31 at.% Si [Mas2]
a = 405.2 ± 0.2 c = 508.55 ± 0.03
[V-C2]
[V-C2] for Mn4FeSi3 [1972Joh] for Mn3Fe2Si3 [1972Joh] for Mn2Fe3Si3 [1972Joh] for Mn1.5Fe3.5Si3 [1972Joh] for MnFe4Si3 [1972Joh]
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Fe–Mn–Si
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
βFeSi2(h) 1220–937
tP3 P4/mmm βFeSi2
a = 269.01 c = 513.4
αFeSi2(r) < 982
oC48 Cmca αFeSi2
a = 986.3 ± 0.7 b = 779.1 ± 0.6 c = 783.3 ± 0.6
(Mn1–xFex)3Si(h)
cF16 Fm 3m BiF3
α1, Fe3Si < 1247
Lattice Parameters [pm]
a = 572.2 a = 566.0 ± 0.1 to 572.1 ± 0.3
a = 566.25 (Mn1–xFex)3Si(r) α2, Fe3Si < 1284 αMn3Si < 677
a = 281.8 ± 0.1 a = 285.8 ± 0.1 oI186 Immm Mn4Si
Mn4Si7
tP44 P 4c2 Mn4Si7
Mn5Si2 < 850
tP56 P41212 Mn5Si2
R, Mn6Si < 880
hR159 R 3 Co5Cr2Mo3
Mn6Si7
hR39 R 3m Fe7W6
66.7 at.% Si [Mas2, V-C2]
10 to 30 at.% Si [Mas2] [V-C2] 25 to 25.6 at.% Si [Mas2] [V-C2] for 0.75 ≤ x ≤ 3. Two linear variations with composition with a break at x = 1.8 [1977Yoo] MnFe2Si [1976Zie] 0≤x≤1
cP2 Fm 3m CsCl
ν, Mn4Si < 1060
53.4 to 57.38 at.% Si [Mas2] Fe0.87Si2.13 [V-C2]
0≤x≤1
a = 565.0
βMn3Si 1070–677
Comments/References
a = 1699.2 ± 0.4 b = 2863.4 ± 0.7 c = 465.6 ± 0.1 a = 552.5 ± 0.1 c = 1746.3 ± 0.3
a = 890.97 ± 0.02 c = 871.53 ± 0.03 a = 886.6 ± 0.1 c = 865.5 ± 0.1 a = 1087.1 ± 0.5 c = 1918.0 ± 0.9 a = 470 ± 0.1 c = 2561 ± 0.5
10 to 22 at.% Si [Mas2] [1956Gla2] 25 to 25.6 at.% Si [Mas2] [1956Gla2] 16.2 to 18.75 at.% Si [Mas2] [V-C2]
[V-C2]
at 28.6 at.% Si [Mas2] [V-C2] for (Mn,Fe)5Si2 with 34 mass% Fe at 750°C [1964Sen] 12 to 17.75 at.% Si [Mas2] [V-C2] [V-C2]
(continued)
DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Fe–Mn–Si
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
MnSi1.75–x < 1155
tP120 P 4n2 Mn11Si19
Mn15Si26
tI164 I 42d Mn15Si26
a = 553.1 c = 6531.1
[V-C2]
Mn27Si47
tP296 P 4n2 Mn27Si47
a = 553 c = 11790
[V-C2]
τ, Mn77Fe4Si19
mC110 C2 Mn77Fe4Si19
a = 1336.2 ± 0.1 b = 1164.5 ± 0.1 c = 873.4 ± 0.1 β = 90.53 ± 0.01
[1977Sho] from Mn74.5Fe6.0Si19.5 to Mn80.2Fe0.8Si19.0 at 1000°C [1964Gup]
Table 3.
Lattice Parameters [pm]
17
Comments/References
63 to 63.4 at.% Si [Mas2] [V-C2]
a = 552 c = 4820
Invariant Equilibria
Reaction
T [°C]
Type
Phase
Composition (at.%) Fe
Mn
Si
L+(δFe) ⇌ α1, γ
1284
D1
L
69.36
12.59
18.05
L+ α1 ⇌ (Mn,Fe)3Si+γ
1200
U1
L
65.37
10.79
23.84
L+(Mn,Fe)Si ⇌ β+(Mn,Fe)5Si3
1100
U2
L (Mn,Fe)Si β (Mn,Fe)5Si3
58.38 44.26 64.24 51.86
7.70 5.26 3.58 8.27
33.92 50.48 32.18 39.87
L+β ⇌ (Mn,Fe)5Si3+(Mn,Fe)3Si
∼ 1100
U3
L β (Mn,Fe)5Si3 (Mn,Fe)3Si
57.55 64.69 47.48 63.89
12.43 4.21 13.26 10.16
30.02 31.10 39.26 25.95
L+(βMn)+ν ⇌ (Mn,Fe)3Si
1040
P1
L
79.00
2.13
18.87
L+(βMn) ⇌ (Mn,Fe)3Si+γ
1020
U5
L (βMn) (Mn,Fe)3Si γ
35.48 31.36 25.20 39.37
44.08 51.99 46.79 44.31
20.44 16.65 28.01 16.32
L ⇌ (Mn,Fe)3Si+γ
< 1020
e9
L
47.56
27.57
24.87
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18 Table 4.
Fe–Mn–Si Thermodynamic Data of Reaction or Transformation
Reaction or Transformation
Temperature [°C]
Quantity, per mol of atoms [kJ, mol, K]
Comments
0.1 Mn+0.9 Fe0.5Si0.5 ⇌ Mn0.1Fe0.45Si0.45
1572
ΔHmix = –7.4 ± 0.7
[1989Bon]
0.2 Mn+0.8 Fe0.5Si0.5 ⇌ Mn0.2Fe0.4Si0.4
1572
ΔHmix = –11.6 ± 1.5
[1989Bon]
0.3 Mn+0.7 Fe0.5Si0.5 ⇌ Mn0.3Fe0.35Si0.35
1572
ΔHmix = –16.6 ± 1.5
[1989Bon]
0.4 Mn+0.6 Fe0.5Si0.5 ⇌ Mn0.4Fe0.3Si0.3
1572
ΔHmix = –21.5 ± 2.0
[1989Bon]
0.5 Mn+0.5 Fe0.5Si0.5 ⇌ Mn0.5Fe0.25Si0.25
1572
ΔHmix = –25.0 ± 2.0
[1989Bon]
0.1 Fe+0.9 Mn0.5Si0.5 ⇌ Fe0.1Mn0.45Si0.45
1572
ΔHmix = –7.4 ± 0.4
[1989Bon]
0.2 Fe+0.8 Mn0.5Si0.5 ⇌ Fe0.2Mn0.4Si0.4
1572
ΔHmix = –13.6 ± 0.6
[1989Bon]
0.3 Fe+0.7 Mn0.5Si0.5 ⇌ Fe0.3Mn0.35Si0.35
1572
ΔHmix = –19.6 ± 0.9
[1989Bon]
0.4 Fe+0.6 Mn0.5Si0.5 ⇌ Fe0.4Mn0.3Si0.3
1572
ΔHmix = –25.5 ± 1.0
[1989Bon]
0.5 Fe+0.5 Mn0.5Si0.5 ⇌ Fe0.5Mn0.25Si0.25
1572
ΔHmix = –29.4 ± 1.5
[1989Bon]
0.015 Mn+0.8 Fe+0.2 Si ⇌ Fe0.8Mn0.015Si0.2
1687
ΔHmix = –16.3
[1984Bon]
0.255 Mn+0.8 Fe+0.2 Si ⇌ Fe0.8Mn0.255Si0.2
1687
ΔHmix = –18.1
[1984Bon]
0.013 Mn+0.7 Fe+ 0.3Si ⇌ Fe0.7Mn0.013Si0.3
1687
ΔHmix = –32.0
[1984Bon]
0.289Mn+0.7 Fe+0.3 Si ⇌ Fe0.7Mn0.289Si0.3
1687
ΔHmix = –29.9
[1984Bon]
0.049Mn+0.4 Fe+0.6 Si ⇌ Fe04Mn0.013Si0.6
1687
ΔHmix = –32.6
[1984Bon]
0.337Mn+0.4 Fe+0.6 Si ⇌ Fe0.4Mn0.337Si0.6
1687
ΔHmix = –34.8
[1984Bon]
0.010Mn+0.35 Fe+0.65 Si ⇌ Fe0.35Mn0.010Si0.65
1687
ΔHmix = –33.0
[1984Bon]
0.294 Mn+0.35 Fe+0.65 Si ⇌ Fe0.35Mn0.294Si0.65
1687
ΔHmix = –36.0
[1984Bon]
Table 5. Phase MnFeSi
Table 6.
Thermodynamic Properties of Single Phases Temperature Range [°C]
Property, per mole of atoms [J, mol, K]
Comments
1162–1536
ΔfG = –105509 –15.31·T
[1996Zai1, 1998Zai]
0
Investigations of the Fe-Mn-Si Materials Properties
Reference
Method/Experimental Technique
Conditions/Type of Property
[1957Tav]
Resistance and hardness measurements
Alloys along MnSi-FeSi, electrical resistivity and hardness
[1970Nar]
X-ray diffraction, Mössbauer spectroscopy, vibrating sample magnetometer, Faraday method
Mn5–xFexSi3 with 1 ≤ x ≤ 5, magnetic moment, Curie and Neel temperatures, susceptibility, magnetization (continued)
DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Fe–Mn–Si
19
Reference
Method/Experimental Technique
Conditions/Type of Property
[1972Joh]
X-ray diffraction, Mössbauer spectroscopy, vibration sample magnetometer
Mn5–xFexSi3 with 0 ≤ x ≤ 5, magnetic hyperfine field, magnetization, magnetic moment, susceptibility
[1973Bin]
Neutron diffraction and Sucksmith balance
Mn5–xFexSi3 with x = 0, 1, 2, 3 and 4, magnetic saturation moment
[1974Yoo]
Vibrating sample magnetometer, Sucksmith ring balance, neutron diffraction
MnxFe3–xSi (0 ≤ x ≤ 2), magnetization, Curie TC and ordering TR temperatures
[1976Lut]
Adiabatic calorimeter, absolute quartz dilatometer, high frequency resonance method
(MnxFe1–x)3Si (0 ≤ x ≤ 0.45), Curie TC and Neel TN temperatures
[1976Nic]
X-ray, magnetization and NMR techniques
MnxFe3–xSi (0 ≤ x ≤ 1), magnetic moment, hyperfine field
[1976Sid]
Vibrating sample magnetometer, pendulum balance, Mössbauer spectroscopy, neutron diffraction
(MnxFe1–x)3Si with 0 ≤ x ≤ 1, magnetization, average magnetic moment, susceptibility
[1976Sud]
Magnetization, heat capacity and thermal expansion measurements
(Mn1–xFex)5Si3 with 0 ≤ x ≤ 1, ordering TR, Neel TN and Curie TC temperatures, thermal expansion coefficient, magnetization
[1976Zie]
Vibrating sample magnetometer, modified Sucksmith ring balance, neutron diffraction
MnFe2Si, spontaneous magnetization, magnetic susceptibility, magnetic moment
[1976Zvi]
Dilatometer, tensile tests, X-ray diffraction
Fe-(20-30)Mn(0-10)Si (mass%), strength, hardness, susceptibility
[1977Nic]
Vibrating sample magnetometer
MnxFe3–xSi, saturation magnetization
[1977Yoo]
Four terminal potentiometric technique Curie temperature, saturation magnetization,
MnxFe3–xSi with 0.75 ≤ x ≤ 3.0, resistivity, magnetic moment
[1978Nar]
Tensile tests, fluxmeter, strain gauge technique, vibrating sample magnetometer
Fe alloys with 6.6 mass% Si and 0.01, 0.12 and 0.2 mass% Mn, tensile strength, saturation magnetization
[1983Bus]
Polar Kerr rotation measurements, vibrating sample magnetometer
MnFe2Si, saturation moment
[1986Sat]
Tensile tests, Faraday magnetic balance, Neel temperature, magnetic susceptibility, yield stress
Fe-(21.2-32)Mn(0.99-6.54) (mass%), martensitic transformations
[1986Mur]
Tensile tests yield stress, Neel temperature
Fe-(25-35)Mn-(0-6)Si (mass%), martensitic transformation
[1987Gel]
Thermal expansion and magnetic susceptibility measurements
MnyFe1-ySi (0 ≤ y ≤ 1), thermal expansion coefficient
[1988Mil]
Microprobe analysis, X-ray diffraction, capacitance dilatometer
–73 to 27°C, MnFe2Si (continued)
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20
Fe–Mn–Si
Reference
Method/Experimental Technique
Conditions/Type of Property
[1990Ots]
Tensile test, vibrating sample magnetometer, electric resistance measurements, yield stress
Fe-(11-3)Mn(5-6)Si with 0 to 13 mass% Cr and 0 to 8 mass% Ni, martensitic transformations
[1990Sad]
Elongation stress, dilatometric measurements, yield stress
Fe-(24-35)Mn(2-9)Si, martensitic transformation and austenitization
[1991Mil]
Vibrating sample magnetometer, threeterminal capacitance dilatometer
Mn0.8Fe2.2Si, thermal expansion
[1991Lan]
Corrosion tester
(22.7–23.2)Mn67.3Fe(10–9.5)Si (mass%), ultimate tensile strength, corrosion resistance
[1993Wat]
Micro Vickers hardness tester, tensile strain tests, shrinkage fraction measurement
Fe-31Mn6Si (mass%) powder deformedheated 13 times, critical resolved shear stress, strain and electrical resistance
[1995Ron]
Electrical resistance, magnetic susceptibility measurements, TEM
Fe-29.9Mn-6.0Si (mass%) prestrained, shape memory effect
[1996Mat]
Tensile prestrain tests, thermomechanical treatment
Fe-28Mn-6Si-5Cr-0.03C (mass%) hot rolled shape memory recovery
[1996Qin]
Magnetic balance, metallography, Neel temperature, magnetic susceptibility, electrical resistivity
Fe-(27.77-32.01)Mn-(4.10-6.10)Si, martensitic transformation and austenitization
[1996Spo]
Soft X-ray emission spectroscopy, XRD, electrical resistivity
Mn1–xFexSi with (0 ≤ x ≤ 1) (at.%), magnetic transition temperature
[1996Zha]
Electrical resistance measurements, XRD analysis, electrical resistance
Fe-29.44Mn-6.2Si (mass%) hot forged and hot rolled, martensitic transformation and austenitization
[1998Cot1]
Dilatometric measurement
Fe-(15–35)Mn-(0–6.5)Si (mass%), martensitic transformation and austenitization
[1998Muk]
Density, thermal, susceptibility measurements, linear thermal expansion, Debye temperature,
MnxFe3–xSi (1.2 ≤ x ≤ 1.8) homogenized one week at 800°C, Grüneisen parameters
[1999Che]
DMA and DME analysis, Neel temperature, internal friction, elastic modulus
Fe-30 mass% Mn-6 mass% Si-5 mass% Cr, forged and rolled, martensitic transformation and austenitization
[2000Aru]
Van der Pauw’s technique, alternating current magnetic method, resistivity and Hall coefficient measurements
MnxFe1–xSi2 (x = 0.01, 0.03, 0.06 and 0.1) powder cold pressing and sintered and then annealed at 800°C for 48 hours
[2000Dai]
Tensile and bending methods
Fe-29.7Mn5.5Si and Fe-18.1Mn5.5Si forged, stress, residual strain
[2000Wu]
Electrostatic audio-frequency internal friction measurements
Fe-30.30Mn6.10Si (mass%), hot-forged and hot-rolled, internal friction, elastic modulus, martensitic and Neel temperatures (continued)
DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Fe–Mn–Si
21
Reference
Method/Experimental Technique
Conditions/Type of Property
[2000Jin1, 2000Jin2]
X-ray diffraction profile analysis, fourprobe electrical resistance method
Fe-(23.37-30.3)Mn(5.87-6.06)Si (in mass%), hot-forged and hot-rolled, martensitic start temperature
[1999Cho, 2001Cho]
Thermoelectric electromotive force measurements, Seebeck coefficient, electrical conductivity
Mn0.1Fe0.9Si2 powders sintered at 1150–1190°C under 230 MPa, annealed at 800 or 865°C for 100 hours
[2002Son]
Quantum magnetometer
FexMn5–xSi3 (x = 0, 1, 2, 3, 4, 5) annealed at 952°C for 5 days, ac-susceptibility, magnetization, magnetic-entropy
[2003Che]
Thermal diffusion and Hall measurements, Seebeck coefficient and electrical conductivity experiments
Mn0.08Fe0.92Six (1.9 ≤ x ≤ 2.5), hot pressed (50 MPa at 975°C) annealed 800°C, thermal conductivity
[2003Zhu]
Seebeck coefficient, electrical resistivity
Mn0.14Fe1.86Si5, cold pressed under 860 MPa, sintered at 1150°C and annealed at 800°C
[2004Can]
Curie temperature and magnetic moment measurements
Annealed at 952°C for 5 days and water quenched, Mn5–xFexSi3 (x = 4, 3.5, 3, 2, 1)
[2005Jia]
Seebeck coefficient measurements
Mn2.36Fe27.14Si70.5 (at.%) alloy annealed at 550°C, Hall coefficient, electrical resistivity
[2005Lin]
Microvickers tester, bending test
Fe-30Mn6Si (mass%) alloy, hardness
[2005Sta]
SEM, bending strain, shape memory effect
Fe, 11 mass% Mn, 4.5 mass% Si, 10 mass% Cr, austenitized at 1000°C
[2007Sta]
X-ray diffraction, tensile strength, hardness, SEM observations,
Fe, 12 to 18 mass% Mn, 4 to 4.5 mass% Si, 10 mass% Cr, austenizing treatment at 1000°C, then annealing 500–800°C.
Landolt-Börnstein New Series IV/11D4
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22
Fig. 1. Fe-Mn-Si.
Fe–Mn–Si
Quasibinary section FeSi-MnSi
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Landolt-Börnstein New Series IV/11D4
23
Fig. 2. Fe-Mn-Si.
Reaction scheme
Fe–Mn–Si
Landolt-Börnstein New Series IV/11D4
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24
Fig. 3. Fe-Mn-Si.
Fe–Mn–Si
Liquidus surface projection
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Fe–Mn–Si
Fig. 4. Fe-Mn-Si.
Landolt-Börnstein New Series IV/11D4
25
Isothermal section at 700°C
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Fig. 5. Fe-Mn-Si.
Fe–Mn–Si
Isothermal section at 900°C
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Fe–Mn–Si
Fig. 6. Fe-Mn-Si.
Landolt-Börnstein New Series IV/11D4
27
Isothermal section at 1000°C
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Fig. 7. Fe-Mn-Si.
Fe–Mn–Si
Vertical section at xMn/xFe = 7.76
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Fe–Mn–Si
Fig. 8. Fe-Mn-Si.
Landolt-Börnstein New Series IV/11D4
29
Vertical section at xMn/xFe = 2.14
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Fig. 9. Fe-Mn-Si.
Fe–Mn–Si
Silicon isoactivity lines in liquid alloys at 1427°C
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Fig. 10. Fe-Mn-Si.
Landolt-Börnstein New Series IV/11D4
31
Manganese isoactivity lines in liquid alloys at 1427°C
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Fig. 11. Fe-Mn-Si.
Fe–Mn–Si
Iron isoactivity lines in liquid alloys at 1427°C
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Fig. 12. Fe-Mn-Si.
Landolt-Börnstein New Series IV/11D4
33
Lines of constant molar Gibbs free energy of formation in liquid alloys at 1427°C, in kJ·mol–1
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34 References [1937Vog]
[1938Jae] [1949Jae] [1956Gla1]
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[1964Sen]
[1966Bar]
[1968Fli]
[1968Set]
[1968Sid]
[1970Nar]
Fe–Mn–Si
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DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
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Fe–Mn–Si
[1972Joh]
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Landolt-Börnstein New Series IV/11D4
35
J. Phys. Chem. Solids, 31, 1511–1524 (1970) (Crys. Structure, Experimental, Magn. Prop., 23) Johnson, V., Weiher, J.F., Frederick, C.G., Rogers, D.B., “Magnetic and Mössbauer Effect Studies of Mn5Si3:Fe5Si3 Solid Solutions”, J. Solid State Chem., 4, 311–323 (1972) (Crys. Structure, Experimental, Magn. Prop., 27) Binczycka, H., Dimitrijevic, Z., Gajic, B., Szytula, A., “Atomic and Magnetic Structure of Mn5–xFexSi3”, Phys. Status Solidi A, 19, K13–K17 (1973) (Crys. Structure, Experimental, Mechan. Prop., 10) Ishida, K., Shibuya, K., Nishizawa, T., “γ Loops in Fe-Mn-V, Fe-Mn-Mo, and Fe-Mn-Si Systems” (in Japanese), J. Jpn. Inst. Met., 37, 1305–1312 (1973) (Experimental, Kinetics, Thermodyn., 22) Abrikosov, N.Kh., Ivanova, L.D., Petrova, L.I., “The System MnSi1.72-FeSi2”, Inorg. Mater., 10(12), 1907–1908 (1974), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 10(12), 2226–2227 (1974) (Experimental, Phase Diagram, Phase Relations, 12) Babanova, E.N., Sidorenko, F.A., Geld, P.V., Basov, N.I., “Ordering of (Fe1–xMnx)3Si Solid Solutions”, Phys. Met. Metallogr., 38(3), 122–126 (1974), translated from Fiz. Met. Metalloved., 38(3), 586–590 (1974) (Crys. Structure, Experimental, 5) Burch, T.J., Litrenta, T., Budnick, J.I., “Hyperfine Studies of Site Occupation in Ternary Systems”, Phys. Rev. Letters, 33(7), 421–424 (1974) (Crys. Structure, Experimental, 12) Yoon, S., Booth, J.G., “Structural and Magnetic Properties of Fe3–xMnxSi Alloys”, Phys. Letters, 48A(5), 381–382 (1974) (Crys. Structure, Experimental, Magn. Prop., 4) Pickart, S., Litrenta, T., Burch, T., Budnick, J.I., “Site Preference of Dilute Transition Metal Solutes in Fe3Si”, Phys. Lett. A, 53A(4), 321–323 (1975) (Crys. Structure, Experimental, 4) Vereshchagin, YuA., Babanova, E.N., Sidorenko, F.A., “Procedure for Synthesis, Homogeneity Regions and Lattice Spacings of (Fe1–xMx)3Si (M = Mn, Co, Ni) Solid Solutions with a D03 Structure” (in Russian), Fiz. Metody Issled. Tverdogo Tela, 1, 40–42 (1975) (Crys. Structure, Experimental, 9) Lutskaya, L.F., Kalishevich, G.I., Geld, P.V., “Magnetic Phase Diagram of (Fe1–xMnx)3Si Solid Solutions”, Sov. Phys. Solid State, 18(6), 1011–1012 (1976), translated from Fiz. Tverd. Tela, 18(6), 1738–1740 (1976) (Experimental, Magn. Prop., 6) Narita, K., Enokizono, M., “Effects of Ni, Al, and Mn Addition on the Mechanical and Magnetic Properties of 6.5% Si-Fe Sheets”, IEEE Trans. Magn., 12(6), 873–873 (1976) (Abstract, 0) Niculescu, V., Raj, K., Budnick, J.I., Burch, T.J., Hines, W.A., Menotti, A.H., “Relating Structural, Magnetisation and Hyperfine Field Studies to a Local Environment Model in Fe3–xVxSi and Fe3–xMnxSi”, Phys. Rev. B, 14(9), 4160–4176 (1976) (Crys. Structure, Magn. Prop., Thermodyn., 26) Sidorenko, F.A., Vereshchagin, Yu.A., Babanova, E.N., Ryzhenko, B.V., Geld, P.V., “Magnetic Transitions in Mutual Solid Solutions of the Lower Iron and Manganese Silicides”, Sov. Phys.- JETP, 43(2), 225–331 (1976), translated from Zhur. Eksp. i Teor. Fiz., 70, 628–638 (1976) (Crys. Structure, Experimental, Magn. Prop., 11) Sudakova, N.P., Kuznetsov, S.I., Mikhelson, A.V., Sychev, N.I., Geld, P.V., “Magnetic Transitions in Mn-Fe Silicide Solid Solutions”, Sov. Phys.-Dokl., 21(5), 273–275 (1976), translated from Dokl. Akad. Nauk SSSR, 228, 582–585 (1976) (Experimental, Magn. Prop., 12) Ziebeck, K.R.A., Webster, P.J., “The Antiferromagnetic and Ferromagnetic Properties of Fe2MnSi”, Philos. Mag., 34(6), 973–982 (1976) (Crys. Structure, Experimental, Magn. Prop., 17) Zvigintseva, G.E., Logunov, V.Ya., “Phase Transformations and Mechanical Properties of Fe-Mn-Si Alloys”, Russ. Metall., (4), 169–172 (1976), translated from Izv. Akad. Nauk SSSR, Met., (4), 197–200 (1976) (Experimental, Mechan. Prop., 11) Niculescu, V., Budnick, J.I., “Limits of Solubility, Magnetic Properties and Electron Concentration in Fe3–xTxSi System”, Solid State Commun., 24(9), 631–634 (1977) (Crys. Structure, Experimental, Magn. Prop., 17)
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36 [1977Sho] [1977Ver]
[1977Yoo]
[1978Gee] [1978Nar]
[1979Ari] [1980Bob]
[1980Smi]
[1981Fuj]
[1983Bus]
[1984Bat]
[1984Bon]
[1986Mur]
[1986Sat] [1987Bal]
[1987Gel]
[1988Mil] [1988Ray] [1988Sud]
Fe–Mn–Si Shoemaker, C.B., Shoemaker, D.P., “The Crystal Structure and Superstructure of the K Phase, Mn77Fe4Si19”, Acta Crystallogr., 33B, 743–754 (1977) (Crys. Structure, Experimental, 14) Vereshchagin, Yu.A., Sidorenko, F.A., Ryzhenko, B.V., Babanova, E.N., Geld, P.V., “Peculiarities of Filling Unequivalent Sites in the Quasi-Binary Solid Solutions with L21 Lattice” (in Russian), Ukrain. Fiz. Zhur., 22(9), 1472–1476 (1977) (Crys. Structure, Experimental, 11) Yoon, S., Booth, J.G., “Magnetic Properties and Structures of Some Ordered (FeMn)3Si Alloys”, J. Phys. F: Met. Phys., 7(6), 1079–1095 (1977) (Crys. Structure, Experimental, Magn. Prop., 19) Gee, R., Rosenqvist T., “Activities in Liquid Fe-Mn-Si Alloys”, Scandin. J. Metall., 7(1), 38–41 (1978) (Experimental, Thermodyn., 10) Narita, K., Enokizono, M., “Effect of Nickel and Manganese Addition on Ductility and Magnetic Properties of 6.5 % Silicon-Iron Alloy”, IEEE Trans. Magn., 14(4), 258–262 (1978) (Experimental, Magn. Prop., Mechan. Prop., Morphology, Phase Relations, 5) Arita, M., Saint Pierre, R., “Activity Measurements in Liquid Fe-Mn-Si Alloys at 1833K”, Trans. Iron Steel Inst. Jpn., 19, 374–376 (1979) (Experimental, Thermodyn., 14) Bobrov, N.P., Sudakova, N.P., Kuznetsov, S.I., Esin, Y.O., Sintez, I, “About Thermodynamics of Alloys of Fe5Si3-Mn5Si3 System” (in Russian), Svoistya Beskiscorod. Neorgan. Materialov, Sverdlovsk, 52–54 (1980) (Experimental, Thermodyn., 5) Smith, T.F., Bristow, G.K., Boughton, C.V., Hudson, P.R.W., “Specific Heat Capacity of Fe3–xMnxSi Compounds”, J. Phys. F, Metal Phys., 10, 2867–2873 (1980) (Experimental, Thermodyn., 8) Fujisawa, T., Imaoka, K., Sakao, H., “Activities of Manganese and Silicon in δ -Iron at Solid-Liquid Equilibrium Temperatures”, Trans. Iron Steel Inst. Jpn., 21, 559–567 (1981) (Experimental, Phase Relations, Thermodyn., 68) Buschow, K.H.J., Van Engen, P.G., Jongebreur, R., “Magneto-Optical Properties of Metals Ferromagnetic Materials”, J. Magn. Magn. Mater., 38, 1–22 (1983) (Experimental, Magn. Prop., Optical Prop., 23) Batalin, G.I., Sudavtsova, V.S., Bondarenko, T.P., Ul’yanov, V.I., “Heat of Mixing of Mn, Si, Al, Ti in rich Iron and its Alloys” (in Russian), Ukr. Khim. Zh., 50, 582–585 (1984) (Experimental, Calculation, Thermodyn., 4) Bondarenko, T.P., Sudavtsova, V.S., Batalin, G.I., Ulyanov, V.I., “Enthalpies of Mixing for Liquid Fe-Si-Mn Alloys”, Russ. Metall., (3), 75–76 (1984), translated from Izv. Akad. Nauk SSSR, Metall., (3), 81–82 (1984) (Calculation, Experimental, Thermodyn., 3) Murakami, M., “Effect of Alloying Content, Phase and Magnetic Transformation on the Shape Memory Effect of Fe-Mn-Si Alloys” (in Japanese), Tetsu To Hagane,, 72(13), 1574–1574 (1986) (Experimental, Magn. Prop., Mechan. Prop., 0) Sato, A., Yamaji, Y., Mori, T., “Physical Properties Controlling Shape Memory Effect in Fe-Mn-Si Alloys”, Acta Metall., 34(2), 287–294 (1986) (Experimental, Magn. Prop., 30) Balkavoi, Yu.V., Aleev, R.A., Manokhin, A.A., “Influence of Carbon, Silicium and Vanadium on the Activity of Mn in Fe” (in Russian), Izv. V.U.Z. Chern. Metall., (9), 128–129 (1987) (Calculation, Thermodyn., Experimental, 3) Geld, P.V., Povzner, A.A., Kortov, S.V., Krentsis, R.P., “Thermal Expansion and Weak Itinerant Magnetism in Fe1–yMnySi and Fe1–xCoxSi Solid Solutions”, Sov. Phys. -Dokl., 32(12), 1006–1008 (1987), translated from Dokl. Akad. Nauk SSSR, 297, 1359–1363, (1987) (Experimental, Thermodyn., Magn. Prop., 10) Miles, J.R., Smith, T.F., Finlayson, T.R., “Thermal Expansion of Fe2MnSi”, Aust. J. Phys., 41, 781–790 (1988) (Experimental, Thermodyn., Magn. Prop., 24) Raynor, G.V., Rivlin, V.G., “Fe-Mn-Si” in “Phase Equilibria in Iron Ternary Alloys”, Inst. Metals, London, 363–377 (1988) (Crys. Structure, Phase Diagram, Phase Relations, Review, 20) Sudavtsova, V.S., Batalin, G.I., Bondarenko, T.P., “The Enthalpies of Formation of Liquid Ferrosilicon-Manganese {(FeSi)-Mn} Alloys”, Russ. J. Phys. Chem., 62, 702–703 (1988), translated from Zh. Fiz. Khim., 62, 1383–1385 (1988) (Experimental, Thermodyn., 5)
DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–Si [1989Bon]
[1990Ots]
[1990Sad]
[1991Lan]
[1991Mil] [1992For]
[1993For]
[1993Wat] [1994Rag] [1995Bou]
[1995Ron]
[1995Tan]
[1996Mat]
[1996Qin]
[1996Spo]
[1996Zai1]
[1996Zai2]
[1996Zha]
Landolt-Börnstein New Series IV/11D4
37
