Plansee Seminar 2 0 0 1

AT0200466

15th INTERNATIONAL PLANSEE SEMINAR 2 0 0 1

Powder Metallurgical High Performance Materials

Proceedings Volume 4: Late Papers .33/46

Editors: Günter Kneringer, Peter Rödhammer and Heiko Wildner PLANSEE The authors themselves are responsible for the contents of their paper. All rights reserved by Plansee Holding AG, Reutte, Tyrol, Austria Copyright 2001 by Plansee Holding AG, Reutte, Tyrol, Austria Printed by RWF, Werbegesellschaft m.b.H., Wartens, Tyrol, Austria Preface

On the eve of the 15th Plansee Seminar we look back on half a century of this conference series founded by Professor Paul Schwarzkopf in 1951. The Proceedings of these Seminars form an impressive chronicle of the continued progress in the understanding of refractory metals and cemented carbides and in their manufacture and application. There the ingenuity and assiduous work of thousands of scientists and engineers striving for progress in the field of powder metallurgy is documented in more than 2000 contributions covering some 30 000 pages.

As we enter the new millennium the basic needs of the markets we are serving remain the same: improved material properties and designs translating into improved product performance and reliability. In view of ever shorter product life cycles, however, the pressure on material scientists and engineers to introduce these new products at reduced time-to-market and with improved benefit/cost ratio is continuously mounting.

Since 1951 the means to meet these demands have greatly improved: novel equipment with improved process control, faster computers for simulation as well as for data acquisition and reduction, analytical tools with undreamed-of accuracy; all coupled with the shrinkage of distances brought about by the revolution in global communication.

The 15th Plansee Seminar was convened under the general theme "Powder Metallurgical High Performance Materials". Under this broadened perspective the Seminar will strive to look beyond the refractory metals and cemented carbides, which remain at its focus, to novel classes of materials, such as intermetallic compounds, with potential for high temperature applications.

It is with great pleasure that I welcome you in the name of Plansee Holding AG here in Reutte/Tyrol and extend to you my sincere wishes for a fruitful, inspiring seminar. Thank you in advance for your personal contribution to making the 15th Plansee Seminar a successful, memorable event.

Reutte, May 2001 Fc r the Executive Board

Dr. Michael Schwarzkopf Organisation

Organized by: Plansee Holding AG, Reutte, Tyrol, Austria

International Liaison Board:

Belgium 0. Van der Biest, Heverlee Brazil H. R. Z. Sandim, Lorena Peoples Republic of China G. Shiju, Beijing W. Yin, Baoji Czech Republic J. Späcil, Sumperk Germany H. E. Exner, Darmstadt D. Stöver, Jülich Great Britain M. Loretto, Birmingham B. Roebuck, Teddington Hungary L. Bartha, Budapest India P. Ramakrishnan, Bombay Israel A. Rosen, Haifa Italy M. Rosso, Torino Japan R. Watanabe, Sendai Liechtenstein W. J. Huppmann, Schaan The Netherlands J. Bressers, Petten F. Mertens, Maarheeze Poland S. Stolarz, Gliwice Russian Federation V. 1. Trefilov, Kiev South Africa S. B. Luyckx, Johannesburg South Korea 1. H. Moon, Seoul Sweden B. Uhrenius, Stockholm USA A. M. Alper, Brant Rock R. M. German, University Park B. North, Latrobe EPMA Europe L. Albano-Müller, Shrewsbury MPIF USA T.E. Friday, Princeton

Scientific Committee: B. Kieback, Dresden G. Kneringer, Reutte G. Leichtfried, Reutte P. Rödhammer, Reutte W. Schintlmeister, Reutte U. Schleinkofer, Reutte W. D. Schubert, Wien R. F. Singer, Erlangen

Local Organization: G. Kneringer P. Rödhammer H. Wildner PLEASE BE AWARE THAT ALL OF THE MISSING PAGES IN THIS DOCUMENT WERE ORIGINALLY BLANK Powder Metallurgical Hiqh Performance Materials VII

Contents

Volume 1: High Performance P/M Metals

Power Engineering Heat Management Novel Refractory Metals Alloys Lighting Technology Transportation

Volume 2: P/M Hard Materials

Ultrafine-grained and Nanocrystalline Hard Materials Hard Metals with Optimized Properties Novel Concepts for Hard Materials Cutting Tools for Extreme Conditions

Volume 3: General Topics

Modelling and Simulation Manufacturing Technologies Special P/M Materials

Volume 4: Post-Deadline Papers

High Performance P/M Metals P/M Hard Materials General Topics Powder Metallurgical High Performance Materials IX

Eröffnungsrede 15. Plansee Seminar

Meine sehr geehrten Damen und Herren, Herzlich willkommen in Reutte und ich freue mich, Sie heute morgen im Namen des Vorstandes der Plansee Holding AG zum 15. Internationalen Plansee Seminar in unserem Haus begrüßen zu dürfen. Wir fühlen uns sehr geehrt, dass sich rund 500 Teilnehmer aus ca. 40 Ländern für diese Konferenz angemeldet haben und ich darf Ihnen allen aufrichtig danken, dass Ihre Wahl auf unser Seminar gefallen ist. Die Plansee Seminare haben eine lange Tradition in der Pulvermetallurgie von Hochschmelzenden Metallen und Hartmetallen und wir waren auch diesmal wiederum sehr bemüht, Ihnen eine Woche mit interessanten Vorträgen und Posterpräsentationen zu diesen Werkstoffgruppen bieten zu können. Die Vorbereitungen für das 15. Seminar haben vor mehr als zwei Jahren begonnen und ich möchte es an dieser Stelle nicht verabsäumen, allen Beteiligten, die ihr Bestes für eine hoffentlich wiederum erfolgreiche Konferenzwoche beigetragen haben, herzlich zu danken. Ohne dem wissenschaftlichen Komitee, dem internationalen Liaison Board und vielen, an dieser Stelle ungenannten Mithelfern wäre es nicht möglich, als Privatunternehmen ein Seminar dieser Größenordnung zu organisieren. Meine Damen und Herren, seit einem halben Jahrhundert hat die Plansee Gruppe erst alle drei, und später alle vier Jahre diese Woche der Begegnung zwischen Wissenschaft und Industrie veranstaltet. Es war nie unser vor- rangiges Ziel, durch diese Seminarwoche das Image und den Bekanntheitsgrad der Unternehmensgruppe zu verbessern, sondern in erster Linie einen entscheidenden Beitrag zum technischen und wirtschaftlichen Fortschritt - durch den Gedankenaustausch und die Zusammenarbeit auf dem Gebiet der Pulvermetallurgie von Hochleistungswerkstoffen - zu leisten. Dieses Seminar sollte aber aus unserer Sicht auch eine Plattform sein, um sich mit den Besten der Welt aus technologischer und innovativer Sicht zu messen. So hoffe ich, dass diese Woche für uns alle ein intellektueller Gewinn sein wird - für Sie als Teilnehmer, für das Organisationsteam - und finanziell hoffentlich auch für die Region Reutte. Wir sind uns aber auch bewusst, dass das Plansee Seminar im weltweiten Wettbewerb mit anderen, äußerst professionell durchgeführten Konferenzen - teilweise sogar mit Meerblick - steht. Von Anfang an hat sich unser Seminar deshalb auf einen kleinen Ausschnitt der pulvermetallurgischen Hochleistungswerkstoffe konzentriert, um sich durch diese Fokussierung von anderen Veranstaltungen zu differenzieren. Unser erklärtes Ziel ist, das Powder Metallurgical High Performance Materials

Plansee Seminar als die weltweit führende, international ausgerichtete und den gesamten Innovationszyklus einbindende Fachtagung für Hochschmelzende Metalle, Hartmetalle und andere neue Hochleistungs- werkstoffe zu positionieren. Haben wir dieses Ziel schon erreicht? Da wir Ihre Meinung dazu hören möchten, haben wir uns erlaubt, Ihren Tagungsunterlagen einen Fragebogen beizulegen. Ich darf Sie heute schon bitten, uns mit ein paar Minuten Ihrer Zeit eine bessere Beurteilung unserer Seminarleistung zu ermöglichen.

Meine sehr geehrten Damen und Herren, Erlauben Sie mir einige Anmerkungen zu den für diese Woche im Mittelpunkt stehenden Werkstoffgruppen. Diese schon in der Reifephase positionierten Materialbereiche haben in den letzten Jahren einen m.E. überraschend erfreulichen Wachstumskurs genommen. Doch nicht nur die stabile weltweite Konjunkturlage, die Globalisierung und damit der Zugang zu neuen Regionalmärkten haben diesen Aufschwung bewirkt. Speziell die Kundenforderungen nach höherer Leistung bei gleichzeitig nachhaltiger Verbesserung der Umwelt-, Energie- und Rohstoffbilanz waren die entscheidenden Innovations- und Wachstumstreiber der letzten Jahre. Nehmen Sie als Beispiel • die Miniaturisierung in der Lichttechnik, • die Anforderungen an Beschichtungsmaterialien oder Leiterplattenbohrer im Umfeld der Elektronik, • die Herausforderungen in der Zerspanung von neuen Leichtmetallen für die Automobilindustrie verbunden mit den ökologischen Auflagen wie die Trockenbearbeitung oder • generell die beschleunigte Substitution von Schnellstahl durch Hartmetall. Nur einige Beispiele die zeigen, dass die industriellen Chancen für diese Werkstoffgruppen nach wie vor gegeben sind. Ein weiterer Trend, der auch durch die Zunahme von Seminarbeiträgen zur numerischen Simulation und Modellierung bestätigt wird, verdeutlicht die generellen Bemühungen zur Optimierung der pulvermetallurgischen Prozesse. Zusammen mit neuen Verfahren der Analytik, der Material- und Qualitätsprüfung sind dies unentbehrliche Werkzeuge für die Weiterentwicklung der Pulvermetallurgie von einer erfahrungsgestützten zu einer wissensgestützten Technologie. Powder Metallurgical High Performance Materials XI

Mit kritischem Blick, meine Damen und Herren, und auch mit Sorge betrachte ich aber die Veränderungen auf dem industriellen Anbieterfeld. Die zunehmende Öffnung der globalen Märkte, verbunden mit der nach wie vor hohen Attraktivität dieser Werkstoffgruppen hat zu einer weiteren Konzentration der Marktteilnehmer, zum Auftritt neuer Lieferanten mit teilweise äußerst aggressivem Marktverhalten und zu einer zumindest zur Zeit instabilen Angebot/Nachfrage-Situation, speziell für Wolframrohstoffe, geführt. Diese veränderten Marktverhältnisse erfordern eine hohe Disziplin von allen Beteiligten, um auch zukünftig die kundenseitige Akzeptanz für diese Werkstoffe zu gewährleisten.

Meine sehr geehrten Damen und Herren, Die Schwerpunkte des 15. Plansee Seminars kreisen traditionsgemäß um Hartmetalle, Hartstoffe und Refraktärmetalle. Dennoch haben wir das Gesamtthema des 15. Plansee Seminares bewusst „Powder Metallurgical High Performance Materials" getauft, um anderen Hochleistungswerkstoffen wie z.B. intermetallischen Verbindungen, die in klassische Anwendungs- bereiche der hochschmelzenden Metalle vordringen, einen Platz einzuräumen. Nur durch die umfassende Betrachtung des konkurrierenden Werkstoffumfeldes, der Darstellung des Innovationszyklus von der Grundlagenforschung bis zur industriellen Anwendung und der gesamten verfahrenstechnischen Prozesskette können schlussendlich die Zukunftschancen dieser pulvermetallurgischen Hochleistungswerkstoffe beurteilt werden. Ich hoffe, dass wir mit den ausgewählten Schwerpunkten für diese Seminarwoche auch Ihre Interessensgebiete abdecken werden. Lassen Sie mich schließen und Ihnen nochmals sehr herzlich für Ihr Interesse am 15. Internationalen Piansee Seminar danken. Ich wünsche dieser Konferenz einen guten Verlauf und Ihnen, meine sehr geehrten Damen und Herren, einen erfolgreichen und angenehmen Aufenthalt in Reutte.

Dr. Michael Schwarzkopf Powder Metallurgical High Performance Materials XIII

Entering the Walter-Schwarzkopf-Haus ...

... to listen to the Welcome Address ... XIV Powder Metallurgical High Performance Materials

... delivered by Dr. Michael Schwarzkopf.

"Is nuclear fusion an answer to global warming?" Critical questions to Prof. Karl Lackner after his Opening Lecture. Powder Metallurgical High Performance Materials XV

Chairmen opening a discussion ...

... to be intensively continued during Coffee Break. XVI Powder Metallurgical High Performance Materials

Refreshing the mind outdoors ...

... and the body at the Snack Bar. Powder Metallurgical High Performance Materials XVII

Any questions? Address the team at the Seminar Office!

Getting seated for a Session XVIII Powder Metallurgical High Performance Materials

Discussions during a Poster Session ...

... and explanations by hand-waving. Powder Metallurgical High Performance Materials XIX

Plansee Dinner, Show & more: Enjoying the dinner

.. to the sounds of the Plansee Big Band. XX Powder Metallurgical High Performance Materials

Closing Remarks 15th International Plansee Seminar

Ladies and Gentlemen, dear friends and colleagues!

Almost 50 years ago Prof. Dr. Paul Schwarzkopf defined the aims of the Plansee Seminar as follows:

• Furtherance of closer cooperation between scientists and engineers • Creation and renewal of personal contacts between powder metallurgists and scientists as well as engineers working in related fields • Exchange of experience among scientific communities working independently in different countries

Looking back to the last five days of communication, discussions and presentations, we can surely recognize that these guiding principles are still valid. The world of academy and industry, the world of the customer in all varieties from the very beginning of the value added chain to the final application has been growing together. Partnership -total and global- has been turning out to be the key also to the materials related challenges of the future, and we all here are the proof for this.

In all 242 presentations have been accepted. Thereof remarkable 50 per cent can be classified as partnership papers: 4 per cent within industrial partners, 20 per cent within academic institutions and 26 per cent within both worlds. In other words every second paper presented at this 15th Plansee Seminar has been based on joint development work as impressive examples of a living task sharing. Powder Metallurgical High Performance Materials XX[

Keeping in mind this growing network of international and interdisciplinary partnership and the high quality of science and technology of the institutions which you are representing we are fully justified in saying that the world of PM high performance materials is in good shape and well prepared to meet the requirements of the next key technologies and markets.

Pleased at this situation, on behalf of the Schwarzkopf family and the executive Board of Plansee, I would like to thank all who have contributed a great deal to the success of the seminar: the authors, the keynote speakers, the session chairmen, the members of the scientific and organizing committee, the members of the liaison board and all participants.

We wish you and your company or institution all the best and great success for your projects, cooperations, and all your business activities.

Once again, thank you very much! Good bye! Auf Wiedersehen!

.-•"7

Dr. Günter Kneringer Powder Metallurgical High Performance Materials XXIII

High Performance P/M Metals

RM 1 Lackner K., Andreani R., Gasparotto M., Pick M. EFDA, Garching, Germany ITER AND THE FUSION REACTOR: STATUS AND CHALLENGE TO TECHNOLOGY 4 471

RM 4 Watanabe R., Li J.F., Saito A., Kurishita H., Igarashi T.*, Okamoto K.**, Kikuchi K.***: Furusaka M.****, Kawai M.**** Tohoku University, Sendai, Japan * Sumitomo Electric Co., Tokyo, Japan ** Allied Material Corp., Tokyo, Japan *** Japan Atomic Energy Research Institute, Tokai, Japan **** High Energy Accelerator Research Organization, Tsukuba, Japan DEVELOPMENT OF SOLID TARGET FOR SPALLATION NEUTRON SOURCE

RM7 Chen Y.L., Jones A.R., Miller U.* The University of Liverpool, Materials Science & Eng., Dept. of Engineering, Liverpool, United Kingdom * Plansee GmbH, Lechbruck, Germany CREEP RESISTANCE OF Fe-BASED ODS TUBES WITH HOOP GRAIN STRUCTURE 12

RM 8 Mahalingam M., Viswanathan V. Motorola Inc., Wireless Infrastructure Systems Division, Tempe, Az., USA MATERIALS REQUIREMENTS FOR THERMAL MANAGEMENT IN WIRELESS COMMUNICATION 4 505

RM 16 Schedler B., Huber T., Zabernig A., Rainer F., Scheiber K.H., Schedle D. Plansee AG, Reutte, Austria THE MANUFACTURE OF CARBON ARMOURED PLASMA-FACING COMPONENTS FOR FUSION DEVICES 4 26

RM 19 Franco T., Henne R., Lang M., Ruckdäschel R., Schiller G. Deutsches Zentrum für Luft- und Raumfahrt (DLR), Institut für Technische Thermodynamik, Stuttgart, Germany INVESTIGATION OF POROUS METALLIC SUBSTRATES FOR PLASMA SPRAYED THIN-FILM SOFC

RM 36 Warren J., Reznikov G.* Coolidge Laboratory and GE Aircraft Engines, Lynn, W. Milwaukee, Wl, USA * X-Ray Tube Target Plant, Warrensville Heights, OH, USA A COMPARISON OF THE MICROSTRUCTURE AND HIGH TEMPERATURE TENSILE PROPERTIES OF A NOVEL P/M Mo-Hf-Zr-Ta-C ALLOY AND TZM 4 79

RM 44 Karpov M.I., Vnukov V.l., Medved N.V., Volkov K.G., Khodoss I.I.* Solid State Physics Institute, Chernogolovka, Moscow distr., Russia * Institute of Microelectronics Technology, Chernogolovka, Moscow distr., Russia NANOLAMINATE - BULK MULTILAYERED Nb-Cu COMPOSITE: TECHNOLOGY, STRUCTURE, PROPERTIES 4 97

RM 57 Dabhade V.V., Godkhindi M.M.*, Rama Mohan T.R., Ramakrishnan P., Gutmanas E.Y.* Dept. of Metallurgical Eng. and Materials Science, Indian Institute of Technology, Bombay, India * Dept. of Metallurgical and Materials Engineering, Indian Inst. of Technology, Kharagpur, India ** Israel Institute of Technology, Haifa, Israel PROCESSING OF NANO SIZED TITANIUM NITRIDE REINFORCED TITANIUM COMPOSITES 4 37

RM 76 Harmat P., Grosz T.*, Rosta L.**, Bartha L. MFA Research Institute for Technical Physics and Materials Science, Budapest, Hungary * Chemical Research Center, Institute of Chemistry, Budapest, Hungary " KFKI SZFKI Research Institute for Solid State Physics and Optics, Budapest, Hungary A NOVEL METHOD FOR SHAPE ANALYSIS: DEFORMATION OF BUBBLES DURING WIRE DRAWING IN c; 7 r DOPED TUNGSTEN 4 108 - J o XXIV Powder Metallurgical High Performance Materials

RM 89 King H.W., Gursan S., Ferguson S., Yildiz M., Kim I.*, Nagy D.R." University of Victoria, Dept. of Mechanical Engineering, Victoria, Canada * School of Advanced Materials, Kum-Oh National Univ. of Technology, Kumi, Kyung Bug, South Korea ** Liburdi Engineering Ltd., Dundas, Canada RESIDUAL STRESS IN TiN COATINGS ON Ti-6AI-4V ALLOY 4 118

RM 98 Schalunov E.P., Dovydenkov V.A.*, Simonov V.S.* Wissenschaftlich-technische Firma TECHMA G.m.b.H., Tscheboksary, Russia * Werk der Komposit- und Metallkeramik-Werkstoffe KERMET AG, Joschkar-Ola, Russia ANWENDUNG DER HOCHEFFIZIENTEN DISPERSIONSGEHÄRTETEN WERKSTOFFE AUF PULVER- KUPFERBASIS IN DEN TEILEN VON MOTOREN UND KRAFTANLAGEN DER TRANSPORTMITTEL 4 126

RM 100 Nojkina A.V., Kolchemanov N.A., Dedkov P.Y.* VNIIALMAZ, Moscow, Russia * VNIITF, Snezhinsk, Russia SINTERING OF NANODIAMOND IN THERMODYNAMIC STABILITY AREA OF DIAMONDS 4 52

RM 102 Reglero V., Ruiz Urien I.*, Velasco T., Santos I.*, Rodrigo J.M., Zarauz J.*, Gasent J.L., Alamo ^-i f •); Chato R. --- ^ o Astronomy and Space Group, Inst, de Ciencias de los Materiales, Univ. of Valencia, Valencia, Spain * Sener, Las Arenas, Spain TUNGSTEN AND OPTICS 4 150

/ RM 103 Reglero V., Ruiz Urien I.*, Velasco T., Santos I.*, Rodrigo J., Zarauz J.*, Gasent J.L., Alamo J., •- b Chato R., Clemente G.**, Sanz-Tudanca C.**, Lopez J.L." Grupo de Astronomia y Ciencias del Espacio, Inst, de Ciencias de los Materiales, Univ. of Valencia, Valencia, Spain * Sener, Las Arenas, Spain ** Mecanizados Gines S.A., Miranda de Ebro, Spain HIGH PRECISION TUNGSTEN CUTTING FOR OPTICS 4 161 P/M Hard Materials

HM 2 Bock A., Zeiler B. Wolfram Bergbau- und Hüttengesellschaft mbH NfG KG, St. Martin, Austria ^ -i b PRODUCTION AND CHARACTERIZATION OF ULTRAFINE WC POWDERS 4 229

HM 12 Sommer M., Schubert W.D., Zobetz E.*, Warbichler P.** Institute for Chemical Technology of Inorganic Materials, Vienna University of Technology, Vienna, Austria * Institue for Mineralogy, Crystallography and Structural Chemistry, Vienna University of Technology, Vienna, Austria ** Research Institute for Electron Microscopy, Graz, Austria

ON THE FORMATION OF VERY LARGE WC CRYSTALS DURING SINTERING OF ULTRAFINE WC-Co ALLOYS 4 299

HM 33 Ryzhkin A.A., llyasov V.V., Lyulko A.V. Don State Technical University, Rostov-on-Don, Russia

c . , WEAR RESISTANCE AND ELECTRONIC STRUCTURE OF CUTTING TOOL MATERIALS ON A BASIS J jo CARBIDES OF TUNGSTEN AND TITANIUM 4 315

HM 35 Schön A., Schubert W.D., Lux B. Institute for Chemical Technology of Inorganic Materials, Vienna University of Technology, Vienna, Austria WC PLATELET-CONTAINING HARDMETALS 4 322

HM 44 Frykholm R.*, Ekroth M., Jansson B.**, Agren J., Andren H.O.* Dept. of Materials Science and Engineering, Royal Institute of Technology, Stockholm, Sweden * Dept. of Experimental Physics, Chalmers University of Techn. and Göteborg University, Göteborg, Sweden ** AB Sandvik Coromant, Stockholm, Sweden THE EFFECT OF BINDER VOLUME FRACTION ON GRADIENT SINTERING OF CEMENTED CARBIDES 4 337 Powder Metallurgical High Performance Materials XXV

HM 46 Korb G., Bulic F.I., Sajgalik P.*, Lences Z.* Austrian Research Centers Seibersdorf, Seibersdorf, Austria * Institute of Inorganic Chemistry, Slovak Academy of Science, Bratislava, Slovakia GRADIENT STRUCTURES IN SiAION'S FOR IMPROVED CUTTING PERFORMANCE 4 386

HM 49 Frage N., Bizaoui M„ Dariel M.P. Ben-Gurion University of the Negev, Dept. of Materials Engineering, Beer-Sheva, Israel THE EFFECT OF Cr2O3 ADDITIONS ON THE PRESSURELESS SINTERING OF B4C 397 S ?/

HM 56 Wollein B., Lengauer W. Institute for Chemical Technology of Inorganic Materials, Vienna University of Technology, Vienna, Austria PHASE REACTIONS IN Ti(C,N)/(Ti,W)C, Ti(C,N)/(Ti,MO)C, (Ti,W)(C,N)/Co AND (Ti,W)(C,N)/Ni DIFFUSION

COUPLES 4 411 C ] 6'

HM 57 Wittmann B., Schubert W.D., Lux B. Institute for Chemical Technology of Inorganic Materials, Vienna University of Technology, Vienna, Austria WC GRAIN GROWTH AND GRAIN GROWTH INHIBITION IN NICKEL AND IRON BINDER HARDMETALS 4 427

HM 65 Jilek M., Holubar P., Sima M. SHM Ltd., Novy Malin, Czech Republic HARD PVD COATINGS nc-fTl,-* AIJN/a-SisN« NOT ONLY FOR DRY MACHINING 4 447

HM 79 Weissenbacher R., Haubner R. Institute for Chemical Technology of Inorganic Materials, Vienna Univ. of Technology, Vienna, Austria DEPOSITION OF BCN LAYERS FROM B-C-N-H SINGLE SOURCE PRECURSORS 4 457

HM 84 Roebuck B., Gee M.G., Morrel! R. NPL Materials Centre, National Physical Laboratory, Teddington, United Kingdom HARDMETALS - MICROSTRUCTURAL DESIGN, TESTING AND PROPERTY MAPS 4 245

HM 88 Friedrichs K. Konrad Friedrichs KG, Kulmbach, Germany TRANSITION FROM HSS TOOLS TO SOLID CARBIDE TOOLS IN THE FIELD OF SHAFT TOOLS AND THE DEVELOPMENT OF NEW TYPES OF CARBIDES 4 267

HM 89 Zinyana J., Broccardo S., Hamar-Thibault S.", Cornish L.A., Witcomb M.J.*, Allibert C.H.**, Luyckx S. School of Process and Materials Engineering, University of the Witwatersrand, Johannesburg, South Africa * Electron Microscope Unit, University of the Witwatersrand, Johannesburg, South Africa " LTPCM (UMR 5614 CNRS-INPG), Grenoble, France EFFECT OF COMPOSITON ON THE (V,W)C CONSTITUTION IN V-W-C-Co ALLOYS 4 291

HM 96 Münz W.D.*, Lembke M.I.*, Lewis D.B.*, Smith I.J.** * Sheffield Hallam University, Materials Research Institute, Sheffield, United Kingdom " Bodycote SHU Coatings Ltd., Sheffield, United Kingdom MICROSTRUCTURE, COMPOSITION AND PERFORMANCE OF PVD COATINGS DESIGNED FOR DRY HIGH SPEED MILLING 4 351 S\(o

HM 97 Westphal H., Van den Berg H., Sottke V., Tabersky R., König U. Widia GmbH, Essen, Germany

NEUE OXID-COMPOSITE BESCHICHTUNGEN FÜR SCHWIERIGE ZERSPANUNGSAUFGABEN 4 362 •*•' '1, &

HM 101 Haubner R., Tang X. Institute for Chemical Technology of Inorganic Materials, Univ. of Techn. Vienna, Vienna, Austria LOW-PRESSURE c-BN DEPOSITION -IS ACVD PROCESS POSSIBLE? 4 372 XXVI Powder Metallurgical High Performance Materials

General Topics

GT 4 German R.M. S 1 -I The Pennsylvania State University, Center for Innovative Sintered Products, University Park, USA UNIQUE OPPORTUNITIES IN POWDER INJECTION MOLDING OF REFRACTORY AND HARD MATERIALS 4 175

GT 8 Shields J.A., Lipetzky P.*, Mueller A.J.** CSM Industries Inc., Cleveland, USA * Pittsburgh Materials Technology, Inc., Pittsburgh, USA ** Bechtel Bettis Inc., West Mifflin, USA < 7 / FRACTURE TOUGHNESS OF 6.4 MM (0.25 INCH) ARC-CAST MOLYBDENUM AND MOLYBDENUM-TZM PLATE AT ROOM TEMPERATURE AND 300°C 4 1B7

GT 43 Smurov I.*, Hurevich V., Gusarov A.*, Kundas S., Kashko T. Belarusian State University of Informatics and Radioelectronics, Minsk, Belarus * Ecole Nationale d'lngenieurs de Saint-Etienne, St-Etienne Cedex, France <- ^ ,

COMPUTER SIMULATION OF ZrO2+8%Y2O3 AND AI2O3 POWDER PARTICLES HEATING UNDER PLASMA SPRAYING 4 200

GT 47 Schintlmeister A.*, Wilhartitz P. Plansee AG, Reutte, Austria * Plansee AG und Institut für Experimentalphysik, Karl Franzens Universität Graz, Graz, Austria <- i / PROTECTION OF MOLYBDENUM AND MOLYBDENUM-ALLOY THIN FILMS AGAINST LOW- TEMPERATURE OXIDATION 4 211

GT 48 Ramminger P., Tessadri R., Krismer R.*, Wilhartitz P.** Institute of Mineralogy, University of Innsbruck, Innsbruck, Austria * Plansee AG, Test Laboratories, Analytics, Reutte, Austria

" Plansee AG, Business Unit Thin Film Technology, Reutte, Austria _ „ / APPLICATION OF HIGH TEMPERATURE DIFFRACTION AS A TOOL OF MATERIAL- ^ ° CHARACTERISATION AND PRODUCT-OPTIMISATION 4 216 Powder Metallurgical High Performance Materials XXVII List of Authors Volume 4: Late Papers

