The main activity directions of the “Metallurgy of Steel and Ferroalloys” department, Moscow Steel and Alloys Institute 1. Educational activities Higher qualification training through post-graduate and doctoral studies. All forms of raising the level of qualification and professional retraining (all programs are forming according to the concrete customer): • Raising the level of qualification up to 72 hours; • Raising the level of qualification more than 72 hours; • Professional retraining more than 500 hours. Master's degree education of "Metallurgy" direction according to the program: «Technological management in ferrous metallurgy” (the program content is corrected depending on requirement of the company-customer). Engineers education of the following specialties: • «Metallurgy of ferrous metals». Specializations: - Steel metallurgy; - Electrometallurgy of steel and ferroalloys; - Resource saving and ecology in metallurgy; - Informatics and business in metallurgy. • «Automation of technological processes and productions». Specializations: - Automation of technological processes; - Computer simulation of technological processes. • «Metallurgy of man-caused and secondary resources». Bachelors education of "Metallurgy" direction according to the following profiles: - Metallurgy of ferrous metals; - Metallurgy of man-caused and secondary resources; - The world market of raw materials and metals.
2. Scientific and professional activities The main directions of research studies: • The theory of steelmaking processes, crystallization and ferroalloys production. • Perfection of steelmaking units working, increase of their productivity and improvement of technical and economic indices. • Steelmaking technology. • Secondary Steelmaking. • Steel casting. • Mathematical simulation of metallurgical processes. • Ecology and resource saving of steelmaking processes. • Theory and practice of complex alloyed and stainless steel production. • High Nitrogen steels. • Research and perfection of continuous steelmaking processes and secondary steelmaking. • Automation and control of steelmaking processes. • Reconstruction of metallurgical plants and designing of mini-mills. • Appraisal of the investment attractiveness of the projected plants. • Quality control system of metal products. Pilot production of steel, alloys and ferroalloys. Technological audit and industrial consulting.
Proceedings of 10-th International Conference on High Nitrogen Steels HNS 2009 6 - 8 July 2009 Moscow, Russia
Edited by A.G. Svyazhin, V.G. Prokoshkina, K.L. Kossyrev
Moscow
• MISIS • 2009
Proceedings of 10-th International Conference on High Nitrogen Steels, HNS 2009
Committees Honorary Chairman
Prof. D.V. LIVANОV, Rector of State Technological University “Moscow Institute of Steel and Alloys”
Conference Chairman
Prof. A.G. SVYAZHIN, State Technological University “Moscow Institute of Steel and Alloys”
International Scientific Committee Nuri AKDUT Chavdar ANDREEV Oleg A. BANNYKH Hans BERNS Guocai CHAI Bruno C. De COOMAN Han DONG Jacques FOCT Karin FRISK Valentin GAVRILJUK Hannu HANNINEN Staffan HERTZMAN Laizhu JIANG Ludmila M. KAPUTKINA Leif KARLSSON Yasuyuki KATADA Yasushi KIKUCHI Konstantin L. KOSSYREV Mats LILJAS Erik MITTEMEIJER U. Kamachi MUDALI Jan-Olof NILSSON Baldev RAJ Tsolo RASHEV James C. RAWERS Jerzy SIWKA Evgeny Kh. SHAKHPAZOV
Belgium Institute of Metal Science, Bulgaria Baikov’s Institute of Metallurgy and Materials Science, Russia Ruhr-Universitat Bochum, Germany AB Sandvik Steel, Sweden POSTECH, GIFT, South Korea Central Iron and Steel Research Institute, National Engineering Research Center of Advanced Steel Technology, China Universite de Lille, France Institute for Metal Research, Sweden Institute of Metal Physics, Ukraine Helsinki University of Technology, Finland Outokumpu Stainless Research Foundation, Sweden Shanghai Baosteel Group Corporation, China Technological University Moscow Institute of Steel and Alloys, Russia ESAB, Sweden National Institute for Materials Science, Japan Osaka University, Japan Technological University Moscow Institute of Steel and Alloys, Russia Outokumpu Stainless AB, Sweden Max Planck Institute for Metals Research, Germany Indira Gandhi Centre for Atomic Research, India AB Sandvik Steel, Sweden Indira Gandhi Centre for Atomic Research, India Institute of Metal Science, Bulgaria U.S. Department of Energy NETL-Albany Center, USA Technical University of Czestochowa, Poland Bardin’s Institute TSNIICHERMET, Russia
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Proceedings of 10-th International Conference on High Nitrogen Steels, HNS 2009
Marcel A.J. SOMERS Markus O. SPEIDEL Jie SU Anatoly G. SVYAZHIN Setsuo TAKAKI Andre Paulo TSCHIPTSCHIN Sybrand van der ZWAAG
Technical University of Denmark, Denmark Swiss Academy of Materials Science, Switzerland Central Iron and Steel Research Institute, National Engineering Research Center of Advanced Steel Technology, China Technological University Moscow Institute of Steel and Alloys, Russia Kyushu University, Japan Escola Politecnica da Universidade de Sao Paulo, Brazil Delft University of Technology, Netherland
Organizing Committee Co-Chairman
Prof. K.L. KOSSYREV, State Technological University “Moscow Institute of Steel and Alloys” Mr. V.V. KOVSHEVNY, Research Company RUSMET
Members
Elena ARGUNOVA, Research Company RUSMET Lev A. CHESALOV, Research Company RUSMET Valentin G. GAVRILJUK, Institute of Metal Physics, Ukraine Ludmila M. KAPUTKINA, State Technological University “Moscow Institute of Steel and Alloys” Vladimir E. KINDOP, State Technological University “Moscow Institute of Steel and Alloys” Olga V. MOROZOVA, State Technological University “Moscow Institute of Steel and Alloys” Vera G. PROKOSHKINA, State Technological University “Moscow Institute of Steel and Alloys” Pavel Yu. SHENDRIKOV, State Technological University “Moscow Institute of Steel And Alloys” Markus O. SPEIDEL, Swiss Academy of Materials Science Dmitry V. KREMIANSKY, State Technological University “Moscow Institute of Steel And Alloys”
Proceedings of 10-th International Conference on High Nitrogen Steels, HNS 2009
Keynote speakers Oleg A. BANNYKH, Russia Russia Progress in the Research and Application of Nitrogen-Alloyed Steels.
Hans BERNS, Germany High Interstitial Stainless Austenitic Steels, Part I: Constitution, Heat Treatment, Properties, Applications.
Han DONG, China The Recent Progress of Product Technologies of High Nitrogen Stainless Steels in China
Jacques FOCT, France Roots and Wings of High Nitrogen Steel.
Valentin G. GAVRILIUK, Ukraine High Interstitial Stainless Austenitic Steels, Part II: Electronic and crystal structure, effect of loading
Guocai CHAI, Sweden Progress in high alloyed duplex stainless steels.
Kamachi MUDALI, India Passive Films and Localized Corrosion – Role of Nitrogen
Tsolo RASHEV, Bulgaria Problems of High Nitrogen Steel Development.
Marcel A.J. SOMERS, Denmark Low temperature surface nitriding of stainless steel.
Markus O. SPEIDEL, Switzerland Commercial Low-Nickel, High-Nitrogen Austenitic Stainless Steels.
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Preface In the given edition, the texts of reports submitted to the 10-th HNS Conference being held in Moscow, July 6 to 8, 2009, are collected. HNS 2009 in Moscow is in a series of international conferences, devoted to research, production and applications of High Nitrogen Steels (HNS), wich have been successfully held in France, Germany, Ukraine, Japan, Finland and Sweden, India, Switzerland, Belgium and China since 1988. High Nitrogen Steels is one of the most prospective concepts and trends in the development of advanced steel materials. Nitrogen is a unique element as it imparts to steels new properties, which can not be obtained by the use of other alloying elements. Nitrogen containing carbon and stainless steels having the unique properties are widely used in aviation, nuclear projects, defense and arms, oil & gas industries. More than 20 years have passed after the first HNS conference (1988, Lille, France). During these years, a new direction in metallurgical science was generated, the international community of researchers and engineers interested in the HNS problems was generated. During the expired years, a significant progress in the knowledge of alloying processes under normal and high pressure for nitrogen-containing steels and in studying the nature of the HNS structure and properties formation has been reached, new application fields for HNS have been found out. Industrial manufacturing of the HNS products for power industry, transport, chemical and food-processing industries and medicine successfully expands. It should be noted, however, that the use of HNS is still not enough in comparison with the possibilities presented by the knowledge of the HNS properties accumulated up to the present time. The possibilities of nitrogen as an alloying element are not studied enough. At the conference, the tendencies of HNS development have been considered, new results of research works in the field of nitrogen alloying, thermodynamics and kinetics of processes, formation of structure and properties in various classes of nitrogen-containing steels as well as their corrosion resistance and surface treatment have been presented and discussed. The HNS 2009 conference has been held in Moscow in accordance with the decision of the International Scientific Committee for High Nitrogen Steels in 2006 (China). The conference has been organized by State Technological University “Moscow Institute of Steel and Alloys” and Research Company “RUSMET”. The HNS 2009 Conference has been sponsored by Department Metallurgy of Steel and FerroAlloys State Technological University “Moscow Institute of Steel and Alloys”, informative support has been rendered by Journals Stal, Izvestia VUZ, Metallurg, Electrometallurgy, and this is gratefully acknowledged. The book can be useful to researchers, engineers and students interested in the HNA problems.
Anatoly G. Svyazhin Chairman of HNS 2009 Conference
CONTENTS Plenary Session Roots and Wings of High Nitrogen Steel Jacques FOCT (University de Lille, France) ………………………………………….........15 The Recent Progress of Product Technologies of High Nitrogen Stainless Steels in China Han Dong, Yuping Lang, Fan Rong and Jie Su (Central Iron and Steel research Group, China) ………………………………………………………..................................................21 Progress in the Research and Application of Nitrogen-Alloyed Steels Oleg BANNYKH (Baikov’s Institute of Metallurgy and Materials Science, Russia) ……....24
Thermodynamics & Kinetics Order-disorder transitions in high-nitrogen steels: from ab-initio to statistical thermodynamics A.J. Böttger, D.E. Nanu, A. Marashdeh ( Delft University of Technology, The Netherlands) ………………………………………………………………………..…..…....31 The thermodynamics investigations of nitrogen in the liquid Fe-N-V alloy A. Hutny, J. Siwka (Chestochova University of Technology, Poland) ……..……....……..35 Alloying of Steels with Nitrogen from a Gas Phase during VOD. S.A. Ivlev1, P.R. Scheller2, A.G. Svyazhin1 (1 State Technological University “Moscow Institute of Steel and Alloys”, Russia; 2 TU Bergakademie Freiberg, Germany)……………41 For the question on interaction of active nitrogen with iron in the plasma heating process. L.M. Simonyan (State Technological University “Moscow Institute of Steel and Alloys”, Russia)………………………………………………………………………………………..47 Classic description of the Nitrogen and Hydrogen solubility in solid iron. Yu. Venets (JSC "Centravis prodaction Ukraine", Ukraine)………………….……………..51 Dependence of the nitrogen solubility in austenite and ferrite on alloying at the elemental level. Yu. Venets (JSC "Centravis production Ukraine", Ukraine)…………………….………….57
Structure and Properties Evolution, Phase transformation Progress in high alloyed duplex stainless steels. G. Chai, U. Kivisakk and S.Ronneteg. (Sandvik Materials Technology, Sweden) ….……67 Effect of Nitrogen on structure properties of thermomechanically strengthened steels. A.G. Svyazhin, L.M. Kaputkina, V.G. Prokoshkina (State Technological University “Moscow Institute of Steel and Alloys”, Russia)………….…………………………………77 Creep Behaviour and Microstructural Evolution in AISI 316LN Steels with Niobium Additions. V. Vodarek, Miroslav Liška, Jaromír Sobotka (Materials and metallurgical Research Ltd., Czech Republic)……………………………….……………………………………………..83 High nitrogen PM tool steels containing niobium. St. Huth, W. Theisen (Ruhr-University Bochum, Germany)…………………….………...90 Effect of Aging and Cold Plastic Deformation on Structure and Stress Corrosion Cracking in High Nitrogen Chromium Steels.
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V.V. Berezovskaya 1, O.A. Bannykh 2, M.V. Kostina 2, M.S. Khadyev 1, A.I. Shestakov1 (1 Ural State Technical University-UPI,Yekaterinburg, Russia; 2 Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences, Moscow, Russia) …..…96 Embrittlement of Stainless Cr-Ni Steels Alloyed with Nitrogen. Balitskii, L. Ivaskevich, V. Mochulskyi (Karpenko Physico-Mechanical Institute, Ukraine) …………………………………………………………………………102 Influence of alloying and temperature strain upon formation of structure, necessary mechanical and corrosion properties in austenitic nitrogen containing steels. V. Rybin, G. Kalinin, S. Mushnikova (FSUE CRISM “Prometey”, Russia)………..….106 Effect of Aging on Mechanical Properties of High Nitrogen Austenitic Stainless Steel. Zu-rui ZHANG, Zhou-hua JIANG, Hua-bing LI, Bao-yu XU (Northeastern University, School of Materials and Metallurgy, Liaoning, Shenyang, China) ……………………….112 Austenitic High Nitrogen Steels Commercial Low-Nickel, High-Nitrogen Austenitic Stainless Steels. Markus O. Speidel and Hannes J. Speidel (Swiss Academy of Materials Science, Switzerland)………………………………………………………………………………..121 High Interstitial Stainless Austenitic Steels, Part I: Constitution, Heat Treatment, Properties, Applications. Hans Berns1, Sascha Riedner1, Valentin Gavriljuk2 (1 Ruhr-Universitat Bochum, Germany; 2 Institute of Metal Physics, Ukraine) ……………………………………………………...129 High Interstitial Stainless Austenitic Steels, Part II: Electronic and crystal structure, effect of loading. Valentin Gavriljuk1, Bela Shanina1, Andriy Tyshchenko1, Hans Berns2, Sascha Riedner2 (1. Institute of Metal Physics, Ukraine; 2.Ruhr-Universitat Bochum, Germany)………......140 Analysis of cavitation-erosion deterioration mechanisms at the micrometer scale in austenitic high nitrogen steels. D. Mesa1, C. Garzón2, A. Tschiptschin3 (1. Metallurgical and Materials Engineering Department, University of Sao Paulo, USP, Brazil; 2. Mechanical Technology Department, Pereira University of Technology, UTP, Colombia; 3. Physics Department, National University of Colombia, UNaL)…………………………………………………………….150 Stainless austenitic steel with 0.85 mass % C + N for castings. S. Riedner, N. Nabiran, H. Berns (Ruhr-Universität-Bochum, Germany)……………….156 Friction Stir Welding of High Nitrogen-containing Austenitic Stainless Steel. Y. Miyano, H. Fujii, S. Yufeng, K. Ieko, Y.Katada, O. Kamiya (Akita University, Japan)………………………………………..………………………………………………162 New grades of austenitic stainless steel high nitrogen content. S. N. Ghali, M. M. Eissa, T. M. Mattar (Central Metallurgical R & D Institute, Egypt)...168 Research of elastic aftereffect, phase transformation and stress relaxation process in deformed stable and metastable height-nitrogen stainless sheet steels. A. Skugorev, L. Kaputkina. (State Technological University “Moscow Institute of Steel and Alloys”, Russia) …………...………………………………………………………………..175 Nitrogen contents and strain ageing behavior in duplex stainless steels. G. Chai, M. Andersson. (Sandvik Materials Technology, Sweden)…………………..…..181
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The effect of interstitial impurity content and extension axis orientation on the parameters of plastic deformation localization in austenitic steel monocrystals. S. Barannikova, L. Zuev (Institute of Strength Physics & Materials Science, Russia) …...189 Research on Hot Working of High Nitrogen Austenitic Stainless Steel. Lang Yuping, Zhou Yong, Rong Fan, Chen Haitao, Weng Yuqing, Su Jie, Dong Han.(Central Iron and Steel Research Group, China)……………………………………...195 Temperature dependence of tensile behavior of nitrogen alloyed austenitic stainless steels. Wei WANG 1, 2), Wei YAN 1), Ke YANG 1), Yiyin SHAN 1), Zhouhua JIANG 3) 1) Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, PR China 2) Graduate University of Chinese Academy of Sciences, Beijing 100080, PR China 3) School of Materials and Metallurgy, Northeastern University, Shenyang 110014, PR China………………………/……………………………………………………………….202 Study of a High Nitrogen Nickel-free Austenitic Stainless Steel for Medical Application. Yibin Ren, Peng Wan, Feng Liu, Ke Yang and Bingchun Zhang (Institute of Metal Research, Chinese Academy of Sciences, China) ……………………………………….…208 The Effect of Nitrogen on Twinning in Single Crystals of Austenitic Stainless Steels. I.V. Kireeva, Yu.I. Chumlyakov, A.V. Tverskov, N.V. Luzginova (OSP “Siberian Physical Technical Institute at Tomsk State University”, Tomsk, Russia)…………………213
Steelmaking and Applications of High Nitrogen Steels Problems of High Nitrogen Steel Development. Ts. Rashev, Ch. Andreev, L. Jekova (Institute of Metal Science, Bulgarian Academy of Sciences, Bulgaria) …………………………………………………………………………221 Nitrogen-Containing Stainless Steels: Prospect of production and application E. Shakhpazov, A. Shlyamnev, O.Novichkova. (Bardin’s Institute TSNIICHERMET, Russia) ……………………………………………………………………………………..233 Peculiarities of High-Chromium Nitrogen Steels Melting and Casting. E.Yu. Kolpishon, E.V. Shytov (St. Petersburg State Polytechnic University, Russia) ......239 Manufacturing of HNS alloys at Energietechnik Essen GmbH – a process overview. R. Ritzenhoff (Volkher Diehl, Energietechnik Essen GmbH, Germany)………………….243 Effect of alloying on the composition-stable nitrogen content and phase composition of the Fe-CrMn-Ni-Mo-V-Nb corrosion-resistant alloys after solidification. Rigina L.G.1, Kostina M.V.2, Bannykhh O.A.2, Blinov V.M.2 (1 OAO NPO "Central Scientific Research Institute of Machine Building Technology", Russia; 2 Baikov Institute of Metallurgy and Materials Science, RAS, Russia)…………………………………………..249 Study of the solidification kinetics and casting-technological properties of new high-nitrogen nonmagnetic Cr-Ni-Mn-Mo-N steel for casting production. V.V. Nazaratin1, M.V. Kostina2, V.D. Gorbatch2, L.G. Rigina2, Е.V. Stetsukovskij3, S.O. Muradjan2 (1 Joint Stock Company Scientific and Production Coporation “Central Research Institute for Machine-building Technology” JSCSPC “TSNIITMASH”, Russia; 2 Baikov Institute of Metallurgy and Materials Science, RAS, Russia; 3 OAO «Centre of Shipbuilding and Ship repair» - Design Office «Armas», Russia)……………………………………….256 Arc Slag remelting of high nitrogen containing Ti-alloys. L. Medovar, V.Saenko (E.O. Paton Electric Welding Institute, Ukraine)………………..263
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Corrosion and Surface Treatment Passive Films and Localised Corrosion – Role of Nitrogen. U.Kamachi Mudali, S.Ningshen, and Baldev Raj (Indira Gandhi Centre for Atomic Research, India) ……………………………………………………………………………271 Study of Corrosion Resistance on 316LN Austenitic Stainless Steel. WANG Mingjia, LIU Xiaocui, WU Siyu, WANG Yan (Chinese Academy of Sciences, China)……………………………………………………………………………………...281 On the Possibilities of Main Gas Pipe Line Construction with Stainless Steels Containing Nitrogen. Ts.Rashev, Ch. Andreev, L.Ts. Jekova (Institute of Metal Science, Bulgarian Academy of Sciences, Bulgaria), L. Nenova – "Balevski" Institute of Metal Science …………………287 Roles of Nitrogen, Cold Work and Sensitization Treatment on Electrochemical Corrosion Behavior of Nickel Free High Nitrogen Stainless Steels. Xinqiang Wu1, Yao Fu1, Enhou Han1, Wei Ke1, Ke Yang1, Zhouhua Jiang2 (1. State Key Laboratory for Corrosion and Protection, Institute of Metal Research, CAS, Shenyang 110016, China; 2. School of Materials and Metallurgy, Northeastern University, Shenyang 110004, China) ……………………………………………………………………………………...294 Effect of the Structure and Phase Composition on the Pitting Corrosion Resistance of the Cr-N Steels with Overequilibrium Nitrogen Content, Russia S.Yu. Mushnikova1, M.V. Kostina2, Ch.A. Andreev3, L.Ts. Zhekova3 (1 FSUE CRISM “Prometey”, Russia; 2 Baikov Institute of Metallurgy and Materials Science, RAS, Russia; 3 Institute of Metal Science, Bulgarian Academy of Sciences, Bulgaria)……….……………….300 Computed Assisted Development of Corrosion Resistant TWIP Steels. L. Mujica1,2, S. Weber 2,3, W. Theisen 3 (1Max-Planck Institut fuer Eisenforschung, Duesseldorf, Germany; 2Ruhr-Universitaet Bochum, Germany; 3Helmholtz-Zentrum Berlin, Germany) …..………………………………………………………………………………306 An Influence of Heat and Deformation Treatment on the Structure and Properties of a Brazed Joint Including Maraging Nitrogen-Containing Steel and Non-Alloyed One. D. E. Kaputkin, M. V. Krasnoshekov, A.V.Gaponov (State Technological University “Moscow Institute of Steel and Alloys”, Russia)………………………………………..…312 Investigation on Pitting Corrosion and Intergranular Corrosion of High Nitrogen Austenitic Stainless Steels Zhouhua Jiang, Yang Cao, Huabing Li, Zurui Zhang (School of Materials and Metallurgy, Northeastern University, Shenyang China)………………………………………………..318 Author Index……………………………………………………………………………………325
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Plenary Session
Roots and Wings of High Nitrogen Steel.
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Roots and Wings of High Nitrogen Steel. J. Foct Laboratoire de Métallurgie Physique et Génie des Matériaux, LMPGM UMR CNRS 8517, Université de Lille 1, Cite scientifique, 59655 Villeneuve d’Ascq, France. ABSTRACT From a retrospect devoted to the nitrogen-metal alloys, solid solutions as well as nitrides, to the previous HNS conferences, to the scientific community involved in Research and Development of HNS, to the available articles, patents and books, from the constraints resulting from economical, environmental, technical nowadays conditions, an attempt to deduce a reasonable conjecture is proposed. According to possible industrial scenario, and to the very specific role of nitrogen alloying, the wings of the future of HNS can be outlined. KEY WORDS high nitrogen steel, state of the art.
1 Introduction To some extent the title of this article suffers a part of vagueness, this results from prudence. ”Retrospect and Prospect of HNS” which would be much more, and even too, classical often forgets that prospect is applying to a future which uses to mock foretelling. Therefore each participant to the HNS 2009 conference organized in Moscow will reconstruct his own view concerning the development of Nitrogen Alloying of Steel. Different points will be examined and discussed : -
The pedestal of knowledge on which the HNS conferences were born. The interstitial compounds and solid solutions, the specific case of each interstitial element. How the previous HNS conferences involved our scientific community. What is the role of ore resources and the coast of alloying elements? Among all possible steel grades which ones may get improved properties from N alloying ? Which major changes will affect steel making, fabrication, thermal treatments, and other processing ? How new tools (simulations, computerized modeling, microscopic and spectroscopic characterization, mechanical and corrosion properties measurements, etc) may contribute to a more efficient R&D strategy. Which data bank will be necessary to face a progressive replacement of Ni by Mn ? Innovation, actual patents and smoke curtain. What are the driving forces for future HNS conferences, does it determine a “modus operandi” for the scientific committee?
2 Evolution of the State of the Art of Nitrogen Alloying, Discussion. 21) Genesis. In contrast with carbon alloying which since some thousands of years ago, had been achieved unconsciously thanks to the combustion of wood and of coal mixed with different ores, nitrogen reaction with other elements was obtained purposely and resulted from scientific choices. The knowledge of chemical properties, of crystallographic structures, of electric and magnetic properties, of mechanical as well as refractory characteristics where studied nearly in tempo with the synthesis of nitrides and nitrogen enriched solid solutions. Very soon it appeared that chemical bonding of N with other atoms can be ionic, covalent, or (and) metallic. Therefore the choice to alloy N in steel can be considered as intentional, except when it resulted from air blowing of low alloyed steel when
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oxygen was not systematically utilized to burn excessive C, P, S and Mn, the main consequence of this involuntary nitrogen alloying was a negative advertising because it induced grain boundary embrittlement. These preliminary remarks are likely to explain that in many parts of the world, nitrogen alloying of stainless steel started much before HNS conferences had been explicitly organized (1988). The very consistent corpus of knowledge which supported initiation as well development of HNS is clearly demonstrated by books such as “Interstitial Alloys” by H.J. Goldschmidt (1967) [1] or ”Transition Metal Carbides and Nitrides” [2].Even if the list of references quoted in more recent monographs [3-11] and HNS conference proceedings [12-21] is far from being exhaustive it is showing how impressive was and is the research on nitrogen alloying. 22) Fingerprints of Interstitial Atoms. In the present case there is no ambiguity between foreign interstitial atoms FIA and self interstitials which may result from energetic irradiation or severe cold working. In transition metals alloys, the possible FIA are light elements : H,B,C,N and O. In steel all of them behave very differently, Hydrogen is intruding the iron alloys without any invitation just taking opportunities provided by electrochemical and chemical potential gradient conditions, this leads to hydrogen induced embrittlement. In contrast with this pernicious role, Hydrogen alloying is studied for making hydrides likely to store and release gaseous H2. Boron plays the Janus game of a double agent likely to take the mask either of substitutional or of interstitial. Oxygen prefers ionic bonding in order to put the coat of Neon. Curiously if metal oxide structure is considered, diffusion of metallic cations in the O-- anion lattice can be compared with diffusion of FIA in metals. Only the behavior of C and N appears to be “honest”, meanwhile an extra 2p electron changes many things which will be examined, analyzed and explained by participants of HNS 2009. The reasons for which Helium can hardly be considered as a usual FIA in steel result from very specific observations and remarks. After bombardment of metallic alloys with energetic He ions, or nuclear-decay of radioactive elements, dramatic blistering phenomena are observed. Qualitatively it means that individual atoms of the inert gas have a too large energy potential reservoir to stay single and this in any position either interstitial (tetrahedral and octahedral) or in substitution (filling a vacancy), this is consistent with the dramatic modification of metallic bonding at short and larger range, shown by ab initio calculations. In addition the activation energy for diffusion of He is extremely low. 23) Human Factor and the involvement of the “HNS Community “. Our late friend Alan Henry identified this sociologic group and wrote excellent lines about the subject (in [9] ). After his PhD, and some industrial activity, Alan was in Newcastle UK achieving research with Kenneth Jack, the author of so many important discoveries in the field of nitrides, N alloyed solid solutions and ceramics as inventor of SiAlONs. During the stay of Alan in Lille as an invited Professor, we profited from advises and encouragements of Ken Jack. In parallel an active research group in ETH Zuerich (CH) leaded by Markus Speidel studied actual industrial nitrogen steel in collaboration with Gerald Stein from VSG, a division of Krupp Steel Industry. Because the see shore favors holidays as well as meeting friendly and clever colleagues, our first brain trust with Markus revealed to be extremely useful to make more precise the frame of the future HNS 88 conference. It became possible to contact well known scientists to become members of the HNS community. Among them those quoted by A. Hendry : Hans Berns, Gerald Stein, M. Hillert, M. Kikuchi, P. Uggowitzer, H. Feichtinger, Eric Mittemeijer, Ken Jack,W. Owen, P. Lacombe, R.P.