Bondarenko, T.P., Sudavtsova, V.S., “Thermochemical Properties in Melts of (MnSi)-Fe, (FeSi)-Mn” (in Russian), Ukr. Khim. Zh., 55(1), 23–24 (1989) (Experimental, Thermodyn., 7) Otsuka, H., Yamada, H., Maruyama, T., Tanahashi, H., Matsuda, S., Murakami, M., “Effects of Alloying Additions on Fe-Mn-Si Shape Memory Alloys”, ISIJ Int., 30(8), 674–679 (1990) (Experimental, Mechan. Prop., Magn. Prop., 19) Sade, M., Halter, K., Hornbogen, E., “Transformation Behaviour and One-Way Shape Memory Effect in Fe-Mn-Si Shape Memory Alloys”, J. Mater. Sci. Lett., 9, 112–115 (1990) (Experimental, Magn. Prop., 8) Langhong, L., Tianxiang, Z., Yaoxiao, Z., “Development of Fe3Si Alloys” in “High-Temp. Ordered Intermetallic Alloys IV ”, Mater. Res. Soc. Symp. Proc., 213, 833–838 (1991) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, Phys. Prop., 8) Miles, J.R., Smith, T.F., Finlayson, T.R., “Thermodynamic Study of the Magnetic Ordering of Fe3–xMnxSi Alloys”, Int. J. Thermoph., 12(4), 617–626 (1991) (Experimental, Magn. Prop., 11) Forsberg, A., Agren, J., “Thermodynamics, Phase Equilibria and Martensitic Transformation in Fe-Mn-Si Alloys” in “Shape-Memory Materials and Phenomena - Fundamental Aspects and Applications”, Mater. Res. Soc., Symp. Proc., Pittsburg, PA, USA, 246, 289–295 (1992) (Phase Diagram, Phase Relations, Thermodyn., 24) Forsberg, A., Agren, J., “Thermodynamic Evaluation of the Fe-Mn-Si System and the γ/ε Martensitic Transformation”, J. Phase Equilib., 14(3), 354–363 (1993) (Calculation, Phase Diagram, Phase Relations, Thermodyn., 29) Watanabe, Y., Mori, Y., Sato, A., “Training Effect in Fe-Mn-Si Shape-Memory Alloys”, J. Mater. Sci., 28, 1509–1514 (1993) (Experimental, Mechan. Prop., 9) Raghavan, V., “Fe-Mn-Si (Iron-Manganese-Silicon)”, J. Phase Equilib., 15(6), 619–620 (1994) (Phase Diagram, Phase Relations, Review, 13) Bouchard, D., Bale, C.W., “Simultaneous Optimization of Thermochemical Data for Liquid Iron Alloys Containing C, N, Ti, Si, Mn, S, and P”, Metall. Mater. Trans. B, 26B(3), 467–484 (1995) (Phase Relations, Theory, Thermodyn., 85) Rong, L.J., Li, Y.Y., Shi, S.X., “An Electron Microscopy Study of the Reverse Transformation Behaviour of Stress-Induced Martensite in an Fe-Mn-Si Alloy”, Mater. Lett., 24, 143–146 (1995) (Experimental, Mechan. Prop., 12) Tanaka, T., Aaronson, H.I., Enomoto, M., “Calculations of α/γ Phase Boundaries in Fe-CX1-X2 Systems from the Central Atoms Model”, Metall. Mater. Trans. A, 26A(3), 535–545 (1995) (Calculation, Phase Relations, Thermodyn., 27) Matsumura, O., Furusako, S., Furukawa, T., Otsuka, H., “Formation of Surface Texture and Anisotropy of Shape Memory Effect in an Fe-Mn-Si Alloy”, ISIJ Int., 36(8), 1103–1108 (1997) (Experimental, Mechan. Prop., 18) Qin, Z., Yu, M., Zhang, Y., “Neel Transition and γ -ɛ Transformation in Polycrystalline FeMn-Si Shape Memory Alloys”, J. Mater. Sci., 31(9), 2311–2315 (1996) (Electr. Prop., Experimental, Morphology, Phys. Prop., Thermodyn., 19) Spolenak, R., Bauer, E., Müller, H., Kirchmayr, H., Blaha, P., Mohn, P., Schwarz, K., “(Fe,Mn) Si: Transitions from a Magnetic to a Nonmagnetic State and from Metallic to Semiconducting Behaviour”, J. Magn. Magn. Mat., 157–158, 715–716 (1996) (Theory, Magn. Prop., 8) Zaitsev, A.I., Litvina, A.D., Shelkova, N.E., Mogutnov, B.M., “Thermodynamic Properties of {xFe+yMn+(1–x–y)Si}(I) and {xFe+yMn+(1–x–y)P(I)”, J. Chem. Thermodyn., 28, 285–306 (1996) (Experimental, Thermodyn., 66) Zaitsev, A.I., Mogutnov, B.M., “The Thermodynamic Properties of Iron-Manganese-Silicon Melts”, Russ. J. Phys. Chem., 70(4), 554–561 (1996), translated from Zh. Fiz. Khim., 70(4), 599–606 (1996) (Experimental, Kinetics, Phys. Prop., Thermodyn., 34) Zhao, Z.T., Liu, T., Liu, G.W., Ma, R.Z., “Reverse Martensite Transformation Induced by Strain in Fe-Mn-Si Alloy”, J. Mater. Sci. Lett., 15(16), 1427–1428 (1996) (Crys. Structure, Electr. Prop., Experimental, Phase Relations, 5)
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DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
38 [1997And] [1997Li1] [1997Li2]
[1998Cot1]
[1998Cot2]
[1998Mie]
[1998Muk] [1998Qin]
[1998Zai]
[1999Che]
[1999Cho]
[1999Cot]
[1999Guo]
[1999Hsu] [1999Jin]
[2000Aru]
[2000Dai] [2000Hsu] [2000Jin1]
Fe–Mn–Si Andersson, M., “Enthalpy Change of the ε→γ Transformation in Fe-Mn-Si Alloys”, Scand. J. Metall., 26, 139–144 (1997) (Experimental, Thermodyn., 10) Li, L., Hsu, T.Y., “Gibbs Free Energy Evaluation of the fcc(γ ) and hcp(ε) Phases in the Fe-Mn-Si Alloys”, Calphad, 21(3), 443–448 (1997) (Calculation, Thermodyn., 17) Li, H., Morris, A., “Evaluation of Unified Interaction Parameter Model Parameters for Calculating Activities of Ferromanganese Alloys: Mn-Fe-C, Mn-Fe-Si, and Mn-Fe-C-Si Systems”, Metall. Mater. Trans. B, 28B(4), 553–562 (1997) (Thermodyn., Calculation, 41) Cotes, S., Fernandez-Guilermet, A., Sade, M., “Phase Stability and FCC/HCP Martensitic Transformation in Fe-Mn-Si Alloys Part 1: Experimental Study and Systematics of the Ms and As Temperatures”, J. Alloys Compd., 278, 231–238 (1998) (Experimental, Phase Relations, 28) Cotes, S., Fernandez-Guillermet, A., Sade, M., “Phase Stability and fcc/hcp Martensitic Transformation in Fe-Mn-Si Alloys: Part II. Thermodynamic Modeling of the Driving Forces and the Ms and As Temperatures”, J. Alloys Compd., 280, 168–177 (1998) (Theory, Phase Relations, Thermodyn., 17) Miettinen, J., “Reassessed Thermodynamic Solution Phase Data for Ternary Fe-Si-C System”, Calphad, 22(2), 231–256 (1998) (Calculation, Assessment, Thermodyn., Phase Relations, 36) Mukherjee, G.D., Bansal, C., Chatterjee, A., “Thermal Expansion Study of Fe-Mn-Si Alloys”, Physica B, 254, 223–233 (1998) (Experimental, Theory, Magn. Prop., 33) Qin, Z., Zhang, Y., “Mössbauer Effect Study of Antiferromagnetic Fe-Mn-Si Alloys”, Hyperfine Interact., 116, 225–235 (1998) (Crys. Structure, Electronic Structure, Experimental, 20) Zaitsev, A.I., Litvina, A.D., Shelkova, N.E., Mogutnov, B.M., “Association in Ternary Metallic Melts Fe-Mn-Si, Fe-Cr-P and Fe-Mn-P”, Thermochim. Acta, 314, 307–315 (1998) (Experimental, Phase Relations, Thermodyn., 31) Chen, S., Chung, C.Y., Yan, C., Hus, T.Y., “Effect of f.c.c. Antiferromagnetism on Martensitic Transformation in Fe-Mn-Si Based Alloys”, Mater. Sci. Eng. A, A264, 262–268 (1999) (Experimental, Mechan. Prop., 28) Cho, W.-S., Choi, S.-W., Park, K., “Microstructure and Thermoelectric Properties of p-type Fe0.9Mn0.1Si2 Compounds Prepared by Pressureless Sintering”, Mater. Sci. Eng. B, B68, 116–122 (1999) (Experimental, Phase Relations, Electr. Prop., 28) Cotes, S., Fernandez-Guillerment, A., Sade, M., “Gibbs Energy Modelling of the Driving Forces and Calculation of the FCC/HCP Martensitic Transformation Temperatures in Fe-Mn and Fe-Mn-Si Alloys”, Mat. Sci. Eng. A, A275, 503–506 (1999) (Calculation, Thermodyn., 21) Guo, Z., Rong, Y., Chen, S., Hsu, T.Y., “Crystallography of fcc(γ)→hcp(ε) Martensitic Transformation in Fe-Mn-Si Based Alloys”, Scr. Mater., 41(2), 153–158 (1999) (Theory, Phase Relations, 15) Hsu, T.Y., Zhuyao, X., “Martensitic Transformation in Fe-Mn-Si Based Alloys”, Mater. Sci. Eng. A, A273–275, 494–497 (1999) (Theory, Phys. Prop., Phase Relations, 40) Jin, X.J., Hsu, T.Y., “Thermodynamic Consideration of Antiferromagnetic Transition on fcc (γ )-hcp(ε) Martensitic Transformation in Fe-Mn-Si Shape Memory Alloys”, Mater. Chem. Phys., 61, 135–138 (1999) (Calculation, Phase Relations, 19) Arushanov, E., Schoen, J.H., Tani, J.-I., Kido, H., “Transport Properties of β-Fe1–xMnxSi2 Alloys”, Phys. Status Solidi A, 181A, 185–191 (2000) (Crys. Structure, Electr. Prop., Experimental, 17) Dai, P., “Effect of Solution Hardening on the Shape Memory Properties of Fe-Mn-Si based Alloys”, J. Mater. Sci. Lett., 19, 111–113 (2000) (Experimental, Mechan. Prop., 9) Hsu, T.Y., “Fe-Mn-Si Based Shape Memory Alloys”, Mater. Sci. Forum, 327–328, 199–206 (2000) (Review, Experimental, Magn. Prop., 62) Jin, X., Hsu, T.Y., “Prediction of Martensitic Transformation Start Temperature Ms in FeMn-Si Shape Memory Alloys”, Mater. Sci. Forum, 327–328, 219–222 (2000) (Experimental, Phase Relations, 14)
DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–Si [2000Jin2] [2000Wu]
[2001Cho]
[2001Wan1] [2001Wan2]
[2002AlN]
[2002Son]
[2002Tan]
[2003Che]
[2003Luo] [2003Zhu]
[2004Can]
[2004Wit]
[2005Jia]
[2005Lin]
[2005Mec]
[2005Sta] [2005Wan]
Landolt-Börnstein New Series IV/11D4
39
Jin, X., Jihua, Z., Hsu, T.Y., “Thermodynamic Calculation of Stacking Fault Energy in Fe-MnSi Shape Memory Alloys”, Mater. Design, 21, 537–539 (2000) (Theory, Thermodyn., 19) Wu, X., Hsu, T.Y., “Effect of the Neel Temperature, TN, on Martensitic Transformation in Fe-Mn-Si-Based Shape Memory Alloys”, Mater. Charact., 45(2), 137–142 (2000) (Experimental, Phase Relations, 20) Cho, W.-S., “Thermodynamic Properties of Fe0.9Mn0.1Si2 Thermoelectric Materials Prepared by Presureless Sintering”, Key Eng. Mater., 189–191, 315–321 (2001) (Crys. Structure, Electr. Prop., Experimental, Phys. Prop., 10) Wan, J., Chen, S., Hsu, T.Y., “The Stability of Transition Phases in Fe-Mn-Si Based Alloys”, Calphad, 25(3), 355–362 (2001) (Calculation, Thermodyn., 21) Wan, J., Chen, S.P., Hsu, T.Y., “Group Theory Analyses of Transition Structures Related to the γ→ε Transformation in Fe-Mn-Si Based Alloys”, Mater. Chem. Phys., 71, 90–93 (2001) (Theory, Phase Relations, 16) Al-Nawashi, G.A., Mahmood, S.H., Lehlooh, A.-F.D., Saleh, A.S., “Mössbauer Spectroscopic Study of Order-Disorder Phenomena in Fe3–xMnxSi”, Physica B, 321, 167–172 (2002) (Crys. Structure, Theory, Electronic Structure, 11) Songlin, D., Tegus, O., Brueck, E., Klaasse, J.C.P., de Boer, F.R., Buschow, K.H.J., “Magnetic Phase Transition and Magnetocaloric Effect in Mn5–xFexSi3”, J. Alloys Compd., 334, 249–252 (2002) (Experimental, Magn. Prop., Phase Relations, 6) Tani, J.I., Kido, H., “First Principle Calculation of the Geometrical and Electronic Structure of Impurity-Doped β-FeSi2 Semiconductors”, J. Solid State Chem., 163, 248–252 (2002) (Calculation, Crys. Structure, Phys. Prop., 23) Chen, H.Y., Zhao, X.B., Lu, Y.F., Mueller, E., Mrotzek, A., “Microstructures and Thermoelectric Properties of Fe0.92Mn0.08Six Alloys Prepared by Rapid Solidification and Hot Pressing”, J. Appl. Phys., 94(10), 6621–6626 (2003) (Crys. Structure, Electr. Prop., Experimental, Morphology, Thermodyn., 30) Luo, S.J., Yao, K.L., Liu, Z.L., “On the Magnetic Phase Transition in FexMn1–xSi”, J. Magn. Magn. Mater., 265, 167–171 (2003) (Crys. Structure, Theory, Magn. Prop., 18) Zhu, T.-J., Zhao, X.B., Lu, L., “Preparation and Thermoelectric Properties of Melt-Spun Fe2Si5 Based Alloys”, Mater. Sci. Forum, 437–438, 471–474 (2003) (Electr. Prop., Experimental, Phase Relations, 7) Candini, A., Moze, O., Kockelmann, W., Cadogan, J.M., Brueck, E., Tegus, O., “Revised Magnetic Phase Diagram for FexMn5–xSi3 Intermetallics”, J. Appl. Phys., 95(11), 6819–6821 (2004) (Crys. Structure, Experimental, Magn. Prop., Phase Relations, 14) Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diff., 25(4), 346–354 (2004) (Experimental, Phase Diagram, Calculation, Thermodyn., #, 34) Jiang, J.X., Matsugi, K., Sasaki, G., Yanagisawa, O., “Conduction Type Evolution During Eutectoid Decomposition of Mn-Added α-Fe2Si5 Alloy”, Scr. Mater., 53(6), 707–711 (2005) (Crys. Structure, Electr. Prop., Experimental, Phase Relations, 22) Lin, H.C., Lin, K.M., Chen, Y.S., Chu, C.L., “The Wire Electro-Discharge Machining Characteristics of Fe-30Mn-6Si and Fe-30Mn-6Si-5Cr Shape Memory Alloys”, J. Mat. Proc. Tech., 161(3), 435–439 (2005) (Crys. Structure, Experimental, Morphology, Phase Relations, Phys. Prop., 17) Meco, H., Napolitano, R.E., “Liquidus and Solidus Boundaries in the Vicinity of OrderDisorder Transitions in Fe-Si System”, Scr. Mater., 52, 221–226 (2005) (Experimental, Calculation, Phase Diagram, Thermodyn., 30) Stanford, N., Dunne, D., “Martensitic Surface Relief in an Fe-Mn-Si-based Alloy Strained by Bending“, Scr. Mater., 53, 739–744 (2005) (Experimental, Mechan. Prop., 8) Wan, J., Chen, S., “Martensitic Transformation and Shape Memory Effect in Fe-Mn-Si Based Alloys”, Curr. Opin. Solid State Mater. Sci., 9, 303–312 (2005) (Theory, Phase Relations, 56)
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DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
40 [2007Sta]
[2007Wit] [Mas2] [V-C2]
Fe–Mn–Si Stanford, N., Dunne, D.P., Monaghan, B.J., “Austenite Stability in Fe-Mn-Si-Based Shape Memory Alloys”, J. Alloys Compd., 430(1–2), 107–115 (2007) (Calculation, Crys. Structure, Experimental, Mechan. Prop., Morphology, Phase Relations, 25) Witusiewicz, V.T., private communication to MSI, (2007) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_21 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–Ti
1
Iron – Manganese – Titanium Volodymyr Ivanchenko, Tetyana Pryadko
Introduction In spite of the practical application of the Fe-Mn-Ti alloys in the development of high strength austenitic steels and nickel free ferritic steels for cryogenic applications, as well as for the development of new Ti based structural alloys and materials for hydrogen storage, the experimental data on the phase equilibria in this system are incomplete and partly contradictory. The liquidus surface of the Fe-Mn-Ti system has been investigated by [1958Mur1] for the Ti contents app. 20 to 100 at.% and reproduced in [1983Riv, 1988Ray]. The 1000°C isothermal section in the same composition region has been presented by [1958Mur2]. Based on these results [1988Ray] presented the content of the ternary (βTi) field as a function of temperature for the temperature range 600–1000°C. The effect of alloying with Fe on the stability of (αMn), (βMn) and TiMn4 phases has been studied by [1969Pan] and reproduced in [1980Bra]. The phase compositions and variation of the cell parameters of the Laves phase (Ti1–yMny)Fe2 with 0 ≤ y ≤ 0.17 have been presented by [1972Mar]. An important contribution was made in the works of [1979Dew, 1979Kau] by a direct experimental investigation of the system, mainly to establish the boundaries of the TiFe phase but providing evidence with regard to the Ti rich boundary of the Ti(Mn,Fe)2 phase and the general nature of the equilibria. Solid phase equilibria of the Fe-Mn-Ti system at 1000°C near the Fe-Mn border were presented by [1988Ray] according to [1969Pan]. The fields of primary separation in this composition region were constructed by [1988Ray] from the data of [1969Pan]. [1990Har] computed the α-γ equilibrium in the Fe rich alloys at 1150, 1050, 950 and 850°C using the sublattice model for TiFe2 and the magnetic single-lattice model for the solution phases. The isothermal sections in the Fe rich part of the phase diagram at 1150 and 850°C have been presented by [1994Rag] according to [1990Har]. Experimental studies of the phase equilibria, thermodynamics and crystal structures are listed in Table 1. Binary Systems The Fe-Ti, Fe-Mn, and Mn-Ti are taken from [Mas2], [2004Wit], and [Mas2], respectively. There is however, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have been written as L ⇌ (δMn) + (γMn,γFe). Solid Phases A complete solid solution forms between TiFe2 and MnFe2. A substantial solid solubility of Mn in TiFe suggests a general replacement of the Fe atoms by Mn atoms up to a limit of approximately 26 at.% Mn [1979Dew]. [1979Dew] reported about the existence of a phase with the composition 54.8Ti-35.9Mn (at.%) at 1000°C. According to [1979Dew] it could be derived from the Mn-Ti ρ phase or it could be an entirely new ternary phase. Crystal structures of the phases in the Fe-Mn-Ti system are presented in Table 2. Quasibinary Systems The quasibinary section TiMn2-TiFe2 is given in Fig. 1 in accordance with [2007Iva]. The melting and freezing points reported by [2007Iva] are lower than the values presented by [1958Mur1]. They are, however, in good agreement with [1988Ray], who omitted the isothermal curves at 1450 and 1500°C presented by [1958Mur1], because they contradict the accepted melting point of TiFe2. Invariant Equilibria Four-phase invariant equilibria in the Fe-Mn-Ti ternary system are presented in Table 3 in accordance with a tentative liquidus surface projection of the partial Ti-TiMn2-TiFe2 system constructed by [1988Ray] using Landolt-Börnstein New Series IV/11D4
MSIT®
DOI: 10.1007/978-3-540-78644-3_22 # Springer 2008
2
Fe–Mn–Ti
experimental data of [1958Mur1, 1979Dew]. A reaction scheme for the partial Ti-TiMn2-TiFe2 system shown in Fig. 2 is taken from [1988Ray]. Liquidus, Solidus and Solvus Surfaces A liquidus surface of the partial system up to 85 mass% Ti is given in Fig. 3 according to [1958Mur1, 1983Riv, 1988Ray], with some corrections. The freezing temperatures of the alloys located along the TiMn2-TiFe2 section measured by [2007Iva] are accepted. They have been obtained using a more precise technique of measurements and are in a better accordance with the freezing point of TiFe2. On the surface of primary separation of the ternary Laves phase only the 1400 and 1350°C isothermal lines are retained from [1958Mur1] and [1988Ray]. Other isothermal lines are excluded as appearing at much too high temperatures compared with [2007Iva]. The primary crystallization fields near the Fe-Mn border are given in Fig. 3 according to a hypothetical construction of [1988Ray]. Using thermal analysis and metallographic and electrical resistance techniques [1958Mur2] derived the position of the (βTi) border at various temperatures, shown in Fig. 4 after minor adjustments for compliance with the accepted binary diagrams. Isothermal Sections The 1000°C isothermal section has been investigated by [1958Mur2] in the area from app. 20 to 100 mass% Ti (22 to 100 (at.%)). A remarkable feature of this section is the complete solid solution (Fe,Mn)2Ti (λ phase). A substantial solubility of Mn in FeTi is also shown, suggesting a general replacement of the Fe atoms by Mn atoms up to a limit of approximately 20 mass% Mn (19 at.% Mn), with a tolerance for a moderate excess or deficit of Ti [1988Ray]. Another version of the 1000°C isothermal section has been proposed by [1979Dew]. Using electron-probe microanalysis [1979Dew] determined at 1000°C a phase with the composition 51.5Ti-38.7Mn (mass%) (55Ti-9.5Fe-36Mn (at.%)). According to [1979Dew] this phase could be derived from the binary Mn-Ti ρ-phase or it could be a new true ternary phase. In the present evaluation we consider it as a ternary extension of the ρ-phase. The electron-probe microanalysis results [1979Dew] show that the Ti rich boundary of the λ phase field bends sharply towards the Ti corner of the diagram when the Fe content decreases to approximately 20 mass% (14 at.%). There are some differences between the diagram of [1979Dew] and that due to [1958Mur2]. The maximum solid solubility of Fe and Mn in (Ti) is substantially higher (32 at.% solute) in [1979Dew] than that in [1958Mur2]. A two-phase equilibrium λ+(βTi) exists in contrast to the equilibrium between FeTi and ρ as presented by [1958Mur2]. Thus, the nature of the equilibria at 1000°C established by [1958Mur2] and by [1979Dew] differs. The electron-probe microanalysis work of [1979Dew] seem to be very carefully carried out and the proposed equilibria may be preferred. Therefore the results of [1979Dew] are accepted and shown in Fig. 5 for the Ti rich area in the version of [1988Rag] who has added a tentative tie triangle ρ+λ+(βTi). The solid phase equilibria in the Ti poor area at 1000°C have been investigated by [1969Pan] and [1972Mar]. In [1972Mar] no mention was made of the heat treatment, and, as it was pointed out by [1988Ray], it must be assumed that the results refer to the constitution of the “as-prepared” specimens, in which the equilibrium is doubtful. Therefore this work is not considered in the present evaluation. The Ti poor area is shown in Fig. 5 mainly from [1969Pan], but after some modifications described below. The shape and the extension of the λ phase shown by [1969Pan] do not agree with that of [1958Mur2] and [1979Dew] and therefore it has been changed in Fig. 5 to bring it in compliance with the latter works. This change resulted in a small shift of the λ corners of the adjacent tie triangles. The homogeneity range of the R phase along the Mn-Ti axis is reduced to bring it in accord with the accepted binary Mn-Ti system. Also a small violation of the Schreinmaker rule has been eliminated in Fig. 5 at the (αFe) corner of the (αFe) +(αMn)+(γFe) tie triangle. The phase denoted as X in [1969Pan] has been interpreted by [1988Ray] as a ternary extension of the TiMn3 phase. This is accepted here. A tie triangle λ+R+TiMn3 is speculatively added in Fig. 5 (marked with a question mark) to merge the Ti rich and Ti poor parts from different works. In fact, this seems to be the only alternative configuration considering all the phases taking part in equilibria
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in this area. In Fig. 5 all the compositions of the binary phases along the binary axes are adjusted to the binary phase diagrams accepted in the present evaluation. Comparison of the 1000°C section calculated by [1979Dew] with the assessed diagram in Fig. 5 reveals the following discrepancies: ternary extensions of the TiFe, (αMn) and (αδFe) phases are much smaller in the calculated diagram; calculation showed an equilibrium (γFe) + λ, which is not established experimentally; λ phase is calculated without a homogeneity range along the binary axis; the ρ phase is not revealed by the calculation. Accept for the above the agreement is good, in particular in the Ti rich region. The isothermal sections at 1150 and 850°C in the Fe corner are shown in Fig. 6 and Fig. 7 from the computations of [1990Har]. Thermodynamics The effect of Ti on the activity coefficient of Mn in Fe-Mn-Ti alloys was analyzed by [1995Bou] who concludes that the presence of Ti leads to a small increase in the activity coefficient of Mn. In liquid Fe, the interactions between Mn and Ti atoms should be similar to those between Ti atoms. Notes on Materials Properties and Applications The Fe-Mn-Ti system is the basis for the development of precipitation strengthened austenitic [1950Cra, 1979Cha] and nickel-free ferritic steels for cryogenic applications [1971Man, 1978Hwa]. Fe and Mn are well known as beta-stabilizers in the titanium industry. Alloying with Fe and Mn in the amount up to 5 mass% Mn and 2 mass% Fe raises the tensile strength but with sacrifice in ductility [1950Cra]. Influence of Mn on sintering process in the Fe-Ti system has been studied by [1985Kiv], and volume changes in sintering of Fe-Mn-Ti compacts were presented. New composite TiMn2-TiFe were developed by the traditional liquid phase sintering [1985Kiv] and by mechanical alloying [2003Cho] with the aim of developing materials with a hydrogen capacity storage. Experimental data regarding works on material properties are summarized in Table 4. Miscellaneous Magnetic phase diagram of the Ti(MnxFe1–x)2 system with x ≤ 0.6 is presented in Fig. 8 according to [2005Yam]. An icosahedral crystal of the Ti61Fe30Mn7Si2 composition has been grown [1992Zha] by electron beam smelting. Unlike most icosahedral phases, the 5-fold axis is the direction of the most rapid growth.
Table 1.
Investigations of the Fe-Mn-Ti Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1958Mur2]
XRD (X-ray diffraction), optical microscopy, electrical resistance
1000°C for 100 h and quenched in water, < 85 mass% Ti
[1969Pan]
XRD, optical microscopy, chemical analysis
1000°C for 72 h and quenched in water, < 30 at.% Ti
[1972Mar]
EPMA (electron probe microanalysis), X-ray diffraction
As cast, λ, (Ti1–xMnx)Fe2 (x < 0.17)
[1979Dew]
EPMA, SEM (scanning electron microscopy)
1000°C, the whole diagram by calculation and experiment
[2007Iva]
XRD, thermal analysis
1325–1450°C, TiMn2-TiFe2 joint
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Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
(αTi) < 882
hP2 P63/mmc Mg
a = 295.06 c = 468.35
αβδ, (Ti,Mn,Fe) (βTi) 1670–882
cI2 Im 3m W
(δMn)(h3) < 1246 (δFe) 1538–1394 (αFe) < 912
Comments/References
pure Ti at 25°C [Mas2] a = 330.65
pure Ti at 25°C [Mas2]
a = 308.0
pure, at >1138°C [Mas2]
a = 293.15 a = 286.65
pure Fe at 1390°C [V-C2, Mas2] pure Fe at 25°C [Mas2]
a = 386.0
pure Mn, at 1100°C [Mas2]
a = 364.67
pure Fe at 915°C [V-C2, Mas2]
γ, (Mn,Fe) (γMn)(h2) < 1138 (γFe) 1394–912
cF4 Fm 3m Cu
(βMn)(h1) < 1100
cP20 P4132 βMn
a = 631.52
pure Mn, at 727°C [Mas2]
(αMn)(r) < 727
cI58 I 43m αMn
a = 891.26
pure Mn, at 25°C [Mas2]
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, 13 GPa [Mas2]
(ωTi)HP
hP3 P6/mmm ωTi
a = 462.5 c = 281.3
at 25°C [Mas2]
TiFe < 1317
cP2 Pm 3m CsCl
a = 298.8 to 297.7
48–50.2 at.% Fe [Mas2]
λ, Ti(Mn,Fe)2
hP12 P63/mmc MgZn2
Laves Phase a = 483.03–5.7 xFe [2007Iva] c = 792.28–16.3 xFe xFe: mole fraction Fe
TiFe2 < 1424
a = 480.4 to 477.7 c = 784.9 to 780.7
65.2–72.8 at.% Fe, annealed at 1000°C [Mas2]
TiMn2 < 1325
a= 485.7 to 481.8 c= 801.7 to 87.9
64–70 at.% Mn, at 1310°C [1976Sve] (continued)
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Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
ρ, βTiMn ≲ 1200
-
-
52 at.% Mn [1961Wat, 1962Wat] 48–52 at.% Mn, [2004Iva]
ϕ, αTiMn ≲ 950
tP30 P42/mnm CrFe
a = 888 c = 454.2
[Mas2], [V-C2]
TiMn3 < 1250
oP* P212121
a = 481.2 c = 789.5
75.5 at.% Mn, at 1000°C [Mas2]
R, TiMn4 1230–930
hR159 R 3m Co5Cr2Mo3
a = 1100.03 c = 1944.6
[V-C2], [Mas2]. Phase referred to as TiMn5 in [V-C2]. Prototype given as ∼δ(Mo,Ni) in [Mas2]
Table 3.
Invariant Equilibria
Reaction
T [°C]
Type
Phase
Composition (at.%) Fe
Mn
Ti
L + λ ⇌ TiFe + ρ
∼ 1100
U1
L
∼17
∼23
∼60
L ⇌ (βTi) + TiFe + ρ
∼ 1050
E1
L
∼16
∼18
∼66
L + (αδFe) ⇌ (γFe) +γ
∼ >1000
U2
L
∼72
∼12
∼16
L + (γFe) + λ ⇌ TiMn4
?
P1
L
∼28
∼52
∼20
L + λ + TiMn4 ⇌ TiMn3
?
P2
L
∼15
∼64
∼21
L + (γFe) ⇌ TiMn4 + (δMn)
?
U3
L
∼12
∼76
∼12
Table 4.