Agren J. HM 44 Jones A.R. RM 7 Ruiz Urien I. RM 102 Alamo J. RM 102 RM 103 RM 103 Karpov M.I. RM 44 Ryzhkin A.A. HM 33 Allibert C.H. HM 89 Kashko T. GT 43 Andreani R. RM 1 Kawai M. RM 4 Saito A. RM 4 Andren H.O. HM 44 Khodoss I.I. RM 44 Sajgalik P. HM 46 Kikuchi K. RM 4 Santos I. RM 102 Bartha L. RM 76 Kim I. RM 89 RM 103 Bizaoui M. HM 49 King H.W. RM 89 Sanz-Tudanca C. RM 103 Bock A. HM 2 Kolchemanov N.A. RM 100 Schalunov E.P. RM 98 Broccardo S. HM 89 König U. HM 97 Schedle D. RM 16 Bulic F.I. HM 46 Korb G. HM 46 Schedler B. RM 16 Krismer R. GT 48 Scheiber K.H. RM 16 Chato R. RM 102 Kundas S. GT 43 Schiller G. RM 19 RM 103 Kurishita H. RM 4 Schintlmeister A. GT 47 Chen Y.L. RM 7 Schön A. HM 35 Clemente G. RM 103 Lackner K. RM 1 Schubert W.D. HM 12 Cornish L.A. HM 89 Lang M. RM 19 HM 35 Lembke M.I. HM 96 HM 57 Dabhade V.V. RM 57 Lences Z. HM 46 Shields J.A. GT 8 Dariel M.P. HM 49 Lengauer W. HM 56 Sima M. HM 65 Dedkov P.Y. RM 100 Lewis D.B. HM 96 Simonov V.S. RM 98 Dovydenkov V.A. RM 98 LiJ.F. RM 4 Smith I.J. HM 96 Lipetzky P. GT 8 Smurov I. GT 43 Ekroth M. HM 44 Lopez J.L. RM 103 Sommer M. HM 12 HM 35 Ferguson S. RM 89 Lux B. Sottke V. HM 97 HM 57 Frage N. HM 49 HM 89 Tabersky R. HM 97 Franco T. RM 19 Luyckx S. Tang X. HM 101 Friedrichs K. HM 88 Lyulko A.V. HM 33 Tessadri R. GT 48 Frykholm R. HM 44 Mahalingam M. RM 8 Furusaka M. RM 4 Medved N.V. RM 44 Van den Berg H. HM 97 Miller U. RM 7 Velasco T. RM 102 Gasent J.L. RM 102 HM 84 RM 103 RM 103 Morrell R. Mueller A.J. GT 8 Viswanathan V. RM 8 Gasparotto M. RM 1 MünzW.D. HM 96 Vnukov V.l. RM 44 Gee M.G. HM 84 Volkov K.G. RM 44 German R.M. GT 4 NagyD.R. RM 89 Godkhindi M.M. RM 57 Nojkina A.V. RM 100 Warbichler P. HM 12 GroszT. RM 76 Warren J. RM 36 Gursan S. RM 89 Okamoto K. RM 4 Watanabe R. RM 4 Gusarov A. GT 43 Weissenbacher R. HM 79 Gutmanas E.Y. RM 57 Pick M. RM 1 Westphal H. HM 97 Wilhartitz P. GT 47 Hamar-Thibault S. HM 89 Rainer F. RM 16 GT 48 Harmat P. RM 76 Rama Mohan T.R. RM 57 WitcombM.J. HM 89 Haubner R. HM 101 Ramakrishnan P. RM 57 Wittmann B. HM 57 HM 79 Ramminger P. GT 48 Wollein B. HM 56 Henne R. RM 19 Reglero V. RM 102 Holubar P. HM 65 RM 103 Yildiz M. RM 89 HuberT. RM 16 Reznikov G. RM 36 Hurevich V. GT 43 Rodrigo J. RM 103 Zabernig A. RM 16 Rodrigo J.M. RM 102 Zarauz J. RM 102 Igarashi T. RM 4 Roebuck B. HM 84 RM 103 llyasov V.V. HM 33 Rosta L. RM 76 Zeiler B. HM 2 Ruckdäschel R. RM 19 Zinyana J. HM 89 Jansson B. HM 44 Zobetz E. HM 12 Jilek M. HM 65 Powder Metallurgical High Performance Materials XXIX

Cumulative List of Presentations Volume 1-4

RM 1 Lackner K., Andreani R., Gasparotto M., Pick M. EFDA, Garching, Germany

ITER AND THE FUSION REACTOR: STATUS AND CHALLENGE TO TECHNOLOGY 4 471

RM 3 Joensson M., Kieback B. University of Technology Dresden, Institute of Material Science, Dresden, Germany P/M GRADIENT MATERIALS - PROCESSING, PROPERTIES AND APPLICATIONS 1 1

RM 4 Watanabe R., Li J.F., Saito A., Kurishita H., Igarashi T.*, Okamoto K.**, Kikuchi K.***, Furusaka M.****, Kawai M.**" Tohoku University, Sendai, Japan * Sumitomo Electric Co., Tokyo, Japan ** Allied Material Corp., Tokyo, Japan *** Japan Atomic Energy Research Institute, Tokai, Japan **** High Energy Accelerator Research Organization, Tsukuba, Japan DEVELOPMENT OF SOLID TARGET FOR SPALLATION NEUTRON SOURCE 4 1

RM 5 Glatz W., Janousek M., Batawi E.*, Doggwiler B.* Plansee AG, Reutte, Austria * Sulzer Hexis Ltd., Winterthur, Switzerland PM NET-SHAPE PROCESSING OF CR-BASED INTERCONNECTS FOR SOLID OXIDE FUEL CELLS 1 16

RM 6 Kim M.J., Doh J.M.*, Park J.K.**, Jung J.P. Micro-Joining Lab., Dept.of Materials Science and Eng., University of Seoul, Seoul, South Korea * Alloy Design Research Center, Div. of Materials, Korea Inst. of Science and Techn., Seoul, South Korea ** Ceramic Proc. Research Center, Div. of Materials, Korea Inst. of Science and Techn., Seoul, South Korea MICROSTRUCTURAL EVOLUTION IN THE Cr-Cu ELECTRIC CONTACT ALLOYS DURING LIQUID PHASE SINTERING 1 29

RM 8 Mahalingam M., Viswanathan V. Motorola Inc., Wireless Infrastructure Systems Division, Tempe, Az., USA MATERIALS REQUIREMENTS FOR THERMAL MANAGEMENT IN WIRELESS COMMUNICATION 4 505

RM 10 Kim D.G., Shim W.S., Kim J.S.*, Kim Y.D. Division of Materials Science and Engineering, Hanyang University, Seoul, South Korea -- t / * Dept. of Materials Engineering, REMM, University of Ulsan, Ulsan, South Korea 5 ,i b FABRICATION OF W-Cu NANOCOMPOSITE BY REDUCTION OF WO3-CUO 1 60

RM 11 Edtmaier C, Eglsäer C. Vienna University of Technology, Institute for Chemical Techn. of Inorganic Materials, Vienna, Austria FROM AQUEOUS METAL-SOLUTIONS TO SUB-MICRON POWDERS BY A CHEMICAL REDUCTION PROCESS AND ODS-PRODUCTS BY CO-PRECIPITATION 1 68

RM 12 Dore F., AMibert C.H., Baccino R.*, Lartigue J.F." Institut National Polytechnique de Grenoble, St. Martin d'Heres, France * CEA/CEREM, Grenoble, France ,- 0,! "" Eurotungstene Poudres, Grenoble, France . -' EFFECTS OF PHASE SIZE AND OF Fe ADDITIONS ON THE SINTERING BEHAVIOUR OF COMPOSITE POWDERS W-CU (20wt%) 1 81 XXX Powder Metallurgical High Performance Materials

RM 13 Schaper M.K., Böhm A., Heilmaier M.*, Klauss HJ., Nganbe M.* University of Technology Dresden, Institute of Materials Science, Dresden, Germany * IFW Dresden, Institute for Metallic Materials, Dresden, Germany FATIGUE LIFE AND FATIGUE CRACK GROWTH OF THE ODS NICKEL-BASE SUPERALLOY PM 1000 94

RM14 GeC, Wu A., CaoW., Li J. „ , Lab. of Special Ceramic and P/M, Univ. of Science and Technology Beijing, Beijing, Peoples Republic of China S •-> ° A NEW SiC/C BULK FGM FOR FUSION REACTOR 1 109

RM 15 Plankensteiner A., Schedler B., Krismer R., Rainer F. Plansee AG, Reutte, Austria DEVELOPMENT OF A LOW TEMPERATURE SOLID HIP PROCESS FOR JOINING CFC-MONOBLOCKS ONTO CuCrZr TUBES 1 116

RM 16 Schedler B., Huber T., Zabernig A., Rainer F., Scheiber K.H., Schedle D. Plansee AG, Reutte, Austria THE MANUFACTURE OF CARBON ARMOURED PLASMA-FACING COMPONENTS FOR FUSION DEVICES 26

RM 17 Nagata S.% Nasu S.**, Moon A.**, Yoshii K.**, Ando T.**, Ohhashi K.**, Sugimata E.**, Kikuchi N.", Kusano E.**, Kimbara A.** * Tohoku University, Sendai, Miyagi, Japan 7 ( " Kanazawa Institute of Technology, Nonoichi, Ishikawa, Japan ° COMPATIBILITY OF LITHIUM TITANATE FUSED CRYSTALS WITH TUNGSTEN SPUTTERED FILMS 1 134

RM 19 Franco T, Henne R., Lang M., Ruckdäschel R., Schiller G. Deutsches Zentrum für Luft- und Raumfahrt (DLR), Institut für Technische Thermodynamik, Stuttgart, Germany INVESTIGATION OF POROUS METALLIC SUBSTRATES FOR PLASMA SPRAYED THIN-FILM SOFC 64

RM 20 Olsson F., Svensson S.A., Duncan R.* Sycon Energikonsult AB, Malmb, Sweden EXTERNALLY FIRED GAS TURBINE CYCLES WITH HIGH TEMPERATURE HEAT EXCHANGERS UTILISING Fe-BASED ODS ALLOY TUBING 1 139

RM 21 Yin W., Tang H., Teng Xiuren Northwest Institute for Nonferrous Metal Research, Xi'an Shaanxi, Peoples Republic of China PERFORMANCES OF NOVEL INTEGRAL TUNGSTEN-COPPER CONTACT UNDER LOAD SERVICE TESTS 1 154

RM 23 Hiraoka Y., Inoue T., Akiyoshi N.*, Yoo M.K.** Graduate School of Science, Okayama University of Science, Okayama, Japan * R & D Division, Toho Kinzoku Co., Neyagawa, Japan *' Metals Division, Korea Institute of Science and Technology, Seoul, South Korea DUCTILE-TO-BRITTLE TRANSITION BEHAVIOR OF TUNGSTEN-COPPER COMPOSITES 1 163

RM 24 Gomes U.U., da Costa F.A.*, da Silva A.G.P. Universidade Fed. R. G. do Norte, Depto. de Fisica, Natal, Brazil t, 11> * IPEN, Depto. de Materials, Cid. Universitaria, Sao Paulo, Brazil ON SINTERING OF W-Cu COMPOSITE ALLOYS 1 177

RM 25 Akiyoshi N., Nakada K., Nakayama M., Kohda K. Toho Kinzoku Co. Ltd., R&D Division, Neyagawa, Japan n! EFFECTS OF PHOSPHOURUS ADDITION ON THE PHYSICAL PROPERTIES AND SURFACE CONDITION OF TUNGSTEN-COPPER COMPOSITES 1 190

RM 26 Frage N., Froumin N., Rubinovich L., Dariel M.P. Dept. of Materials Engineering, Ben-Gurion University of the Negev, Beer-Sheva, Israel INFILTRATED TiC/Cu COMPOSITES 1 202 Powder Metallurgical High Performance Materials XXXI

RM 29 Martinz H.P., Matej J.*, Leichtfried G. Plansee AG, Reutte, Austria * Laboratory of Inorganic Materials of IIC AS CR and ICT, Prague, Czech Republic THE CORROSION OF COATED AND UNCOATED MO AND OF UNCOATED MO-ZRO2 IN GLASS MELTS J6 UNDER A.C. LOADING 217

RM 34 Nagae M., Takemoto Y.*, Takada J.*, Hiraoka Y.**, Yoshio T. Faculty of Environmental Science and Technology, Okayama University, Okayama, Japan * Faculty of Engineering, Okayama University, Okayama, Japan ** Faculty of Science, Okayama University of Science, Okayama, Japan MECHANICAL PROPERTIES OF MOLYBDENUM-TITANIUM ALLOYS MICRO-STRUCTURALLY CONTROLLED BY MULTI-STEP INTERNAL NITRIDING 1 248

RM 35 Nagae M., Fukumitu T., Takada J.*, Hiraoka Y**, Yoshio T. Faculty of Environmental Science and Technology, Okayama University, Okayama, Japan * Faculty of Engineering, Okayama University, Okayama, Japan ** Faculty of Science, Okayama University of Science, Okayama, Japan GRAIN GROWTH CONTROL OF DILUTE MOLYBDENUM-TITANIUM ALLOYS BY MULTI-STEP INTERNAL NITRIDING 1 257

RM 40a Gallet D., Dhers J., Levoy R., Polcik P.* CEA, Centre de la vallee du Rhone, Service STME, Pierrelatte, France * Plansee AG, Reutte, Austria CREEP LAWS FOR REFRACTORY TUNGSTEN ALLOYS BETWEEN 900 AND 1100°C UNDER LOW STRESS 1 277 ssc

RM 41 Takida T., Kurishita H.*, Mabuchi M.**, Igarashi T.*", Doi Y."**, Nagae T."** Allied Material Corp., Toyama, Japan * The Oarai Branch, Institute for Materials Research, Tohoku Univ., Oarai, Japan ** National Industrial Research Institute of Nagoya, Nagoya, Japan *** Osaka Prefecture University, Sakai, Japan **** Toyama Industrial Technology Center, Takaoka, Japan MECHANICAL PROPERTIES OF FINE-GRAINED SINTERED MOLYBDENUM ALLOY PROCESSED BY MECHANICAL ALLOYING 1 293

RM42 Vasile T., Dobrescu M. University Politehnica Bucharest, Material Science Department, Bucharest, Romania 1 305

RM 43 Rao D., Upadhyaya G.S. Dept. of Materials & Metallurgical Engineering, Indian Institute of Technology, Kanpur, India SINTERING OF Mo2FeB2 BASED CERMET AND ITS LAYERED COMPOSITES CONTAINING SiC FIBERS 1 312 5

RM 47 Zhang H., Liu X.*, Ning A.** Dept. of Mech. Eng. & Autom., Xiangtan Polyt. Univ. & Central South Univ., Peoples Republik of China * Dept.of Materials Science and Engineering, Central South Univ., Hunan Changsha, Peoples Republic of China ** Dept. of Mech. Engineering, Shaoyang High Technical College, Hunan Shaoyang, Peoples Republic of China MICROSTRUCTURES AND PROPERTIES OF RARE EARTH-PARTICLES REINFORCED MoSi; COMPOSITES 1 339 XXXII Powder Metallurgical High Performance Materials

RM 48 Lyulko V.G., Shalunov E.P.*, Oleinikov D.V.**, Antropov V.V.*** Don State Technical University, Rostow-an-Don, Russia r * Wissenschaftliche-Produktionsfirma Techma GmbH, Tscheboksari, Russia ! / ** z.Z. Techn. Universität Wien, Institut für chemische Technologie anorg. Stoffe, Wien, Austria ••JO •*• AG "Atommasch", Volgodonsk, Russia

3D/2D-MODELLE FÜR GEFÜGE- UND EIGENSCHAFTEN: INTERPRETATION VON AI2O3-TiO2-Ti(C,N) - DISPERSIONVERFESTIGNER KUPFERVERBUNDWERKSTOFFEN 1 352

RM 50 Filgueira M., Pinatti D.G.* Universidade Estadual do Norte Fluminense, CCT-LAMAV, Campos, Brazil * FAENQUIL, Dept. Materials Engineering, Lorena, Brazil PRODUCTION OF DIAMOND WIRE BY Cu15v-% Nb "IN SITU" PROCESS 1 360

RM51 KangS.H., NamT.H. Division of Materials Engineering and RECAPT, Gyeongsang National University, Gyeongnam, South Korea FABRICATION OF Ti-Ni-Cu SHAPE MEMORY ALLOY POWDERS BY BALL MILLING METHOD 1 375

RM 52 Briant C.L. Brown University, Division of Engineering, Providence, USA NEW APPLICATIONS AND NOVEL PROCESSING OF REFRACTORY METAL ALLOYS 1 389

RM 54 Bewlay B.P., Briant C.L.*, Jackson M.R., Subramanian P.R., Subramanian P.R. General Electric Company, Corporate Research and Development Center, Schenectady, USA * Division of Engineering, Brown University, Providence, USA RECENT ADVANCES IN Nb-SILICIDE IN-S1TU COMPOSITES 1 404

RM 56 Bram M., Ahmad-Khanlou A., Buchkremer H.P., Stöver D. Forschungszentrum Jülich GmbH, Institut f. Werkstoffe und Verfahren der Energietechnik IWV-1, Julien, Germany

POWDER METALLURGY OF NiTi-ALLOYS WITH DEFINED SHAPE MEMORY PROPERTIES 1 434

PROCESSING OF NANO SIZED TITANIUM NITRIDE REINFORCED TITANIUM COMPOSITES 4 37

RM58 Kumar P., Uhlenhut H. H.C. Starck Inc., Newton, USA RECENT ADVANCES IN P/M-TANTALUM PRODUCTS 1 448

RM 59 Makarov P.V., Povarova K.B.* Leco Instrumente GmbH, Moscow, Russia * Baikov Institute of Metallurgy and Material Science, Russian Academy of Science, Moscow, Russia PRINCIPLES OF THE ALLOYING OF TUNGSTEN AND DEVELOPMENT OF THE MANUFACTURING TECHNOLOGY FOR THE TUNGSTEN ALLOYS 1 463

RM 60 Mueller A.J., Shields, Jr. J.A.*, Buckman, Jr. R.W.** Bechtel Bettis, Inc., West Mifflin, USA * CSM Industries, Inc., Cleveland, USA ~- \ L " Refractory Metals Technology, Pittsburgh, USA THE EFFECT OF THERMO-MECHANICAL PROCESSING ON THE MECHANICAL PROPERTIES OF MOLYBDENUM -2 VOLUME % LANTHAN A 1 485

RM 61 Suzuki T., Nomura N., Hanada S. Tohoku University, Institute for Materials Research, Sendai, Japan ^ 3 ( MICROSTRUCTURE AND MECHANICAL PROPERTIES OF Mo/ZrC IN-SITU COMPOSITES 1 498 Powder Metallurgical High Performance Materials XXXIII

RM 62 Brady M.P., Wright I.G., Anderson I.M., Sikka V.K., Ohriner E.K., Walls C, Westmoreland G., Weaver M.L.* Oak Ridge National Laboratory, Oak Ridge, USA * University of Alabama, Tusoaloosa, USA DUCTILIZATION OF Cr VIA OXIDE DISPERSIONS 1 506

RM 64 Zissis G. CPAT - Univ. Toulouse III, Toulouse, France

ELECTRICAL DISCHARGE LIGHT SOURCES: A CHALLENGE FOR THE FUTURE 1 521

RM 65 Van Hees G., Fischer E.*, Derra G.* Philips Lighting Turnhout, Turnhout, Belgium * Philips Forschungslabor, Aachen, Germany UHP: A PHILIPS LIGHTING INVENTION FOR PROJECTION APPLICATIONS 1 538

RM 66 Schlichtherle S., Leichtfried G.*, Martinz H.P.*, Retter B.* University of Graz, Graz, Austria * Plansee AG, Reutte, Austria CHARACTERIZATION OF THE INTERFACE MOLYBDENUM SEALING FOIL / VITREOUS SILICA 1 545

RM 67 Selcuk C, Wood J.V., Morley N.*, Bentham R.* University of Nottingham, Nottingham, United Kingdom * Tungsten Manufacturing Ltd., Brighton, United Kingdom CHARACTERISATION OF POROUS TUNGSTEN BY MICROHARDNESS 1 557 5 !> '

RM68 Nie Z., Chen Y., Zhou M., Zuo T. Beijing Polytechnic University, School of Materials Science & Eng., Beijing, Peoples Republic of China RESEARCH AND DEVELOPMENT OF TUNGSTEN ELECTRODES ADDED WITH RARE EARTH OXIDES 1 568 ^ °

RM 69 de Cachard J., Millot F.*, Cadoret K., Martinez L., Veillet D. Thales Electron Devices, Thonon-Les-Bains, France * Centre de Recherche sur les Materiaux Haute Temperature, CNRS, Orleans, France NEW DOPED TUNGSTEN CATHODES. APPLICATIONS TO POWER GRID TUBES. 1 573

RM 70 Appel F., Clemens H., Oehring M. Institute for Materials Research, GKSS Research Centre, Geesthacht, Germany DESIGN AND PROPERTIES OF ADVANCED y(TiAI) ALLOYS 1 588

RM 71 Terauchi S., Teraoka T., Shinkuma T., Sugimoto T.*, Ahida Y.** OsakaYakin Heat-Treating Co. Ltd., Osaka, Japan * Kansai University, Faculty of Engineering, Kansai, Japan ** Kobelco Research Institute Inc., Japan DEVELOPMENT OF PRODUCTION TECHNOLOGY BY METALLIC POWDER INJECTION MOLDING FOR TiAI-TYPE INTERMETALL1C COMPOUND WITH HIGH EFFICIENCY 1 609

RM 72 Böhm A., Scholl R., Kieback B. Fraunhofer Institute for Manufacturing and Advanced Materials, Dresden, Germany REACTION SINTERING OF COMPONENTS FROM y-TiAI-BASED ALLOYS 1 623

RM 73 Povarova K.B., Mileiko S.T.*, Sarkissyan N.S.*, Antonova A.V., Serebryakov A.V.*, Kolchin A.A.*, Korzhov V.P.* A.A.Baikov Institute of Metallurgy and Material Science RASc (IMET RAS), Moscow, Russia * Solid State Physics Institute RAS {SSPI RAS), Russia SAPPHIRE/TiAl COMPOSITES - STRUCTURE AND PROPERTIES 1 633

RM 74 Leonhardt T., Downs J. Rhenium Alloys Inc., Elyria, USA 0 / NEAR NET SHAPE OF POWDER METALLURGY RHENIUM 1 644 XXXIV Powder Metallurgical High Performance Materials

RM 76 Harmat P., Grosz T.*, Rosta L.**, Bartha L. MFA Research Institute for Technical Physics and Materials Science, Budapest, Hungary * Chemical Research Center, Institute of Chemistry, Budapest, Hungary ** KFKI SZFKI Research Institute for Solid State Physics and Optics, Budapest, Hungary A NOVEL METHOD FOR SHAPE ANALYSIS: DEFORMATION OF BUBBLES DURING WIRE DRAWING IN DOPED TUNGSTEN 4 108

RM 77 Horacsek 0., Bartha L. Research Institute for Technical Physics and Material Science of HAS, Budapest, Hungary -j i ö POTASSIUM INCORPORATION INTO TUNGSTEN BY ENTRAPMENT OF DOPANT PARTICLES FORMED ON THE SURFACE OF AKS-DOPED TBO 1 655

RM 82 Cadoret K., Millot F.*, de Cachard J., Martinez L., Hennet L.*, Douy A.*, Licheron M.* Thales Electron Devices, Thonon-Les-Bains, France " Centre de Recherche sur les Materiaux Haute Temperature, CNRS, Orleans, France '•_, ~\ 4 CHEMICAL BEHAVIOUR OF LANTHANIDES -TUNGSTEN COMPOSITE MATERIALS USED IN THERMO- EMISSIVE CATHODES 1 666

RM 83 Wang J., Zhou M., Zuo T., Zhang J., Nie Z. School of Materials Science and Engineering, Beijing Polytechnic University, Beijing, Peoples Republic of China S s4 CARBONIZATION KINETICS OF LA2O3-MO CATHODE MATERIALS 1 681

RM 84 Kablov E.N., Povarova K.B.*, Buntushkin V.P., Kasanskaya N.K.*, Basyleva O.A. VIAM, All-Russian Inst. of Aviation Materials, Russia * IMET A.A.Baikov, Inst. of Metallurgy and Materials Science RASc, Moscow, Russia DEVELOPMENT OF AERO-SPASE STRUCTURAL Ni3AI-BASED ALLOYS FOR SERVICE AT TEMPERATURES ABOVE 1000°C IN AIR WITHOUT PROTECTION COATING 1 692

RM 85 Povarova K.B., Lomberg B.S.*, Kazanskaya N.K., Gerasimov V.V.*, Drozdov A.A. IMET A.A.Baikov, Inst. of Metallurgy and Material Science RASc, Moscow, Russia * VIAM, All-Russian Inst. of Aviation Materials, Russia STRUCTURAL HIGH-TEMPERATURE AND (ßNiAI+y)-ALLOYS BASED ON Ni-AI-Co-Me SYSTEMS WITH AN IMPROVED LOW-TEMPERATURE DUCTILITY 1 707

RM 86 Kim J.S., Jang Y.I., Kim Y.D.*, Kwon Y.S. ReMM, University of Ulsan, Ulsan, South Korea * Division of Material Science & Engineering, Hanyang University, Hanyang, South Korea " Dept.of Met. Engineering, Gyeongsang University, Gyeongsang, South Korea SPARK-PLASMA SINTERING AND MECHANICAL PROPERTY OF MECHANICALLY ALLOYED NiAl POWDER COMPACT AND BALL-MILLED (Ni+AI) MIXED POWDER COMPACT 1 723

RM 87 Lucaci M., Vidu C, Vasile E. Metav S.A.- Aviation Metallurgy, Bucharest, Romania

THE Ni3AI AND NiAl ALLOYS: A CLASS OF INTERMETALLICS WHICH CAN REPLACE THE Ni-BASE SUPERALLOYS FOR THE AEROSPACE HIGH TEMPERATURE STRUCTURAL APPLICATIONS 1 733

RM 89 King H.W., Gursan S., Ferguson S., Yildiz M., Kim I.*, Nagy D.R.** University of Victoria, Dept. of Mechanical Engineering, Victoria, Canada * School of Advanced Materials, Kum-Oh National Univ. of Technology, Kumi, Kyung Bug, South Kore< ** Liburdi Engineering Ltd., Dundas, Canada RESIDUAL STRESS IN TiN COATINGS ON TN6AI-4V ALLOY 4 118

RM 90 Tabernig B., Thierfelder W., Alber H., Sudmeijer K.* Plansee AG, Reutte, Austria * Fokker Space, Leiden, The Netherlands FABRICATION OF METALLIC HONEYCOMB PANELS FOR REUSABLE TPS-STRUCTURES 1 743 Powder Metallurgical High Performance Materials XXXV

RIM 92 Yi W., German R.M., Lu P.K. Pennsylvania State University, Center for Innovative Sintered Products, University Park, USA SHAPE DISTORTION AND DIMENSIONAL PRECISION IN TUNGSTEN HEAVY ALLOY LIQUID PHASE SINTERING 1 758

RM 93 Nicolas G., Voltz M. Cime Bocuze, St.Pierre-en-Faucigny, France TUNGSTEN HEAVY METAL ALLOYS - RELATIONS BETWEEN THE CRYSTALLOGRAPHIC TEXTURE S 1 AND THE INTERNAL STRESS DISTRIBUTION 1 777

RM 95 Sfat C, Vasile T., Vasilescu M. Material Science Department, Politehnica University of Bucharest, Bucharest, Romania RESEARCHES FOCUSED ON STRUCTURE OF ALUMINIUM ALLOYS PROCESSED BY RAPID SOLIDIFACTION, USED IN AUTOMOTIVE INDUSTRY 1 789

RM 96 Motoiu P., Rosso M.* Institute for Non Ferrous & Rare Metals, Bucharest, Romania * Politecnico di Torino, Torino, Italy CHARACTERISATION AND PROPERTIES OF SHOCK AND CORROSION RESISTANT OF TITANIUM BASED COATINGS 1 798

RM 97 Terentieva V. Moscow State Aviation Institute, Techn. University, Moscow, Russia CHARACTERIZATION OF NOVEL HETEROPHASIC POWDERED SILICIDE-TYPE MATERIAL FOR HIGH - TEMPERATURE PROTECTION SYSTEMS 1 808

RM 100 Nojkina A.V., Kolchemanov N.A., Dedkov P.Y.* VNIIALMAZ, Moscow, Russia * VNIITF, Snezhinsk, Russia SINTERING OF NANODIAMOND IN THERMODYNAMIC STABILITY AREA OF DIAMONDS 4 52

RM 101 Labusca-Lazarescu E. Institute "Optoelectronica - 2001", Bucharest, Romania

SINTERING OF HIGH-ALUMINA TRANSLUCENT ARC-DISCHARGE TUBES FOR HPSV-LAMPS 1 820

RM 102 Reglero V., Ruiz Urien I.*, Velasco T., Santos I.*, Rodrigo J.M., Zarauz J.*, Gasent J.L., Alamo J., Chato R. Astronomy and Space Group, Inst, de Ciencias de los Materiales, Univ. of Valencia, Valencia, Spain * Sener, Las Arenas, Spain TUNGSTEN AND OPTICS 4 150

RM 103 Reglero V., Ruiz Urien I.*, Velasco T., Santos I.*, Rodrigo J., Zarauz J.*, Gasent J.L., Alamo J., Chato R., Clemente G.**, Sanz-Tudanca C.**, Lopez J.L.** Grupo de Astronomia y Ciencias del Espacio, Inst, de Ciencias de los Materiales, Univ. of Valencia, Valencia, Spain * Sener, Las Arenas, Spain ** Mecanizados Gines S.A., Miranda de Ebro, Spain HIGH PRECISION TUNGSTEN CUTTING FOR OPTICS 4 161 XXXVI Powder Metallurgical High Performance Materials

HM 1 Gille G., Szesny B., Dreyer K.*, Van den Berg H.*, Schmidt J.**, Gestrich T.***, Leitner G.*** H.C. Starck, Goslar, Germany * Widia Valenite, Essen, Germany ** HPTec GmbH, Ravensburg, Germany *** Fraunhofer Institute IKTS, Dresden, Germany SUBMICRON AND ULTRAFINE GRAINED HARDMETALS FOR MICRODRILLS AND METAL CUTTING INSERTS 2 733