Roots and Wings of High Nitrogen Steel.
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Reed, T. Rashev and many other eminent colleagues. It was a great satisfaction for the author of the present paper to discover that despite the iron curtain, scientific choices developed in a powerful institute such as MPI Kurdjumov in Kiev, by V.G. Gavriljuk, were very similar to those made in much more modest Laboratories such as the one I was founding in Lille. This example demonstrates how, once over-passed the nucleation barrier of the initial conference, the community could continue to grow in quantity as well as in quality. It is a common observation for scientists to measure how poor relative is the History of Sciences and Techniques in comparison of the political, economical, military ones. Although the evolution of techniques determines the society, the study of techniques is often more devoted to anecdotes than to the history of ideas and concepts. This remark does not deserve any trace of bitterness, but just a comment, pertinent concepts of R&D, very often if not always, do not originate from a unique well identified brain, but from the cross fertilization of different intellectual sources. It is precisely the aim and the result which the sequence of HNS conferences aimed and reached. In 1990 the HNS conference [13] organized by H. Witulski and G. Stein in Aachen (D), the first actual capital of Western Europe, demonstrated the impact of nitrogen alloying in the world steel industry thanks to a solid and convincing group of participants involved in steel making, properties, performances, applications etc. coming from all the different places of a world becoming more and more global. The success of the steel making route ESPR (electro-slag pressure remelting) made possible thanks to the exceptional engineering abilities of our late friend Gerald Stein and the application to hyper-resistant retaining rings, stimulated competing processes based on optimized alloying, plasma nitiding, counter-pressure casting, powder metallurgy, mechanical alloying etc. Three years later, V.G. Gavriljuk and V.M. Nadutov organized HNS 93 in the prestigious Kurdjumov Institute of Kiev (at that time CCCP) which provided a new opportunity to enrich the theoretical corpus of the domain as well as the field of actual and possible applications [14]. The fruitful collaborations between institutes of different countries which where initiated were shown to be stable as proven by further publication in particular [9]. A special issue of the well known ISIJ International Journal from “Iron and Steel Institute of Japan” was devoted to HNS 95 which took part in Kyoto under the chairmanship of Makoto Kikuchi. Beside the continuity in the deepening understanding of physical metallurgy of nitrogen alloying shown during the meeting, it appeared that very well focused applications were under development in Japan. Because nitrogen alloying is most of the time concerning corrosion resistant steel, the review by H.G. Grabke and the companion papers [15] constitute an important element of a “HNS corrosion data bank”. As a result of the very good iron ore mines in Scandinavia, and of not less excellent talent of the Wiking metallurgists, the science of metals in these countries was highly developed, therefore the joined Finnish-Swedish HNS 98 organized in Espoo and Stockholm by Hannu Hanninen and Steffan Hertzman brought unforgettable souvenirs and indispensable data collected in a key proceedings volume [16]. As a ( may be: the) cradle of the steel civilization, India became in 2002 the privileged place (Chennai) for the organization of HNS 2002,hindered to participate to this conference I just imagine how fascinating it could be to discuss about nitrogen alloying in this immense garden of the human culture. The quality of the papers collected in a special issue of the Transactions of the Indian Institute of Metals edited by Baldev Raj, Markus Speidel, U. Kamachi Mudali, V.S. Srinivasan gives an unambiguous evidence of the high level of this meeting [17]. The Rhine valley irrigates the Western European culture. Near the triple point formed by the meeting of the borders between Germany, France and Switzerland a picturesque town,
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Plenary Session
Schaffhausen did welcome the HNS 2003 conference organized in honor of the continuous contribution of Professor Markus O. Speidel to the Research and Teaching of Metallurgy. The corresponding proceedings edited by M.O. Speidel, C. Kowenda and M. Diener [18] constitute an other milestone on the progressing route of HNS. In 2004, thanks to the active and efficient involvement of Nuri Akdut and the scientific expertise of Bruno De Cooman the dense HNS 2004 conference in Ostend (Be) was organized, the weight of the international committee increased, but sadly suffered the loss of Alan Hendry and of Gerald Stein. Those who examined the proceedings [19] may also remark the increasing participation of academic and industrial researchers at an actual global scale. It can be considered as the announcement of the organization of further HNS conferences in other giant ( by the talent of their scientists) countries: China and Russia . The reader of the HNS 2006 proceedings [21] (edited by Han Dong, Jie Su and Markus O. Speidel) who could not listen the lectures in the magic landscape of Jiuzhaigou, (Sichuan, China), may identify different lines of force : 1) the continuity of the scientific strategy :2 ) the displacement of the center of gravity of the science and industry of metals at the global scale,3)the awareness of the role of non scientific parameters. The soil which did breed so many geniuses was pregnant with the HNS 2009 conference, the time of delivery is sounding, the prominent research achievements of Russian Institutes ( Moscow, Petersbourg, Ekaterinebourg, Tomsk, Novosibirsk etc.) are making the HNS community convinced of the success of the HNS 2009 conference organized in Moscow under the chairmanship of Professors A.G. Sviashin , D.V. Livanov, K.L. Kossyrev. It would be a mistake to suggest that the sequence of HNS conferences could hold a monopolistic duty to install the milestones of the knowledge of nitrogen alloying, many very interesting results have been discussed in other well (and less well) known conferences [20]. The impressive ability of our colleagues to discover possible nuggets in the WEB exempts from further quoting. 24) The constraints of steel alloying and consequences on HNS development. Nitrogen alloying of steel needs processes which are a bit more sophisticated than usual ones, even if many routes may become easier and therefore cheaper. In order to counterbalance the coast of fabrication the outlay devoted to alloying elements has to be reduced. For stainless steel Cr will still be necessary, it is not the case for nickel. Ni is needed as a gamma former, but it is a too expensive metal. The probability to observe a still growing market price of Ni ores is supported by the strategic demand of Ni based super alloys. Concerning the allergic properties of Ni, which should not impede the fabrication of retaining rings, it will reduce its usage for prosthetic and jewelry. Reduction or suppression of Ni in some steel grades constitutes the most convincing argument in favor of N alloying. This unavoidable progressive and sometime partial replacement will demand a large effort of R&D. In reality the substitution will not be simply Ni by N but Ni by N+Mn. Manganese is a tricky element, it uses to make the job of researchers more difficult from the point of view crystal structure; of phase transformations; of reactivity; of oxides, hydroxides, compounds; of corrosion; of mechanical properties; of thermo-mechanical treatments etc. Therefore new steel grades will feed the curiosity of many doctorants, engineers, and academics…These grades, such as Z0NiSS, LNiSS, defined in the FeCrMnN and FeCrMn NiN systems would be stainless, tool, as well as construction steels. There is a priori no reason to exclude neither any type of morphology, austenitic, martensitic,
Roots and Wings of High Nitrogen Steel.
duplex,TRIP,TWIP, pearlitic etc; nor any of the bainitic, bake-hardening, ausforming, etc.
19
different treatments such as mar-tempering,
25) Consequences on R&D strategy. Beside the Royal (or Revolutionary !) way of inspired innovation concerning: steel making, new grades and treatments, which is unpredictable, systematic and to some extent repetitive works have to be carried on: thermodynamics measurements and calculations about multi-component diagrams, phases transformations TTT and CCT curves, rolling and forging conditions and textures, influence of nitrogen alloying on recovery and recrystallization, round robin programs on mechanical and corrosion properties. Many breakthrough concerning hard means such as thermo-mechanical equipments, new microscopies, spectroscopies, testing machines, treatment simulators, as well as soft intellectual or virtual codes and concepts are likely to meet efficiently any gap in our knowledge. The engineering demand of HNS supposes the exact knowledge of data concerning fatigue life, corrosion fatigue, stress corrosion, toughness, critical stress intensity factor, etc. The problem of welding nitrogen supersaturated steel may be overstated, many results show that degassing is not as severe as faired. Curiously the reasons for this good surprise did not seem to stimulate critical analysis [22]. A possible explanation could result from answers to questions such as : 1)What happens when a single NH3 molecule is in perfect vacuum far from any other molecule or material? 2) What is the driving force for a single N atom to leave a metallic alloy towards the gas phase when other N neighbors are screened by metal atoms ? Up to now the necessary “HNS data bank“ does not seem to be achieved. It is not excluded that “here or there” would exist discrete small banks, meanwhile it shares with icebergs the property to be 90 percent hidden. The sequence of HNS did contribute to underline the specific role of nitrogen alloying of steel and its dual role with carbon (which originates promising C+N alloying and new H(C+N)S grades) , but it does not obey to any monopolistic view. HNS has been and will be studied in the frame of many other meetings, was critically presented in excellent text books, and international journals, this spreading interest to the present subject reinforces also the need of further focused HNS labeled conferences. Distinguished and outstanding members of scientific committee as well participants of this meeting will take advantage of their own expertise to deepen or to correct the ideas discussed above, and to reinforce the “HNS Spirit”. 3 Conclusion. The economical conditions which reign in the world of 2009 and later make the usage of any expensive and frequently rare alloying elements more critical. Nitrogen alloying of steel appears for many strategic cases unavoidable. The R&D investments which are necessary to induce a dramatic and coast effective breakthrough in HNS production has to be critically analyzed. Articles presented, discussed, and published in the frame of HNS 2009 will be extremely useful to achieve this task and to determine the future of focused HNS conferences.
Plenary Session
20 References [01] [02] [03] [04] [05] [06] [07] [08] [09] [10] [11] [12] [13] [14] [15]
[16]
[17] [18] [19] [20] [21] [22]
Goldschmidt H. J., Interstitial Alloys, Butterworth, Pub. U.K., (1967). Louis E. Toth, Transition Metal Carbides and Nitrides, Academic Press, (1971). V. A. Gubanov, A. L. Ivanovski, V. P. Zhukov, Electronic structure of refractory Carbides and Nitrides, Cambridge University Press, (1994). Kunze J., Nitrogen and Carbon in Iron and Steel, Akademie – Verlag Berlin, (1990). Grigorova N., Carbonitrides and High Speed steels – Chemical phase analysis, Interlsoft, Sofia, (1995). Rashev T., High Nitrogen Steels, Metallurgy Under Pressure, Publishing House of the Bulgarian Academy of Sciences, Sofia, (1995). Murata T. and Sakamoto M., Nitrogen-Alloyed Steels – Fundamentals and Applications – Edited by Imay Y., AGNE Publishing Inc., (1997). Gavriljuk V. G. and Berns H., High Nitrogen Steels, Springer-Verlag Berlin Heidelberg, (1999). High Nitrogen Steels and Stainless Steels, Edited by Kamachi Mudali U. and RAJ B. Manufacturing, Properties and Applications, Narosa Publishing House, Chennai, Inde, (2004). Medovar B. I., Marinskii G. S., Latash Yu. V., Tikhonovskii A. L. and Torkhov G. F., Special Electrometallurgy, Vol. 1, No 5, Sov. Tech. Rev. C, (1989). U. K. Mudali, B. Raj., Corrosion Science and Technology, C.RC Press, (2009). High Nitrogen Steels, HNS’88, Proceedings of the 1st International HNS Conference, Lille France, Edited by Foct J. and Hendry A., The Institute of Metal, (1988). High Nitrogen Steels, HNS’90, Verein Deutscher Eisenhüttenleute (VDEh) and Deutsche Gesellschaft für Metallkunde e.V., Edited by Stein G. and Witulski H., Aachen, Germany, (1990). High Nitrogen Steels, HNS’93, Proceedings of the 3rd International Conference, Edited by Gavriljuk V. G. and Nadutov V. M., Part I and Part II, Kiev, Ukraine, (1993). Special Issue on High Nitrogen Steels, Proceeding of 4th HNS Conference, Kyoto (Japan), Edited By Kikuchi M. and Mishima Y., ISIJ International, The Iron and Steel Institute of Japan, Vol. 36, Number 7, (1996). High Nitrogen Steels, HNS’98, Proceedings of the 5th International Conference on High Nitrogen Steels, Editors Hänninen H., Hertzman S., Romu J., Trans. Tech. Publications LTD, Espoo, Finland, and Stockholm, Sweden, (1998). High Nitrogen Steels, Part A, Part B, Edited by Baldev Raj, Speidel M., Kamachi Mudali U., Srinivasan V. S., Trans. Indian Inst. Met., Vol. 55, No 4 and No 5, Chennay, Inde, (2002). High Nitrogen Steels, HNS’2003, Edited by Speidel M. O., Kowanda C. and Diener M., Institute of Metallurgy, ETH Zürich, Switzerland, (2003). High Nitrogen Steels, HNS’2004, Edited by Akdut N., De Coomand B. C. and Foct J., International Conference, GRIPS media, Ostend, Belgium, (2004). Speidel M. O. und Uggowitzer P. J., Ergebnisse der Werkstoff – Forschung, Band 4, Verlag der Schweizerischen Akademie der Werkstoff-Wissenschaften, (1999). Proceedings of International Conference HNS’2006, Han Dong, Jie Su, M. U. Speidel Eds., M.I.P., Beijing, (2006). Foct J., Steel Research, in the press, ( July 2009).
The Recent Progress of Product Technologies of High Nitrogen Stainless Steels in China
21
The Recent Progress of Product Technologies of High Nitrogen Stainless Steels in China Han Dong, Yuping Lang, Fan Rong and Jie Su Central Iron and Steel Research Institute, Beijing 100081, China (Extended Abstract) Overview of R&D activities on high nitrogen steel The requirements for high strength, high toughness and corrosion resistance have promoted the development of high nitrogen steels (HNS). And people have also born an idea that the addition of nitrogen to stainless steel could lead to the reduction in the consumption of nickel. This kind of stainless steel has attracted the interests from academy, government, and industry. Since the late of 1980’s, retaining ring steel Mn18Cr18N forgings for power station application has been studied by researchers in China. They devoted their efforts especially on ESR melting, hot forging, properties and applications. During 1990’s, steel people in Shanghai metallurgy industry carried out a joint R&D activity on HNS with Prof. T.Rachev, and resulted in a good foundation for understanding HNS until now. From the beginning of 21st century, people in China have further realized the potential wide applications of HNS, due to the extended exchange of people, especially Prof. M.O.Speidel’s multi-visits to China and other international HNS society people’s influences. Thanks to HNS 2006, this really promoted the interests and activities on HNS in China. Other important motivations for the development of HNS are the requirements for high performance stainless steel by the market and sustainable development, and the shortage of nickel due to the fast growing of stainless steel production in China. In 2008, 6.94 million tons of stainless steel has been produced in China (small reduction compared with the output of 7.20 million tons in 2007). The situation of market, production and sustainable development mean a good potential for the development of HNS nowadays and in the foreseeable future. So, quite a bit of people are now involved in the R&D of HNS. Nowadays, several research groups have been led by the famous Chinese professors to develop HNS technologies and products, Table 1. The understanding of behaviors of high nitrogen steels and the related mechanism The ductile to brittle transition phenomena occurred in 22Cr-16Mn-N steels has been considered to be related to the content of nitrogen in steels. A regression relation has been established to describe the DBTT, DBTT(oC)=300CN-303. As the nitrogen content is over 1%, the DBTT could be around room temperature. The addition of nitrogen in steel leads to the reduction of stacking fault energy, and promote the formation of twins during deformation and decrease of effective slip systems. The interaction of octahedral interstices of nitrogen atoms with dislocations and stacking faults results in the difficulty of lattice slipping, and leads to tendency of brittleness. The isothermal experiment was conducted to study the precipitation of nitrides in HNS. There are two kinds of nitrides in 22Cr-16Mn-N steels: Cr2N and (CrFe)2N1-X. The derivation of PTT curves was the result of precipitation study and will give a help to the people in the development of steel products. The mapping of hot deformation of 22Cr-16Mn-N steels was carried out. The DRex, unstable, hard areas were identified in the deformation map. The increased of stress and the decrease of hot ductility is the result of the addition of nitrogen. Deformation activation energy could be raised by the increase of nitrogen content. The recommended window for hot deformation is in 1000 – 1150oC. An equation of hot deformation was derived based on the experiments: 746500 ε& = 2.62 ×10 28 [sinh(ασ p )]6.45 exp(− ) RT Table 1 Research groups in China concerning with the R&D of high nitrogen steels
Plenary Session
22 Institute
Topic
Professor
Funding
Steel CrMnN, CrMnNiN 0Cr21Ni6Mn9N 0.1C-22Cr-16Mn1.5Ni-0.6N 0.1C-16Cr-8Mn-1.5Ni0.22N Mn18Cr18N Melting technology Fe-18Cr-15Mn-2Mo0.62N Fe-17Cr-14Mn-2Mo1Cu-0.43N Fe-18Cr-16Mn-2Mo(0.52-0.81)N Fe-17Cr-1V-(0.601.02)N Fe-24Mn-18Cr-3Ni0.62N Fe-24Mn-13Cr-1Ni0.44N
Central Iron & Steel Res. Inst.
Steel products, property, welding
Jie Su
MOST973
Northeast Uni.
Melting, ASR
Zhouhua Jiang
NSF
Institute of Metal Res, CAS
Biomedical SS, mechanical property
Ke Yang
MOST863 NSF
Wuhan Uni. of Sci. & Techn.
Steel preparation, corrosion
Guangqing Li
NSF
Jiangsu Uni.
Wear, corrosion and fatigue
Qixun Dai
NSF
Yanshan Uni.
Hot deformation Forging
Wantang Fu
NSF
Mn18Cr18N
Taiyuan Heavy Machinery Inst.
Hot deformation
Huiguang Guo
NSF
Mn18Cr18N
The progress of high nitrogen steel product technologies The baseline for the design of new HNS is at 0.2% proof yield strength level of 500 – 600MPa with the cost reduction of 20% at least compared with AISI 304 steel. It means the yield strength of long products and flat products will be doubled compared with conventional both stainless steels and plain carbon steels. There are strong demands for steel products with the target strength level. In the alloy design, one of the key points is to reduce the temperature range of delta phase and increase the solubility of nitrogen during the solidification. Two steels have been designed based on the Thermo-Calc software: 22Cr-16Mn-2Ni high nitrogen steel and 17Cr-8Mn-2Ni metastable steel. The welding of HNS was investigated by using thermal simulation technique, gas tungsten arc welding (GTAW), metal inert gas welding (MIG) and laser welding. The results of the thermal simulation indicated that the hardness of the HAZ was higher than that of the base metal. The impact toughness of coarse grained HAZ was improved at first and then decreased with the increase of the cooling rate. In the welding of HNS by GTAW and laser welding, the nitrogen content of the weld metal increased as the nitrogen in the shielding gas containing Ar and N2 was increased under the same heat input. For laser welding of HNS, The higher the heat input, the less the porosity in the weld metal existed. A Cr-Mn-Ni-N steel welding wire was developed to weld HNS. The strength of the weld metal obviously was raised by adding N2 into the shielding gas in GTAW. The strength of the weld joint by MIG matched with that of the base metal. The weld joint by GTAW showed good toughness, and the same case occurred in the weld joint with single pass by MIG. It should be born in the mind that low toughness was the result of welding cycles in the HAZ of the multi-pass weld joint by MIG. Retaining ring steel Mn18Cr18N (equivalent to P900) has been drawing the attention of the people for decades. By the efforts in research on the melting technology, hot ductility,
The Recent Progress of Product Technologies of High Nitrogen Stainless Steels in China
23
recrystalization, FEM analysis, forging and cooling technique, the problems concerning with hot cracking, coarse grain and abnormal grain growth, uniformity have been solved. The forged retaining rings are now used in the power station. The construction code requires the life of the bridges, the high-rise building and the important building to be over 100 years. Another urged demand is the application of high strength steel rebar to reduce the consumption of steel. HNS rebar developed certainly meets the both requirements. 1Cr22Mn15N0.56 steel rebar was produced by the process of AOD melting, ingot casting, and hot rolling. The rebar presents a good combination of strength and ductility: Rp0.2=550MPa, A=62%, Y/T=0.58. The yield strength to ultimate tensile strength ratio meets the standard requirement of less than 0.80, which is the key index for the building to withstand earthquake. For conventional microalloyed steel rebar and fine grained steel rebar, the Y/T ratio is around 0.8. The HNS rebar possesses the characteristics of grade IV rebar (high strength), anti-earthquake (lower Y/T ratio), and long duration (corrosion resistance). The future work is to lower both alloying cost and processing cost. Combined with cold drawing, a higher strength level can be achieved in HNS wire. The attempts have been made to the applications of HNS wire rod. One of these is the high strength bolt with good heat resistance property at 350oC compared with AISI 304. The bolt made of HNS possesses high strength both at room temperature and at intermediate temperature: Rm=1100MPa at RT and Rm=725MPa at 350oC, compared with AISI 304 steel bolt of Rm=700MPa at RT and Rm=495MPa at 350oC respectively. Cold drawing wires of HNS have also been used in the riddle wire mesh. The Cr17Mn12Ni1.3N steel meshes were used instead of AISI 205 and AISI 304. The hot rolled wire rods in diameter of 5mm were cold drawn into the triangle wires. The application showed the benefits of non-magnetic, wear resistant and corrosion resistant behaviors of HNS. The HNS plate has found its way to the applications in security: the container to protect explosion blast, the equipment to proof bullet, and strongbox to withstand impact and flame cutting. These are all due to the high strength, high ductility, high strain hardening exponent even at high strain rate, and hard to be cut by conventional flame. Accompanied with the deep understanding of HNS, the innovation of affordable processing, the requirement for high performances and the sustainable development of the society, HNS may certainly be produced in a variety of products and be applied widely in the foreseeable future. Acknowledgements The authors would like to present sincere thanks to the colleagues of our HNS team: Prof. M.O.Speidel, Prof. T.Rachev, Mr. Lixin Wang, Prof. Xifan Kang, Mr. Jie Shi, Mr. Haitao chen, Dr. Lin Zhao, Dr Yuxi Ma, Ms Yuli Gu, for their great contributions in the R&D of HNS in CISRI and Chinese steel plants. The great respect should be given to the professors listed in the Table 1 who involved their efforts to study HNS and made excellent progresses in this field.