Investigations of the Fe-Mn-Ti Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1950Cra]
Tensile tests, hardness measurements
Tensile strength, elongation, Vickers hardness
[1971Man]
Tensile tests, hardness measurements
Tensile strength, elongation, Vickers hardness
[1978Hwa]
Tensile tests at cryogenic temperatures, Auger electron spectroscopy, SEM
Tensile strength, elongation
[1979Cha]
Optical microscopy, EDAX, X-ray diffraction, TEM, precipitation strengthening
Annealed 1.5 h at 1150°C, then water quenched and at 700–900°C 20 and 28 mass% Mn, 2 mass% Ti
[1985Kiv]
XRD, specific surface and volume change measurements
1050–1250°C, TiFe0.8Mn0.2, shrinkage under sintering (continued)
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Reference
Method/Experimental Technique
Type of Property
[1992Zha]
XRD, TEM, EDS (electron dispersive X-ray spectroscopy)
Ti61Fe30Mn7Si2, electron beam melting, icosahedral crystal
[2003Cho]
XRD, EDS, hydrogen adsorption by the Sievert’s method
Ball milled, then annealed 1H at 800°C, hydrogen storage
[2005Yam]
XRD, Magnetization, 55Mn nuclear magnetic resonance (NMR)
4.2–300 K, Ti(Fe1–xMnx)2 (x < 0.6), magnetic ordering
Fig. 1. Fe-Mn-Ti. Quasibinary section TiMn2-TiFe2
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Fig. 2. Fe-Mn-Ti. Partial reaction scheme
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Fig. 3. Fe-Mn-Ti. Liquidus surface projection
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Fig. 4. Fe-Mn-Ti. Solid solubility of Fe and Mn in (βTi) at 1000–600°C
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Fig. 5. Fe-Mn-Ti. Isothermal section at 1000°C
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Fig. 6. Fe-Mn-Ti. Computed isothermal section at 1150°C
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Fig. 7. Fe-Mn-Ti. Computed isothermal section at 850°C
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Fig. 8. Fe-Mn-Ti. Magnetic phase diagram of the Ti(Fe1–xMnx)2 system
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14 References [1950Cra] [1958Mur1]
[1958Mur2]
[1961Wat] [1962Wat]
[1969Pan]
[1971Man]
[1972Mar]
[1976Sve]
[1978Hwa]
[1979Cha]
[1979Dew]
[1979Kau]
[1980Bra] [1983Riv]
[1985Kiv]
[1986Cen]
Fe–Mn–Ti
Craighead, C.M., Simmons, G.W., Eastwood, L.W., “Ternary Alloys of Titanium”, Trans. Met. Soc. AIME, 188, 514–538 (1950) (Experimental, Morphology, Mechan. Prop., 1) Murakami, Y., Yukawa, Y., Enjyo, T., “Investigation of the Ti-Rich Ti-Fe-Mn Alloy System. II. On Liquidus Surface of the Ti-Fe-Mn System” (in Japanese), Nippon Kinzoku Gakkai-Shi, 22, 265–269 (1958) (Crys. Structure, Experimental, Morphology, Phase Relations, Phys. Prop., 2) Murakami, Y., Enjyo, T., “Investigation of the Ti-Rich Ti-Fe-Mn Alloy System III. On the Solid Equilibrium Relation in Ti-Fe-Mn System” (in Japanese), Nippon Kinzoku Gakkai Shi, 22, 328–332 (1958) (Experimental, Morphology, Phase Diagram, Phase Relations, Phys. Prop., 3) 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 Diagram, 17) Panigrahy, P.C., Gupta, K.P., “Some Observations on α-Mn, β-Mn, and R Phases in the Mn-Ti-Fe and Mn-Ti-Co Systems”, Trans. Met. Soc. AIME, 245, 1533–1536 (1969) (Crys. Structure, Phase Relations, Experimental, #, 13) Manenc, J., “Study of Hardening of Ferritic and Martensitic Fe-Mn-Ti Alloys” (in French), Mem. Scient. Rev Metallurg., 68(11), 761–769 (1971) (Phase Relations, Experimental, Mechan. Prop., 7) Marandel, J., Schmitt, B., Gantois, M., “Structure and Composition of Certain Laves Phases and Observation of the Fe-Mn-Ti” (in French), Compt. Rend. Acad. Sci. Paris, 275C, 1479–1481 (1972) (Crys. Structure, Phase Relations, Experimental, 3) 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 Metallofiz., 64, 24–29 (1976) (Crys. Structure, Experimental, 7) Hwang, S.K., “The Cryogenical Properties and the Intergranular Brittleness in Fe-12%Mn(0.2%Ti) Ferritic Alloy System”, J. Korean Inst. Met., 16(5), 420–427 (1978) (Phase Realtions, Experimental, Mechan. Prop., 19) Chang, K.M., Morris, J.W., Jr., “Precipitation Strengthened Austenitic Fe-Mn-Ti Alloys”, Metallurg. Trans., A, 10A(9), 1377–1387 (1979) (Crys. Structure, Phase Relations, Experimental, Mechan. Prop., 14) Dew-Hughes, D., Kaufman, L., “Ternary Phase Diagrams of the Manganese-Titanium-Iron and the Aluminum-Titanium-Iron Systems: a Comparison of Computer Calculations with Experiment”, Calphad, 3(3), 175–203 (1979) (Phase Diagram, Thermodyn., Calculation, #, 23) Kaufman, L., Dew-Hughes, D., “Illustration of Ternary Diagram Synthesis Mn-Ti-Fe and AlTi-Fe” in “Calculation of Phase Diagrams and Thermochemistry of Alloy Phases”, Div. Basic Energy Sci., Nat. Sci. Found., Washington, 46–71 (1979) (Phase Diagram, Thermodyn., Calculation, 23) Brandes, E.A., Flint, R.F., “Mn-Fe-Ti” in “Manganese Phase Diagrams”, Manganese Centre, Paris, France, 121 (1980) (Phase Relations, 1) Rivlin, V.G., Raynor, G.V., “Phase Equilibriums in Iron Ternary Alloys 9: Critical Review of Constitution of Ternary Systems Fe-Mn-X (X = Ti, V, Cr, Co, Ni, Cu)”, Int. Met. Rev., 28(1), 23–64 (1983) (Phase Diagram, Phase Relations, Review, 68) Kivalo, L.I., Skorokhod, V.V., “Influence of Manganese on Sintering Process in the Ti-Fe System. 1. Volume Changes in Sintering of Ti-Fe-Mn Compacts”, Sov. Pow. Metall. Met. Ceram., 24(12), 914–917 (Crys. Structure, Phase Relations, Experimental, Interface Phenomena, Kinetics, Transport Phenomena, 70) Cenzual, K., Parthe, E., Waterstrat, R.M., “Zr21Re25, a New Rhombohedral Structure Type Containing 12 Å-Thick Infinite MgZn2 (Laves)-Type Columns”, Acta Crystallogr. C, 42C, 261–266 (1986) (Crys. Structure, Experimental, Phase Relations, 24)
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[1990Har]
[1992Zha]
[1994Rag] [1995Bou]
[2003Cho]
[2004Iva]
[2005Yam]
[2007Iva]
[2007Wit] [Mas2] [V-C2]
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Raynor, G.V., Rivlin, V.G., “Fe-Mn-Ti” in “Phase Diagrams Ternary Iron Alloys”, Inst. Metals, London, 4, 378–388 (1988) (Phase Diagram, Phase Relations, Review, Assessment, *, 7) Hari Kumar, K.C., Raghavan, V., “BCC-FCC Equilibrium in Ternary Iron Alloys-IV”, J. Alloy Phase Diagram, 6(1), 31–49 (1990) (Phase Diagram, Thermodyn., Calculation, Review, #) as quoted by [1994Rag] Zhang, X., Kelton, K.F., “Vertex Growth of a Ti-Mn-Fe Icosahedral Phase”, Scr. Metall. Mater., 26(3), 393–398 (1992) (Crys. Structure, Experimental, Morphology, Phase Relations, 23) Raghavan, V., “Fe-Mn-Ti (Iron-Manganese-Titanium)”, J. Phase Equilib., 15(6), 620–621 (1994) (Phase Diagram, Review, 3) Bouchard, D., Bale, C.W., “Simultaneous Optimization of Thermochemical Data for Liquid Iron Alloys Containing C, N, Ti, Si, Mn, S, and P”, Metall. Mater. Trans. B, 26B, 467–484 (1995) (Phase Relations, Theory, Thermodyn., 85) Choi, M.J., Hong, H.S., Lee, K.S., “Electrochemical Characteristics of the Composite Metal Hydride of TiFe and TiMn2 Synthesized by Mechanical Alloying”, J. Alloys Compd., 358, 306–311 (2003) (Crys. Structure, Electrochem., Experimental, 15) 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, Phase Diagram, 14) Yamada, Y., Nakamura, K., Kitagawa, K., Obara, G., Nakamura, T., “Magnetic Properties of C14 Laves Phase Ti(Fe1–xTx)2 with T = Mn, Co and Ni (x<0:6)”, J. Magn. Magn. Mater., 285 (1–2), 28–38 (2005) (Crys. Structure, Experimental, Magn. Prop., Optical Prop., 10) Ivanchenko, V., Dekhtyarenko, V., Kosorukova, T., “Phase Equilibria on the TiMn2-TiFe2 Polythermal Section”, Abstracts of X International Conference on Crystal Chemistry of Intermetallic Compounds, Lviv, Ukraine, 17–20 September, 48 (2007) (Crys. Structure, Phase Diagram, Experimental, Abstract, #, 2) Witusiewicz, V.T., private communication to MSI, (2007) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Fe–Mn–V
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Iron – Manganese – Vanadium Lesley Cornish, Andy Watson
Introduction The Fe-Mn-V system is important for steels; Mn is a solid solution strengthener and vanadium is useful to combine with carbon or/and nitrogen to form carbides, nitrides and/or carbonitrides, which can be used to pin grain boundaries, and hence produce steels with a finer grain size. The control of grain size is vital in rolling, and knowledge of the effects of additions to steels is vital for improving productivity. Vanadium is beneficial in promoting fine, mainly nitride, precipitates which give dispersion strengthening. Additionally, since both the Fe-V and Mn-V systems form the σ phase, which imparts deleterious mechanical properties, knowledge of the stability range of this phase is important so that its precipitation can be avoided in commercial steels. There have been few experimental studies of the ternary system. The σ phase and its stability in the Fe-Mn-V system was first studied by [1957Dar]. Electrolytic grade Fe and Mn were vacuum induction melted with V chips in alumina crucibles and portions of the cast material were homogenized at 1000°C. XRD and metallographic analysis revealed a continuous series of solid solutions of the σ phase across the system. An early investigation of the phase relationships of the system was made by [1973Ish] who studied alloys containing up to 8 at.% Mn and 6 at.% V at 900, 1000, 1100 and 1200°C to derive the boundaries of the ‘γ-loop’ and to obtain stabilizing parameters. The ordering effects of the δ’ CsCl phase were investigated by [1975Hag] using DTA and XRD. Four partial isothermal sections were determined by [1992Bud]. Compositions for study were selected lying along three lines of constant Mn:V ratio (3:1, 1:1 and 1:3) across the phase diagram. The samples were prepared by arc-melting on a water-cooled copper hearth under argon using Armco iron (∼99% purity), and electrolytic grade Mn and V, with a total purity in excess of 99.95%. The samples were analyzed using metallography, X-ray diffraction, chemical analysis and hardness tests. The system was assessed by [1991Hua1] using the Calphad technique. The calculated phase diagrams agree well with the experimental work of [1957Dar]. Ordering of the δ’ phase was not included in the modeling because of the difficulty in extending the binary phase description into the ternary system without there being any experimental data available for the optimization. Substitutional solution models were used for all except for the σ phase. The magnetic terms for the Mn-V system were set to zero owing to a lack of data, and similarly, no ternary interactions were included in the modeling. The σ phase was modeled with three sublattices: (Fe,Mn)8:V4:(Fe,Mn,V)18, and some of the parameters were estimated from experimental data [1978Hil]. The system has been reviewed by [1980Bra, 1983Riv, 1988Ray, 1994Rag]. The experimental works are summarized in Table 1. Binary Systems The binary Fe-Mn system has been assessed by [1989Hua] and the calculated diagram is very similar in appearance to that given in [Mas2]. The main difference between the two is a predicted eutectoid decomposition of the γ phase to (αFe) and (αMn) occurring at 247.9°C (521 K). The system has been reassessed more recently, however, incorporating new experimental information on thermodynamic properties [2004Wit]. There are slight differences in the positions of the phase boundaries, and the agreement with experimental Fe rich phase equilibria is improved over the earlier assessment. The γ phase exists as a continuous series of solid solutions right across the phase diagram. There is however, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have been written as L ⇌ (δMn) + (γMn,γFe). The phase diagram of [2004Wit] is accepted here. The Fe-V system has been assessed by [1991Hua2] but [2006Oka] has reported recent work by [2005Ust] who studied the extent of the σ phase in the binary Fe-V system using XRD and electron microscopy. They Landolt-Börnstein New Series IV/11D4
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suggest that the σ phase decomposes below about 650°C accompanied by phase separation of the δ (bcc) phase. However, in the phase diagram shown by [2006Oka], there would seem to be a narrow strip of single phase bcc between the σ phase and the region of phase separation. As pointed out by [2006Oka], this would seem to be unlikely, the sigma phase most probably decomposing eutectoidally to give Fe rich and V rich bcc phases, much like in the Cr-Fe system. As this remains uncertain, the feature has been ignored in the present work. The version of the phase diagram give in [Mas2] is accepted here. The binary Mn-V system is characterized by a wide series of solid solutions based on (V) and (δMn). The σ phase has a homogeneity range of approximately 8 at.% at 600°C, which extends to about 18 at.% at 1020°C, but as pointed out by [1982Smi], the width of the phase field is not well defined. The system was assessed by [1991Hua1] resulting in a phase diagram that is somewhat different to the version given in [Mas2], particularly with respect to the congruent transformation of the σ phase. The version given by [1982Smi] (reproduced in [Mas2]) shows a transformation from the σ phase and eutectoid decomposition of (V,δMn), the shapes of which are unlikely. The assessment of [1991Hua1] gives a congruent transformation of the σ phase that is more symmetrical and occurring at a higher temperature resulting in a more reasonable looking eutectoid reaction. The result is in better agreement with the experimental information in this region of the diagram. This version of the phase diagram is also preferred by [1992Oka] and is accepted here. The ordered phase indicated in the diagram of [Mas2] was not modeled in the work of [1991Hua1]. Solid Phases No ternary phases have been found in this system. The major feature of the phase diagram is the presence of a continuous σ phase field stretching across the diagram between the Fe-V and the Mn-V edge binary systems [1957Dar]. The lattice parameters of the σ phase in the Fe-V and Mn-V systems are very close to each other, so it should be no surprise that there is a continuous series of solid solutions between them. The binary terminal solid solutions dissolve appreciable amounts of the third element resulting in phase fields that penetrate deep into the ternary system. [1975Hag] demonstrated the presence of the δ’ CsCl type superlattice and also that the order-disorder transition temperature, Tc, of the δ’ VMn superlattice was increased by the addition of Fe. An estimated value for the Tc of the FeV superlattice, which is not measurable because of the intervention of σ, was extrapolated to be between 850 and 880°C. Table 2 gives crystallographic details of the solid phases. Liquidus, Solidus and Solvus Surfaces According to the work of [1992Bud], a study of the liquidus surface has been undertaken by [1990Ban], but unfortunately, these results are not available. In any case, it is unlikely that the liquidus surface would show any distinguishing features as there are only 2 binary invariant reactions taking place and both lie in the FeMn system. It is expected that a monovariant line would extend from the binary L + (δFe) ⇌ (γMn,γFe) peritectic into the ternary system, eventually encountering the lower temperature L ⇌ (δMn) + (γMn,γFe) eutectic reaction in the same binary system. The nature of this monovariant reaction would therefore switch from peritectic to eutectic at some unknown composition in the ternary system. Using the thermodynamic model parameters from [1991Hua1] and the updated parameters for the Fe-Mn system from [2004Wit] a liquidus surface has been calculated, and is shown in Fig. 1. It should be stressed, however, that there has been no experimental confirmation for this liquidus surface. Isothermal Sections [1957Dar] made one of the earliest studies of the solid phase equilibria in the system. Using manganese and iron of electrolytic grade, and vanadium chips with an impurity content of 0.25%, samples were prepared by vacuum induction melting using recrystallized alumina crucibles. Weight losses of the ingots were less that 1 mass%, and the compositions were obtained by chemical analysis. XRD was carried out on crushed samples after they had been homogenized for “a prolonged period” at 1000°C in evacuated and sealed fused quartz tubes, followed by quenching in water by breaking the submerged tubes. Identification of the σ phase was achieved using metallography, with XRD being used for confirmation. Since the σ-containing specimens exhibited a certain degree of cracking, this was an aid to interpretation, however the other phases DOI: 10.1007/978-3-540-78644-3_23 # Springer 2008
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in the specimens were not identified. This enabled the extent of the σ phase field to be determined. [1973Ish] studied equilibria in the Fe rich corner of the phase diagram at 900, 1000, 1100 and 1200°C (up to 8 at.% Mn and 6 at.% V) using X-ray microanalysis of annealed samples. The phase boundaries were found to agree well with the calculated sections of [1991Hua1], which were based mainly on the extrapolation of the thermodynamic descriptions of the binary systems into the ternary, but with some adjustment of the model parameters of the σ phase to give agreement with the experimental data of [1957Dar]. The experimental data of [1973Ish] are compared with the calculated phase boundaries for 900, 1000, 1100 and 1200°C in Figs. 2 to 5 indicating the level of agreement. Full isothermal sections calculated for 800, 1000 and 1200°C are given in Figs. 6 to 8. Alterations have been made where necessary to ensure agreement with the accepted binary phase diagrams. The binary Mn-V system given by [Mas2] shows the existence of an ordered CsCl type phase, denoted δ’, although its exact location in the phase diagram is not clear, as indicated by dotted lines in the Figure. [1975Hag] used DTA and XRD to study the order/disorder transition in ternary alloys. They found that the addition of Fe raised the critical temperature of the ordering. Using their experimental data, they were able to derive thermodynamic interaction parameters for the order/disorder reaction and produce a set of isotherms, which are reproduced in Fig. 9. The results were also shown in [1980Bra]. Partial isothermal sections (from 40 mass% Fe) were determined for 600, 800, 1000 and 1200°C by [1992Bud]. Armco Fe and electrolytic Mn and V were arc-melted on a water-cooled copper hearth under argon. Alloys were annealed at 1200, 1000, 800 and 600°C for times between 24 and 250 h depending on the heat treatment temperature. The alloys were investigated by XRD, microstructural and chemical analysis and hardness measurement. However, in the partial isothermal sections shown, there is no evidence of the σ phase, which is known to exist across the diagram, certainly at the lower temperatures. Temperature – Composition Sections An experimental section for xMn/xV = 1 at high Fe contents was determined by [1973Ish] which showed a closed γ loop. Their own thermodynamic calculations gave a slightly smaller loop and later calculations by [1991Hua1] did not show any closure at all. Introducing a small ternary interaction parameter for the bcc δαffl phase (–10kJ·mol–1) enabled a closed loop to be produced, but then it was found that the fit of the calculated phase boundaries to the experimental data suffered. However, [1991Hua1] suggested that observing small amounts of the second phase would be experimentally challenging as some of the observed compositions were essentially on the δ/δ+γ phase boundary. The experimentally determined γ loop is depicted in Fig. 10. A vertical section was produced by [1992Bud] at a Mn content of 10 mass%, but as this does not agree with the phase fields or the peritectic temperature shown in the accepted Fe-Mn phase diagram, and therefore it is not reproduced here. Thermodynamics From the experimentally determined δ/γ phase boundaries, [1973Ish] derived a series of thermodynamic parameters and showed that for low alloy contents, they could be used to derive the phase relationships. For higher alloying contents however, there were no reliable relationships. From their order-disorder transition studies, [1975Hag] derived interaction parameters for the δ phase Fe-Mn (200 ± 300), Fe-V (–4520 ± 60) and Mn-V (–4440). Using Hamiltonian cluster models, [1977Bra] calculated the extent of the σ phase in the ternary system but found that the fit to the available experimental data was poor. It was acknowledged that there may be issues with the simplicity of the cluster model employed, and they also stated that data was required for the different bandwidths in the metal, charge transfer, magnetic and temperature effects, as well as the need to include of all the phases in the calculations. [1987Bal] studied the effect of vanadium on the activity of the manganese in dilute Fe-alloys (Fe3.14 mass% Mn) at 1873 K (1600°C) using a continuous weighing technique. They found that the evaporation rate constant for Mn decreased from 11.5·10–3 to 7·10–3 cm·s–1 leading to a pairwise interaction coefficient of εVMn of 1.55. Landolt-Börnstein New Series IV/11D4
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Fe–Mn–V
Notes on Materials Properties and Applications Since the Fe-Mn-V system is important for steels, mechanical property measurement such as hardness testing has been used to identify useful steel compositions. As well as avoiding the σ phase, which is detrimental to the mechanical properties of steel thus limiting the amounts of V and Mn in a given material, alloys with more than 50 mass% (Mn + V) were also found to be embrittled by the precipitation of the δ phase (αFe) [1992Bud]. A beneficial martensitic transformation occurred in alloys where (Mn + V) was between 10 and 12 mass%. [1981Kho] studied the elasticity modulus of Fe-Mn dispersion hardened alloys and the effects of V additions as an aging accelerator. The maximum strength of a Fe-Mn Elinvar alloy (Fe-28.8% Mn-1.9% V-0.5% C) was achieved after aging at 650–750°C Miscellaneous In a study involving the effects of a variety microalloying additions to steel, [2004Li] showed the beneficial affects of both V and Mn in controlling the grain size of thin slab cast product. The addition of P promotes the stability of amorphous Fe-Mn alloys. The processes occurring in amorphous Fe-P-Mn-V alloys at elevated temperatures, up to the crystallization temperature, were studied by [2001Vav, 2004Vav]. Results from internal friction measurements and XRD were used to evaluate the activation energy of defect migration and the particle size of the crystalline phases precipitating in the alloys during annealing. Increasing the vanadium content of amorphous Fe-P-Mn-V alloys increases the energy of defect migration reducing the particle size of the precipitating phases. The dissolution of Mn and V was found to inhibit the growth of Fe3P particles during annealing. Table 1.
Investigations of the Fe-Mn-V Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1957Dar]
Metallography, XRD
Isothermal section at 1000°C, 10–60 mass% V, extent of σ phase stability
[1973Ish]
Thermodynamic analysis, X-ray microanalysis
Partial isothermal sections at 800°C–1200°C, up to 8 at.% Mn and 6 at.% V
[1975Hag]
XRD, DTA
Partial isothermal sections at 500°C–900°C, 50 at.% V with 5, 10, 15, 20, 25, 30, 40 at.% Fe, as well as 50V:50Fe, measurement of Tc related to ordering of bcc phase
[1977Bra]
Cluster model calculation
1000°C, extent of σ phase stability
[1991Hua1]
Thermodynamic modeling by the Calphad technique
Complete system
[1992Bud]
Metallography, XRD, chemical analysis, hardness and microhardness
Isothermal sections at 600°C, 800°C, 1000°C and 1200°C, from 40 mass% Fe, (Mn, Fe) + (αFe)
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5
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
δ, (δαFe, δMn,V)
cI2 Im 3m W
(δFe) 1538–1394 (V) < 1910 (δMn) 1246–1138 (αFe) < 912 γ, (γMn,γFe) (γFe) 1394–912 (γMn) 1138–1100
cF4 Fm 3m Cu
Lattice Parameters [pm]
Comments/References
a = 293.15
pure Fe at 1390°C [V-C2, Mas2]
a = 302.4
pure V [Mas2]
a = 308.0
pure Mn [Mas2]
a = 286.65
pure Fe at 25°C [Mas2]
a = 364.67
pure Fe at 915°C [V-C2, Mas2]
a = 386.0
pure Mn [Mas2]
δ’
cP2 Pm 3m CsCl
a = 291.0
VMn. Metastable in Fe-V system
(αMn)
cI58 I 43m αMn
a = 891.26
[Mas2]
(βMn)
cP20 P 4132 βMn
a = 631.52
[Mas2]
σ
tP30 P42/mnm σCrFe
σVFe < 1252 σVMn < 1058
a = 896.5 c = 463.3 a = 898.5 c = 462.1
[Mas2] [V-C2] [1991Hua1, V-C2]
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, 13 GPa [Mas2]
Martensite
tI4 I4/mmm
-
[Mas2]
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Fig. 1. Fe-Mn-V.
Fe–Mn–V
Calculated liquidus surface projection
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Fig. 2. Fe-Mn-V.
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Calculated boundaries for (δFe)/γ(Fe) at high Fe contents, at 900°C
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Fig. 3. Fe-Mn-V.
Fe–Mn–V
Calculated boundaries for (δFe)/γ(Fe) at high Fe contents, at 1000°C
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Fig. 4. Fe-Mn-V.
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Calculated boundaries for (δFe)/γ(Fe) at high Fe contents, at 1100°C
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Fig. 5. Fe-Mn-V.
Fe–Mn–V
Calculated boundaries for (δFe)/γ(Fe) at high Fe contents, at 1200°C
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Fig. 6. Fe-Mn-V.
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Calculated isothermal section at 800°C
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Fig. 7. Fe-Mn-V.
Fe–Mn–V
Calculated isothermal section at 1000°C
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Fig. 8. Fe-Mn-V.
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Calculated isothermal section at 1200°C
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Fig. 9. Fe-Mn-V.
Fe–Mn–V
Calculated order-disorder transition temperatures
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Fig. 10. Fe-Mn-V.
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Vertical section at xMn/xV = 1 at high Fe content
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16 References [1957Dar] [1973Ish]
[1975Hag]
[1977Bra] [1978Hil] [1980Bra] [1981Kho]
[1982Smi] [1983Riv]
[1987Bal]
[1988Ray] [1989Hua] [1990Ban]
[1991Hua1] [1991Hua2] [1992Bud]
[1992Oka] [1994Rag] [2001Vav]
[2004Li]
Fe–Mn–V
Darby, J.B., Beck, P.A., “σ Phase in Certain Ternary Systems with Vanadium”, Trans. A.I.M. E., 209, 69–72 (1957) (Phase Diagram, Experimental, 8, #) Ishida, K., Shibuya, K. and Nishizawa, T., “γ -Loops in Fe-Mn-V, Fe-Mn-Mo and Fe-Mn-Si Systems” (in Japanese), J. Jpn. Inst. Met., 37(2), 1305–1312 (1973) (Phase Diagram, Thermodyn., Calculation, Experimental, 22) Hagiwara, M., Seki, J., Suzuki, T., “CsCl type Superlattice Formation in V-Fe-Mn, Fe-V-Cr and V-Mn-Cr Ternary Alloys” (in Japanese), J. Jpn. Inst. Met., 39(4), 402–408 (1975) (Thermodyn., Experimental, Phys. Prop., 27) Brauwers, M, “Occurrence of the σ Phase Computed from a Cluster Model”, J. Phys. F: Met. Phys., 7(6), 921–927 (1977) (Calculation, 17) Hillert, M., Jarl, M., “A model for alloying in ferromagnetic metals”, Calphad, 2(3), 227–238 (1978) cited by [1991Hua1] Brandes, E.A., Flint, R.F., “Mn-Fe-V” in “Manganese Phase Diagrams”, Manganese Centre, Paris, France, 122 (1980) (Phase Relations, Review, 1) Khomenko, O.A., Khil'kevich, I.F., “Characteristics of `Nonmagnetic' Dispersion-hardened Fe-Mn-Elinvar”, Met. Sci. Heat Treat., 23(3–4), 211–214 (1981) (Phase Relations, Mechan. Prop., 8) Smith, J.F., Carlson, O.N., “The Mn-V (Manganese-Vanadium) System”, Bull. Alloy Phase Diagrams, 2(4), 469–478 (1982) (Crys. Structure, Phase Diagram, Thermodyn., Review, 19) Rivlin, V.G., Raynor, G.V., “Phase Equilibriums in Iron Ternary Alloys 9: Critical Review of Constitution of Ternary Systems Fe-Mn-X (X = Ti, V, Cr, Co, Ni, Cu)”, Int. Met. Rev., 28(1), 23–64 (1983) (Phase Diagram, Phase Relations, Review, 68) Balkovoy, Yu.V., Aleev, R.A., Manohin, A.A., “Influence of Carbon, Silicon and Vanadium on the Activity of the Manganese in Iron” (in Russian), Izv. Vyss. Uchebn. Zaved., Chern. Metall., 9, 128–129 (1987) (Experimental, Thermodynamics, 1) Raynor, G.V., Rivlin, V.G., “Fe-Mn-V” in “Phase Equilibria in Iron Ternary Alloys”, Inst. Metals, London, 388–392 (1988) (Review, #,2) Huang, W., “Assessment of the Fe-Mn System”, Calphad, 13(3), 243–252 (1989) (Phase Diagram, Thermodyn., Assessment, 35) Bannykh, O.A., Budberg, A.P., Potemkin, A.Ya., “Special Features of the Interaction of the Components in Ternary Fe-Mn-V” in “Problems in Establishing the Optimum Structure of Aviation Materials”, Sergo Ordzhonikidze Aviation Inst, Moscow (1990) cited by [1992Bud]. Huang, W., “A Thermodynamic Analysis of the Mn-V and Fe-Mn-V Systems”, Calphad, 15 (2), 195–208 (1991) (Phase Diagram, Assessment, Calculation, #, *, 23) Huang, W., “A Thermodynamic Evaluation of the Fe-V-C System”, Z. Metallkd., 82(5), 391–401 (1991) (Phase Diagram, Thermodyn., Assessment, Calculation, 54) Budberg, A., Potemkin, A.Ya., “Phase Structure of Alloys of the Iron - Manganese Vanadium System over the Temperature Range 600–1200°C”, Isv. Akad. Nauk SSR, Met., 5, 133–136 (1992) (Experimental, Phys. Prop., 1 #) Okamoto, H., “Mn-V (Manganese-Vanadium)”, J. Phase. Equilib., 13(5), 575 (1992) (Crys. Structure, Phase Diagram, Review, 8) Raghavan, V., “Fe-Mn-V (Iron-Manganese-Vanadium)”, J. Phase Equlib., 15(6), 621–622 (1994) (Review, 2) Vavilova, V.V., Kovneristyi, Y.K., Palii, N.A., “Physicochemical Properties of Amorphous and Crystalline Fe-P-Mn and Fe-P-Mn-V Alloys”, Inorg. Mater., 37(2), 166–169 (2001) (Phase Relations, Experimental, Thermodyn., 5) Li, Z., Wilson, J.A., Crowther, D.N., Mitchell, P.S., Craven, A.J., Baker, T.N., “The Effects of Vanadium, Niobium, Titanium and Zirconium on the Microstructure and Mechanical Properties of Thin Slab Cast Steels”, ISIJ Int., 44(6), 1093–1102 (2004) (Crys. Structure, Experimental, Morphology, 35)
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[2004Wit]
[2005Ust]
[2006Oka] [2007Wit] [Mas2] [V-C2]
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Vavilova, V.V., Vygovskii, D.A., Kalinin, Y.E., Kovneristyi, Y.K., Palii, N.A., “Inelastic Properties of Amorphous and Nanocrystalline Fe-P-Mn-V Alloys”, Inorg. Mater., 40(8), 815–821 (2004) (Phase Relations, Experimental, Thermodyn., 13). Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diff., 25(4), 346–354 (2004) (Phase Diagram, Thermodyn., Assessment, Calculation, Experimental, 34) Ustinovshikov, Y., Pushkarev, B., Sapegina, I., “Phase Transformations in Alloys of the Fe-V System”, J. Alloys Compd., 398, 133–138 (2005) (Crys. Structure, Phase Diagram, Experimental, 9) Okamoto, H., “Fe-V (Iron-Vanadium)”, J. Phase Equilib. Diff., 27(5), 542 (2006) (Review, Phase Diagram, 3) Witusiewicz, V.T., private communication to MSI, (2007) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson's Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Fe–Mn–Zr
1
Iron – Manganese – Zirconium Nathalie Lebrun, Pierre Perrot
Introduction The main interest in the Fe-Mn-Zr alloys rests in its use as hydrogen storage material. The experimental works done on this system deal mainly with its hydrogen storage capacity: characteristics hydrogen content vs imposed hydrogen pressures, thermodynamics related to the hydrogen absorption. The only part of the diagram investigated is the vertical section ZrMn2-ZrFe2. The entire ternary diagram has never been investigated and no Calphad assessment is known. Main experimental works are reported in Table 1. Binary Systems Using new experimental thermodynamic data, the Fe-Mn system has been recently reassessed and the Calphad description was updated by [2004Wit]. This new reevaluation of the phase equilibria leads to consistently better fits to the available experimental data. Consequently, the Fe-Mn system is accepted from [2004Wit]. There is however, a typographical error in [2004Wit] in that the Mn rich invariant reaction involving the liquid phase is given as a peritectic type reaction in the table of invariants. This reaction should be denoted as eutectic, as confirmed by [2007Wit], and have been written as L ↾ (δMn) + (γMn,γFe). The Fe-Zr binary system is accepted from [2002Ste]. The Mn-Zr system is accepted from [1999Sch]. It presents a non stoichiometry for the ZrMn2 compound which is lower than that reported by [Mas2]. Solid Phases All the data on unary, binary and ternary phases are indicated in Table 2. The Laves phases ZrMn2 and ZrFe2 give a solid solution Zr(Mn1–xFex)2 whose structure varies from C14 (x < 0.6) to C15 (x > 0.8) [1968Kan]. For 0.6 < x < 0.8 [1972Pet], reproduced by [1992Rag], proposes a two-phase domain C14 + C15. Actually, it is more probable that the solid solution presents a C36 structure (hP24, P63/ mmc, MgNi2 type), a transition already observed in ZrFe2 based solid solutions [1973Sve] such as (Nb1–xZrx)Fe2. A strong argument comes from the fact that the iron magnetic moment of the solid solution grows monotonically from 0 to 1.2 μB when x varies from 0.5 to 1 [1968Kan]. X-ray investigations carried out by [1971Kan] confirm the existence of the C36 phase. The pattern diffraction of Zr(Mn0.3Fe0.7)2 is similar to that of Nb0.4Zr0.6Fe2 and may be indexed as a six hexagonal layer structure. It is not clear whether the transitions C14-C36 and C36-C15 are of the first or second order. The crystal parameters of the solid solutions Zr(Mn1–xFex)2 are plotted in Fig. 1 according to the data of [1999Sch] and [2004Li]. As the C14-C36 and C36-C15 transitions have not been formally recognized, the domain between the C14 and C15 regions is indicted by a question mark. Quasibinary Systems [1972Pet] presents a phase diagram through the quasibinary ZrMn2-ZrFe2 section, which is questionable, because he proposes for ZrMn2 a melting point at 1570°C, which is 120°C higher than the accepted value and a two-phase domain around 20 mol% ZrMn2, which contradicts the observations of [1971Kan]. Based on the phase diagram proposed by [1972Pet], it is probable that the liquidus and the solidus lines of the ZrMn2-ZrFe2 phase diagram present a cigar like shape between 1673 and 1450°C, the respective melting points of ZrFe2 and ZrMn2. The phase equilibria are shown in Fig. 2. The liquidus line, mainly from the DTA experiment carried out by [1972Pet], has been modified to take into account the accepted melting point of ZrMn2. Thermodynamics The molar partial enthalpy of dissolution of H2 in ZrMnFeHy (1 < y < 2.5) was measured by [2000Pra] from the isobaric curves and found at ΔdHH = +9.25 ± 0.25 kJ·(mol⊊H)–1 and its molar partial entropy of Landolt-Börnstein New Series IV/11D4
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Fe–Mn–Zr
dissolution at ΔdSH = +26 ± 1.0 J·(mol⊊H)–1·K–1. The partial replacement of Fe by Co gives, for ZrMnFe0.75Co0.25Hy (1 < y < 1.5) ΔdHH = – 9.3 ± 0.5 kJ·(mol⊊H)–1 and ΔdSH = – 36.8 ± 0.6 J·(mol⊊H)–1·K–1. This effect is due to the contraction of the lattice induced by the substitution of Fe by Co [2000Pra]. Notes on Materials Properties and Applications Zr based Laves phases such as the Zr(Fe,Mn)2 solid solution have attracted much attention as hydrogen storage materials [2004Li]. ZnMn2 is a representative member of the series, which absorbs hydrogen up to the composition of ZrMn2H4 [1984Pou]. However, its hydrides are too stable so that further partial substitutions (Zr by Ti and Mn by Fe) are needed to provide materials stable within a wide range of hydrogen pressures. For instance the ZrMn2.7 compound presents, on the hydrogen pressure - hydrogen content curve, a plateau at 0.03 MPa of hydrogen pressure for 20°C and at 0.5 MPa of hydrogen pressure for 150°C. These plateaus are raised when substituting Mn by Fe [1986Uch], by Fe and Co [2000Pra, 2000Ram, 2002Man2] or by Fe and Ni [2002Man1] in the solid solution. For the ZrMn1.11Fe1.22 compound, the respective positions of the plateaus are: 1 MPa of hydrogen pressure at 20°C and at 5 MPa of hydrogen pressure at 150°C. The crystal structure remains unchanged upon hydrogenation [2003Dor], but the relation between the volume change induced by hydrogenation and the hydrogen uptake is not linear. The reducing power of Zr combined with the hydrogen storage capacity of the ternary alloy have been used to reduce HTO (tritiated water) to HT (tritiated hydrogen) at 400–600°C with a conversion efficiency of 90 % for residence time of the order of the second [1995Ghe, 1997Ghe]. The conversion rate is greater than 6 μmol·min–1·(g of alloy)–1. The oxygen to alloy mole ratio may be as high as 2.6 without significant loss of efficiency. The HT/HTO conversion efficiency may go up to 99% with a working temperature of 800–900°C. The main experimental investigations on material properties are gathered in Table 3. Miscellaneous The influence of Mn on the magnetic properties of (Fe,Zr) alloys have been widely investigated [1987Nic, 1994Sri, 1997She, 2001Per2, 2006Per]. The amorphous Fe90Zr10 alloy presents a ferromagnetic transition at 230 K followed by the loss of long range order into a spin glass like state at 12 K [1987Nic]. The substitution of Fe by Mn up to 10 at.% Mn (Fe90–xMnxZr10, with x < 10) decreases the spin freezing temperature. The Curie temperature decreases from 230 K (for x = 0) down to 150 K (for x = 12) and the spin freezing temperature increases from 35 K (for x = 0) up to 65 K (for x = 12) [1997She, 2001Per3, 2002Per]. Hydrogenation enhances the magnetization of the materials [1987Zha], probably as a consequence of the expansion of the lattice. By lowering temperature, the amorphous alloy Fe80Mn10Zr10 presents a paramagnetic-ferromagnetic transition which depends on the applied field, then, by further cooling, a transition, observed at 40 K from the ferromagnetic state to a new state referred as a “reentrant spin glass phase” [2001Per1, 2002Per] which is characterized by the vanishing of the ferromagnetic order. Some features of this new phase look like spin glass state, but its true nature is not well established. The substitution of Mn by B in the amorphous Fe82Mn82–xBxZr10 (x < 8) increases the Curie temperature (368 K for x = 8) [2005Uly]. A paramagnetic state in the amorphous matrix is observed together with a superparamagnetic state due to nanoparticles embedded in the matrix.