HM 2 Bock A., Zeiler B. Wolfram Bergbau- und Hüttengesellschaft mbH NfG KG, St. Martin, Austria PRODUCTION AND CHARACTERIZATION OF ULTRAFINE WC POWDERS 4 229

HM 3 Natter H., Lackner A.*, Kniinz G.*, Hempelmann R. Physikalische Chemie, Universität des Saarlandes, Saarbrücken, Germany s -. { " F'lansee TIZIT AG, Reutte, Austria J -1 '° IN SITU HIGH TEMPERATURE X-RAY DIFFRACTION STUDY ON TUNGSTEN CARBIDE FORMATION 2 1

HM 4 Carroll D.F., Järvi J.* OMG Americas, Research Triangle Park, USA * OMG Kokkola Chemicals Oy, Kokkola, Finland GRANULATED COBALT POWDERS FOR WC/Co MATERIALS 2 15

HM 5 Falkovsky V.A., Klyachko L.I., Glushkov V.N., Khokhlov A.M., Eiduk O.N., Lukashova N.M. State-owned Unitary Enterprise All Russia Res. and Design. Inst. for Refractory&Hard Metals, Moscow, Russia MULTICARBIDE HARDMETALS 2 29

HM 6 Azcona I., Ordonez A., Dominguez L.*, Sanchez J.M., Castro F. CEIT-Centro de Estudios e Invest. Tecnicas de Guipuzcoa/Escuela Superior de Ing., Univ.de Navarra, San Sebastian, Spain * Escuela Superior de Ingenieros de San Sebastian, Univ. de Navarra, San Sebastian, Spain HOT ISOSTATIC PRESSING OF NANOSIZED WC-Co HARDMETALS 2 35

HM 7 Kishino J., Nomura H.*, Shin S.G.*, Matsubara H.*, Tanase T. Mitsubishi Materials Corp., Yuuki-gun, Japan " Japan Fine Ceramics Center, Nagoya, Japan COMPUTATIONAL STUDY ON GRAIN GROWTH IN CEMENTED CARBIDES 2 50

HM 9 Warbichler P., Hofer F., Grogger W., Lackner A.* Forschungsinstiut für Elektronenmikroskopie Graz, Techn. Universität Graz, Graz, Austria * Plansee Tizit AG, Reutte, Austria EFTEM-EELS CHARACTERIZATION OF VC AND Cr3C2 DOPED CEMENTED CARBIDES 2 65

HM 10 Milman Y.V., Luyckx S.*, Goncharuck V.A., Northrop J.T.** Institute for Problems of Materials Science, Kiev, Ukraine * University of the Witwatersrand, Johannesburg, South Africa C ~i 4 ** Boart Longer Research Centre, South Africa MECHANICAL PROPERTIES IN BENDING TESTS AND MECHANICAL BEHAVIOUR OF SUBMICRON AND MICRON WC-Co GRADES AT ELEVATED TEMPERATURES 2 75

HM 11 Falkovsky V., Blagoveschenski Y.*, Glushkov V.N., Klyachko L., Khokhlov A. All-Russia Institute for Refractory and Hard Metals, Moscow, Russia * Institute of Metallurgy and Materials Science RAS, Moscow, Russia ^ > b NANOCRYSTALLINE WC-Co HARDMETALS PRODUCED BY PLASMOCHEMICAL MEHTOD 2 91 Powder Metallurgical High Performance Materials XXXVII

HM 12 Sommer M., Schubert W.D., Zobetz E.*, Warbichler P.** Institute for Chemical Technology of Inorganic Materials, Vienna University of Technology, Vienna, Austria * Institue for Mineralogy, Crystallography and Structural Chemistry, Vienna University of Technology, Vienna, Austria " Research Institute for Electron Microscopy, Graz, Austria ON THE FORMATION OF VERY LARGE WC CRYSTALS DURING SINTERING OF ULTRAFINE WC-Co ALLOYS 4 299

HM 13 Bartha L., Kotsis I.*, Laczko L.*, Harmat P. Research Institute for Technical Physics and Material Sciences, Budapest, Hungary * Inst. f. Silicate and Materials Engineering, University of Veszprem, Veszprem, Hungary C i •' HIGH TEMPERATURE REACTIONS OF WC-Co NANOPOWDERS IN VARIOUS ATMOSPHERES 2 97 - i o

HM 15 Lay S., Hamar-Thibault S., Lackner A.* LTPCM, INPG/CNRS/UJF, St. Martin d'Heres cedex, France * Plansee TIZIT AG, Reutte, Austria HREM CHARACTERISATION OF VC IN DOPED WC-Co CERMETS 2 106 > ']

HM 16 Cha S.I., Hong S.H., Ha G.H.*, Kim B.K.* Korea Advanced Institute of Science and Technology, Dept. of Materials Science and Engineering, Taejon, South Korea * Korea Institute of Machinery & Materials, Kyungnam, South Korea MICROSTRUCTURE AND MECHANICAL PROPERTIES OF ULTRA-FINE WC-10Co HARDMETALS 2 119 ^ '

HM 17 Zgalat-Lozynskyy O.B., Ragulya A.V., Herrmann M.* Institute for Problems of Materials Science, Ukrainian National Academy of Science, Kiev, Ukraine * Fraunhofer Institute of Ceramics Technology and Sintering, Germany HIGH MELTING POINT NANOCRYSTALLINE CERAMICS 2 131

HM18 Morita S., Shimizu H„ Sayama Y. Japan New Metals Co. Ltd., Osaka, Japan SYNTHESIS OF CHROMIUM NITRIDE POWDER BY CARBO-THERMAL NITRIDING 2 139

HM 21 Vasile T., Sfat C, Vasilescu M., Zafiu G. Physical Metallurgy Dept., Politehnica University of Bucharest, Bucharest, Romania THE IMPROVEMENT OF THE WC-TiC-Co ALLOY PERFORMANCES, REALISED BY ADDITION OF - 1 (Ta,Nb)C 2 152 •-> •>

HM 22 Vasile T., Zafiu G., Vasilescu M., Vasilescu I., Sfat C. Physical Metallurgy Dept., Politehnica University of Bucharest, Bucharest, Romania THE IMPROVEMENT OF CLASSICAL ALLOYS BY ADDING Cr3C2 CARBIDES 2 160

HM23 LisovskyA.F. Institute for Superhard Materials, Kiev, Ukraine PROPERTIES OF CEMENTED CARBIDES ALLOYED BY METAL MELT TREATMENT 2 168

HM25 Chen H., Zhang D., Li Y., Chen J. Central Iron & Steel Research Institute, Beijing, Peoples Republic of China HIGH PERFORMANCE SINTER-HIP FOR HARD METALS 2 180

HM 26 Bondarenko V., Pavlotskaya E., Martynova L., Epik I. V. Bakul Institute for Superhard Materials of the National Academy, Kiev, Ukraine ADVANCED TECHNOLOGIES OF PRODUCTION OF CEMENTED CARBIDES AND COMPOSITE MATERIALS BASED ON THEM 2 189 XXXVIII Powder Metallurgical High Performance Materials

HM 27 Vleugels J., Volders I., Put S., Zhao C, Van der Biest O., Groffils C.*, Luypaert P.J.*, Barbier G.**, Bourgeois L.** Katholieke Universiteit Leuven, Dept. of Metallurgy and Materials Engineering, Leuven, Belgium * Microwave Energy Applications Company, MEAC, Leuven, Belgium " Cerametal Sari, Mamer, Luxembourg HYBRID-MICROWAVE SINTERING OF HARDMETALS AND GRADED OXIDE COMPOSITES 2 204

HM 28 Zhou J.C., Huang B.Y. State Key Laboratory for Powder Metallurgy, Central South University, Changsha, Peoples Republic of China A NOVEL PLASTIFICATION AGENT FOR CEMENTED CARBIDES EXTRUSION MOLDING 2 216

HM 29 Mori G., Zitter H., Lackner A.*, Schretter M.* Dept. for General and Analytical Chemistry, University of Leoben, Leoben, Austria * Plansee Tizit AG, Reutte, Austria INFLUENCING THE CORROSION RESISTANCE OF CEMENTED CARBIDES BY ADDITION OF CR2C3, TiC AND TaC 2 222

HM 30 Roebuck B., Stewart M., Gee M.G., Bennett E.G. NPL Materials Centre, National Physical Laboratory, Teddington, United Kingdom A SCANNING DEPTH SENSING HARDNESS TEST SYSTEM TO CHARACTERISE HARDMETALS 2 237

HM 31 Missiaen J.M., Lackner A.*, Schmid L* Lab. de Thermodynamique et de Physico-Chimie Metallurgiques, St. Martin d'Heres, France * Plansee Tizit AG, Reutte, Austria > i b QUANTITATIVE ANALYSIS OF BIMODAL GRAIN SIZE DISTRIBUTIONS IN WC-CO HARD MATERIALS 2 252

HM 32 Engqvist H., Ederyd S., Hogmark S.*, Uhrenius B. AB Sandvik Hard Materials, Stockholm, Sweden * Dept. of Materials Science, Uppsala, Sweden SLIDING WEAR OF CEMENTED CARBIDES 2 267

HM 33 Ryzhkin A.A., llyasov V.V., Lyulko A.V. Don State Technical University, Rostov-on-Don, Russia WEAR RESISTANCE AND ELECTRONIC STRUCTURE OF CUTTING TOOL MATERIALS ON A BASIS CARBIDES OF TUNGSTEN AND TITANIUM 4 315

HM 34 Kitamura K., Kobayashi M., Hayashi K.* Toshiba Tungaloy Co. Ltd., Iwaki, Japan * Institute of Industrial Science, The University of Tokyo, Tokyo, Japan PROPERTIES OF NEW WC-Co BASE HARDMETAL PREPARED FROM CoxWyCz COMPLEX CARBIDE POWDER INSTEAD OF WC+Co MIXED POWDER 2 279

HM 35 Schön A., Schubert W.D., Lux B. Institute for Chemical Technology of Inorganic Materials, Vienna University of Technology, Vienna, Austria WC PLATELET-CONTAINING HARDMETALS 4 322

HM 36 Kim S., Han S.H.*, Park J.K.\ Kim H.E. Center for Microstructure Science of Materials, School of Materials Science & Engineering, Seoul, South Korea * Ceramic Processing Res. Center, Div. of Materials, Korea Inst. of Science and Techn., Seoul, South Korea OBSERVATION OF WC GRAIN SHAPES DETERMINED BY CARBON CONTENT DURING LIQUID PHASE SINTERING OF WC-Co ALLOYS 2 294

HM 37 Toth R.E., Smid I.*, Sherman A.**, Ettmayer P.*", Kladler G.*, Korb G.* TCHP Inventor, En Dur Aloy Corp., Savannah, USA * Austrian Research Center Seibersdorf, High Performance Materials, Seibersdorf, Austria " Powdermet Inc., Sun Valley, USA *** Technical University of Vienna, Vienna, Austria TOUGH-COATED HARD POWDERS FOR HARDMETALS OF NOVEL PROPERTIES 2 306 Powder Metallurgical High Performance Materials XXXIX

HM 38 Moriguchi H., Tsuzuki K., Ikegaya A. Itami Research Laboratories, Sumitomo Electric Industries Ltd., Itami, Hyogo, Japan DIAMOND DISPERSED CEMENTED CARBIDE PRODUCED WITHOUT USING ULTRA HIGH PRESSURE EQUIPMENT 2 326

HM 39 Scrivani A., Rosso M.*, Salvarani L.** Turbocoating S.p.A. and Universita die Parma, Dip. di Ingegneria Industriale, Rubbiano, Parma, Italy * Politecnico di Torino, Dip. di Scienze dei Materiali e Ingegneria Chimica, Torino, Italy ** Flametal S.p.A., Fornovo Taro, Parma, Italy PERFORMANCES AND RELIABILITY OF WC BASED THERMAL SPRAY COATINGS 2 337

HM 40 Dreyer K., Kassel D., Daub H.W., Van den Berg H., Lengauer W.*, Garcia J.*, Ucakar V.* Widia GmbH, Essen, Germany * Institute for Chemical Techn. of Inorganic Materials, Vienna University of Technology, Vienna, Austria FUNCTIONALLY GRADED HARDMETALS AND CERMETS: PREPARATION, PERFORMANCE AND PRODUCTION SCALE UP 2 768

HM 41 Gasik M.M., Zhang B„ Kaskiala M., Ylikerälä J. Helsinki University of Technology, Espoo, Finland SINTERING OF HARDMETALS IN DIFFERENT CONDITIONS: EXPERIMENTAL RESULTS OF 2-D DILATOMETRY AND COMPUTER SIMULATIONS 2 351

HM 42 Put S., Vleugels J., Van der Biest O. Katholieke Universiteit Leuven, Dept. of Metallurgy and Materials Engineering, Leuven, Belgium FUNCTIONALLY GRADED WC-Co HARDMETALS 2 364

HM 43 Ucakar V., Krai C, Dreyer K.*, Lengauer W. Institute for Chemical Technology of Inorganic Materials, Vienna University of Technoloy, Vienna, Austria * Widia GmbH, Essen, Germany NEAR-SURFACE MICROSTRUCTURAL MODIFICATION OF (Ti,W)(C,N)-BASED COMPACTS WITH NITROGEN 2 784

HM 47 Wu A., Cao W., Ge C, Li J. Lab. of Special Ceramic and P/M, University of Science and Technology Beijing, Beijing, Peoples Republic of China

FABRICATION OF Alpha/Beta-SIALON SYMMETRICALLY COMPOSITIONAL FGM 2 375

HM 48 Li J.T., Djuricic B.*, Korb G.*, Ge C.C. Lab. of Special Ceram. & P/M, University of Science and Technology Beijing, Beijing, Peoples Republic of China * Austrian Research Center Seibersdorf, Seibersdorf, Austria ALPHA-SiAION-NEW OPPORTUNITIES VIA NEW MICROSTRUCTURE 2 383

HM 49 Frage N., Bizaoui M., Dariel M.P. Ben-Gurion University of the Negev, Dept. of Materials Engineering, Beer-Sheva, Israel THE EFFECT OF Cr2O3 ADDITIONS ON THE PRESSURELESS SINTERING OF B4C 4 397 XL Powder Metallurgical High Performance Materials

HM 52 Takagi K.I., Yonetsu M., Yamasaki Y. Toyo Kohan Co. Ltd., Kudamatsu, Japan NOVEL BORIDE BASE CERMETS WITH VERY HIGH STRENGTH 2 403

HM 53 Richter V., Von Ruthendorf M. Fraunhofer-Institut Keramische Technologien und Sinterwerkstoffe, Dresden, Germany GRADIENT STRUCTURES IN HARDMETALS 2 800

HM 58 Ordanyan S.S., Zalite I.*, Andronova T.E., Vladimirova M.A., Pantelejev I.B. St.-Petersburg State Institute of Technology, St.-Petersburg, Russia * Riga Technical University, Institute of inorganic Chemistry, Salaspils, Latvia > 3 b PROPERTIES OF HARD ALLOYS ON THE BASIS OF WC-Co WITH THE ADDITIVES OF NANODISPERSE TiN 2 434

HM 59 Garcia J., Lengauer W. Institute for Chemical Technology of Inorganic Materials, Vienna University of Technology, Vienna, Austria QUANTITATIVE MASS SPECTROMETRY OF DECARBURISATION AND DENITRIDATION OF CEMENTED CARBONITRIDES DURING SINTERING 2 814

HM 62 Holzschuh H. Walter AG, Tübingen, Germany CHEMICAL-VAPOR DEPOSITION OF WEAR RESISTANT HARD COATINGS IN THE Ti-B-C-N SYSTEM: PROPERTIES AND METAL-CUTTING TESTS 2 441

HM 63 Borisova A., Borisov Y., Shavlovsky E., Mits I., Castermans L.*, Jongbloed R.* E.O. Paton Electric Welding Institute, Kiev, Ukraine * Chromin Maastricht bv., Maastricht, The Netherlands ^ ' ° VANADIUM CARBIDE COATINGS: DEPOSITION PROCESS AND PROPERTIES 2 452

HM65 JilekM., HolubarP., SimaM. SHM Ltd., Novy Malin, Czech Republic HARD PVD COATINGS nc-(Ti1-xAlx)N/a-Si3N4 NOT ONLY FOR DRY MACHINING 4 447

HM 69 Shevchenko A.D. Institute of Metal Physics Ukrainian AS, Kyiv, Ukraine PRODUCTION OF MATERIAL WITH HIGHER DAMPING ABILITY FOR SUPERHARD CUTTING TOOLS 2 463

HM70 Rosso M., Ugues D., Valle A.* Politecnico di Torino, Torino, Italy * MG Utensili Diamantati, Torino, Italy ADVANCES IN DIAMOND TOOLS FOR WORKING LITHOID MATERIALS 2 469

HM 72 Liu S., Yu Z., Yi D., Li Y* Dept. of Materials Science & Engineering, Central South University, Hunan, Peoples Republic of China * Hunan Yin Zhou Nonferrous Metals Hi-Tech. Ltd. Company, Hunan, Peoples Republic of China CHEMICAL PRETREATMENTS AT THE SURFACE OF WC-15 wt% Co FOR DIAMOND COATINGS 2 484 Powder Metallurgical High Performance Materials XLI

HM 73 Zhou K.S., Dai M.J., Song J.B., Kuang T.C.*, Liu Z.Y.* Guangzhou Research Institute of Non-Ferrous Metals, Guangzhou, Peoples Republic of China "South China University of Technology, Guangzhou, Peoples Republic of China THE STUDY ON DIAMOND-COATED INSERT BY DC PLASMA JET CVD 2 492

HM 74 Köpf A., Haubner R., Lux B. Institute for Chemical Technology of Inorganic Materials, TU - Vienna, Vienna, Austria MULTILAYER COATINGS CONTAINING DIAMOND AND OTHER HARD MATERIALS ON HARDMETAL SUBSTRATES 2 503

HM 75 Haubner R., Lux B. Institute for Chemical Technology of Inorganic Materials, Univ. of Technology Vienna, Vienna, Austria DEPOSITION OF BALLAS DIAMOND AND NANO-CRYSTALL1NE DIAMOND 2 827

HM 76 Buck V., Deuerler F.*, Kluwe H., Schmiler B. Universität GH Essen, Essen, Germany * Bergische Universität GH Wuppertal, Wuppertal, Germany INFLUENCE OF CHEMICAL PRETREATMENTS OF HARD METAL SUBSTRATES FOR DIAMOND DEPOSITION 2 517

HM 77 Dittrich K.H., Oelsner D* Roth & Rau Oberflächentechnik GmbH, Wüstenbrand, Germany * Hochschule für Technik und Wirtschaft Mittweida, Mittweida, Germany HERSTELLUNG UND CHARAKTERISIERUNG VON TROCKENSCHMIERSTOFFSCHICHTEN FÜR WERKZEUGE AUF DER BASIS VON KOHLENSTOFF 2 529

HM 78 Yedave S.N., Sundaram N., Brown W.D., Russell W.C.* MRL, Dept. of Mechanical Engineering, University of Arkansas, Fayetteville, USA * Valenite Inc., Troy, USA A NOVEL COMBINATORIAL APPROACH FOR THE REALIZATION OF ADVANCED cBN COMPOSITE COATING 2 544

HM 79 Weissenbacher R., Haubner R. Institute for Chemical Technology of Inorganic Materials. Vienna Univ. of Technology, Vienna, Austria DEPOSITION OF BCN LAYERS FROM B-C-N-H SINGLE SOURCE PRECURSORS 4 457

HM 80 Fleischer W., Baranski N.; van der Kolk G.J.*, Stockmann Y.*, Kunen H.**, Hoppe S.*** Bodycote Coating Centrum BV, Venlo, The Netherlands * Hauzer Techno Coating Europe BV, Venlo, The Netherlands ** Jabro Tools BV, Lottum, The Netherlands *** WZL, Aachen, Germany

HARD MACHINING UNDER DRY CONDITIONS WITH HARD PVD COATINGS ON CEMENTED CARBIDE ENDMILLS 2 559

HM 82 Richter V. Fraunhofer-Institut Keramische Technologien und Sinterwerkstoffe, Dresden, Germany VERSCHLEISSMECHANISMEN BEI DER BEARBEITUNG VOM GRAUGUSS 2 841

HM 83 Mertens T., Engering G., Lahres M., Pecher U.*, Damm S.*, Dörr J.**, Hühsam A,*** DaimlerChrysler AG, Ulm, Germany * Institute of Production Engineering and Logistics (IPL), Kassel, Germany " Institute of Production Engineering and Machine Tools (PtW), Darmstadt, Germany *** Institute of Machine Tools and Production Science (wbk), Karlsruhe, Germany IN-SITU TEMPERATURE MEASUREMENT TO DETERMINE THE MACHINING BEHAVIOUR OF DIFFERENT TOOL COATINGS 2 571

HM 84 Roebuck B., Gee M.G., Morrell R. NPL Materials Centre, National Physical Laboratory, Teddington, United Kingdom HARDMETALS - MICROSTRUCTURAL DESIGN, TESTING AND PROPERTY MAPS 4 245 XLII Powder Metallurgical High Performance Materials

HM 85 Upadhyaya A., Sarathy D.*, Wagner G.* Materials & Metallurgical Engineering Dept., Indian Inst. of Technology, Kanpur, India * Widia (India) Ltd., Bangalore, India IMPROVEMENT IN PROCESSING OF METAL-CUTTING INSERTS THROUGH SINTERING TIME REDUCTION AND ALLOY DESIGN 2 585

HM 86 Schwenke G.K. Osram Sylvania Inc., Chemical & Metallurgical Products, Towanda, USA

THERMODYNAMICS OF THE HYDROGEN-CARBON-OXYGEN-TUNGSTEN SYSTEM, AS APPLIED TO THE MANUFACTURE OF TUNGSTEN AND TUNGSTEN CARBIDE 2 598

HM 87 Andersson G., Jansson B., Föjer C. AB Sandvik Coromant, Stockholm, Sweden THE SOLUBILITY OF CUBIC CARBIDE FORMERS IN LIQUID COBALT 2 613

HM 88 Friedrichs K. Konrad Friedrichs KG, Kulmbach, Germany TRANSITION FROM HSS TOOLS TO SOLID CARBIDE TOOLS IN THE FIELD OF SHAFT TOOLS AND THE DEVELOPMENT OF NEW TYPES OF CARBIDES 4 267

HM 89 Zinyana J., Broccardo S., Hamar-Thibault S.**, Cornish L.A., Witcomb M.J.*, Allibert C.H.", Luyckx S. School of Process and Materials Engineering, University of the Witwatersrand, Johannesburg, South Africa * Electron Microscope Unit, University of the Witwatersrand, Johannesburg, South Africa ** LTPCM (UMR 5614 CNRS-INPG), Grenoble, France EFFECT OF COMPOSITON ON THE (V,W)C CONSTITUTION IN V-W-C-Co ALLOYS 4 291

HM 90 Agrawal D.K., Papworth A.J.*, Cheng J., Jain H.*, Williams D.B.* Pennsylvania State University, Materials Research Laboratory, University Park, USA * Lehigh University, Dept. of Materials Science and Engineering, Betheleham, USA (•- <ä / MICROSTRUCTURAL EXAMINATION BY TEM OF WC/Co COMPOSITES PREPARED BY CONVENTIONAL 1 b AND MICROWAVE PROCESSES 2 628

HM91 Yong Z., Weihao X. Huazong University of Science and Technology, Wuhan, Peoples Republic of China SOME PECULIARITIES OF THE RIM STRUCTURE OF Ti (C.N)-BASED CERMETSHM 92

HM 92 Beste U., Engqvist H., Jacobson S. Tribomaterials Group at the Angstrom Laboratory, Uppsala University, Uppsala, Sweden PRESSURE CYCLING INDUCED MODIFICATION OF A CEMENTED CARBIDE 2 636

HM 93 Klocke F., Bausch S. Fraunhofer-Institute of Production Technology IPT, Aachen, Germany LASER-ASSISTED TURNING OF COMPONENTS MADE OF SILICON-NITRIDE CERAMICS 2 649

HM94 Collins J.L. De Beers Industrial Diamonds (UK) Ltd., Ascot, United Kingdom HIGH SPEED DRY MACHINING OF MMCs WITH DIAMOND TOOLS 2 662

HM 96 Münz W.D.*, Lembke M.I.*, Lewis D.B.*, Smith I.J." * Sheffield Hallam University, Materials Research Institute, Sheffield, United Kingdom ** Bodycote SHU Coatings Ltd., Sheffield, United Kingdom MICROSTRUCTURE, COMPOSITION AND PERFORMANCE OF PVD COATINGS DESIGNED FOR DRY HIGH SPEED MILLING 4 351 Powder Metallurgical High Performance Materials XLIII

HM 97 Westphal H., Van den Berg H., Sottke V., Tabersky R., König U. Widia GmbH, Essen, Germany NEUE OXID-COMPOS1TE BESCHICHTUNGEN FÜR SCHWIERIGE ZERSPANUNGSAUFGABEN 4 362

HM 98 Kathrein M., Heiss M., Rofner R., Schleinkofer U., Schintlmeister W., Schatte J., Mitterer C* Plansee TIZIT AG, Reutte, Austria * Dept. of Physical Metallurgy and Materials Testing, University of Leoben, Leoben, Austria WEAR PROTECTION IN CUTTING TOOL APPLICATIONS BY PACVD (Ti,AI)N AND AI2O3 COATINGS 2 677

HM 99 Glushkov V.N., Anikeev A.I., Anikin V.N., Vereschaka A.* Unitary State-Owned Enterprise (GUP) VNIITS, Moscow, Russia * State Technological University "Stankin", Moscow, Russia MULTILAYERED AND COMPOSITE PVD-CVD COATINGS IN CEMENTED CARBIDES MANUFACTURE 2 691

HM 100 Russell W.C., Malshe A.P.*, Yedave S.N.", Brown W.D.* Valenite Inc., Troy, USA * MRL-MEEG, University of Arkansas, Fayetteville, USA NOVEL COMPOSITE cBN-TiN COATING DEPOSITION METHOD: STRUCTURE AND PERFORMANCE IN METAL CUTTING 2 705

HM101 Haubner R., Tang X. Institute for Chemical Technology of Inorganic Materials, Univ. of Techn. Vienna, Vienna, Austria LOW-PRESSURE c-BN DEPOSITION-IS A CVD PROCESS POSSIBLE? 4 372

HM 103 Lugscheider E., Hornig T., Kienitz S., Klocke F.*, Krieg T.* RWTH Aachen, Lehr- u. Forschungsgebiet Werkstoffwissenschaften, Aachen, Germany * RWTH Aachen, Lab. f. Werkzeugmaschinen, Lehrst.f.Techn. d. Fertigungsverfahren, Aachen, Germany WEITERENTWICKLUNG UMWELTVERTRÄGLICHER TRIBOSYSTEME IN DER ZERSPANUNG MITTELS PVD-TECHNOLOGIE 2 717

GT 2 Wieters K.P., Kieback B. Dresden University of Technology, Institute of Materials Science, Dresden, Germany COATED METAL POWDER, SINTERING BEHAVIOUR AND APPLICATION IN ALTERNATIVE FORMING PROCESSES 3 1

GT 4 German R.M. The Pennsylvania State University, Center for Innovative Sintered Products, University Park, USA UNIQUE OPPORTUNITIES IN POWDER INJECTION MOLDING OF REFRACTORY AND HARD MATERIALS 4 175

GT 6 Ge C, Li J., Cao W., Shen W. Lab. of Special Ceramic and P/M, Univ. of Science and Technology Beijing, Beijing, Peoples Republic of China SHS COMPOSITES-A NEW EMERGING FAMILY OF ADVANCED COMPOSITES 3 16

GT 8 Shields J.A., Lipetzky P.*, Mueller A.J.** CSM Industries Inc., Cleveland, USA * Pittsburgh Materials Technology, Inc., Pittsburgh, USA " Bechtel Bettis Inc., West Mifflin, USA FRACTURE TOUGHNESS OF 6.4 MM (0.25 INCH) ARC-CAST MOLYBDENUM AND MOLYBDENUM-TZM PLATE AT ROOM TEMPERATURE AND 300°C 4 187

GT 9 Zhang J.L., Zhang Y.M., Lie M.Q., Su Y.Sh., Duan Q.H., Xie H.J. Northwest Institute for Nonferrous Metal Research, Xi'an, Peoples Republic of China THE INFLUENCE OF WORKING WAYS ON THE MICROSTRUCTURE AND PERFORMANCE OF Mo BAR 3 26 XLIV Powder Metallurgical High Performance Materials

GT 10 Stickler C, Knabl W.*, Stickler R.**, Weiss B.*** MIBA-Nuova Mersintersrl, Napoli, Italy * Plansee AG, Reutte, Austria ** Institut für Physikalische Chemie-Werkstoffwissenschaften, Universität Wien, Wien, Austria C ~1 f *** Institut für Materialphysik, Universiät Wien, Wien, Austria CYCLIC BEHAVIOR OF Ta AT LOW TEMPERATURES UNDER LOW STRESSES AND STRAIN RATES 3 34

GT11 Wolske M., Kopp R., Parteder E.* Institute of Metal Forming, RWTH Aachen, Aachen, Germany * Plansee AG, Technology Center, Reutte, Austria C 1 ( INVESTIGATIONS ON THE FORM ABILITY OF A TUNGSTEN ALLOY BY USE OF ACOUSTIC EMISSION J 3 LJ ANALYSIS 3 49