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Plenary Session
Progress in the Research and Application of Nitrogen-Alloyed Steels O.A. Bannykh Baikov Institute of Metallurgy and Materials Science, Russian Academy of Sciences ABSTRACT The basic problems of physical metallurgy, which, from author’s point of view, are of interest for the successful development of nitrogen-alloyed steels of different structural classes, are considered. The importance of the studies of such steels is noted, since they are generally characterized not only by low cost and high corrosion resistance, but also by good mechanical properties. It is noted also that high-strength nitrogen-alloyed steels (first of all, of martensitic and martensitic-austenitic classes) can be considered as the possible substitutes for aluminum and titanium alloys without any deterioration of specific strength and with a gain in technological properties. For this reason, there is a possibility of using nitrogencontaining steels in the articles of the modern equipment, requiring lowered weight and high service reliability. KEY WORDS nitrogen-alloyed steel, structure, mechanical properties
It has long been known that nitrogen belongs to the strongest austenite-forming elements and it is possible to substitute nitrogen for nickel in corrosion-resistant austenitic steels. Such substitution makes it possible to reduce the cost of steels. Just this specific feature was a stimulus to the progress in the research of nitrogen-alloyed steels. The following additional important features of nitrogencontaining steels were revealed upon the accumulation of the knowledge about their properties: increased ability to strengthening upon plastic deformation; high static and, especially, cyclic strength; hardness; wear resistance. In this case, the resistance to most types of corrosion for the nitrogen-alloyed steels compared with the conventional chromium-nickel steels does not decrease, but even increases in some chemical media. It became clear that not only cost effectiveness, but also mechanical properties should be considered as an important stimulus for the search of the new types of nitrogen-alloyed steels. It was shown that high-strength nitrogen-containing steels can be used as substitutes for titanium and aluminum alloys since they are as good as these alloys in specific strength and surpass them in technological effectiveness. In this case, it is possible to reduce the weight of the articles of modern equipment without any loss of reliability. The combination of high corrosion resistance and good mechanical properties makes it possible also to successfully substitute the nitrogen-alloyed steels for bronze. It is noteworthy that the above substitutions are advantageous also from the economic point of view. Thus, the field of the reasonable application of nitrogen-alloyed steels was substantially enlarged. Nitrogen at present offers wide prospects for the alloying of high-chromium austenitic, martensitic, martensitic-austenitic, and austenitic-ferritic steels. In these steels, nitrogen is one of the main alloying elements, determining their physicochemical properties. For the realization of the above characteristics of the nitrogen-alloyed steels, the metallurgists and the metallographers had to solve several complex problems. The first of the most important problems, which define the development of the steels containing nitrogen as the alloying element rather than as a technological impurity, is the search for the methods to introduce the necessary quantity of nitrogen into steel. We can state with satisfaction that there are significant successes in the development of a whole series of technological processes of the nitrogen introduction both into molten and into solid steel. Nevertheless, there sill remains a wide field for new solutions of the problem of the steel saturation with nitrogen, in particular, by the methods of powder metallurgy, plasma technology, and ion implantation. The second problem, which is of special importance for two-phase austenitic-ferritic steels, is to establish the regularities of the nitrogen distribution in the crystal lattice of each phase present in the
Progress in the Research and Application of Nitrogen-Alloyed Steels
25
steel. This distribution affects the mechanism and kinetics of the phase transformations under the thermal actions on the metal both in the initial state, for example, after solidification or sintering, and after heat treatment and plastic deformation. This problem is closely related to the action of nitrogen dissolved in austenite on the formation and distribution of lattice defects, which not only determine the strength, plasticity, and ductility of the steel, but also effectively influence its chemical properties. In this association, it is noteworthy that the nitrogen-alloyed austenitic and austenitic-ferritic steels compared with the conventional chromium-nickel steels are more intensely strengthened upon plastic deformation. This is caused by the reduction in the austenite stacking fault energy, which hinders the cross slip of dislocations. The theoretical and experimental studies of the specific features of the interaction between nitrogen dissolved in steel and the alloying elements, in particular, determining the solubility of nitrogen in the phases present in the steel are already fulfilled for the solution of the problems related to the development of the steels with required properties. It should be noted that the studies of such specific features were performed both for the tetriary systems “Fe-Cr-N-alloying element” and for multicomponent systems. Substantial advances were achieved in this direction, but there are still many “blind spots”. Only Fe-Cr-N ternary system is reliably investigated, and only Fe-Cr-MnN tertiary system is studied in some detail. The Fe-Cr-Mn-Ni-N and Fe-Cr-Mn-Ni-Mo-N systems, which are of great practical importance, require further detailed investigation. It is also very important to establish the character and intensity of the effect of nitrogen in combination with other alloying elements both on individual physicochemical properties of steels and on the combinations of these properties. As a result of numerous studies performed in many countries, the basic knowledge is obtained about the possibilities of the alloying of steels with nitrogen for the directed action on their properties. A great contribution into this information is introduced by the proceedings of the traditional conferences on high-nitrogen steels. The greatest attention of researchers was paid to the study of nitrogen-containing austenitic steels. They occupied the leading position due to a relatively simple possibility to reduce the nickel content in the conventional chromium-nickel corrosion-resistant steels, and even a simpler task of the substitution of nitrogen for carbon in such steels without any reduction in the nickel content. The application of melting under high pressure of nitrogen made it possible to obtain austenitic chromium-nitrogen steels, which are free from nickel and other austenite-forming chemical elements (for example, manganese, cobalt, or copper). Such chromium-nitrogen austenitic steels have a high biological compatibility with the human body, and this makes it possible to recommend them for the implants, which can compete with titanium alloys. If we consider the austenitic nitrogen-containing steels from the positions of their cost effectiveness, we can conclude that the basic difficulties at present are already overcome, and it is hard to expect any fundamentally new scientific solutions. However, there is a wide field for the studies directed toward the rational use of high-nitrogen steels, which have enhanced mechanical properties and are nonmagnetic. For such steels, the new fields of application are opened, for example, in shipbuilding industry. It seems to be reasonable to study high-strength steels with the maximum attainable nitrogen content or close to it. It should be taken into account that, under thermal effects, the nitride formation begins to play an important role in high-nitrogen steels. The size, distribution, and volume fraction of nitrides are of great importance for the realization of the required strength, plasticity, ductility, and corrosion resistance. The simplest and most reliable method to obtain high nitrogen content in steels is to increase the chromium content in its composition. In this case, the precipitation temperature of chromium nitrides decreases, and their dissociation temperature increases. At temperatures ranging between ~500 and ~900°C, the steels with the so-called "overequilibrium" nitrogen content undergo the discontinuous decomposition of austenite. The resulting pearlite-like structure consists of the austenite depleted of nitrogen and nitrides at the first decomposition stage and of ferrite saturated with nitrogen and nitrides at the final stage. In this case, the remaining non-decomposed austenite is
26
Plenary Session
maximally saturated by nitrogen. The steel with the structure formed by the discontinuous decomposition is characterized by decreased toughness and corrosion resistance. Generally, it is very difficult to suppress the discontinuous decomposition in massive parts or semifinished products. For such products, the requirements for the selection of the homogenizing heat treatment, in particular, the cooling regime after homogenization are especially severe. The elevated temperature of this treatment causes an undesirable austenite grain growth and in many instances leads to the formation of delta ferrite. These processes can be controlled, first of all, by an increase in the contents of the austenite stabilizing elements, for example, Mn, in the steel composition. One of the difficult tasks, which should be solved for the processing of high-nitrogen steels by ingot technology, is the formation of pores upon the ingot solidification. The high-nitrogen steel ingots free from developed porosity can be obtained due to the selection of the combination of alloying elements suppressing the formation of delta ferrite, in which the nitrogen solubility is substantially lower than in austenite, during solidification from the melt. An additional series of difficulties arise during the development of high-nitrogen steels providing not only high mechanical properties, but also good technological properties, which are of more specific character and require the special consideration (for example, the realization of the properties, which either are uniformly distributed, or change according to a predetermined temperature gradient over the volume of large semifinished products and articles). In recent decade, many studies of so-called corrosion-resistant steels with the dual (austenite + ferrite) structure were carried out. Note that such steels alloyed with nitrogen are the most complex objects for the optimization of chemical composition and regimes of heat or thermoplastic treatment. The achievement of high corrosion resistance is due to the introduction of significant quantities of elements such as molybdenum, tungsten, niobium, or titanium into the steel composition. Various intermetallic phases and nitrides can form under the thermal action, for example, upon welding. I will not consider the processes of the precipitation and decomposition of intermetallic phases. It is possible to note that the distribution of nitrogen between ferrite and austenite in the steels with the dual structure after homogenizing heat treatment depends not only on the homogenization temperature, but also on the chemical composition of the steel. At the same relationship between ferrite and austenite contents, the relationship between nitrogen concentrations in ferrite and austenite significantly differs in the steels of different chemical compositions. This relationship affects the development of corrosion processes. In the first approximation, it is possible to note that the steels with minimum difference in the nitrogen contents in ferrite and austenite should exhibit the best corrosion resistance. In this association, for the steels with the dual structure, one should critically examine the empirical polynomials, which reflect the resistance to pitting corrosion (PRE) and slit corrosion (MARC) as a function of the contents of chromium, molybdenum, tungsten, nitrogen and carbon, and other elements. It is evident that the coefficients preceding each element in these polynomials for the same steel are constants, but they change with changing heat-treatment regimes (for example, homogenization temperature). The steels with martensitic or martensitic-austenitic structures hold a prominent place among nitrogen-alloyed steels. At present, the advantages of the martensitic steels containing 12-16% Cr, 2-5% Ni, and up to 0.2% N over nitrogen-free steels containing up to 0.2% C can be considered as reliably established. In addition to the noticeable increase in strength, the nitrogen-alloyed steels demonstrate a higher tprocessing plasticity. The most obvious cause for the increased plasticity is the presence of up to 10-15% retained austenite in the structure of nitrogen-containing steel. The formation of martensite from this austenite upon cold plastic deformation (for example, under the conditions of cold upsetting) can cause the TRIP effect. Another advantage of the nitrogen-alloyed steel is the increased corrosion resistance. In this association, the weld joints of the nitrogen-alloyed martensitic steel should be noted. Even without subsequent heat treatment, the weld joints in thin articles from such steel retain the satisfactory resistance to the action of some corrosion-active
Progress in the Research and Application of Nitrogen-Alloyed Steels
27
media, for example, sea water. This is caused by the exception of the possibility of the Me23C6 carbide formation in the heat-affected zone of the weld joint. The nucleation and growth of the particles of this carbide occur rapidly: this process is controlled by carbon diffusion, since the metal part of this carbide can have the chemical composition corresponding to the composition of the metal elements contained in the steel. The process of nitride formation occurs substantially more slowly, since it is controlled by the chromium diffusion into the volume of the forming particle. For steels of these types, the strength up to 1800 MPa can be attained at a relative elongation of up to 25-30%, and the specific fatigue strength after 10 million loading cycles is comparable to the strength of light alloys upon single loading. One should note also that the specific rigidity, which is characterized by the relationship between the modulus of elasticity and density, for the steels under consideration is more than twice as high as that of conventional aluminum or magnesium alloys. An increase in the chromium content above 16% in the steels causes the formation of deltaferrite. In this case, the corrosion resistance can remain high, but the ductile-brittle transition temperature increases, and hot processing plasticity decreases. Nevertheless, the martensitic-ferritic steels alloyed with nitrogen can be considered as promising object for further studies. Such steels can exhibit a substantial grain refinement and the formation of the structure close or even corresponding to nano-dimensional one. The possibilities of the development of nitrogen-alloyed steels are far from being exhausted. The fields of their practical use will be extended, and it seems real that such steels in overall production will occupy the leading position among corrosion-resistant steels.
Thermodynamics & Kinetics
- 30 -
Order-Disorder Transitions in High-Nitrogen Steels: From Ab Initio to Statistical Thermodynamics
31
Order-Disorder Transitions in High-Nitrogen Steels: From Ab Initio to Statistical Thermodynamics A.J. Böttger, D.E. Nanu, A. Marashdeh Delft University of Technology, Materials Science and Engineering, Mekelweg 2, 2628 CD Delft, The Netherlands ABSTRACT Interstitial nitrogen atoms in Fe-N show short - or long - range order depending on the nitrogen content. This is due to elastic interactions between interstitial N atoms in the metal matrix. The tendency of interstitial atoms to order needs to be taken into account in order to predict phase stabilities and order-disorder transitions in high-nitrogen steels. The method utilized to predict phase stabilities combines ab initio calculations with statistical thermodynamics (cluster variation method). The cluster approximation and vibrational contributions issues and their effect on describing the phase boundaries between the γ and γ’ Fe-N phases is discussed. KEYWORDS Fe-N, ab initio, cluster variation method, order-disorder transitions, vibrational energy
1 Introduction
The formation and stability of interstitial solid solutions containing nitrogen, carbon, boron, and hydrogen in a metal host matrix is an important issue in the field of steel making. Since thermochemical data are often hard to obtain as (carbo)nitrides appear as (metastable) precipitates or surface layers, this work aims at providing a fundamental approach to obtain the thermodynamic properties of Fe nitrides starting from first principles calculations. The elastic interactions between interstitial atoms in a metal matrix generally do not favor a random distribution of the atoms over the interstitial sublattice. As observed for N in Fe-N in γ, γ’ - Fe4N1-x and ε - Fe2N1-z short-range order and/or or long-range order can occur depending on the nitrogen content. This tendency of interstitial atoms to order needs to be taken into account in order to predict phase stabilities in highnitrogen steels. However, if the Fe - matrix contains alloying elements, the presence of interstitial atoms can also induce order-disorder transitions on the metal sublattice. Two different approaches to account for this effect are discussed and compared in the current paper. Note that the first approach focuses on ordering on the interstitial sublattice and the second approach allows including both the ordering on the interstitial sublattice and the interstitial-induced ordering on the metal sublattice. 2 Phase boundary calculations
Phase-boundary calculations are based on statistical thermodynamics, i.e. the cluster variation method (CVM) [1]. The energy functional (grand potential) used in the CVM is expressed in terms of small groups of atoms (clusters) for which correlated occupancies are considered. The orderdisorder transition between the austenite γ and nitride γ’ - Fe4N1-x phase is modeled here. For the purpose of modeling, the structure is described as consisting of a fully occupied FCC host matrix, formed by metal atoms, with nitrogen occupying the octahedral sites only (see Fig. 1). The sublattice consisting of the octahedral sites also forms a FCC structure with its sites occupied by N and/or vacancies. The two approaches that will be discussed in this paper are (i) the tetrahedron approximation: in this case the interstitials are described in the mean field of the metal host and the system is described using as the basic cluster a tetrahedron formed on the interstitial FCC sublattice (see Fig. 1) and (ii) the cube cluster approximation: in this case the description includes both the metal and the interstitial sublattices (see Fig. 1). In the current paper these approaches will be compared and discussed. In addition it will be shown that vibrational contributions, to both the internal energy and the entropy, affect the calculated phase equilibriums.
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Thervjdenamics & Kinetics
2.1 Ab initio calculations
For the tetrahedral approximation the internal energy at 0 K has been calculated using the Cluster Expansion Method (CEM) [2-5]. The CEM connects the internal energies of ordered ground-state structures calculated by first principles with that of any structure. In this approach the internal energy of a substitutional system (i.e. here, the interstitial sublattice) is expanded as a rapidly convergent series of multisite cluster interaction energies (from empty to the tetrahedron cluster energies), the cluster interaction energies being determined from first principles calculated internal energies of the ordered structures. For the cube approximation the internal energy at 0 K has been calculated using the cube approach [6] in which the energy is written in terms of energies of cubes consisting of four interstitial and four metal sites; the energies of all cubes being directly determined from first principle calculations. For a binary system such as Fe-N five different clusters are considered i.e. the interstitial sublattice site occupation ranges from 0 to 4 N atoms while the metal host lattice is fully occupied by Fe atoms.
}
Fig. 1. The unit cell of the cubic compounds γ and γ’ - Fe4N1-x: the metal (Fe) sites are indicated by and the interstitial (N or vacancy) sites by {. The tetrahedron cluster (comprising interstitial sites only) and the cube cluster (comprising metal and interstitial sites) are indicated.
In the present work the structure optimizations for the ground-state structures are given in [3] and the optimizations of the cube clusters, i.e. the equilibrium Fe atomic positions and the lattice parameters, are performed by the projector-augmented wave method [7,8] as implemented in the VASP code [9,10], where the electronic structure calculations are carried out using the PW91 generalized gradient approximation. The complex ferromagnetic nature of Fe-N compounds requires that the first principles calculations are performed within the spin-polarized approximation. The presence of two kinds of FM (ferromagnetic) states [11,12] i.e. the low spin FM and the high spin FM state needs to be considered. The actual state depends on the nitrogen content. For pure Fe the low spin state is of lower volume (and magnetic moment) than the high spin state; since N increases the volume the high spin state is used for the calculations.
Order-Disorder Transitions in High-Nitrogen Steels: From Ab Initio to Statistical Thermodynamics
33
2.2 Cluster Variation Method
The CVM enables to take into account the correlated occupancies of sites. To this end the configurational entropy of the ordering of N and vacancies on the interstitial sublattices of γ and γ' Fe4N1-x is described as a function of the cluster distribution probabilities. The frequency of occurrence of all possible arrangements of N and vacancies on the tetrahedron and its subclusters (points, pairs, etc.) is considered. Note that since in the case of the binary Fe-N system the metal sublattice is fully occupied by one type of metal atoms, the calculation of the entropy in the cube cluster approach and tetrahedron approach are the same. The total energy consists of the internal energy in terms of cluster interactions based on the ab initio calculations and the configurational entropy. In order to include temperature effects the vibrational internal energy and vibrational entropy, both functions of the volume and temperature, are estimated using the Debye - Grüneisen model according to [13]. The properties required to estimate the Debye temperature, such as the bulk modulus and its derivative, and the equilibrium volumes are obtained from the ab initio calculations via an equation of state [3]. 3. Results and Discussion
The calculated γ and γ' - phase boundaries are compared with literature values (obtained from experiments) in Fig. 2. The results show that vibrational contributions have different effects on the stability of the γ - phase and the ordered γ' - phase. Clearly the calculated (γ + γ')/γ' - phase boundary agrees well with the experimental data regardless of the consideration of vibrational effects. The γ/(γ + γ') - phase boundary is strongly affected by the vibrational contributions. Both the tetrahedron CEM and the cube approximation underestimate the N solubility in the γ - phase in the case that temperature effects (i.e. vibrational contributions) are not taken into account. Including vibrational contributions is necessary to get a good agreement between the calculated and experimental phase boundaries.
Fig. 2 (a) the Fe - N phase diagram (b) Enlargement of the γ and γ' - phase boundaries: : literature …… data, - - - : cube approximation, ⎯ : tetrahedron CEM approximation, : cube approximation including vibrational contributions, –xx– : tetrahedron CEM approximation including vibrational contributions.
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Thervjdenamics & Kinetics
Although the correspondence between the calculations and experimental data of the phase boundaries is good for both approaches some differences occur (see Fig. 2). These differences can be caused by the approximations made in the calculations. The most important approximations are: the limitation in the number of basic structures used in the CEM, the accuracy of the ab initio simulations because of the approximations made in electron energies and the consequences of that on the volume dependencies (Grüneisen constant), and the way the magnetic contributions are treated, i.e., at high temperature Fe - N phases are paramagnetic whereas the low temperature magnetic state has been used in the calculations. It can be concluded that for this system, the CEM approximation is slightly better than the cube approximation. This is maybe caused by the fact that the tetrahedron CEM approach includes other than nearest neighbour interstitial interactions whereas the cube approximation includes nearest neighbour interactions only. In particular for systems such as Fe-N, in which long-range elastic interactions between interstitial N atoms play a role, the cube approximation could underestimate the energy related to long - range interactions. References [1] Kikuchi R.A., Phys. Rev. 81 (1951) 988 [2] Sanchez J.M., Ducastelle F., Gratias D., Physica 128A (1984) 334 [3] Shang S., Böttger A.J., Acta Mat. 53 (2005) 255 [4] Drautz R., Diaz-Ortiz A., Fahnle M., Dosch H., Phys. Rev. Lett. 93 (2004) 067202 [5] Zarkevich N.A., Johnson D. D., Phys. Rev. Lett. 93 (2004) 255702 [6] Nanu D.E., Deng Y., Böttger A.J., Phys. Rev. B74 (2006) 014113 [7] Blöchl P.E., Phys. Rev. B50 (1994) 17953 [8] Kresse G., Joubert D., Phys. Rev. B59 (1999) 1758 [9] Kresse G., Hafner J., Phys. Rev. B49 (1994) 14251 [10] Kresse G., Furthmüller J., Phys. Rev. B54 (1996) 11169 [11] Acet M., Gehrmann B., Wassermann E.F., et al., J. Magn. Magn. Mater. 232 (2001) 221 [12] Herper H.C., Hoffmann E., Entel P., Phys. Rev. B60 (1999) 3839 [13] Moruzzi V.L., Janak J.F., Schwarz K., Phys. Rev. B37 (1988) 790
The thermodynamics investigations of nitrogen in the liquid Fe-N-V alloy
35
The thermodynamics investigations of nitrogen in the liquid Fe-N-V alloy Artur Hutny, Jerzy Siwka Czestochowa University of Technology, 42-200 Czestochowa, Poland ABSTRACT The values of nitrogen solubility in the liquid Fe-N-V alloy and in liquid vanadium were determined experimentally. The alloy of a vanadium content of 1.5; 2.5; 4.0; 5.6; 7.8; 12.0; and 45.0, respectively, and vanadium of purity of 99.7 wt % were subjected to experiment. Using the experimental data and the findings of the previous study on the liquid Fe-N alloy, temperature relationships of inter- and self-reaction parameters have been determined. Using the coefficient values for the Fe-V alloy, calculated on the basis of our own investigation results, the values of the energy parameter h (non-random solubility model by Schmid-Fetzer) were determined. The determined energy parameter h in the form of a temperature function relationship will enable the determination of the numerical values of the for the Fe-V alloy of any arbitrary composition. KEYWORDS iron-vanadium alloys, liquid state, nitrogen activity
1 Introduction
Presently, the source of thermodynamic data are computer databases. It should be stressed that the data on the solubility of nitrogen in liquid vanadium and the Fe-N-V alloy have been derived from works published in the 60-80s of the past century. Therefore, the thermodynamic characteristics obtained from the authors’ own studies can be used for complementing and, in some cases, making the existing databases more accurate. This assumption was fully confirmed in previous works by [1-3], where the values of nitrogen solubility in the liquid Fe-N-V alloy and in liquid vanadium were determined experimentally. The composition of the metal was varied from pure iron, through the Fe-V alloy to pure vanadium. The partial pressure of nitrogen in the gaseous phase was set starting from a very low level of 0.0002 MPa in the case of liquid vanadium to 3.2 MPa in the case of liquid iron. The measurements of nitrogen solubility were taken at temperatures of 1903, 2053, 2073, 2173, 2203 and 2273 K, respectively. A basis for developing thermodynamic characteristics was the description of the nitrogen activity coefficient conforming to the principles of phenomenological thermodynamics, with the formal treating of the Wagner-Chipman equation. The following relationship was used: lgf N = e (NN ) ⋅ [%N] + e (NV ) ⋅ [%V] + rN(N ) ⋅ [%N]2 + rN(V ) ⋅ [%V]2 + rN(N,V ) ⋅ [%N]⋅ [%V]
(1)
The results of the above-mentioned works can, on the one hand, be used for the calculation of nitrogen solubility in a liquid iron alloy containing vanadium and, on the other hand, from the theoretical point of view, can be treated as original source data providing the spur for universalizing of existing physical models based, for example, on statistical thermodynamics. The aim of this study was to adapt the Schmid-Fetzer model to enable the calculation the activity coefficient γ 0N in the liquid Fe-N-V alloy for the whole range of variation of chemical composition. 2 Materials and methods
The measurements were conducted in the homogeneous range of liquid Fe-N-V and V-N solutions. In preliminary tests, the thermodynamic parameters of the point of transition from the liquid homogeneous solution to the heterogeneous solution were identified and the kinetic time constant of the test system was determined. Fe-V alloys with a vanadium content of 1.5; 2.5; 4.0; 5.8; 7.8; 12.2 and 45 wt% were subjected to testing. The saturation with nitrogen was carried out at three temperatures: 2073 K, 2173 K, 2273 K and 2373 K and within the range of nitrogen partial pressure variations in the gaseous phase of the liquid metal–gas system, which spanned from 0.001 to 3.2 MPa.
Thervjdenamics & Kinetics
36
Factors that are decisive to the adequacy and significance of the experiment are: the possibility of producing a “pure” liquid metal–gaseous phase system, the intensive agitation of liquid metal, and the rapid cooling of metal samples at a given point of the experiment. The experiment proper involved the melting of a metal sample in a state of suspension in an electromagnetic field in an appropriate reaction gas at a pressure slightly higher than 0.1 MPa, and the rapid filling of the metal reactor body with the same working gas until the preset nitrogen pressure level was attained, while heating the liquid metal up to the appropriate temperature. In this state, the specimen was held for a duration of 180 s, which was considered satisfactory for the liquid metal-gaseous phase system to reach the thermodynamic equilibrium in the conditions of the experiment. After the power supply of the levitation coil had been switched off, a sample of a mass of approx. 1 g dropped by gravity to a copper testing mould. The metal solidified under the conditions of cooling at a rate of 103 - 104 K×s-1 [4]. 3 Results and discussion
Immediately after the experiment, the equilibrium nitrogen concentration in the liquid Fe-N-V alloy was determined at the preset thermodynamic parameters of the liquid metal–gas system. For the determination of the thermodynamic characteristics of the alloy examined, and primarily the solubility of nitrogen in this alloy, a thermodynamic relationship in the following form was used, which enables the determination of the nitrogen activity coefficient for any single experimental point: lg f N = lg K N (Fe) - lg
[%N] PN 2
(2)
Figure 1 shows sample results of the measurements of nitrogen solubility in the Fe-V-N alloy for the temperature 2273 K. From the distribution of point in Figure 1 it can be seen that with increasing partial pressure of nitrogen, its solubility also increases. The character of the nitrogen solubility increase in not linear. The two above-mentioned trends correspond to the thermodynamic properties of the liquid Fe-V-N alloy. The experiment results show that the nitrogen activity coefficient, fN, takes on values less than 1.0; moreover, it depends on the nitrogen content of the alloy. This relationship is more visible at a higher vanadium content (above 4 wt%). So, the thermodynamic description of the Fe-N-V should allow for the interactions between different atom pairs, not only between the Fig. 1. Solubility of nitrogen in the liquid Fe-V-N alloy at a nitrogen-vanadium, nitrogen-vanadiumtemperature of 2273 K [5] vanadium pairs, but also between nitrogennitrogen, nitrogen-nitrogen-vanadium. After rearranging according to the substances dissolved and taking the experimental coefficient EN(N) [6], the nitrogen activity coefficient, fN, will take on the following form:
lg f N = lg f
(V) N
+ E (N) N ⋅ [%N]
(3)
The coefficient EN(N) is calculated based on the results of specific experiments, so it has an experimental character. Whereas, its internal structure contains phenomenological universal factors
The thermodynamics investigations of nitrogen in the liquid Fe-N-V alloy
37
that can be determined by the analysis of the values of EN(N) with formal treating the Wagner's theory and the Wagner-Chipman equation for their determination. Using Equation (3), the primary data were worked out with the use of the mathematical linear regression function. Discussion of results. The broad range of variations of thermodynamic parameters and the extensive database of the obtained results of experimental studies on the solubility of nitrogen in the liquid Fe-N-V alloy enabled the thermodynamic characteristics to be developed. Nitrogen-vanadium interaction in the Fe-N-V alloy. On the basis of the distribution of experimental points it was assumed that the relationship had a curvilinear character. Hence, the following form of functional relationship was taken for calculation:
lg f N(V) = e
(V) N
⋅ [%V] + r
(V) N
⋅ [%V]2
(4)
The effect of temperature on the values of the interaction parameters was determined with the use of an equation analogous to the Van’t Hoff equation, as shown in Table 1 below. Tabele 1. Temperature dependencies of the interactions and self-interactions parameters [7]
Interaction parameters
− 424 ± 42 + 0,106 m 0,02 e = T 6,1 ± 1,4 rN(V) = − 0,00195 m 6 ⋅ 10 −4 T (V) N
Self-interaction parameters
15,86 ± 15 − 0,005 m 0,007 T 0,37 ± 0,47 = − 4,2 ⋅10 −4 m 2 ⋅10 −4 T
rN(NV) = t (NVV) N
By evaluating the temperature relationships along with their result dispersions it can be stated that, in respect of the 1st order parameter, the authors' studies have yielded a definitely satisfactory description. In the case of the 2nd order parameter, greater random dispersions of the coefficients are noticeable. On this account, the functional relation rN(V) = f [%V] is less accurate, which can also be observed in the case of studies published in literature, related to vanadium and many other alloying constituents. Figure 2 shows comparison of the thermodynamic characteristics determined by the authors with corresponding literature data.
Fig. Fig. 2. The effect of temperature on interaction parameters in the liquid Fe-N-V alloy [7] a) the 1st order parameter; b) the 2nd order parameter
As shown by this comparison, the straight line defining the values of the parameter e (V) , determined N from the authors' studies, lies in the lower part of the set of straight lines determined by other authors. It follows from the above that the interactions being described contribute to the increase in solubility to a greater extent than reported by authors in earlier works. Due to the fact that the nitrogen-nitrogen and nitrogen-nitrogen-vanadium self-interaction parameters were not determined . So, the values of the literature in those works, this interaction is "included" in the parameter e (V) N parameters need to be verified.
Thervjdenamics & Kinetics
38
Similar is true for the comparison of the second-order parameters. The enthalpy and entropy parameters determined in the authors' studies are distinguished by higher dispersions. Other authors do not specify the deviations with which they determined their relationships. It should be emphasized that the parameters determined in the authors' studies, along with their deviations, do not change the parameter sign, which is consistent with the physical assumptions. In the thermodynamic description of the Fe-N-V alloy, both parameters should be considered. Nitrogen-nitrogen and nitrogen-nitrogen-vanadium self-interactions in the Fe-N-V alloy. To , an equation in the following make the thermodynamic analysis of the experimental coefficient E (N) N form was used:
E (N) N = e
(N) N
+r
(N, V) N
V, V) ⋅ [%V] + t (N, ⋅ [%V]2 N
(5)
included in Equation (5) was fundamentally derived in work [1] The self-interaction parameter e (N) N when studying the equilibrium of the liquid iron–nitrogen system in the range of PN pressure from 2
0.1 to 3.2 MPa. The temperature relationships of the parameters of self-interaction for nitrogen and nitrogenvanadium, and nitrogen and nitrogen-vanadium-vanadium are shown in Table 1. The 2nd order parameter rN(N,V) is a quantity that is determined with a specific probability (which V, V) corresponds to the standard error 1σ), while the 3rd order parameter t (N, is a quantity with an N estimated temperature trend. The latter quantity has a more mathematical nature than the physical meaning. It is, however, indispensable in the final model equation of the solubility characteristics. relative to the Figure 3 represents the distribution of the values of the experimental coefficient E (N) N vanadium content of the Fe-N-V alloy for the temperature of 2273 K. For the quantitative description of the activity of nitrogen, as has been said in the introduction, the Schmid-Fetzer nonrandom model was used, according to [18]. The model is described in detail in work [5]. Individual factors of this equation must, however, be known. Their values were determined using literature data and the results of the authors' own studies. The activity of iron and vanadium in the Fe-V alloy was determined using the commercial program THERMOCALC. Fig. 3. Distribution of the values of the experimental coefficient Then, the activity coefficients γ Fe and
in relation to the content of vanadium in the Fe-N-V alloy (T=2273 K) [7]
γ V were calculated.