Table 1.
Investigations of the Fe-Mn-Zr Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1972Pet]
XRD, differential thermal analysis
900–1675°C, Zr(FexMn1–x)2 (0 < x < 1), quasibinary section
[1984Pou]
XRD, hydrogen absorption, hydrogen intake - pressure characteristics
23–190°C, ZrMn2.8, ZrMn2Fe0.8, ZrMn2Fe1.2, < 5 MPa H2, partial quantities of H2 dissolution (continued)
DOI: 10.1007/978-3-540-78644-3_24 # Springer 2008
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Fe–Mn–Zr
3
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[2000Pra]
XRD, hydrogen absorption, hydrogen intake - pressure characteristics
30–100°C, ZrMnFe and ZrMnFe0.7Co0.3, < 3 MPa H2, partial quantities of H2 dissolution
[2000Ram]
XRD, hydrogen absorption, hydrogen intake - pressure characteristics
30–100°C, ZrMnFe1–xCox (x = 0.2, 0.3, 0.5), < 0.1 MPa H2, partial quantities of H2 dissolution
[2002Man1]
XRD, hydrogen absorption, hydrogen intake - pressure characteristics
30–100°C, ZrMnFe1–xNix (x < 0.6), < 4 MPa H2
[2002Man2]
XRD, hydrogen absorption, hydrogen intake - pressure characteristics
30–100°C, ZrMnFe1–xCox (x < 0.6), < 4 MPa H2
[2003Dor]
XRD, volume change upon hydrogenation
25°C, Zr(FexMn1–x)2 (x < 1)
[2004Li]
XRD, calorimetry, hydrogen absorption
25°C, Zr(FexMn1–x)2 (x < 0.8), < 0.1 MPa H2
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
αδ(Fe,Mn)
cI2 Im 3m W
(δFe) 1538–1394 (δMn) 1246–1138
Comments/References
a = 293.15
pure Fe at 1390°C [V-C2, Mas2] Dissolves 10 at.% Mn at 1474.5°C [2004Wit] [Mas2] Dissolves up to 12.2 at.% Fe at 1236.6°C [2004Wit] and 1.25 at.% Zr at 1155°C [1999Sch] pure Fe at 25°C [Mas2, V-C2]. Dissolves 5 at.% Mn at 527°C [2004Wit]
a = 308.0
(αFe) (Ferrite) < 912 (εFe)
Lattice Parameters [pm]
a = 286.65
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, > 13 GPa [Mas2]
(continued)
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DOI: 10.1007/978-3-540-78644-3_24 # Springer 2008
4
Fe–Mn–Zr
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
γ(Fe,Mn)
cF4 Fm 3m Cu
(γFe) (Austenite) 1394–912 (γMn) 1138–1100
Lattice Parameters [pm]
Comments/References
a = 364.67
pure Fe at 915°C [V-C2, Mas2] Complete solubility with (γMn) [2004Wit] [Mas2] complete solubility with (γFe) [2004Wit]. Dissolves 0.5 at.% Zr at 1120°C [1999Sch]
a = 386.0
(βMn) 1100–727
cP20 P4132 βMn
a = 631.52
[Mas2]. Dissolves up to 32.9 at.% Fe [2004Wit] and 1.7 at.% Zr at 1140°C [1999Sch]
(αMn) < 727
cI58 I 43m (αMn)
a = 891.26
pure Mn at 25°C [Mas2]. Dissolves up to 35.2 at.% Fe [2004Wit] and 1.25 at.% Zr at 730°C [1999Sch]
(βZr) 1855–863
cI2 Im 3m W
a = 360.90
[Mas2] dissolves up to 4.2 at.% Fe at 928°C [2002Ste] and 10.2 at.% Mn at 1090°C [1999Sch]
(αZr) < 863
hP2 P63/mmc Mg
a = 323.16 c = 514.75
[Mas2] dissolves up 10.2 at.% Mn at 1090°C [1999Sch]
Zr3Fe < 851
oC16 Cmcm BRe3
a = 332 b = 1100 c = 882.0
from 74.8 to 75.4 at.% Zr [2002Ste]
Zr2Fe 951–780
tI12 I4/mcm Al2Cu
a = 638 c = 560
from 66.7 to 67.2 at.% Zr (C16 structure) [2002Ste]
ZrFe2 (C15) < 1673
cF24 Fd 3m MgCu2
a = 702 to 709
from 27.5 to 34.5 at.% Zr [2002Ste], Laves phase. Dissolves up to 30 mol% ZrMn2 [1968Kan, 2004Li] at 33.3 at.% Zr (ZrFe2) [2004Li] at 33.3 at.% Zr and 20.0 at.% Mn (x = 0.7) [2004Li]
Zr(FexMn1–x)2 (x > 0.7)
a = 707.0 a = 708.5
ZrFe2 (C36) 1345–1240 ± 10
hP24 P63/mmc MgNi2
a = 495 c = 1614
from 26.5 to 27 at.% Zr [2002Ste] Laves phase
Zr6Fe23 1482–1175 ?
cF116 Fm 3m Th6Mn23
a =1169
metastable, from 20.6 to 21.6 at.% Zr [2002Ste] sometimes labelled ZrFe3 (continued)
DOI: 10.1007/978-3-540-78644-3_24 # Springer 2008
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Fe–Mn–Zr
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
ZrMn2 (C14) < 1450
hP12 P63/mmc MgZn2
Zr(FexMn1–x)2 (x < 0.5) Zr1–y(FexMn1–x)2 (x < 0.5, y < 0.07)
Table 3.
5
Lattice Parameters [pm]
Comments/References
28 to 40 at.% Zr [1999Sch] Laves phase. Dissolves up to 60 mol% ZrFe2 [1968Kan, 2004Li] at 28.1 at.% Zr [1999Sch]
a = 499.7 c = 820.7 a = 502.6 c = 825.8 a = 506.8 c = 830.2 a = 501.3 c = 819.6 a = 498.4 c = 817.1
at 33.3 at.% Zr [1987Zha, 1999Sch] at 40.0 at.% Zr [1999Sch] at 33.3 at Zr and 33.3 at.% Fe (x = 0.5) [2004Li] x = 0.40, y = 0.07 [1987Zha]
Investigations of the Fe-Mn-Zr Materials Properties
Reference
Method/Experimental Technique
Conditions, Type of Property
[1968Kan]
Saturation magnetization, magnetic moment measurements
Zr(Mn1–xFex)2, (0 < x < 1), evolution from C14 to C15 structure with composition
[1971Kan]
XRD, magnetic moment, Curie temperature
77–500 K, Zr(Mn1–xFex)2, (0 < x < 1), structural and magnetic evolution
[1986Uch]
XRD, hydrogen intake vs hydrogen pressure characteristics
20–250°C, ZrMn2.7, ZrMn2.0Fe0.8, ZrMn1.22Fe1.11, ZrMn1.11Fe1.22
[1987Nic]
XRD, magnetic susceptibility, Curie temperature
4.2–120 K, amorphous Fe90–xMnxZr10 (x < 22 at.% Mn), spin glass behavior
[1987Zha]
XRD, effective magnetic moment, magnetization, Curie temperature
4.2–273 K, ZrMn2, ZrMn2.8, ZrMn2Fe0.8 and hydrogenated products, spin glass state
[1993Sou]
XRD, SEM, neutron diffraction, EDX
1000°C, ZrMn1.6Fe0.4 and hydrogenated products
[1994Sri]
Electrical conductivity, magnetization, Curie temperature
4.2–300 K, amorphous Fe90–xMnxZr10 (x < 12 at.% Mn)
[1995Ghe, 1997Ghe]
Activation analysis, reduction of tritiated water
400–600°C, composition not given, tritium recuperation
[1997She]
XRD, magnetic moment, Curie temperature, magnetization
1.5–1200 K, amorphous Fe90–xMnxZr10 (x < 15 at.% Mn), spin glass behavior
[2001Per1]
XRD, magnetic susceptibility under alternative current
12–300 K, amorphous Fe80Mn10Zr10, reentrant spin glass state
[2001Per2]
XRD, spontaneous magnetization, zero field and low field susceptibility
12–300 K, amorphous Fe90–xMnxZr10 (x < 8), critical exponents determination (continued)
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DOI: 10.1007/978-3-540-78644-3_24 # Springer 2008
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Fe–Mn–Zr
Reference
Method/Experimental Technique
Conditions, Type of Property
[2001Per3]
Low field susceptibility under alternative current
77–300 K, amorphous Fe90–xMnxZr10 (x < 16), critical exponents determination
[2002Per]
Magnetization, high field susceptibility, coercivity
4.2–300 K, amorphous Fe90–xMnxZr10 (x < 12), spin glass behavior
[2005Uly]
Magnetization, electron spin resonance, Curie temperatures
Amorphous Fe82Mn6–xBxZr10 (x < 8), superparamagnetism, nanoparticles
[2006Per]
Magnetization, magnetic susceptibility, coercivity
4.2–300 K, amorphous Fe90–xMnxZr10 (x < 16 at.% Mn), reentrant magnetic transition
Fig. 1. Fe-Mn-Zr. Crystal parameters of the solid solutions Zr(Fe,Mn)2
DOI: 10.1007/978-3-540-78644-3_24 # Springer 2008
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Fe–Mn–Zr
7
Fig. 2. Fe-Mn-Zr. The ZrMn2-ZrFe2 quasibinary system
Landolt-Börnstein New Series IV/11D4
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DOI: 10.1007/978-3-540-78644-3_24 # Springer 2008
8 References [1968Kan]
[1971Kan]
[1972Pet]
[1973Sve]
[1984Pou]
[1986Uch]
[1987Nic]
[1987Zha]
[1992Rag]
[1993Sou]
[1994Sri]
[1995Ghe]
[1997Ghe]
[1997She]
[1999Sch] [2000Pra]
[2000Ram]
Fe–Mn–Zr
Kanematsu, K., “Magnetism and Crystal Structures of Zirconium Compounds with Laves Structure”, J. Appl. Phys., 39(2), 465–466 (1968) (Crys. Structure, Experimental, Magn. Prop., 2) Kanematsu, K., “Magnetic Moment in Laves Phase Compounds. II. Zr(Fe1–xMnx)2 and Zr (Fe1–xCox)2”, J. Phys. Soc. Jpn., 31(5), 1355–1360 (1971) (Crys. Structure, Experimental, Magn. Prop., 8) Petkov, V.V., “Examination of Laves Phases in the Zr(Fe, Co, Ni)2-ZrMn2 Systems”, Russ. Metall., (5), 113–115 (1972), translated from Izv. Akad. Nauk SSSR, Met., (5), 155–157 (1972) (Crys. Structure, Experimental, Phase Diagram, 8) Svechnikov, V.N., Kocherzhinsky, Yu.A., Markiv, V.Ya., Pet’kov, V.V., “Laves Phases in Transition Metal Systems of the IV-VII Groups of Periodic Systems” (in Russian), Akad. Nauk Ukr. SSR, Metallofizika, 46, 35–45 (1973) (Experimental, Phase Diagram, Review, 74) Pourarian, F., Sinha, V.K., Wallace, W.E., “Hydrogen Sorption Properties of Nonstoichiometric ZrMn2-Based Systems”, J. Less-Common Met., 96, 237–248 (1984) (Experimental, Phase Relations, 16) Uchida, M., Bjurstroem, H., Suda, S., Matsubara, Y., “On the Equilibrium Properties of Some ZrMn2-Related Hydride-Forming Alloys”, J. Less-Common Met., 119, 63–74 (1986) (Crys. Structure, Experimental, Phase Relations, Thermodyn., 16) Nicolaides, G.K., Hadjipanayis, G.C., Rao, K.V., “Effect of Mn on the Collapse of “Ferromagnetic” Order in Melt-Spun Fe90Zr10 Amorphous Alloys”, J. Appl. Phys., 61(8), 3642 (1987) (Abstract, Magn. Prop., 0) Zhang, L.Y., Pedziwiatr, A.T., Pourarian, F., Wallace, W.E., “Influence of Hydrogen on the Magnetic Properties of Hyperstoichiometric ZrMn2T0.5 (T = Fe, Co, Ni and Cu)”, J. Magn. Magn. Mater., 68, 309–313 (1987) (Experimental, Magn. Prop., 9) Raghavan, V., “The Fe-Mn-Zr (Iron-Manganese-Zirconium) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Institute of Metals, Calcutta, 6B, 982–983 (1992) (Crys. Structure, Phase Diagram, Review, 5) Soubeyroux, J.L., Pontonnier, L., Miraglia, S., Isnard, O., Fruchart, D., Akiba, E., Hayakawa, H., Fujitani, S., Yonezu, I., “Crystal Structure and Microstructure Analysis of Alloys Zr(Mn1–xMx)2Hy with M = V, Fe, Ni, Co, Al and their Hydrides”, Z. Phys. Chem. (Munich), 179, 181–198 (1993) (Crys. Structure, Experimental, 12) Srinivas, V., Nigam, A.K., Chandra, G., Lawther, D.W., Yewondwossen, M., Dunlap, R.A., “Electrical Transport in Amorphous Fe-Mn-Zr Alloys”, J. Appl. Phys., 76(10/2), 6501–6503 (1994) (Experimental, Electr. Prop., Magn. Prop., 6) Ghezzi, F., Shmayda, W.T., Venkataramani, N., Bonizzoni, G., “Efficient HTO Reduction Using a Zr-Fe-Mn Alloy”, Fusion Eng. Design., 28, 367–372 (1995) (Experimental, Kinetics, 8) Ghezzi, F., Shmayda, W.T., Bonizzoni, G., “An Experimental Investigation of the Efficiency of HTO Reduction by Zr-Fe-Mn Getter Alloy”, Fusion Technol., 31(1), 75–105 (1997) (Experimental, Kinetics, Review, 50) Shen, B., Guo, H., Zhang, J., Zhan, W., Zhao, J., “Magnetic Properties and Crystallization of Amorphous Fe90–xMnxZr10 Alloys with x = 0–15”, Mater. Sci. Eng. A, 226A–228A, 668–671 (1997) (Experimental, Magn. Prop., 15) Schlesinger, M.E., “The Mn-Zr (Manganese-Zirconium) System”, J. Phase Equilib., 20(1), 79–83 (1999) (Crys. Structure, Phase Diagram, Review, 42) Prakash, M., Ramaprabhu, S., “Hydrogen Absorption Studies in ZrMnFe0.7Co0.3”, Int. J. Hydrogen Energy, 25, 871–878 (2000) (Crys. Structure, Experimental, Phase Relations, Thermodyn., 15) Ramaprabhu, S., “Structure and Hydrogen Solubility Studies on ZrMnFe1–xCox (x = 0.2, 0.3 and 0.5) Alloys”, Int. J. Hydrogen Energy, 25, 879–889 (2000) (Crys. Structure, Experimental, Phase Relations, Thermodyn., 14)
DOI: 10.1007/978-3-540-78644-3_24 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mn–Zr [2001Per1]
[2001Per2]
[2001Per3]
[2002Man1]
[2002Man2]
[2002Per]
[2002Ste]
[2003Dor]
[2004Li]
[2004Wit]
[2005Uly]
[2006Per] [2007Wit] [Mas2] [V-C2]
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Perumal, A., Srinivas, V., Rao, V.V., Dunlap, R.A., “Dynamics of Magnetic Susceptibility in Amorphous Fe80Mn10Zr10”, J. Alloys Compd., 326, 288–291 (2001) (Experimental, Magn. Prop., 6) Perumal, A., Srinivas, V., Kim, K.S., Yu, S.C., Rao, V.V., Dunlap, R.A., “Correction to Scaling Critical Exponents and Amplitudes for Amorphous Fe-Mn-Zr Alloys”, J. Magn. Magn. Mater., 233(3), 280–289 (2001) (Experimental, Phase Relations, Magn. Prop., 26) Perumal, A., Srinivas, V., Rao, V.V., Dunlap, R.A., “Low Field ac Response in Mn Substituted Amorphous Fe-Zr Alloys”, Mater. Sci. Eng. A, 304A–306A, 1004–1007 (2001) (Experimental, Magn. Prop., 12) Mani, N., Sivakumar, R., Ramaprabhu, S., “Hydrogen Storage Properties of ZrMnFe1–xNix (x = 0.2, 0.4, 0.5 and 0.6) Alloys”, J. Alloys Compd., 337, 148–154 (2002) (Crys. Structure, Experimental, Phase Relations, Thermodyn., 16) Mani, N., Kesavan, T.R., Ramaprabhu, S., “Study of Hydrogen Absorption by ZrMnFe1–xCox (x = 0.2, 0.4, 0.5 and 0.6) Alloys”, J. Phys.: Condens. Matter, 14(15), 3939–3949 (2002) (Crys. Structure, Experimental, Phase Relations, Thermodyn., 15) Perumal, A., Srinivas, V., Kim, K.S., Yu, S.C., Rao, V.V., Dunlap, R.A., “Magnetic Properties of Amorphous Fe90–xMnxZr10 (0 < x < 12) Alloys”, Phys. Rev. B, 65B(6), 064428_1–064428_15 (2002) (Experimental, Magn. Prop., 59) Stein, F., Sauthoff, G., Palm, M., “Experimental Determination of Intermetallic Phases, Phase Equilibria, and Invariant Reaction Temperatures in the Fe-Zr System”, J. Phase Equilib., 23(6), 480–494 (2002) (Crys. Structure, Experimental, Phase Relations, Phase Diagram, Morphology, 88) Dorogova, M., Hirata, T., Filipek, S.M., “Hydrogen-Induced Volume Changes in ZrCr2 and Pseudo-Binary Compounds of ZrCr2, ZrMn2 and ZrV2”, Phys. Status Solidi A, 198A(1), 38–42 (2003) (Crys. Structure, Experimental, 13) Li, G., Nishimiya, N., “Phase Relations of Zr(FexMn1–x)2-H2 Systems Studied by XRD and Isothermal Calorimetry Measurements”, J. Alloys Compd., 375, 205–211 (2004) (Crys. Structure, Thermodyn., Experimental, 18) Witusiewicz, V.T., Sommer, F., Mittemeijer, E.J., “Reevaluation of the Fe-Mn Phase Diagram”, J. Phase Equilib. Diff., 25(4), 346–354 (2004) (Experimental, Phase Diagram, Calculation, Thermodyn., #, 34) Ulyanov, A.N., Yu, S.C., Kang, Y.M., Yoo, S.I., “High Temperature Superparamagnetism in Boron Substituted Fe-Zr-Mn Alloys”, IEEE Trans. Magn., 41(10), 3265–3267 (2005) (Experimental, Magn. Prop., Nano, 12) Perumal, A., “Study of Magnetic Properties of Amorphous Fe-Mn-Zr Alloys”, Proc. Indian Natl. Sci. Acad., Part A, 72A(1), 43–48 (2006) (Experimental, Magn. Prop., 20) Witusiewicz, V.T., private communication to MSI, (2007) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
MSIT®
DOI: 10.1007/978-3-540-78644-3_24 # Springer 2008
Fe–Mo–N
1
Iron – Molybdenum – Nitrogen Vasyl Tomashik
Introduction The Fe-Mo-N ternary system was investigated already since fifty years but not much is known about the phase relationships of this system [1984Rag, 1987Rag1]. The solubilities of nitrogen in liquid and solid Fe-Mo alloys were measured by [1958Kas, 1960Mae, 1960Peh, 1963Sch, 1968Uda, 1973Pip, 1973Sak1, 1973Sak2, 1975Bel, 1981Yos]. All available data on nitrogen solubility in Fe rich alloys were evaluated by [1970Kun]. It was established that for alloys containing between 3 and 10 mass% Mo the solubility is independent on temperature within the interval 1500 to 1900°C. At 1700°C the solubility appears to increase approximately as a linear function of Mo content up to 10 mass% Mo [1958Kas]. Nitrogen solubility was found to obey Sieverts’ law for Fe-Mo alloys [1960Peh]. In arc-melted samples the nitrogen solubility appears about 20% higher than in levitation melted ones, most probably due to cracking the N2 molecules in the arc and the resulting non-equilibrium content of free N atoms in the gas. In a low nitrogen-content atmosphere (0.14 vol% N2 + Ar) the influence of Mo on the solubility of nitrogen is quite anomalous [1968Uda]. At higher N2 contents (3 or 5 vol% N2 + Ar) it is influenced by the alloying Mo in a manner similar to the equilibrium solubility found by levitation melting. The formation of a two-dimensional surface compound MoN was found on the (100) surfaces of a Fe-Mo single crystal, containing 3.5 mass% Mo [1997Bar]. This compound is epitaxed to the bcc (100) surface. It was found that this compound consists of a double Mo and a single N layer [1997Bar, 1998Vil]. The thickness of the MoN surface compound is about 230 pm [1997Bar]. [1973Sak1] and [1973Sak2] reported data on the (αFe)/((αFe) + Fe8N) and (αFe)/((αFe) + Fe4N) phase boundaries between 100 and 450°C in alloys containing up to 1.26 mass% Mo. The ternary compounds Fe3Mo3N [1957Eva, 1997Pan], Fe7Mo13N4 [1957Eva], (Fe0.8Mo0.2)MoN2 [1995Bem] and FeMoN2 [1998Pan] were found in this system. Thermodynamic properties of solid and liquid alloys in the Fe-Mo-N ternary system were investigated by [1958Kas, 1960Mae, 1960Peh, 1963Sch, 1968Uda, 1973Pip, 1973Sak1, 1973Sak2, 1983Pul2, 1991Sat, 1994Tan]. Table 1 lists the numerous experimental works on phase equilibria, crystal structure and thermodynamics of the Fe-Mo-N system. Binary Systems The Fe-Mo, Fe-N and Mo-N binary system are accepted from [Mas2]. Solid Phases Crystallographic data of all unary, binary and ternary compounds are listed in Table 2. The solubility of nitrogen in Fe-Mo alloys is increased by the presence of Mo [1958Kas, 1960Mae, 1960Peh, 1968Uda, 1973Pip, 1973Sak1, 1973Sak2, 1975Bel, 1981Yos]. According to [1975Huf] nitrided alloys contain a stoichiometric phase MoN and no mixed metal nitrides containing Fe were detected. The solubility of nitrogen in Fe- Mo alloy, containing up to 25 mass% Mo, at 1200–1600°C obeys Sieverts’ law up to 1 MPa N2 pressure [1963Sch, 1991Sat]. The alloy, containing 2.9 at.% Mo, dissolves 2.5 at.% of nitrogen at 580°C [1972Dri]. An analysis of the data of [1981Yos] by [1984Rag, 1987Rag1] gave the following empirical equation for the nitrogen solubility at 101.3 kPa N2 up to 1200°C: log(mass% N) = – 1046/T – 1.58 + (– 55/T + 0.064)(mass% Mo). At temperatures above 1200°C this equation gives lower solubilities than those found experimentally by [1981Yos].
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DOI: 10.1007/978-3-540-78644-3_25 # Springer 2008
2
Fe–Mo–N
From evaluation of all available published data on nitrogen solubility in Fe rich alloys at various temperatures and nitrogen partial pressures [1970Kun] concluded, that the nitrogen solubility at constant temperature (log(mass% NT)), the heat of solution (H ’) and the entropy of solution (S ’) of nitrogen are practically linearly independent irrespective of melt composition. The following semiempirical equation was proposed for determination of nitrogen solubility: log (mass% NT) = 12log p(N2) – 1.290 – H’/4.575 (1/T – 1/3746). The solubility of nitrogen at 0.1 MPa pressure in Fe-Mo alloys at 1606 and 2240 ± 50°C according to [1958Kas, 1960Peh, 1968Uda] is shown in Fig. 1. The concentration dependence of the lattice parameter of the Fe-Mo-N solid solutions at 23°C can be expressed by a (nm) = (2.4 ± 0.1)·10– 4(at.% Mo) [1983Pul2]. The (αFe)/((αFe) + Fe8N) and (αFe)/((αFe) + Fe4N) phase boundaries computed from the results of [1973Sak1, 1973Sak2] by [1984Rag, 1987Rag1] are shown in Figs. 2 and 3. The limiting solubility of nitrogen in ferrite generally increases with increasing Mo content and increasing temperature. At the overlapping temperature range 200 to 350°C, where both, Fe8N and Fe4N, can form, the solubility in the matrix in equilibrium with the metastable precipitate Fe8N is larger than in that with the stable Fe4N. Ternary compounds Fe3Mo3N [1957Eva, 1997Pan], Fe7Mo13N4 [1957Eva], (Fe0.8Mo0.2)MoN2 [1995Bem] and FeMoN2 [1998Pan] are formed in the Fe-Mo-N ternary system. Fe3Mo3N has a η1 carbide structure with a cubic unit cell. Fe7Mo13N4 is also cubic and has the βMn type atom arrangement. (Fe0.8Mo0.2) MoN2 and FeMoN2 are hexagonal [1995Bem]. Thermodynamics Molybdenum has a pronounced negative effect on the activity coefficient of nitrogen in liquid Fe (eMoN = – 0.013 at 1600–1700°C up to 7 mass% Mo [1960Mae, 1991Sat]; eMoN = – 0.011 at 1600°C up to 10.25 mass % Mo [1960Peh]; eMoN = – 0.05 at 1200°C at 0.56 at.% Mo [1960Peh]; eMoN = – 0.0051 at 2140–2240°C up to 8 mass% Mo and eMoN = – 0.0079 at 3.04 kPa of N2 up to 8 mass% Mo [1968Uda]; eMoN = – 0.19 ± 0.06 at 400–535°C up to 4 mass% Mo). According to [1973Pip] eMoN = – 260/T + 0.13 in the temperature range 400 to 500°C. According to the calculation of [1994Tan] εMoN = – 6.2. The interaction parameters are: eMoN = 0.13 and rMoN = 7.9·10– 5 [1991Sat]. ðxS Þ=k ) of nitrogen in Fe-Mo-N solid solutions are i ) and excess entropies (S Partial energies (E i – 435.9 ± 21.1 kJ·mol–1 and 9.7 ± 1.1 at 3.1 at.% Mo, – 436.7 ± 19.1 kJ·mol–1 and 10.1 ± 0.9 at 4.2 at.% Mo and – 440.0 ± 22.3 kJ·mol–1 and 10.7 ± 0.9 at 5.3 at.% Mo, respectively [1981Yos]. Assessed values of enthalpy and entropy interaction parameters are equal ηMoN = – 183 kJ·mol–1 and σMoN = – 46.1 J·(mol·K) –1 [1994Tan]. The addition of Mo decreases the heat of solution of Fe8N and Fe4N in (αFe) [1973Sak1, 1973Sak2]. The activity coefficient of nitrogen in liquid Fe decreases with decreasing Mo content [1958Kas].
Notes on Materials Properties and Applications In general, nitrogen is unwanted in Fe-Mo alloys [1991Sat]. During solidification by the involved decrease of solubility the gas may segregate and, depending on the boundary conditions, nitrogen bubbles can be produced. This may lead to refinishing operations or waste of work pieces. In some special cases, however, nitrogen is used as an alloying element. Investigations of some Fe-Mo-N material properties are given in Table 3. Fe3Mo3N shows metallic behavior [1997Pan]. Changes in the temperature coefficient of electric resistance at 120 K corroborate a magnetic phase transition. The material shows paramagnetic behavior and orders antiferromagnetically at a Neel temperature, TN = 120 K. The magnetic susceptibility displays a maximum for (Fe0.8Mo0.2)MoN2 at 20 K and four-probe conductivity measurements indicate poor metallic conductivity with a very small temperature dependence [1995Bem]. Ultrafine FeMoN2 is Pauli-paramagnetic in the temperature range 77–300 K having a magnetic susceptibility of 2.3·10–4 emu·g–1 at room temperature [1998Pan]. This compound exhibits also metallic behavior.
DOI: 10.1007/978-3-540-78644-3_25 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mo–N
3
Damping measurements on nitrided Fe alloys with Mo contents between 0.7 and 4 mass% were carried out by [1966Koe]. The effects of nitrogen content, heat treatment and degree of deformation on the formation of various maximum damping values were discussed. Miscellaneous The rate of total nitridization of Fe-Mo alloys is controlled by the diffusion of the nitrogen [1973Kun]. Diffusion coefficients of the nitrogen in such alloys depending on the Mo content at different temperatures are given in Fig. 4. According to [1973Pip] Mo increases the precipitate density at the nitriding of the Fe-Mo alloys. The results of [1972Dri] show that the early stages of precipitation of a nitride from ferrite are associated with the formation of zones, as indicated by the appearance of side bands in the XRD patterns. According to [1975Bel] the Mo nitrides are not formed in the nitrated layer. But according to [1975Huf] the nitrided alloys contain precipitates of the stoichiometry MoN and no ternary compounds were detected. [1987Mor] noted that since the value of the standard Gibbs energy of Mo2N formation at 1600°C in liquid iron alloys is positive this nitride can not be precipitated in such alloys under normal conditions. The hardening of the diffusion layers of the Fe-Mo alloys is determined by the quantity of dissolved nitrogen [1975Bel]. Mo was found to segregate to interfaces in such alloys [1983Pul1]. Segregation of Mo is partly reduced by the formation of clusters or precipitates between nitrogen and solute atoms. It was found that the presence of small concentrations of bulk dissolved nitrogen in Fe-Mo alloys, containing 2 and 3.5 mass% Mo, induce a strong synergetic segregation (cosegregation) of Mo and N [1998Vil, 2001Vil]. This effect is strongest for (100)-oriented alloy, where an epitaxially stabilized MoN surface compound is formed [1998Vil]. The rate of Mo segregation is only slightly enhanced by the presence of nitrogen [1999Vil]. The diffusion parameters of Mo in the first 20 layers of a Fe matrix were determined as D = 7·10–3 m2·s–1 and E = 290 kJ·mol–1. Fe3Mo3N was synthesized through an intermediate nitride having WC type structure by nitridation of a FeMoO4 precursor at 700°C [1997Pan] by ammonolysis of this precursor as well as by plasma nitridation of oxide precursors [1999Jac]. This compound is thermally stable up to 360°C [1997Pan]. The layered ternary nitride (Fe0.8Mo0.2)MoN2 could be synthesized in a single-step reaction by ammonolysis of Fe2(MoM4)3·8H2O [1995Bem]. This nitride decomposes by heating it in oxygen to 700°C, where the products of the oxidation were identified as Fe2O3 and MoO3. FeMoN2 is formed when Fe3Mo3N is heat treated at 300°C under flowing NH3 for 6 h [1998Pan]. The crystallite size was found to be 12 nm.
Table 1.