GT 12 Parteder E., Riedel H.*, Sun D.Z.* Plansee AG, Reutte, Austria * Fraunhofer Institute for Mechanics of Materials, Freiburg, Germany SIMULATION OF HOT FORMING PROCESSES OF REFRACTORY METALS USING POROUS METAL PLASTICITY MODELS 3 60

GT13 Schulz G. Wideflow Metal Powder Production and Engineering GmbH, Ohrdruf, Germany ULTRAFINE METAL POWDERS FOR HIGH TEMPERATURE APPLICATIONS MADE BY GAS ATOMIZATION 3 74

GT 14 Leonhardt T., Moore N., Hamister M. Rhenium Alloys Inc., Elyria, USA S 1 b SPHERICAL RHENIUM METAL POWDER 3 85

GT 15 Nutsch G., Blum J., Herrmann A., Schwenk A., Weiss K.H. TU Ilmenau, Fachgebiet Plasma- und Oberflächentechnik, Ilmenau, Germany ERZEUGUNG ULTRAFEINER PULVER UND HERSTELLUNG NANOSTRUKTURIERTER SCHICHTEN MIT DEM THERMISCHEN INDUKTIONSPLASMA 3 97

GT 16 Nutsch G„ Linke P., Zakharian S., Dzur B., Weiss K.H. TU Ilmenau, Fachgebiet Plasma- und Oberflächentechnik, Ilmenau, Germany PULVERBEHANDLUNG UND SPHÄROIDISIERUNG MIT DEM THERMISCHEN INDUKTIONSPLASMA 3 113

GT 17 Schulmeyer W.V., Ortner H.M.* Plansee AG, Reutte, Austria * Darmstadt University of Technology, Inst, of Material Science, Dept. of Chem. Analytics, Darmstadt, Germany MECHANISMS OF THE HYDROGEN REDUCTION OF MOLYBDENUM OXIDES 3 129

GT 18 Bucki J.J., Rozniatowski K., Kurzydlowski K.J. Warsaw University of Technology, Dept. of Materials Science, Warszawa, Poland QUANTITATIVE DESCRIPTION OF THE MICROSTRUCTURE OF SINTERED MATERIALS 3 147

GT 23 Bobrovnichii G.S., Holanda J.N.F. Laboratory of Advanced Materials, UENF-CCT, Campos dos Goytacazes-RJ, Brazil EFFECT OF THE HIGH PRESSURE ON THE DENSIFICATION OF Nb-ATR POWDER COMPACTS 3 178

GT 24 Cathala B., Nicolas G. Cime Bocuze, St.Pierre-en-Faucigny, France CHARACTERISATION OF POWDERS UP TO DATE LAWS OF COMPRESSION 3 186

GT 25 Boginsky L, Reut O., Piatsiushyk Y., Zabolotny O., Kupryianov I. Interbranch Institute for Improvement of Qualification at the BSPA, Minsk, Belarus THE DEVELOPMENT OF PROCESSES OF PRESSING OF ARTICLES FROM POWDERS ON THE BASES OF METALS, CERAMICS AND GRAPHITE 3 197 Powder Metallurgical High Performance Materials XLV

GT 26 Kern A.Y., Lyulko V.G., Kern Y.A. Don State Technological University, Rostov-on-Don, Russia CHARACTERISTICS AND PRODUCTION OF COMPOUND POWDER ARTICLES 3 210

GT27 SemrauW.M. IWS - Ingenieurbureau Wolfgang Semrau, Bremen, Germany LOW TEMPERATURE PROCESSING OF TUNGSTEN-FIBRE HIGH-STRENGTH COMPOSITE 3 219

GT 29 Kenik E.A., Miller M.K., Russell K.F., Bryhan A.J.* * Applied Materials, Santa Clara, Canada Microscopy and Microanalytical Sciences Group, Oak Ridge National Laboratory, Oak Ridge, USA IMPROVEMENTS IN THE DUCTILITY OF MOLYBDENUM WELDMENTS BY ALLOYING ADDITIONS OF Zr, BANDC 3 226

GT 31 Es-Said O.S., Abouriaiy N., Carlen J.C.*, Leonhardt T.*, Garrrmstani H.**, Foyos J., Ocegueda A., Smith T., Omidghaemi S., Suchit R. NSF Research for Undergraduates Program, Loyola Marymount University, Los Angeles, USA * Rhenium Alloys Inc., Elyria, USA " FAMU-FSU College of Engineering, Dept. of Mech. Eng. and Center for Materials Research and Techn., Tallahassee, USA THE EFFECTS OF THE VARIOUS INTERMITTENT ANNEALING TEMPERATURES ON THE MECHANICAL PROPERTIES OF THE 0, 45, AND 90° ROLLED RHENIUM SAMPLES 3 235

GT 33 HyuscHenko A., Pilinevich L. Powder Metallurgy Research Institute with Pilot Plant, Minsk, Belarus POROUS MATERIALS WITH GRADIENT AND BIPOROUS STRUCTURE, METHODS OF THEIR PRODUCTION 3 248

GT 34 Gasik M., Zhang B. Helsinki University of Technology, Espoo, Finland MODELLING OF SINTERING OF FUNCTIONALLY GRADATED MATERIALS 3 262

GT 35 Reut O., Piatsiushyk Y., Makarchuk D., Yakubouski A. Interbranch Institute for Improvement of Qualification at the BSPA, Minsk, Belarus THEORETICAL AND TECHNOLOGICAL FUNDAMENTALS OF PRESSING POROUS POWDER ARTICLES OF THE COMPLEX SHAPE 3 271

GT 36 Piatsiushyk Y., Reut O., Yakubouski A., Boginsky L. Interbranch Institute for Improvement of Qualification at the BSPA, Minsk, Belarus MAIN ASPECTS OF THE THEORY AND TECHNOLOGY OF PRODUCING PERMEABLE MATERIALS WITH THE ORGANIZED POROUS STRUCTURE THROUGH DEFORMATION PROCESSING 3 285

GT 37 Strobl S., Danninger H. Institut für Chemische Technologie anorganischer Stoffe, Techn. Universität Wien, Wien, Austria PM - VERFAHREN ZUR HERSTELLUNG METALLISCHER ZELLULARWERKSTOFFE 3 300

GT 38 Strobl S., Danninger H. Institut für Chemische Technologie anorganischer Stoffe, Techn. Universität Wien, Wien, Austria PRODUCTION AND PROPERTIES OF BRONZE BASED CELLULAR MATERIALS 3 314

GT 39 Fesenko I. Institut for Superhard Materials, Kyiv, Ukraine SINTERING AND THERMAL CONDUCTIVITY OF AIN BASED CERAMICS CONTAINING REFRACTORY COMPOUNDS 3 329

GT 40 Kuzenkova M., Fesenko I., Kisly P. Institute for Superhard Materials, Kyiv, Ukraine INVESTIGATIONS OF VARIOUS METHODS FOR PRODUCTION OF NANOSTRUCTURED AIN CERAMICS 3 333 XLVI Powder Metallurgical High Performance Materials

GT41 Zalite I., Krastins J. Institute of Inorganic Chemistry, Riga Technical University, Salaspils, Latvia PREPARATION AND PROPERTIES OF NANODISPERSE TRANSITION METAL CARBONITRIDES 3 340

GT 42 Weimar P., Prümmer R.* Forschungszentrum Karlsruhe, Institut für Materialforschung, Karlsruhe, Germany * Karlsruhe University, Institut für Werkstoffkunde, Karlsruhe, Germany EXPLOSIVE COMPACTION OF NANOCRYSTALLINE ALUMINA POWDER 3 352

GT 44 Kundas S., Kashko T., Lugscheider E.*, von-Hayn G.*, Hyuschenko A.** Belarusian State University of Informatics and Radioelectronics, Minsk, Belarus * Technical University Aachen, Aachen, Germany ** Research Institute of Powder Metallurgy, Minsk, Belarus

3D-SIMULATION OF RESIDUAL STRESSES IN TBC PLASMA SPRAYED COATING 3 360

GT 46 Borisov Y., Myakota I.*, Polyakov S.* Paton Welding Institute, Kiev, Ukraine * Cherkassy Engineering and Technological Institute, Cherkassy, Ukraine PRODUCTION OF PRESS MOULDS BY PLASMA SPRAY FORMING PROCESS 3 375

GT 47 Schintlmeister A.*, Wilhartitz P. Plansee AG, Reutte, Austria * Plansee AG und Institut für Experimentalphysik, Karl Franzens Universität Graz, Graz, Austria PROTECTION OF MOLYBDENUM AND MOLYBDENUM-ALLOY THIN FILMS AGAINST LOW- TEMPERATURE OXIDATION 4 211

GT 48 Ramminger P., Tessadri R., Krismer R.*, Wilhartitz P.** Institute of Mineralogy, University of Innsbruck, Innsbruck, Austria * Plansee AG, Test Laboratories, Analytics, Reutte, Austria " Plansee AG, Business Unit Thin Film Technology, Reutte, Austria APPLICATION OF HIGH TEMPERATURE DIFFRACTION AS A TOOL OF MATERIAL- CHARACTERISATION AND PRODUCT-OPTIMISATION 4 216

GT 49 Traxler H., Arnold W.", Resch J., Knabl W., Leichtfried G. Plansee AG, Technology Center, Reutte, Austria f * Fraunhofer-Institut für zerstörungsfreie Prüfverfahren (IZFP), Saarbrücken, Germany T, '-O CHARACTERIZATION OF THE DELAMINATION BEHAVIOR OF TUNGSTEN BASED ALLOYS BY MEANS J OF ACOUSTIC EMISSION 3 387 Powder Metallurgical High Performance Materials XLVII Cumulative List of Authors Volume 1-4

Abourialy N. 3 GT 31 Brady M.P. 1 RM 62 Deuerler F. 2 HM 76 Agrawal D.K. 2 HM 90 Bram M. 1 RM 56 Dhers J. 1 RM 40a Agren J. 4 HM 44 Briant C.L. 1 RM 52 Dittrich K.H. 2 HM 77 Ahida Y. 1 RM 71 1 RM 54 Djuricic B. 2 HM 48 Ahmad-Khanlou A. 1 RM 56 Broccardo S. 4 HM 89 Dobrescu M. 1 RM 42 Ahn I.S. 1 RM 86 Brown W.D. 2 HM 100 Doggwiler B. 1 RM 5 Akiyoshi N. 1 RM 23 2 HM 78 Doh J.M. 1 RM 6 1 RM 25 Bryhan A.J. 3 GT 29 DoiY. 1 RM 41 Alamo J. 4 RM 102 Bryskin B. 1 RM 31 Dominguez L. 2 HM 6 4 RM 103 Buchkremer HP. 1 RM 56 Dore F. 1 RM 12 Alber H. 1 RM 90 Buck V. 2 HM 76 Dörr J. 2 HM 83 AllibertC.H. 4 HM 89 Bucki J.J. 3 GT 18 Douy A. 1 RM 82 1 RM 12 Buckman, Jr. R.W. 1 RM 60 Dovydenkov V.A. 4 RM 98 Amato I. 2 HM 50 Bulic F.I. 4 HM 46 Downs J. 1 RM 74 Anderson I.M. 1 RM 62 Buntushkin V.P. 1 RM 84 Dreyer K. 2 HM 1 Andersson G. 2 HM 87 2 HM 40 Ando T. 1 RM 17 Cadoret K. 1 RM 69 2 HM 43 Andreani R. 4 RM 1 1 RM 82 2 HM 55 Andren H.O. 4 HM 44 CaoW. 3 GT 6 Drozdov A.A. 1 RM 85 Andronova T.E. 2 HM 58 2 HM 47 Duan Q.H. 3 GT 9 Anikeev A.I. 2 HM 99 1 RM 14 Duncan R. 1 RM 20 Anikin V.N. 2 HM 99 Carlen J.C. 3 GT 31 Dzur B. 3 GT 16 Antonova A.V. 1 RM 73 Carroll D.F. 2 HM 4 Antropov V.V. 1 RM 48 Castermans L. 2 HM 63 Ederyd S. 2 HM 32 Appel F. 1 RM 70 Castro F. 2 HM 6 Edtmaier C. 1 RM 11 Arnold W. 3 GT 49 Cathala B. 3 GT 24 Eglsäer C. 1 RM 11 Azcona I. 2 HM 6 ChaS.I. 2 HM 16 EidukO.N. 2 HM 5 Chato R. 4 RM 102 Ekroth M. 4 HM 44 Baccino R. 1 RM 12 4 RM 103 Engering G. 2 HM 83 Baranski N. 2 HM 80 Chen H. 2 HM 25 Engqvist H. 2 HM 32 Barbier G. 2 HM 27 Chen J. 2 HM 25 2 HM 92 Bartha L. 2 HM 13 Chen L. 2 HM 55 Epik I. 2 HM 26 4 RM 76 Chen Y. 1 RM 68 Es-Said O.S. 3 GT 31 1 RM 77 Chen Y.L. 4 RM 7 Ettmayer P. 2 HM 37 Basyleva O.A. 1 RM 84 Cheng J. 2 HM 90 2 HM 55 Batawi E. 1 RM 5 Clemens H. 1 RM 70 Bausch S. 2 HM 93 Clemente G. 4 RM 103 Falkovsky V.A. 2 HM 5 Bennett E.G. 2 HM 30 Collins J.L. 2 HM 94 2 HM 11 Bentham R. 1 RM 67 Cornish L.A. 4 HM 89 Ferguson S. 4 RM 89 Beste U. 2 HM 92 Fesenko I. 3 GT 39 Bewlay B.P. 1 RM 54 da Costa F.A. 1 RM 24 3 GT 40 Bizaoui M. 4 HM 49 da Silva A.G.P. 1 RM 24 Filgueira M. 1 RM 50 Blagoveschenski Y. 2 HM 11 Dabhade V.V. 4 RM 57 Fischer E. 1 RM 65 Blum J. 3 GT 15 Dai M.J. 2 HM 73 Fleischer W. 2 HM 80 Bobrovnichii G.S. 3 GT 23 Damm S. 2 HM 83 Föjer C. 2 HM 87 Bock A. 4 HM 2 Danninger H. 3 GT 37 Foyos J. 3 GT 31 Boginsky L. 3 GT 25 3 GT 38 Frage N. 4 HM 49 3 GT 36 Dariel M.P. 4 HM 49 1 RM 26 Böhm A. 1 RM 13 1 RM 26 Franco T. 4 RM 19 1 RM 72 DaubH.W. 2 HM 40 Friedrichs K. 4 HM 88 Bondarenko V. 2 HM 26 2 HM 55 Froumin N. 1 RM 26 Borisenko E.B. 1 RM 55 de Cachard J. 1 RM 69 Frykholm R. 4 HM 44 Borisov Y. 3 GT 46 1 RM 82 Fukumitu T. 1 RM 35 2 HM 63 De Cola P. 2 HM 50 Furusaka M. 4 RM 4 Borisova A. 2 HM 63 Dedkov P.Y. 4 RM 100 Bourgeois L. 2 HM 27 Derra G. 1 RM 65 Gallet D. 1 RM 40a XLVIII Powder Metallurgical High Performance Materials Cumulative List of Authors Volume 1-4

Garcia J. 2 HM 40 Hennet L. 1 RM 82 Kazior J. 1 RM 46 2 HM 59 Herrmann A. 3 GT 15 Kem A.Y. 3 GT 26 Garmestani H. 3 GT 31 Herrmann M. 2 HM 17 Kem Y.A. 3 GT 26 Gasent J.L. 4 RM 102 Hiraoka Y. 1 RM 23 Kenik E.A. 3 GT 29 4 RM 103 1 RM 34 Khodoss I.I. 4 RM 44 1 RM 35 KhokhlovA. 2 HM 11 Gasik M.M. 3 GT 34 Hofer F. 2 HM 9 KhokhlovA.M. 2 HM 5 2 HM 41 Hogmark S. 2 HM 32 Kieback B. 3 GT 2 Gasparotto M. 4 RM 1 Holanda J.N.F. 3 GT 23 1 RM 3 GeC.C. 3 GT 6 Holubar P. 4 HM 65 1 RM 72 2 HM 47 Holzschuh H. 2 HM 62 Kienitz S. 2 HM 103 2 HM 48 Hong S.H. 2 HM 16 Kikuchi K. 4 RM 4 1 RM 14 Hoppe S. 2 HM 80 Kikuchi N. 1 RM 17 Gee M.G. 2 HM 30 Horacsek 0. 1 RM 77 Kim B.K. 2 HM 16 4 HM 84 Hornig T. 2 HM 103 Kim D.G. 1 RM 10 Gerasimov V.V. 1 RM 85 Huang B.Y. 2 HM 28 Kim H.E. 2 HM 36 German R.M. 4 GT 4 HuberT. 4 RM 16 Kim I. 4 RM 89 1 RM 92 Hühsam A. 2 HM 83 Kim J.S. 1 RM 10 Gestrich T. 2 HM 1 Hurevich V. 4 GT 43 1 RM 86 Gialanella S. 1 RM 46 3 GT 44 Kim M.J. 1 RM 6 Gille G. 2 HM 1 KimS. 2 HM 36 GlatzW. 1 RM 5 Igarashi T. 4 RM 4 Kim Y.D. 1 RM 10 GlushkovV.N. 2 HM 11 1 RM 41 1 RM 86 2 HM 5 Ikegaya A. 2 HM 38 Kimbara A. 1 RM 17 2 HM 99 llyasov V.V. 4 HM 33 King H.W. 4 RM 89 Gnesin B. 3 GT 19 llyuschenko A. 3 GT 33 Kireiko V.V. 3 GT 19 1 RM 55 3 GT 44 Kishino J. 2 HM 7 Godkhindi M.M. 4 RM 57 InoueT. 1 RM 23 Kisly P. 3 GT 40 Gomes U.U. 1 RM 24 Kitamura K. 2 HM 34 Goncharuck V.A. 2 HM 10 Jackson M.R. 1 RM 54 KladlerG. 2 HM 37 Gonzalez-Rodrig. G. 3 GT 21 Jacobson S. 2 HM 92 Klauss H.J. 1 RM 13 Groffils C. 2 HM 27 Jain H. 2 HM 90 Klocke F. 2 HM 103 Grogger W. 2 HM 9 Jang Y.I. 1 RM 86 2 HM 93 Grosz T. 4 RM 76 Janousek M. 1 RM 5 Kluwe H. 2 HM 76 Gurjiyants P.A. 1 RM 55 Jansson B. 4 HM 44 Klyachko L. 2 HM 11 Gursan S. 4 RM 89 2 HM 87 Klyachko L.I. 2 HM 5 Gusarov A. 4 GT 43 Järvi J. 2 HM 4 Knabl W. 3 GT 10 Gutmanas E.Y. 4 RM 57 JilekM. 4 HM 65 3 GT 49 Joensson M. 1 RM 3 Knünz G. 2 HM 3 HaG.H. 2 HM 16 Jones A.R. 4 RM 7 Knüwer M. 1 RM 9 Hamar-Thibault S. 2 HM 15 Jongbloed R. 2 HM 63 Kobayashi M. 2 HM 34 4 HM 89 Jung J.P. 1 RM 6 Kohda K. 1 RM 25 Hamister M. 3 GT 14 Kolchemanov N.A. 4 RM 100 Han S.H. 2 HM 36 Kablov E.N. 1 RM 84 Kolchin A.A. 1 RM 73 Hanada S. 1 RM 61 Kakazey M. 3 GT 21 König U. 4 HM 97 Harmat P. 2 HM 13 Kang S.H. 1 RM 51 Köpf A. 2 HM 74 4 RM 76 Karpov M.I. 4 RM 44 Kopp R. 3 GT 11 Haubner R. 4 HM 101 Kasanskaya N.K. 1 RM 84 Korb G. 2 HM 37 2 HM 74 Kashko T. 4 GT 43 4 HM 46 2 HM 75 3 GT 44 2 HM 48 4 HM 79 Kaskiala M. 2 HM 41 Korzhov V.P. 1 RM 73 Hayashi K. 2 HM 34 Kassel D. 2 HM 40 Kotsis I. 2 HM 13 Heilmaier M. 1 RM 13 C M HM 55 Krai C. 2 HM 43 Heiss M. 2 HM 98 Kathrein M. 2 HM 98 Krastins J. 3 GT 41 Hempelmann R. 2 HM 3 Kawai M. 4 RM 4 Krieg T. 2 HM 103 Henne R. 4 RM 19 Kazanskaya N.K. 1 RM 85 Powder Metallurgical High Performance Materials XLIX Cumulative List of Authors Volume 1-4

Krismer R. 4 GT 48 LiuX. 1 RM 47 MünzW.D. 4 HM 96 1 RM 15 Liu Z.Y. 2 HM 73 Myakota I. 3 GT 46 Kuang T.C. 2 HM 73 Lomberg B.S. 1 RM 85 Kumar P. 1 RM 58 Lopez J.L. 4 RM 103 Nagae M. 1 RM 34 Kundas S. 4 GT 43 Lu P.K. 1 RM 92 1 RM 35 3 GT 44 Lucaci M. 1 RM 87 Nagae T. 1 RM 41 Kunen H. 2 HM 80 Lugscheider E. 3 GT 44 Nagata S. 1 RM 17 Kupryianov I. 3 GT 25 2 HM 103 Nagy D.R. 4 RM 89 Kurishita H. 4 RM 4 Lukashova N.M. 2 HM 5 Nakada K. 1 RM 25 1 RM 41 Lux B. 4 HM 35 Nakayama M. 1 RM 25 Kurzydlowski K.J. 3 GT 18 4 HM 57 NamT.H. 1 RM 51 Kusano E. 1 RM 17 2 HM 74 Nasu S. 1 RM 17 Kuzenkova M. 3 GT 40 2 HM 75 Natter H. 2 HM 3 Kwon Y.S. 1 RM 86 Luyckx S. 2 HM 10 Nganbe M. 1 RM 13 4 HM 89 Nicolas G. 3 GT 24 Labusca-Lazar. E. 1 RM 101 Luypaert P.J. 2 HM 27 1 RM 93 Lackner A. 2 HM 15 Lyulko A.V. 4 HM 33 NieZ. 1 RM 68 2 HM 29 LyulkoV.G. 3 GT 26 1 RM 83 2 HM 3 1 RM 48 Ning A. 1 RM 47 2 HM 31 Nojkina A.V. 4 RM 100 2 HM 9 Mabuchi M. 1 RM 41 Nomura H. 2 HM 7 Lackner K. 4 RM 1 Mahalingam M. 4 RM 8 Nomura N. 1 RM 61 Laczko L. 2 HM 13 Makarchuk D. 3 GT 35 Northrop J.T. 2 HM 10 Lahres M. 2 HM 83 Makarov P.V. 1 RM 59 Nutsch G. 3 GT 15 Lang M. 4 RM 19 MalsheA.P. 2 HM 100 3 GT 16 Lartigue J.F. 1 RM 12 2 HM 78 LayS. 2 HM 15 Martinez L. 1 RM 69 Ocegueda A. 3 GT 31 Leichtfried G. 3 GT 49 1 RM 82 Oehring M. 1 RM 70 1 RM 29 Martinz H.P. 1 RM 29 Oelsner D. 2 HM 77 1 RM 66 1 RM 66 Ohhashi K. 1 RM 17 Leitner G. 2 HM 1 Martynova L. 2 HM 26 Ohriner E.K. 1 RM 62 Lembke M.I. 4 HM 96 Matej J. 1 RM 29 Okamoto K. 4 RM 4 Lences Z. 4 HM 46 Matsubara H. 2 HM 7 Oleinikov D.V. 1 RM 48 Lengauer W. 2 HM 40 Medved N.V. 4 RM 44 Olsson F. 1 RM 20 2 HM 43 Meinhardt H. 1 RM 9 Omidghaemi S. 3 GT 31 2 HM 55 Mertens T. 2 HM 83 Ordanyan S.S. 2 HM 58 4 HM 56 Mileiko S.T. 1 RM 73 Ordonez A. 2 HM 6 2 HM 59 Miller M.K. 3 GT 29 OrtnerH.M. 3 GT 17 Leonhardt T. 3 GT 14 Miller U. 4 RM 7 3 GT 31 Millot F. 1 RM 69 Pantelejev I.B. 2 HM 58 1 RM 74 1 RM 82 Papworth A.J. 2 HM 90 Levoy R. 1 RM 40a Milman Y.V. 2 HM 10 Park J.K. 2 HM 36 Lewis D.B. 4 HM 96 Missiaen J.M. 2 HM 31 1 RM 6 LiJ. 3 GT 6 Mits I. 2 HM 63 Parteder E. 3 GT 11 2 HM 47 Mitterer C. 2 HM 98 3 GT 12 1 RM 14 Molinari A. 1 RM 46 Pavlotskaya E. 2 HM 26 Li J.F. 4 RM 4 Moon A. 1 RM 17 Pecher U. 2 HM 83 Li J.T. 2 HM 48 Moore N. 3 GT 14 Peltan P. 1 RM 46 LiY. 2 HM 25 Mori G. 2 HM 29 Piatsiushyk Y. 3 GT 25 2 HM 72 Moriguchi H. 2 HM 38 3 GT 35 Licheron M. 1 RM 82 Morita S. 2 HM 18 3 GT 36 Lie M.Q. 3 GT 9 Morley N. 1 RM 67 Pick M. 4 RM 1 Linke P. 3 GT 16 Morrell R. 4 HM 84 Pieczonka T. 1 RM 46 Lipetzky P. 4 GT 8 Motoiu P. 1 RM 96 Pilinevich L. 3 GT 33 Lisovsky A.F. 2 HM 23 Mueller A.J. 4 GT 8 Pinatti D.G. 1 RM 50 LiuS. 2 HM 72 1 RM 60 Piscone F. 2 HM 50 Powder Metallurgical High Performance Materials Cumulative List of Authors Volume 1-4

Plankensteiner A. 1 RM 15 Santos I. 4 RM 102 Smurov I. 4 GT 43 Polcik P. 1 RM 40a 4 RM 103 Sommer M. 4 HM 12 Polyakov S. 3 GT 46 Sanz-Tudanca C. 4 RM 103 Song J.B. 2 HM 73 Povarova K.B. 1 RM 59 Sarathy D. 2 HM 85 Sottke V. 4 HM 97 1 RM 73 Sarkissyan N.S. 1 RM 73 Sreckovic T. 3 GT 21 1 RM 84 Savich V. 3 GT 33 Stewart M. 2 HM 30 1 RM 85 Sayama Y. 2 HM 18 Stickler C. 3 GT 10 Prümmer R. 3 GT 42 Schalunov E.P. 4 RM 98 Stickler R. 3 GT 10 PutS. 2 HM 27 Schaper M.K. 1 RM 13 Stockmann Y. 2 HM 80 2 HM 42 Schatte J. 2 HM 98 Stöver D. 1 RM 56 Schedle D. 4 RM 16 Strobl S. 3 GT 37 Ragulya A.V. 2 HM 17 Schedler B. 1 RM 15 3 GT 38 Rainer F. 1 RM 15 4 RM 16 Su Y.Sh. 3 GT 9 4 RM 16 Scheiber K.H. 4 RM 16 Subramanian P.R. 1 RM 54 RakA. 3 GT 33 SchillerG. 4 RM 19 1 RM 54 Rama Mohan T.R. 4 RM 57 Schintlmeister A. 4 GT 47 Suchit R. 3 GT 31 Ramakrishnan P. 4 RM 57 SchintlmeisterW. 2 HM 98 Sudmeijer K. 1 RM 90 Ramminger P. 4 GT 48 Schleinkofer U. 2 HM 98 Sugimata E. 1 RM 17 Rao D. 1 RM 43 Schlichtherle S. 1 RM 66 Sugimoto T. 1 RM 71 Reglero V. 4 RM 102 Schmid L 2 HM 31 Sun D.Z. 3 GT 12 4 RM 103 Schmidt J. 2 HM 1 Sundaram N. 2 HM 78 Resch J. 3 GT 49 Schmiler B. 2 HM 76 Suzuki T. 1 RM 61 Retter B. 1 RM 66 Scholl R. 1 RM 72 Svensson S.A. 1 RM 20 Reut 0. 3 GT 25 Schön A. 4 HM 35 Szesny B. 2 HM 1 3 GT 35 Schretter M. 2 HM 29 3 GT 36 Schubert W.D. 4 HM 12 Tabernig B. 1 RM 90 Reznikov G. 4 RM 36 4 HM 35 Tabersky R. 4 HM 97 Richter V. 2 HM 53 4 HM 57 Takada J. 1 RM 34 2 HM 82 SchulmeyerW.V. 3 GT 17 1 RM 35 Riedel H. 3 GT 12 Schulz G. 3 GT 13 Takagi K.I. 2 HM 52 Ristic M.M. 3 GT 21 Schwenk A. 3 GT 15 Takemoto Y. 1 RM 34 Rodrigo J. 4 RM 103 Schwenke G.K. 2 HM 86 Takida T. 1 RM 41 Rodrigo J.M. 4 RM 102 Scrivani A. 2 HM 39 Tanase T. 2 HM 7 Roebuck B. 2 HM 30 Selcuk C. 1 RM 67 Tang H. 1 RM 21 4 HM 84 Semrau W.M. 3 GT 27 Tang X. 4 HM 10 Rofner R. 2 HM 98 Serebryakov A.V. 1 RM 73 Teng Xiuren 1 RM 21 Rosso M. 2 HM 39 Sfat C. 2 HM 21 Teraoka T. 1 RM 71 2 HM 70 2 HM 22 Terauchi S. 1 RM 71 1 RM 96 1 RM 95 Terentieva V. 1 RM 97 Rosta L. 4 RM 76 Shalunov E.P. 1 RM 48 Tessadri R. 4 GT 48 Rozniatowski K. 3 GT 18 Shavlovsky E. 2 HM 63 ThierfelderW. 1 RM 90 Rubinovich L. 1 RM 26 Shen W. 3 GT 6 Toth R.E. 2 HM 37 Ruckdäschel R. 4 RM 19 Sherman A. 2 HM 37 Traxler H. 3 GT 49 Ruiz Urien I. 4 RM 102 Shevchenko A.D. 2 HM 69 Tsuzuki K. 2 HM 38 4 RM 103 Shields J.A. 4 GT 8 Russell K.F. 3 GT 29 1 RM 60 Ucakar V. 2 HM 40 Russell W.C. 2 HM 100 Shim W.S. 1 RM 10 2 HM 43 2 HM 78 Shimizu H. 2 HM 18 Ugues D. 2 HM 70 Ryzhkin A.A. 4 HM 33 Shin S.G. 2 HM 7 UhlenhutH. 1 RM 58 Shinkuma T. 1 RM 71 Uhrenius B. 2 HM 32 Saito A. 4 RM 4 SikkaV.K. 1 RM 62 Upadhyaya A. 2 HM 85 Sajgalik P. 4 HM 46 Sima M. 4 HM 65 Upadhyaya G.S. 1 RM 43 Salvarani L. 2 HM 39 Simonov V.S. 4 RM 98 Sanches Mondr. J.J. 3 GT 21 Smid I. 2 HM 37 Vallauri D. 2 HM 50 Sanchez J.M. 2 HM 6 Smith I.J. 4 HM 96 Valle A. 2 HM 70 Smith T. 3 GT 31 Powder Metallurgical Hiqh Performance Materials LI Cumulative List of Authors Volume 1-4