The original own data on the solubility of nitrogen in liquid iron, determined experimentally by [1], and the own data on the solubility of nitrogen in liquid vanadium, obtained by [3], were converted with reference to the new standard state. As the standard state, an infinitely diluted solution of nitrogen in liquid metal was assumed, in which the concentration is expressed in mole fractions. The equilibrium contents of nitrogen in liquid iron and vanadium, expressed in wt%, were converted into atom fraction. Then the values of ln γ 0 were calculated for individual temperatures using the following expression: Ν
lnγ 0N
= −ln(
XN PN 2
)
(6)
The thermodynamics investigations of nitrogen in the liquid Fe-N-V alloy
39
0
Having the values of ln γ Ν available for specific temperatures, temperature relationships of ΔG 0N were developed for nitrogen dissolved in liquid iron and in liquid vanadium [5]: ⎡ ⎤ J ⎡ ⎤ J 0 ΔG0N(Fe) = 7167 + 48.77 ⋅ T ⎢ ⎥ ΔG N(V) = −291885 + 129.66 ⋅ T ⎢ ⎥ ⎣ g ⋅ atom N ⎦ ⎣ g ⋅ atom N ⎦
(7)
Next, in order to calculated the values of ln γ 0 concerning iron and vanadium, the following Ν(Me)
relationship was used: ΔG 0N(Me) (8) lnγ 0 = R ⋅T Considering the fact that the melting point of pure vanadium is 2163K, for the purpose of comparison within the whole range of variation of chemical composition of the Fe-V alloy, a liquid alloy test temperature of 2273K was chosen for subsequent calculations. The energy parameter h Fe-V was determined based on the work of [19]. It is constant at a given temperature and does not depend on the kind of dissolved substance (N2, O2, etc.). Its value is negative, and for the temperature of 2273K is –1335. Ν(Me)
0
Using the values of the coefficient γ Ν for the Fe-V alloy, calculated on the basis of our own investigation results, the values of the energy parameter h were determined. Figure 4 represents the results of calculations of ln γ 0Ν for liquid iron, vanadium and the Fe-V alloy, respectively. The own data are compared with the data given in the work of [20]. The calculated value of the matched energy parameter, determined from the results of our own investigation, is: h= 4 607
⎡ ⎤ J ⎢ ⎥ g ⋅ atom N ⎣ ⎦
Fig. 4. Dependence of ln γ 0Ν on the atom fraction of
(9)
vanadium in the liquid Fe-V-N alloy [5]
SUMMARY
Investigation of the solubility of nitrogen in the liquid Fe-V-N alloy was carried out, and a database with the authors' own experimental thermodynamic data was created. Comparison with literature data was made. The investigation results have shown a non-linear dependence of the coefficient fN not only on the vanadium content, but also on the nitrogen content of the alloy. The data served for the determination of the values of the matched energy parameter h in the Schmid-Fetzer model. This parameters enables the determination of nitrogen activity in the Fe-V-N alloy of any arbitrary chemical composition. The results of the investigation carried out within the present work show the necessity for the continuation of the investigation in respect of the V-N alloy and, from the point of view of the bases of the HNS steel melting technology, it should be extended to cover alloys of iron with titanium and niobium. In summary, it should be noted that the results of the present work have considerably contributed to the expansion of the original database of thermodynamic data that have been determined experimentally for years at the Department of the Extraction and Recycling of Metals of the Czestochowa University of Technology. REFERENCES
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[1] Siwka, J. D.Sc. Thessis, Diss. TUCz, serie Metallurgy, Czestochowa, 8, 1999. [2] Hutny, A., Siwka, J., Fitzner, K. The thermodynamic characteristic of liquid binary alloy Fe-V with nitrogen in hyperbaric conditions. Journal of Steel and Related Materials, vol. 2, pp. 623-627, 2004. [3] Hutny A., Siwka J.. The Thermodynamic Characteristics of Nitrogen in Liquid Vanadium. Proceed. Inter. Conf. on „High Nitrogen Steels 2006”, Eds. Dong H., Su J., Speidel M.O, Jiuzhaigou – Sichuan, China – August 2006, pp. 401-407. [4] Svyazhin, A. G., Siwka, J., Skuza Z, et al. Crystallization, the Cast Structure and Bubble Formation in High Nitrogen Steels. Transaction. Indian Institute of Metals, vol. 55(4A),pp. 183-191, 2002. [5] Siwka J., Hutny A.,. The activity of nitrogen in the liquid Fe-N-V alloy based on the non-random solubility model by Schmid and Fetzer, Proceed. Inter. Conf. on „Interstitially Alloyed Steels”, Eds. Akdut, De Coman, Kim – Pohang, South Korea – September 2008, pp. 33-36. [6] Svyazhin A.G.: Metod parametrow vzaimodejstvija i parametr samovzaimodejstvija azota v rasplavach zeleza, Izv VUZ Cernaja Metallurgija, 1996, Nr 5, s. 1-7. [7] Hutny, A.: Aktywność azotu w ciekłym stopie Fe-V w warunkach hiperbarycznych, Hutnik-Wiadomości Hutnicze, 2008, Nr 7, s. 349-355. [8] Wada H., Pehlke R. D.: Nitrogen Solubility in Liquid Fe-V and Fe-Cr-Ni-V Alloys, Metall. Trans. B, 1981, vol. 12B, pp. 333-339 [9] Pomarin Ju. M., Grigorenko G. M., Lakomskij V. I.: Rastvorimost azota v splavach żelaza s vanadiem i niobiem, Izv. A. N. SSSR Metally, 1975, Nr 5, 75-79. [10] El Tayeb N. M., Parlee N. A. D.: The Solubility of Nitrogen and the Precipitation of Vanadium Nitride in Liquid Iron-Vanadium Alloys, AIME Trans., 1963, vol. 227, pp. 929. [11] Evans D.B, Pehlke R.D.:Equilibria of Nitrogen with the Refractory Metals Titanium, Zirconium, Columbium, Vanadium and Tantalum in Liquid Iron, Trans. Metall. Society AIME, vol. 223, august 1965, 1620-1624. [12] Brick M. R., Creevy J. A.: Die Löslichkeit von Stickstoff in flüssigen Eisen-Chrom umd Eisen-VanadinLegierungen, Stahl und Eisen, 1940, Nr 49, pp. 1114-1115. [13] Kashyap V, Parlee N: Solubility of Nitrogen in Liquid Iron And Iron Alloys, Trans. Metall. Society of AIME, 1958, February, 86-91. [14] Pehlke R. D., Elliot J. F.: Solubility of Nitrogen in Liquid Iron Alloys. Thermodynamics, AIME Trans., 1960, vol. 218, pp. 1088-1101 [15] Rao M. R., Parlee N.: La solubilité de l’azote dans les alliages liquides fer-vanadium et fer-titane et l’équilibre dans la réaction x Ti + N = Tix N (d), Memoires Scientif. Rev. Metall., LVIII, Nr 1, 1961, pp. 52. [16] Korolev L. G., Morozov A. H.: Rastvorimost azota v zidkich splavach żeleza s vanadiem, Izv. VUZ. Cernaja Metallurgija, 1962, Nr 7, 27-29. [17] Morita Z., Tanaka T., Yanai T.; Equilibria of Nitride Forming Reactions in Liquid Iron Alloys, 1987, vol. 18B, pp.195. [18] Schmid R., Lin Jen-Chwen, Chang A.Y. The Activity Coefficient of Nonmetalilic Elements in Binary Liquid Alloys from a Nonrandom Solvation Shell Model. Z. Metallkunde, vol. 75, pp.730-736, 1984. [19] Jen-Chwen L., Schmid R., Chang A.Y. Comparision of Solution Models for Nonmetalilic Solutes in Binary Liquid Alloys: Nitrogen in Fe-Cr and Fe-Ni, Metallurgical Trans B, vol. 17B, pp.785-789, 1986. [20] Wada, H. Solubility of Nitrogen in Molten Fe-V Alloy, Transactions ISIJ, vol. 9, pp. 399-403, 1969. The work has been carried in the framework of Grant of Polish Government No. N N508 391935
Alloying of Steels with Nitrogen from a Gas Phase during VOD
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Alloying of Steels with Nitrogen from a Gas Phase during VOD S.A. Ivlev1, P.R. Scheller2, A.G. Svyazhin1 1. Technological University “Moscow Institute of Steel and Alloys”, 119049 Moscow, Russia; 2. Freiberg University of Mining and Technology, D-09596 Freiberg, Germany ABSTRACT It has been investigated of nitrogen absorption kinetics Cr-Ni-Mo steels by blow in a ladle with weight 30 t. The mathematical model of process of steel saturation is offered became by nitrogen from a gas phase during VOD operation. KEY WORDS nitrogen, kinetics, alloying, high nitrogen steels
1. Introduction
The nitrogen content in metallic materials is determined by it´s partial pressure during alloying and solidification and is independent on the kind of alloying procedure. The adsorption of nitrogen through the free surface in the furnace or ladle in technical processes is very low and can be neglected. In contrast to this the dissolution rate when blowing the nitrogen directly into the melt using the lance or porous plugs in the vessel bottom is much higher as fresh interface between liquid metal and nitrogen gas is generated. Such alloying technique is widely used in the AOD process or ladle treatment e.g. in the VOD process. The nitrogen alloying from gas phase at high pressure is technically difficult and shows poor efficiency [1]. 2. Experimental
In previous papers [2, 3] the kinetics of nitrogen alloying in Fe-Cr-Ni-Mo 30 t VOD steel melts was investigated and process model developed. The nitrogen was blown into the deoxidised melts containing less then 50 mass % of total oxygen through the porous plugs in the ladle bottom. The chemical composition of the melts varied within the following range (mass %): Cr 17-24, Ni 5-25, Mo 0-6 with the corresponding activity coefficient of nitrogen between 0.13 and 0.30. The further process parameter were as follows: nitrogen flow rate 28-65 m3/h s.t.p. (Nm3/h); steel temperature at the start of nitrogen injection 1500 – 1660°C; nitrogen content at blowing start 0.010 – 0.029 mass % and the final content of 0.05 – 0.17 mass %. 3. Mathematical model
The evaluation of experimental results and the activation energy [2, 3] show that the investigated reaction was of 1st order. Starting from the differential equation d [N ] S = ⋅ β ⋅ f N ([N ]S − [N ]) (1) dτ V S and the definition of the reaction rate constant k = β ⋅ ⋅ f N , V S V where - gas-metal surface, - metal volume, β - mass transfer coefficient in the melt in the plume, τ - time, [ N ]S , [ N ] ] – equilibrium and actual nitrogen concentration, % mass and fN activity coefficient of nitrogen the following regression equation for the reaction rate constant was derived [2,3] :
k = (1,695 ⋅ V&N 2 − 2,262) ⋅ f N ⋅ 10 −3 min −1
(2)
- 42 -
Thervjdenamics & Kinetics
Using this dependency the actual nitrogen content or the saturation of the melt when alloying the nitrogen by gas injection can be calculated: N = 0,691( k ⋅ t ) 0,630 , NS
(3)
Where N , N S -the actual and the equilibrium concentration of nitrogen in metal, % mass.; V&N 2 nitrogen blowing rate, Nm3/h. The equations (2) and (3) are valid for ladle treatment as the geometry of the reactor and the blowing conditions influence the value of S/V and therefore the regression coefficients. In the present work an other method of the mathematical description of the kinetics of nitrogen absorption in liquid steel is used [4,5]. The structure of the plume when injecting gas in liquids is reasonably well studied at room temperature. In liquid steel in the industrial scale there are very few information until now [6]. At high blow rates of gases which are typical in the ladle processing of steels the structure of the gasliquid plume is determined by the coalescence and dispersion of gas bubbles depending on turbulence grade. In a small distance above the tuyere or porous plug the average bubble size is determined by the interfacial tension and the shear stresses generated by the local turbulent flow and increases while ascending in the melt. With increasing of the gas flow the size of the bubbles change slightly. The concept of the bubble size has a statistical sense as they change continuously during their flow in the plume. Regarding the described process the main resistance in the mass transfer lies in the diffusion boundary layer at the melt side. Therefore the variation of fluid dynamics of a liquid in a vicinity of the bubbles rather slightly influences the total mass transfer. Taking these into account the gas-metal interface can be assumed as nearly constant and can be calculated as a spherical segment with equivalent radius req for the given flow conditions [4, 7]. The total interfacial area between all gas bubbles and the metal changes with blowing time due to continuously increasing saturation of the melt and therefore decreasing bubble consumption during its rise at given melt height [3]. Assuming the main mass transfer resistance is placed in the diffusion boundary layer on the metal side the equation (1) can be used. After the integration one obtain: [N ]0 − [N ]S = S ⋅ β ⋅ f ⋅ τ ln (4) N [N ] − [N ]S V where [N ]0 -initial concentration of nitrogen in the metal; % mass. The surface of a individual bubble can be defined as a surface of a spherical segment with the radius req of a bubble with equivalent volume: si = 18.033 ⋅ req2 .
(5)
The average number of bubbles assuming the normal ambient pressure, liquid metal temperature and average metallostatic pressure as well as the dynamic behaviour of bubbles in the plume is: V&N 2 ⋅ 10 6 H T n= , (6) ⋅ ⋅ 4 H ⎞ ⎛ 3 ξ ⋅ g ⋅ req 3600 ⋅ ⋅ π ⋅ req 273 ⋅ ⎜1 + ⎟ 3 ⎝ 290 ⎠ where V&N 2 - gas flow rate, Nm3/h; H - distance from the injection point to the melt surface, cm; ξ the hydrodynamic factor in value close to 1. Using the equation (6) the total area of bubbles is S=si·n and specific area of an interface between gas and melt is defined by the equation:
Alloying of Steels with Nitrogen from a Gas Phase during VOD
- 43 -
S T ⋅Q = 15,3 ⋅ 10 − 4 ⋅ 3 1 ⎞ V ⎛1 (7) ⎜ + ⎟ ⋅ ξ ⋅ req2 , 290 H ⎝ ⎠ 3 Q where - specific gas consumption, Nm /t.; the coefficient 15,3·10-4 unites numerical factors from equations (5) and (6). The equilibrium concentration of nitrogen changes with the distance between the injection point of gas and the bath surface due to pressure change according to: KN H [ N ]S = ⋅( 1+ + 1). (8) 2⋅ fN 145
where H is the distance between the bubble and the surface of the bath and the value of K N , f N are taken from [7]. Introducing the equation (7) and (8) in the equation (4) in the beginning without taking into account activity coefficient fN and grouping the a priori unknown values, we develop a parameter describing the mass transfer taking into account the geometrical factors. This mass transfer
β
parameter is defined as
3 ξ ⋅ req2 . Thus the equation (4) become convenient for further practical use
[4,5]: lg Y =
β0 3
ξ ⋅ req
⋅Х ,
(9)
2
⎞ ⎛ ⎛ ⎞ ⎜ K N ⋅ ⎜1 + 1 + H ⎟ − 2 ⋅ f N ⋅ [N ]0 ⎟ ⎜ ⎟ ⎜ 145 ⎟⎠ ⎝ ⎠ Y=⎝ , ⎞ ⎛ ⎛ ⎞ ⎜ K N ⋅ ⎜1 + 1 + H ⎟ − 2 ⋅ f N ⋅ [N ]⎟ ⎜ ⎟ ⎜ 145 ⎟⎠ ⎝ ⎠ ⎝ 0 , 02 T Q ⋅ X = 1,53 ⋅10−3 ⋅ ⋅ ([O] + 0,5 ⋅ [S ]) ⎛⎜ 1 + 1 ⎞⎟ , ⎝ 290 H ⎠
(10)
(11)
As the mass transfer rate is affected by the surface active elements as [O] and [S] it will be corrected by the expression ([O] + 0,5 [S]) which is introduced into the value of X. Where
β0
3 ξ ⋅ req2 -
mass transfer parameter of nitrogen in the metal, led some content of surface-active elements in melt, chosen as standard ([O] + 0,5 [S]) 0=0,02 mass %; [O] and [S] - actual concentration of oxygen and sulfur in the metal, mass %. The previously published industrial trials data [2,3] were evaluated using the equations (9-11). For different gas flow rates indicated by different points the values of X and Y are plotted in the figures 1 and 2 for different alloys. The X-Y values for different flow rates are dissipated over the whole range for both steels and fit each to one curve. This plots show that the mass transfer parameter
β0
3 ξ ⋅ req2 is independent of gas
flow rate within the investigated range. It can be concluded that the flow conditions near the gas bubbles do not markedly change with gas flow rate. The effect of temperature via the diffusion coefficient could not be evaluated due to small temperature range at which the trials were
Thervjdenamics & Kinetics
- 44 -
performed. In the previous paper [4] no any influence of the chemical composition of steel (activity coefficient of nitrogen fN) on the mass transfer parameter were observed. At data processing [2, 3] on the equation (9) content of steel (activity coefficient fN) noticeably influences factor mass transfer, fig. 3.
0,16 lgY = 9,8E-05*X R2 = 0,8752
lgY
0,12 0,08 0,04
- 35;
- 45;
- 55.
0 0
500
1000
1500
X
Fig. 1 lgY as a function of X for the steel Cr22Ni5Mo3Mn at different V&N 2 , Nm3/h
0.3 lgY = 0.00015*X
0.25
2
R = 0.7051
lgY
0.2 0.15 0.1 0.05
- 30; - 45;
0 0
500
- 33; - 48; 1000
- 36; - 51;
- 39; - 54; 1500
- 42; -63 2000
X Fig. 2 lgY as a function of X for the steel Cr18Ni13Mo4Mn at different V&N 2 , Nm3/h
Alloying of Steels with Nitrogen from a Gas Phase during VOD
- 45 -
In the equation (9) there is no of nitrogen activity coefficient fN. This equation is confirmed by experiments on various steels [4]. Figs 3 shows, that for Cr-Ni-Mo steels is necessary to consider also of activity coefficient fN and to present the equation (9) in a following type: lg Y =
β0 3
ξ ⋅ req2
fN ⋅Х 0.2
⋅
(12) For further evaluation the trials data [2,3] with the value of fN=0,2 were used. Also in the plot of the mass transfer parameter versus gas flow rate, fig. 4, no dependency of the mass transfer parameter on the gas flow rate can be stated. The data processing in this case completely correspond to the equation (9) in the form (12).
0.0003 β/ξ r
0.00025
3/2
-4
= 7,8.10 fN
2
R = 0,8131
0.0002 0.00015 0.0001 0.00005 0 0
0.05
0.1
0.15
0.2
0.25
0.3
fN Fig. 3 Dependence of mass transfer coefficient from nitrogen activity coefficient
0.0004 y = 9E-06x + 0.0001 2
R = 0.0116
0.0003
0.0002
0.0001
0 0
0.5
1
1.5
2
2.5
3 -1
QΣ , m t
Fig. 4 Relation between mass transfer coefficient and gas consumption Nm3/t steel by fN=0,2
Thervjdenamics & Kinetics
- 46 The is
β0
average 3
ξ ⋅ req2
value
of
the
mass
transfer
parameter
derived
from
this
evaluation
= (1,57 ± 0,38) ⋅ 10 − 4 cm-1/2·s-1 . Within the deviation range this value correspond very good
with the value of 1,2·10-4 predicted by the equation (13) according to [4] for injection of nitrogen through a porous element in the bottom of the ladle:
β0
3
ξ ⋅ req
= 6 ⋅ 10 − 4 ⋅ G
−1
3
cm-1/2·s-1 ,
(13)
2
where G-weight of metal in a ladle, t. 4. Conclusion Using previously developed relationships [4] for nitrogen absorption in steels the industrial trials data were evaluated in other procedure. The trials were performed in 30 t VOD ladles where nitrogen gas was injected into the high alloyed steels. The results show that the new defined mass transfer parameter is independent on the gas flow rate. The relationships presented can be applied for a wide range of steel melt weights. References 1.Svyazhin A.G., Kaputkina L.M. Steels, alloyed with Nitrogen. VUZ. Chernaya Metallurgia, 2005(10): 36-41. (Steel in Translation. Allerton Press. N.Y., 2006, Vol. : 64-76). 2.Steinmetz E., Scheller P.R. Fundamentals of nitrogen transfer in gas-metal bath systems during nitrogen injection. Steel research, 1987, 58(7): 303-309. 3.Steinmetz E., Scheller P.R. Kinetics of mass transfer during nitrogen injection in industrial Fe-Cr-Ni-Mo-melts. Steel research, 1987, 58(7): 310-318. 4.Svyazhin A.G., Khalek Shakhin M.A., Shevchenko A.D. Mass exchange by injection of nitrogen in liquid steel. VUZ. Chernaya Metallurgia, 1984(6): 37-42. (Steel in USSR, 1984(6): 274-278). 5.Svyazhin A.G. The kinetics of nitrogen absorption when nitrogen is injected into molten steel. High nitrogen steels. HNS 90. Conference Proceed.: Stahleisen Verlag, 1990, 117-121. 6.Steinmetz E., Scheller P.R. Beitrag zu den Strömungsverhältnissen in einer Spühlsteinpfanne. Stahl u. Eisen, 1987, 107(9): 417-425 7.Svyazhin A.G. The theory of steel alloying with gaseous nitrogen. Proceedings of Intern. Conf. on High Nitrogen Steels 2006.Beijing: Metallurgical Industry Press, 2006, 353-359.
For the question on interaction of active nitrogen with iron in the plasma heating process
- 47 -
For the question on interaction of active nitrogen with iron in the plasma heating process L.M.Simonyan Moscow Steel and Alloys Institute, Moscow, 119049, Russia ABSTRACT One of the factors influencing process of dissolution of nitrogen at plasma smelt is the characteristic anode areas (temperature and structure), which temperature exceeds temperature on a metal surface . KEY WORDS high nitrogen metal, atomic nitrogen, plasma swimming
At plasma smelt the high contents of nitrogen in metal is observed. It depends on parameters of a gas phase in a border zone of the discharge, a condition of a melt surface and structure of metal At the same time, on the basis of available experimental data about solubility of gases it is possible to draw a conclusion that as a result of their activation in a zone of an arc there is an acceleration of course of all physical and chemical processes. In particular, reaction of dissolution of nitrogen in liquid iron and steels proceeds with speeds on the order big, than in other melting units. It is known that the dissolution of molecular nitrogen in ferrous fusion leads to a reaction 1/2N2→[N] and in equilibrium conditions is governed by the Sieverts law
[N ] =
k N 2 p N 2 , where
kN2 is the dissolution constant of nitrogen dependent on temperature, and PN2 – the partial pressure of molecular nitrogen above the bath. For the temperature interval between 1809 and 1873 K the magnitude of kN2 can be established by means jf an empirical formula /1/, [N] 560 lg k N2 = lg =- 1,06 (1) T PN 2 lg [N] = -
560 1 - 1,06 + lgPN2 T 2
(2)
at 1973 1973 К in a equilibrium conditions it is possible to use both the equation (3) and the equation (8). In these conditions, plasma swimming trunks thermal balance between metal and a gas phase is absent. Because the temperature of gas is higher than temperature of metal, gas possesses some superfluous energy. This superfluous energy is distributed between components of a gas phase, particularly between particles of nitrogen (atoms, molecules, ions). Therefore, gas will be enriched with atoms of nitrogen, in which solubility in liquid iron is essentially higher than in molecules (see the equations 3 and 8). During the interaction of such gas with metal, dissolution of nitrogen will proceed more easily, as it will not be necessary to spend energy on the breaking of connections in a molecule. In metal will be achieved higher, so-called “super equilibrium”, concentration of nitrogen [N]*, in comparison with equilibrium (for same values РN2 and temperature of metal). Therefore it is necessary for plasma heating to use not the equation (3), which does not take into account superfluous energy of molecular nitrogen, but the equation (8). It should be taken into account that the value of [N] depends on two temperatures: it depends less on the temperature of metal Тме and more on the temperature of gas Тг, whilst Тг> Тмe. In this case the equation (8) is expedient for writing down as lg [N]
T Me
=
24004,1 - 4,28 + lg(P N) Т г TMe ,
(9)7 where Тг> Тме. Knowing average values of temperature of metal and gas during plasma melting, it is possible to calculate the average mass-concentration of nitrogen in metal [N]cp and the contrary. We shall estimate the probable temperature of gas in the settlement derived from experimental data. At РN2=0,02 atm (2 kPa), concentration of nitrogen in metal makes [N]cp = 0,1 %. The temperature of a melting surface can be different values: from 1810 on periphery up to 2700 K in the center /4/. For example, Тме(ср.) = 2500 K. From the equation (9), we find the partial pressure of nitrogen according to the given value [N]cp: PN=4,7⋅10-7 atm (470 Pa)
and
lg PN = - 6,32
(10)
- 480016,5 + 66,8Tг 25102,6 =+ 2,64 2,3RTг Tг
(11)
On the other hand, derived from equation (5) at РN2=0,02 atm, we find
lg PN = 1/2 lg 0,02 +
Substituting the value lgPN= - 6,32 into the equation (11), we shall receive the temperature of gas as Тг(ср.)=2800 К, adequate to the given composition of plasma, or given PN. The difference between Тме and Тг is insignificant and makes only 300 K. Let's assume, that gas at a surface of the melted solution is diluted with pairs of metal (see reactions 4) and the partial pressure of nitrogen in boundary area of the discharge is on the order less set. Substituting in the equation (11) PN2= 0,002 atm (instead of 0,02 atm), we shall obtain the temperature of gas from the surface of the metal, not exceeding 3000 K, and overheating will make no more than 500 K. From here, it is possible to make a conclusion that even low overheating of a gas stream in comparison with metal creates preconditions for "super equilibrium" solubility of nitrogen in metal.