Investigations of the Fe-Mo-N Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1957Eva]
XRD
Room temperature/Fe3Mo3N and Fe7Mo13N4
[1958Kas]
Sieverts’ method, N-solubility
1500–1900°C/Fe-Mo-N up to 10 mass% Mo
[1960Mae]
Sieverts’ method, N-solubility
1600–1750°C/Fe-Mo-N up to 7 mass% Mo
[1960Peh]
Sieverts’ method, N-solubility
1600°C/Fe-Mo-N up to 10.25 mass% Mo
[1963Sch]
Chemical analysis, N-solubility
1200°C/Fe-Mo-N at 0.56 at.% Mo
[1968Uda]
Levitation melting, chemical analysis arc melting, N-solubility
2140–22400°C/(Fe-Mo) up to 8 mass% Mo + N2/Ar
[1970Kun]
Calculation
1400–1800°C/Fe-Mo-N (continued)
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Fe–Mo–N
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1972Dri]
XRD, microhardness measurements, gravimetry
580°C/Fe-Mo-N at 2.9 at.% Mo
[1973Kun]
Steady-state method, N-solubility
1550–1700°C/Fe-Mo-N
[1973Pip]
XRD, metallography
400–500°C/Fe-Mo-N
[1973Sak1]
Internal-friction technique
100–250°C/(Fe-Mo) at 0.65 and 0.92 mass% Mo + NHs/H2
[1973Sak2]
Internal-friction technique
250–450°C/(Fe-Mo) at 0.36, 0.65, 0.92 and 1.26 mass% Mo + NH3/H2
[1975Bel]
XRD, microhardness measurements, metallography
450–1000°C/(Fe-Mo) at 1.05, 1.82, 3.02 and 5.0 mass% Mo + N2
[1975Huf]
TEM, Mössbauer spectroscopy
400–600°C/(Fe-Mo) at 0.28 and 2.97 at.% Mo + NH3/H2
[1981Yos]
Chemical analysis, elastic properties measurements
660–1372°C/(Fe-Mo) at 3.1, 4.2 and 5.3 at.% Mo + N2
[1983Pul1]
SEM, AES
Room temperature/(Fe-Mo) at 2.4 at.% Mo + N2
[1983Pul2]
TEM, spectroscopic analysis
400–535°C/Fe-Mo-N up to 4 mass% Mo
[1991Sat]
Chemical analysis, N-solubility at elevated pressure
1600°C/Fe-Mo-N up to 25 mass% Mo
[1995Bem]
XRD, TGA, SEM, chemical analysis, crystal structure determination
Up to 700°C/(Fe0.8Mo0.2)MoN2
[1997Bar]
XPD, XPS, LEED, surface structure
620°C/Fe-Mo-N at 3.5 mass% Mo
[1997Pan]
XRD, DTA, SEM, chemical analysis, magnetic and electrical resistivity measurements, Mössbauer spectroscopy
Up to 700°C/Fe3Mo3N
[1998Pan]
XRD, morphology, Mössbauer spectroscopy
Up to 300°C/FeMoN2
[1998Vil]
XRD, AES, LEED, chemical analysis
500–900°C/Fe-Mo-N at 3.5 mass% Mo
[1999Jac]
XRD, TGA, neutron diffraction, chemical analysis, crystal structure determination
Room temperature/Fe3Mo3N
[1999Vil]
AES, surface segregation
500–600°C/Fe-Mo-N at 3.5 mass% Mo
[2001Vil]
AES, surface segregation
517–602°C/Fe-Mo-N at 2 mass% Mo
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Fe–Mo–N Table 2.
5
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(αδFe) (δFe) 1538–1394 (αFe) < 912
cI2 Im 3m W
a = 293.15
pure Fe at 1390°C [V-C2, Mas2]
a = 286.65
pure Fe at 25°C [Mas2] dissolves up to 25 at.% Mo at 1449°C
(γFe) 1394–912
cF4 Fm 3m Cu
a = 364.67
pure Fe at 915°C [V-C2, Mas2]
(Mo) < 2623
cI2 Im 3m W
a = 314.70
at 25°C [Mas2]
MoFe (σ phase) 1611–1235
tP30 P42/mnm CrFe
a = 921.8 ± 0.2 c = 481.3 ± 0.2
[Mas2, V-C2] contains 43–57 at.% Mo
MoFe2 (λ phase) < 927
hP12 P63/mmc MgZn2
a = 474.5 c = 773.4
[Mas2, V-C2]
Mo6Fe7 (μ phase) < 1370
hR39 R 3m W6Fe7
a = 475.46 ± 0.05 c = 2571.6 ± 0.3
[Mas2, V-C2] contains 39–44 at.% Mo
Mo2Fe3 (R phase) 1488–1200
hR159 R 3 Co5Cr2Mo3
a = 1095.6 c = 1935.3
[Mas2, V-C2] contains 33.9–38.5 at.% Mo
α”Fe16N2
tI* I4/mmn
a = 572 c = 629
ordered fcc structure, metastable [1987Rag2]
γ’Fe4N < 680
cP5 Pm 3m Fe4N
a = 378.7
19.4 to 20.6 at.% N. Ordered fcc structure [1987Rag2]
ε Fe3N < 580
hP10 P6322 Fe3N
a = 469.96 ± 0.03 c = 438.04 ± 0.03 a = 471.8 c = 438.8 a = 479.1 c = 441.9
ζFe2N < 500
oP12 Pbcn Fe2N
a = 551.2 b = 482.0 c = 441.6
15.8 to 33.2 at.% N [1987Rag2] εFe3N at RT [1999Lei] εFe3N1.10 [2001Lei] Lattice parameters decrease slightly with decrease in nitrogen content [2001Lei] εFe3N1.39 [2001Lei] at 25°C [1987Rag2]
(continued)
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6
Fe–Mo–N
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
MoN ?
hP16 P63/mmc MoN
a = 572.5 c = 560.8
[Mas2, V-C2]
βMo2N < 850
tI12 I41/amd Mo2N
a = 420.0 c = 801.0
[Mas2, V-C2] contains 28.7-35 at.% N
Mo3N2 ?
cP8 Pm 3m Mo3N2
a = 416.5
[Mas2, V-C2]
* τ1, Fe3Mo3N
cF112 Fd 3m W3Fe3C
a = 1108.0 ± 1.5 a = 1106.7 a = 1108.59 ± 0.03
[1957Eva] Stable in air up to 560°C [1997Pan] [1999Jac]
* τ2, Fe7Mo13N4
cP24 P4132 βMn
a = 669.84 ± 0.36
[1957Eva]
* τ3, (Fe0.8Mo0.2) MoN2
hP8 P 31c Fe0.8Mo1.2N2
a = 285.62 ± 0.01 c = 1099.97 ± 0.04
Stable in oxygen up to 700°C [1995Bem]
* τ4, FeMoN2
hP2 P 6m2 WC
a = 284.3 c = 273.5
[1998Pan]
Table 3.
Investigations of the Fe-Mo-N Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1966Koe]
Torsion pendulum
Damping
[1995Bem]
Four-probe conductivity measurements, magnetometry
Electric conductivity and magnetic susceptibility for (Fe0.8Mo0.2)MoN2
[1997Pan]
Four-probe conductivity measurements, magnetometry, AC susceptibility measurements
Electric conductivity, magnetic susceptibility, temperature coefficient of resistance and Neel temperature for Fe3Mo3N
[1998Pan]
Magnetometry, AC susceptibility measurements, resistivity study
Magnetization, magnetic susceptibility, and temperature coefficient of resistance for FeMoN2
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Fe–Mo–N
7
Fig. 1. Fe-Mo-N. Solubility of nitrogen at 0.1 MPa pressure in Fe-Mo alloys at 1606 and 2240 ± 50°C
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Fe–Mo–N
Fig. 2. Fe-Mo-N. (αFe)/((αFe) + Fe4N) phase boundaries at 125, 175 and 225°C
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Fe–Mo–N
9
Fig. 3. Fe-Mo-N. (αFe)/((αFe) + Fe8N) phase boundaries at 250, 350 and 450°C
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Fe–Mo–N
Fig. 4. Fe-Mo-N. Diffusion coefficients of nitrogen in Fe-Mo alloys depending on the Mo content at different temperatures
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Fe–Mo–N References [1957Eva] [1958Kas] [1960Mae]
[1960Peh]
[1963Sch]
[1966Koe]
[1968Uda]
[1970Kun]
[1972Dri] [1973Kun]
[1973Pip]
[1973Sak1]
[1973Sak2]
[1975Bel]
[1975Huf] [1981Yos]
[1983Pul1]
[1983Pul2]
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11
Evans, D.A., Jack, K.H., “Interstitial Alloys with the Structure of β -Manganese”, Acta Crystallogr., 10, 769–770 (1957) (Experimental, Crys. Structure, 2) Kashyap, V.C., Parlee, N., “Solubility of Nitrogen in Liquid Iron and Iron Alloys”, Trans. Met. Soc. AIME, 212, 86–91 (1958) (Experimental, Phase Relations, *, 19) Maekawa, S., Nakagawa, Y., “Solubility of Nitrogen in Liquid Iron and Iron Alloys. II. Effect of Niickel, Cobalt, Molybdenum, Chromium and Vanadium on the Solubility in Liquid Iron”, Tetsu to Hagane, 46(9), 972–976 (1960) (Experimental, Phase Relations, Thermodyn., 8) Pehlke, R.D., Elliott, J.F., “Solubility of Nitrogen in Liquid Iron Alloys. I. Thermodynamics”, Trans Met. Soc. AIME, 218(6), 1088–1101 (1960) (Experimental, Phase Relations, Thermodyn., *, 32) Schenck, H., Frohberg, M.G., Reinders, F., “Contribution to the Study of the Solubility of Nitrogen in the Iron Alloys over the Temperature Range from 700 to 1200°C” (in German), Stahl u. Eisen, 83(2), 93–99 (1963) (Experimental, Phase Relations, Thermodyn., 36) Koester, W., Horn, W., “Damping Investigation on Nitrided Iron-Molybdenum and IronVanadium Alloys” (in German), Arch. Eisenhuettenwes., 37(3), 245–252 (1966) (Experimental, Morphology, Thermodyn., 22) Uda, M., Wada, T., “Solubility of Nitrogen in Arc-Melted and Levitation-Melted Iron and Iron Alloys”, Trans. Nat. Res. Inst. Met. (Jpn.), 10(2), 79–91 (1968) (Experimental, Phase Relations, Thermodyn., *, 40) Kunze, H.-D., Schuermann, E., Parlee, N.A.D., “Influence of Temperature and Equivalent Effect of Added Elements on the Solubility, Activity and Activity Coefficient of Nitrogen in Liquid Iron”, Met. Trans., 1(1), 281–290 (1970) (Calculation, Phase Relations, Thermodyn., 45) Driver, J.H., Unthank, D.C., Jack, K.H., “Substitutional-Interstitial G.P. Zones in Nitrided FeMo Alloys”, Philos. Mag., 26(5), 1227–1231 (1972) (Experimental, Phase Relations, 8) Kunze, H.-D., “Effect of the Elements Chromium, Manganese, Cobalt, Nickel, Molybdenum and Tungsten on the Diffusion of Nitrogen in Liquid Iron Alloys” (in German), Arch. Eisenhuettenwes., 44(2), 71–80 (1973) (Experimental, Kinetics, 50) Pipkin, N.J., Roberts, W., Speirs, D.L., Grieveson, P., Jack, K.H., “Effect of Substitutional Alloying Elements on the Activity Coefficients and Behaviour of Interstitial Solutes in Iron”, Internat. Symposium on Metallurgical Chemistry - Applications to Ferrous Metall., 351–352 (1973) (Experimental, Phase Relations, 7) Sakamoto, M., Imai, Y., “Effects of Cobalt, Nickel, Molybdenum and Tungsten on the Solubility of Fe16N2 in α-Iron”, J. Jpn. Inst. Met., 37(7), 708–714 (1973) (Experimental, Phase Relations, Thermodyn., *, #, 9) Sakamoto, M., Masumoto, T., Imai, Y., “The Effect of Cobalt, Molybdenum and Tungsten on the Solubility of Fe4N in α-Iron”, J. Jpn. Inst. Met., 37(3), 343–349 (1973) (Experimental, Phase Relations, Thermodyn., *, #, 19) Belotsky, A.V., Nikitina, O.V., Permayakov, V.G., “Structure and Some Properties of the Nitrided Layer in Iron-Molybdenum Alloys” (in Russian), Izv. Akad. Nauk SSSR, Met., (1), 112 (1975) (Experimental, Phase Relations, Mechan. Prop., 9) Huffman, G.P., Podgurski, H.H., “Mössbauer Study of Nitrided Fe-Mo and Fe-Ti Alloys”, Acta Metall., 23(11), 1367–1379 (1975) (Experimental, Phase Relations, 21) Yoshihara, M., McLellan, R.B., “Thermodynamics of Body-Centered Cubic Iron- MolybdenumNitrogen Solid Solutions”, Acta Metall., 29(3), 471–478 (1981) (Experimental, Thermodyn., *, 12) Pulkkinen, R.E.E., Laehdeniemi, M., “Grain-Boundary Segregation in Bright Nitrided α-Irons Alloyed with Chromium, Molybdenum and Silicon”, J. Mater. Sci., 18(11), 3421– 3426 (1983) (Experimental, Morphology, 12) Pulkkinen, R., “The Activity Coefficients of Nitrogen in α -Irons Alloyed with Chromium, Molybdenum and Silicon”, Scand. J. Metall., 12, 87–92 (1983) (Experimental, Crys. Structure, Thermodyn., 24)
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12 [1984Rag] [1987Mor] [1987Rag1]
[1987Rag2]
[1991Sat]
[1994Tan]
[1995Bem]
[1997Bar]
[1997Pan]
[1998Pan]
[1998Vil]
[1999Jac]
[1999Lei]
[1999Vil]
[2001Lei]
[2001Vil] [Mas2] [V-C2]
Fe–Mo–N Raghavan, V., “The Fe-Mo-N (Iron-Molybdenum-Nitrogen) System”, Trans. Indian Inst. Met., 37(4), 308–312, (1984) (Review, Crys. Structure, Phase Diagram, Phase Relations, *, #, 23) Morita, Z., Tanaka, T., Yanai, T., “Equilibrium of Nitride Forming Reactions in Liquid Iron Alloys”, Metall. Trans., B, 18B(3), 195–202 (1987) (Review, Phase Relations, Thermodyn., 29) Raghavan, V., “The Fe-Mo-N (Iron- Molybdenum- Nitrogen) System” in “Phase Diagrams of Ternary Iron Alloys”, Ind. Inst. Techn., Delhi, 1, 191–194 (1987) (Review, Crys. Structure, Phase Diagram, Phase Relations, *, #, 20) Raghavan, V., “The Fe-N (Iron-Nitrogen) System” in “Phase Diagrams of Ternary Iron Alloys”, Ind. Inst. Techn., Delhi, 1, 143–144 (1987) (Crys. Structure, Experimental, Phase Diagram, Phase Relations, 7) Satir-Kolorz, A.H., Feichtinger, H.K., “On the Solubility of Nitrogen in Liquid Iron and Steel Alloys Using Elevated Pressure”, Z. Metallkd., 82(9), 689–697 (1991) (Experimental, Phase Relations, Thermodyn., 24) Tanaka, T., Gokcen, N.A., Iida, T., Morita, Z.-I., “Thermodynamic Relationship between the Enthalpy Interaction Parameter and the Entropy Interaction Parameter in Liquid Iron-Nitrogen Based Ternary Alloys”, Z. Metallkd., 85(10), 696–700 (1994) (Calculation, Thermodyn., 17) Bem, D.S., Olsen, H.P., zur Loye, H.-C., “Synthesis and Electronic and Magnetic Characterization of the Ternary Nitride (Fe0.8Mo0.2)MoN2”, Chem. Mater., 7(10), 1824–1828 (1995) (Experimental, Crys. Structure, Magn. Prop., 46) Baraldi, A., Brena, B., Cocco, D., Comelli, G., Lizzit, S., Paolucci, G., Baumann, P., Scheuch, V., Uebing, C., “The Structure of the MoN Surface Compound on Fe-3.5%Mo-N (100) Studied by X-Ray Photoelectron Diffraction: First Results from ELETTRA”, Vacuum, 48(3–4), 351–355, (1997) (Experimental, Crys. Structure, Interface Phenomena, 34) Panda, R.N., Gajbhiye, N.S., “Electronic and Magnetic Properties of Fe3Mo3N”, J. Alloys Compd., 256(1–2), 102–107 (1997) (Experimental, Crys. Structure, Magn. Prop., Morphology, 25) Panda, R.N., Gajbhiye, N.S., “Chemical Synthesis and Magnetic Properties of Nanocrystalline FeMoN2”, J. Cryst. Growth, 191(1–2), 92–96, (1998) (Experimental, Crys. Structure, Morphology, 15) Viljoen, E.C., Uebing, C., “Effect of Nitrogen and Sulphur on the Segregation of Molybdenum on Single Crystalline Fe3.5%Mo-N,S Alloys”, Surf. Sci., 410(1), 123–131, (1998) (Experimental, Interface Phenomena, 24) Jackson, S.K., Layland, R.C., zur Loye, H.-C., “The Simultaneous Powder X-Ray and Neutron Diffraction Refinement of Two η-Carbide Type Nitrides, Fe3Mo3N and Co3Mo3N, Prepared by Ammonolysis and by Plasma Nitridation of Oxide Precursors”, J. Alloys Compd., 291(1–2), 94–101 (1999) (Experimental, Crys. Structure, 25) Leineweber, A., Jacobs, H., Hüning, F., Lueken, H., Schilder, H., Kockelmann, W., “ε-Fe3N: Magnetic Structure, Magnetization and Temperature Dependent Disorder of Nitrogen”, J. Alloys Compd., 288(1–2), 79–87 (1999) (Experimental, Crys. Structure, 40) Viljoen, E.C., Jordaan, W.A., Uebing, C., du Plessis, J., “The Effect of N on the Segregation Kinetics of Mo in a Fe-3.5%Mo-N(100) Single Crystal Studied by Constant Temperature Heating Method”, Mater. Sci. Forum, 294–296, 461–464 (1999) (Experimental, Interface Phenomena, 10) Leineweber, A., Jacobs, H., Hünning, F., Luecken, H., Kockelmann, W., “Nitrogen Ordering and Ferromagnetic Properties of ε-Fe3N1+x (0.10 < x < 0.39) and ε-Fe3(N0.80C0.20)1.38”, J. Alloys Compd., 316, 21–38 (2001) (Experimental, Crys. Structure, 47) Viljoen, E.C., Jordaan, W.A., du Plessis, J., “On the Determination of Ternary Segregation Parameters”, Vacuum, 61(2–4), 141–144, (2001) (Experimental, Interface Phenomena, 5) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
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Fe–Mo–Ni
1
Iron – Molybdenum – Nickel Volodymyr Ivanchenko, Tetyana Pryadko
Introduction Magnetically soft alloys of the Fe-Ni system are used extensively in instrument fabrication owing to a favorable combination of a number of properties: single-phase alloys in a stable condition over a wide concentration range, an absence of allotropic transformations which can cause development of internal stresses, and the existence of ordering with the possibility of controlling the degree of ordering through alloying. However, alloys of the Fe-Ni system have low strength and wear resistance which limit their use, for example as magnetic recording head material. One of the most widespread methods for improving the mechanical properties of these alloys is by alloying with a third element, such as Mo, in order to form solid solutions and to impart an improvement in mechanical properties [1992Sos]. The Fe-Mo-Ni system is the basis for the development of new permanent magnetic materials with excellent magnetic shielding performance up to 4.2 K. Alloys of the Fe-Mo-Ni system exhibit high ductility. The constitutional diagram was first examined by [1934Koe], reproduced by [1949Jae], then by [1952Das] at 1200°C, diagram reproduced by [1973Wad]. Primary crystallization fields, and the resultant equilibria in the solid state after solidification were critically evaluated by [1984Ray, 1988Ray] which pointed out that knowledge of the binary systems Fe-Mo and Mo-Ni upon which the interpretation of the equilibria was based, was at that time seriously incomplete. Isothermal sections between 800 and 1250°C, and vertical sections were investigated by [1989Oda, 1990Gan] and the results reviewed by [1994Rag] who constructed the Fe-Mo-Ni reaction scheme from the solidification range down to 900°C. [1993Goz] reported about the melting temperature and the composition of a new γ+σ+τ ternary eutectic. [1997Sch] carried out isothermal holding tests in the two-phase region between liquidus and solidus. The liquidus line defining the threephase equilibrium (αδ+γ+L) was obtained from the intersecting points of the isotherms on the γ and αδ liquidus surfaces. [1997Sch] extrapolated their results to obtain a liquidus projection for the Fe“Mo2Fe3”-MoNi-Ni region. The data that [1997Sch] assumed for the Fe-Mo intermediate phases, and the liquid-solid equilibria are obsolete and do not agree with the accepted binary diagrams. Based on these results and on the known liquidus lines in the three boundary binary systems, as well as on the known concentration fields of the MoFe3-MoNi solid solutions, the melting equilibria were developed. To obtain the homogeneity region of the γ phase, series of specimens in which the concentration of Mo slightly exceeded the estimated solubility range for γ austenite were prepared and studied by [1989Gol]. The experimental examination has been combined with the Calphad method with the aim to get consistent phase diagrams. But there are serious discrepancies between calculations and experimental data of the previous authors [1952Das, 1953Das, 1980Loo]. The isothermal section at 1100°C was determined by [2006Yan]. It does not agree with the results of [1980Loo] and [1992Fri] due to a very low solubility of Ni in the μ phase and an absence of a ternary phase. The undercooling of Fe-Mo-Ni alloys under different cooling rates was studied by [1994ElM]. Investigations of the Fe-Mo-Ni phase relations, structures and thermodynamics are listed in Table 1. Binary Systems The Fe-Mo system is accepted from [Mas2]. The Fe-Ni and Mo-Ni systems are accepted from [2007Kuz] and [1991Sin], respectively. It should be noted, that some ternary diagrams in the present evaluation are accepted from the thermodynamic calculations of [1992Fri] and are based on the following binaries: FeMo from [1984Gui], Mo-Ni from [1990Fri], and Fe-Ni from [1990Xin]. The calculated Fe-Ni and Mo-Ni phase diagrams are practically the same as those accepted by us. In Fe-Mo the main difference between the calculated [1984Gui] and evaluated [Mas2] diagrams is in the temperature of the peritectoid formation of MoFe2. In other respects these two phase diagrams are nearly the same.
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2
Fe–Mo–Ni
Solid Phases The effect of Mo on the ordered phase FeNi3 was studied by [1951Che, 1968Art, 1970Gom]. Using neutron-diffraction techniques [1970Gom] found that solution of Mo in FeNi3 caused a remarkable decrease in the long-range-order parameter S. Along the compositional line representing the formula (Fe,Mo)Ni3, which runs from 75.9Ni-24.lFe (mass%) on the Fe-Ni axis, the order parameter S decreases from unity to 0.8 at 3.2 mass% Mo, 21.9 mass% Fe, and falls to zero at approximately 6.4 mass% Mo, 19.7 mass% Fe. These results are in general agreement with studies of [1951Che, 1968Art]. In accordance with [1980Loo], an approximated composition of the ternary phase τ is Mo53Fe11Ni36 (7.9Fe-65.lMo-27Ni, mass%). Crystal structure data of the unary, binary and ternary phases are presented in Table 2. Invariant Equilibria A reaction scheme has been derived by [1994Rag] who used the results of computations of [1992Fri]. Since the computed temperatures of most of the invariant reactions were not reported by [1992Fri], they must be regarded only as tentative values. The reaction scheme shown in Fig. 1 is taken mainly from [1994Rag], however the temperatures of the invariant reactions E1, U11 and P2 have been modified to be consistent with the isothermal sections accepted in the present evaluation, see section “Isothermal Sections” below. Liquidus, Solidus and Solvus Surfaces A liquidus projection has been computed by [1992Fri]. As the experimental data on the liquid equilibria at that time were scarce, the computed liquidus projection should be considered as tentative. It is given in Fig. 2. [1997Sch] extrapolated the results of their experimental study of the liquidus surface in the Fe rich corner to obtain the liquidus surface in the Fe-Ni-Fe3Mo2-MoNi subsystem. Since it did not included equilibria with the σ and R phases and was in contradiction with the results of [1993Goz], who reported the formation of the γ + σ + τ ternary eutectic with the melting point of 1312°C and the composition of 16.5Fe-34.5Mo-49Ni (at.%), results of the extrapolations obtained by [1997Sch] are omitted. The way of formation of this eutectic may be derived from [1992Fri], if the reaction sequence of L + σ⇌(Mo) + τ → L + (Mo)⇌MoNi + τ → L + τ⇌MoNi + γ would be replaced by the sequence: L + MoNi⇌(Mo) + γ → L + (Mo)⇌σ + γ → L⇌γ + σ + τ. Since this assumption needs experimental verification, the liquidus projection calculated by [1992Fri] as well as the reaction scheme constructed by [1994Rag] are not changed. The spinodal region was determined in Fe-Mo-Ni ternary alloys at 450 and 500°C within a composition range 10 to 16 at.% Mo and 2 to 14 at.% Ni by [1992Sat]. Isothermal Sections Constitutional diagrams proposed by [1934Koe] are in contradiction with the accepted binary phase diagrams. Therefore neither the high-temperature nor the low-temperature phase diagrams are reproduced in the present evaluation. The isothermal section at 1270°C was computed by [1992Fri] and given in Fig. 3. This section is in good agreement with the isothermal section at 1200°C presented by [1952Das]. The major features of the latter are the extensive γ solid solution, the considerable solid solubility of Ni in Mo6Fe7, and the existence, between μ and MoNi, of the ternary compound τ. The 1200°C isothermal section from [1952Das] is shown in Fig. 4 after small modifications to bring it in agreement with the accepted binary phase diagrams. A very close location of the γ corners of the τ + MoNi + γ and τ + μ + γ triangles presented by [1952Das] seems to be improbable. Therefore in Fig. 4 they are slightly separated. [1952Das] found τ to be in equilibrium with γ and μ, thus preventing μ from coexisting with MoNi. [1980Loo], however, with a τ composition slightly richer in Mo than that given by [1952Das], proposed that at 1100°C a two-phase region μ + MoNi is established, thus preventing equilibrium between τ and γ. In this connection, an alloy containing 9Fe-56.5Mo-34.5Ni (mass%) (12Fe-44Mo-44Ni (at.%)) was identified unambiguously by [1980Loo] as containing the μ, δ, and γ phases, so that the equilibria summarized in Fig. 5 appear to be well established. [1984Ray, 1988Ray] noted that, assuming the more Mo rich composition for τ, the equilibrium described at DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
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1100°C by [1980Loo] may be derived from that proposed at 1200°C by [1952Das] by the univariant reaction τ + γ ⇌ μ + MoNi and it is not impossible that such a reaction should occur within the temperature interval involved. This assumption is confirmed by the results of computation of the isothermal sections at 1270°C (Fig. 3) and 950°C (Fig. 6) performed by [1992Fri]. [1991Goz] and [1992Fri] showed that τ was stable at 1000°C. [1991Goz] did not find τ at 850°C and below. [1992Fri] found τ to be stable at 950°C, with a decreased range of homogeneity. The computed isothermal section at 950°C by [1992Fri] validated by the experimental results indicates that τ becomes unstable just below 950°C and decomposes through a ternary eutectoid reaction (E3). [1989Oda] presented the isothermal section at 800°C, it is redrawn from [1994Rag] in Fig. 7. The isothermal section at 1100°C presented by [2006Yan] is characterized by an unlikely low μ phase region and by an absence of the ternary phase. For this reason it is not included in the present evaluation. Temperature – Composition Sections The temperature-composition section presented by [1957Bec] at a constant Fe content of 70 mass% is doubtful whether it represents equilibrium state and is in contradiction with the results of [1952Das] and [1980Loo]. [1984Ray, 1988Ray] assumed that [1957Bec] mistook some non-equilibrium constituent for Fe2Mo. This assumption, combined with the postulate that Fe2Mo forms by the peritectoid reaction (P2) eliminates the wide γ + Fe2Mo + μ area presented on the polythermal section, given by [1957Bec] and leads to the observed sequence of phase fields below the peritectoid transformation αδ + μ, αδ + μ + Fe2Mo, αδ + Fe2Mo, αδ + γ + Fe2Mo, γ + Fe2Mo, γ + Fe2Mo + μ, γ + μ). The hypothetical polythermal section proposed by [1984Ray, 1988Ray] is shown in Fig. 8, however the temperature of the P2 transformation is shown below 950°C to be consistent with the isothermal section at 950°C from [1992Fri]. Thermodynamics [1983Kub] measured calorimetrically the heats of formation of ternary alloys in the temperature range of 900 to 1050°C. These results are presented in Table 3. Thermodynamic evaluation of the Fe-Mo-Ni system was performed by [1992Fri]. Based on the mixing volumes of alloys in the boundary binary systems [1993Leo] made an attempt to estimate mixing volume of the γ phase in ternary alloys. Notes on Materials Properties and Applications The effect of composition and heat treatment on mechanical properties of Fe-Mo-Ni alloys was studied by [1975Roa, 1978Ish, 1980Nos]. The Fe-Mo-Ni maraging steels have evoked a great deal of interest in the aerospace and aeronautical industries where they are used for missile framework and in landing chassis and in various other industries such as those manufacturing extrusion trains for light alloys, injection moulds for plastics, storage tanks for liquid gas, mechanical parts subjected to low temperatures, etc. [1977Ser]. Fe-Ni-Mo steels (Fe-13Ni-3Mo (mass%)) are used for cryogenic service [1978Ish]. Magnetically soft iron-nickel alloys containing 76–80% Ni have maximum magnetic permeability in weak fields. They are used for the fabrication of magnetic screens, small transformers, relays, sound-recorder heads, pulse transformers, and other contactless magnetic elements. The most widely used is the alloy 79NM, in which added molybdenum increases the specific electrical resistance, substantially raises the initial magnetic permeability and decreases its sensitivity to cooling rate, but in doing so lowers the magnetic induction [2001Vla]. An alloy containing 34.5 mass% Ni and 2 mass% Mo, exhibited permeability of 55000 at 25°C [1973Pfe]. The permeability and saturation magnetization of this alloy are comparable to nickel-iron alloys in the higher nickel range. The specific resistivity of the investigated alloy has the significantly higher value of 0.9 Ωαffl m. The obtained Curie temperature of 160°C is comparable with that of highly permeable ferrites. The achieved level of characteristics permits recommending alloys for production of precise instrument engineering parts, electronic devices and other facilities of automation and computer engineering. The electrodeposition of alloys has been gaining in importance owing to its practical utility in modern technology for special items which cannot be produced in any other way. Recent efforts have been directed to preparing Landolt-Börnstein New Series IV/11D4
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magnetic alloys which are of great interest in the sophisticated electronics industry related to rocketry, computers, space technology etc. Strain-insensitive magnetic foil of Fe-Mo-Ni alloy electrodeposited from a sulfamate bath has been reported to contain 18 mass% Fe, 9 mass% Mo and 73 mass% Ni with a coercive force of 0.05 Oe (~ 4 A·m–1) and a magnetic flux of 8200 G [1980Sri]. New, ductile, and cobalt-free magnet alloys based on the Fe-Mo-Ni system have been developed for potential substitution of some Co-Fe base magnet alloys. The new alloys are very ductile and can be processed into thin tape or fine wire. A Fe-20Mo-5Ni (mass%) alloy can be processed to exhibit magnetic properties of Br ≈ 10500 G, Hc ≈ 210 Oe (~16 800 A·m–1), and (BH)max ≈ l.l MGOe which are comparable to those of Vicalloy I (52Co38Fe-10V). A Fe-7 mass% Mo-5 mass% Ni alloy exhibits high remanence, square-loop, anisotropic magnet properties of Br ≈ 19200 G, Hc ≈ 30 Oe (~ 2 400 A·m–1), and Br/Bs ≈ 0.97 which are superior to those of classical Fe-Co-Ni magnetic alloys. The new Fe-Mo-Ni alloys consist of abundant elements and could substitute for some of the Co-Fe base magnet alloys which contain large amounts of expansive and strategic cobalt [1981Jin, 1984Tie]. Ni-14.5 mass% Fe-4.5 mass% Mo has excellent magnetic shielding performance at 4.2 K. It is required for superconducting quantum interference devices (SQUID) to avoid influences of electromagnetic noises [1984Suz]. [1988Mag] reported, that plastic deformation of Fe-20 mass% Mo-5 mass% Ni before annealing raises Br as well as Hc at higher degrees of deformation owing to the formation of a crystallographic texture and diminution in size of precipitates of the intermetallic phase. Miscellaneous The effect of alloying with Mo on the driving force of the martensitic transformation was examined by [2000Rag]. The atomic ordering in the FeNi3 alloyed with Mo was studied by [1968Ega, 1969Yeg1, 1969Yeg2, 1979Wen] by measuring temperature dependences of resistivity, saturation magnetization and temperature and composition dependencies of the Young modulus (E). It was shown that alloying with Mo leads to a depression of E with a minimum at 3.8 at.% Mo. At the Mo content below this value Mo atoms take part in ordering, due to the formation of a three-component superstructure [1969Yeg1]. The magnetic properties and hardness of Fe-Mo-Ni alloys has been studied by [1979Wen]. The paramagnetic-magnetically ordered transition and the effect of Mo content on the magnetic properties of martensite phase in Fe-30% Ni-x% Mo alloys have been investigated by in [2006Yas] AC susceptibility and Mössbauer spectroscopy. The results show that martensitic transformation temperature (Ms) indicates the magnetic transition from paramagnetic to magnetically ordered state. It is found that volume fraction of thermally induced martensite and Ms temperature decreased with the increasing Mo content [2006Yas]. Opportunities of the production of Fe-Mo-Ni alloys by means of powder metallurgy including mechanical milling were studied by [1986Rod, 1986Sar, 2001Vla, 2004Chi, 2004She]. Nanocrystalline and amorphous Fe40Ni38B18 Mo4, Fe49Ni46Mo5 and Fe42Ni40B18 alloys were synthesized by ball milling to study the effect of boron and molybdenum on the milling process and on the magnetic properties of Fe-Ni based alloys [2005Du]. [1969Yed] studied an maraging Fe-16.8Ni-5Mo (mass%) alloy and showed that under aging of martensite at 300–650°C the Laves phase Mo(Fe,Ni)2 is formed in the forms of dispersed particles. The rise of carbon content in Fe-Mo-Ni austenitic steel as well as rise of the austenitizing temperature leads to a decrease of the Ms (Martensite start) temperature [1996Gol, 1998Gol, 2000Gol]. Structural evolution of the alloy Fe10Ni-10Mo (mass%) by returning to a lower temperature from the point As was studied by [1975Kno]. The nuclear gamma resonance (NGR) has been used to study the influence of annealing on the fine atomic structure of the austenite of alloys Fe-(28-36) % Ni with additions of Mo and Si. Below 500°C treatment reduces the hyperfine magnetic field on the iron nucleus. A comparison of these results with known neutron diffraction and Mössbauer data leads to the conclusion that annealing at 300–500°C causes redistribution of the alloying atoms with formation of regions enriched with nickel and molybdenum atoms. After treatment at temperatures above 500°C the hyperfine magnetic field in those alloys has increased, which could either be due to a breakdown of concentration heterogeneities, or precipitation of a second phase [1980Rod]. A series of solidification experiments with different heat extraction methods were performed on an iron alloy 12 mass% Ni and 10 mass% Mo by [1994ELM]. The undercooling during the solidification process was measured as a function of the cooling rate. An undercooling of 550°C was achieved at a cooling rate of 7000 K·s–1. This undercooling is far below the solidus curve in the ternary phase diagram. DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
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Redistribution of alloying elements under aging in the Fe-16Ni-5Mo (mass%) alloy was studied by [1973Gru, 1980Fra]. The changes in the fine atomic structure of austenitic alloys Fe-(28-36) at.% Ni with additions of Mo were studied by NGR after annealing in the range 200–800°C, and the activation energy of the processes occurring below 500°C determined [1980Rod]. The main processes which take place in Fe-16 mass% Ni-6 mass% Mo at the early stages of decomposition of the supersaturated solid solution, filling of Cottrell atmospheres, formation of concentration inhomogeneities of molybdenum and their transformation into particles of a strengthening phase were investigated by [1987Val]. A quantitative analysis of the 57Fe hyperfine field distribution P(H) has been performed in quenched bcc disordered Fe-Mo-Ni alloys in order to determine the matrix depletion of Ni and Mo atoms that occurs after aging [1995Dje]. The small atomic fraction assigned to the precipitates, observed on the Mössbauer spectra after a 72-h aging at 450°C of Fe78.3Ni18.3Mo3.4 and Fe79.3Ni15.0Mo5.7 alloys, is corroborated by the small volume fraction of the nanocrystalline precipitates revealed by transmission electron microscopy (TEM). The two precipitated phases, (Fe-Ni)3Mo and ω (Fe-Ni)7Mo2, detected by TEM are consistent with the Mössbauer approach, with an overall Fe to (Ni,Mo) atomic ratio twice smaller for Fe79.3Ni15.0Mo5.7 than for Fe78.3Ni18.3Mo3.4, giving evidence for an evolution to depleted-Fe species. Table 1.