Van den Berg H. 2 HM 1 Wolske M. 3 GT 11 2 HM 40 Wood J.V. 1 RM 67 4 HM 97 Wright I.G. 1 RM 62 Van der Biest O. 2 HM 27 WuA. 2 HM 47 2 HM 42 1 RM 14 van der Kolk G.J. 2 HM 80 Van Hees G. 1 RM 65 Xie H.J. 3 GT 9 Vasile E. 1 RM 87 Yakubouski A. 3 GT 35 Vasile T. 2 HM 21 3 GT 36 2 HM 22 Yamasaki Y. 2 HM 52 1 RM 42 YedaveS.N. 2 HM 100 1 RM 95 2 HM 78 Vasilescu I. 2 HM 22 YiD. 2 HM 72 Vasilescu M. 2 HM 21 YiW. 1 RM 92 2 HM 22 Yildiz M. 4 RM 89 1 RM 95 YinW. 1 RM 21 Veillet D. 1 RM 69 Ylikerälä J. 2 HM 41 Velasco T. 4 RM 102 Yonetsu M. 2 HM 52 4 RM 103 Yoo M.K. 1 RM 23 Vereschaka A. 2 HM 99 Yoshii K. 1 RM 17 Vidu C. 1 RM 87 Yoshio T. 1 RM 34 Viswanathan V. 4 RM 8 1 RM 35 Vladimirova M.A. 2 HM 58 YuZ. 2 HM 72 Vlasova M. 3 GT 21 Vleugels J. 2 HM 27 Zabernig A. 4 RM 16 2 HM 42 Zabolotny O. 3 GT 25 VnukovV.I. 4 RM 44 Zafiu G. 2 HM 21 Volders I. 2 HM 27 2 HM 22 Volkov K.G. 4 RM 44 Zakharian S. 3 GT 16 Voltz M. 1 RM 93 Zalite I. 3 GT 41 Von Ruthendorf M. 2 HM 53 2 HM 58 von-Hayn G. 3 GT 44 Zarauz J. 4 RM 102 4 RM 103 Wagner G. 2 HM 85 Zeiler B. 4 HM 2 Walls C. 1 RM 62 Zgalat-Loz. O.B. 2 HM 17 Wang J. 1 RM 83 Zhang B. 3 GT 34 Warbichler P. 4 HM 12 2 HM 41 2 HM 9 Zhang D. 2 HM 25 Warren J. 4 RM 36 Zhang H. 1 RM 47 Watanabe R. 4 RM 4 Zhang J. 1 RM 83 Weaver M.L. 1 RM 62 Zhang J.L. 3 GT 9 Weimar P. 3 GT 42 Zhang Y.M. 3 GT 9 Weiss B. 3 GT 10 Zhao C. 2 HM 27 Weiss K.H. 3 GT 15 Zhou J.C. 2 HM 28 3 GT 16 Zhou K.S. 2 HM 73 Weissenbacher R. 4 HM 79 Zhou M. 1 RM 68 Westmoreland G. 1 RM 62 1 RM 83 Westphal H. 4 HM 97 Zinyana J. 4 HM 89 Wichmann K.H. 1 RM 9 Zissis G. 1 RM 64 Wieters K.P. 3 GT 2 Zitter H. 2 HM 29 Wilhartitz P. 4 GT 47 Zobetz E. 4 HM 12 4 GT 48 ZuevA.P. 3 GT 19 Williams D.B. 2 HM 90 ZuoT. 1 RM 68 Witcomb M.J. 4 HM 89 1 RM 83 Wittmann B. 4 HM 57 Wollein B. 4 HM 56 LII Powder Metallurgical High Performance Materials

ERRATUM

Volume 3: General Topics

GT24 Cathala B., Nicolas G. Cime Bocuze, St. Pierre-en-Faucigny, France

CHARACTERISATION OF POWDERS UP TO DATE LAWS OF COMPRESSION

The relationship p. 191-192 and 194 is written as follows:

|(d-do) / (dth/do) = KPri

The correct relationship is:

l(d-do) / (dth-do) = KPn AT0200467

R. Watanabe et al. RM 4

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

DEVELOPMENT OF SOLID TARGET FOR SPALLATION NEUTRON SOURCE

R. WATANABE13, J.-F. LI1', A.SAITO1}, H. KURISHITA0, T. IGARASHI2) K. OKAMOTO3', K. KDCUCHI4), M. FURUSAKA5), M. KAWAI5) i:>Tohoku University, Sendai, Japan, 2:>Sumitomo Electric Co., Tokyo, Japan 3)Allied Material Corp. Tokyo, Japan 4)Japan Atomic Energy Research Institute, Tokai, Japan 5)High Energy Accelerator Research Organization, Tsukuba, Japan

SUMMARY Manufacturing of tungsten based materials for the target of spallation neutron source for the next generation is described. A sintered tungsten block cladded with tantalum sheet and a stainless-steel bonded heavy alloy were selected as candidates for the long-life target materials having enough resistance to corrosion in hot water and less intense susceptibility to radiation damage. The process conditions and pertaining metallurgical problems for the target fabrication are briefly described and discussed. The tantalum-clad tungsten target has been used in the current proton-accelerator-driven pulsed neutron source of the High Energy Accelerator Research Organization, Tsukuba, Japan, and has proved to give neutron yield of 20% higher than that of the current Ta target. The stainless-steel bonded heavy alloy still needs further investigation on the mechanical and corrosion properties.

KEYWORDS: spallation neutron source, solid target, radiation damage, corrosion resistance, tantalum-clad tungsten, HEP, heavy alloy, stainless steel binder

1. INTRODUCTION

The construction of high intensity proton accelerators is being planned in Japan to pursue frontier science in the fields of particle physics, nuclear physics, materials science and life science, as well as in nuclear technology1-1'1. This is a joint project of the Japan Atomic Energy Research Institute (JAERI), Tokai, and the High Energy Accelerator Research Organization (KEK: Japanese abbreviation), Tsukuba, Japan, and is divided into two phases. The first is the construction of proton accelerator complex consisted of 400 MeV Linac(linear accelerator) and 3-50GeV synchrotrons with highest level of beam power, which is accompanied by facilities of spallation neutron source and mason probe for materials and life science, and approved this year (in 2001) by Japanese government and will be completed in six years by 2006. The second is the construction of accelerator driven system (ADS) for nuclear transmutation and neutrino science, which will start in 2007. Both of the neutron scattering facility and ADS need high-flux spallation neutron source composed of RM 4 Ft. Watanabe et al.

15'* International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

metal target and cooling system. One of the key technologies to promote this construction plan is the development of the target that generates high-flux neutrons by the irradiation of high-energy proton beams(2). To develop a long-life target a research project was organized by KEK, under the support of Japanese Ministry of Education and Science, in which people from seven Japanese universities and the Japan Atomic Energy Research Institute have participated together with the KEK researchers. The technical requirements for the spallation target are(3): (1) high neutron yield, (2) structural and thermal stability to withstand high energy proton beam, (3) radiation damage resistant, (4) corrosion resistant in cooling water (actually heavy water is used as a coolant), (5) compatibility with cooling pipe materials (usually stainless steel is used). High-density metals are candidate target materials because they have a high neutron yield. Liquid target, such as mercury and Pb-Bi alloy, is the first option because they will simultaneously serve as a coolant, and solid target is the second option. But in the present project the research efforts were focused on the solid target because the liquid target has still many unknown factors to be investigated in the target design, while solid target is currently in use in the neutron source with a low power proton beam. As a target material tungsten was selected because it has a high neutron production ability, and its post irradiation lifetime is much shorter compared to its neighboring elements (3>4). However, tungsten shows poor corrosion resistance against water coolants due to the formation of ragged tungsten hydroxide (5), and intense susceptibility to radiation embrittlement. Two approaches for the solid target fabrication will be described in this paper: one is a stainless-steel-bonded tungsten alloy and the other is tantalum-clad tungsten. The former is fabricated by liquid phase sintering, the latter by hot isostatic pressing (HIP).

2. TUNGSTEN-STAINLESS STEEL HEAVY ALLOY

2.1 Specimen preparation Critical issues for the spallation target materials are toughness degradation due to intense radiation damage and corrosion in hot water. Pure tungsten is known to corrode easily in water during radiation (6>7). Tungsten heavy alloy (W-Ni-Cu-Fe) is also subjected to corrosion due to the formation of porous film of tungsten- and iron hydroxides. Kawamura (5) has reported that the corrosion of the tungsten heavy alloy (97W-Ni-Cu-Fe) in a wet condition is remarkably improved by the addition of small amount of chromium. It seemed that the chromium addition improved the corrosion resistance of the binder phase, a stainless steel of 304 type was used as corrosion resistant binder. R. Watanabe et al. RM4

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Tungsten powders of 5 and 9 urn in average size and a stainless steel powder (SUS304L: 0.017C, 08Si, O.IMn, 0.025P, 0.005S, 18.2Cr, 10.8Ni, Fe(bal)) of 8.2 ^im were used as starting materials, as shown in Fig. 1. The stainless steel powder was water-atomized and has an oxygen content of 4300ppm mainly as a surface oxide. They were mixed in weight % of 93 and 97W by wet milling. The dried and disintegrated powder was die-compacted at 150MPa to a cylindrical form of 10mm in diameter and 15mm in height, and subsequently CIPed at 200MPa to eliminate density inhomogeneity generated by die wall friction. Vacuum sintering was conducted in a temperature range of 1350-1550C and time interval of 5 min-5h in a tungsten-mesh heater furnace using boron nitride crucible as a sample holder. The sintered compacts were HIPed to eliminate residual pores at 200MPa and 1500C. Some samples were liquid-phase-sintered in hydrogen, or directly HIPed below the melting point of the stainless steel (1400-1450C) to see microstructure variations.

Fig. 1 Tungsten and stainless powders used in the present investigation (a)W powder, 5 urn, (b) Stainless steel powder, 8.2 u.m

2.2 Microstructure Figure 2 shows macro views of the cross section of the sample vacuum sintered lh at 1450C. Three zones having different microstructures can be seen in the micrograph. They are mono-phase surface layer, dense intermediate layer and porous core. The surface shell structure would have been formed by the vaporization of binder phase during liquid phase sintering. Pore elimination in the center core is supposed to be suppressed by the formation of this dense shell structure. On the other hand a homogeneously distributed heavy alloy structure was obtained by hydrogen sintering. RM4 R. Watanabe et al.

15'" Internationa! Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Fig.2 Cross section of the 97W-Stainless steel heavy alloy vacuum sintered for 1 h at 1450C

Figure 3 shows typical microstructures of vacuum and hydrogen sintered samples that were HIP-treated at 1500C after sintering at 1450C. Both exhibit typical heavy alloy type microstructure. The microphotograph of the vacuum-sintered sample was taken from the binder-depleted intermediate zone where small pores still remain mainly in the binder phase. On the other hand in the binder phase of the hydrogen sintered sample angular voids are seen to exist, which would be solidification blow holes formed during cooling from sintering temperature. The relative densities of the

(a) (b) Fig. 3 Typical heavy alloy structures of the specimens HEP-treated for 3 h at 1500C after normal sintering. (a)Vacuum sintered for lh at 1450C, (b)Hydrogen sintered for lh at 1450C. R. Watanabe et al. RM 4

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HEP-treated samples did not reach 100%(about 98% as determined by image analysis) because of such residual pores. Composition analysis by EPMA showed that W dissolved into the binder phase by 72wt%, whereas iron and chromium in the binder phase were found to dissolve into W phase by about 1 % or less. The densities of se dissolution behaviors must have influenced the densification and microstructure evolution during sintering. Judging from the phase diagrams of W-Fe, W-Ni, and W-Cr, tungsten dissolution into binder phase by several tens of wt% would not increase the melting point of the binder. But the dissolution above these values would increase the melting point, and it is likely to occur sintering retardation and microstructural change with increasing sintering time. This situation was actually observed in the vacuum sintering of 93W-7stainless steel at 1500C as shown in Fig. 4. The sintering structure was observed to change with sintering time, a dispersion structure in the first stage, a coagulated structure in the middle stage and again dispersion structure in the final stage. This suggests that along with the increase in the melting point due to the W-dissolution into binder phase the dihedral angle, i.e. the interface energies, also changed during sintering. The complete densification and a typical heavy alloy structure would require sintering temperature of higher than 1500C for this alloy. Considering that the composition of the binder phase changes to a great extent by the dissolution of W, HIP consolidation at 1300C under the melting point of stainless steel was conducted. The microstructure of the HEP-consolidated compact is shown in Fig. 5, which exhibits a skeletal structure of tungsten phase. This HIPed alloy possesses a microstructural homogeneity, and is to be studied as the alternative to the liquid-phase-sintered alloy. The binder content of the solid state consolidated compact was measured to be 25.3 vol%, a little smaller than the liquid-phase-sintered compact with binder content of 26.9 vol%.

; , " * .". L . . . M '

• - 20p.in <;

Fig. 4 Microstructure change in 93 W-7stainless steel heavy alloy during sintering at 1500C in vacuum, (left) hold for 0 min, (right) hold for 20 min. RM4 R. Watanabe et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

50(am

Fig. 5 A 93W-7stainless steel alloy HEP-consolidated at 1300C and 180 MPa.

2.3 Thermal and mechanical properties The thermal conductivity of the present 93 W alloy was measured by laser flash method (Netzsch LFA427), as shown in Fig. 6, along with the data for tungsten and stainless steel. The present alloy has lower thermal conductivity as compared with

93%W (predicted value)

93%W (measured value)

SUS 304 20

0 200 400 600 800 1000 1200 Temperature (°C)

Fig. 6 Thermal conductivity of the present 93W-7stainless steel alloy as compared with those of tungsten and stainless steel. R. Watanabe et al. RM4

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reurte (2001), Vol. 4

the conventional alloy, and its temperature dependence is rather similar to the stainless steel, that is, the thermal conductivity increases with temperature. These data will be used in the thermal design of the spallation target. In general, high thermal conductivity is preferred to reduce the thermal stress that will be generated by proton bombardment in the pulse neutron generator. Preliminary calculation by finite element method has shown that this thermal stress does not exceed 200 MPa^8 ^ , but nevertheless it will be critical for the present alloy under the irradiation. The strength and toughness of the present alloy was evaluated by small punch test(9>10) with disk-shaped specimens (see Fig. 7). The stress is defined by Eq. (1), (but strictly only for the onset of loading when the plain contact is present between the puncher and the specimen).

1-v2 h2 -(1 + v) (1) 2m where, amax is the nominal maximum fracture stress, P is load at fracture, t is specimen thickness, v is Poisson's ratio, a and b are radii of the puncher and the loading area, respectively. Figure 8 shows the load-deflection curves for the

puncher 200 180 HIPedat 1300°C upper die Vac-Sintered at 145O°C. 160 then H[Ps

Trffr 23~ t

Fig. 7 Schematic illustration of small punch Fig.8 Load-deflection curves for the set-up and the loading condition. liquid- phase-sintered 93W-stainless steel and solid- state consolidated alloys. RM 4 R. Watanabe et al.

15:h International Piansee Seminar, Eds. G. Kneringer, P. Ro'dhammer and H. Wildner, Plansee Holding AG, Reutte (2001J, Vol. 4

liquid-phase-sintered 93W- stainless steel and solid-state consolidated alloys. The liquid-phase-sintered specimen exhibits serrated deflection curve that indicates a multiple mode of crack propagation, whereas the solid-state consolidated specimen is seen to be brittle. The multiple-mode fracture observed in the liquid-phase sintered alloy would be preferable to alleviate the radiation embrittlement.

3. TANTALUM-CLAD TUNGSTEN

3.1 HIP cladding of tantalum onto tungsten(U) Tungsten is the first candidate for the solid target because of high neutron yield, but it is very corrosive in hot water, particularly under proton radiation(6'7). Therefore, cladding tungsten with corrosion-resistant metals is necessary for the long-life target with high neutron yield. Tantalum has been considered to be the best material for the cladding because it is corrosion resistant, and has relatively high neutron yield and better metallurgical compatibility with tungsten compared to the other corrosion-resistant metals, such as stainless steel, zircalloy, niobium and so on. ISIS(12 ^as developed a tantalum-cladded tungsten target by HIP at 1400C and at a pressure of 1.45GPa. In the present investigation HIP cladding under conventional HIP pressure (<200MPa) has been studied for the practical application of this method to a large-sized target of the order of about 100 mm. The cladded target had also been expected for use in the current neutron source installed in the 500 MeV synchrotron at KEK, Tsukuba, Japan. The sintered and hot-rolled tungsten block that had been canned with tantalum coverage of 0.6mm thick vacuum-sealed by electron beam welding was HEPed at 1400-1800C at a pressure of 180 MPa.

3.2 Evaluation of bonding interface The bonding integrity of the interface between the tantalum overlayer and the tungsten block is of first concern for the spallation target because bonding defects, if any, would result in peeling off the corrosion resistant surface layer, thus leading to shortening of the target life. Raising HIPing temperature is effective in endowing the interface with high bonding strength. But the recrystallization and grain coarsening of tungsten and tantalum may be enhanced simultaneously during HEP at high temperature, which would result in the decrease in the fracture toughness of the target. It is therefore necessary to optimize HIP cladding temperature. In the present study, the bonding strength of the W-Ta interface was evaluated using the small punch test(SP test), and the diffusion thickness across the interface was measured by EPMA. The SP samples were prepared by the following procedures. First, tungsten and tantalum billets with a semicircle cross section (radius= 5 mm) were sampled from the same kind of materials as the target. The rectangular longitudinal cross section was polished to the same finish as the tantalum can and tungsten block for the R. Watanabe et al. RM4 5" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

targets. The polished sections of the both billets were put together and inserted into a tantalum tube (inner diameter = 10 mm, thickness = 1 mm). The tube was vacuum-sealed by electron beam welding, and HIPed under the same conditions as in the HIP cladding. Finally, the HIPed sample was sliced into thin disk specimens, one half of which is tungsten and another half is tantalum. After polishing the disk surface, the specimens were subjected to SP testing, in which the load was applied to the specimen center, i.e., at the center of the bonded interface. Figure 9 shows typical curves of the SP test of the bonded specimens. For all the specimens, a load drop was observed before the eventual fracture. By SEM observation, it was found that such a load drop corresponded to the crack initiation. Regardless of the HDPing temperature, crack was formed and propagated within tungsten, as shown in Fig. 10. This observation suggests that the interface formed under the present HIPing conditions was well bonded, with a bonding strength at least higher than the fracture strength of the tungsten block. However, when the HIPing temperature exceeded 1500 C, the load for the crack initiation, as well as the fracture strength, decreased due to recrystallization. Therefore, the optimal HIPing temperature was determined to be 1500C. The diffusion thickness formed at this temperature was measured by EMPA to be 3.5 jam.

700

* Test interruptted 600 1 Crack initiation

500 HIPed at 1400°C

400

gj 300 HIPed at 1500t

-HIPed at 1700°C 200 *

100 -.

HIPed at 1800°C

50 100 150 200 250 Displacement (micrometer)

Fig. 9 Load deflection curves in small punch test for the bonded specimens of tungsten and tantalum. 10 RM4 R. Watanabe et al.

15:h International Plansee Seminar, Eds. G. Kneringer, P. Ro'dhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Bonded interface

7 a

Bonded interface

Fig. 10 Crack initiation and propagation in HEP-bonded W-Ta (The load was applied at the center of the bonded interface.)

3.3 Fabrication of practical target Carburization and oxidation of the tantalum layer during HIPing was found to result in the crack formation during cooling. Carbon source is graphite element that is conventionally used as HIP heater, and oxygen source is argon gas as a pressure medium. To avoid these difficulties the canned block was carefully enveloped with zirconium foil as a getter of carbon and oxygen. The overall view of the tantalum-cladded tungsten target is shown in Fig. 11. This target has been used in the

Fig.ll Overall view of the tantalum-cladded tungsten target. (L= 76.89mm, W= 56.16mm, H= 29.18mm) R. Watanabe et al. RM 4 V\_

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current proton-accelerator-driven pulsed neutron source of the High Energy Accelerator Research Organization, Tsukuba, Japan, in place of Ta target, and has proved to give a neutron yield of 20% higher than that of the Ta target(13).

4. CONCLUSION

Stainless-steel-bonded tungsten heavy alloy still needs study of the mechanical and corrosion properties. But a multiple crack propagation in this alloy would alleviate a catastrophic fracture of the radiation-damaged target. Tantalum-cladded tungsten target for high-density neutron source is successfully fabricated by HIPing at 1500C and 180 MPa with special care of protecting from carburization and oxidation.

ACKNOWLEDGEMENT

The authors would like to express their sincere thanks to Ms. Ritsuko Yoshimura, Graduate School of Engineering, Tohoku University, for her preliminary work on the process exploration for W-stainless steel heavy alloy.

REFERENCES (1) S. Nagamiya, J.Nucl.Sci.Technol., Suppl.l, (2000), pp.40-48. (2) R. Hino et al., KEK Proceedings of Workshop on the Materials for Spallation Neutron Source (December 9-10, 1999, KEK, Tsukuba, Japan), pp. 19-32. (3) IPNS Upgrade: A feasibility Study, ANL-95/13(April 1995). (4) K. Kitazaki and Y. Kiyanagi: Proc. of International Workshop on JHF Science, KEK, Tsukuba, Japan, March 3-7, 1998, Vol.3, pp.68. (5) T. Kawamura: Sumitomo Technical Report, 3(146) (1995), pp.78-81. (6) P.D.Ferguson, Private communication (ACC-APP98). (7) T.A.Broome, Proc.Workshop, JHFScience, KEK Proc, 98-5, III, p.III-82,ANL-95/13,(1999), pp.IV-2-58. (8) N.Takenaka, 2°" Workshop, Spallation Materials Technology, 21-22, March (2001),Tsukuba, Japan. (9)M. Saito, H. Takahashi, H. D. Jeong, A. Kawasaki and R. Watanabe: JSME, 57(1991), pp.522-526. (10) R. Watanabe, A. Kawasaki, M. Tanaka and J.-F. Li, Int. J. of Refractory Metals & Hard Materials, 12 (1993-1994), pp.187-193. (11) M.Kawai, K.Kikuchi, H.Kurishita, J.-F.Li and M.Furusaka, To be published in J. Nucl. Mater.,(2001). (12) T.A.Broome, Proc.2nd Intl.Workshop on Spallation Materials Techonology, Ancona, Italy, Sept. 19-22, (1997), jul-3450, pp.55. (13) M.Kawai et al., Proc. 15th Meeting, Intl. Collaboration on Advanced Neutron Sources, ICANS-XV, 6-9 Nov., (2000), Tsukuba, Vol.11,(2001), pp.747. AT0200468 12 RM7 Y.L Chen et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

CREEP RESISTANCE OF Fe-BASED ODS TUBES WITH HOOP GRAIN STRUCTURE

Y. L. Chen, A. R. Jones, Materials Science and Engineering, Department of Engineering The University of Liverpool, Liverpool, L69 3GH, England

U. Miller Plansee GmbH / Lechbruck, D-86983 Lechbruck, Germany

Abstract

To overcome the shortcoming in creep performance of conventional Fe- based oxide dispersion strengthened (ODS) alloy tubing and release the true potential for the use of these materials in high temperature heat exchanger applications, PM2000 tubes with coarse hoop grain structures have been produced by flow forming techniques. In this paper the creep resistance of these tubes at 1100°C is reported. It was found that the hoop-grained tubes had improved creep properties compared to conventionally extruded tubes. In fact, creep performance approached that of extruded bar materials. Post- creep metallographic tests suggested that tube material failed predominantly by intergranular creep cavity link-up, a common mechanism for this material. Elongated hoop grains together with cleaner grain boundaries are believed to be principal factors contributing to improved creep performance. Further practical issues together with the way in which these tubes can be exploited for a new generation of high temperature heat exchanger are also discussed in this paper.

Keywords

Fe-based ODS alloys, Creep resistance, Powder metallurgy, Mechanical alloying, Fractography Y.L Chen et al. RMJ ]3_

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1. Introduction

Oxide dispersion strengthened (ODS) alloys produced by the powder metallurgy technique mechanical alloying can exhibit an excellent combination of high temperature creep strength and oxidation resistance that cannot be obtained in other alloys [1]. Their unique properties make ODS alloys an important class of materials for potential engineering applications at high temperatures, and often in aggressive gaseous environments. Coarse grained Fe-based ODS alloys such as PM2000 are attractive for potential application in advanced, indirect combined cycle gas turbine systems currently been investigated to enable overall energy conversion efficiencies of 45% and above in biomass systems. To reach such high efficiencies, thermodynamics dictate (Brayton Cycle) that it will be necessary to develop a heat exchanger capable of gas operating temperatures and pressures of around 1100°C and 15-30 bar, respectively, for efficient entry heating of the gas turbine working fluid. This, in turn, demands metal operating temperatures up to 1150°C for heat exchanger tubing. In such pressurised tubes, the maximum principal creep stress is oriented in the hoop direction and conventional Fe-based ODS alloy tubing currently exhibits substantially poorer hoop than axial creep resistance. The reason for this is clear. Conventional Fe-based ODS alloy tubing is processed by unidirectional extrusion followed by heat treatment. This results in tubes with anisotropic, very coarse grained, axially aligned microstructures which exhibit excellent axial creep properties. To overcome this problem and release the true potential for application of Fe-based ODS alloys in high temperature heat exchangers, a European funded BRITE-EuRAM project has been undertaken in an effort to develop PM2000 alloy tubes with a coarse grained, high grain aspect ration (GAR) microstructure aligned in the hoop rather then the axial direction using commercial flow forming techniques. The latter is a chipless machining technique with great precision and improved material economy. PM2000 tubes with coarse hoop grain structure have been successfully formed and the detailed results have been reported elsewhere [2, 3].

In this paper, the hoop direction creep resistance of two of these flow formed tubes has been examined. Results indicate that such tubes can exhibit much improved creep peformance compared with the normal extruded tubes. U RM 7 Y.L Chen et al.

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2. Experimental Methods

PM2000 is a ferritic ODS superalloy (Fe-20Cr-5AI-0.5Ti-0.5Y2O3) produced by mechanical alloying. Two flow formed tubes with overall levels of deformation of 80.80% (Tube 1) and 89.60% (Tube 2), respectively, have been examined in this work. Details of tube production processes can be found elsewhere [2,3].

To prepare hoop creep samples for testing from flow formed tubes, ring samples were cut from as-flow formed tubes. These were then slit axially and flattened at a temperature of 800°C. The flattened strips were annealed at 1380°C for 1h to give a fully recrystallised coarse grain structure. To ensure a flat sample was obtained, a load was applied to the strip during recrystallisation annealing. Hoop orientation creep samples were sectioned from tube after annealing and prepared for testing.

Creep tests were carried out at 1100°C with a constant load of either 40MPa or 50MPa, respectively. Details of samples tested in this work are summarised in Table"!.

Fractography of crept samples was performed using a Hitachi S-2460N SEM operated at 25kV.

To further understand creep behaviour, longitudinal sections of crept samples (i.e. transverse sections of tube) were mounted and mechanically polished to a 0.5|im finish using an alumina suspension and then were examined optically using differential interference contrast (Nomarski) techniques. One important advantage of this technique is that chemical etching is not required, obviating the contrast differences that can complicate simultaneous observation of porosity and grain structures in chemically etched samples.

3. Results and Discussion

Tablei summarises the creep results from the two tubes tested at 1100°C. For comparison, these creep data together with creep data from several Fe- based ODS alloys and product forms tested at 1100°C have been presented in Fig.1. It is clear that the hoop creep strength of the flow formed and recrystallised tubes is enhanced compared with traditionally extruded PM2000 and MA956 tubes. This was particularly true for tube 2, where creep performance approached that of PM2000 bar material. These results indicate Y.L. Chen et al. RM7 15

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

that hoop direction creep performance can be improved by realigning the coarse grain structure from the axial to the hoop direction.