For the question on interaction of active nitrogen with iron in the plasma heating process
- 49 -
This is due to the fact that the dependence of partial pressure of atoms of nitrogen from temperature is exponential, and from РN2 is degree: − ΔG 0 1 1/2 PN = Р N 2 e RT (12) Let's estimate the superfluous energy of nitrogen in a border zone. From experimental data according to dependence: [N] * ]*=К*эксп.√PN2, (13) have obtained К*эксп.=0,71 %/атм1/2 /5/. At Тме=2500 K, the constant of balance is 0,06. In the equation (13), if if К*эксп Кр⋅Кизб , the following equation results: [N]*=Кр⋅Кизб. √PN2, where Кизб. = К*эксп/Кр=0,71/0,06=11,85.
is replaced by (14) (15)
The ration К*эксп/Кр usually designates ϒ and the name factor of nonequilibrium /6/. It specifies as far as the chemical potential of nitrogen in a gas phase exceeds the chemical potential adequate to an equilibrium condition: μi=Gi= μi 0+RTlnPi+RTlnγi
(16)
The factor nonequilibrium ϒ reflects (indicated) cumulative influence on Gi (и μi) as a change of internal energy of particles of nitrogen, and as a change of entropy which occurs when their partial pressure changes in the considered system /7/. For the equilibrium condition μi=Gi= μi 0+RTlnPi (17) dependence μi from Pi is entirely caused by corresponding dependence on an entropy from pressure since internal energy does not depend on pressure. For nitrogen, the chemical potential in a solution is equal to chemical potential of particles of nitrogen in a gas phase. For molecular nitrogen, taking into account superfluous energy, it can be said that μ[N]*= 1/2μN2=1/2 μN2 0+1/2RTlnPN2+RTlnγ[N]* (18) In an equilibrium condition μ[N]0 = 1/2μN20+1/2RTlnP N2
(19)
If the second equation is subtracted from the first, the following results: μ[N]*- μ[N]0 = RTlnγ[N]*
(20)
The received expression represents superfluous Gibbs energy for the molecular nitrogen interacting with metal which can be caused by a change of internal energy, change of entropy, or both. Similar expressions can be received for other forms of nitrogen - atoms and ions. The value γ[N]* shows how many times concentration of active particles of nitrogen at plasma (that is nonequilibrium) heating exceeds the concentration of active particles at nonarc heating. For above mentioned conditions of experiment, superfluous energy of nitrogen at Тме=2500 К is equal μ[N]*- μ[N]0 = RTlgγ[N]* = 8,314⋅2500⋅ln11,85 = 51387 J/mol (21)
Thervjdenamics & Kinetics
- 50 -
It should be noted that, in some works, superfluous Gibbs’s energy is attributed to only separate particles. In work /4/ the superfluous energy equal to 6,82 kcal/g-at=28,5 kJ/mol is received. Authors assume that activation is caused by dissociation nitrogen and receive a constant of reaction dissociation КD=e-28500/RT=0,15 (при Тме=1800 0С), which is achieved at a temperature of gas 5100 0 С. However; as it has been shown above, calculation on the equation (11) on the basis of the relation of experimental and equilibrium solubilities [N]*/[N] gives temperature of activation of nitrogen 2800 K. The detailed analysis shows that in the scheme ln[N]*/[N]⇒ lnγизб⇒lnK*эксп/Кр⇒ -ΔGизб/RT⇒ lnKD=ln(PN/PN2)1/2 -ΔGизб/RT≠lnKD is not taken into account, and it is equal ln(K*D/KD) that is superfluous Gibbs energy of reaction of dissociation (7). Because the pressure in the system does not change, and only the temperature changes, there is a redistribution of energy between particles of gas. For atomic nitrogen the change ΔGизб basically is contributed due to increase in value of PN and speed of progressive movement as electronic excitation of atoms in considered processes can be neglected. It means, that for our experiment РN(изб)/РN=11,85. For molecular nitrogen (РN2(изб)/РN2)1/2=11,85, hence РN2(изб)/РN2=140. However, contrary to atoms, molecules have a more complex structure with rotary, oscillatory and electronic excited conditions. Therefore superfluous Gibbs’s energy is distributed also between the internal excited conditions of a molecule. Due to РN2≈РN2(исх), in can be said hat there is a surplus of molecules in excited conditions. Then it is more correct to write РN2*(изб)/РN2* =140, that specifies that the share of the molecules, capable to be dissolved in iron at plasma smelt, is more than two orders higher than at usual heating. If superfluous Gibbs energy value is known, it is possible to use any of the above mentioned equations (3) and (8) for calculations, having added in the right part lgγ[N]*. Then calculations can be made at one temperature, specifically at the temperature of the metal: lg [N] = lg [N] =
1100 - 0,79 + 0,5 lgPN2 + lgγ[N]* T
24004,1 - 4,28 + lg PN + lgγ[N]* Т
(22) (23)
The equation (22) takes into account the excited conditions of molecules of nitrogen. Thus from the thermodynamic point of view, it doesn’t matter what state the activated particle is in the gas phase. The only thing that matters is that the activated particle is the carrier of the superfluous energy transmitted during dissolution of the smelt, which causes an increase in the solubility of nitrogen in metal.
References 1. Grigorjan V.A., Beljanchikov L.N., Stomahin A.J. Theoretical of a basis of electrosteel-smelting processes. - M: Metallurgy, 1987, 272 c. 2. Lakomskyi V.I. Interaction diatomic gases with liquid metals at high temperatures. Kiev: Наук.думка, 1992.-232 p. 3. Dembovsky V. Plasma metallurgy. Amst.-Oxford-N.Y.-Tokio, 1985, 476 p. 4. Simonjan L.M. Distributing of temperature on the thickness of boundary layer at the vakuum-plazma-heating of metal fusions. Physics and chemistry of treatment of materials. M., 2000, № 3, p. 91-92. 5. Simonyan L.M. The new concept of a boundary layer at plasma heating metal. Isvestija vusov, ChМ, 1999, №9, p. 78. 6. Yerokhin A.A. Plasma-arc smelt of metals and alloys.-М., Nauka, 1975. 7. Eremin E.N. Bases of chemical thermodynamic. M., "Vyssh..shkola", 1974, 341 p.
CLASSIC DESCRIPTION OF THE NITROGEN AND HYDROGEN SOLUBILITY IN SOLID IRON
- 51 -
CLASSIC DESCRIPTION OF THE NITROGEN AND HYDROGEN SOLUBILITY IN SOLID IRON. Yu.S.Venets CJSC «CENTRAVIS PRODUCTION UKRAINE», Nikopol, Dnepropetrovsk region, Ukraine, 53201. SUMMARY: There are adduced classical functions of the nitrogen and hydrogen solubility in ferrite and austenite caused by electronic configurations of iron. For its calculation the empiric thermodynamic model, reflecting through the whole solid metal temperature range was suggested, which enables the determination of maximum iron alloying level. The model was built according to the Hume-Rothery's principle on the grounds of common functional thermodynamic descriptions of iron electronic configurations, being uniform for austenite and ferrite, as well as dependence of size factor on its configurations. Theoretically the ideality of the electronic gas in the solid iron as validity for it the Mendeleev-Clapeyron classical equation and distribution of all orbital electrons as Fermi-Dirac quantum-mechanical distribution, as well as derivation of its type from Maxwell-Boltzmann classical distribution for part of them are provided. Analysis of coefficients in the proposed ferrite and austenite solubility model enables the evaluation of difference of the nitrogen introduction structural nature. KEY WORDS: Nitrogen, solubility, ferrite, austenite, electronic configuration. “Everything moves, everything passes, and there is no end, Where did it all disappear? From where did it all come? Both the fool and the wise man know nothing.” 07.04.1841, St. Petersburg – T. Shevchenko
1. INTRODUCTION. The gaseous nitrogen penetration into liquid and solid steels for the practical purposes at
metallurgical processes temperatures was described only to the atomic level [1, 2]. In this case three phases of different substances at the same temperature (solid, liquid and gaseous) in this field can be explained only on the grounds of physical sense of number values of electronic configuration parameters within the Mendeleev's table. At the same time six possible conversions from one phase into the other one for every substance at the temperature changing (melting-crystallization, boilcondensation and sublimation-desublimation) are described by the thermodynamic also only to the atomic level. That is why owing to the found collision and on the grounds of available experimental data in this work we propose the thermodynamic description of the nitrogen solubility in steels at the electronic structure level. 2. ANALYSIS OF REGULARITY FOR CLASSICAL AND ELECTRONIC GAS According to the previously published classical regularity on the intratomic level, there was
suggested to understand the t2g–quantity of electrons as the dependence (1) under the parameter a(T) there. ln(a (T )) =
A + B; t + 273
(1).
where for iron А = – 1909; В = 2,528. The value a(Т) for iron under normal conditions almost coincides with the value of half relation of one collective electron to their aggregate number of 26, that testifies to the necessity of an additional analysis of its value.
Thervjdenamics & Kinetics
- 52 -
It was considered, that at the temperatures when iron was in solid condition, the electronic gas in iron was degenerated, and i.e. it didn’t obey the regularities for the ideal gas. But, as in iron the presence of the classical thermodynamic dependence for part of d-electrons was revealed, it is obvious to check the validity for it of the perfect gas condition equation – (the MendeleevClapeyron equation) with regard for the proposed quantum idea about redistribution of electrons in the solid iron by the temperature changing. By general constancy of number of electrons with the temperature elevation, let’s suppose that the volume will be changed owing to the allocation difference within the volume of electrons on eg–and t2g–suborbitals. Let’s take the linear suborbital size t2g as the iron radius and designate eg with the coefficient r(Р). Ten the d-orbital size will be equal to: RTr =
j a (T ) [amax − a(T )] ⋅r + j j amax amax
(2)
j where amax – quantity of electrons on the d-orbital of the j-element;
The following equation (3) represents the function of correlations between linear size growth as the cube root from the volume according to the Mendeleev-Clapeyron equation under condition of uniform pressure and substance quantity related to the growth of the d-orbital linear size according to the equation (2): r T ;Tc
L
(T =
1
Tc )3 r RT RTcr
(3)
where T and Tc – current temperature and temperature of comparison. Proceeding from the known distortion of the electronic suborbital eg arrangement in the disintegration (109,5o angle instead of 90o), supposed being caused by the p6-shell influence and under condition of permanent dimension of the electronic orbital eg projection in the plane (1;1;0) towards [11] over the p6-shell, its size in the distorted volumetric direction was found as: r ( P) =
(4).
0,6 ⎛ 109,5 − 90 ⎞ Cos ⎜ ⎟ 2 ⎠ ⎝
2⋅P
where Р – pressure, atm; 0,6 — size of the electronic orbital eg projection in the plane (1;1;0) towards [11] in atom size fractions. The degree
2 in (4) transfers the probability of suborbital eg enlargement from the plane
dimensionality of the pair of orbitals into the volume dimensionality for the whole atom, increasing the probability of changing of suborbital eg linear dimension by 2 times in the perpendicular section plane for two planes of suborbital eg proportionally to the linear dimension increasing. Under the atmospheric pressure r(Р) is equal to 0,6125, being almost practically equal to
CLASSIC DESCRIPTION OF THE NITROGEN AND HYDROGEN SOLUBILITY IN SOLID IRON
- 53 -
3 (0,6124). 2 2 Practical coincidence r(1) with the relation of intervals along crystallographic directions [0,5;0,5;0,5] and [1;1;0] in the elemental cubic cell under the atmospheric pressure apparently shows the directivity of links of suborbitals eg and t2g for these directions correspondingly. If the supposed physical sense of the relationship (4) is true, we can theoretically define more exactly the angle of the electronic orbital eg arrangement distortion in the body-centered cubic lattice, which for full coincidence with the relation
3 2 2
should be equal to 109,42° instead of 109,5°. In this case we can
affirm that the electronic orbital eg arrangement distortion is caused not by the p6-shell influence, but pre-defined by the body-centered cubic lattice origin in iron. The calculation error for the atom linear dimension using the d-orbital’s dimension in comparison with the calculation using the Mendeleev-Clapeyron equation is appr. 0,3% (Fig.1) in dependence on the temperature within the range of γ–iron existence at r(1) in the equation (3) at the comparison temperature of 910°С and atmospheric pressure within the range of γ–iron existence is about 0.3% (Fig.1) as well as allows to take the value r(1) in the body-centered cubic lattice at the atmospheric 0.6
pressure equal to 0,6125, and the same value in the
0 .45
single atom of iron equal to 0,6. The considered
0.3
comparability of the electronic gas volume with the
0 .15 0 600
80 0
1000 1 200 1400 1 600
t for (3)average linear Fig. 1. The calculating error dimension of the iron atoms electronic shell using the d-orbital dimension, calculated under (2), in comparison with the calculation using the Mendeleev-Clapeyron equation regarding temperature 910°С at the continuous pressure and r=0,6125, %.
calculation, carried out using the MendeleevClapeyron equation can be interpreted as its base caused by principles of the atomic structure. Amount of all t2g & eg–electrons with regard for the relationship of their dimensions
3 2 2
and after
substitution and elementary transformations at the atmospheric pressure has an appearance of the Fermi-Dirac distribution: a(T ) + [6 − a(T )] ⋅
3 3 ⎛ − 1909 ⎞ = exp⎜ + 1,581⎟ + 6 ⋅ 2 2 2 2 ⎝ t + 273 ⎠
(1')
3. MODEL OF GASES SOLUBILITY IN IRON Let’s suppose the introduction of nitrogen and hydrogen atoms into the lattice as practicable owing
to distortion of the suborbital eg. Then, firstly, we will take the increment of fracture of all electronic orbitals at the temperature T in comparison with the temperature Tc, taken as 0,5°К, as proportional to the cube of the increment of linear dimensions amount of eg– and t2g–suborbitals:
Thervjdenamics & Kinetics
- 54 -
ΔVel.orb.( j ) Tr ( Р )
⎛ Rr =⎜ rT ⎜R ⎝ 0,5° K
⎞ ⎟ ⎟ ⎠
3
(5)
Second, let’s suppose the volume of pores, into which the nitrogen introduces, as proportional to the change of electronic orbitals volume with the distortion of eg-orbital over the p6-shell in comparison with the volume without distortion, taking into account changes from 0.5°К: 0,6 r
V PIN , ( j ) T
ΔVel .orb ( j ).T0 , 6 = −1 ΔVel .orb.( j ) Tr
(6)
But the pores volume dimensionality is evaluated in atomic fractures of iron. The solubility of nitrogen and hydrogen (% mass), determined here as the ratio of quantities of gases to all iron atoms in α,β,δ–Fe and γ–Fe is shown in the Fig. 2 and found as equations: 0,6 r
[ N ]α ,β ,δ − Fe = V PIN ( Fe ),T ⋅ a (T ) ⋅ 0,6 r
100 14 ⋅ 55,8 ⋅ 10 55,8
(7)
2 2 100 14 ⋅ ⋅3⋅ 55,8 ⋅10 55,8 3 100 1 ⋅ a (T ) ⋅ ⋅ 55,8 ⋅ 10 55,8
(8)
[ N ]γ − Fe = V PIN ( Fe ),T ⋅ [6 − a(T )] ⋅ 0,6 r
[ H ]α ,β ,δ − Fe = V PIN ( Fe ),T 0,6 r
[ H ]γ − Fe = V PIN ( Fe ),T ⋅ a (T ) ⋅
(9)
100 1 2 2 ⋅ ⋅1⋅ 55,8 ⋅ 10 55,8 3
(10)
where 1, 14, 55,8 – are atomic weights of hydrogen, nitrogen and iron, g; 100 – coefficient for fractions conversion into per cents; 10 – quantity of fillets on the d–orbital; 1 and 3 – quantity of electrons on external orbitals of hydrogen and nitrogen; Thereupon, as, according to the preliminary calculations to 625atm, and when r(Р) is equal to 1 and, correspondingly, the dimensions of d-suborbitals are equal and further changing is apparently limited by the dimensions of atoms, the solubility of gases in the solid iron solutions is specified for the atmospheric pressure, and in case of its variation up to the Fig. 2. Nitrogen and hydrogen solubility in solid iron according to equations (7) - (10) at the atmospheric pressure.
specified value is supposed as the true one according to the Siverts law. However, the actual variation of pressure oscillations in
the analyzed metallurgical processes, i.e. gas-oxygen refining (GOR, being the analogue to the Argon-Oxygen Decarburizing) and solid-phase decarburizing combined with the nitrogen alloying is lower. Besides iron dimensional factor and its- and gas electronic parameters, reflecting their electronic
CLASSIC DESCRIPTION OF THE NITROGEN AND HYDROGEN SOLUBILITY IN SOLID IRON
- 55 -
interactions, the hydrogen and nitrogen solubility proved to be inverse proportional to the sum of protons and neutrons of iron by 55.8 and quantity of its electronic fillers on the d–orbitals. Moreover, in the austenite it is proportional to the electrons quantity on the external orbital of gases, and its level proved to be higher by 2 2 times. It could be explained with changes of arrangement 3
of links of gases with iron from ferrite to austenite from the direction [0,5;0,5;0,5] by t2g-electrons in the body-centered cubic lattice iron to the direction [1,1,0] in the same coordinates or [1,0,0] in the coordinates of the face-centered cubic lattice iron by eg-electrons. As the quantity of electrons in the iron is invariable, the length reduction apparently brings to the increasing of electronic density in the nitrogen and hydrogen links in the iron austenite, reflected by coefficients
2 2 in equations 3
(8) and (10) in comparison with equations (7) and (9). The coincidence of the gases solubility in the solid iron by values of the parameter a(Т) lower than 2 and at temperatures lower than 768°С, i.e. in cases, when the value [6–a(Т)] for the iron exceeds the electrons capacity on the suborbital eg, shows the necessity of revision of the a(Т) parameter sense. Apparently, its value, previously being supposed as the quantity of electrons on the t2gsuborbital, should be considered as the part of all electrons on the d–orbital and called the «classical» one, as well as the resting ones [6–a(Т)] should be called as the «basic» ones. This fact is confirmed with existence of the earlier found thermodynamic regularity of the electronic configuration parameter for silicon, p-orbital of which is not divided onto suborbitals. 4. Deductions: • The part of the electronic gas in the solid iron obeys the classical rules, i.e. Mendeleev-
Clapeyron equation and Maxwell-Bolzman distribution for “classical” electrons; the derivative from which sum of “classical” and “base” electrons in the solid iron have the form of the FermiDirac distribution. •
The electronic eg-suborbital distortion apparently can be explained with the crystal structure formation in iron.
•
The nitrogen and hydrogen solubility in steel depends only on its electronic-configurational energetic interactions while its structural factors as crystallographic directivity and dimensional factor of the Hume-Rothery's principle are derivative from it, as well as heat capacity, electrical resistance, heat conduction, not described in this article. References:
1 2 3
V.V.Averin, A.V. Revyakin, V.I. Fedorchenko, L.N. Kozina. Nitrogen in Metals. (M: Metallurgy, 1976). P.V. Geld, R.A. Ryabov. Hydrogen in Metals and Alloys. (M: Metallurgy, 1974). Y.S. Venets. Correlation of Gas Solubility in Solid Iron and its Electron Configuration. (Metallurgy Theory and Practice. №1, 2, 2006).
- 56 4 5
Thervjdenamics & Kinetics
B.M. Yavorsky, A.A. Detlaf. Guide on Physics. (M: Science, 1980). V.K. Grygorovych. Metallic Linkage and Structure of Metals. (M: Science, 1988).
DEPENDENCE OF THE NITROGEN SOLUBILITY IN AUSTENITE AND FERRITE ON ALLOYING AT THE ELEMENTAL LEVEL
- 57 -
DEPENDENCE OF THE NITROGEN SOLUBILITY IN AUSTENITE AND FERRITE ON ALLOYING AT THE ELEMENTAL LEVEL. Yu.S.Venets CJSC «CENTRAVIS PRODUCTION UKRAINE», Nikopol, Dnepropetrovsk region, Ukraine, 53201. SUMMARY: There are adduced classical functions of the nitrogen solubility in ferrite and austenite caused by alloying at the intratomic level there. For its calculation the empiric thermodynamic model, reflecting the accumulated elemental influences by the example of iron, chrome, manganese, nickel through the whole solid metal temperature range was suggested, which enables the determination of maximum metal alloying level, ensuring the non-defect structure by crystallization. The model was built according to the Hume-Rothery's principle on the grounds of common functional thermodynamic descriptions of alloying elements and iron electronic configurations, being uniform for austenite and ferrite, as well as dependence of size factor for each component on these configurations. Analysis of coefficients in the proposed ferrite and austenite solubility model enables the evaluation of difference of the nitrogen introduction structural nature for these metals by alloying with considered elements. KEY WORDS: Nitrogen, solubility, ferrite, austenite, electronic configuration.
5. Introduction. The nitrogen solubility in the high-chromium steels, as also in iron, for practical purposes described
by thermodynamic methods by Taylor series expansion of coefficients of alloying elements influence. Given approach allows description of solubility with the sufficient accuracy within the concentration and temperature range being studied, but it is limited with application of determined rules out of limits of the being analyzed and rather narrow temperature interval, e.g. near the crystallization temperature. 6. The Model of nitrogen solubility in high-chromium ferrite and austenite According to the previously published classical regularity on the intratomic level, there was
suggested to understand the t2g–quantity of electrons dependence (1) under the parameter a(T) there. ln(a (T )) =
A + B; t + 273
(1).
where for iron А = – 1909; В = 2,528. Let’s take the linear suborbital size t2g as the atomics radius and designate eg with the coefficient r(Р). Ten the d-orbital size will be equal to: RTr =
j a (T ) [amax − a(T )] ⋅r + j j amax amax
(2)
j where amax – quantity of electrons on the d-orbital of the j-element; T and Tc – current
temperature and temperature of comparison 0,5ºК. Proceeding from the known distortion of the electronic suborbital eg arrangement in the disintegration (109,5o angle instead of 90o), supposed being caused by the p6-shell influence and under condition of permanent dimension of the electronic orbital eg projection in the plane (1;1;0) towards [11] over the p6-shell, its size in the distorted volumetric direction was found as:
Thervjdenamics & Kinetics
- 58 -
(3).
0,6
r ( P) =
⎛ 109,5 − 90 ⎞ Cos ⎜ ⎟ 2 ⎠ ⎝
2⋅P
where Р – pressure, atm; 0,6 — size of the electronic orbital eg projection in the plane (1;1;0) towards [11] in atom size fractions. Let’s suppose the introduction of nitrogen and hydrogen atoms into the lattice as practicable owing to distortion of the suborbital eg. Then, firstly, we will take the increment of fracture of all electronic orbitals at the temperature T in comparison with the temperature Tc, taken as 0,5°К, as proportional to the cube of the increment of linear dimensions amount of eg– and t2g–suborbitals:
ΔVel.orb.( j )
r (Р) T
⎛ RTr =⎜ r ⎜R ⎝ 0,5° K
⎞ ⎟ ⎟ ⎠
3
(4)
Second, let’s suppose the volume of pores, into which the nitrogen introduces, as proportional to the change of electronic orbitals volume with the distortion of eg-orbital over the p6-shell in comparison with the volume without distortion, taking into account changes from 0.5°К:
ΔVel .orb ( j ).T0 , 6 V PIN , ( j ) T = −1 ΔVel .orb.( j ) Tr But the pores volume dimensionality is evaluated in atomic fractures of iron. 0,6 r
(5)
The electronic configuration parameter for the manganese in the solid state а(Т) was found proceeding from the temperatures of its allotropic transformations, proportional to expressions of whole numbers (Tables 1 and 2). The correlating factor value and graphical mapping of the obtained equation (Fig. 1) testify to the functional but not the static character of dependence. the liquid state, its electronic configuration above its melting point was found in comparison with other elements: iron, silicon, manganese, nickel and aluminum, shown in the equation (1) (Table 2 and 3, Fig. 2).
β—γ
1087
7,35
γ—δ
1137
7,09
δ — liquid
1245
6,59
2
1 2
2 2 3 3
0,916 0,981 1,099
т
Mn(T)]
1,0 0,8
As at the steels crystallizing temperatures the pure manganese is in
Table 1. Correspondence of the manganese electronic configuration parameter in solid state to temperatures of allotropic changes s s 4 a Mn (T ) ] otropic change /Т, К-1 a Mn (T ) t, °С 1 1 0,288 727 10,00 α—β 3
ln[a 1,2
0,6 0,4 0,2 6
7
8
9
4
Fig. 1. Dependence of the electronic parameter configuration of manganese in solid state on temperature.
Table 2. Energetic parameters of the equation (1), showing the electronic configuration of parameters in solid, as well in liquid and gaseous states Element Phase А В R –1909 2,528 1,0000 S[3] Fe Si Mn Ni Cr Al
L, G S[3] L, G S L, G S L, G S L, G
-1
10/Т, К11 10
–1005 –912 –721,2 –2378,2 –3258 –341,6 –1648 –935,2 –724,9
1,941 1,871 1,526 2,6659 2,149 2,2787 2,341 1,9052 1,182
1,00000 1,0000 0,99999 1,00000 1,00000 — 1,00000 — 1,00000
DEPENDENCE OF THE NITROGEN SOLUBILITY IN AUSTENITE AND FERRITE ON ALLOYING AT THE ELEMENTAL LEVEL
Table 3. Correspondence of elements electronic configuration parameter to temperatures of their changes in liquid and gaseous state according to the equation (1). Element. Iron. Silicon. Manganese. Nickel. Aluminum. Typical point t, ˚С a(Т) t, ˚С a(Т) t, ˚С a(Т) t, ˚C a(Т) t, ˚C a(Т) Crystallizing. 1539 4,00 1414 3,00 1244 1,00 1453 4,00 660,5 1,50 Boiling 2750 5,00 2355 3,50 1962 2,00 2732 6,00 2467 2,50 T critical. 6477 6,00 4886 4,00 5777 5,01 6021 8,00 8377 3,00
* — temperatures are taken from [7.1].
- 59 -
ln[a(T)] 2
Mn Fe Ni Si Al
1,5 1 0,5 0 4
-1
1 2 3 4 5 6 7 8 10 9 /T, 10 К11
Fig.2 Dependence of parameter of electronic configurations of elements in liquid and gaseous state on the temperature.
The atomic electronic configuration parameter for liquid and gaseous state was supposed to be proportional to whole numbers at melting, boiling temperatures and critical temperature, above which the steams of the element can not be condensed by means of volume compression (Table 3). The correlating coefficient value testifies, as well as for solid iron, silicon and manganese, the dependence functionality. The obtained parameter for liquid and gaseous phases should be according to its nature classified (like the analogous parameter for the solid state) as the electronic configuration, as at the critical temperature for mentioned elements obtained parameter is equal to the total quantity of their external electrons. As qualitative confirmation of abnormally low value a(Т) for manganese at the temperature of steelmelting processes in comparison with iron, silicon, nickel and aluminum, we can see abnormal pressure of manganese steam. I. e. the relatively permanent energy influences the less quantity of manganese atoms and, correspondingly, the greater quantity of atoms turns into steam. Lower temperatures of melting and boiling, as well as steam heat of manganese were previously asserted among d-elements of the 4th period of Mendeleev’s table owing to assumption of relatively low concentration of collective electrons, being confirmed by relatively low a(Т) value for manganese in liquid state, approximating to the melting point. This fact also results the stepwise increasing of ferrite quantity in the austenitic high-chromium manganese metal owing to drop of its charge over the manganese melting temperature. As the quantitative confirmation of obtained electronic-configurational data for solid and liquid state we adduce the comparison of calculated on their base melting entalphies and entropies according to (3) and (4) with empiric atomic values (Table 4). ΔSmelt = −( BL − ( BS − ln X )) ⋅ R ⋅ 2
ΔHmelt = AL − AS ⋅ R ⋅ 2
(6) (7)
By calculation of melting entropies for silicon not taking into account in the melting energies forming the liquid state parameters can testify the absence of orbital interactions in the molten silicon; this fact is confirmed with the data regarding full destruction of directed links during silicon melting process. The electronic configuration parameters for nickel and chrome in the solid state are obtained by determining of solubility in the high-chromium austenitic metal.
Thervjdenamics & Kinetics
- 60 -
Table 4. Comparative values of melting entropies (S) and entalpies (H). ∆Sempir[О ∆Нempir[О шибка! шибка! ∆Ssolv, Источник Δ, ∆Нsolv, Источник Δ, Elem Вs ВL X* As AL ent ссылки не % ссылки не % найден.] найден.] Joule/mole·К. Joule/mole. Fe Вs ВL 4,36/4 8,340 8,383 0,5 As AL 15091 15190 0,7 1/2 Mn Вs ВL (2,67/3) 9,570 9,657 0,9 As AL 14626 14640 0,1 — — — — As ⋅ 3 AL 17610 0,3 Ni — — 17556 1/2 — — — — — Si Вs 0 (3,5/3) 29,82 30,05 0,8 * Х – is the expression of elements electronic configuration parameters for elements in the solid state in the nearest to the melting temperature allotropic change and its nearest stable (whole) value.