Investigations of the Fe-Mo-Ni Phase Relations, Structures and Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1934Koe]
Thermal analysis, Brinell hardness
1200, 800°C, Fe-Mo Fe3-MoNi-Ni, primary crystallization fields
[1952Das] [1953Das]
Light microscopy, XRD (X-ray diffraction)
1200°C, < 70 at.% Mo, isothermal section from the analysis of 105 alloys
[1957Bec]
Light microscopy, XRD
< 1350°C, 70 mass% Fe, 0-20 mass% Ni, isothermal sections
[1968Art]
Resistivity, Young modulus, XRD
Ni3(Fe,Mo) annealed at 800°C then cooled at the rate of 10°C·h–1
[1969Geo1, 1969Geo2]
Optical microscopy, ballistic magnetometry
24.3 mass% Ni-5.2 mass% Mo, annealing at 1260°C, then quenched. Martensitic transformation at (77–293 K)
[1969Mir]
Magnetometer, DTA, thermodynamic calculations
1100°C, 9.6 at.% Ni, < 3 at.% Mo
[1969Yed]
TEM and electron microdiffraction
16.8 mass% Ni, 5 mass% Mo, cooling in air from 1000°C, then aged at 300–650°C. Austenitization
[1969Yeg1]
Young modulus, saturation magnetization
75 at.% Ni, 17.1-24.3 at.% Fe, annealing at 250–600°C. Ordering in (Mo,Fe)Ni3
[1969Yeg2]
Resistivity, Young modulus
350–400°C, (Mo1–xFex)Ni3, ordering
[1970Cou]
TEM, electron microdiffraction
Ageing between 100 and 570°C, Fe-19 at.% Ni-5 at.% Mo, α→γ transformation
[1970Yed]
X-ray diffraction, Vickers hardness, resistivity
15.8 to 16.5 mass% Ni, < 7.8 mass% Mo. Annealed at 1000°C, then aged at 410–500°C. Austenitization (continued)
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Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1970Gom]
Neutron diffraction
Mox(FeNi3)1–x, x = 0.01–0.05, temperature not given, superstructure in FeNi3
[1973Gru]
Nuclear gamma resonance
16 mass% Ni, 5 mass% Mo Quenching from 1000°C, then aged at 420°C, redistribution of Mo atoms
[1974Gen]
Nuclear gamma resonance
12.2 at.%Ni, 6.2 at.% Mo quenched from 1000°C and aged at 450°C. Austenitization
[1975Kno]
XRD, dilatometry, DTA
10 mass% Ni, 10 mass% Mo, homogenized at 1200°C then quenched, α-γ transformation
[1976Cou]
Mössbauer spectroscopy
Fe-19 at.% Ni-3 at.% Mo, α→γ transformation
[1977Ser]
Dilatometry, DTA thermo magnetometry, thermoresistivity, XRD, small angle X-ray scattering, SEM, optical microscopy.
Fe-Mo-Ni alloys with mass% Mo up to 20 and mass% Ni up to 21.66 were prepared by sintering at 1400°C, αδ, μ, martensite
[1980Geo]
X-ray diffraction, dilatometry, temperature dependence of resistivity
Water quenching from 1150 and 1100°C, aging at 200–650°C for 2–30 min and up to 120 h, Fe70+xNi30–2xMox (x=0.5, 1, 1.5, 2, 2.5), γ→α transformation
[1980Loo]
Diffusion couple technique, optical microscopy, EPMA, X-ray diffraction
1100°C, phase equilibria Mo-FexNi1–x, structure of the τ phase
[1980Zam]
NGR, internal friction, optical microscopy, TEM, magnetometry
Quenching from 1150°C followed by tempering at 350, 450, 550, and 600°C for 1–300 h, Fe-24 mass% Ni-5 mass% Mo, γ→α
[1982Nik]
Optical microscopy, magnetic measurements using ballistic magnetometer
25.5 mass% Ni, 4.5 mass% Mo annealing at 1150°C then aged at 400–600°C. Kinetics of martensite transformation
[1983Sat]
TEM, X-ray diffraction, Vickers hardness
(12–18) at.% Mo, (2–14) at.% Ni, homogenized at 1000°C, hot-rolled at 1100°C, annealed at 1300–1350°C then quenched
[1987Val]
Method of positron lifetime measurement using positron source 22Na on a mylar substrate with activity of 2·108 Bq
16 mass% Ni, 6 mass% Mo, quenched from 950°C, heated at 100–550°C for 30 min, early stages of decomposition of supersaturated solid solution
[1989Gol]
XRD, DTA, EPMA, calculations
1550–1650°C, solid-liquid equilibria
[1989Oda]
Optical microscopy, SEM, EDAX, XRD
800–1250°C, 50 at.% Mo, diffusion couple
[1990Gan]
Optical microscopy, EMPA
1000°C, diffusion couple
[1991Goz]
Metallography, XRD, SEM, EDAX
700–1200°C, Fe-40 mass% Mo, 60 mass% Ni (continued)
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Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1992Fri]
Diffusion couple, optical microscopy, SEM, EDAX, XRD
Liquidus surface and isothermal sections at 950 to 1270°C
[1992Sat]
Hardness measurements, optical microscopy
< 16 at.% Ni, < 40at.%Mo, 450–600°C, spinodal decomposition
[1993Goz]
SEM, DTA, XRD
16.5 at.% Fe, 34.5 at.% Mo, 1300°C for 10 h, rapidly cooled to 1000°C and quenched
[1993Leo]
XRD
XNi=0.93, 0.90, 0.87, 0.84, 0.81, 0.78; XMo/(XFe+XMo) = from 0 to 1
[1994ELM]
Differential thermal analysis (DTA), optical microscopy, EPMA
Fe-12 mass% Ni-9.5 mass% Mo, influence of the cooling rate
[1995Dje]
Mössbauer, EPMA, TEM, influence of the cooling rate on the Mo segregation
20 to 30 at.% Ni, 3 to 5 at.% Mo, x = 0.2– 0.3, annealed at 1400°C, cold-rolled, aged at 450°C and cooled
[1996Gol]
Internal friction in the ultrasonic and hertz frequency spans, induction magnetometric analysis, DTA, mechanical tests
Fe-24.5Ni-5Mo (mass%) annealed at 1150°C, then quenched and aged at 350–600°C. Kinetics of martensitic transformation
[1997Sch]
Isothermal holding tests in the liquid + solid region, EPMA
1450–1530°C, partial liquidus surface
[1998Gol, 2000Gol]
Temperature dependent and amplitude dependent internal friction in temperature interval from 100 to 900 K
Fe-24.5Ni-5Mo (mass%) annealed at 900–1150°C then quenched. Martensitic transformations
[2006Yan]
Diffusion triple, EMPA
Fe-Mo-Ni, 1100°C
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
αδ, (αδFe,Mo)
cI2 Im 3m W
(δFe) 1538–1394 (αFe) < 912 (αMo) < 2623
Lattice Parameters [pm]
Comments/References
a = 293.15
dissolves up to 24.4 at.% Mo at 1449°C and 5.5 at.% Ni at 1514°C pure Fe at 1390°C [V-C2, Mas2]
a = 286.65
pure Fe at 25°C [Mas2]
a = 314.7
pure Mo at 25°C [Mas2] dissolves up to 35 at.% Fe at 1611°C and 3 at.% Ni at 1362°C 0.9 at.% Ni [1991Sin] ~20 at.% Fe [1982Fer]
a = 314.5 a = 311.8
(continued)
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Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
γ, (γFe,Ni)
cF4 Fm 3m Cu
(γFe) 1394–912
Lattice Parameters [pm]
Comments/References
a = 364.67
pure Fe at 915°C [Mas2] dissolves up to 2 at.% Mo at 1100°C pure Ni at 25°C [Mas2] dissolves up to 28 at.% Mo at 1317°C
a = 352.40 (Ni) < 1455 (εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, 13 GPa [Mas2]
λ, MoFe2 < 927
hP12 P63/mmc MgZn2
a = 474.4 c = 772.5
33.3 at.% Mo [Mas2]
R, Mo2Fe3 1488–1200
hR159 R 3 Co5Cr2Mo3
μ, Mo6Fe7 < 1370
hR39 R 3m Fe7W6
σ, MoFe 1611–1235
tP30 P42/mnm σCrFe
FeNi3 < 517
cP4 Pm 3m AuCu3
a = 355.23
63–85 at.% Ni [2007Kuz]
FeNi
tP2 P4/mmm AuCu
a = 357.9 c = 357.9
Metastable [2007Kuz]
MoNi < 1362
oP56 P212121 MoNi
MoNi3 < 910
oP8 Pmmm βCu3Ti
a = 506.4 b = 422.4 c = 444.8
75.1–75.6 at.% Ni [1991Sin]
MoNi4 < 870
tI10 I4/m MoNi4
a = 572.7 c = 356.6
80.1–80.6 at.% Ni [1991Sin]
MoNi2 (m)
oI6
a = 1091.0 c = 1935.4 a = 475.1 c = 2568
a = 921.8 c = 481.3
a = 910.8 ± 0.5 b = 910.8 ± 0.5 c = 885.2 ± 0.5
pffiffiffi a = 2p2ffiffiffia (fcc) b = 3 2a (fcc) c = a (fcc)
33.9–38.5 at.% Mo [1992Rag] [V-C2] 39.0–44.0 at.% Mo [1982Fer]
42.9–56.7 at.% Mo [Mas2] 50 at.% Mo [1992Rag]
46–48 at.% Ni [1991Sin] [V-C2]
Metastable 66.7 at.% Ni [1991Sin] (continued)
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Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
MoNi3 (m)
tI8 I4/mmm Al3Ti
-
Metastable 75 at.% Ni [1991Sin]
MoNi4 (m)
cF4
-
Metastable 80 at.% Ni [1991Sin]
Mo6Ni17 (m)
o**
a = 5.5 a (MoNi3) b = b (MoNi3) c = c (MoNi3)
Metastable 77 at.% Ni [1991Sin]
* τ, Mo53Fe11Ni36
o**
a = 909.1 b =1700.2 c = 479.5
[1980Loo]
Table 3.
Thermodynamic Data of Reaction or Transformation
Reaction or Transformation
Temperature [°C]
Quantity, per mol of atoms [kJ, mol, K]
Comments
0.4556Fe + 0.1213Ni + 0.4231Mo→μ
900
ΔmixH = 0.715
[1983Kub]
0.1519Fe + 0.3637Ni + 0.4844Mo→τ
1000
ΔmixH = 1.347
[1983Kub]
850
ΔmixH = 0.985
[1983Kub]
1000
ΔmixH = 2.180
[1983Kub]
0.080Fe + 0.425Ni + 0.495Mo→ (MoNi) 0.10Fe + 0.80Ni + 0.10Mo→γ
Table 4.
Investigations of the Fe-Mo-Ni Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1972Bor]
Ballistic method
Magnetic properties and texture of hot rolled Fe-Mo-Ni industrial alloys
[1973Pfe]
Ballistic method, Thompson bridge method
Permeability, Curie temperature, magnetostriction coefficient, specific resistivity
[1975Roa]
SEM, Vickers hardness, Instron machine
Hardness and tensile properties of Hastelloy B in temperature interval from room to 1200°C
[1978Ish]
SEM, tensile test
Mechanical properties of the welded specimens at ambient and low temperatures
[1980Nos]
XRD, TEM, mechanical tests
Ni-22.3Mo-5.9Fe (at.%) single crystal quenching from 1100°C, then aged at 700°C, formation of ordered MoNi2 (continued)
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Fe–Mo–Ni
Reference
Method/Experimental Technique
Type of Property
[1981Jin]
Magnetic permeability, saturation magnetization
Magnetic properties
[1983Sat]
SEM, standard ballistic technique, dilatometry, electrical resistivity
Magnetic, heat and electrical properties
[1984Luz]
Measurements of HC, magnetization and remanent magnetization
Magnetic properties of Fe-20Mo-5Ni (mass%)
[1984Tie]
Yield strength, coercitivity, saturation magnetization
20 mass% Ni, 4 to 6 mass% Mo, Cobalt free semi hard magnet
[1988Mag]
Optical microscopy, TEM, X-ray diffraction, ballistic method in solenoid
coercive force HC, residual induction Br, energy product (BH)m
[2004She]
XRD, SEM, TEM, grain size, coercitivity
15 mass% Fe, 5 mass% Mo, high energy ball milling, soft magnetic material
[2005Cha]
Light microscopy, SEM, mechanical tests: tensile and axial fatigue tests on a servohydraulic frame.
The microstructure and mechanical properties of sintered Fe-0.85 mass% Mo-Ni steels
[2005Du]
XRD, DSC, SEM, TEM, coercitivity, saturation magnetization
Fe49Ni46Mo5 (in at.%), mechanical alloying, Nanostructured magnetic material
[2006Yas, 2007Yas]
AC magnetic susceptibility, Mössbauer spectroscopy, TEM
Fe, 30 mass% Ni, < 5 mass% Mo annealed at 1100°C and quenched. Martensitic transformations
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Fig. 1a. Fe-Mo-Ni.
Reaction scheme, part 1
Fe–Mo–Ni
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Fe–Mo–Ni
Fig. 1b. Fe-Mo-Ni.
Reaction scheme, part 2
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Landolt-Börnstein New Series IV/11D4
Fe–Mo–Ni
Fig. 2. Fe-Mo-Ni.
Landolt-Börnstein New Series IV/11D4
13
Computed tentative liquidus projection
MSIT®
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
14
Fig. 3. Fe-Mo-Ni.
Fe–Mo–Ni
Isothermal section at 1270°C
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mo–Ni
Fig. 4. Fe-Mo-Ni.
Landolt-Börnstein New Series IV/11D4
15
Isothermal section at 1200°C
MSIT®
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
16
Fig. 5. Fe-Mo-Ni.
Fe–Mo–Ni
Computed isothermal section at 1100°C
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mo–Ni
Fig. 6. Fe-Mo-Ni.
Landolt-Börnstein New Series IV/11D4
17
Computed isothermal section at 950°C
MSIT®
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
18
Fig. 7. Fe-Mo-Ni.
Fe–Mo–Ni
Isothermal section at 800°C
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mo–Ni
Fig. 8. Fe-Mo-Ni.
Landolt-Börnstein New Series IV/11D4
19
Hypothetical equilibria involving α, γ, μ and Fe2Mo phases at 70 mass% Fe
MSIT®
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
20 References [1934Koe] [1949Jae] [1951Che] [1952Das]
[1953Das]
[1957Bec]
[1968Art]
[1968Ega]
[1969Geo1]
[1969Geo2]
[1969Mir]
[1969Yed]
[1969Yeg1]
[1969Yeg2]
[1970Cou]
[1970Yed]
Fe–Mo–Ni
Köster, W., “The Iron-Molybden-Nickel System” (in German), Arch. Eisenhuettenwes., 8, 169–171 (1934) (Experimental, Phase Diagram, Phase Relations, 5) Jänecke, E., “Ni-Fe-Mo” (in German) in “Kurzgefasstes Handbuch Aller Legierungen”, Winter Verlag, Heidelberg, 640–641 (1949) (Phase Diagram, Phase Relations) Chevenard, P., Josso, E., (in French), Compt. Rend. Acad. Sci. Paris, 233, 539–541 (1951), as quoted in [1984Ray] Das, D.K., Rideout, S.P., Beck, P.A., “Intermediate Phases in the Mo-Fe-Co, Mo-Fe-Ni and Mo-Ni-Co Ternary Systems”, Trans. AIME, 194(10), 1071–1075 (1952) (Crys. Structure, Experimental, Morphology, Phase Diagram, 23) Das, D.K., Beck, P.A., “Survey of Portions of the Iron-Nickel-Molybdenum and Cobalt-IronMolybdenum Ternary Systems at 1200°C”, Nat. Adv. Com. Aeronautic, Techn. Note, 2896, 1–56 (1953) (Crys. Structure, Experimental, Morphology, Phase Diagram, 15) Bechtoldt, C.J., Vacher, H.C., “Phase Diagram Study of Alloys in the Fe-Cr-Mo-Ni System”, J. Res. Nat. Bur. Stand., 58(1), 7–19 (1957) (Crys. Structure, Experimental, Morphology, Phase Diagram, 29) Artsishevskaya, L.F., Ibragimov, E.A., Selissky, Ya.P., Sorokin, M.N., “Investigation of the Atomic Ordering in Ni-Fe-Cr and Ni-Fe-Mo Alloys.”, Russian Met. (4), 102–107 (1968), translated from Izv. Akad. Nauk SSSR, Met., (4), 158–164 (1968) (Crys. Structure, Phase Relations, Electr. Prop., Mechan. Prop., 21) Eganyan, I.L., Selissky, Ya.P., “Investigation of the Effects of Atomic Ordering in Ternary Solid Solutions on the Basis of Ni3Fe Alloyed by VI Group Elements - Chromium, Molybdenum, Tungsten” (in Russian), Akad. Nauk Ukr. SSR, Metallofizika, (20), 100–104 (1968) (Crys. Structure, Phase Relatiuons, Physical Prop., Experimental, 3) Georgieva, I.Ya., Izotov, V.I., Nikitina, I.I., Khandarov, P.A., “Structural Singularities of Athermic and Isothermic Martensite in an Fe-Ni-Mo Alloy” (in Russian), Fiz. Met. Metalloved., 27(6) 1129–1130 (1969) (Morphology, Phase Relations, Experimental, 6) Georgieva, I.Ya., Nikitina, I.I., “Isothermal and Athermal Martensite Transformation in Fe-Ni-Mo Alloy”, Dokl. Akad. Nauk SSSR, 186(1), 85–88 (1969) (Morphology, Phase Relations, Experimental, 6) Mirzayev, D.A., Shteynbeg, M.M., Goykhenberg, Yu.N., “Influence of Manganese, Chromium and Molybdenum on Martensitic Transformation in Alloys on Iron-Nickel Base”, Phys. Met. Metallogr., 28(1), 162–169 (1969), translated from Fiz. Met. Metalloved., 28 (1), 152–159, 1969 (Phase Relations, Thermodyn., Experimental, Calculation, 12) Yedneral, A.F., Perkas, M.D., “Structural Changes in Martensite of an Fe-Ni-Mo Alloy as a Result of Ageing”, Phys. Met. Metallogr., 28(5), 101 (1969), translated from Fiz. Met. Metalloved., 28(5), 862–871, (1969) (Morphology, Phase Relations, Experimental, 11) Yeganyan, I.L., Selisskiy, Ya.P., “A Study of the Effects of Atomic Ordering in Ternary Solid So1utions of Cr, Mo, and W in Ni3Fe”, Phys. Met. Metallogr., 27(2), 18–25 (1969), translated from Fiz. Met. Metalloved., 27(2), 210–218 (Phase Relations, Phys. Prop., Experimental, 8) Yeganyan, I.L., Selissky, Ya.P., “ΔE Effect in the Ni3(Fe,Cr) and Ni3(Fe,Mo) Alloys”, Phys. Met. Metallogr., 27(4), 199–201 (1969), translated from Fiz. Met. Metalloved., 27(4), 763– 764, 1969 (Phase Relations, Phys. Prop., Experimental, 5) Courrier, R., “Kinetic Study of the Isothermal Transformation of the Martensite During Ageing of a Ternary Alloy with 76% Fe, 19% Ni and 5% Mo” (in French), Compt. Rend. Acad. Sciences, Serie C, 271(25) 1514–1517 (1970) (Crys. Structure, Morphology, Phase Relations, Experimental, Kinetics, 12) Yedneral, A.F., Perkas, M.D., “Investigation of Kinetics of Ageing of Martensite in Fe-NiMo Alloys” (in Russian), Fiz. Met. Metalloved., 30(2), 418–425 (1970) (Crys. Structure, Phase Relations, Experimental, Kinetics, Electr. Prop., Mechan. Prop., 17)
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mo–Ni [1970Gom]
[1972Bor]
[1973Gru]
[1973Pfe]
[1973Wad] [1974Gen]
[1975Roa]
[1975Kno]
[1976Cou]
[1977Ser]
[1978Ish]
[1979Wen]
[1980Fra]
[1980Geo]
[1980Loo]
[1980Nos]
Landolt-Börnstein New Series IV/11D4
21
Goman’kov, V.I., Puzey, I.M., Mal’tsev, E.I., “Effect of Alloying Elements on the Superstructure of Ni3Fe” (in Russian), Dokl. Akad. Nauk SSSR, 194(2), 309–311 (1970) (Crys. Structure, Experimental, Magn. Prop., 6) Borodkina, M.M., Pesin, V.S., Smirnova, L.G., Sosnin, V.V., Starostin, Yu.V., “Magnetic Properties and Texture of Thin Strips of Nickel-Iron-Molybdenum Alloys” (in Russian), Fiz. Met. Metalloved., 33(6), 1188–1194 (1972) (Morphology, Magn. Prop., Experimental, 12) Gruzin, P.L., Rodionov, Yu.L., Li, Yu.A., Edneral, A.F., Zhukov, O.P., Perkas, M.D., “Redistribution of Alloying Elements During the Recovery in Maraging Fe-Ni-Mo and Fe-Ni-Co-Mo Alloys” (in Russian), Fiz. Met. Metalloved. (in Russian), 36(2), 423–427 (1973) (Phase Relations, Experimental, 10) Pfeifer, F., Cremer, R., “Highly Permeable Nickel-Iron-Molybdenum Alloys Containing 33 to 37% Nickel” (in German), Z. Metallkd., 64(5), 362–365 (1973) (Phase Relations, Magn. Prop., Experimental, 9) Wada, T., “Iron-Molybdenum-Nickel” in “Metals Handbook”, Lyman, T. (Ed), ASM, Metals Park, Ohio, 8, 431–431 (1973) (Phase Diagram, Review, 1) Genin, J.M., Lecaer, G., Maitrepi, P., Thomas, B.J., “Mössbauer Study of Precipitation in Fe-Ni-Mo and Fe-Ni-Co-Mo Maraging Alloys”, Scr. Metall., 8(1), 15–22 (1974) (Electronic Structure, Phase Relations, Experimental, 17) Roan, D.-Y., Brooks, C.R., “The Effect of Ageing on the Mechanical Properties of a NickelMolybdenum-Iron Alloy (Hastelloy B)”, Metal. Trans. A, 6A(9), 1829–1831 (1975) (Morphology, Mechan. Prop., Experimental, 7) Knosp, B., Lepine, G., Cizeron, G., “Structural Evolution of the Alloy Fe-10Ni-10Mo by Returning to a Lower Temperature from the Point As” (in French), Scr. Metall., 9(9), 901– 906 (1975) (Crys. Structure, Phase Relations, Experimental, Kinetics, 7) Courrier, R., “Study of the Structure Changes of Maraging Alloy Fe-19 at.% Ni-3 at.%Mo” (in French), Mem. Sci. Rev. Metall., 73(2), 101–116 (1976) (Phase Relations, Experimental, Kinetics, 26) Servant, C., Lacombe, P., “Structural Transformations Produced During Tempering of Fe-NiCo-Mo Alloys”, J. Mater. Sci., 12, 1807–1826 (1977) (Experimental, Morphology, Phase Diagram, 42) Ishikawa, K., Maruyama, N., “Fracture and Strength of MIG Welded Fe-13%Ni-3%Mo Alloy for Cryogenic Service”, Cryogenics, 18(10), 585–589 (1978) (Morphology, Experimental, Mechan. Prop., 11) Wenchong, X., Junjian, W.,. Xiujin, S., Zhihua, L., “An Investigation of Magnetic Ni-Fe-Mo and Ni-Fe-Nb Alloys” (in Chinese), Acta Metall. Sin., 15(2), 252–258 (1979) (Morphology, Phase Relations, Experimental, Mechan. Prop., 8) cited from abstract Frackowiak, J.E., Morawiec, H., Panek, T.J., Uhlig, M.R., “Mössbauer Effect Study of K-effect in Ni-Fe-Mo Alloy”, J. Phys. Colloq., 41(C-1) C1373–1374 (1980) (Phase Relations, Experimental, Magn. Prop., 6) Georgieva, I.Ya., Matyushenko, L.A., “Effect of Heat Treatment on the Kinetics of the TwoStage Martensitic Transformation in Fe-Ni-Mn and Fe-Ni-Mo Alloys”, Met. Sci. Heat Treat., 22(5–6), 311–313 (1980), translated from Metall. Term. Obrab. Met., 22(5), 3–5 (1980) (Phase Relations, Experimental, Kinetics, Electr. Prop., 7) Van Loo, F.J.J., Bastin, G.F., Vrolijk, J.W.G.A., Hendricks, J.J.M., “Phase Relations in the Systems Fe-Ni-Mo, Fe-Co-Mo and Ni-Co-Mo at 1100°C”, J. Less-Common Met., 72, 225–230 (1980) (Crys. Structure, Experimental, Kinetics, Phase Diagram, 11) Nosova, G.I., Polyakova, N.A., “Phase Composition, Strength Parameters and Dislocation Structure of Alloy Ni-Mo-Fe”, Phys. Met. Metallogr., 50(5), 168–171 (1980), translated from Fiz. Met. Metalloved., 50(5), 1090–1093 (1980) (Crys. Structure, Morphology, Phase Relations, Experimental, Mechan. Prop., 9)
MSIT®
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
22 [1980Rod]
[1980Sri] [1980Zam]
[1981Jin] [1982Fer] [1982Nik]
[1983Kub]
[1983Sat]
[1984Gui]
[1984Luz]
[1984Ray]
[1984Suz]
[1984Tie] [1985Xin] [1986Rod]
[1986Sar]
[1987Val]
Fe–Mo–Ni Rodionov, Y.L., Isfandiyarov, G.G., Zambrzhitskiy, V.N., “Influence of Annealing on the Redistribution of Atoms in the Austenite of Fe-Ni-Mo and Fe-Ni-Si”, Phys. Met. Metallogr., 49(2), 94–100 (1980), translated from Fiz. Metal. Metalloved., 49(2), 335–341 (1980) (Crys. Structure, Experimental, 8) Srivastava, S.C., “Electrodeposition of Ternary Alloys: Developments in 1972–1978”, Surf. Technol., 10, 237–257 (1980) (Electrochemistry, Magn. Prop., Review, 121) Zambrzhitskiy, V.N., Maksimova, O.P., Gruzin, P.L., Rodionov, Y.L., Utevskiy, L.M., Isfandiyarov, G.G., Pankova, M.N., Seleznev, V.N., “Annealing Influence on Austenite Structure and Martensite-Transformation in Iron-Nickel-Molybdenum Alloy with Isothermal Transformation Kinetics” (in Russian), Fiz. Met. Metalloved., 49(4), 776–787 (1980) (Morphology, Phase Relations, Experimental, Kinetics, 20) Jin, S., Tiefel, T.H., “New Ductile Fe-Mo-Ni Magnet Alloys”, J. Appl. Phys., 52(3), 2503– 2505 (1981) (Experimental, Magn. Prop., 7) Fernandez-Guillermet, A., “The Fe-Mo (Iron-Molybdenum) System”, Bull. Alloy Phase Diagrams, 3(3), 359–367 (1982) (Crys. Structure, Phase Diagram, Review, 40) Nikitina, I.I., Rozhkova, A.S., “Martensite γ-α Transformations in Fe-Ni-Mo Alloys with Various Purity Degrees” (in Russian), Fiz. Met. Metalloved., 53(6), 1138–1142 (1982) (Morphology, Phase Relations, Kinetics, Magn. Prop., Experimental, 14) Kubaschewski, O., Hoster, T., “Enthalpy of Formation and Transformation in Binary and Ternary Systems of the Metals Iron, Cobalt, Nickel and Molybdenum” (in German), Z. Metallkd., 74(9), 607–609 (1983) (Experimental, Phase Diagram, Thermodyn., 8) Sato, S., Urushikara, F., Suzuki, T., “Effect of Nickel Addition on the Phase Decomposition of Iron-Molybdenum Alloys”, Scr. Metall., 17, 1391–1394 (1983) (Crys. Structure, Morphology, Phase Diagram, Experimental, 11) Guillermet, A.F., Consejo National de Investigationes Cientificas y Tecnicas, Centro Atomico Bariloche, San Carlos de Bariloche, Argentina, personal communication, 1984, as quoted in [1992Fri] Luzhinskaya, M.G., Shilova, N.F., Shur, Ya.S., “Investigation of Magnetization Reversal Processes in Fe-Mo-Ni Alloy” (in Russian), Fiz. Met. Metalloved., 57(4), 821–824 (1984) (Morphology, Experimental, Magn. Prop., 3) Raynor, G.V., Rivlin, V.G., “Phase Equilibria in Iron Ternary Alloys. XV. Critical Evaluation of Ternary Alloys of Iron and Molybdenum with Cobalt, Chromium, Manganese, and Nickel”, Int. Metall. Rev., 29(5), 329–375 (1984) (Assessment, Phase Diagram, Review, 70) Suzuki, Y., Horikoshi, E., Niwa, K., “Magnetic Properties and Ferromagnetic Shielding of Ni-Mo-Fe Alloys at Cryogenic Temperatures”, Fujitsu Sci. Tech. J., 20(2), 167–177 (1984) (Phase Relations, Experimental, Magn. Prop., 7) Tiefel, T.H., Jin, S., “Microduplex Fe-Ni-Mo Semihard Magnet Alloys”, J. Appl. Phys., 55(6), 2112–2114 (1984) (Phase Relations, Experimental, Magn. Prop., 3) Xing, Z.S., Gohil, D.D, Dinsdale, A.T., Chart, T.G., DMA(A)103, Report National Physics Laboratory, London (1985) as quoted in [1992Fri] Rodzinak, D., “Contact Fatigue of Materials Produced by Powder Metallurgy”, Met. Mater., 24(1), 50–56 (1986), translated from Kovove Mater., 24(1), 102–113 (1986) (Morphology, Experimental, Mechan. Prop., 22) Sarkisyan, L.E., “Structure and Properties of Iron-Nickel Powder Metallurgy Alloys”, Sov. Powder Met.-Metal Cer., 25(11), 931–936 (1986), translated from Poroshk. Metall., Ukr. SSR, 25(11), 79–84 (1986) (Morphology, Experimental, Electr. Prop., Magn. Prop., Phys. Prop., 4) Valuyev, N.P., Zhikharev, A.N., Moysh, Yu.V., Perkas, M.D., Rusanenko, V.V., Shcherbedinskiy, G.V., “Position Lifetime Investigation of the Early Stages of Ageing of Martensite of an Fe-Ni-Mo Alloy”, Phys. Met. Metallogr. 64(5), 117–120 (1987), translated from Fiz. Met. Metalloved., 64(5), 957–960 (1987) (Phase Relations, Experimental, Kinetics, 8)
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mo–Ni [1988Mag]
[1988Ray] [1989Gol]
[1989Oda]
[1990Fri] [1990Gan]
[1991Goz]
[1991Sin]
[1992Fri]
[1992Rag]
[1992Sat]
[1992Sos]
[1993Goz]
[1993Leo]
[1994ElM] [1994Rag] [1995Dje]
Landolt-Börnstein New Series IV/11D4
23