Table 1. Hoop Direction Creep Life of the Flow Formed PM2000 Tubes

Tube Number Sample Stress (MPa) Temperature Creep Life Number (°C) (h) 1 HC11 40 1100 42 1 HC12 50 1100 9.6 2 HC21 50 1100 136 2 HC22 50 1100 282

Creep strength at 110CTC

100

^»—PM2000 sheet -&- PM2000 bar X v $• tube 1 hoop -! ;-PM2000tube X "&• tube 2 hoop X MA956tube

10 100 1000 10000 100000 time [h]

Fig.1 Hoop creep resistance of the flow formed tubes together with other Fe-based ODS materials.

Fig.2 is a SEM image from the fracture surface of creep sample HC21, while Fig.3 shows an optical micrograph from the same sample which reveals coarse grains elongated in the hoop direction. Comparison of Figs. 2 and 3 indicates that an essentially 3-grain microstructure exists through the wall of 16 RM7 Y.L. Chen et al.

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the tube and this is reflected in the fracture surface. Grain boundary sliding striations (Fig. 4) were observed on the intergranular surfaces between the two grains seen in region B at the centre of Fig.2. It was also observed that these grains were elongated in the tensile direction, suggesting necking had taken place. The separation between large grains in the longitudinal-'hoop' plane in Fig.2 also suggested that necking had occurred. Such intergranular failure with plastic deformation is typical of grain boundary separation at high temperatures, with no evidence of the flat, smooth facets that would be expected with low temperature intergranular fracture [4].

There is a ridge at the 'top' of the two grains exhibiting creep plasticity in sample HC 21. This can be seen in the upper right hand corner of Fig.4; local shearing can be seen near the ridge. This suggests that the creep crack initiated from the 'bottom' and propagated to the 'top'. Crack propagation at this late stage in the failure process amounted to local tensile shearing failure under the increasing net section stress.

Besides grain boundary sliding, creep cavitaties (dimples) can be seen at the surface of the grain in the lower right hand corner of Fig.2. Fig.5 is a SEM image of another region of the creep failure surface in sample HC21, showing predominantly creep cavitation and indicating that microvoid coalescence is a further important creep crack growth mechanism in this sample. The original coarse grain structure is also apparent in Fig.5, as is slight separation or tearing along the grain boundaries.

Most creep cavities observed in the tested samples were equiaxed, consistent with the fact that most of these fracture surface are perpendicular to the loading direction. However, elongated shear cavities were observed in some areas where shear stress is possible and Fig.6 show such an example, near the area indicated by L. The slope of the area is not very obvious because of the large depth of focus of SEM.

The fractography of sample HC22 was very similar to that of sample HC21. Y.L. Chen et al. RM7 17

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Fig. 2 A SEM micrograph from the fracture surface of sample HC21.

0.2mm

Fig.3 An optical micrograph from longitudinal section of sample HC21showing the coarse grain structure. 18 RM 7 Y.L. Chen et al.

15:h International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

25|xm

F/g.4 >A higher magnification SEM image from the area indicated by letter B in Fig.2 showing sliding striation.

Fig.7 is a montage of the creep failure surface observed in sample HC11. Creep cavitation was the principal mechanism evident on the creep fracture, with grain boundary sliding observed in two areas in the grains observed on the left hand side of the figure. It is particularly interesting to note that the grains in Fig.7 seem to be 'stacked' in groups through the tube wall, with virtually common boundary edges in the radial- 'hoop' plane, bounded by local tearing. This periodicity in the fracture surface is consistent with the periodicity in the deformation pattern induced by the regular passage of the flow forming rollers over the outer tube surface prior to recrystallisation.

In fact, previous work has already shown that the longitudinal section of this tube comprises a recrystallised grain structure that reflects the periodicity of the deformation process, appearing as 'stacks' of grains along the tube axis. The common grain boundary in the radial-hoop plane exhibits shearing during creep crack growth giving rise to the peculiar feature observed in Fig.7. Y.L Chen et al. RM7 19

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Fig.5 Another SEM fractography from sample HC21showing creep cavitations.

Fig.8 is an optical micrograph showing a longitudinal section of sample HC21 through the region of the creep failure. It is clear that in areas of the fracture surface indicated by the arrows, the fracture has followed an intergranular path where grain boundary sliding appeared to be the predominant creep mechanism.

One concern for this work is whether the present method can really show the hoop creep resistance of the flow formed tubes because Figures 3 and 8 also clearly show that the grains in the flattened creep samples were both coarse and elongated in the hoop direction, in other words, similar to the microstructure observed in flow formed and recrystallised tube. This obviates the concern that the extra deformation introduced by flattening the tube to produce a creep sample may have substantially changed the original grain structure.

Fig.9 is an optical micrograph from the longitudinal section of the fracture in sample HC12, showing creep cavities at the fracture face. This confirmed that cavitation was an important creep crack growth mechanism in the current tests on the ODS tubes. 20 RM7 Y.L. Chen et al.

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Fig.6 Another SEM image from HC21 showing elongated cavitations as that in the area indicated by L,

An optical image of a region close to the fracture face in sample HC11 (tube 1) is shown in Fig. 10. Comparison with Figs.3 or 8 show that the grain size in tube 1 is smaller than that in tube 2, as reported previously [3]. More particularly, creep cavities can be seen in Fig. 10, mainly concentrated along a grain boundary perpendicular to the tensile direction. This suggests that creep cavitation may be principally intergranular in nature.

The above results indicate that there are two main micro-mechanisms of creep crack propagation in the PM2000 tubes: grain boundary sliding and intergranular creep cavity link-up. This confirms the important role of grain boundaries in determining creep performance at high temperature in ODS alloys. On this basis the improved creep performance found in tubes with a Y.L. Chen et a!. RM7 21

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Fig.7 A SEM montage from sample HC11 showing the typical morphology of this sample. 22 RM7 Y.L Chen et al.

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Fig.8 Optical micrograph from longitudinal section of sample HC21showing that grain boundary sliding might have happened during creep.

Fig.9 Optical micrograph from longitudinal section of sample HCHnear the fracture tip showing creep cavitations link-up at grain boundary. Y.L. Chen et al. RM 7 23

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t»J&}T

0.2mm

Fig.10 Optical micrograph from longitudinal section of the fracture tip of sample HC12 showing that the sample might fail by creep cavitations link-up. coarser hoop grain structure than found in conventionally processed material is easy to understand.

First of all, the flow formed tubes have a coarse grain structure with a high GAR in the hoop direction e.g. as shown in Fig.3. In fact, in tube 2, hoop oriented grains were found to have dimensions which could exceed half the circumference of the tube. This contrasts sharply with traditionally extruded ODS tubes where grains are elongated in the axial rather than the hoop direction. With high aspect ratio grains aligned in the hoop direction, which is the orientation of maximum principal stress for internally pressurised tubes, higher creep lives are to be expected.

Moreover, flow formed and recrystallised tubes may contain grain boundaries less decorated with oxide particles than in conventional tubes, imparting inherently improved resistance to intergranular creep cavity nucleation. In the present case the YAG dispersoid or (particularly) entrained alumina will be favoured sites for creep cavity nucleation [5], thus 'cleaner' grain boundaries will reduce the incidence of microvoid initiation, extending creep life. In traditionally extruded ODS tubes, the oxide particles tend to align in stringers in the axial direction, causing high local Zener pinning and trapping 24 RM 7 Y.L. Chen et al.

15th International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

recrystallised grain boundaries. This particular particle alignment is in fact an important reason for the anisotropic grain structures obtained in extruded ODS tubes [6]. The high density of oxide particles at grain boundaries has been confirmed experimentally [7]. In contrast, extensive TEM work on the current, flow formed tubes has confirmed a more homogeneous oxide dispersion with no obvious particle alignment [3]. Therefore, a cleaner grain boundary will be expected in the flow formed tubes.

The results from Table 1 also suggest that tube 2 has better creep resistance than tube 1. This is consistent with the smaller grain size found in tube 1 compared with tube 2, as shown by Figs. 3 and 10. The reason for this has been discussed in a recent work [3]. The higher level of deformation used to produce tube 2 may also result in a more homogeneous particle distribution and, thus, cleaner recrystallised grain boundaries.

As mentioned in previous section, the lack of hoop direction creep resistance is an important factor preventing Fe-based ODS tubing from being applied in next generation high temperature heat exchanger. This work has shown that by re-aligning the high GAR direction in the ODS tubes from the axial direction to the hoop direction, tube hoop creep resistance can indeed be improved. These tubes have a significantly enhanced potential for application as high temperature internally pressured tubing in heat exchangers in future commercial power generation systems.

4. Summary

The creep resistance of flow formed PM2000 tubes has been explored through 1100°C, hoop oriented tests on flattened samples. Metallographic assessment was subsequently performed on these crept samples. The results can be summarised as follows:

(i) The hoop creep resistance of Fe-based ODS tubes can be improved by flow forming techniques. Tubes produced by this method can have creep endurance close to that of axially tested bar materials.

(ii) Operating creep mechanisms in the ODS tubes included grain boundary sliding and intergranular creep cavity link-up. Improved creep resistance in flow formed tube has been explained in terms of these creep mechanisms and the unique microstructure of the flow formed tubes. Y.L Chen et al.RM 7

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5. Acknowledgements

Dr. Chen is very grateful for constructive discussions with Prof. D. Hull, The University of Liverpool. Financial support from the European Commission is acknowledged via the BRITE-EuRAM initiative. It is also a pleasure to acknowledge our project partners: Cambridge University, Metall-Spezialrohr GmbH (MSR), MBEL, Ris0 National Laboratory and Sydkraft AB.

REFERENCES

1. E. F. Bradley, Superalloys: A Technical Guide, ASM INTERNATIONAL, Metals Park, OH, 1988 2. Y. L. Chen and A. R. Jones, Recrystalllization-Fundamental Aspects and relations to Deformation Microstructure, Eds: N. Hansen et al., Roskilde, Denmark (2000), pp.291-296 3. Y. L. Chen, A. R. Jones, R. C. Pond and U. Miller, Mater. Metall. Trans. A, (Submitted) 4. Metal Handbook, Materials Park, OH, ASM International, 10th ed, Vol.12, pp.64-78 5. R. E. Smallman, Modern Physical Metallurgy (Third edition), The Butterworth Group, pp.481-498 6. H.K.D.H.Bhadeshia, Mater. Sei. Engr. 1997, Vol. A223, pp.64-77 7. J. Ritherdon, unpublished work AT0200469

26 RM 16 B. Schedler et al.

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The Manufacture of Carbon Armoured Plasma-Facing Components for Fusion Devices

B.Schedler, Th.Huber, A.Zabernig, F.Rainer, K.H.Scheiber, D.Schedle Plansee AG, A-6600 Reutte, Austria

Summary:

Within the last decade Plansee has been active in the development and manufacture of different Plasma-Facing-Components for nuclear fusion experiments consisting in a tungsten or CFC-armour joined onto metallic substrates like TZM, stainless steel or copper-alloys. The manufacture of these components requires unique joining technologies in order to obtain reliable thermo mechanical stable joints able to withstand highest heat fluxes without any deterioration of the joint. In an overview the different techniques will be presented by some examples of components already manufactured and successfully tested under high heat flux conditions. Furthermore an overview will be given on the manufacture of different high heat flux components for TORE SUPRA, Wendelstein 7-X and ITER.

Keywords: Joining, C/C, CuCrZr, TZM, Active Metal Casting, Laser treatment, Brazing, ITER, Tore Supra, Wendelstein 7-X, RFX, Divertor, High Heat Flux Components, NDT, radiographic examination, ultrasonic examination, thermography

Introduction: In the last decade an extensive technology development program has been carried out by Plansee on the manufacture of carbon armoured high heat flux components for nuclear fusion experimental reactors. Beside the development and manufacture of prototypes for the german stellerator experiment Wendelstein 7-X and ITER1, which is a triplate, collaborative project carried out by the European Union, Japan and the Russian Federation, Plansee has also supplied the thermally highest loaded components for the French Tokamak experiment Tore Supra, namely a part B. Schedler et al. RM 16 27

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of the Inner First Wall, Antennae Limiter and Toroidal Pumped Limiter. All these components consist in at least one portion of C/C composite that has to be joined onto metals like molybdenum, stainless steel or copper alloys. The main obstacle to overcome is the realisation of a reliable joint dealing with heat fluxes up to 30MW/m2 and the associated thermo mechanical loads. The approach of Plansee comprising elements like laser structuring2 of the joint area or Active Metal Casting3 has proven its reliability under to above described thermal load and is regarded as the reference solution for ITER [1].

Design of Relevant Components In the frame of the ITER EDA4 several plasma facing component designs have been proposed [2].

•I- -" - i Each of them has specific flat tile _._"_L advantages and extensive testing Un- has shown that they could sustain saddle tile 1 - _ . the requirements. The three main options depend on the bonding of

mono- / macroblock the plasma facing material to the heat sink. A sketch of each is shown fig. 1: basic designs of high heat flux in fig. 1. The flat tile design is the components for divertors most common and easiest to manufacture. The joining interface is flat and allows an easy fitting with rectangular heat sinks having different types of cooling concepts. Straight tubes with or without turbulance promoters (fins, swirl tapes etc.) are the simplest cooling structures with reasonable performance. Hypervapotrons are more sophisticated options revealing improved performance. In the flat tile geometry a stress singularity occurs at the free edge of the interface. The strength of this singularity depends on the elastic properties and on the coefficient of thermal expansion of the dissimilar materials, on the shape of the interface and on the temperature field. The saddle block consists of a series of separate heat sink square blocks connected by a concentric tube and protected by a saddle shaped armour. It is an improvement of the proceeding one with the bonding interface closer to the coolant fluid.

2 patent: EP 0 476 772 Bl 3 patent: EP 0663 670 Bl 4 EDA...Engineering Design Activities 28 RM 16 B. Schedler et al.

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The monoblock, consisting of a series of individual armour square blocks connected by a concentric tube, removes the stress singularity by eliminating the flat interface. The monoblock or macroblock design takes the advantage of the carbon composites thermomechanical properties. The bonding interface is the circular cooling tube which gives the lowest possible temperature for the joint. The high thermal conductivity, the low elastic modulus and the low thermal expansion coefficient reduce drastically the thermomechanical stresses of this concept. This design allows for large massive blocks, having structural properties, which reduce local peaking heat fluxes by spreading out the heat flow through the carbon monoblock. By now the components based on the monoblock design outperformed by far the flat tile and saddle designs.

Manufacturing Route according to the Plansee Approach

Reliable metal-C/C composites require beside a perfect wetting of the braze alloy also the mastering of stresses generated due to the mismatch of properties of the different materials to be joined. In the manufacturing approach applied by Plansee both issues are considered, combining a LASER treatment of the C/C joint surface with the so called active metal casting process. The reliability and outstanding performance of components manufactured with these two techniques have been demonstrated within the devolpment activities of carbon armoured plasma facing components for nuclear fusion experiments [1,3,8,9,10,11, 12].

Laser Structuring

An important consideration in making dissimilar materials joints relates to the thermal mismatch stresses and strains. In any joining of portions of two dissimilar materials the mismatch in CTE is extremely important. From the mismatch of their thermal expansion thermal stresses result. The magnitude of these mismatch stresses and strains determines the failure probability of the joint. These mismatch stresses and strains must be carefully controlled. According to Hagy and Ritland [5] the CTE mismatch differentials of within 100ppm are considered as allowable for reliable dissimilar material joints. An important problem with common joining processes is the understanding and control over a period of cooling time the mismatches in CTE and thermal strain and stress gradients in the joint region. At Plansee these considerations have led to the development of a microengineered joint interface. With this process a LASER drills conical B. Schedler et al. RM 16 29

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holes into the joining interface C/C part. A typical metallographic cut of a brazed sample with such a structured interface shall be demonstrated in fig. 2. Depending on the metallic substrate such laser treated Graphite or C/C parts can be joined directely by usual brazing processes onto metallic substrates like molybdenum, stainless steel or copper.

The outstanding positive influence of such a treatment can be easily demonstrated by thermal shock experiments. Under standardized quenching conditions there is a one-to-one correspondence between the joint strength and the allowable initial quenching temperature [6]. This temperature can be regarded as a direct measure of the joint strength for a specific joint configuration.

In order to test the integrity of the joint such quench tests have been performed in the frame of the qualification procedure for the inner first wall of Tore Supra [3]. Metal-C/C composites samples have been prepared representing the joint being used for the inner first wall of Tore Supra. These samples consisted of a C/C joined onto stainless steel via an intermediate 2mm thick OFHC-copper layer joined with a Ti-Cu-Ag braze alloy. One part of the samples has been joined without laser treatment whereas the other part of the samples revealed a LASER treated surface. Samples with microengineered surfaces survived 20 thermal shock cycles from 600°C to cold water without destruction whereas samples without laser treatment have been damaged after a few cycles. During later exposure to the plasma the components manufactured with this qualified technology could sustain heat fluxes up to 1 MW/m2 without any deterioration [12] during a several years lasting operation period. The positive influence of the structuring might be explained by different effects. By such a structure a graded-material is created in the joint region smoothening the transition of properties. Finite Element Calculations [4] further revealed that the stress in the structured CFC portion is reduced by a transfer of the load from the CFC into the metallic portion cast into the cones. Furthermore also the propagation of cracks in the brittle carbide interface fig. 2: laser structure brazing interface is impeded. (C/C joined onto TZM with TiCuAg 30 RM 16 B. Schedler et al.

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Active Metal Casting

In general the joining of metals to graphite or ceramic materials is not only a problem due to the residual stresses caused by the mismatch of properties, but also due to metallurgical reasons. Usually active Brazing is the prime choice in order to obtain metal-C/C joints. Carbide forming elements out of the group IV and V are constituents of such active braze alloys. These elements react with the C/C joint surface and make it wettable for the liquid metals of the braze alloy. However, due to the materials involved such brazed components often suffer from difficult to control processes resulting in braze flaws. The probability of braze flaws increases with increasing joint area, leading to the rejection of the whole component. In order to overcome this critical issue Plansee has developed the Active Metal Casting process. With this process a non-carbide forming metal like copper is cast onto the C/C composite in the presence of Titanium. During this process the Titanium reacts with the C/C composite by the formation of a TiC interlayer, which then can be wet be the liquid copper. This process can be applied to a wide range of geometries such as straight and bent flat tiles and monoblocks (fig. 3). Usually this process in combined with the above described LASER structured interface in order to deal with the stress occurring during the cooling from the casting temperature (fig. 4).

fig. 3 active metal cast C/C flat tile and fig. 4: active metal cast interface with monoblock laser structured C/C With such laser treated, active metal cast C/C flat tiles or monoblocks a copper layer is generated on the graphite substrate, which then acts as metallic interface for further less critical and more reliable metal/metal joints by applying techniques like brazing, Electron Beam Welding, Hot Isostatic Pressing or Diffusion Bonding for the joining of the C/C armour onto the heat sink. This possibility opened by the Active Metal Casting Process is interesting for Heat Flux Components consisting of materials sensitive to heat treatments like age-hardening Cu-alloys. B. Schedler et al. RM 16 31

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Compared to a direct brazing operation of the CFC onto a metallic heat sink the intermediate active metal casting step allows a 100% non destructive inspection of the critical metal-C/C joint prior to final joining onto the heat sink. At Plansee this non-destructive inspection is usually done by radiography, judging the wetting and penetration of the cast metal into the C/C composite. Such a pre-selection of the critical metal to C/C joints is a big advantage with regard to large components, where even one bad metal-C/C joint results in the rejection of the whole component.

Manufacture of Antennae and Toroidal Pumped Limiter for TORE Supra

fig. 5: Antennae Limiters fig. 6: Toroidal Pumped Limiter

For the French nuclear fusion experiment TORE SUPRA Plansee manufactured beside a part of the inner first wall also 250 Antennae Protection Limiters and the 600 Toroidal Pump Limiter (see fig. 5 and 6). Both types of components consist in laser structured, active metal cast C/C flat tiles joined onto a CuCrZr heat sink by means of deep penetration electron beam welding [7]. As it can be seen on the images it is possible to apply the described techniques not only to flat but also to bent C/C tiles. The quality of the active metal cast interface is judged by means of radiographic inspection whereas the electron beam weld is judged by ultrasonic inspection. Flaws with maximum 03mm are accepted for both cases. After successfully passing these non destructive examination the cooling channels are machined into the CuCrZr armour. Subsequently the cooling system is closed by plugs joined by electron beam welding. For the transition of the copper cooling system to the stainless steel cooling system of the TORE SUPRA machine an electron beam welded dog-bone joint geometry comprising an intermediate Ni-layer has been applied. For final acceptance the components undergo a series of non-destructive examinations like a hot-helium leak test operating at 250°C/40 bar internal He-pressure, followed by a transient thermographic 32 RM 16 B. Schedler et al.

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inspection carried out by flooding the cooling system with hot steam (250°C/40 bar) and measuring the thermal response of the C/C surface. Fig. 7 shall summarize all the destructive and non-destructive inspection steps necessary within the manufacture of these critical components.

Pfematorial Inspection SEPNIt Density Measurement OFHC Chemical Analysis CuCrZr Hardness-, Electrical Cond, Measurement. Chem. Analysis Nickel Chemical Analysis AISI 316L Clicrriicat Anaiysis.'Slructure Evaluation

1 /OFH<" Flat Tiles e ola L sting X RT In reel on

Cu Z r ZU CFC+ÄNIC-aJ'? CÜCrZr HeaVsSriiT [____ _E 8 m c Electron Beam Welding i Holiu nte kT h e Ultrasonic Inspection

CuCrZr / CuCfZr Rear End Plugs Electron Bcarn_Welding Helium Leati Tighlnessjnspection Toroidal Pumped Limitor Eleme | Hol_Hclium Leak Test | Inlrarod Thermography Test I CuCrZr Hardness MeasuremG 'CuCrZr Electrical Conduc'ivit/ Meas: fig. 7: Manufacture and Test Sequence of Toroidal Pumped Limiters

Manufacture of prototypical High Heat Flux Components for ITER with C/C monoblock armour joined onto CuCrZr-tubes

For areas like the divertor for ITER the C/C armoured components are envisaged as the reference solution [8]. In the lower area of these components heat loads up to 20MW/m2 are expected during so-called "slow transients" As described above monoblocks proved to be the most robust solution to cope with these extremely demanding operating conditions. fig. 8 ITER Vertical Target Medium Scale Mock-Up B. Schedler et al. RM 16 33_

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The first generation of C/C monoblock components was developed in the mid 90th. The reference manufacturing technology has been active metal casting of a 0,5 - 1,0mm thick OFHC copper interlayer onto the laser structured surface of a drilled bore of a CFC tile [9]. The AMC C/C monoblock is then joined onto the copper alloy tube (PH-Cu alloy CuCrZr or DS-Cu alloy GlidCop®) by an eutectic Ti-Cu brazing at 880°C. Following this manufacturing route a medium scale vertical target prototype has been manufactured and tested (fig. 8). This component contained all the main features of the ITER divertor design and had an overall length of about 600mm. The high heat flux part had a lower straight section made of C/C monoblocks and an upper bent region with a tungsten macrobrush flat tile armour. The heat sink has been manufactured from the DS-Cu alloy GlidCop® by HIP of a bent GlidCop® tube into two GlidCop® half shells. Onto this heat sink the tungsten macrobrush flat tiles comprising a cast OFHC layer have been joined by electron beam welding. The AMC C/C monoblocks have then been joined onto the GlidCop®-tube by an eutectic Ti- Cu brazing at 880°C. A twisted tape with a twist ratio of 2 was inserted in the monoblock region to enhance the critical heat flux limit. This component has then been tested at FE200 electron beam facility at FRAMATOME in Le Creusot, France on both the tungsten and the C/C armour at 10MW/m2 without any failure. The C/C monoblocks has further been tested up to 30 MW/m2 without any failure.

The results of the R&D activities carried out on copper alloys revealed that DS-Cu has a very poor fracture toughness which together with low creep resistivity and weldability has triggered the choice of the reference heat sink material to the precipitation hardened material CuCrZr. Furthermore components with a bent C/C monoblock section are required. These considerations led to the need for developing a new manufacturing technology because the brazing technique at 880°C was not compatible with the use of CuCrZr as heat sink material. An extensive R&D has been carried out by Plansee to develop an optimised manufacturing process for active metal cast CFC monoblocks fully compatible with the use of CuCrZr alloy. The selected manufacturing process was Hot Isostatic Pressing of OFHC copper active metal cast C/C monoblocks onto solution annealed water quenched CuCrZr tubes, which are age hardened during the HIP cycle. With this process it has been possible for the first time to manufacture bent C/C monoblock components as it is shown in fig. 9 and 10. 34 RM 16 B. Schedler et al.

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fig.9: Curved C/C monoblock mock-up fig. 10:ITER Divertor Full Scale Prototype

The components of fig. 9 represent small scale mock-ups manufactured by HIP in the frame of the qualification of this technology. These mock-ups have been successfully tested under high heat flux with electron beam up to 18MW/m2. Based on this technology a full scale ITER Divertor high heat flux component (length 1200mm) shown in fig. 10 has been manufactured. This component consists in a section of tungsten and a CFC monoblocks, respectively. Both monoblock types comprise a OFHC copper interlayer in the bore hole, which has been used for the joining onto the CuCrZr tube by Hot Isostatic Pressing.

Manufacture of prototypical High Heat Flux Components for Wendelstein 7-X

The open divertor design proposed for the stellerator experiment Wendelstein 7-X, currently under construction at Greifswald (D), consists of 10 independend divertor units with an overall target plate area of 22m2, formed of approximately 1500 actively cooled elements [10,11]. The target plates are designed for steady state operation with a local heat load of up to 10MW/m2. Because of the geometrical boundary conditions in the W7-X vessel, the target plates have to be very thin, which is why a flat tile design with C/C flat tiles brazed onto a TZM water cooled heat sink was selected. The thickness of the C/C tiles was restricted with 6mm in order to keep a surface temperature limit of maximum 1200°C for the highest power load. In order to increase the thermal transfer coefficient between the heat sink and the cooling water a fin design with 5 fins per cooling channel has been chosen. In fig. 11a target element cross section is shown. It consists of the laser structured C/C flat tiles, the TZM fin plate heat sink and the TZM back plate with for parallel cooling channels. B. Schedler et al. RM 16 35

15"1 International Plansee Seminar, Eds. G. Kneringer, P. Rödhamrner and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

fig. 11: Cross section of W7-X prototypical fig.12: Prototypical W7-X divertor target divertor target plates plates C/C joined onto TZM

The manufacturing route of the target elements was determined by the requirement of a complete 100% non-destructive inspection of all braze interfaces. This led to a specific feature of the manufacturing route, i.e. the split of the brazing operation in two cycles. The first between the laser structured C/C tiles and the TZM-fin plates (TiCuSil / 850°C), the second between the TZM fin plate and the TZM back plate (AgCuPd / 810°C). For such a sequence high precesion is demanded in order to compensate for the warpage induced by the different thermal expansion coefficients of the C/C composite and the TZM. Such a way manufactured elements have then been tested using the electron beam device JUDITH of the KFA Jiilich with a stationary power load up to 12MW/m2 without any detachment of the C/C tiles. 36 RM 16 B. Schedler et al.

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References:

[I] International Thermonuclear Experimental Reactor, Technical Basis for the ITER Final Design Report, Cost Review and Safety Analysis, IAEA, Vienna 1998, Chapter Il-Section 4.1 - Page 34

[2] Ph.Chappuis et al., Possible Divertor Solutions for a Fusion Reactor. Part 2. Technical Aspects of a Possible Divertor, Fusion Engineering and Design, Volume 36/1 (1997), page 109-117

[3] M.Lipa et al., Development and fabrication of improved CFC-brazed components for the inner first wall of Tore Supra, Proceedings of the 19th Symposium on Fusion Technolgy, Lisbon, Portugal, 16-20 September 1996

[4] Plansee Internal Report

[5] Hagy and Ritland, Viscosity Flow in Glass-to-Metal Seals, J. Amer.Ceram.Soc., Vol. 40, pp 58-62, 1957

[6] Chou H.Li, Ceramic Composite, United States Patent, 5.874,175,

[7] Kneringer et al., EP 0 741 116 B1, Process for Manufacturing of components with high thermal load capability

[8] International Thermonuclear Experimental Reactor, Technical Basis for the ITER final design report, cost review and safety analysis (FDR), chapter ll-section 4.1-page 33, IAEA, Vienna 1998

[9] G.Vieider et al., European development of the ITER Divertor target, Fusion Eng. Des. 46, 221 (1999)

[10] Renner H. et al., J.Nucl.Mater., 241-243, 946-949, (1997)

[II] Greuner H. et al., Design, Manufacture and Testing of the W7-X Target Element Prototypes, Proceedings of the 19th Symposium on Fusion Technology, Fusion Technology, Vol. 1, 463-466 (1996)

[12] Schlosser J. et al., Experience Feedback from High Heat Flux Component Manufacturing for Tore Supra, Proceedings of the 21th SOFT conference, 11.-15. September 2000 Madrid, Spain, in press V.V. Dabhade et al. RM 57 37

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PROCESSING OF NANO SIZED TITANIUM NITRIDE

REINFORCED TITANIUM COMPOSITES

V.V. Dabhade, M.M. Godkhindi*, T.R. Rama Mohan, P. Ramakrishnan and E. Y. Gutmanas.**

Department of Metallurgical Engineering and Materials Science Indian Institute of Technology, Bombay, Mumbai, India. *Department of Metallurgical and Materials Engineering Indian Institute of Technology, Kharagpur, India.