The dimension of j-element orbitals, volume increment of all electronic orbitals and pores volume is determined analogously to the iron using equations (2), (4) and (5) correspondingly with the insertion of corresponding parameter of j-element electronic configuration. We will describe the nitrogen solubility in ferrite and austenite of the high-chromium metal with the dependence analogous to the pure solid iron, taking into account the basic components, i. e. iron, chrome, manganese and nickel and deducing to their concentration the dimensional and electronic factors: 0,6 r
100 14 2 ⋅ 2 [Fe] ⋅ A 2 ⋅ ⋅ E Nfaz ⋅ ⋅ 10 ⋅ A ⋅ A 55,8 2 3 [ j ]2 ⋅ A2 14 faz faz ⋅ E j ⋅ 100 ⋅ ⋅ EN ⋅ x j ⋅ A A 2j 2
faz ⋅ [ N ] Fe faz = V PIN , ( Fe ) T ⋅ E J 0, 6 r
[ N ] jfaz = V PIN,( j ) T
(8) (9)
where j –mass concentration of the j–element; 55,8 and A — average atomic weigh for iron and steel, g; Aj – atomic mass of the j–element; E Nfaz , E jfaz – quantity of electrons for different phases (for nitrogen in the ferrite - 1 and in
the austenite – 3, for Сr in ferrite - ECrferrit = [amax .Fe − aCr (T )] , and for j-element in austenite E austenit = [amax j − a j (T )] ) j
xj –nitrogen solubility structural coefficient, caused by j–element. The nitrogen solubility in austenite and ferrite of high-chromium steel Na was found as sum: Cr Ni Mn [ N ] faz = [ N ] Fe faz + [ N ] faz + [ N ] faz + [ N ] faz
(10)
Experimental values of nitrogen solubility in the high-chromium steel and comparative solubility calculated values according to the Taylor series expansion of chrome, manganese and nickel influence up to the parameters of the second order and double parameters are taken from works (Fig. 3). For solubility determining we restrict ourselves to 11 austenitic high-chromium steels at temperatures 1200 и 1300°С with chrome content (16,2-18,0%), nickel content (up to 14,6%) and manganese content (up to 16,5%), as well as 4 compositions of ferritic high-chromium (16,8; 22,6;
DEPENDENCE OF THE NITROGEN SOLUBILITY IN AUSTENITE AND FERRITE ON ALLOYING AT THE ELEMENTAL LEVEL
- 61 -
24,1; 31,6%weigh) steels at temperatures 1250, 1300 and 1350°С. The search of solubility dependence was carried out by minimization of the square deviations sum of experimental and A, B for chrome and nickel obtained using the equation (1) and xj obtained using the equation (9), specified in tables 2 and 5
solve
calculated values according to the formula (10). The coefficients
correspondingly apparently demonstrate the vector sum of interactions of alloying elements masses with iron and nitrogen.
0 0,1 0,2 0,3 0,4 0,5 0,6 0,7 0,8 0,9 1 1,1
As the calculation using the formula (10) is based on physical
1,1 y = 1,003x + 0,002 1 R2 = 0,988 0,9 0,8 0,7 0,6 0,5 0,4 y = 0,949x + 0,01 0,3 0,2 R2 = 0,991 0,1 0 empiric
nature of the nitrogen-into-steel introduction process, we obviously should extrapolate this formula out of the temperature range 1200 up to 1350°С. The major importance has the calculation of solubility at the liquidus temperature for
Fig. 3 Convergence of nitrogen solubility in austenite and ferrite. Ο, «—» –
[ N ] max[Tailor ] (R=0,995);
×, «- - -» –
determining of maximum nitrogen alloying level of the liquid
[N ] max(10) (R=0,994)
points – reference data, lines – correlative right lines.
bath with nitrogen with the goal to prevent nitrogen porosity during crystallizing. The check of formulas
Эле-мент
complexes was carried out
Fe
during
Table 5. Equation coefficients (1) xj for austenite 2⋅ 2 3
horizontal
centrifugal casting of the
Cr
8,06 =
experimental metal, grade 03Х18Н5АГ4 into
(Table
the
Ø103х35х1000mm
6)
Mn
tube
Ni
3 ⋅ (55,85 − 52) 2
2,7 = 3 ⋅ (55,85 − 54,9) − 7,53 = 3 ⋅
under
elevated pressure, stipulated by production process. The casting was carried out into the water-cooled horizontal metal mold with coating of lunkerite 2mm
3 ⋅ (55,85 − 58,7) 2
xj for ferrite 1
3 (Cr σП
A 1 - σ< σП i2
i
Fig. 8. Polarization diagram of corrosion at SCC: ϕO2ABC – cathode curve; ϕFe/Fe2+x11 and ϕFe/Fe2+x22 – anode curves at a different stress.
b
a −
Zone axis [ 1 03 ]γ c
−
3 11 Fig. 6. Kinetics of electric resistance change in investigated steels Cr21N1 (1), Cr18Ni2N1 (2) и Cr18C1 (3) during the aging at: a – 300 °С; b – 400 °С; c – 500 °С.
b 200
000 110α, 020γ
220 311
310 − 002 Zone axis [ 001 ]α −
Zone axis [ 11 0 ]α
c 222 Fig. 9. Structure of deformed at an 8% reduction before aging at 350 °С steel Cr18Ni2N1: а – martensite ασ; b – selected-area electron diffraction pattern taken from region (a); c – key pattern.
The results of studies an effect of cold plastic deformation on the phase content and corrosion properties is presented in Table 2. CPD at an 8-24% reduction was combined with aging at 350 °C by two ways: before (lines 6-10 in Table 2) and after aging (lines 11-14) as compared to the undeformed state (lines 1-5). It follows from Tab. 2 that CPD at 8-12% reduction before aging at 350 °C decreases to thrice a corrosion rate, but the 20% deformation on the contrary to five time increases it because of martensite forming from depleted austenite after strain aging and cooling. Probably a small deformation causes a stress relaxation but a considerable one increases the level of strength and the corrosion current as a result (Fig. 7). After 24% deformation the austenite has more stability than after 20% deformation because of strain recovery, that’s confirmed by dependence of
Effect of Aging and Cold Plastic Deformation on Structure and Stress Corrosion Cracking in High Nitrogen Chromium Steels
- 101 -
micro hardness from deformation degree (Fig. 8). Fig. 9, b shows the electron diffraction pattern with less diffusion effects than on the same one after an aging without CPD (Fig. 3 e). A treatment with CPD at 20% reduction after aging at 350°C is found to be more effective: the corrosion rate decreased by an order and the reduction of strength at SCC was about a quarter of magnitudes of undeformed state [6]. This positive influence of CPD at an 8-12% reduction probably is connected with breaking the order in metastable clusters which could transform to stress-induced martensite at SCC. CONCLUSIONS 1. The structure of steels Cr21N1, Cr18Ni2N1 and Cr18C1 with high content of nitrogen and carbon, respectively 1.021, 0.899 and 0.920% consists of strain metastable austenite and nitrides/carbides settled down the grain boundaries. Also a small content of α-phase takes place near the nitrides in high nitrogen steels. 2. The highest hardness, corrosion rate and susceptibility to SCC were observed at the state after aging at 350 °C. The formation of ordered clusters enriched in chromium and nitrogen, coherently connected with matrix, precedes homogeneous decomposition of austenite with the formation of the CrN nitrides. These clusters have an ability to transform into a stress induced martensite which gives evidence of high strength level in the steels and has a negative influence on resistance to SCC. 3. Cold plastic deformation at 8-12% reduction preceding the aging at 350 °C is shown to decrease three times the corrosion rate of the steel Cr18Ni2N1 as compared to the undeformed state. A thermoplastic treatment by 20% CPD after aging at 350 °C is more effective. ACKNOWLEDGMENTS This work was supported by the Russian Foundation for Basic Research, project no. 07-0300062a. REFERENCES 1. Bannykh O. A., Blinov V. M., Kostina M. V. Nitrogen as an Alloying Element in Iron Based Alloys: Trudy shkolyseminara. Vyp. 3. Magnitogorsk: Auspher, 2003. 2. Shpaidel M. O. New Nitrogen Austenitic Stainless Steels with High Strength and Plasticity. Metalloved. Term. Obrab. Met., 2004, 11(605): 9-14. 3. Berezovskaya V. V., Kostina M. V., Blinov E. V., et. al. Corrosion Properties of Austenitic Cr-MnNi-N Steels with Various Manganese Concentration. Izv. Ross. Akad. Nauk, Ser. Met., 2008, 1: 36-41. [Russian Metallurgy (Metally), 2008, 1: 29-33]. 4. Kostina M. V., Bannykh O. A., Blinov V. M., Berezovskaya V. V., et. al. Effect of Chemical Composition and Heat Treatment on the Corrosion Properties in High Nitrigen Iron-Based Alloys Containing 15-24%Cr. Izv. Ross. Akad. Nauk, Ser. Met., 2001, 3: 26-34. [Russian Metallurgy (Metally), 2001, 3: 243-250. 5. Bannykh O. A., Blinov V. M., Berezovskaya V. V., et. al. Effect of the γ→α Martensite Transformation in Fe-Cr-N Alloys on Their Stress Corrosion Cracking. Izv. Ross. Akad. Nauk, Ser. Met., 2005, 4: 26-31. [Russian Metallurgy (Metally), 2005, 4: 310-314. 6. Berezovskaya V. V., Golyakov I. V., Bannykh O. A., Blinov V. M. Effect of Cold Plastic Deformation on the Structure and Corrosion Resistance of Austenitic Aging 0Kh18N2A Alloy. Izv. Ross. Akad. Nauk, Ser. Met., 2006, 5: 29-32. [Russian Metallurgy (Metally), 2006, 5: 390-393. 7. Livshits B. G., Kraposhin V. S., Linetskiy Ya. L. Physical Properties of Metals. Moscow: Metallurgiya, 1980.
- 102 -
Structure and properties Evolution, Phase Transformation
HYDROGEN EMBRITTLEMENT OF STAINLESS STEELS ALLOYED WITH NITROGEN. A.I. Balitskii, L.M. Ivaskevich, V.M. Mochulskyi. Karpenko Physiko-Mechanical institute, Lviv, Ukraine.
[email protected]
The resistance to hydrogen degradation of Cr-Ni and Cr-Mn stainless steels with different content of nitrogen (from 0.1 to 0.7 wt. %) after plasma arc remelting was investigated. The dependence of the mechanical properties of steels on nitrogen content is of a complicated character. It is established, that hydrogen caused the changes of phase composition of steels. In hydrogen under the pressure of 10 MPa the strength of unstable Cr18-Ni9-N and Cr18Mn10-N steels and the relative elongation of all materials except Cr20-Ni15-N steel significantly decrease. As the pressure of hydrogen increases to 35 MPa, the intensity of hydrogen embrittlement remains constant. Among the tested steels the best combination of the characteristics of strength and plasticite in hydrogen is exhibited by Cr18-Mn15-N steel with 0.32 % of nitrogen (σu = 850 MPa, σ0,2 = 510 MPa and δ = 60 %). The embrittlement of this material with high content of nitrogen is caused by the presence of nitride particles. KEY WORDS: hydrogen degradation, stainless steels, nitrogen content. ABSTRACT
Nitrogen-containing steels with overequilibrium concentrations of nitrogen form a new class of structural materials with improved corrosion and mechanical characteristics and, in particular, with high resistance to hydrogen embrittlement [1, 2]. However, the procedure of overequilibrium alloying with nitrogen requires special industrial equipment for melting metals under high pressure which restricts its applicability. Therefore, it is reasonable to study the resistance to hydrogen degradation for steels with different equilibrium contents of nitrogen (attained by using standard metallurgical equipment). We studied two types of corrosion-resistant austenitic steels alloyed with nickel or manganese. Nitrogen was introduced in the metal after melting in an induction furnace from the ambient medium in the course of plasma-arc remelting. We obtained ingots with the following concentrations of nitrogen (wt. %): (1) Cr18-Ni9-N (0.077, 0.094, 0.116, and 0.157), (2) Cr20-Ni15-N (0.07, 0.139, 0.195, and 0.21), (3) Cr18-Mn10-N (0.056, 0.198. 0.29, and 0.40), (4) Cr18-Mn15-N (0.176, 0.32, 0.595, and 0.685). The Cr20-Ni15-N and Cr18-Mn15-N alloys have stable austenitic structure. At the same time, the structure of the Cr18-Ni9-N and Cr18-Mn10-N alloys is metastable and contains the residual a-phase the amount of which increases in the process of deformation as a result of the γ→α transformation. After hot rolling, the plates were held for 30min at 1420°C, hardened in water, and subjected to sandblasting treatment to remove scale. The blanks obtained as a result were used to stamp microspecimens with a working part 10 mm in length and 1x3 mm in the cross-sectional dimensions. The specimens were ground, annealed in a vacuum furnace at 1200 K for 30 min, and cooled together with the furnace down to room temperature. The specimens were tested for static tension in gaseous hydrogen in an R(P)TAr-1.0 installation [3] at room temperature for a velocity of motion of an active clamp of 0.1 mm/min. It was discovered that the strength and plasticity of steels attain their minimum values in hydrogen under a pressure of l0 MPa. For this reason, our tests were carried out under the indicated conditions (see Fig. 1). 1. Cr18-Ni9-N steel. In the intact state, this type of steel contains 0.07, 0.094, or 0.116 wt.% of nitrogen and the amount of the residual α-phase in the specimens varies within the range 4-9 wt.%. In
Hydrogen Embrittelement of Stainless Steels Alloyed with Nitrogen
- 103 -
addition, in the process of tension, the specimens undergo the deformation martensitic transformation. This transformation is noticeably affected by nitrogen and hydrogen [4, 5]. The combined action of these factors is responsible for the complicated character of the influence of nitrogen on the mechanical properties of steel. Thus, in air, the ultimate short-term strength of Cr18-Ni9-N steel decreases and the conventional yield limit significantly (by a factor of 1.4) increases as the concentration of nitrogen increases from 0.07 to 0.157 wt. % (Fig 1a). σb,MPa 600
σb,MPa
1
2 2
a
400 200 100 δ,% 75 50 25
b
400
3 4
3
200
6
0.08
0.12
1000
0.16
N, %
60
6
50
0.08
0.20
d
4
400 80
60
5
60
5
6
40
6
0.1
N, %
4
δ,%
20
0.16
2 3
600
δ,%
40
0.12
1
c
2 3
5
800
800 600
70
σb,MPa
1
400
4
δ,%
5
σb,MPa
0
1
600
0.2
0.3
N, %
0.4
20 0.2
0.3
0.4
0.5
0.6
N, %
0.7
Fig. 1. Dependences of the ultimate strength (1, 2), yield limit (3, 4) and relative elongation (5, 6) of the specimens on the concentration of nitrogen in Cr18-Ni9-N (a), Cr20-Ni15-N (b), Cr18-Mn10-N (c), and Cr18-Mn15-N steels (d): (1, 3, 5) - in air, (3, 4, 6) - in hydrogen under a pressure of l0 MPa. It should be emphasized that the plasticity of steel is very high, especially for the concentrations of nitrogen equal to 0.116 and 0.157 wt. %. In this case, the relative elongation of the specimens can be as high as 105-107% (Fig 1a). This effect is well known for the materials suffering the deformation phase transformations [6, 7]. It is attained for a certain kinetics of the processes. Therefore, it is not surprising that the plasticity of specimens significantly decreases in the presence of hydrogen which affects the intensity of the y-cc transformation and promotes the process of cracking along the interfaces. At the same time, the fact that the ultimate short-term strength of steel decreases in hydrogen is quite unexpected. Indeed, this is typical of high-strength nickel alloys. At the same time, the strength of Cr18-Ni10-T steel (whose chemical composition and properties are close to the corresponding characteristics of the analyzed steel) is practically identical in hydrogen and neutral media [5]. 2. Cr20-Ni15-N steel. As the concentration of nitrogen increases from 0.77 to 0.21 wt. %, the ultimate strength σu increases by 20%, the quantity σ0,2 increases by 50%, and the relative elongation of specimens decreases from 80 to 60%. Hydrogen does not affect the mechanical properties of steel even in the case of testing of preliminarily hydrogenated specimens (at a temperature of 623°C, a pressure of 35 MPa, and a time of holding of 5 h). Unfortunately, even in the presence of 0.21% of nitrogen, the strength properties of this type of steel are poor.
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Structure and properties Evolution, Phase Transformation
It seems likely that, in order to strengthen the nitrided iron-nickel austenite with preservation of its hydrogen resistance, it is necessary to introduce additional amounts of nitrogen, possibly together with other hardeners of a solid solution, such as substitutional Mo-type atoms. The procedure of alloying of chromium-nickel austenitic steels with nitrogen is quite promising. This is confirmed by the conclusions made in [8] according to which the negative effect of nitrogen on the hydrogen resistance of materials with high-energy packing defects becomes weaker and our data on the solubility and diffusion mobility of hydrogen. Indeed, as the concentration of nitrogen in Cr20Ni15-N steel increases from 0.077 to 0.21wt.% at 573, 773, or 973°C, the solubility of hydrogen becomes 1.3-1.45 times lower. At the same time, its permeability and the diffusion coefficient become half as large as before. 3. Cr18-Mn10-N steel. After hardening, the Cr18-Mn10-N steel contains a certain amount of the α-phase [from 76 (0.056% N2) to 8% (0.40% N2)]. Just as in the nickel austenite (Cr18-Ni9-N steel), the presence of nitrogen strengthens the solid solution, stabilizes austenite, and (possibly) leads to the formation of nitrides. As a result of introducing 0.4% of nitrogen, the short-term ultimate strength of steel in neutral media increases by 150 MPa and its relative elongation first also increases (from 27 to 75%) but then decreases (Fig. 1c). Nitrogen and hydrogen affect the conditional yield limit relatively weakly. Gaseous hydrogen noticeably decreases the characteristics of plasticity and strength. At the same time, both these parameters remain practically constant in hydrogen-containing media in the entire investigated interval of concentrations of nitrogen (0.0560.40 wt. %). Thus, nitrogen and hydrogen strongly and ambiguously affect the properties of unstable chromium-nickel and chromium-manganese steels (such as Cr18-Ni9-N and Cr18-Mn10-N). 4. Cr18-Mn15-N steel. After hardening, this type of steel contains, possibly, a certain amount of δ-ferrite. In the specimens subjected to annealing at 1200°C followed by gradual cooling in a vacuum furnace, the formation of nitride phases is possible, at least if the concentration of nitrogen is equal to 0.595 or 0.685 wt.%. This assumption is confirmed by the fact that the strength of the specimens decreases as the concentration of nitrogen increases (Fig. 1d). It is known [9-16] that, within certain limits, the redistribution of nitrogen from the solid solution into nitride particles is accompanied by the process of softening of the material. Thus, the behavior of Cr18-Mn15-N steel in hydrogen can be explained as follows: The phenomenon of noticeable softening and decrease in the plasticity of specimens with 0.176% of nitrogen in hydrogen is explained by the presence of residual δ-ferrite. As the content of nitrogen increases to 0.32%, austenite stabilizes. For this reason, hydrogen weakly affects the characteristics of strength and relative elongation
(δ = 60%) (Fig. 1d). In high-nitrogen modifications of Cr18-Mn15-N steel, hydrogen embrittlement manifests itself only in the decrease in plasticity, which is, in fact, typical of steels with stable austenite. CONCLUSIONS
The procedure alloying of Cr20-Ni15-N -type steels is accompanied by a monotonic increase in the characteristics of strength and a decrease in the degree of relative elongation at room temperature in neutral media. The dependences of the mechanical properties of materials on the concentration of nitrogen are complicated, which is explained by changes in the phase composition of steels. In hydrogen, under a pressure of 10 MPa, the strength of unstable Cr18-Mn10-N and Cr18Ni9-N steels and the relative elongation of all materials except Cr20-Ni15-N steel significantly decrease. As the pressure of hydrogen increases to 35 MPa, the intensity of hydrogen embrittlement remains constant. In the collection of materials studied in the present work, the best combination of the characteristics of strength and plasticity in hydrogen is exhibited by Cr18-Mn15-N steel with 0.32% of
Hydrogen Embrittelement of Stainless Steels Alloyed with Nitrogen
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nitrogen (σu = 850 MPa, σ0,2 = 510 MPa and δ = 60%). The embrittlement of this material with high content of nitrogen is caused by the presence of nitride particles.
REFERENCES 1. Ts. V. Rashev, Production of Alloyed Steels, Publ. House of the Bulgarian Academy of Sciences, Sofia, 1980. 2. V. I. Vytvyts'kyi, V. I. Tkachov, and S. O. Hrebenyuk, Alloying of steel to overequilibrium concentrations of nitrogen from gaseous atmosphere in the course electroslag remelting, Fiz.-Khim. Mekh. Mater., 2000, 36, No. 3, 115-116 . 3. G. G. Maksimovich, O. N. Voznichak, I. Yu. Tret'yak, et al., High-temperature tensile testing in gaseous high-pressure media, Probl. Prochn., 1984, No. 9, 97-99 . 4. V. I. Shapovalov, Influence of Hydrogen on the Structure and Properties of Iron-Carbon Alloys, Moscow: Metallurgiya, 1982. 5. G. G. Maksimovich, I. Yu. Tret'yak, L. M. Ivas'kevich, and T. V. Slipchenko, On the role of the martensitic transformation in the process of hydrogen embrittlement of unstable austenitic steels, Fiz.-Khim. Mekh. Mater.,1985, 21, No. 4, 29-32 . 6. M. M. Shteinberg, L. G. Zhuravlev, and O. P. Chernogorova, Formation of martensite under loading and its influence on the mechanical properties of metastable austenitic alloys, Fiz. Met. Metalloved., 1977, No. 1, 217-220 . 7. B. A. Potekhin, Contribution of the martensitic transformation in the process of deformation to the plasticity of metastable austenitic steels, Fiz. Met. Metalloved., 1979, No. 5, 1065-1075 . 8. A. U. Thompson and I. M. Bernstein, The role of metallurgical factors in the processes of hydrogen-assisted fracture, in: Advances in Corrosion Science and the Technology of Corrosion Protection. Corrosion Cracking of Metals, Moscow: Metallurgiya, 1985, 47-159. 9. L. D. Afanas'ev, V. G. Gavrilyuk, V. A. Duz', and V. L. Svechnikov, Aging of cold-worked nitrogen-containing austenitic steels, Fiz. Met. Metalloved., 1990, No. 7, 106-110 . 10. Balitskii A.I. 3rd International Conference “High Nitrogen Steels”, Physicochemical Mechanics of Materials, 1994, vol.30, No 1, 145 . 11. Balitskii A., Pokhmurskii V., Khoma M., Balitskii O. Investigation of Deformation Influence on Morphology of High Nitrogen Steel Surface Layers, High Nitrogen Steels: HNS-95 (Proceedings of 4th International Conference, Sept.27-29, 1995, Kyoto, Japan, 63. 12. Balitskii A.I. Long Term Strength of High Nitrogen Steels in Water and Chloride Solution, High Nitrogen Steels (Abstracts of 5th International Conference, Espoo-Finland, May 24-26, Stockholm-Sweden, May 27-28, 1998, 143. 13. Balitskii A.I., Diener M., Magdowski R.M., Pokhmurski V.I., Speidel M.O. Anisotropy of Fracture Toughness of Austenitic High Nitrogen Chromium-Manganece Steel, High Nitrogen Steels (Abstracts of 5th International Conference, Espoo-Finland, May 24-26, Stockholm-Sweden, May 27-28, 1998), 35. 14. Balitskii A.I. Increasing of Exploitation Characteristic of High-Nitrogen Cr-Mn Steels, High Nitrogen Steels.Institute of Metallurgy Swiss Federal Institute of Technology, ETH: VDF Hochschulerlag AG an der ETH Zurich, 2003, 323-331. 15. Balitsky A. I. International Conference of High-Nitrogen Steels HNS-2004, Electrometallurgy Today, 2004. № 4, 56 . 16. Balitskii A.I. Applications of high nitrogen steels in nuclear power plants equipments , Proceedings of the International Conference on High Nitrogen Steels 2006, August 29 to 31 2006, Jiuzhaigou, Sichuan, Chine.Editede by Han Dong, Jie Su, Markus O.Speidel.- Beijing: Metallurgical Industry Press, 2006, 295-302.
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Structure and properties Evolution, Phase Transformation
Influence of alloying and temperature strain upon formation of structure, necessary mechanical and corrosion properties in austenitic nitrogen containing steels V.V Rybin., G.Y. Kalinin, S.Y.Mushnikova Central Research Institute of Structural Materials “Prometey”, St.Petersburg, 191015, Russia ABSTRACT Results of studies of the influence of alloying elements (nitrogen, chrome, manganese, nickel, molybdenum, vanadium and niobium) upon structure, mechanical properties and pitting corrosion resistance of steel in chloride water solutions are presented in this paper. On the basis of these results a chemical composition for new nitrogen containing steel NS-5T (0.04C-20Cr-14Ni-6Mn-2Mo-N-Nb) was selected. While devising a process of hot rolling of NS-5T steel it was found out that after hot plastic deformation steel had a clearly distinct texture which included nitrides and carbides of chrome and niobium, besides austenite. Heating up beyond a recrystallization temperature, with subsequent water quenching, provides for dissolving of secondary phases and formation of a uniform solid solution structure with σ0,2 = 400-500 MPa. To obtain a corrosion-resistant austenitic nitrogen containing steel of higher strength with fcc lattice being kept under active thermal and deformation influence the following chemical composition was recommended: ≤ 0,05% C; 0,45÷ 0,55% N; 18-20% Cr; ~18%(Ni + Mn); 0,2-0,45%(V+Nb); 1,5÷ 1,8% Mo. Optimum combination of mechanical properties (σ0,2 ≥ 690 MPa, KCV+20ºC ≥ 80 J/sm2) and high corrosion resistance in the same composition is achieved after special high temperature thermomechanical processing (HTMP) which provides for formation of a mixed structure consisting of re-crystallized, fragmented and polygonized components without nitrides precipitation at grain boundaries. KEY WORDS austenitic steels, pitting corrosion, pitting formation potential, repassivation potential, thermomechanical treatment.
1 Introduction
The greater part of austenitic steels is characterized by yield strength not higher than 400 MPa, what restricts essentially their usage as materials for heavily loaded structures. A known fact [1-7] of nitrogen favorable influence upon strength and corrosion properties of Cr-Ni-Mn steels of austenitic class extends the range of nitrogen containing steels application in different fields, including shipbuilding structures of different purposes (hulls, oil and gas production platforms, marine engineering equipment and so on). The aim of this work consisted in development of stainless shipbuilding steels of different strength levels: the first one – with the yield strength 500600 MPa, and second – with the yield strength more than 690MPa. The investigation itself consisted in selection of alloying compositions, study of alloying elements influence upon nitrogen solubility in steel, formation of structure, obtaining of mechanical and corrosion properties, determination of principal process diagram of rolled plates manufacturing suitable for promotion at metallurgical mills. 2 Methods of investigations
In laboratory conditions steels with different content of alloying elements were melted in induction furnaces and teemed into 25 kg moulds. Steel ingots weighting up to 5÷50 t were manufactured in electric-arc furnaces for production of pilot-industrial batches. Nitrogen containing chrome and manganese ferroalloys were used for nitrogen alloying. Manufactured ingots were subjected to hot deformation (forging and rolling into plates) at temperature 1200-850°C. If special mechanical properties were required, plates were subjected to heat treatment at temperature 1050-1100°C. Mechanical properties of plates at tension were determined on cylindrical specimens. Impact bending tests were carried out on V-notched Charpy specimens. Pitting corrosion resistance of steel (kind of corrosion which is most dangerous for stainless steels used in sea water) was evaluated by potentiodynamic polarization curves in 3,5% NaCl solution at room temperature (with determination of pitting potential Ep and repassivation
Influence of alloying and temperature strain upon formation of structure, necessary mechanical and corrosion properties in austenitic nitrogen containing steels
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potential Erep) and chemical method by immersion in 6% FeCl3 solution. Steel resistance to intervrystalline corrosion was determined by boiling in copper sulphate and sulfur acid solution with copper cuttings. Steel resistance to stress corrosion cracking was determined by slow strain rate testing (SSRT) and by cantilever bending of precracking specimens. Structure was studied with an optical microscope Neophot-2, and with X-ray electronic microscope JEM-200CX. 3 Investigations results and discussion 3.1 Austenitic steel with yield strength 400-500 MPa.