Magat, L.M., Makarova, G.M., Lapina, T.P., Belozerov, Ye.V., Maykov, V.G. et. al., “Influence of Plastic Deformation and Alloying on the Magnetic Properties and Structure of Alloys for Permanent Magnets Based on the Fe-Mo-Ni System”, Phys. Met. Metallogr., 65(6), 28– 33 (1988), translated from Fiz. Met. Metalloved., 65(6), 1075–1080 (1988) (Crys. Structure, Experimental, Magn. Prop., Morphology, 3) Raynor, G.V., Rivlin, V.G., “Fe-Mo-Ni” in “Phase Diagrams of Ternary Iron Alloys”, Inst. Metals, London, 4, 392–397 (1988) (Phase Diagram, Phase Relations, Review, 13) Golczewski, J.A., Lukas, H.L., Bamberger, M., Dirnfeld S.F., Gozlan, E., Klodt, J., Prinz, B., “Phase Diagrams of the Ni-Fe-Mo and Ni-Cr-Mo Ternary Systems-Experiments and Thermodynamic Calculations as a Basis of Superalloy Development” in “Adv. Mater. Processes”, Proc. 1st Eur. Conf., 1, 365–370 (1989) (Thermodyn., Experimental, Calculation, Phase Diagram, 7) Odagiri, S., “Phase Diagram of the Fe-Rich Side in the Fe-Mo-Ni System” (in Japanese), Nippon Kogyo Daigaku Kenkyu Hokoku, 79(1), 91–94 (1989) (Crys. Structure, Morphology, Phase Diagram, Experimental), as quoted in [1994Rag] Frisk, K., “A Thermodynamic Evaluation of the Mo-Ni System”, Calphad, 14(3), 311–320 (1990) (Thermodyn., Calculation, Assessment, #) as quoted in [1992Fri] Gan, W., Jin, Z., “Phase Equilibria in the Mo-Ni-Co and Mo-Ni-Fe System at 1000°C”, J. Less-Common Met., 160, 29–33 (1990) (Crys. Structure, Morphology, Phase Diagram, Experimental, 5) Gozlan, E., Bamberger, M., Dirnfeld, S.F., Prinz, B., Klodt, J., “Topologically Close-Packed Precipitations and Phase Diagrams of Ni-Mo-Cr and Ni-Mo-Fe and of Ni-Mo-Fe with Constant Additions of Chromium”, Mater. Sci. Eng. A, 141, 85–95 (1991) (Crys. Structure, Experimental, Morphology, Phase Diagram, 12) Singleton, M.F., Nash, P., “Mo-Ni (Molybdenum-Nickel)” in “Phase Diagrams of Binary Nickel Alloys”, Nash P., (Ed.), ASM Intl., Materials Park, OH, 207–212 (1991) (Crys. Structure, Phase Diagram, Thermodyn., Assessment, #, 48) Frisk, K., “An Experimental and Theoretical Study of the Phase Equilibria in the Fe-Mo-Ni System”, Metall. Trans. A, 23A(2), 639–649 (1992) (Crys. Structure, Morphology, Phase Diagram, Thermodyn., Experimental, #, 26) Raghavan, V., “The Fe-Mo (Iron-Molybdenum) System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Metals, Calcutta, 6A, 36–36 (1992) (Crys. Structure, Phase Diagram, Review, 1) Sato, S., Urushihara, F., Suzuki, T., “Metastable Miscibility Gap of Fe-Mo-Ni Ternary Alloy System” (in Japanese), J. Jpn. Inst. Met., 56(4), 378–383 (1992) (Experimental, Phase Diagram, Thermodyn., 20) Sosnin, V.V., Glezer, A.M., Zhigalina, O.M., “Structure and Properties of Microcrystalline Alloys of the System Ni-Fe-Nb(Mo)”, Met. Sci. Heat Treat., 34(3–4), 195–202 (1992) (Crys. Structure, Morphology, Phase Relations, Experimental, 12) Gozlan, E., Dirnfeld, S.F., Bamberger, M., Klodt, J., Prinz, B., “Liquid/Solid Equilibrium and Segregation Behavior of Ni-Fe-Mo-Cr Alloys”, Z. Metallkd., 84, 776–780 (1993) (Morphology, Phase Diagram, Phase Relations, Experimental, 9) Leonov, V.V., Nikiforov, G.A., Bel’mach, E.Yu., “Relationships of Mixing Volumes in Binary and Ternary Ni-Fe-Mo Alloys”, Russ. Metall., (1) 41–44 (1993), translated from Dokl. Ross. Akad. Nauk, Metally, (1) 49–52 (1993) (Crys. Structure, Thermodyn., Experimental, 3) El-Mahallawy, N., Taha, M., Fredriksson, H., “Rapid Solidification of Fe-Ni-Mo Alloys”, Mater. Sci. Eng. A, 179, 587–591 (1994) (Thermodyn., Experimental, 6) Raghavan, V., “Fe-Mo-Ni (Iron Molybdenum-Nickel)”, J. Phase Equilib., 15(6), 622–626 (1994) (Phase Relations, Phase Diagram, Review, 7) Djega-Mariadassou, C., Bessais, L., Servant, C., “Nanocrystalline Precipitates Formed by Aging of BCC Disordered Fe-Ni-Mo Alloys”, Phys. Rev. B: Condens. Matter, 51(14), 8830–8840 (1995) (Crys. Structure, Thermodyn., Experimental, Phys. Prop., 16)
MSIT®
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
24 [1996Gol]
[1997Sch] [1998Gol]
[2000Gol]
[2000Rag]
[2001Vla]
[2003Rag] [2004Chi]
[2004She]
[2005Cha]
[2005Du]
[2006Yan]
[2006Yas]
[2007Yas]
[2007Kuz]
[Mas2] [V-C2]
Fe–Mo–Ni Golovin, S.A., Golovin, I.S., “Inelastic Effects and Martensite Transformations in Fe-Ni-Mo Alloys” (in Russian), Fiz. Met. Metalloved., 82(2), 71–81 (1996) (Phase Relations, Experimental, Kinetics, 23) Schürmann, E., Djurdjevic, M., Nedeljikovic, L., “Melting Equilibria in the Iron-Rich Corner of the Fe-Ni-Mo System”, Steel Research, 68, 383–387 (1997) (Phase Relations, Experimental, 8) Golovin, S.A., Sergantova, G.V., Troitsky, I.V., “Internal Friction in Fe-Ni-Mo Steels” (in Russian), Izv. Akad. Nauk., Ser. Fiz., 62(7), 1369–1376 (1998) (Phase Relations, Experimental, Kinetics, 19) Golovin, S.A., “Diffusive and Dislocational Anelasticity in Alloyed Austenite of Fe-Ni-Mo Steels”, Bull. Russ. Acad. Sci., Physics, 64(9), 1336–1341 (2000) (Phase Relations, Experimental, Kinetics, 14) Raghavan, V., “The Martensitic Fransformation in Fe-Ni and Fe-Ni-X Alloys at Subzero Temperatures”, Trans. Indian Inst. Met., 53(6), 597–603 (2000) (Phase Relations, Theory, Kinetics, 44) Vlasova, O.V., Panasyuk, O.A., Maslyuk, V.A., “Sintered Metals and Alloys/Preparation and Properties of Magnetically Soft Powder Metallurgy Alloy 79NM”, Powder Metall. Met. Ceram., 40(9–10), 497–500 (2001) (Phase Relations, Experimental, Magn. Prop., 3) Raghavan, V., “Fe-Mo-Ni (Iron-Molybdenum-Nickel)”, J. Phase Equilib., 24(3), 267–268 (2003) (Crys. Structure, Phase Diagram, Review, 8) Chicinas, I., Pop, V., Isnard, O., “Synthesis of the Supermalloy Powders by Mechanical Alloying”, J. Mater. Sci., 39(16–17), 5305–5309 (2004) (Crys. Structure, Morphology, Phase Relations, Experimental, 16) Shen, Y., Hng, H.H., Oh, J.T., “Synthesis and Characterization of High-Energy Ball Milled Ni-15%Fe-5%Mo”, J. Alloys Compd., 379(1–2), 266–271 (2004) (Morphology, Experimental, Thermodyn., Phys. Prop., 24) Chawla, N., Deng, X., “Microstructure and Mechanical Behavior of Porous Sintered Steels”, Mater. Sci. Eng. A, 390(1–2), 98–112 (2005) (Morphology, Phase Relations, Thermodyn., Experimental, Kinetics, Mechan. Prop., Phys. Prop., 33) Du, S.W., Ramanujan, R.V., “Mechanical Alloying of Fe-Ni Based Nanostructured Magnetic Materials”, J. Magn. Magn. Mater., 292, 286–298 (2005) (Crys. Structure, Morphology, Experimental, Magn. Prop., Nano, 27) Yan, F., Richu, W., Danhua, W., Qingyong, L., Shengming, W., “Determination of Isothermal Sections of the Ni-Fe-Mo Ternary System at 1373 K” (in Chinese), Rare Met. (China), 35(8), 1228–1230 (2006) (Phase Relations, Phase Diagram, Experimental, 8) Yasar, E., Gungunes, H., Kilic, A., Durlu, T.N., “Effect of Mo on the Magnetic Properties of Martensitic Phase in Fe-Ni-Mo Alloys”, J. Alloys Compd., 424, 51–54 (2006) (Crys. Structure, Morphology, Phase Relations, Experimental, Magn. Prop., 22) Yasar, E., Gungunes, H., Akturk, S., Durlu, T.N., “New Observations on the Formation of Athermal Martensite in Fe-Ni-Mo Alloys”, J. Alloys Compd., 428, 125–129 (2007) (Crys. Structure, Experimental, Morphology, Phase Relations, 30) Kuznetsov, V., “Fe-Ni (Iron-Nickel)”, Binary Evaluation Program as Published in MSIT Workplace, Effenberg, G. (Ed.), MSI, Materials Science International Services GmbH, Stuttgart; to be published, (2007) (Phase Diagram, Phase Relations, Crys. Structure, 41) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_26 # Springer 2008
MSIT®
Landolt-Börnstein New Series IV/11D4
Fe–Mo–O
1
Iron – Molybdenum – Oxygen Pierre Perrot
Introduction The Fe-Mo-O system presents an interest because it appears in the preparation of ferromolybdenum alloys by metallothermic reduction of MoO2. Some Fe-Mo oxides have also important industrial applications by forming a protective coating over molybdenum steels. On an another hand, the molybdate Fe2(MoO4)3 is the active component in Fe-Mo based oxidation catalysts. Iron and molybdenum oxides are “easily reduced” oxides, which means that equilibria between solid phases are easily obtained under H2-H2O or CO-CO2 atmospheres. Such atmospheres give access not only to the phase compositions, but to the oxygen potential at equilibrium with an accuracy better than that obtained from electromotive force measurements. The main experimental investigations on phase equilibria and thermodynamics are gathered in Table 1. No Calphad assessment was carried out. Binary Systems The Fe-Mo binary system is accepted from the thermodynamical assessment carried out by [1982Fer] and reproduced in [Mas2]. The Fe-O system is accepted from [Mas2], which reproduces the fundamental work of [1945Dar, 1946Dar]. Two careful assessments may be found in [1991Sun, 1995Kow]. They reproduce well the experimental equilibria and differ by the description of the (Fe,O) liquid solution. The Mo-O binary system is accepted from [1980Bre]. Between MoO2 and MoO3 several oxides of the general formula MonO3n–m have been prepared. For m = 1, oxides with n = 4, 8, 9, 14 are known; for m = 2, oxides with n = 18 to 22 have been prepared; for m = 3, oxides with n = 12, 17 and 26 exist; for m = 4, Mo17O57 is known. These oxides are stable with respect to the (MoO2 + MoO3) mixture, but it is not clear whether all of them belong to the stable phase diagram. Their structure is derived from the ReO3 type structure. Mo4O11, the most stable of these oxides, disproportionates into MoO2 + liquid at 818 ± 7°C. The oxides Mo8O23 and Mo9O26 are probably stable. Solid Phases The compounds whose existence is widely recognized are presented in Table 2. There is no solid solutions between iron and molybdenum oxides. The wüstite has been reported to dissolve up to 3 at.% Mo at 1140°C [1972Abe]. Several oxides are known: MoFeO0.33 (τ1) phase of the “high speed steel” structure, Ti2Ni like according to [1967Hol] and MoFeOx (τ2) (0.4 < x < 1), known as the Hägg compound may be obtained by mixing one metal with the oxide of the other metal [1954Sch1, 1954Sch2]. τ2 is widely non stoichiometric, but with an oxygen content which never exceeds 33 at.% which corresponds to the ideal formula MoFeO. MoFe2O4 (τ3) is an inverse spinel which gives a solid solution with magnetite Fe3O4. The crystal parameter of the solid solution (MoxFe1–x)Fe2O4 presents a break in the slope for x = 0.18 [1993Mor]. The iron (II) molybdate (IV) Mo3Fe2O8 (τ4) has been reported by [1956McC, 1957McC] and was later recognized as the mineral Kamiokite. MoFeO4 (τ5) is easily obtained by solid state reaction between FeO and MoO3 oxides at 500°C [1925Tam]. The iron (III) molybdate (VI) Fe2(MoO4)3 (τ6) may be obtained by oxidation of Mo3Fe2O8 (τ4) with the intermediate formation of MoFeO4 (τ5). The compounds in which molybdenum is hexavalent (MoFeO5 (τ5), Fe2(MoO4)3 (τ6) and MoO3) are liquid at 1100°C whereas the compounds in which molybdenum is tetravalent (MoFe2O4 (τ3), Mo3Fe2O8 (τ4) and MoO2) have a melting point higher than 1100°C. Single crystals of the τ3, τ4, τ5 and τ6 compounds may be grown by chemical transport reaction [1993Mor], from the hot zone at 1050°C to the cold zone at 1000°C, using TeCl4 as transporting agent. Besides τ1 and τ2, another low valence compound, namely Mo4Fe2O7 (τ7) has been described by [2001Hai]. Discovered by serendipity on the occasion of synthesis of NaMo4O6 in steel crucible, it may be obtained by sintering Fe + Mo + MoO3 powders at 1200°C. Below 200 K, it presents a ferrimagnetism due to the different magnetic moments of iron in two different crystallographic sites. τ7 phase is probably metastable, it has not been introduced in the phase diagrams presented. Landolt-Börnstein New Series IV/11D4
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2
Fe–Mo–O
Quasibinary Systems The Fe2O3-MoO3 quasibinary system determined by [1959Jae] is shown in Fig. 1. The oxygen pressure at equilibrium is not determined because it depends on the composition of the mixture and a more informative diagram would show isobaric lines. The quasibinary systems MoFeO4-MoO3 and Fe2(MoO4)3-MoFeO4 are shown in Figs. 2 and 3, respectively. The melting points of solid phases given by [1959Jae] being 10°C lower than the accepted ones, the position of the eutectic shown in Figs. 1, 2 and 3 have been raised by 10°C. Isothermal Sections The phase equilibria in the solid state, that is below 800°C, are shown in Fig. 4. The equilibria in the triangle FeO-MoO2-O are taken from [1959Jae, 1972Sch]; the reduced part of the diagram (FeO-MoO2-Mo-Fe domain) is accepted from [1989Rag]. This diagram may also be used at 900°C, but it must be pointed out that, at 900°C, MoO3 becomes liquid with a vapor pressure of the order of 0.1 MPa. The phase equilibria, investigated at 1000°C by [1981Kle, 1993Mor] are represented in Fig. 5. The oxidized part of the diagram (MoO2-τ4-τ3-O region) is from [1993Mor], but it must be pointed out that, at 1000°C, Fe2(MoO4)3 (τ6) (which melts at 965°C) and MoO3 are liquid and the vapor pressure of MoO3 exceeds 0.1 MPa. The phase diagram proposed at 900 and 1000°C by [1981Kle, 1994Koy] differs by the nature of the phases in equilibrium with τ3 and τ4. According to [1983Kle], (Fe) and μ phases are in equilibrium with Mo3Fe2O8 (τ4) below 916°C and with MoFe2O4 (τ3) above 916°C. That temperature is the inversion temperature of the equilibrium: τ4 + (Fe) ⇌ τ3 + μ. Below 916°C, τ4 is in equilibrium with (Fe); above 916°C, τ3 is in equilibrium with μ. The same observation was reported by [1994Koy] with an invariant temperature of 957°C. The phase equilibria in the triangle MoO2-FeO-Fe3O4, investigated at 1140°C by [1972Abe] under CO2-H2 atmospheres, corresponding to 10–12.7 < pO2/bar < 10–7.8, is shown in Fig. 6. The spinel solid solution (Mo1–xFex)Fe2O4 is in equilibrium with Mo3Fe2O8 (τ4) in its τ3 rich part (0.36 < x < 1) and with a liquid phase of undetermined composition in its Fe3O4 rich part (0 < x < 0.36). The phase equilibria at 1600°C, experimentally determined by [1970Tho] is shown in Fig. 7, after slight modifications to take into account the accepted Fe-Mo diagram, particularly the presence of the σ phase which decomposes into (Mo) + L1 above 1611°C. Thermodynamics The thermodynamic quantities related to the formation from the oxides of MoFe2O4 (τ3), MoFeO4 (τ5) and to the phase transition of Fe2(MoO4)3 (τ6) are given in Table 3; the thermodynamic quantities related to single phases are given in Table 4. The Gibbs energy of formation of τ3, τ4, τ5 and τ6 together with that of μ, Mo0.42Fe0.58 phases are accepted from [2003Koy] because they are internally consistent, being obtained from the same equipment. Some older measurements are presented for comparison. The enthalpy of formation of τ5 from its elements, measured at –1053.5 kJ·mol–1 by [2003Koy] in the temperature range 839–1066°C must be compared to an earlier experimental value of –1063 kJ·mol–1 reported earlier by [1984Kag] at 25°C. On another hand, the Gibbs energy of formation of τ3 and τ4 proposed by [2003Koy] are in good agreement with the earlier measurements of [1981Kle] obtained also from electromotive force measurements in the same temperature range. The activities of Fe3O4 in the spinel solid solution have been evaluated at 900°C [1972Sch] and at 1200°C [1975Kat] from the oxygen pressures at equilibrium spinel/wüstite under various CO-CO2 or H2-CO2 gaseous atmospheres. The activities of Fe3O4 in the spinel present a small negative departure from ideal behavior at 900°C and an ideal behavior at 1200°C. The activity of MoFe2O4 were calculated by integration of the Gibbs-Duhem equation. The interaction coefficients of Mo and O in liquid iron at 1600°C, (< 20 mass% Mo) have been evaluated by [1965Tan] by equilibrating liquid alloys under H2-H2O atmospheres: eO(Mo) = {∂ log10 fO/∂ (mass% Mo)} = 0.0025 at 1600°C, with fO = {(mass% O in pure Fe)/(mass% O in alloy)}. An higher value eO(Mo) = 0.007 at 1600°C was measured by [1971Fis] from emf measurements, with stabilized zirconia used as solid electrolyte. The most probable value eO(Mo) = 0.0035 at 1600°C is proposed in more recent evaluations [1966Sch, 1974Sig, DOI: 10.1007/978-3-540-78644-3_27 # Springer 2008
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Fe–Mo–O
3
2003Per] in which the liquid alloy is equilibrated with oxygen. This positive value, not far from zero, means that Mo decreases slightly the oxygen solubility in Fe. Notes on Materials Properties and Applications The main experimental works are summarized in Table 5. When a Mo steel is heated in air at high temperature, the surface is coated by a layer composed of τ3, τ4 and τ5 which prevents further oxidization of the steel. The spinel τ3,MoFe2O4 exhibits a weak spontaneous magnetization (0.15 to 0.20 μB·mol–1) below the Neel temperature of 348 K (75°C) [1972Abe], with an unexplained change of sign below 160 K. The molybdate τ6 is widely used for its catalytic properties. For instance, its catalyses the selective oxidation of methanol into formaldehyde, the ammoxidation of propylene into acrylonitrile [1979Che] or the oxidation of propane into formaldehyde [2003Tay]. It has been claimed that the ability of τ6 to incorporate MoO3, inducing a possible non stoichiometry, is important for its catalytic properties. However, [1981Mas] found no evidence of a possible excess of MoO3 in the structure. The linear thermal expansion coefficient α = (1/ℓ)(∂ℓ/∂T)P of τ6 was measured by [2002Tya]. For the low temperature modification, monoclinic, α = 9.74 · 10–6 K–1; for the high temperature modification, after the monoclinic-orthorhombic transition, α = – 14.82 · 10–6 K–1 < 0, a feature attributed to the bending of the Fe-Mo-O bonds. Miscellaneous The partition coefficient of Mo between Mo saturated liquid slag and liquid metal in the Fe-Mo-Si-O system ranges from 0.34 at 0 % SiO2 and 1600°C to less than 0.01 at silica saturated [1970Tho], which means that SiO2 in the slag displaces Mo from the slag to the metal. The partition coefficient increases with temperature. The reduction of MoFe2O4 by H2, carried out by [2002Mor] between 850 and 1050°C, leads to MoFe2 in one step with an activation energy of 173.5 kJ·mol–1. The electrical conductivity and magnetic susceptibility of τ4, measured by [1982Str] below room temperature show that τ4 presents an antiferromagnetic ordering at 59.5 K. τ4 presents also a strong anisotropy. For H || c (magnetic field parallel to the c axis), μeff = 5.8 ± 0.3 μB per Fe atoms; for H⊥ c (magnetic field perpendicular to the c axis), μeff = 4.4 ± 0.2 μB per Fe atom. The ratio of the electrical resistivities σ⊥ c / σ || c varies from ∼ 40 at room temperature to ∼ 300 at 150 K. Below 140 K, resistivity cannot be measured because the resistances are of the order of 10 GΩ.
Table 1.
Investigations of the Fe-Mo-O system Phase Relations, Structures, Thermodynamics
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1925Tam]
XRD, calorimetry
< 600°C, FeO + MoO3
[1954Sch1, 1954Sch2]
XRD, solid-state reaction
MoFeO0.33 and MoFeOx structures (0.4 < x < 1)
[1956McC, 1957McC]
XRD, solid-state reaction
Fe2Mo3O8 prepared from 3 Fe2O3 +4 Mo + 5 MoO3 at 1150°C
[1959Jae]
XRD, equilibria measurements under H2-H2O atmospheres
800–1100°C, FeO-Fe3O3-MoO3-MoO2 phase equilibria
[1965Tan]
Oxygen solubility in (Fe,Mo) liquid alloy, chemical analysis
1550–1700°C, < 20 mass% Mo, H2-H2O atmospheres
[1968Sle]
XRD, pressure, phase transitions
FeMoO4, < 700°C, < 300 MPa (continued)
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Fe–Mo–O
Reference
Method/Experimental Technique
Temperature/Composition/Phase Range Studied
[1970Tho]
XRD, thermal analysis
1600–1800°C, the Fe-Mo-O phase diagram, O2 pressure undetermined
[1971Fis]
EMF, stabilized ZrO2 as solid electrolyte
1600°C, < 23 mass% Mo, < 0.1 mass% O
[1972Abe]
XRD, thermogravimetry, chemical analysis
1140°C, MoO2-FeO-Fe3O4 diagram, CO2-H2 atmospheres
[1972Sch]
XRD, thermogravimetry, chemical analysis
900°C, Fe-Mo-O diagram, CO-CO2 atmospheres
[1975Kat]
XRD, thermogravimetry, chemical analysis
1200°C, tie lines in the spinel-wüstite twophase domain
[1981Kle]
XRD, EMF, ThO2 stabilized with Y2O3 as solid electrolyte
900–1000°C, Fe-Mo-MoO2-Fe3O4 diagram
[1981Mas]
XRD, DSC
< 600°C, Fe2(MoO4)3, crystal structure, polymorphic transition
[1993Mor]
XRD, IR spectroscopy, thermal analysis
1000°C, Fe-Mo-O diagram, single crystal growth
[1994Koy]
XRD, EMF, ZrO2 stabilized with Y2O3 as solid electrolyte
900–1200°C, Fe-Mo-O diagram
[2003Koy]
EMF, ZrO2 stabilized with Y2O3 as solid electrolyte
767–1193°C, Gibbs energy of formation of Mo-Fe oxides
Table 2.
Crystallographic Data of Solid Phases
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
α, (αFe,δFe) (δFe) 1538–1394
cI2 Im 3m W
(αFe) (Ferrite) < 912
Lattice Parameters [pm]
Comments/References
a = 293.15
pure Fe at 1390°C [V-C2, Mas2] Dissolves 24.4 at.% Mo at 1449°C [1982Fer] pure Fe at 25°C [Mas2] Dissolves 13.4 at.% Mo at 1200°C [1982Fer]
a = 286.65
(Mo) < 2623
cI2 Im 3m W
a =314.70
[Mas2]. Dissolves 31.3 at.% Fe at 1611°C [1982Fer]
(εFe)
hP2 P63/mmc Mg
a = 246.8 c = 396.0
at 25°C, > 13 GPa [Mas2]
(continued)
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Fe–Mo–O
5
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
(γFe) (Austenite) 1394–912
cF4 Fm 3m Cu
a = 364.67
pure Fe at 915°C [Mas2] Dissolves 1.7 at.% Mo at 1140°C [1982Fer]
λ, MoFe2 < 927
hP12 P63/mmc MgZn2
a = 474.4 c = 772.5
[1982Fer]
R, Mo2Fe3 1488–1200
hR159 R 3 Co5Cr2Mo3
μ, Mo6Fe7 < 1370
hR39 R 3m Fe7W6
σ, MoFe 1611–1235
tP30 P42/mnm σCrFe
FeO (Wüstite) 1422–569
33.9 to 38.5 at.% Mo [1982Fer] a = 1091.0 ± 0.3 c = 1935.4 ± 0.5
[V-C2] 39.0 to 44.0 at.% Mo [1982Fer]
a = 474.8 ± 0.2 c = 257.0 ± 0.2 a = 478.0 ± 0.2 c = 258.4 ± 0.2
at 39.0 at.% Mo [1981Kle] at 44.0 at.% Mo [1981Kle]
a = 921.8 c = 481.3
42.9 to 56.7 at.% Mo [1982Fer]
cF8 Fm 3m NaCl
a = 431.0 a = 429.3
actually Fe1–xO with 0.05 < x < 0.12 x = 0.05 x = 0.12
αFe3O4 (r) < 580
oP56 Pbcm Fe3O4 I
a = 1186.8 b = 1185.1 c = 1675.2
[V-C2]
αFe2O3 (Hematite) < 1457
hR30 R 3c αAl2O3 (Corundum)
a = 503.42 ± 0.03 c = 1374.83 ± 0.04
[V-C2]
βFe2O3
cI80 Ia 3 βMn2O3 (Bixbyite)
metastable phase a = 939.3 ± 0.2 [V-C2]
γFe2O3 (Maghemite)
cF56 Fd 3m MgAl2O4
a = 834
metastable phase [1989Rag]
Mo3O
cF16 Fm 3m Anti-BiF3
a = 554.9
low temperature phase, probably metastable [1980Bre]
MoO2 < 2300
mP12 P21/c VO2
a = 560.7 ± 0.1 b = 486.0 ± 0.1 c = 562.4 ± 0.1 β = 120.94°
[V-C2]
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6
Fe–Mo–O
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
αMo4O11
mP60 P21/a
a = 2454 b = 543.9 c = 670.1 β = 94.28°
[1989Rag]
βMo4O11 < 818 ± 7
oP60 Pna21 Mo4O11
a = 2449 b = 675.2 c = 545.7
[1989Rag] Disproportionates into MoO2 + Liquid [1980Bre]
Mo17O47
oP128 Pba2 Mo17O47
a = 2161.5 b = 1963.2 c = 395.1
[1980Bre] probably metastable
Mo8O23 < 780
mP62 P2/c Mo8O23
a = 1340 b = 404 c = 1680 β = 106.1°
[1989Rag] Disproportionates into Mo4O11 + MoO3 [1980Bre]
Mo9O26
mP70 P2/c Mo9O26
a = 1445 b = 403.0 c = 1675 β = 96.0°
[1989Rag]
Mo18O52
aP140 P 1 Mo18O52
a = 814.5 b = 1189 c = 1966 α = 95.47° β = 90.39° γ = 109.97°
[1980Bre] probably metastable
MoO3 < 804
oP16 Pnma MoO3
a = 1385.5 b = 369.64 c = 396.28
[1989Rag]
* τ1, MoFeO0.33
cF112 Fd 3m Ti2Ni
a = 1094 ± 1
[1954Sch2, 1989Rag] probably metastable
* τ2, MoFeOx
hP3 P63/mmc Fe2N
a = 275 c = 432
0.4 < x < 1 [1989Rag] probably metastable Hägg type structure [1954Sch1]
spinel (sp) τ3, MoFe2O4
cF56 Fd 3m MgAl2O4
a = 849.9 ± 0.1
inverse spinel [1993Koe]
a = 840.4 ± 0.1
Mo0.18Fe2.82O4, break in the slope [1993Mor] at 1000°C at 25°C
βFe3O4 (h) (Magnetite) 1597–580
a = 854.5 a = 839.6
(continued)
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Fe–Mo–O
7
Phase/ Temperature Range [°C]
Pearson Symbol/ Space Group/ Prototype
Lattice Parameters [pm]
Comments/References
* τ4, Mo3Fe2O8 (Kamiokite)
hP26 P63mc Zn2Mo3O8
a = 577.32 ± 0.06 c = 1005.42 ± 0.11
[1957McC, 1982LeP]
a = 572.39 ± 0.02 c = 1005.48 ± 0.02
[1993Koe]
* τ5, βMoFeO4 1068–400
mC12 C2/m αMnMoO4
a = 980.5 b = 895.0 c = 766.0 β = 114.05°
[1968Sle, 1989Rag]
* τ5´, αMoFeO4 < 400
mC12 C2/m αCoMoO4
a = 1029.0 b = 939.4 c = 707.2 β = 106.31°
[1989Rag]
* τ5(I), γMoFeO4
aP6 Structure related to monoclinic NiWO4
a = 494.43 b = 570.06 c = 470.78 α = 92.32° β = 90.27° γ = 89.33°
[1972Sle] High pressure modification (300 MPa, 600°C) [1968Sle]
* τ6, βFe2(MoO4)3 965–513
oP68 Pnca -
a = 933.0 ± 0.3 b = 1286.8 ± 0.5 c = 924.2 ± 0.3
at 546°C paraelastic [1981Mas, 1993Koe]
* τ6´, αFe2(MoO4)3 < 513
mP136
a = 1573.7 ± 0.8
at 22°C
P21/a -
b = 923.1 ± 0.5 c = 1822.4 ± 0.9 β = 125.46 ± 0.02°
Pale green crystal, ferroelastic [1979Che, 1981Mas]
oI52 Imma Sc0.75Zn1.25Mo4O7
a = 601.8 ± 0.1 b = 578.2 ± 0.1 c = 1690.1 ± 0.5
[2001Hai] probably metastable
* τ7, Mo4Fe2O7
Table 3.
Thermodynamic Data of Reaction or Transformation
Reaction or Transformation
Temperature [°C]
Quantity, per mol of atoms [kJ, mol, K]
Comments
MoO2 + 2FeO ⇌ τ3
900
ΔrH° = – 3.4 ± 3.4
[1972Sch]
MoO3 + FeO ⇌ τ5
500
ΔrH° = – 56.5
[1925Tam]
ΔfH° = – 1063 (Experimental) ΔfH° = – 1084 (Calculated)
[1984Kag]
ΔtrH° = 1.866 ± 0.084
[1981Mas]
Mo + Fe + 2O2 ⇌ τ5 τ6´, αFe2(MoO4)3 ⇌ τ6, βFe2(MoO4)3
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8 Table 4.
Fe–Mo–O Thermodynamic Properties of Single Phases
Phase
Temperature Range [°C]
Property, per mole of molecules [kJ, mol, K]
Comments
τ3, MoFe2O4
900–1050
ΔfG° = – 1081.7 + 0.2667 T (± 2.6 kJ·mol–1)
[1981Kle]
τ3, MoFe2O4
970–1193
ΔfG° = – 1174 + 0.342 T (± 1 kJ·mol–1)
[2003Koy]
τ4, Mo3Fe2O8
900–1050
ΔfG° = – 2267.8 + 0.6165 T (± 4.5 kJ·mol–1)
[1981Kle]
τ4, Mo3Fe2O8
839–1066
ΔfG° = – 2347 + 0.6814 T (± 1 kJ·mol–1)
[2003Koy]
τ5, MoFeO4
25
H298 – H0 = 20.58 S°298 = 129.30
[1988Bag]
τ5, MoFeO4
839–1066
ΔfG° = – 1053.5 + 0.2983 T (± 0.4 kJ·mol–1)
[2003Koy]
τ6, Mo3Fe2O12
767–872
ΔfG° = – 2993 + 0.9105 T (± 2 kJ·mol–1)
[2003Koy]
μ, Mo0.42Fe0.58
889–950
ΔfG° = – 18.7 + 0.01117 T (± 0.1 kJ·mol–1)
[2003Koy]
Table 5.