** Israel Institute of Technology, Haifa, Israel.

SUMMARY:

Attempts have been made to study the effect of nano sized TiN reinforcement loading, sintering temperature and atmosphere on the densification behavior of titanium - titanium nitride nanocomposites. Titanium with average particle size of 2 microns and titanium nitride with a particle size of the order of 50 nanometers were used in the present investigation. The compacts showed good sinterabiüty both in argon and nitrogen atmospheres. The initial stage of sintering of pure Ti as well as Ti-TiN nanocomposites has been studied by using the constant rate of heating technique. The frequency factor (Do), activation energy for sintering (Q) and diffusion coefficient (D) have been evaluated and discussed for the Ti and Ti-TiN nanocomposites sintered in argon and nitrogen atmospheres. Sintering of Ti and Ti-TiN nanaocomposites in nitrogen atmosphere is affected by two competing mechanisms such as endothermic activation energy required for mass transport mechanism and an exothermic energy due to insitu nitridation reaction. Microstructural studies have been carried out to supplement the densification data.

KEYWORDS:

Titanium, Titanium Nitride, Nanocomposite, Sintering, Frequency factor, Activation energy, Diffusion coefficient. 38 RM 57 V.V. Dabhade et al.

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1. INTRODUCTION:

Titanium is the design choice for many aerospace, automotive, medical, marine, chemical processing and biomedical applications because it offers low density, outstanding corrosion resistance and good mechanical properties at room and moderately elevated temperatures [1-3]. Despite the fact that titanium is the best choice for many systems from the stand point of performance durability, less titanium is used in the final production because of the high cost of the components which includes the initial cost of metal, the cost of processing and forging and the cost of machining. As a result, net shape or near-net shape technologies are one of the major avenues for reducing the cost of titanium components. These processes include isothermal forging, casting, superplastic forming and powder metallurgy [3]. Titanium powder metallurgy offers the potential for true net shape capability, an efficient material utilization combined with mechanical properties that are equal to or exceed cast and wrought products [4,3]. It also allows a four to five fold increase in metal utilization and a one to two fold decrease in labor consumption [5].

Titanium matrix composites (TMC's) are processed by the incorporation of high modulus and high strength reinforcements into titanium. Titanium itself possesses a high specific strength at room and moderately elevated temperature and due to the incorporation of these reinforcements there is a significant improvement in its specific modulus, specific strength and creep resistance [6]. Titanium nitride (TiN) is an important reinforcement candidate material for TMC's. Titanium nitride possesses unique properties like high hardness, stability at high temperatures, and electrical and optical properties that are highly desirable for a variety of applications. Titanium nitride is used for cutting tools, tool coatings, solar-control films, and microelectronics applications. It is harder than alumina, thermally stable to about 3300K, chemically stable with respect to most etching solutions and provides an excellent diffusion barrier against metals [7].

Particulate reinforced TMC's have been mainly processed by solidification/casting process and powder metallurgy. The high melting point of titanium (1933 K) and its high reactivity in the liquid state are the major limitations in processing particulate TMC's by the solidification/casting process [6]. Powder metallurgy provides an excellent route for the processing V.V. Dabhade et al. RM 57 39_

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of particulate TMC's provided homogeneous distribution of the reinforcing phase takes place [8].

A nano-composite is a composite consisting of two or more phases where atleast one of the phases is in the nano range (< 100 nm). Because of the extremely small dimensions of the particles a large volume fraction of the atoms are located at the grain boundaries, and this confers special attributes to these materials. The properties of titanium nano-composites are very often superior to those of conventional coarse-grained TMC's. Titanium nano- composites exhibit higher strength/hardness, enhanced diffusivity, improved ductility/toughness when compared with conventional coarse-grained TMC's [9], Several novel processing methods have been developed to synthesis nanosized powders from which nano-composites are processed. These processes can be classified broadly as mechanical synthesis [2,10], thermochemical synthesis [11] and chemical synthesis [12]. Titanium nano- composites can be synthesized from these powders by consolidation techniques like pressureless sintering, sinter-forging, hot pressing and hot isostatic pressing [13].

In the present investigation, attempts have been made to study the effect of nano sized TiN reinforcement loading, sintering temperature and atmosphere on the densification behavior of Ti-TiN nano-composites, processed by the powder metallurgy route.

2. EXPERIMENTAL:

Elemental titanium powder (99.8% purity) of average particle size of 2 ^m and nano sized titanium nitride powder of average particle size of 50 nm were used in the present investigation. The particle size was measured using TEM. Three sets of powders were prepared. First set of elemental titanium, second set of powder mixture of elemental titanium with 8-wt % of titanium nitride and third set of powder mixture of elemental titanium with 15-wt % of titanium nitride. The powder mixtures were prepared by manual mixing in a mortar and pestle followed by ultrasonic mixing, while the same treatment was also given to the set of elemental titanium powder. The powder sets were then compacted uniaxially in a rigid die of 12 mm diameter at a pressure of 435 MPa. The compacts were then sintered in a rapid heating tubular furnace at five different temperatures in the range of 900-1100°C for 30 min. Sintering experiments were repeated in two different sintering atmospheres of argon and nitrogen. 40 RM57 V.V. Dabhade et al.

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3. RESULTS AND DISCUSSION:

3.1 Characterization of Powders Figures 1(a) and 1(b) show the TEM photographs of pure titanium and titanium nitride powders respectively. The reported values of the particle size of titanium and titanium nitride powders measured by TEM are the average of at least 25 readings. The powders were also subjected to X-ray diffraction (XRD) analysis and they revealed that titanium was in a form and both titanium and titanium nitride were single-phase structures (no peaks corresponding to any other TixN phase were observed). TEM photographs of the starting powders clearly reveal that Ti powder is micron-sized and TiN powder is of nanometric size though both the powders do show a tendency to form clusters and agglomerates. 4 (a)

Figure 1: TEM photographs of (a) pure Ti powder and (b) TiN powder.

3.2 Sintering of elemental titanium compacts Figure 2(a) indicates the variation of sintered density as a function of sintering temperature for pure titanium samples sintered in argon and nitrogen atmospheres. It can be seen that the sintered density increases in both the atmospheres. Within the narrow temperature range studied, the sintered densities seem to increase linearly as a function of sintering temperature. The rate of sintering seems to be higher in argon than in nitrogen atmosphere.

3.3 Sintering of Ti-TiN nano-composites Figures 2(b) to 2(c) show the variation in sintered density as a function of sintering temperature for compacts of Ti-8wt%TiN and Ti-15 wt%TiN sintered in argon and nitrogen atmospheres. The presence of denser TiN (p = 5.5 V.V. Dabhade et al. RM57 41

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

gm/cm3 for TiN as compared to p = 4.5 gm/cm3 for Ti) tend to exhibit a higher sintered density.

Figure 3. shows the XRD pattern of elemental Ti sample, Ti-8wt%TiN and Ti- 15wt %TiN nano-composites sintered in nitrogen atmosphere at 1000°C. The XRD shows that the nitrogen dissolution/reaction at the sintering temperature leads to precipitation/formation of not only the equilibrium TiN and Ti2N phases but also various metastable phases. Due to the uncertainty of quantitative contribution of each of these phases to the overall increase in the density, it has not been possible to find the contribution of heat of reaction to the sintering rates of titanium nor is it possible to ascertain the exact level of densification of these samples.

Samples sintered in nitrogen atmosphere exhibited a golden yellow color and considerable weight increase. The X-ray diffraction analysis (Figure.3) indicated the weight increase of the samples sintered in nitrogen atmosphere could be attributed to the formation of various nitrides. A weight increase of the order of 13% -17% for samples sintered at 1000°C, according to the Ti-N equilibrium diagram [14] should result in the formation of TiN and Ti2N equilibrium phases, but XRD revealed the formation of other non-equilibrium phases like TiN03 and Ti3N129 along with the equilibrium ones. These phases may have formed due to the non-equilibrium cooling from the sintering temperature.

Samples sintered in argon atmosphere showed an overall decrease in volume and considerable densification. Densities of the order of 85% could be reached at 1100°C which could be attributed to the fine particle size of the starting powders.

Table 1: Percentage change in radius values (AR/RO x 100) as a function of sintering temperature for samples sintered in argon as well as in nitrogen atmosphere (negative indicates shrinkage while positive indicates expansion of the sintered samples).

Sr. Sample 900°C 950°C 100C°C 1050°C 1100°C no Ar N? Ar N2 Ar N2 Ar N2 Ar N2 1. Ti -0.5 1.6 -1.2 1.9 -1.2 2.3 -2.3 2.7 -3.3 3.1 2. Ti-8wt%TiN -0.8 1.5 -2.6 1.7 -2.0 2.1 -2.8 2.8 -4.3 3.1 3. Ti-15wt%TiN -1.2 1.3 -2.1 1.8 -2.1 2.3 -2.9 2.4 -3.8 3.2 42 RM57 V.V. Dabhade et al. 15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

4.0-1

£3.9H ü Ar atmosphere "E 3.8- s N2 atmosphere .|^3.7H c •a 3.6- •o 2 Ä 3.5- c 60 3.4 900 950 1000 1050 1100 Sintering temperature (°C) Figure 2(a): Variation of sintered density as a function of sintering temperature for pure titanium samples sintered in argon and nitrogen atmosphere.

4.0n Ar atmosphere

N2 atmosphere

900 950 1000 1050 1100 Sintering temperature (°C)

Figure 2(b): Variation of sintered density as a function of sintering temperature for Ti-8wt%TiN samples sintered in argon and nitrogen atmosphere. V.V. Dabhade et al. RM57 43 15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

4.0-, Ar atmosphere

N2 atmosphere

900 950 1000 1050 1100 Sintering temperature (°C) Figure 2©: Variation of sintered density as a function of sintering temperature forTi-15wt%TiN samples sintered in argon and nitrogen atmosphere.

Ti-15w!%TiN

Ti-8wt%TiN

48 68 26 (degree) • Figure 3: XRD pattern of pure Ti, Ti-8wt%TiN and Ti-15wt%TiN samples sintered at 1000°C in nitrogen atmosphere. 44 RM 57 V.V. Dabhade et al.

15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Table. 1 gives the percentage change in radius values as a function of sintering temperature for samples sintered in argon and nitrogen atmosphere. The shrinkage values from table. 1 show a maximum shrinkage below 4% thus indicating that the sintering is confined to the initial stages of sintering.

Taking into consideration the most general expression for volume diffusion in order to study the sintering kinetics,

2 3 y = 5.23yQDvt / KTa (1)

Where: y = AL/LQs 1/2(AR/R0) = fractional radial shrinkage, Y = surface energy (J/m2), Q = volume of Ti vacancy (m3), QV/RT 2 Dv = volume diffusion coefficient = Dvoe" (m /sec), K = boltzmans constant, T = temperature (°K), a = particle radius (m), t = time (sec).

Since all the experiments have been performed with a rate of 1° / 2 sec, T /1 = C is taken as 0.5 °K / sec. Substituting the latter, eqn. (1) can now be modified as:

y = (5.23yQ / CKa3)a5exp-(0.5Q / RT) (2)

Figure 4(a) and 4(b) show the typical Arrhenius plots i.e. Iny Vs 1/T for pure Ti samples sintered in argon and nitrogen atmospheres respectively. The slope of the plot gives 0.5QD / KT while the intercept on the y-axis gives 3 0.5ln(5.25yfiDo / CKa ). The computed values of Do and Q are given in table. 2 for all the specimens. V.V. Dabhade et al. RM57 45

15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

-3.5

-4.0

C -4-5 H

-5.0-

-5.5 12 7A 7.6 7.8 8.0 8.2 ' 8.4 ' 8.6 1/T°K

Figure 4(a): Arrhenius plot of pure Ti sintered in argon atmosphere.

-3.6-

-3.8-

-4.0-

-4.2 7.2 7.4 7.6 7.8 8.0 8.2 8.4 8.6 1/T°K

Figure 4(b): Arrhenius plot of pure Ti sintered in nitrogen atmosphere. 46 RM57 V.V. Dabhade et al. 15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Table 2: Frequency factor (Do), activation energy (Q) and diffusion coefficients (D) for samples sintered in Ar and N2 atmospheres.

Sr Ti Ti-8wt%TiN Ti-15wt%TiN no Ar N2 Ar N2 Ar N2 D„ 1. (m2/s) 7.9 x 10"8 1.7 xlO"13 3.4 xlO"8 6.1 x 10"13 2.2 x 10"" 1.2 xlO"12

Q 2. (KJ/mol) 237.4 89.7 221.83 104 140.82 112.12

3. D 1.3 x 10" 2.4 x 10"'7 3.4 x 1017 3.4 x 10"17 3.1 x 1017 3.0 x 10"17 (m2/sec)

Sintering of Ti in argon atmosphere required an activation energy of 237.4 8 2 KJ/mol while the Do value obtained is 7.9 x 10" m /sec. The reported value in 2 6 2 literature [15] of Do varies from 1.09 x 10" to 4.54 x 10" m / sec and Q values for self-diffusion of Ti vary from 131-251.2 KJ / mol. The present values obtained are in this range although both Do and Q values are lower than the reported values. Change in the sintering atmosphere from Ar to N2 had decreased the sintering activation energy drastically to 89.7 KJ/mol. Similar trend was noted in Ti-8wt%TiN and Ti-15wt%TiN samples.

The lower values of activation energy in general were observed in systems where the thickness of the defective interparticle region is likely to be larger than the neck dimensions [16], in such a case it is difficult to distinguish between volume diffusion and grain boundary diffusion mechanisms as both these are taking place through a defective interparticle region and hence can be classified as boundary enhanced diffusion.

Addition of TiN to samples that are sintered in Ar in general decreased the activation energy of sintering. This could be due to the active nano sized TiN particles (may be even metastable in composition) in contact with Ti causing diffusion of nitrogen along the surfaces.

Sintering of samples in N2 atmosphere is likely to be dominated more by the in situ nitride formation which is exothermic in nature. Almost all the samples V.V. Dabhade et al. RM 57 47^

15lh International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

have expanded during sintering possibly due to formation of several nitrides as confirmed by XRD. Thus the observed activation energy is a combination of energy required by the system for sintering and energy given out during nitridation. At this stage it is difficult to determine the separate contribution of each of these mechanisms.

Using the Do and Q values the diffusion coefficient D at 1000°C was computed, the values of which are shown in table.2. It is evident from the table that although the Do and Q values are varying, D remained more or less constant at 3.0 - 3.4 x 1CT17 m2/sec for all the systems excepting in the case of Ti sintered in pure Ar where it is 1.3 x 10~17rn2/sec.

The observed D values at 1000°C are smaller than those reported for self- diffusion of titanium in ß-titanium [15]. Further the magnitude of D suggests that it is more likely to be related to the self diffusion of nitrogen in titanium nitrides (this is perhaps the reason that the computed D values for systems sintered both in N2 as well as in Ar are nearly the same). During the early stages of sintering the nitrogen in the pores is likely to react with the titanium particles to form titanium nitrides. Further growth of this layer depends on the self diffusion of N2/Ti through the layer either way. Over a period of time sufficient amounts of nitrides will be formed. The lower value of activation energy were observed in systems where the thickness of the defective interparticle region is likely to be larger than the neck dimension (as mentioned earlier). In the present case the micron sized titanium particles are associated with a large free surface which is likely to contain adsorbed species such as O2 and N2 from the atmosphere and also surface defects. When such particles are brought into contact, the defective region at the point of contact is much larger in thickness than the interparticle neck width. In such a case the large sized nitrogen ion diffuses through the grain boundary to the neck surface while the smaller titanium ion is transported to the neck surface by lattice diffusion. However, since the neck width is much smaller than the width of the interparticle defective region both the diffusion processes take place through this defective region. In fact this process can now be called boundary enhanced diffusion [16].

The lack in variation in D values suggests that the contribution of diffusion process towards sintering of Ti-TiN mixtures and Ti in nitrogen atmosphere is minimal as compared to formation of titanium nitride by chemical reaction. Since the latter is an exothermic reaction the decrease in observed activation energies can also be justified by the same reason. 48 RM 57 V.V. Dabhade et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Figures 5(a) to 5(c) shows the SEM micrographs of the Ti sample, Ti-8 wt%TiN and Ti-15wt% TiN nano-composites sintered in argon atmosphere at 1000°C while figures 5(d) to 5(f) show the SEM micrographs of the Ti sample, Ti-8wt%TiN and Ti-15wt%TiN nano-composites sintered in nitrogen atmosphere at 1000°C. The SEM micrographs of samples sintered in N2 and Ar atmospheres reveal distinct features. The Ar sintered samples show increasingly well dispersed TiN particles among the micron sized Ti particles. Since the sintering process is in the initial stages considerable amount of porosity can be seen. The N2 sintered samples on the other hand showed the entire sample surface to be covered by TiN filling where this film has got ruptured the inside structure is seen consisting of coarse TiN particles and porosity.

4. CONCLUSION:

1. Initial stage of sintering of pure Ti as well as of Ti-TiN nano-composites has been studied by using constant rate of heating technique. The frequency factor evaluated varied between 7.9 x 10"8 to 2.2 x 10"11rn2/s in Ar 13 12 2 atmosphere and between 1.7 x 10~ to 1.2 x 10~ m /s in N2 atmosphere. The values are in the same range as those reported in literature. 2. The activation energy for sintering of Ti specimens in N2 atmosphere is somewhat lower than the reported values and varied between 89.7 - 112.12 KJ / mol. It is believed that this is most likely due to grain boundary enhanced diffusion predominating in such fine particulate system. 3. Sintering of Ti and Ti-TiN nanocomposite samples in N2 atmosphere is affected by two competing mechanisms such as endothermic energy requirement for mass transport and an exothermic in situ nitridation reaction. More work is required to evaluate the individual contribution for each of these mechanisms. 4. The lack of variation in the D values suggest that the sintering of Ti in nitrogen atmosphere and Ti-TiN nanocomposites in argon atmopshere as well as in nitrogen atmosphere is most likely controlled by the nitridation reaction rather than by the diffusion process. V.V. Dabhade et al. RM57 49

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

"l N

«f. • -fci

(a) (b)

20p.m

^ S

•(e)

20|im Figure 5: SEM micrograph of (a) pure Ti, (b) Ti-8wt%TiN, (c) Ti- 15wt%TiN samples sintered at 1000°C in argon atmosphere and SEM micrographs of (d) pure Ti, (e) Ti-8wt%TiN, (f) Ti-15wt%TiN samples sintered at 1000°C in nitrogen atmosphere. 50 RM57 V.V. Dabhade et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

REFERENCES:

1. K.H. Kramer, "Titanium applications - a critical review", Proceedings of The Sixth World Conference on titanium 1988, Part 1, Edited by P. Lacombe, B.Trieot, G.Beronger, pp.523-529. 2. T.R. Rama Mohan, T.S. Murty, P. Ramakrishnan and Elazer Y. Gutmanas,"Titanium nano-composite powders by attribution milling", Advances in Powder Metallurgy and particulate Materials, APMI, Compiled by Charles L. Rose and Martin H. Thibadeau, Metal Powder Industries Federation, Vol. 1, 1999, pp. 159-164. 3. Peter C. Eloff, Sintering of titanium, Metals Handbook (Ninth Edition), Vol. 7, Powder Metallurgy, ASM Handbook Committee, Coordinator Erhard Klar, 1984, pp. 393-399. 4. V.S. Arunachalam, "Powder Metallurgy of Titanium", Titanium and Alloys, vol. 3, The Metallurgical Society of AIME Proceedings, Edited by J.C. Williams and A.F. Below, pp. 2305-2314. 5. V.S. Ustinov, A.M. Petrunka, Yu G. Olesov and R.K. Ogner, "A novelty in The field of titanium powder metallurgy", Titanium and Titanium Alloys, Vol.3, The Metallurgical Society of AIME Proceedings, Edited by J.C. Williamsand A.F. Belov, pp.2315-2321. 6. S. Ranganath, "A review on particulate reinforced titanium matrix composites", Journal of Materials Science, 32, 1977, pp. 1-16. 7. Shierif M. El-Eskandarany, Omari, M., KonnoT.J., Sumiyama K., Hirai T. and Suzuki, K., "Synthesis of full-density nanocrystalline titanium nitride compacts by plasma-activated sintering of mechanically reacted powder", Metallurgical and Materials Transactions A, Vol. 29A, July 1998, pp.1973- 1981. 8. P.S. Goodwin, A. Wisbey, H.S. Ubhi, Z. Kulikowski, P. Gasson, and CM. Wardclose, "Synthesis of particulate reinforced titanium MMC by Mechanical alloying", Titanium 95: Science and Technology, Vol. Ill, Proceedings of the Eight World Conference on titanium, Edited by P.A. Blenkinsop, W.J. Evans, H.M. Flower, pp. 2874-2881. 9. C. Suryanarayana, "Nanocrystalline Materials", International Materials Reviews, 1995, Vol. 40, pp.41-115. 10. V.V Dabhade, M.M Godkhindi, T.R. Rama Mohan and P. Ramakrishnan, "Nitridation reaction during sintering of mechanically milled titanium powders," communicated. 11. B. Günther and A. Kumpmann, "Ultrafine oxide powders prepared by Inert gas evaporation", Nanostructured Materials, Vol. 1, 1992, pp. 27-30. V.V. Dabhade et al. RM 57 51_

15"1 International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

12. S. Amarchand, T.R. Rama Mohan and P. Ramakrishnan," A novel chemical solution technique for the preparation of nano size titanium powders from titanium dioxide," Advanced Powder Technology," Vol.2, No.4, (2000), pp. 415-422. 13. M.J. Mayo,"Processing of nanocrystalline ceramics from ultrafine particles," International Materials Reviews, 1996,Vol.41,No.3, pp.85-113. 14. "ASM Engineered Materials Reference Book," Reference publications, ASM INTERNATIONAL, (1989), Materials Park, Ohio, USA, p.257. 15. R. Sundaresan, A.C. Raghuram, R.M. Mally and K.I. Vasu," Sintering of ß Titanium: role of dislocation pipe diffusion," Powder Metallurgy, 1996,Vol 39, No.2, pp. 138-142. 16. D,L Johnson and L. Berrin, "Grain boundary diffusion in the sintering of oxides," Sintering and related phenomena, Ed. G.C.Kucznsk, N.A. Hooton and C.F.Cubbor, Gordon and Breach, New York, 1967. 52 RM 100 A.V. Nojkina et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Title: SINTERING OF NANODIAMOND IN THERMODYNAMIC STABILITY AREA OF DIAMONDS

Authors: A.V. Nojkina1, N.A. Kolchemanov1, P.Ya. Dedkov2

Affiliation: * VNIIALMAZ, Moskow, Russia; 2 VNIITF, Snezhinsk, Russia

Diamond powders UDA have mentioned below physico-chemical propeerties: - light-grey colour - grain size 0,02-0,004 mkm (biggest grain size is 20nm, smallest - 4nm), it is permitted the presence of the lumps. It is found the separate mi- crocrystals with size 60-90nm. - density - 2.97 - 3.0 g/cub.sm. - specific surface - 250-350 sq.m/g. - moisture - 3.00% mass. - admixtures (incombustible residue) - less than 2.0% mass. - admixture mass, less then: iron - 0.025% mass., manganese - 0.01% mass., molybdenum - 0.01% mass., lead - 0.04% mass., sulphur - 0.38% mass., chlorine - 0.33% mass. - temperature of beginning of intensive gas emission under heating in air - 753K - crystal structure of the UDA - cubic add hexagon (hexagular) - element composition of UDA: carbon - 86.91% mass., nitrogen - 1.92% mass., hydrogen -0.78%o mass., oxygen - 9.2% mass., organic and inorganic functional groups, water. A.V. Nojkina et al. RM 100 53^

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According results of emissive spectrum analysis the admixtures content does not exceed 2%. Ni and Mn are the active catalyst of grafitization, but their content is two degrees lower then in the diamonds of stationary synthesis. Specific surface of nanodiamond powders is ten times more than she- cific surface of diamond micropowders of stationary synthesis. That's why it is possible to consider that UDA contains the less quanlity of the admixtures than the diamonds of stationary synthesis. It was studied the surfaces of ten types of diamond powders of dynamic synthesis and natural diamond powders. Method of the thermopro- grammed desorption and mass-spectrometric analysis of desorption products were used for this study of the surfaces. Fig. 1-4 show the kinetic curves of H?O and CO2 desorption. It was found that the composition and the state of the surface of the powders UDA and BDA differ essentially. The mass- spectrums corresponding to temperature 400°C witness this. It was found that mainly the water and carbon dioxid are on the surface of UDA powders, but besides of the water and carbon dioxid the hydrocarbons with M=58 a.w. are on the surface of BDA powders. It has been studied kinetics of desorption of H2O and CO2 under linear heating of the powders to 400°C. The characteristic curves of H2O and CO2 desorption are shown at fig. 1. It may be seen that the surfaces of studied types of powders have the strong sorption activity to H2O. Desorption curves pass through two or three maximums of desorption rate. It depends on the types of powders and witnesses of presence of few active adsorption centers on the surface. Physically adsorbed water has poor bond with the surface and is moved away under low temperature. We have studied the surfaces of natural diamond powders too. The characteristic features of the last ones are the biggest content of CO with ac- 54 RM 100 A.V. Nojkina et al.

15m International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wüdner, Plansee Holding AG, Reutte (2001), Vol. 4

tivation energy of 7.9 kJ/mol and nearly absolute lack of CO2 sorption on the surface. Superhard materials of different sorts of detonation diamond powders have been made at the press-device DO-137A in the high-pressure chamber "toroid" - type. The same eqwipment was used for making of the composite diamond-content materials of£.DA and UDA mixes with static synthesis diamonds and impregnating material. Sinters density rises with increase of pressure. This fact was established in the course of sintering of the different marks of acto nation diamond powders. The maximum density that was obtained under the pressure of 90 kbar in the AVD "toroid" - type is different for tested marks. It is: for BDA - 24 g/sq. sm.; LDA - 2.4 g/sq. sm.; UDA - 2.35 g/sq. sm. Density of the sinters that were obtained in polypuncheon apparatus under the pressure 80 kbar over 18 hours is 3.1 for LDA and 2.9 for BDA. It is higher than density of the sinters that were obtained in AVD"toroid" - type, but it is somewhat lower than density of the sinters natural submicron powders. Density decreases in the row: LDA-BDA-UDA. It corresponds with density volume of the tested powders ( , 3,23; 3,0; 2,9) g/cub. sm. and corresponds with available data of 0ES. It was determined that density rises with increase of the lonsdeylite content. So rise of the sample density depends of pressure and percentage content of the lonsdeylite. It was explained not only by the porosity decrease with the pressure rise but by increase of the share of diamond bonds in the sample. It corresponds with available data of 0ES (fig. 5) and with results of ratio-phase analysis of the sinters. It was studied the carbon states in the experimental samples (Fig.5). Quantitative measuring has shown that carbon state, corresponding to SP3 - A.V. Nojkina et al. RM 100 55^

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hybridization, is less expressed for the samples of UDA type than for the samples of LDA type. Influence of the pressure on transformation of the graphite to diamond was studied too. It was found out that the pressure increase from 75 to 115 atm leads to displacement of CIS maximum the direction of the maximum in the natural diamond spectrum (fig.6). We think that this fact is conditioned by increase of the portion of carbon that has turned into diamond. Displacement of CIS maximum for sinter of the powder UDA (Pi=75 atm; U=3,5 V) is 0,8 mV and for the same sinter under P=115 atm and U=3,5 V the displacement is 0,5 mV. It was studied the temperature influence on the sinters properties. It was determined that the density rises with the temperature increase under sintering in the area of the diamond stability. Density decrease under the temperature corresponding to the area of the diamond metastability (fig.7). The impregnation increases the density to a greater degree than the mixture sintering (up to 3.2 g/cub. sm.), and temperature does not influence on the density . It witness about high thermal stability of the samples. Den- sity of the sinters, that were obtained of the original diamond powders of detonation synthesis, decreases at the temperature higher than 1800°C. But density of the sinters, that were obtained of diamond powders of detonation synthesis with silicon impregnation does not decrease up to the temperature 2400°C. Thus, it was determined that PCD samples of LDA with the density more than 3.2 g/cub. sm. may be obtained by means of impregnation with liquid metals in Russian apparatus AVD "to'roid" - type. Made PCD are characterized by thermal stability more than 1200°C. Their hardness is more than 6000kg/mm2, cracking resistivity - 10 MPa/m°>5. 56 RM 100 A.V. Nojkina et al.

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rb'dhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Experimental sintering of LDA, UDA, BDA has determined the correlation between quality of the sinter and quality of original diamonds of detonation synthesis. The same correlation has been obtained for diamonds of statical syn- thetics in comparison with natural diamonds. In shot, the more perfect is the structure of original raw materials, the higher are the physico-mechanical properties of obtained sinters. The strength of diamond powder of 60/70 mesh obtained from crushed sinter samples was measured to estimate the quality of the diamonds. After the test was complated it was found that the strength of the powders obtained from the nanodiamond sinters corresponds with the strength of the analogous fraction of the polycrystalline powders obtained from the diamonds made under static synthesis conditions. In the conclusion we consider sintering of nanodiamonds as the good option to use them in the wearproof tools and composites A.V. Nojkina et al. RM 100 57

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400 T°C

Fig. 1 Kinetic of H2O and CO2 desoprtion from the surface of power UDA. 58 RM 100 A.V. Nojkina et al.

15:h International Plansee Seminar, Eds. G. Kneringer, P. Rfidhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

M d.W.

300 T°C

Fig. 2 Kinetic of H2O and CO2 desorption from the surlace of the UDA (a) 11 and mass - spectrograms of destruction under 400°C (b). A.V. Nojkina et al. RM 100 59

15:" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

.Htß

Cy Hio.

Fig. 3 Kinetic of H2O and CO2 desorption from the surface of the BDA (a) and mass - spectrograms of destruction under 400°C (b). 60 RM 100 A.V. Nojkina et al.