In development of chemical composition of stainless steel with enhanced strength it was selected a basic composition 20Cr-14Ni-6Mn which provided formation of a stable austenitic structure of steel and high solubility of nitrogen in γ-hard solution. For optimization of steel structure and properties alloying elements content was varied within the following limits: 0,250,60% N, 17-26% Cr, 3-12% Mn, 10-18% Ni, 0-3,5% Mo, 0,1-0,6% V, 0,1-0,6% Nb. As nitrogen concentration was growing up, other alloying elements being kept at a basic level, it was found out that (fig. 1a) strength characteristics (σ0,2 and σB) were enhanced, level of plasticity characteristics remained constant and impact toughness was somewhat decreased (as a result of nitrides formation along grains boundaries). Increase of resistance to pitting corrosion (fig. 1b, c), when nitrogen content grows up from 0,25% up to a value of maximum dissolution of nitrogen in austenite (Nmax. dis. ~ 0,27-0,29%) calculated by known empiric formulae [8], is explained by nitrogen concentration in hard solution [9]. With further increase of nitrogen quantity resistance to pitting corrosion is kept at a constant high level, and at nitrogen content equal to 0,5-0,6% a decrease of pitting corrosion resistance is observed what is explained by nitrogen and chrome combination (linking) into nitrides while formation of their grain boundary precipitations. Maximum recommended nitrogen content in steel (0,40%) is also limited to avoid a possibility to obtain non-porous ingot in melting of nitrogen containing steel. As chrome content is growing from 17% up to 23% steel strength increases, plasticity decreases slightly, resistance to pitting corrosion increases. If chrome content is higher than 23%, steel plasticity and impact toughness decreases sharply, resistance to pitting corrosion slightly decreases, what is explained by appearance of δ-ferrite and σ-phase in steel structure. Thus, at 18-22% chrome content steel possesses optimum complex of mechanical properties (σ0,2 = 470-570 MPa, KCV+20°C = 130-180 J/sm2) and high resistance to pitting corrosion. Molybdenum addition into steel increases essentially resistance to pitting corrosion and strength level. But if Mo content is higher than 3,5% ageing processes are being developed in steel with χphase formation (precipitation) similar to σ-phase by its negative influence upon plasticity and impact toughness. Hence, Mo content within the limits 2-3% is optimum. Investigations of heats with variable manganese and nickel content showed that these elements stabilize austenitic structure of steel but slightly affect its mechanical and corrosion properties. However, at nickel content less than 10% and manganese content higher than 12% mechanical properties and corrosion resistance deteriorate because of unsatisfactory stability of γ-phase and formation (precipitation) of α-phase and ε-martensite respectively. Besides, in the whole range of manganese concentrations it was observed a stability of potentials Ep and Erep which testify that there was no bad influence of this element. Lower boundaries of nickel (13%) and manganese (5%) alloying, selected as a result of this investigation, provide a stable austenitic structure of steel, and top boundaries of nickel (16%) and manganese (7%) alloying are limited by increase in price, in the first case, and by possibility of martensitic phases formation and ecological danger in melting and welding, in the second case.
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Structure and properties Evolution, Phase Transformation
Ep, mV (she) 1500
Erep, mV (she) 1500
σB,σ0,2, МPa
1400
1400
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1200
1200
1100
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δ5 ,ψ,% 80 60
ψ
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0
0,2
0,3
0,4
0,5 0,6 N, wt.-%
a) b) c) Fig. 1 Nitrogen influence upon of pitting potential Ep (a), repassivation potential Erep (b) and mechanical properties (c) of steel 0.04C-20Cr-14Ni-6Mn-2Mo-N-Nb.
Introduction vanadium and niobium being strong nitride generating elements does not allow for a considerable increase of strength of the new steel in condition after heat treatment (homogenization). But, taking into account a necessity of niobium alloying in order to avoid intercrystalline corrosion, on the one hand, and possibility of separation of a large quantity of niobium carbides and nitrides which causes a decrease of plasticity and impact toughness, on the other hand, the following limits of these elements concentrations were accepted: ≤ 0,1% V and 0,15-0,25% Nb. If carbon content increases more than 0,04%, steel plasticity, impact toughness and resistance to corrosion slightly decreases in condition after heat treatment. However, after soakings at temperature 850-900°C, which cause formation of special carbides and carbonitrides, resistance of steel with 0,09-0,13% C content to pitting and intercrystalline corrosion decreases sharply. On the basis of results of investigation of laboratory and industrial heats a new chemical composition of new nitrogen containing steel of NS-5T grade (0.04C-20Cr-14Ni-6Mn-2Mo-N-Nb) was selected: ≤ 0,04% C; 0,30-0,40%N; 18-22%Cr; 14-16% Ni; 5-7% Mn; ≤ 0,1% V; 0,15-0,25% Nb; 2,0-3,0% Mo; ≤ 0,8% Si. In the process of working out of heat treatment conditions for NS-5T steel [10] it was found out that heating to temperatures lower than 1050°C wasn’t enough for stress relieving (relaxation), dissolution of secondary phases and obtaining of a structure of hard uniform solution. At heating above 1050ºC undesirable growth of grain was observed. Soaking of metal at temperature 10501150°C for 2-4 min/mm with accelerated water quenching was optimum. 3.2
Austenitic steel with yield strength 690 MPa.
CRISM “Prometey” together with A.A.Baikov IMET carries out a complex of investigations with the aim to develop austenitic steel with yield strength equal to 690 MPa and higher [11-19]. A promising method to obtain high-strength nitrogen containing steel (σ0,2≥ 690 MPa) which combines different hardening mechanisms (solid solution, dispersed hardening at precipitation of
Influence of alloying and temperature strain upon formation of structure, necessary mechanical and corrosion properties in austenitic nitrogen containing steels
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hardening phase particles, grain refinement, work hardening) is a high temperature thermomechanical processing (HTMP) [20] consisting in heating of steel to temperature 12001260°C, what provides formation of homogeneous solid solution of austenite, and hot rolling as per certain conditions with subsequent water quenching. On the basis of principle that at plastic deformation austenite Cr-Mn-N is hardened to a larger extent than austenite Cr-Ni-C it was selected a composition 20Cr-6Ni-11Mn. It was found out that for maintenance of γ-phase stability in steel nitrogen content equal to 0,4% was enough at carbon content equal to 0,04% and lower. At carbon content > 0,09% C and nitrogen content > 0,55% N it is difficult to obtain satisfactory plasticity and impact toughness (especially in welded joints) because of formation of chrome carbides of Me23C6 type and chrome nitrides of Cr2N type at thermal soakings. Addition of 19-21% Cr into steel is necessary to provide corrosion resistance and nitrogen solubility within the indicated limits. If chrome content is more than 21% and nitrogen content is less than 0,45% steel could be characterized by a lowered plasticity because of δ-ferrite and α-phase formation. As nickel content is increasing above 9% nitrogen solubility in metal is decreasing. Manganese content equal to 8-11% increases nitrogen solubility and provides stability of austenite. Introduction of niobium into steel at a quantity not less than 0,10% and vanadium more than 0,10% provides fine grained structure and better strength thanks to formation of finedispersed nitrides of vanadium and niobium. At a lower V and Nb concentration positive effect of their presence is not considerable. Increase of niobium content above 0,30% and vanadium – above 0,60% results in decrease of impact toughness of steel because of hard solution leaning with nitrogen as a result of formation of difficult-to-dissolve nitrides and carbonitrides of vanadium and niobium mostly along the grain boundaries. Molybdenum at a quantity of 1,5% increases steel resistance to local kinds of corrosion, but if molybdenum content exceeds 2% there is a danger of ferro-magnetic phases formation (martensite, ferrite). Thus, it was determined a chemical composition of high-strength steel of 0.04C-20Cr-6Ni-11Mn-N-V-Nb grade as follows: ≤ 0,04%C; 0,45-0,55% N; 18-20% Cr; ~ 18% (Ni+Mn); 0,2-0,45% (V+Nb); 1,5-1,8% Mo. In the course of development of high temperature thermo-mechanical treatment process for the new steel the influence of temperature-deformation conditions while hot rolling and subsequent cooling down (in air and water) was studied. It was found out that [21, 22] as a degree of hot plastic deformation is increasing steel strength increases and plasticity decreases. Using microstructural and X-ray spectral methods it was found out that decrease of plasticity and impact toughness is connected, to a large extent, with precipitation of nitride phases containing chrome (Cr2N) and the top boundary of this chrome separation (precipitation) is located in the range (area) 950-1000°C. Fig. 2 shows metallographic and fine structures with the indication of basic mechanical properties of nitrogen containing steel depending of finish rolling temperature. As it follows from the obtained data, if finish rolling temperature is lower than 950°C (Tfr < 950°C) steel has a clearly marked texture (stripped structure), grains boundaries, in this case, are decorated with nitrides what testifies that cold worked steel is characterized by high strength, low plasticity and decrease of pitting corrosion resistant. Water quenching after rolling heating (Tfr=950-1000°C) provides formation of a mixed structure with recrystallized, fragmented and poligonized components. With such a structure and in the absence of secondary phase’s precipitations along grains boundaries best combination of mechanical and corrosion resistance properties is provided. If a finish rolling temperature is higher than 1000°C (Tfr> 1000°C) metal cooled in water will have a completely recrystallized structure, but if only a solid solution component is available, steel strength will decrease to σ0,2=500 MPa.
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Structure and properties Evolution, Phase Transformation
25µm
0,5μm
Stripped structure Tend of rolling < 950оС
25µm
0,5μm
Mixed structure Tend of rolling = 950-1000оС
25µm
0,5μm
Recrystallized structure Tend of rolling > 1000о С
a) b) c) Fig. 2 Metallographic and fine structures of steel 0.04C-20Cr-6Ni-11Mn-N-V-Nb hardened by HTMP depending on finish rolling temperature.
On the basis of many experimental results it has been developed a technology of high temperature thermomechanical processing (HTMP) of high-strength austenitic nitrogen containing steel, with basic requirements made to this process being as follows: a total degree of hot deformation should be not less than 70% (15-20% per a pass); finish rolling temperature should be about 950°C; accelerated water quenching after rolling. 3 Conclusion
Alloying compositions have been selected and it has been determined optimum content of alloying elements for nitrogen containing steels with yield strength 400-500 MPa (0.04C-20Cr14Ni-6Mn-2Mo-N-Nb) and 690 MPa (0.04C-20Cr-6Ni-11Mn-N-V-Nb). Manufacturing processes of new steels which provide high levels of pitting and intercrystalline corrosion resistance and mechanical properties have been developed. It has been determined that regulation of temperature-deformation conditions of high temperature thermomechanical processing (HTMP) allows for obtaining a certain structural condition of steel 0.04C-20Cr-6Ni-11Mn-N-V-Nb which provide a necessary mechanical and corrosion resistance properties: recrystallized structure of austenite, characterized by low density of nitrides precipitations inside grains and on grains boundaries, provides σ0,2 =500-600 MPa, KCV+20°C=100-200 J/cm2 and high resistance to corrosion; mixed structure having recrystallized, fragmented and polygonized components without nitrides precipitations provides σ0,2 ≥ 690 MPa, KCV+20ºC ~ 100 J/cm2 and high corrosion resistance; stripped structure with dislocations density increased strength, but reduced impact toughness and corrosion resistance.
Influence of alloying and temperature strain upon formation of structure, necessary mechanical and corrosion properties in austenitic nitrogen containing steels
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References 1. ASM Specialty Handbook Stainless Steels: Ed. J. R. Davis. The materials information society .USA, 1994. 2. Катада Э., Ванишцу Н., Бабак Х. Стали с повышенным содержанием азота, разработанные в национальном институте материаловедения. Металловедение и термическая обработка металлов, 2005, 11 (605): 14-16. 3. Шпайдель М.О. Новые азотосодержащие аустенитные нержавеющие стали с высокими прочностью и пластичностью. Металловедение и термическая обработка металлов, 2005, 11 (605): 9-13. 4. Gavriljuk V.B. Nitrogen in Iron and Steel. ISIJ International, 1996, 36 (7): 738-745. 5. Гаврилюк В.Г. Физические основы конструирования азотистых сталей. Известия РАН. Серия физическая, 2005,. 69 (10): 1470-1474. 6. Козлов Э.В., Теплякова Л.А., Конева Н.А., Гаврилюк В.Г., Попова Н.А., Игнатенко Л.Н., Федосеева Г.Л., Смук С.Ю., Пауль А.В., Подковка В.П. Роль твёрдорастворного упрочнения и взаимодействий в дислокационном ансамбле в формировании напряжения течения азотсодержащей аустенитной стали. Известия высших учебных заведений. Физика, 1996, 3: 33-56. 7. Corrosion Test and Standards: Application and Interpretation. Ed. R.Baboian.- USA: ASTM, 1995. 8. Приданцев М. В., Талов Н. П., Левин Ф. Л. Высокопрочные аустенитные стали. М.: Металлургия, 1969. 9. Мушникова С.Ю., Легостаев Ю.Л., Харьков А.А., Петров С.Н., Калинин Г.Ю. Исследование влияния азота на стойкость к питтинговой коррозии аустенитных сталей. Вопросы материаловедения, 2004, 2(38): 126-135. 10. Калинин Г. Ю., Легостаев Ю. Л., Малышевский В. А., Мушникова С. Ю., Харьков А. А. Новая коррозионностойкая азотсодержащая аустенитная сталь НС−5Т, Вопросы материаловедения, 1996, 3(6): 5−15. 11. Костина М.В., Банных О.А., Блинов М.В., Дымов А.А. Легированные азотом хромистые коррозионностойкие стали нового поколения. Материаловедение, 2001, 2 (47): 35-44. 12. Блинов В.М., Банных О.А., Костина М.В., Ригина Л.Г., Блинов Е.В. О влиянии легирования на предельную растворимость азота в коррозионно-стойких низкоуглеродистых сплавах Fe-Cr-Mn-Ni-Mo. Металлы, 2004, 4: 42-49. 13. Устиновщиков Ю.И., Рац А.В., Банных А.О., Блинов В.М. Структура высокоазотистых сплавов Fe-18%Cr. Металлы, 1996, 1: 67-73. 14. Устиновщиков Ю.И., Рац А.В., Банных А.О., Блинов В.М., Костина М.В., Морозова Е.И. Структура и свойства высокоазотистых аустенитных сплавов Fe-18%Cr, содержащих до 2%Ni. Металлы, 1998, 2: 38-43. 15. Калинин Г. Ю., Мушникова С. Ю., Фомина О. В., Харьков А. А. Исследование структуры и свойств высокопрочной коррозионностойкой азотистой стали 04Х20Н6Г11М2АФБ. Вопросы материаловедения, 2006, 1(45): 45−54. 16. Горынин И. В., Рыбин В. В., Малышевский В. А., Калинин Г. Ю., Мушникова С. Ю., Малахов Н. В., Ямпольский В. Д. Создание перспективных принципиально новых коррозионно-стойких корпусных сталей, легированных азотом. Вопросы материаловедения, 2005, 2(42): 40−54. 17. Банных О.А., Блинов В.М., Костина М.В. Исследование эволюции структуры азотистой коррозионностойкой аустенитной стали 06Х21АГ10Н7МФБ при термодеформационном и термическом воздействии. Вопросы материаловедения, 2006, 1(45): 9−20. 18. Сагарадзе В.В., Уваров А.И., Печеркина Н.Л., Калинин Г. Ю., Малышевский В. А., Ямпольский В. Д. Структура и механические свойства толстолистовой азотосодержащей аустенитной стали, 04Х20Н6Г11М2АФБ. Физика металлов и металловедение, 2006, 102 (3): 250-256. 19. Костина М. В., Банных О. А., Блинов В. М. Особенности сталей, легированных азотом. Металловедение и термическая обработка металлов, 2000, 12: 3−6. 20. Коджаспиров Г.Е., Сулягин Р.В., Карьялайнен Л.П. Влияние температурно–деформационных условий на упрочнение и разупрочнение азотосодержащих коррозионно-стойких сталей. Металловедение и термическая обработка металлов, 2005, 11(605): 22-26. 21. Калинин Г. Ю., Малышевский В. А., Мушникова С. Ю., Петров С. Н., Ямпольский В. Д. Влияние степени горячей пластической деформации на микроструктуру и механические свойства аустенитной высокопрочной коррозионностойкой стали 05Х19Н5Г12АМ2БФ. Вопросы материаловедения, 2003, 4(36): 5−11. 22. Банных А.О., Блинов В.М., Костина М.В., Калинин Г.Ю. Влияние режимов горячей прокатки и термической обработки на структуру, механические и технологические свойства аустенитной азотосодержащей стали 05Х22АГ15Н8М2Ф-Ш. Металлы, 2006, 4: 33-41.
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Structure and properties Evolution, Phase Transformation
Effect of Aging on Mechanical Properties of High Nitrogen Austenitic Stainless Steel Zu-rui ZHANG, Zhou-hua JIANG, Hua-bing LI, Bao-yu XU (Northeastern University, School of Materials and Metallurgy, Liaoning, Shenyang, 110004, China) Abstract: High nitrogen austenitic stainless steel 18Cr-18Mn-2Mo-0.77N in aging treatment was investigated by using optical microscopy (OM), scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Results showed that intergranular, cellular, intragranular Cr2N and Chi (χ) precipitated gradually during the aging treatment. 0.05% precipitation volume fraction was determined as the initial curve of time-temperature-precipitation (TTP). At 900℃, Vickers hardness firstly decreased because the intergranular precipitation made the matrix nitrogen depletion. With increasing the aging time, more precipitation including cellular, intragranular Cr2N and χ presented a larger effect on hardness than nitrogen depletion in the matrix. It was the same reason for aging treatment at 800 and 850℃ for 10 hours. Impact toughness reduced sharply with the prolonged aging treatment, and SEM morphologies showed leading brittle intergranular fracture. Tensile strength deteriorated gradually with aging time. However, yield strength firstly increased then decreased because of the reduction of resistance on mobile dislocation in the matrix for nitrogen depletion, whereas larger number of precipitation made more resistance on mobile dislocation with increasing aging time. During the isochronal aging treatment for 10 hours, cellular Cr2N can be only observed at 900℃, which is main reason for the highest Vickers hardness and lowest impact toughness. Key words: High nitrogen austenitic stainless steel; aging treatment; Cr2N; Chi (χ) phase; mechanical properties.
1. Introduction
High nitrogen austenitic stainless steels with a favorable combination of strength and toughness at room temperature have been developed as a type of structural materials [1]. Nitrogen as an interstitial solution alloy element increases strength level of the austenite matrix without deteriorating ductility and toughness [2-4]. Nitrogen is instead of the expensive Ni as a strong stabilizing element in HNS [5]. However, the high nitrogen content will induce the formation of brittle precipitation in the range of 650℃ to 900℃ in aging treatment or hot forming process [6]. Moreover, the precipitated behavior in the range of 650℃ to 900℃ is the most sensitive. The investigation on mechanical properties of solution-treated HNS and precipitated behavior of HNS has been possessed in recent years [7]. Meanwhile, the effect of aging on mechanical properties of HNS has not been widely reported yet. This paper investigated the TTP behavior and the effect of aged precipitation on mechanical properties including Vickers hardness, impact and tensile toughness of aged HNS. 2. Experimental procedure The present investigated material was 18Cr-18Mn-2Mo-0.77N high nitrogen steel (HNS) with the following composition in wt pct: 19.15Cr, 18.83Mn, 2.25Mo, 0.77N, 0.05C, and balance Fe. The ingots used in this experiment were manufactured using vacuum induction furnace by adding nitrided ferroalloy under nitrogen atmosphere, and using the electroslag furnace for remelting under nitrogen atmosphere [8]. The plates were hot rolled to 6 mm to prepare the V-notched Charpy impact test specimens from the longest part of these as-quenched specimens which is parallel to the rolling direction with a size of 5×10×55mm3. Some hot rolled plates were cold rolled to 1.5mm to prepare the tensile test specimens with a size as shown in Fig.1. All specimens were solution-treated at
1100℃ for 60 minutes, followed by water quenching. Then the impact and tensile test specimens were isothermally aged at 800 and 850℃ for 10 hours, and at 900℃ for 1, 10 and 50 hours, and quenched into water.
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Fig.1 Schematic of tension specimens used in present study, mm
Uniform tensile tests were performed with 3mm/min cross speed at room temperature. The microstructural observations of each aged, impact and tensile-tested specimen were carried out on a Carl Zeiss OM and SSX-550 SEM operating at 15 kV. Thin foils for TEM were prepared by mechanical grinding down to 80 μm, and finished by a twin-jet electrolytic polishing technique using an electrolyte at 70V of 8% perchloric acid in ethanol at -25℃. The electron microscopy and analysis was used by TECNCI G2 20 TEM at 200 kV. 3. Results and discussion 3.1 Time-Temperature-Precipitation Behavior
This paper used quantitative metallographic analysis with OM and Image-pro Plus 5.0 software to determine the time-temperature-precipitation curve of aged HNS as shown in Fig.2. 0.05% precipitation volume fraction was determined as the initial curve and the nose temperature of precipitation was found to be 900℃. Fig.3 shows a series of OM micrographs of aged HNS at 900 ℃ for various time. Arrow A shows Cr2N which firstly nucleate along grain boundaries. With increasing the aging time, Cr2N grows up and inward austenitic grains by cellular form. Then χ phase was observed in the long time aged specimens. The identification of the precipitation was processed by TEM analysis as shown in Fig.4, Fig.5 and Fig.6. Arrow B, C and D point at the intergranular, cellular Cr2N and χ phase, respectively.
Fig.2 Time-temperature-precipitation curve of aged 18Cr-18Mn-2Mo-0.77N
(b)
(a)
A
(c)
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Structure and properties Evolution, Phase Transformation
(d)
(e)
Fig.3 OM micrographs of aged 18Cr-18Mn-2Mo-0.77N at 900℃ for (a) 30s, (b) 300s, (c) 3.6ks, (d) 36ks and (e) 180ks
(b)
(a)
[ 1 1 0 3] (1 1 2 0)
B
(2 1 1 1) Fig.4 TEM micrograph of intergranular Cr2N. (a) BF image and (b) SAD pattern of the aged HNS at 900℃ for 10h
(b)
(a)
[5 7 2 3] (1 011)
C
(2 1 3 1)
Fig.5 TEM micrograph of cellular Cr2N. (a) BF image and (b) SAD pattern of the aged HNS at 900℃ for 10h
(b)
(a)
[ 1 3 5] (2 1 1)
D
(1 2 1)
Fig.6 TEM micrograph of intermetallic χ phase (a) BF image and Fig.6 TEM micrograph of intermetallic χ phase (a) BF image and (b) SAD pattern of the aged HNS at 900℃ for 50h
3.2 Effect of precipitation on mechanical properties To examine the effects of precipitation on the mechanical properties of HNS, Vickers hardness, Charpy impact and tensile tests were performed. The results of Vickers hardness tests to the aged
HNS are shown in Fig.7. With increasing the aging time at 900℃ (Fig.7 (a)), the value of Vickers hardness firstly decreased because of the matrix nitrogen depletion through intergranular Cr2N precipitated. Then the value increased to more than the solution-treated (ST) specimens because the
Effect of Aging on Mechanical Properties of High Nitrogen Austenitic Stainless Steel
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cellular Cr2N and χ phase which presented a larger effect on hardness than nitrogen depletion in the matrix precipitated during the 10 to 50h aging treatment. The values Vickers hardness of the specimens which were aged for 10h at various temperature as shown in Fig.7 (b) present a similar tendency to Fig.7 (a). Although intergranular Cr2N which made the matrix nitrogen depletion to decrease the hardness can be observed in aged specimens at 800 and 850℃ for 10h, precipitation of both intergranular and cellular Cr2N in aged specimens for 10h at 900℃made the increment more than ST specimens. (a)
(b)
Fig.7 Vickers hardness with various (a) aging time at 900℃and (b) aging temperature for 10h
Fig.8 shows the correlation between the precipitation and impact toughness of aged HNS. Impact toughness reduced sharply with the prolonged aging treatment, and SEM morphologies presented leading brittle intergranular fracture at room temperature as shown in Fig.9. With prolonging aging time or increasing aging time, more precipitation precipitated along grain boundaries and grew inward grains to cellular form, which made deformation along grain boundaries prior than grain interior. The lamellar stripping structure was due to the formation of many small cleavage steps which nucleated along the precipitation. (a)
(b)
Fig.8 Impact toughness with various (a) aging time at 900℃and (b) aging temperature for 10h
Tensile tests were performed on an Instron tensile-tested machine with 3mm/min crosshead speed at room temperature. Table 1 and Fig. 10 reveal the relationship between tensile properties (stress-strain curve) and precipitated behavior aged at 900℃ for various aging time. Tensile strength and elongation deteriorated gradually with prolonging the aging time which meant forming more intergranular Cr2N precipitation, cellular Cr2N and χ phase. However, yield strength firstly increased because of the reduction of resistance on mobile dislocation in the matrix for nitrogen depletion, and then decreased because larger number of precipitation made more resistance on mobile dislocation than the nitrogen depletion in the matrix with increasing aging time. Fig.11 presents a series of SEM micrographs of tensile fracture surfaces of aged HNS at 900℃ for various time. The SEM morphology of ST specimen shows a favorable ductility and equiaxial dimples in Fig.11 (a). Dimples and intergranular crack exist together in the specimen aged for 1h as shown in Fig.11 (b). With increasing the aging temperature, more precipitation made deformation along grain
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Structure and properties Evolution, Phase Transformation
boundaries prior than grain interior and many small cleavage steps which nucleated along the precipitation insect to form lamellar stripping structure as shown in Fig.11 (c) and Fig.11 (d).