Investigations of the Fe-Mo-O Materials Properties
Reference
Method/Experimental Technique
Type of Property
[1982Str]
Magnetic susceptibility, electric conductivity
τ4, Mo3Fe2O8, < 300 K, anisotropy, antiferromagnetic ordering
[1993Koe]
XRD (Rietveld refinement), Mössbauer
Bond valence, bond length, clusters in the Mo-Fe oxides
[2001Hai]
XRD, magnetic measurements (Faraday magnetometer)
τ7, Mo4Fe2O7, crystal structure and ferrimagnetic properties
[2002Mor]
XRD, thermogravimetry, reduction of MoFe2O4 by H2
800–1050°C, preparation of MoFe2, activation energy determination
[2002Tya]
XRD, thermodilatometry
25–800°C, phase transition and thermal expansion of Fe2(MoO4)3
[2003Ehr]
Neutron powder diffraction, magnetization
Fe2(MoO4)3
[2003Tay]
XRD, Raman spectroscopy, surface measurement (BET)
300-500°C, Fe2(MoO3)3, catalytic properties
[2006Ha]
XRD, TEM, field emission SEM
Fe-Mo nanoparticles preparation
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Fe–Mo–O
Fig. 1. Fe-Mo-O.
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The MoO3-Fe2O3 quasibinary system
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Fig. 2. Fe-Mo-O.
Fe–Mo–O
The MoO3-τ5,MoFeO4 quasibinary system
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Fe–Mo–O
Fig. 3. Fe-Mo-O.
Landolt-Börnstein New Series IV/11D4
11
The τ5,MoFeO4- τ6,Fe2(MoO4)3 quasibinary system
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Fig. 4. Fe-Mo-O.
Fe–Mo–O
The phase equilibria in the solid state (T < 800°C)
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Fe–Mo–O
Fig. 5. Fe-Mo-O.
Landolt-Börnstein New Series IV/11D4
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Isothermal section at 1000°C
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Fig. 6. Fe-Mo-O.
Fe–Mo–O
Partial isothermal section at 1140°C. Values give log10(pO2/bar)
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Fe–Mo–O
Fig. 7. Fe-Mo-O.
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Isothermal section at 1600°C
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16 References [1925Tam] [1945Dar]
[1946Dar]
[1954Sch1]
[1954Sch2] [1956McC] [1957McC] [1959Jae] [1965Tan]
[1966Sch]
[1967Hol] [1968Sle] [1970Tho]
[1971Fis]
[1972Abe] [1972Sch]
[1972Sle] [1974Sig] [1975Kat]
[1979Che] [1980Bre]
Fe–Mo–O
Tammann, G., “Chemical Reaction between two Powdered Crystals” (in German), Z. Anorg. Allg. Chem., 149, 21–98 (1925) (Phase Relations, Kinetics, Experimental, 41) Darken, L.S., Gurry, G.W., “The System Iron-Oxygen. I. The Wuestite Field and Related Equilibria”, J. Am. Chem. Soc., 67, 1398–1412 (1945) (Experimental, Phase Diagram, Thermodyn., *, 26) Darken, L.S., Gurry, G.W., “The System Iron-Oxygen. II. Equilibria and Thermodynamics of Liquid Oxides and Other Phases”, J. Am. Chem. Soc., 68, 798–816 (1946) (Experimental, Phase Diagram, Thermodyn., *, 24) Schoenberg, N., “On the Existence of Metallic Ternary Oxides Me’Me’’O with the Metal Atoms in Hexagonal Close-Packing”, Acta Chem. Scand., 8(4), 630–632 (1954) (Crys. Structure, Experimental, 5) Schoenberg, N., “On the Existence of Ternary Transition Metal Oxides”, Acta Chem. Scand., 8(6), 932–936 (1954) (Crys. Structure, Experimental, 13) McCarroll, W.H., Ward, R., Katz, L., “Ternary Oxides of Tetravalent Molybdenum”, J. Am. Chem. Soc., 78, 2909–2910 (1956) (Crys. Structure, Experimental, 3) McCarroll, W.H., Katz, L., Ward, R., “Some Ternary Oxides of Tetravalent Molybdenum”, J. Am. Chem. Soc., 79, 5410–5414 (1957) (Crys. Structure, Experimental, 13) Jäger, W., Rahmel, A., Becker, K., “The Ternary System Iron-Molybdenum-Oxygen” (in German), Arch. Eisenhuettenwes., 30, 435–439 (1959) (Crys. Structure, Experimental, Phase Diagram, 22) Tankins, E.S., Thomas, M.K., Erthal, J.F., Williams, F.S., “The Activity of Oxygen in Liquid Fe-Mo and Fe-W Alloys”, Trans. Am. Soc. Met., 58(3), 245–252 (1965) (Experimental, Phase Relations, Thermodyn., 10) Schenck, H., Steinmetz, E., “Activity, Standard Conditions and Coefficients of Activity” (in German), Stahleisen-Sonderberichte, Duesseldorf: Verlag Stahleisen, (7), 1–36 (1966) (Phase Relations, Thermodyn., Review, 161) Holleck, H., Thuemmler, F., “Investigations on the Formation of Metalloid-Stabilized Zr-Rich Transition” (in German), J. Nucl. Mater., 23, 88–94 (1967) (Crys. Structure, Review, 12) Sleight, A.W., Chamberland, B.L., “Transition Metal Molybdates of the Type AMoO4”, Inorg. Chem., 7(8), 1672–1675 (1968) (Crys. Structure, Experimental, 10) Thorne, J.K., Dahl, J.M., Van Vlack, L.H., “Partition Coefficient of Molybdenum in the Fe-Mo-Si-O System”, Metall. Trans., 1(8), 2125–2132 (1970) (Experimental, Phase Diagram, Phase Relations, 9) Fischer, W.A., Janke, D., “The Activity of Oxygen in Iron Melts Containing Molybdenum, Tungsten, Niobium or Tantalum” (in German), Arch. Eisenhuettenwes., 42(10), 695–698 (1971) (Thermodyn., Experimental, 15) Abe, M., “Phase Equilibrium Relations in the MoO2-FeO-Fe3O4 System at 1140°C”, Mater. Res. Bull., 7, 1443–1453 (1972) (Experimental, Phase Diagram, Phase Relations, 14) Schmahl, N.G., Dillenburg, H., “Phase Equilibria and Thermodynamics of the Ternary Systems Fe-Mo-O and Fe-W-O” (in German), Z. Physik. Chemie. Neue Folge, 77, 113–126 (1972) (Experimental, Phase Diagram, Thermodyn., 40) Sleight, A.W., “Accurate Cell Dimensions for ABO4 Molybdates and Tungstates”, Acta Crystallogr., B28, 2899–2902 (1972) (Crys. Structure, Experimental, 19) Sigworth, G.K., Elliott, J.F., “The Thermodynamics of Liquid Dilute Iron Alloys”, Met. Sci., 8, 298–310 (1974) (Review, Thermodyn., 249) Katsura, T., Wakihara, M., Hara, S.I., Sugihara, T., “Some Thermodynamic Properties in Spinel Solid Solutions with the Fe3O4 Component”, J. Solid State Chem., 13(1–2), 107–113 (1975) (Experimental, Phase Diagram, Thermodyn., 11) Chen, H., “The Crystal Structure and Twining Behaviour of Ferric Molybdate, Fe2(MoO4)3”, Mater. Res. Bull., 14, 1583–1590 (1979) (Crys. Structure, Experimental, 7) Brewer, L., Lamoreaux, R.H., “The Mo-O System (Molybdenum-Oxygen)”, Bull. Alloys Phase Diagrams, 1(2), 85–89 (1980) (Crys. Structure, Thermodyn., Assessment, #, 12)
DOI: 10.1007/978-3-540-78644-3_27 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mo–O [1981Kle]
[1981Mas] [1982Fer] [1982LeP] [1982Str] [1983Kle]
[1984Kag]
[1988Bag]
[1989Rag]
[1991Sun] [1993Koe]
[1993Mor]
[1994Koy]
[1995Kow]
[2001Hai]
[2002Mor]
[2002Tya]
[2003Ehr]
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Kleykamp, H., Schauer, V., “Phase Equilibria and Thermodynamics in the Fe-Mo and Fe-Mo-O Systems”, J. Less-Common Met., 81(2), 229–238 (1981) (Experimental, Phase Diagram, Thermodyn., 35) Massarotti, V., Flor, G., Marini, A., “Crystal Data for Ferric Molybdate: Fe2(MoO4)3”, J. Appl. Crystallogr., 14, 64–65 (1981) (Crys. Structure, Experimental, 9) Fernandez-Guillermet, A., “The Fe-Mo (Iron-Molybdenum) System”, Bull. Alloy Phase Diagrams, 3(3), 359–367 (1982) (Assessment, Crys. Structure, Thermodyn., #, 40) Le Page, Y., Strobel, P., “Structure of Iron Molybdenum (IV) Oxide Fe2Mo3O8”, Acta Crystallogr., 38B(4), 1265–1267 (1982) (Crys. Structure, Experimental, 7) Strobel, P., “Growth and Physical Properties of Single Crystals of Fe2(II)Mo3(IV)O8”, J. Solid State Chem., 42, 242–250 (1982) (Crys. Structure, Electr. Prop., Magn. Prop., Experimental, 26) Kleykamp, H., “Experimental Aspects of Solid Galvanic Cell Methods for Thermodynamic Studies on Alloys”, Ber. Bunsen-Ges. Phys. Chem., 87(9), 777–781 (1983) (Phase Diagram, Thermodyn., Theory, 32) Kaganyuk, D.S., Perepelitsa, A.P., “Estimating the Enthalpy of Formation for Compounds of MR(EO4)2 Type”, Inorg. Mater., 20(4), 566–571 (1984), translated from Izv. Akad. Nauk SSSR, Neorg. Mater., 20(4), 653–658 (1984) (Thermodyn., Calculation, 14) Bagdavadze, D.I., Tsagareishvili, D.Sh., Tskhadaya, R.A., Gvelesiani, G.G., “Method of Computation of Enthalpy Increment of Crystalline Inorganic Compounds in the 0-298.15 K Temperature Range” (in Russian), Izv. Akad. Nauk Gruz. SSR, Ser. Khim, 14(3), 199–206 (1988) (Calculation, Thermodyn., Review, 8) Raghavan, V., “The Fe-Mo-O System” in “Phase Diagrams of Ternary Iron Alloys”, Indian Inst. Metals, Calcutta, 5, 193–204 (1989) (Crys. Structure, Phase Diagram, Phase Relations, Review, 15) Sundman, B., “An Assessment of the Fe-O System“, J. Phase Equilib., 12(1), 127–140 (1991) (Phase Diagram, Phase Relations, Thermodyn., Assessment, 53) König, U., Morgenstern, T., Försterling, G., “A Study of Structural Crystallography on Ternary Metal Oxides in the System Fe-Mo-O”, Mater. Sci. Forum, 133–136, 687–692 (1993) (Crys. Structure, Electronic Structure, Experimental, 7) Morgenstern, T., Lienhardt, J.L., Reichelt, J.L., König, U., Opperman, H., “Investigations on the Ternary Systems Fe-Mo-O and Ni-Mo-O”, Mater. Sci. Forum, 133–136, 627–632 (1993) (Phase Diagram, Crys. Structure, Transport Phenomena, Experimental, 7) Koyama, K., Harada, T., “Phase Diagram of the Fe-Mo-O System at 1173 K similar to 1473 K”, J. Jap. Inst. Metals, 58(12), 1401–1407 (1994) (Phase Diagram, Thermodyn., Experimental, 14) Kowalski, M., Spencer, P.J., “Thermodynamic Revaluation of the Cr-O, Fe-O and Ni-O Systems: Remodelling the Liquid, BCC and FCC Phases”, Calphad, 19(3), 229–243 (1995) (Assessment, Phase Diagram, Thermodyn., Review, 47) Hainz, M., Boller, H., “GaMgMo4O7 and Fe2Mo4O7 - Two Low Valent Molybdenum Oxides with a Fully Ordered Sc0.75Zn1.25Mo4O7 Type Structure”, J. Alloys Compd., 317–318, 132–135 (2001) (Crys. Structure, Experimental, Magn. Prop., 9) Morales, R., Arvanitidis, I., Sichen, D., Seetharaman, S., “Reduction of the Fe2MoO4 by Hydrogen Gas”, Metall. Mater. Trans. B, 33b(4), 589–594 (2002) (Experimental, Kinetics, Morphology, 23) Tyagi, A.K., Achary, S.N., Mathews, M.D., “Phase Transition and Negative Thermal Expansion in A2(MoO4)3 System (A = Fe3+, Cr3+ and Al3+)”, J. Alloys Compd., 339, 207–210 (2002) (Crys. Structure, Phys. Prop., Experimental, 14) Ehrenberg, H., Bramnik, K.G., Müssig, E., Buhrmeister, T., Weitzel, H., Ritter, C., “The Ferrimagnetic Structure of Fe2Mo3O12: Dependence of the Fe-O-O-Fe Superexchange Coupling on Geometry”, J. Magn. Magn. Mater., 261, 353–359 (2003) (Crys. Structure, Magn. Prop., Experimental, 5)
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DOI: 10.1007/978-3-540-78644-3_27 # Springer 2008
18 [2003Koy]
[2003Per] [2003Tay]
[2006Ha]
[Mas2] [V-C2]
Fe–Mo–O Koyama, K., Morishita, M., Harada, T., Maekawa, N., “Determination of Standard Gibbs Energies of Formation of Fe2Mo3O12, Fe2Mo3O8, Fe2MoO4, and FeMoO4 of the Fe-Mo-O Ternary System and μ-Phase of the Fe-Mo Binary System by Electromotive Force Measurement Using a Y2O3-Stabilized ZrO2 Solid Electrolyte”, Metall. Mater. Trans B, 34B, 653–659 (2003) (Thermodyn., Experimental, 15) Perrot, P., Foct, J., “Gases other than Hydrogen in Iron and Steels” (in French), Techniques de l’IngJnieur M4275, 1–23 (2003) (Thermodyn., Kinetics, Phase Relations, Review, 127) Taylor, S.H., Hutchings, G.J., Palacios, M.-L., Lee, D.F., “The Partial Oxidation of Propane to Formaldehyde Using Uranium Mixed Oxide Catalysts”, Catal. Today, 81, 171–178 (2003) (Catalysis, Experimental, Interface Phenomena, 9) Ha, J.K., Cho, K.K., “Characterization of Fe-Mo Alloyed Nanoparticles Synthesized by Chemical Vapor Condensation Process”, Mat. Sci. Forum, 510–511, 466–469 (2006) (Crys. Structure, Experimental, Morphology, Nano, 7) Massalski, T.B. (Ed.), Binary Alloy Phase Diagrams, 2nd edition, ASM International, Metals Park, Ohio (1990) Villars, P. and Calvert, L.D., Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, 2nd edition, ASM, Metals Park, Ohio (1991)
DOI: 10.1007/978-3-540-78644-3_27 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mo–Si
1
Iron – Molybdenum – Silicon Elena Semenova
Introduction The constitution of the Fe-Mo-Si system has been of interest since 1955, owing to the presence of the MoSi2 compound in the Mo-Si binary system. This compound is known to have high melting point and excellent oxidation resistance, serving as a basis for high temperature structural applications. The Fe rich ternary Fe-Mo-Si alloys were considered by the researches in the following decades as prospective for the development of materials with soft magnetic characteristics, greater hardenability, better mechanical properties. In order to investigate the solubility of iron in MoSi2, [1956Fit] studied the phase relationships in the FeMo-Si ternary system in the region of up to about 40 at.% Mo. The results were presented as an isothermal section at 1400°C. From this section, it can be seen that up to 10 at.% Fe can be dissolved in αMoSi2. Two ternary phases were also presented in this section, one of which having close to equiatomic composition. A number of the binary phases shown in the section, such as β, η, α and ζα, should not be observed at 1400°C if the Fe-Si phase diagram as presented by [Mas2] is taken into account. [1961Vog] studied the constitution of the Fe-Mo-Si system in the Fe-Mo-MoSi2-FeSi region at the temperature of the liquid-solid transformation. Six vertical sections were constructed from which [1961Vog] derived the liquidus surface projection. A ternary phase of the MoFe2Si2 composition was found. The microhardness of some alloys was also measured. [1985Ray, 1988Ray], in their thorough assessment of the work of [1961Vog], removed some of the discrepancies in the liquidus projection that arose through conflicts with more recent data relating to the binary systems and proposed hypothetical reactions for the uninvestigated region of the system. Along with the work of [1956Fit] and [1961Vog], the assessment also reviewed the studies reported by [1960Gla, 1962Gla, 1963Bar, 1965Sko, 1966Sko, 1968Gla, 1969Kar]. All of these publications concern the phase relationships in the solid state of the ternary system. [1960Gla, 1962Gla, 1963Bar] studied alloys with compositions lying along the 33 at.% Mo section to investigate the possibility of the formation of a ternary Laves phase. The existence of this phase, with a crystal structure of the MgZn2 type, was shown at an equiatomic composition as single phase; it being found as a second phase in other alloys along the 33 at.% Mo section up to the composition MoFeSi. The phase was interpreted by [1963Bar] as a solid solution of Si in the binary MoFe2 phase. [1969Tes] considered the electronic and structural factors leading to the formation of a Laves phase of this type. [1965Sko] investigated whether a phase with a crystal structure similar to the R phase was present in the Fe-Mo-Si system and reported that the Mo5Fe3Si2 ternary phase, having the hexagonal crystal structure attributed to R type phases, was observed in as cast alloys. The intermediate ternary phases forming in the Fe-Mo-Si system and their relationship with other phases of the system were studied at 800°C by [1966Sko, 1968Gla]. The existence of the Mo5Fe3Si2 [1965Sko] and MoFe2Si2 [1961Vog] ternary phases was confirmed. The Mo3FeSi and (Mo0.17Fe0.33)5Si3 ternary phases were found for the first time. An isothermal section for the Fe-Mo-Si system at 800°C was constructed. [1998Yi] examined the possibility of modifying the crystal structure and microstructure of MoSi2 by alloying with transition elements, including iron, in an attempt to improve mechanical properties. Compositions of the Fe-Mo-Si system that are enriched in iron are also of practical interest [1969Kar, 1988Nak, 1988Kum, 2002You, 1998Mie]. [1969Kar] reported on the substantial hardening observed after ageing alloys in the Fe corner of the system at 550–750°C. [1988Nak] studied the effects of rapid quenching on the suppression of ordering of the α phase. Taking into account the importance of information regarding the influence of alloying element on the stability of the γ phase during heat treatment, [1988Kum] used a subregular solution model with a contribution to the Gibbs energy from magnetic ordering to calculate the effect of Mo and Si on the stability of the γ phase in the ternary Fe-Mo-Si system at various temperatures. [2002You] studied the effect of Si additions on sintering and mechanical properties, as well as the effect of composition and cooling rate on sinter hardening characteristics of Fe-1.5 mass% Mo powders. Landolt-Börnstein New Series IV/11D4
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2
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[1998Mie] performed a Calphad optimization of thermodynamic parameters for the iron rich part of the ternary system based on the information given in the review of [1988Ray]. The wettability of a MoSi2 substrate by an iron melt was investigated by [1966Yas]. The review of the phase equilibria in the Fe-Mo-Si by [1994Rag] is unfortunately incomplete and so does not add any new information. Data on the temperature and composition ranges used in the investigations described above, as well as the experimental techniques employed, are summarized in Table 1. Binary Systems The Fe-Si and Mo-Si phase diagrams are accepted from [1982Kub] and [Mas2] respectively. The temperatures of the two invariant reactions in the Mo-Si system, l ⇌ Mo5Si3 + βMoSi2 (αMoSi2) and βMoSi2 ⇌ αMoSi2, are the same at 1900°C. The βMoSi2 ⇌ αMoSi2 degenerate reaction can be of metatectic or peritectic type depending on whether it precedes or follows the eutectic reaction. Some uncertainties are observed in the representation of the Fe-Mo phase diagram given by [Mas2] in the region of the α phase. [Mas2], citing [1982Gui2] (which is essentially the same work as presented in [1982Gui1]), presents what looks like a eutectic reaction l ⇌ α + R at 1449°C in the Fe rich region of the phase diagram (although this feature is by no means clear) but without any indication as to the composition of the phases. It is worth noting that the text indicates that the Fe rich liquidus falls to a minimum with the addition of Mo, contradicting the phase diagram presented in the accompanying figure. In the original Calphad assessment work [1982Gui1, 1982Gui2], experimental evidence is cited of a peritectic reaction taking place at 1449°C and a minimum in the liquidus and solidus occurring at a temperature of 1448°C. It is also indicated that ‘The agreement is excellent’ between calculated and experimentally determined phase boundaries. A later modeling of the system by the same author (cited in 1988And]) gives an invariant temperature of ‘1724 K’; 2 degrees higher than in the earlier work, ensuring a peritectic formation of the (αFe) phase; although this is higher than the experimental measurement quoted in [1982Gui1, 1982Gui2]. It is clear that this part of the diagram is in much need of confirmation, but in the meantime, the version given in [1982Gui2] is accepted here. Solid Phases The crystallographic data on the solid phases of the Fe-Mo-Si system are listed in Table 2. There are reports of four ternary phases present in the system. The work of [1960Gla, 1961Vog, 1962Gla, 1963Bar, 1956Fit, 1966Sko, 1968Gla, 1969Tes] refer to the presence of a ternary Laves type phase at an alloy of equiatomic composition, but it is considered here to be an extended solid solution of Si in the binary MoFe2 Laves phase over a considerable range of Si content. The experimental data of [1963Bar] and [1956Fit] suggest that the λ1 phase of the MoFeSi composition still exists at 1200 and 1400°C, respectively, even though it decomposes at 927°C in the Fe-Mo binary system. The τ1 ternary phase forms at 1440°C and has the composition of 34.5Mo-42.5Fe-23 Si (mass%) (18.54Mo-39.24Fe-42.22 Si (at.%)) according to [1961Vog]. Its existence was confirmed at 800°C by [1966Sko, 1968Gla]. [1966Sko] noted that its range of homogeneity at 800°C was insignificant, and it was found by [1961Vog] at subsolidus temperatures. The ternary phase observed by [1966Sko] at 800°C having the Mo3FeSi composition did not have a noticeable homogeneity range. Two other ternary phases with compositions of Mo2Fe2Si and Mo2Fe3Si observed by [1961Vog] were interpreted by [1966Sko] as ternary extensions of intermediate phases in the Fe-Mo binary system. While the Mo2Fe2Si phase was attributed to the endpoint of the σ phase in the ternary system, the Mo2Fe3Si phase, which forms congruently [1961Vog], was assumed to be related to the phase known at that time as Mo2Fe3. This was later shown by [Mas2] to be the R phase, and Si additions enable its temperature stability range to be increased to 1610°C [1961Vog]. The ternary phase with the composition Mo5Fe3Si2 found at 800°C by [1965Sko, 1966Sko] has a crystal structure similar to that of the R phase of the binary Fe-Mo system. In a critical analysis of the available ternary phase equilibrium data, [1985Ray] suggested that the Mo5Fe3Si2 phase was the same phase as the binary R phase, stabilized by Si additions to lower temperatures and higher Mo concentrations than it is in the binary system. The Mo10Fe52Si38 phase first observed by [1966Sko] was assumed to be a solid solution, based on the η binary phase stabilized by Mo to lower temperatures. DOI: 10.1007/978-3-540-78644-3_28 # Springer 2008
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Fe–Mo–Si
3
All intermediate phases of the Fe-Si and Mo-Si system dissolve about 2-4 at.% of the third component at 800°C [1966Sko]. The exception is the MoSi2 phase for which [1998Yi] found little solubility for Fe. [1969Kar] reported that additions of Si decreased the Mo solubility in the α phase, but without indicating an absolute value. The phase strengthening of this alloy at ≤800°C occurs mainly by the precipitation of the λ1 phase with a small amount of the μ phase. This statement given by [1969Kar] seems to be unlikely, it being in contradiction with the data of [1966Sko, 1968Gla] which would suggest that the μ phase cannot be in equilibrium with the α phase at these temperatures. Quasibinary Systems According to [1961Vog], the FeSi-MoSi2 section of the Fe-Mo-Si ternary system is quasibinary of eutectic type. Taking into account that the eutectic point is at the same time a saddle point on the ternary liquidus, and that the components of the FeSi-MoSi2 section melt congruently, this section can be considered as a true quasibinary, at least at subsolidus temperatures. It is constructed here (Fig. 1) based on the melting points of the FeSi and MoSi2 phases taken from [Mas2], the coordinates of the eutectic point from [1961Vog] and information on their mutual solubility [1966Sko, 1998Yi]. Invariant Equilibria According to [1961Vog], there are 8 invariant four-phase reactions with the participation of liquid occurring in the Fe-Mo-MoSi2-FeSi part of the ternary system. Taking into consideration the constitution of the Fe-Si system, where two eutectic reactions, l ⇌ ε + β and l ⇌ β + α, are shown, [1985Ray, 1988Ray] added one more invariant equilibrium of the transition type (U7), in which the ε, β, α phases and liquid are involved. In agreement with the version of Mo-Si phase diagram given in [Mas2], where the βMoSi2 ⇌ αMoSi2 transformation is shown, it is necessary to assume the existence of an invariant four-phase reaction resulting from the low-temperature modification of the MoSi2 compound, L + βMoSi2 ⇌ αMoSi2 + Mo5Si3. According to [1961Vog, 1985Ray, 1988Ray], the R phase forms in the ternary system congruently at 1610°C and the composition of 49Mo-42Fe-9Si (mass%) (47.5Mo-32.26Fe-20.24Si (at.%)). Therefore the maximum temperature on the monovariant line between the primary crystallization fields of the R and σ phases (1605°C, at the approximate composition of 50.5Mo-42Fe-7.5Si (mass%)) should correspond to the eutectic reaction L ⇌ σ + R, on the contrary to the peritectic reaction L + σ ⇌ R reported by [1985Ray, 1988Ray]. All of these reactions are listed in Table 3 together with the compositions of the liquid phase involved. Five temperature maxima were determined experimentally in the Fe-Mo-MoSi2-FeSi part of the ternary system by [1961Vog]. These five reactions are also included in Table 3. A reaction scheme for the system is derived in Figs. 2a, b. It contains, along with the above mentioned invariant reactions in the ternary system, two more reactions assumed by [1985Ray, 1988Ray] for the composition region not investigated by [1961Vog]. As these are hypothetical they are drawn with dashed lines. Based on the liquidus surface given by [1961Vog] and the boundary binary phase diagrams, and keeping in mind the constitution of the isothermal section at 800°C, [1988Ray] deduced the sequence of reactions that could proceed in the solid state. This is reasonable speculation that is consistent with the available experimental data relating to the liquidus surface and the alloy constitution at 800°C, proposing the probable scheme for the transition from a high to a low temperature state in the system. However it cannot be unconditionally accepted because of the lacking data on the high-temperature constitution of the system. Liquidus, Solidus and Solvus Surfaces The projection of the liquidus surface for the Fe-Mo-MoSi2-FeSi part of the ternary system was derived by [1961Vog] using experimental data for alloys located along certain vertical sections. According to [1961Vog], the τ1 phase forms by the peritectic reaction l + Mo5Si3 ⇌ τ1 at 1440°C. Nothing is known about the formation of the τ2 phase [1966Sko, 1968Gla]. If it forms from the solid, the liquidus projection of the system as shown by [1961Vog] for the Fe-Mo-MoSi2-FeSi region will remain unchanged. The projection was modified by [1985Ray] to conform with more recent versions of the binary diagrams, which coincide with those accepted here. An updated version of the liquidus surface of the system for the whole composition range is shown in Fig. 3. It includes the hypothetical invariant lines (dashed lines) separating Landolt-Börnstein New Series IV/11D4
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the fields of primary crystallization of the ε, βMoSi2, ζα and (Si) phases which were added by [1985Ray, 1988Ray]. Dashed lines separating tentatively the primary crystallization fields of βMoSi2 and αMoSi2 phases have also been added. The solidus and liquidus temperatures, (∼1045 and 1180°C, respectively, obtained by [2002You] for a sintered alloy of composition Fe-1.5Mo-3Si (mass%) can be considered as underestimated, as do not agree with those measured by [1961Vog] for alloys of similar composition. Isothermal Sections The isothermal section at 800°C constructed by [1966Sko] and reproduced in [1968Gla] is shown in Fig. 4 with small corrections related to the lack of solubility of iron in αMoSi2 in accordance with the work of [1998Yi]. The solubility range of the τ2 phase is shown as taken from the drawing in [1966Sko], however most likely it does not have a noticeable range of homogeneity as follows from the experiment data of [1966Sko]. Figures 5a to 5d present a set of the isothermal sections in the Fe rich corner at 950, 1050, 1150 and 1250°C calculated by [1988Kum]. Although the locations of the α and γ phase boundaries on the Fe-Mo and Fe-Si binary sides somewhat differ from those in [Mas2] the sections give a notion as to the variation of the γ/α+γ and α/α+γ phase boundaries with changing temperature. A probable isothermal section of the Fe-Mo-Si phase diagram at room temperature was estimated and presented by [1988Ray]. Temperature – Composition Sections Six vertical sections were constructed by [1961Vog] based on experimental data. Two of them, FeSi-Mo5Si3 and Fe75Si25-τ1, are shown in Figs. 6 and 7 with some modifications relating to the stoichiometry of the most thermostable silicide (Mo5Si3 instead of Mo3Si) and the homogeneity range of the ε phase. Other sections constructed by [1961Vog] are not shown here since they pass through phases whose composition changes substantially with temperature and therefore may lie out of the plane of the section at temperatures other than those studied (e.g. μ or R phases). Thermodynamics [1998Mie] performed a Calphad optimization of thermodynamic parameters describing the liquid, α and γ phases of the iron rich part of the Fe-Mo-Si system in order to calculate phase equilibria in multicomponent steels. Data used in the assessment were taken from the review given by [1988Ray]. The temperatures of two invariant equilibria involving the liquid and α phases; L + R ⇌ α + τ and L ⇌ τ1 + α + ε, were calculated by [1998Mie] to be 1188 and 1135°C, which are reasonably close to the experimental values of 1186 and 1170°C, respectively, given by [1961Vog]. Notes on Materials Properties and Applications The microhardness of the τ1 phase was measured to be HV 1080 (10.59 GPa) and that of the ternary R phase in a Mo2Fe3Si alloy was about HV 1180 (11.57 GPa) [1961Vog]. The hardness of a Fe-Mo-2.5Si (mass%) alloy increased to HV 190 (1.863 GPa) [1969Kar] on increasing the Mo concentration from 3 to 10 mass%; the hardness of aged specimens increasing noticeably. High silicon alloys containing around 6 mass% Si are well known as excellent soft magnetic materials [1988Nak]. Using iron single crystals containing 5.5 mass% Si and 0.5-1.5 mass% Mo, [1988Nak] examined the effect of quenching on the magnetic anisotropy and saturation magnetization. The latter was found to decrease with increasing Mo concentration while the magnetic anisotropy decreased with increasing quench temperature. The behavior of the anisotropy was attributed to the degree of ordering in the alloys. An addition of 3 mass% Si to Fe-1.5Mo powder increases ultimate tensile strength up to 445 MPa, hardness up to HV 10 (0.098 GPa) as well as an improvement in the three point bend strength. These improvements in the mechanical properties are due to solid solution strengthening, enhanced α phase sintering and activated liquid phase sintering. All three properties increase with temperature [2002You]. Elongation was shown to decrease on increasing the Si content from 2 to 5 mass%. This was mainly related by [2002You] to an embrittling effect caused by Si. Complete wetting of a MoSi2 substrate by molten iron was observed by [1966Yas]. DOI: 10.1007/978-3-540-78644-3_28 # Springer 2008
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Landolt-Börnstein New Series IV/11D4
Fe–Mo–Si
5
Miscellaneous The effective radius of silicon in the ternary Laves phase alloys was shown by [1963Bar] to be about 118 pm, which is noticeably smaller than its atomic radius of 134 pm. Comparing the α phase lattice parameter of two aged alloys, Fe-10Mo (mass%) and Fe-10Mo-1Si (mass%) [1969Kar] reported that Si addition resulted in a decrease in the lattice parameter. Additions of Mo suppress the ordering of the α solid solution in the ternary system [1966Sko]. On increasing the Si content (2–5 mass%) in Fe-1.5Mo (mass%) alloy powder and increasing the temperature of sintering (1140–1250°C), shrinkage increases, and sintering under vacuum results in higher densities than sintering under pure hydrogen [2002You]. The Fe-1.5Mo-3Si-1.2C (mass%) composition exhibits high sintered density and good mechanical properties. A single stage compaction and sintering process have been developed.
Table 1.
Investigations of the Fe-Mo-S