(b)
(a)
(c)
(e)
(d)
Fig. 9 Morphologies of fracture surfaces of aged HNS (a) 900℃, 1h, (b) 900℃, 10h, (c) 900℃, 50h, (d) 800℃, 10h and (e) 850℃, 10h
Table 1 Effects of aged 18Cr-18Mn-2Mo-0.77N at 900℃ on tensile properties Aging time (h) 0 (ST) 1 10 50
Rm, MPa 970 960 870 810
Rp0.2, MPa 586 595 569 550
Elongation, % 51.2 50.4 15.2 7.2
Fig.10 Stress-strain curve of aged HNS at 900℃ for various aging time (a)
(b)
(c)
(d)
Fig. 11 SEM morphologies of tensile fracture of aged HNS at 900℃ for (a) 0h (ST), (b) 1h, (c) 10h and (d) 50h
Effect of Aging on Mechanical Properties of High Nitrogen Austenitic Stainless Steel
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Conclusions
(1) Initial time-temperature-precipitation curve of aged 18Cr-18Mn-2Mo-0.77N HNS starts with 0.05% precipitation volume fraction and the nose temperature of precipitation was found to be 900 ℃. Based on TEM analysis, intergranular, cellular, intragranular Cr2N and χ phase precipitated gradually during the aging treatment. (2) With prolonging the aging time or increasing the aging temperature, Vickers hardness firstly decreased then increased because the intergranular precipitation made the matrix nitrogen depletion, whereas more precipitation including cellular, intragranular Cr2N and χ presented a larger effect on hardness than nitrogen depletion in the matrix. (3) Impact toughness reduced sharply with the prolonged aging treatment, and SEM morphologies showed leading brittle intergranular fracture. Tensile strength and elongation deteriorated gradually with the prolonged aging treatment. However, yield strength firstly increased then decreased because of the reduction of the resistance on mobile dislocation in the matrix for nitrogen depletion, whereas larger number of precipitation made more resistance on mobile dislocation with increasing aging time. (4) SEM morphologies of impact and tensile fracture surfaces present the lamellar stripping structure because the formation of many small cleavage steps which nucleated along the precipitation. With increasing the aging temperature, more precipitation made deformation along grain boundaries prior than grain interior. Acknowledgments The authors would like to thank the financed support of National Natural Science Foundation of China and Baosteel Group Company No. 50534010. References 1. Uggowitzer P J, Magdowski R and Speidel M O. Nickel Free High Nitrogen Austenitic Steels [J]. ISIJ International, 1996, 36 (7): 901-908. 2. Speidel M O. Properties and Applications of High Nitrogen Steels, in J. Foct and A. Hendry (eds.), HNS 88, Lille, France, May, 1988, The Institute of Metals, London, 1989: 92 3. Simmons J W. High-nitrogen alloying of stainless steels [J]. Materials Science and Engineering A, 1996, 207: 159-169. 4. Gavriljuk V G, Berns H. High Nitrogen Steels (structure manufacture applications), Springer-Verlag Berlin, 1999: 135 5. Mudali U K, Raj B. High Nitrogen Steels and Stainless Steels, ASM International, 2004, Chapter 6. 6. Lee T H, Oh C S, and Lee C G et al. Precipitation characteristics of the second phases in high-nitrogen austenitic 18Cr-18Mn-2Mo-0.9N steel during isothermal aging [J]. Metals and Materials International, 2004, 10 (3): 231-236. 7. Lee T H, Kim S J and Takaki S. Time-Temperature-Precipitation Characteristics of High-Nitrogen Austenitic Fe-18Cr-18Mn-2Mo-0.9N Steel [J]. Metallurgical and Materials Transactions, 2006, 37A (12): 3445-3454. 8. Li H B, Jiang Z H, Shen M H et al. Manufacturing high nitrogen austenitic stainless steels by nitrogen gas alloying and adding nitrided ferroalloys [J]. Journal of Iron and Steel Research International, 2007, 14(3): 6368.
Effect of Aging on Mechanical Properties of High Nitrogen Austenitic Stainless Steel
Austenitic High Nitrogen Steels
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Commercial Low-Nickel and High Nitrogen Steels Markus O. Speidel, and Hannes J. Speidel Swiss Academy of Materials Science
ABSTRACT. Developments and applications of commercial stainless steels with low nickel and high nitrogen contents are presented and discussed. Austenitic stainless steels with four weight-percent nickel have a good chance to economically replace Type 304 austenitic stainless steel, particularly if the requirements on corrosion resistance are moderate. Austenitic stainless steels with one weight percent nickel can be economically useful if either the requirements on deepdrawability are moderate or high strength and good ductility are required. Most of the newer ferritic stainless steels and most of the low nickel austenitic stainless steels are not suitable for applications in marine atmosphere. KEY WORDS high nitrogen steels, low nickel, austenite
1. Composition and Microstructure of Stainless Steels
For a steel to be stainless, a minimum of about 12 weight-percent chromium content is necessary. This will ensure that the steel will not visibly rust in a few years time in a very mild atmosphere such as indoors in a temperate climate. In Figure 1, such steels are represented by the commercial grades 409 and 410, and it is also seen that their microstructure is ferritic, i.e. they have a body centered cubic crystal lattice.
Fig. 1. Crystal structures of commercial stainless steels as a function of the nickel equivalent and the chromium equivalent.
It is easily possible to increase the corrosion resistance of such ferritic stainless steels by increasing their chromium content and/or their molybdenum content and this has been done in the 400 series commercial stainless steels as also seen in Figure 1. A recent development, for example, is 443 ferritic stainless steel with 21 weight- percent chromium. This highly corrosion resistant and reasonably ductile material (Ref.1) consists essentially only of iron and chromium and thus avoids
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Austenitic High Nitrogen Steels
the two most expensive alloying additions which we find often in other types of stainless steels, namely nickel and molybdenum, Figure 2.
Fig. 2. Nickel and molybdenum prices in recent years.
As ferritic stainless steels of the 400 series can thus deliver a specific corrosion resistance at the lowest possible cost of all stainless steel types, the question may be asked why we should add costly nickel or its equivalent alloying additions to obtain duplex or austenitic microstructures as shown in Figure 1. There are at least five practical reasons for this: ductility, strength, weldability, deepdrawability and ferromagnetism. 2. Ductility and Strength of Stainless Steels
Yield strength, ultimate tensile strength and elongation to fracture (ductility) of commercial austenitic stainless steels are presented in Figure 3, (Ref 2).
Fig. 3. Effect of interstitials on strength and ductility in austenitic stainless steels.
numerous
The austenitic structure and composition of those steels are shown in Figure 1. It is generally agreed that, compared to the strong effect of carbon and nitrogen in solid solution on yield strength and tensile strength , the effect of substitutional elements like Cr, Ni, Mn and Mo is negligible. The point to be made here with Figure 3, however, is that the elongation to fracture (A) remains fairly constant, independent of the C+N concentration and independent of the consequently strong differences in yield strength and tensile strength. The scatter of the ductility values (A) seen in Figure 3 is due to the fact that the data reported there are individual test results from many different commercial alloys with different grain sizes and different impurity levels. Nevertheless, these actual test data scatter around an average of 55 percent elongation to fracture. For a more fundamental, systematic investigation of the mechanical properties it is indispensable to evaluate additionally the effect of grain size, Ref.3), Ref.4). In contrast to Figure 3 with ACTUAL individual test data, Figure 4 presents the published MINIMUM elongation to fracture of commercial stainless steels.
Fig. 4. Minimum elongation to fracture (ductility) of ferritic and austenitic stainless steels.
It is quite apparent that ferritic stainless steels have only half the ductility of austenitic stainless steels. This is the basic drawback of the body-centered cubic crystal lattice of ferritics with its lower number of slip systems compared to the face centered cubic crystal lattice of the austenites. The higher ductility of the austenitic steels and the lower ductility of the ferritic steels is representative of the two distinct groups of materials. Ductility depends on the microstructure and not on the content of the major alloying elements as shown in Figures 3 and 4. Here we have a major reason why alloying with the expensive nickel-equivalent to reach the austenite region in Figure 1 can be justified, because in this way we can control the essential properties of strength and ductility, according to Figures 3 and 4. 3. Applications Requiring Low Strength and High Ductility
Deep drawn components account for a very significant part of the overall use of stainless steel sheet in the world today. Examples include many appliances used in kitchens, such as sinks, refrigerators,
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Austenitic High Nitrogen Steels
dishwashers, washing machines and so on. Some of these parts can be made of ferritic stainless steels and thus relatively cheap, but whenever intricately shaped parts and complicated deep drawing operations are required, there is the chance that austenitic stainless steels are the material of choice because of their superior ductility. Materials of excellent deep-drawability should fulfill, among many others, the following four prerequisites: a) low strength (low hardness), resulting in low drawing forces, b) low coefficient of work hardening, again resulting in low drawing forces c) high ductility to avoid the formation of cracks during the drawing operation d) high austenite stability to avoid delayed cracking which can result from transformation martensite. All these requirements can be fulfilled by standard austenites of the 300 series, such as Type 304 or 316, especially if their nickel content is on the high side ( say, 304 with 9 weight-percent nickel) and their interstitial content is on the low side. The excellent ductility of 300 series stainless steels at low strength is also obvious in Figure 5.
Fig. 5. Elongation to fracture and strength-to-density ratio of different classes of structural materials.
The additional requirement of austenite stability ( prevention of excessive formation of transformation martensite during deep drawing) can be achieved, in line with Figure 1, when both, nickel-equivalent and chromium equivalent are reasonably high. Such alloys, while still widely used, however are expensive, according to Figure 2, particularly because of their nickel content of at least 8 weight-percent. The requirement of a low work hardening coefficient is achieved in two ways, namely the prevention of excessive transformation martensite formation and a low interstitial content. The latter is demonstrated in Figure 6, where it is seen that lower nitrogen contents in solid solution result not only in lower strength but also in lower work hardening. The high and volatile price of nickel, according to Figure 2, has been a powerful driving force for the development of those 200 series stainless steels which are intended for deep drawing. Two groups of materials have evolved: One group with about four percent nickel and one with about one percent nickel or less. Here we see conflicting aims at work. On the one hand is the economic driving force which would push the nickel content below one percent. On the other hand, replacing the nickel by the other elements of the nickel equivalent in Figure 1, in order to place the overall alloy composition above the austenite line generally includes a higher amount of nitrogen ( or interstitials) and this results in unwanted higher strength.
Fig. 6. Effect of nitrogen on strength and work hardening in austenitic stainless steels.
One escape from this problem is to reduce the chromium equivalent, because this results in a lower nickel equivalent needed to stay above the austenite line. The negative consequences of such a move are a reduction of the corrosion resistance, as indicated by the Figures 7 and 8, as well as a reduction of the austenite stability or resistance to transformation martensite.
Fig. 7. Effect of alloying elements on the corrosion resistance of austenitic
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Austenitic High Nitrogen Steels
Figure 8. Weight loss due to corrosion upon exposure to marine atmosphere.
The net result is at this time the existence of two groups of commercial stainless steels of the 200 series. The “ 1 percent Ni group” e.g. alloys J4 and 204Cu with about 0.18% N and the “4 percent Ni group” ,e.g. alloys J1 and 201 with about 0.08 to 0.13 %N. The “1 percent Ni group” is more economic but has less deep drawability and less corrosion resistance, the “4 percent Ni group” has higher cost, but better deep drawability and better corrosion resistance. Members of both groups may contain up to 2% Cu to stabilize against transformation martensite and thus keep the work hardening coefficient low. It is apparent from Figure 5 that at the lower strength end, the 200 series stainless steels with 4% Ni approach the low strength and the high ductility of the 300 series stainless steels with 8% Ni. This has been a successful alloy development, carried out by steel makers on several continents around the world: cost-effective 200 series stainless steel sheets for deep drawing. 4. Applications Requiring High Strength
Both, Figure 3 and Figure 5 indicate the wide range of strength levels which can be achieved with nitrogen-containing austenitic stainless steels. For comparison, Figure 5 also contains ductility and strength data for carbon steels, Al-, Mg- and Ti- alloys, Ref.5). We have already discussed above the similar strength and ductility of 300 series and 200 series stainless steels at he low strength end. We shall now discuss the competing carbon steels and light alloys before focusing on high strength stainless steels. The car body steels listed in Figure 5 include all modern types of ferritic and quenched and tempered carbon and low alloy steels, which are presently used or considered for cars and thus represent a huge volume of annual steel consumption. Their ductility-strength combination lies below that of all other material types which will compete in the future for applications in light-weight, fuel-efficient cars. For crash-worthiness of cars, a minimum ductility of their structural materials is stipulated, say, no less than 20 percent elongation to fracture. According to Figure 5, this limits the strength-to-density ratio achievable with ferritic steels to less than 100. At the same level of ductility, Al-, Mg-, and Ti-alloys can achieve a strength-to density ratio of about 150, and this is the reason, why some of these materials have replaced carbon steels in modern cars. Austenitic stainless steels, however, particularly those with high nitrogen content can
be far superior to all other materials as shown by the line labeled 200 in Figure 5. Herein lies a huge potential for future applications of austenitic stainless steels. We have in the past already reported about the highest strength-to-density values achievable with forgings and plate materials, Ref.5). These data are marked by the open circles and open squares in Figure 5. We shall therefore concentrate here and now on materials which have been produced in sheet form , primarily the 200 series of austenitic stainless steels. The corresponding line in Figure 5 indicates what is possible today and what may be the goal for tomorrow. Already there is a fair number of commercial austenitic stainless steels available in sheet form with ductilities around 50 percent elongation to fracture and strength to density larger than 100. If the future development of this class of materials parallels that which has already been achieved with forged materials, (Figure 5), the following combination of properties can be envisioned as an extrapolation of the present-day 200 series : 30% elongation to fracture at a strength-to-density ratio of 150. This would deliver, at equal ductility, three times the strength of presently used car steels with that same ductility. The higher cost of the stainless austenite, compared to the presently used carbon steel, would be set off in three ways: a) lower material consumption because of higher strength b) lower fuel consumption because of lower weight c) higher scrap (recycling) value because of the more valuable elements contained. At the time of this writing, the worlds automotive industry is in deep economic trouble. Many companies and governments hope for better positions of their car industries through the development of lighter, more fuel efficient cars. Here is a chance to develop at the same time a better material to achieve this goal: nitrogen containing high-strength austenitic sheet material. If materials with a high strength-to-density ratio are important for the car industry, they are even more important for the aerospace industry. Again, Figure 5 may be seen as a guide for future structural materials of the aerospace industry. The same holds true for shipbuilding , military applications, architecture and others. 5. Corrosion Resistance
It is well known that higher Cr, Mo and N concentrations in solid solution improve the corrosion resistance of stainless steels. This is demonstrated in Figure 7 where the corrosion resistance in concentrated NaCl solution is measured by the critical pitting corrosion temperature. The data in Figure 7 give a quantitative relation between the corrosion resistance and the alloy content. The lowest point on the curve in Figure 7 corresponds to Type 304 stainless steel, and lower critical pitting corrosion temperatures cannot be measured since the test solution would then freeze. This is regrettable since many recently developed stainless steels have just that much lower corrosion resistance. On the other hand, accelerated laboratory corrosion test results are often difficult to translate into predictions of service behavior. We have therefore several years ago started an extensive investigation to measure and to compare newer stainless steels with respect to their relative corrosion resistance in various relevant atmospheres, such as indoors, outdoors in central Europe and in marine atmosphere at the coast of the Mediterranean sea. The tests are now running for the second year and each coming year the results will be more reliable, because the scatter will be reduced. The steels investigated include the majority of those listed in Figures 1 and 3. The results for the first year are presented in Figure 8. Both, the reliability of the test results and the number of alloys involved will increase in the years to come, but already after the first full year a number of significant conclusions can be drawn: a) alloys 316L, 216L, LDX, 317 and 2205 exhibit NO measurable weight loss.
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Austenitic High Nitrogen Steels
They can be considered resistant to corrosion in the marine atmosphere close to the Mediterranean sea. b) alloys 304 and 301 are NOT corrosion resistant in the marine atmosphere. They develop visible rust and a measurable weight loss. c) all other steels measured so far (and marked in Figure 8) are less corrosion resistant than 304. This is true not only for the ferritic steels mentioned, but also for the more economic “4% nickel “ and the “1% nickel “ austenitic stainless steels containing nitrogen. Thus, neither 304 stainless steel nor its more economic relatives should be used in outdoors atmospheres near the sea. In contrast, all these steels have been free of staining or rust after two years exposure to indoors atmosphere in central Europe. In that sense they are truly “stainless”. All of the above remarks should be tempered by the fact that corrosion resistance is not only dependent on alloy composition but also on the homogeneity of the microstructure and on the surface condition. Thus, while Type 316 stainless steel, according to Figure 8, can be perfectly resistant to corrosion in marine atmosphere, it can also develop a rusty surface and a measurable weight loss in that same atmosphere, if it has a rough or contaminated surface. This is represented in Figure 8 by the one data point (without label), just above the data point for steel 316. 6. Summary
Austenitic stainless steels containing 8 to 10 weight-percent nickel have been in commercial use since almost 100 years. In order to save the cost of nickel, they are now partially substituted by ferritic stainless steels and by austenitic stainless steels containing reduced nickel content and increased nitrogen content. Relative advantages and disadvantages of these steels are demonstrated here. The huge future potential of high-nitrogen, low-nickel austenitic stainless steels is based not only on its relatively low cost but also on its excellent combination of strength and ductility. References 1. Dong Wenbo, Ma Li, Jiang Laizhu, Ou Xiangbo and Lu Weijiang. Development of corrosion-resistant ferritic stainless steel B443NT. The Third Baosteel Biennial Academic Conference, 2008, F68- F72. 2. Markus O. Speidel and Hannes J.C.Speidel. Development of new austenitic stainless steels. The Third Baosteel Biennial Academic Conference, 2008, F1-F4. 3. Markus O. Speidel and Mingling Zheng-Cui. High-Nitrogen Austenitic Stainless Steels. HNS Conference Proceedings. Zurich: VDF Hochschulverlag ETH Zürich, 2003, 63–73. 4. Markus O. Speidel. Nanograin size, high-nitrogen austenitic stainless steels. Zeitschrift Metallkunde, 2003,(94): 719-722. 5. Markus O. Speidel and Hannes J. C. Speidel. Austenitic stainless steels of high strength and ductility. Zeitschrift Metallkunde, 2004(95): 596-600.
High Interstitial Stainless Austenitic Steels, Part I: Constitution, Heat Treatment, Properties, Applications
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High Interstitial Stainless Austenitic Steels, Part I: Constitution, Heat Treatment, Properties, Applications Hans Berns1, Sascha Riedner1, Valentin Gavriljuk2 1
2
Ruhr University Bochum, Germany, Institute for Metal Physics, Kiev, Ukraine
Abstract: Austenitic steels with 18 to 19 mass% Cr and Mn each and 0.85 to 1.1 mass% C + N cold work harden from a 0.2 % proof strength of ≈ 600 MPa to a true fracture strength of ≈ 2500 MPa at ≈ 70 % elongation. Their impact wear resistance equals that of Hadfield manganese steel and their corrosion resistance is improved by adding Mo. The new grades were melted under normal pressure of air. Harmful precipitation during quenching from solution anneal temperature is subdued by optimizing the C/N ratio. Some applications are discussed. Keywords: Strength, toughness, wear, corrosion, nonmagnetic
1. Introduction
Austenitic high nitrogen steels with 0.8 to 1 mass% N produced by pressure metallurgy have not experienced much of a market pull in spite of excellent properties [1, 2]. The powder metallurgical production of respective CrMn grades was accompanied by detrimental surface layers on the atomized powder grains [3]. However, by raising the CrMn content a steel with (mass%) Cr21Mn23Ni1.5N0.85 was manufactured by standard ingot metallurgy e.g. for application in drill collars [4]. Replacing some of the high nitrogen content by carbon allowed to produce leaner stainless austenitic CrMn steels by ingot metallurgy, e.g. Cr15Mn17C0.39N0.43 [5, 6]. It was shown [7] that the addition of up to 0.4 mass% C to steels with 0.5 mass% N improved the resistance to pitting corrosion. It is the aim of the present study to further increase the strength of stainless austenitic steels by raising the interstitial content from 0.85 to 1.1 mass% and adjust the C/N ratio so that melting under atmospheric pressure is feasible. Doing without pressure metallurgy and nickel is supposed to reduce costs and thus add to the market push. 2. Methods of investigation
The constitution was analyzed by thermodynamic calculations along the ThermoCalc program [8] and the TCFE4 database. The results led to the steels CN0.85, CN0.96, CN1.07 and respective grades alloyed with copper and molybdenum to enhance the corrosion resistance (Table 1). All steels were molten in air of normal pressure, CN0.96 and CN1.07 electro-slag-remelted (ESR), followed by hot working. Specimens were taken in parallel to the direction of hot working by electro-spark-erosion, quenched from solution anneal temperature TSA (Table 1) in water and ground to size. Specimens of Ø 4x10 mm were gas cooled at pre-selected rates expressed by the t10/7 cooling time between 1000 and 700°C to derive the critical tc10/7 cooling time at which grain boundary precipitation becomes detectable by intercrystalline corrosion tests according to Strauss and EN ISO 3651-2-A (Fig.1). The cooled specimens were ground to 4x1x10mm, etched, bend to 90° and inspected for cracks at 10x magnification. Tensile tests were performed along EN 10002 with Ø 5 mm specimens and ISO-V notch impact bending tests along EN 10045 using a 300 or 450 J hammer. To measure impact wear two wear plates 50x35x10 mm were mounted on a rotor in adjacent positions and impacted by mineral particles of greywacke of hardness 760 HV0.1 and grid size 11 to 8 mm, which fell slowly, at a rate of about 1/s and individually counted by a light barrier,
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Austenitic High Nitrogen Steels
parallel to the vertical rotor axis. They were hit by the plates rotating under an impact angle of 90° at a velocity of 25 or 37 m/s. The mass loss of the two plates was measured after every 1000 impacts and cleansing in an ultrasonic bath of alcohol. The hardness penetration below the wear surface and the hardness on the fracture of tensile specimens were derived by the Vickers test under a load of 100 g (HV0.1). The relative magnetic permeability µrel of the plates was measured by a Förster Magnetoscop 1.067-103 with the handheld sensor 1.005-1532. The fatigue resistance was measured up to N=107 cycles of four-point rotating bending at 40 Hz using hourglass-shaped specimens of Ø 5 mm waist, ground in axial direction. Light optical microscopy (LOM) and scanning electron microscopy (SEM) with electron backscattered diffraction (EBSD) were used to visualize the appearance of the fracture and wear surfaces and subsurface zones. Current density/potential tests along ASTM-G5 were carried out to describe the corrosive properties in aqueous solutions. The potential is that of the standard hydrogen electrode (SHE). Table 1 Chemical composition and solution anneal temperature TSA
3. Results
steel
CN1.07
CN0.96
CN0.85
3.1. Constitution
0.489 0.578 1.07 0.846 18.8 18.9 0.43 1175
0.344 0.614 0.96 0.544 18.2 18.9 0.30 1100
0.256 0.596 0.85 0.430 18.3 18.4 0.54 1100
At first the conditions for homogenous austenite at a suitable TSA of e.g. < 1150°C were evaluated by means of isothermal phase diagrams plotting the Cr content in dependence of the Mn content for fixed C+N contents and C/N ratios. In addition isothermal plots of Cr over C+N at 18 mass% Mn revealed that 1 mass% C+N and 18 mass% Cr are soluble at C/N=0.5 and 1100°C (Fig. 2). However, a ratio of C/N=0.7 would require an increase of TSA to prevent M23C6 carbides. At 18 mass% Cr and Mn each no range of homogenous austenite exists, if carbon is added solely (Fig. 3, left side), but a wide range opens in pure nitrogen steels (bottom). Joint alloying with C+N enlarges, i.e. stabilises, the austenite phase field which is encased by M23C6, ferrite and M2N. Judging from the inclination of the respective borderlines, ferrite is more subdued by N than by C, even in atom%. M23C6 is subdued by N while the M2N field is expanded by C. The background of austenite stability is discussed in chapter 3 of Part II.
1)
mass% C N C+N C/N Cr Mn Si TSA[°C] 1)
P: 0.017-0.026, S50 000m3), etc. Considerable quantities were exported to Russia, Greece and NATO. The aforementioned complex of properties provided the grounds for the construction of the “Russia – Bulgaria South” main gas pipeline and a significant part of the internal gas conduit network. External coating was applied. The testing of time since 1974 till present days gives a good testimony to these materials. 2.2. Nitrogen stainless steels (0.2%N) Results from the investigation [14] are shown (Table 2) for the laboratory nitrogen steel NS brand 04Cr16Ni4N0.2 of the martensite or austenite-martensite class, obtained under atmospheric conditions. This NS is not sensitive to the cooling rate after rolling or thermal treatment and combines high strength with high plasticity and fracture toughness. When the steel is in the most strengthened state, its yield strength is 1300МРа, the tensile strength is up to 1700МРа, the cross section contraction is not lower than 60% with 5-6 times higher fracture toughness than that of the X60 black steel. The welded connections obtained by argon-arc, manual or automatic welding exhibit high quality. The authors point out that the susceptibility to hydrogen brittleness and the corrosion resistance of this steel in aqueous solutions (containing hydrogen sulphide) are not sufficiently well studied and that detailed economic assessment and complex experiments under typical conditions for main pipelines in marine environment are necessary. 2.3. Stainless high nitrogen austenitic nickel-free steels. These steels need specialized furnaces. This preliminary study was aimed at obtaining information about the basic characteristics of steels and about the relevance of the research problem. We tried to develop a stable mono-phase austenitic structure on the basis of the following principal considerations: high plasticity and toughness, combined with big difference between the plasticity and impact elasticity values, maximum coefficient of deformation strengthening compared to the other phase containing steels, thus providing the possibility of ignoring certain pipe deformations resulting from internal or external pressure, which increases the risk of immediate or postponed failure, etc. The favorable aspects of this investigation in Bulgaria are that for the present the only special furnaces under pressure in the world and technologies for high nitrogen stainless austenitic steels have been developed in the country [15, 16] and they solve two basic problems: - sufficiently high concentration of nitrogen for ensuring entirely the austenitic structure, without alloying with nickel and practically without introducing carbon and manganese, i.e. alloying only with nitrogen; - metallurgical production of high nitrogen sheet ingots as initial billets for the production of sheets for welded main pipes.
Corrosion and Surface Treatment
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1800 1600 1400 10Mn2SiVN0.02
1200
X70
1000
12Cr18Ni10Ti
800
04Cr23N1.2
600
04Cr16Ni4N0.2
400 200 0 A, %
Re, Mpa
Pm, Mpa
Fig. 3 Basic mechanical characteristics of the conventional black steel (X70) and stainless steel alloyed with nitrogen.
Using the proved possibilities of the Big Steelmaking Bath (BSB) method [1517], we have manufactured laboratory sheet ingots with weight of 8 kg, which were further on subjected to hot deformation to a thickness from 40 mm to 12 mm and to initial testing. It is seen in Table 2 and Fig. 3 that the basic mechanical characteristics of the 04Cr23N1.2 high nitrogen steel meet with a reserve the demands according to API 5LX towards its analogue with the highest requirements X70.
HNS 04Cr23N1.2 has proved high characteristics: Rm=1250MPa, Re=815MPa, A=51% [16]. The corrosion resistance is high, compared to the classical nickel stainless steel 04Cr18Ni10Ti, as well as one order higher corrosion resistance compared to that of black steel of the X70 type and comparable technological qualities in hot and cold state. There are publications [12, 13, 17] indicating that alloyed stainless chromium-manganese austenitic steels with high nitrogen concentration (0.6–0.8%) impart positive effect to low-cycle fatigue and to a number of other characteristics. These HNS possess high corrosion resistance are also technological for metallurgical production and mechanical processing. The obtained high parameters are explained mainly with the successful replacement of classical nickel austenite in stainless steels by a new (nitrogen) austenite. After quenching nitrogen is transferred entirely to the solid solution and is not released under conditions of the common technologies applied in metallurgy and machine building [14, 16, 17]. In the scale of the obtained under laboratory conditions sheet thickness of 12 mm, and with the high mechanical characteristics of the new steel, it was not possible to make correct experimental determination of K1С. At the same time, in our opinion, the programs available in reference literature for evaluating the resources of main pipelines contain two essential shortcomings: - the programs do not take under consideration the new and most important factor – the resistance of steel against corrosion, which can increase the operation period in the order of times. - the programs consider that the steel pipelines possess lower characteristics Re