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15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

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15" International Plansee Seminar, Eds. G. Kneringer, P. Ro'dhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. •

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15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

p, rria

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1000 2000 3000 4000 7, "C

to ? Phase diagram for element carbon (Bandy, 1989) 1 - area of the catalytic in metal presence diamond synthesis of graphite; 2 - area of direct transformation between graphite and diamond; 3 - area of direct transformation between diamond and graphite; 4 - area of direct transformation between graphite and lonsdaylit; 5 - area of the catalytic transformation between diamond and graphite. AT0200470

64 RM 19 Th. Franco et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Investigation of Porous Metallic Substrates for Plasma Sprayed Thin-Film SOFC

Thomas Franco, Rudolf Henne, Michael Lang, Robert Ruckdäschel and Günter Schiller

Deutsches Zentrum für Luft- und Raumfahrt (DLR) Institut für Technische Thermodynamik Pfaffenwaldring 38 - 40, D - 70569 Stuttgart, Germany

Summary:

For plasma sprayed thin-film SOFCs porous metallic substrates were investigated regarding the required properties. The main essential properties are a sufficient electrical conductivity, a high temperature corrosion resistance and an adapted thermal expansion coefficient with respect to the other SOFC components. Proceeding from these requirements metallic materials such as a porous Cr-ODS alloy (Cr-5Fe-1Y2O3), a stainless steel plate (Fe-18Cr-12Ni- 2Mo) and a Ni felt were tested. Additionally, to investigate the metallic substrates and their interaction with the layers of the cells completely plasma sprayed SOFCs were fabricated and electrochemically characterised.

Keywords:

Porous metallic substrates, SOFC, VPS, porosity, corrosion resistance, electrical conductivity, thermal expansion

1. Introduction

Solid Oxide Fuel Cells (SOFCs) convert electrochemically the chemical energy of fuel gases (H2, CO, CH4 and other hydrocarbons) with oxygen or air directly into electrical power. In comparison to the power generation by means of combustion processes, the electrochemical oxidation of the fuel gases in a SOFC leads to higher overall efficiencies with very low harmful emissions. Th. Franco et al. RM 19 65

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The central unit of a SOFC is consisting of two electrodes (a porous (La,Sr)MnO3 cathode, and a porous ZrO2/Ni anode) separated by a gastight solid ZrO2 electrolyte (Fig. 1). This combination is also called "PEN" (positive electrode, electrolyte, negative electrode) or "MEA" (membrane electrodes assembly).

Cathode

o, •» o p 4 2 Solid electrolyt" e o2-!

H2O H2O

Anode

Fig. 1: Principle of a Solid Oxide Fuel Cell

An oxidant (in this case oxygen) is fed to the porous cathode, where it ac- cepts electrons from the external circuit, and undergoes a reduction reaction:

2_ (Eq. 1.1)

The electrolyte conducts the oxygen ions due to the difference of concentration. Fuel (in this case H2) is fed to the Anode, where it undergoes an oxidation reaction with the oxygen ions, and releases electrons to the external circuit:

(Eq. 1.2)

The overall reaction can be written as:

(Eq. 1.3) 66 RM19 Th. Franco et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

This is a strong exothermal reaction. The Gibbs free enthalpy is AG°= -228,7 kJ/mol.

In the first generation of planar SOFCs, an approximately 150 urn thick solid electrolyte has been necessary in order to support the cell mechanically (1,2). Because of this high thickness and the strong dependence of the ionic conductivity of the ZrO2 with temperature, the cells are usually operated between 900-1000 °C. The high thermal load of the materials can result in detrimental diffusion and evaporation processes with a strong reduction of the cell performance and of long-term stability. Another disadvantage of the electrolyte supported cells is the low fracture toughness of zirconia, the danger of crack formation and hence the very limited size of the cells which can be fabricated and operated reliably.

In order to increase stability, life time and economy, the operating temperature of the cells has to be reduced to about 700-800 °C. For this purpose it is necessary to decrease the overall thickness of the SOFC and particularly of the electrolyte. This means that no longer a relatively thick electrolyte serves as "backbone" of a cell. With this second-generation SOFC characterised by a thin-film electrolyte either one of the electrodes (3-5) or an additional component - a porous metallic substrate - has the function to mechanically stabi- lise the cell.

2. The DLR Plasma Spray SOFC Concept

A novel concept for a metallic substrate supported thin-film SOFC has been developed at DLR Stuttgart (6,7). In the DLR concept, the entire cell is deposited onto a porous metallic substrate by an integrated multistep vacuum plasma spray process (see chapter 3). Its characteristic properties such as short process time, high material deposition rate and the ability to be transferred to an automated production line promise a fast and cost-effective fabrication of cells with large active cell areas. Because of this substrate support, the electrolyte layer may have a significantly reduced thickness of only 20-30 urn, resulting in a thin-film cell with a total thickness of less than 100-120 urn.

The principle and design of the substrate supported DLR spray concept is schematically shown in Fig. 2. In this concept the substrate supported MEA is fitted into a recess within the bipolar plate and sealed by a glass sealant layer. Th. Franco et al. RM 19 67

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Bipolar Plate Oxidant Gas Supply Contact Layer, Cathode Protective Layer Electrolyte Anode Sealant Porous Layer Substrate

Fuel Gas Supply

Fig. 2: Principle of SOFC design according to the DLR spray concept

Fig. 3: Vacuum plasma sprayed SOFCs on porous metallic substrates with active cell areas between 5 and 320 cm2 68 RM 19 Th. Franco et al.

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

For practical reasons the porous substrate is fixed within the bipolar plate prior to the deposition process and the sealing procedure is performed after the manufacture of the entire MEA structure. The fuel gas is fed to the porous substrate (gas distributor) through gas channels within the bipolar plate in counter flow to the air on the cathode side. In general, for the assembling of SOFC stacks square-shaped cells with 10 cm x 10 cm and 20 cm x 20 cm with active areas up to 320 cm2 were produced (Fig. 3).

3. The Plasma Spray Process (VPS)

The plasma spray process is based on the generation of a plasma jet consisting of argon or argon with admixtures of H2 and He which are ionised by a high current arc discharge in a plasma torch. Fig. 4 shows schematically the principle of the DLR plasma torch with a Laval-like nozzle contour.

L val Nozz e Plasma- Electrical High ?u ; ' ,J Substrate gases Current Arc with integrated powder injection |_ayer Water Cathode Anode Cooling Plasma Jet

(•A"*1 • i

Fig. 4: Principle of the DLR plasma torch

Powders to be sprayed are injected into the plasma where they are accelerated, melted and finally projected onto a substrate. The coating is formed by solidification and flattening of the particles at impact on the substrate. Carry- ing out the spray process in a vacuum chamber with reduced pressure a long and laminar plasma jet with high velocity and reduced interaction with the surrounding cold gas is formed resulting in improved spray conditions. The plasma torch is moved by means of a robot system to cover the whole surface of the component to be coated in a reproducible way. Th. Franco et al. RM 19 69

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4. Experimental

A material to be used as a porous metallic substrate for SOFC applications has to fulfil a variety of required properties. Beside high mechanical strength and a sufficient gas permeability some further requirements have to be met. Those are an excellent electrical conductivity, an adequate corrosion stability at SOFC operating conditions and an adapted thermal expansion behaviour with respect to the MEA layers. Proceeding from these requirements the following substrate materials have been investigated and tested:

Thickness Porosity Substrate Material composition Supplier [mm] [vol. %] Plansee, Cr-ODS plate Cr-5%Fe-1%Y O -1.0 -35 2 3 Austria Stainless steel Fe-18%Cr-12% Ni- GKN, -2.0 -45 plate 2%Mo Germany Medicoat, Ni felt Ni -2.0 -80 Switzerland Table 1 : Substrates studied for the plasma sprayed SOFC

In a second step, to investigate the porous metallic substrates and their interaction with the MEA layers, SOFCs have been fabricated and subsequently electrochemically characterised. For this purpose cells with active areas of 12 cm2 were completely plasma sprayed onto different metallic substrates under same producing conditions. The selection of the substrates depended on the results of the previous material investigations. The powder materials, which have been used for the spray process are listed in table 2.

ZrO - Powder NiO 2 (Lao.8Sr0.2)o.98 7 mol% Y2O3 MnO3 Short name NiO YSZ LSM Particle size <28 urn <22 urn 20 - 50 urn Medicoat, Swit- Medicoat, EMPA, Supplier zerland Switzerland Switzerland Table 2: Powders used for spraying of SOFC layers 70 RM 19 Th. Franco et al.

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5. Results and Discussion

5.1 Electrical Conductivity

For the use of porous metallic materials as SOFC substrates, the ohmic resistance should be as low as possible. Therefore, a high electronic conductivity of the substrates is required. To investigate the conductivity, a 4-point measurement which offers the possibility to measure in SOFC relevant atmospheres was used. Fig. 5 shows the electrical conductivity of the porous metallic substrates as a function of temperature, measured in Ar-5% H2-2% H2O atmosphere.

1.0E+07 |

1.0E+06

• D O Cr-ODS D O O a D ©•© O o o O Ni-felt 1.0E+05 A A A A \ Stainless steel

1.0E+04 200 400 600 800 1000 1200 T[°C]

Fig. 5: Electrical conductivity of different porous metallic substrates as a function of temperature, measured in Ar-5% H2-2% H2O

The conductivity of the Cr-ODS substrate is nearly one order of magnitude higher than the conductivity of the stainless steel plate and about half an order of magnitude higher than the conductivity of the Ni felt. According to reference (8) the value of bulk nickel should be higher compared to the electrical conductivity of the Cr-ODS and the stainless steel plate. But, due to the high porosity (Table 1) the electrical conductivity of the Ni felt is significantly re- Th. Franco et al. RM 19 71_

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duced. In accordance with reference (8) the stainless steel plate shows the lowest conductivity in the measured temperature range.

As expected for metals, the electronic conductivity decreases when the temperature is rising. The Cr-ODS alloy shows the highest and the stainless steel plate the lowest temperature dependency. In the case of the Ni felt a bend in the conductivity curve can be observed at about 360 °C (Curie Point), which is caused by the ferromagnetic to paramagnetic transition. The same phenomenon can be observed with a smaller extent in the stainless steel curve because of 12 wt. % of Nickel in the alloy (Table 1).

5.2 High Temperature Corrosion Resistance

In general, metallic substrates to be used for SOFC application have to possess an excellent high temperature corrosion resistance in humid and reducing atmospheres. To investigate the corrosion behaviour of the metallic substrates they have been annealed under SOFC relevant conditions and subsequently characterised by optical microscopy and X-ray diffraction. Fig. 6 shows the cross sections and the corresponding X-ray diffraction patterns of the investigated substrates after annealing at 900 °C for 50 hours in Ar-5% H2-2% H2O atmosphere.

The Cr-ODS and the stainless steel plate show a significant formation of oxide (bulk material: white, oxide: grey). The Cr-ODS plate has formed an 5 -10 urn thick Cr2O3 layer on the surface of its Cr particles, whereas the stainless steel plate shows a significant formation of FeCr2O4 inside its bulk material. The reasons of these relatively strong oxide formations are given on the one hand in an oxygen pollution of the gas atmosphere where substrates have been annealed. On the other hand, according to literature (9), an adsorptive dissociation of water molecules of the 2% H2O steam on the surface of the metallic substrates is possible in the given atmosphere. In this case the oxygen of the water steam (dissociation product) reacts directly to oxide with the substrate material.

However, the Ni felt (Fig. 6b) shows an excellent oxidation stability in the used atmosphere. According to reference (10) a reversible oxidation reaction in atmospheres with hydrogen partial pressures > 10~6 bar is possible. In the annealing atmosphere a hydrogen partial pressure of about 50 x 10"3 bar was existing. 72 RM 19 Th. Franco et al.

15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

350

300 x Cr

o Cr2O3 (/} 200 c CD 150

100

0 *MV, 20 30 40 50 60 70 80 a) Cr-ODS plate Angle 2 0 / [°]

x Ni

>, ',-tw; 6000 c

^ 4000

2000

20 30 40 50 60 70 80 90 b) Nickel felt Angle 2 0 / [°]

20 30 40 50 60 70 80 90 c) Stainless steel plate Angle 2 0 / [°] Fig. 6: Cross sections and corresponding X-ray diffraction patterns of the investigated substrates after annealing at 900 °C in Ar-5% H2-2% H2O Th. Franco et al. RM 19 73

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

In general, the formation of oxide on porous metals is essentially higher compared with bulk materials because of its larger surface. Such oxide formation as observed in the case of the Cr-ODS and the stainless steel plate during SOFC operation can considerably reduce the electrical conductivity, especially when the sintered contacts between the bulk particles were oxidised (Fig. 6a).

5.3 Thermal Expansion Behaviour

For crack-free deposition of MEA layers onto large substrates (10x10 cm2) and for using them in a SOFC stack, a mismatch of the thermal expansion coefficients between the substrate and the MEA layers has to be kept as low as possible. The ideal expansion coefficient of metallic SOFC components with respect to the MEA layers is about 11 x 10~6 K"1 in the temperature range 30 - 1000 °C. To investigate the thermal expansion coefficients of the substrates a dilatometer (Netzsch, Selb) was used. Fig. 7 shows the temperature dependence of the thermal expansion coefficient of different porous metallic substrates in Ar-5% H2 atmosphere.

2,0E-05 A- A A A^/*

1.6E-05 • * ^ &

? ,r- 1.2E-05 •"~/ A •-•••• "ö 8,0E-06 •A A Stainless steel plate Nickel (>99,9%, lit. [11]) 4,0E-06 • • Cr-ODS plate (dense) Cr-ODS plate (porous) O.OE+00 i "• "i - () 200 400 600 800 1( Temperature / [°C] Fig. 7: Temperature dependence of thermal expansion coefficient of the different porous metallic substrates in Ar-5% H2-atmosphere 74 RM 19 Th. Franco et al,

15" International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The stainless steel plate shows a thermal expansion coefficient of 19 x 10"6 K"1 in the SOFC operating temperature range (700 - 900 °C), which is about 8 x 10"6 K"1 higher than the required expansion coefficient. For SOFC application such a thermal mismatch is not acceptable.

In the case of the Ni felt, a measurement with the dilatometer was not possible. Because of its structure (thin nickel wires) a deflection of the felt at about 400 °C occured during the measurement. Thus, for a rough estimation of the thermal expansion behaviour correlative values of bulk nickel from literature (8) were used. The curve of the bulk nickel shows an expansion coefficient of about 16 - 17x 10"6 K"1 in the relevant temperature range. On the one hand, this expansion coefficient is significantly too high and problems could arise especially with large cell areas (10x10 cm2 or 20 x 20 cm2). But on the other hand, in the case of the Ni felt the inherent flexibility of its structure is able to compensate this mismatch to a certain extent (11).

However, the Cr-ODS alloy (here a porous plate compared with a dense plate) shows an excellent adapted thermal expansion coefficient of about 10 - 11 x 10~6 K"1 between 700 - 900 °C. The reason for that is the content of 5 wt. % iron in the Cr based alloy (Table 1), which shifts the thermal expansion coefficient close to YSZ (yttria stabilised zirconia) (12). The porous Cr-ODS plate shows in spite of its porosity of 35 vol. % (Table 1) the same expansion characteristics than the dense alloy. This means, that the thermal expansion behaviour of this substrate material is not dependent of its porosity.

5.4 Electrochemical Performance of Substrate Supported Cells

Because of the high thermal expansion coefficient and the relatively bad oxidation resistance of the stainless steel plate the plasma sprayed cells have been produced using Ni felts and Cr-ODS plates. Fig. 8 shows the cross section of a completely plasma sprayed SOFC, consisting of a NiO + YSZ (ZrO2- 7 mol % Y2O3) anode, a YSZ electrolyte and a double layered LSM ((Lao.8Sr0.2)o.98Mn03) + YSZ / LSM cathode, which were consecutively sprayed onto a porous Cr-ODS plate. The plasma sprayed YSZ electrolyte exhibits a dense lamellar microstructure, whereas the anode has a fine open porosity of 21 vol. % after the reduction of the NiO with H2 to pure Ni. The overall porosity of the cathode reaches only about 10 vol. % with relatively coarse pores. Th. Franco et al. RM 19 75

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Recent developments aim at improving the cathode's microstructure and at optimisation of the thicknesses of the layers (13).

Fig. 8: Metallographic cross section of an entirely plasma sprayed cell on a porous Cr-ODS plate

The electrochemical performance of the substrate supported cells was mainly determined by performing I-V characteristics and impedance spectroscopy measurements. For a direct comparison of the electrochemical behaviour of the substrate supported cells it was important that the MEA layers were fabricated under identical conditions. This means, that the plasma sprayed MEA layers had nearly the same composition, the same thickness and especially the same structure.

Fig. 9 shows a comparison of the l-V-characteristics of different substrate supported cells with identical MEA layers (A: YSZ/NiO, E: YSZ, C: LSM/YSZ). SOFC test equipment which was developed at DLR-Stuttgart was used for the characterisation. The operating gases used were 0.5 SLPM H2 and 0.5 2 SLPM O2, the temperature was 900 °C and the active cell area 12 cm . 76 RM 19 Th. Franco et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

1200 600

500

400

300

200 a.

100

0 200 400 600 800 1000 1200 i / [mA/cm2]

Fig. 9: Comparison of the l-V-characteristics of different substrate supported cells, fabricated with identical MEA layers (900°C, 0.5 SLPM H2, 0.5 SLPM O2 as operating gases)

The OCV ("open circuit voltage") of the Cr-ODS cell is 900 mV, whereas the Ni felt cell exhibits an OCV of about 980 mV. In the case of the Ni felt cell a power density of about 420 mW/cm2 at a cell voltage of 0.7 V and a current density of 600 mA/cm2 was obtained. In contrast, the Cr-ODS cell has reached a power density of only 260 mW/cm2 at a current density of 380 mA/cm2 and the same cell voltage.

In general, the decrease of the OCV is mainly caused by a decrease of the activities of the operating gases on the electrodes. For that reason, two causes seem to be responsible. On the one hand, the lower porosity of the Cr-ODS substrate (Table 1) can cause an inhibition of diffusion processes, in which H2 is fed to the anode and H2O (as reaction product) is released from the anode (14). On the other hand, a volume extension can arise through the higher oxidation of the CR-ODS plate (see chapter 5.2) which causes cracks in the MEA. In this case, a strong diffusion of the process gases H2 and O2 through the MEA can occur and hence decrease the concentration of operating gases on the electrodes. Th. Franco et al. RM 19 T]_

15'h International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

6. Conclusion

The presented work has shown the results of material investigations on porous metallic substrates for plasma sprayed thin-film SOFCs. The investigated Cr-ODS alloy possesses a high conductivity and an excellent adapted thermal expansion coefficient to the SOFC layers. But in the SOFC relevant atmosphere it shows a relatively low oxidation resistance. In contrast, the Ni felt exhibits an adequate oxidation stability, but a relatively high thermal expansion behaviour. In general, it is difficult to find a material, which fulfils the required properties for this high temperature process. In the further development a combination of the excellent adapted thermal expansion coefficient of the Cr-ODS alloy and the adequate corrosion resistance of the Ni felt is currently investigated.

7. References

(1) H.J. Beie, L. Blum, B. W.Drenckhahn, H. Greiner, H. Schichl, Proc. 3th Eu- rop. Solid Oxide Fuel Cell Forum, (Philippe Stevens, ed.), Nantes, Vol. 1, pp. 3-14, 1998 (2) G.M Christie, R.C. Huiberts, E.J. Siewers, J.P.P Huiberts, Proc. 3th Europ. Solid Oxide Fuel Cell Forum, (Philippe Stevens, ed.), Nantes, Vol. 1, pp. 133-141, 1998. (3) D. Stöver, U. Diekmann, U. Flesch, H. Kabs, W. J. Quadakkers, F. Tietz, I. C. Vinke, Proc. 6th Intern. Symp. Solid Oxide Fuel Cells (SOFC VI), (S. C. Singhai, M. Dokiya, eds.), Honolulu, Hawaii, The Electrochemical Society, Proc. Vol. 99-19, pp. 812-821, 1999. (4) S.C. Singhal, 6th Intern. Symp. Solid Oxide Fuel Cells (SOFC VI), (S. C. Singhai, M. Dokiya, eds.), Honolulu, Hawaii, The Electrochemical Society, Proc. Vol. 99-19, pp. 39-50, 1999. (5) R. Bolden, K. Föger, T. Pham, Proc. 6th Intern. Symp. Solid Oxide Fuel Cells (SOFC VI), (S. C. Singhal, M. Dokiya, eds.), Honolulu, Hawaii, The Electrochemical Society, Proc. Vol. 99-19, pp. 80-87, 1999. (6) G. Schiller, R. Henne, M. Lang, S. Schaper, Proc. 6th Intern. Symp. Solid Oxide Fuel Cells (SOFC VI), (S. C. Singhal, M. Dokiya, eds.), Honolulu, Hawaii, The Electrochemical Society, Proc. Vol. 99-19, pp. 893-903, 1999. 78 RM 19 Th. Franco et al,

15lh International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutle (2001), Vol. 4

(7) M. Lang, R. Henne, S. Schaper, G. Schiller: to be published in Journal of Thermal Spray Technology, ASM International, (C.C. Bemdt, ed.), 2001. (8) Handbook of Chemistry and Physics, D.R. Lide (ed.), CRC-Press, New York, Tokyo, London, pp. 12-43 and 12-44, 1997. (9) M. Oertel, PhD-Thesis, Universtity of Stuttgart, performed at DLR Stutt- gart, 1993. (10) M. Lorenz, PhD-Thesis, Universtity of Essen, performed at DLR Stutt- gart, 1999. (11)M. Lang, T. Franco, R. Henne, R. Ruckdäschel, G. Schiller, P. Szabo Joint Topical Meeting IEA (Annex XIII: Solid Oxide Fuel Cells) and Euro- pean Science Foundation (OSSEP), Les Diablerets, Switzerland, pp. 92- 98,2001. (12) M. Janousek, W. Köck, M. Baumgärtner, H. Greiner, Proc. 5th, Intern. Symp. Solid Oxide Fuel Cells (SOFC-V), (U. Stimming, S.C. Singhai, H. Tagawa, W. Lehnert, eds.), Aachen, Germany, The Electrochemical So- siety, Proc. Vol. 97-18, pp. 1225-1233, 1997. (13) G. Schiller, R. Henne, M. Lang, R. Ruckdäschel, S. Schaper, Proc. 4th European SOFC Forum (A.J. Me Evoy, ed.), Lucerne, Vol.1, pp. 37-46, 2000. (14) M. Lang, PhD-Thesis, Universtity of Stuttgart, performed at DLR Stutt- gart, 1999. AT0200471

J. Warren et al. RM 36 79^

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A COMPARISON OF THE MICROSTRUCTURE AND HIGH TEMPERATURE TENSILE PROPERTIES OF A NOVEL P/M Mo-Hf-Zr-Ta-C ALLOY AND TZM

J. Warren® and G. Reznikov* GE Medical Systems Coolidge Laboratory, 4855 W. Electric Avenue, W. Milwaukee, Wl 53219 *X-Ray Tube Target Plant, 18683 South Miles Road, Warrensville Heights, OH 44128

Abstract

The microstructure and elevated temperature quasi-static tensile yield and ultimate strength observed in a novel, forged Mo-based alloy (Mo-0.25Hf- 0.25Zr-0.25Ta-0.025C) has been analyzed and compared to a standard forged TZM composition (Mo-0.50Ti-0.08Zr-0.02C). The novel material exhibits the desirable forging characteristics typical of the widely used TZM composition yet possess a higher ultimate strength and 0.2% offset yield strength in both the stress-relieved and recrystallized conditions over a 400°- 1200°C temperature range. The greater strength measured in the novel composition has been attributed to the combined effects of precipitation of Hf, Zr and Mo-(carbide) precipitates that strengthen the matrix in the classical Orowan fashion and improved resistance to recrystallization after high temperature exposure. Elevated temperature creep behavior, not addressed in the study presented here, will be reported on in a subsequent analysis.

Introduction

The development of precipitation strengthened molybdenum alloys such as MHC (Mo-1.0Hf-0.10C), TZC (Mo-1.25Ti-0.15Zr-0.02C) and TZM (Mo-0.5Ti- 0.08Zr-0.025C) have resulted in alloys that exhibit outstanding high temperature tensile strength and creep resistance. These alloy derive high temperature strength from a combination of mechanisms which include (1) the precipitation and dispersion of discrete carbide and oxide particles, (2) the retention of a heavily worked molybdenum matrix and (3) solid solution strengthening. Of the three alloy systems described above, TZM has gained the widest commercial acceptance due principally to this alloys relative ease of manufacture into wrought product forms. For ease of fabrication however, 80 RM36 J. Warren et al.

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a compromise in both elevated temperature strength and resistance to recrystallization at temperature extremes are realized.

This paper describes a novel powder metallurgical alloy system designated Alloy-2 with a nominal composition of Mo-0.25Hf-0.25Zr-0.25Ta-0.03C [1]. This particular composition was down-selected from a larger group of alloys fabricated within the constraints of an experimental design matrix; the goal of which was the development of an alloy that both possessed the favorable hot fabrication characteristics associated with the TZM composition and simultaneously exhibited a significant improvement in elevated temperature strength when compared to the TZM composition processed under identical conditions. Concentrations of the alloying agents hafnium and zirconium (carbide and oxide formers, respectively), tantalum (solid solution strengthened and carbon where systematically adjusted within the master mix, blended with pure molybdenum powder, pressed and forged into flat disks for subsequent mechanical analysis. Based on the hot yield and ultimate tensile strength data accumulated for all of the compositional variations investigated in the experimental matrix, the material analyzed in this report, Alloy-2, was selected as a promising candidate for continued evaluation.

Manufacturing cost is a significant consideration in the industrial setting and improvements in elevated temperature mechanical behavior must be assessed against both an alloys ease of fabrication and the potential advances in components manufactured from the alloy, especially with the added expenses associated with alloying agents rich in hafnium and tantalum. For example, in the manufacture of high power TZM rotating anode x-ray tube targets, the elevated temperature yield and creep resistance of TZM restricts the size and rotational speed of the hot target thus limiting useful x-ray output power. Higher hot strength target alloys that incur greater fabrication costs enable greater x-ray power output which, from the customers perspective, adds value that justifies the premium paid.

In this report, the elevated temperature yield and ultimate strength comparisons between Alloy-2 and TZM are made. Each alloy system was tested in two heat-treated conditions; a partially recrystallized condition denoted in this report as stress relieved and fully recrystallized. Recrystallization is especially of interest since the thermal processing normally associated with x-ray tube target manufacture (i.e. brazing and J. Warren et al. RM 36 81_

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vacuum degassing) usually results in a target alloy that is placed in service in a condition that has undergone recrystallization and grain growth.

Materials

Rectangular tensile test specimen blanks of both the TZM (Mo-0.5Ti-0.08Zr- 0.03C) and Alloy-2 (Mo-0.25Hf-0.25Zr-0.25Ta-0.03C) compositions were wire EDM cut from forged target disks 185mm in diameter and 18mm thick as shown in Fig. 1. The disks for this study were manufactured at the GE X-ray Tube Target Plant (Warrensville Heights, OH) and processed using parameters suited for the fabrication of molybdenum based x-ray tube target assemblies. Unlike monolithic wrought molybdenum alloys, x-ray tube targets are pressed and forged with an integral tungsten-rhenium (W-Re) outer annulus (i.e. track) about 1mm thick which serves as the focal spot for electron beam impingement and subsequent x-ray generation. Thus, the presence of the track in contact with the substrate, must be accounted for during processing. Theoretically, to achieve the highest possible strength, a solution heat treatment that effectively dissolves preexisting carbide inclusions and homogenizes the alloying elements within the matrix for subsequent controlled precipitation by aging is necessary. For example, solution treatment temperatures greater than 2300°C followed by rapid furnace cooling favor the kinetics of complete dissolution of the (Zr, Ti)C and Mo2C carbides in TZM. This heat treatment would then be followed by controlled aging, elevated temperature forging and final stress relieving at temperatures typically between 1100°-1200°C. However, to avoid the deleterious diffusion affects (Kirkendall porosity) associated with the W-Re track in contact with the molybdenum substrate at temperatures greater than 2300°C, sintering temperatures are limited in this production program to 2000°-2100°C; a temperature range where the kinetics of grain boundary Mo2C formation is favored [2]. 82 RM36 J. Warren et al.

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Figure 1. Method of specimen removal from a test forging.

The standard processing route used at GE for the production of TZM x-ray targets was used for the fabrication of the samples manufactured from the Alloy-2 material. The process begins by first cold pressing blended powder compositions of each alloy to a green density of approximately 0.65 theoretical. The density is further increased to nearly 0.95 theoretical by vacuum or hydrogen sintering between 2000° and 2100°. Using the same die and identical forging parameters, the sintered targets are preheated to 1500°-1550°C and subsequently hot forged. A moderate reduction in the through thickness of about 20% is achieved during forging resulting in a final density of about 0.99 theoretical for both the TZM and the Alloy-2 substrates. As discussed previously, stress relieving is accomplished at a relatively high temperature; 1500°-1550°C for 30 minutes in a dry hydrogen batch furnace. The forgings exit the hot-zone and are immediately cooled by convection in a hydrogen purged cold-zone outfitted with water-cooled muffles. In addition to having cooling rates faster than those obtained by air cooling, hydrogen cooling results in part surfaces that are oxide free. After deformation processing, the carbon content in each alloy was measured to be approximately 300-600 ppm by wt. Residual oxygen content is typically less than 30 ppm by weight for both alloy compositions.

Two separate heat treatments of each alloy were tensile tested between 400° and 1200°C. First, tensile specimens were extracted from forgings in the stress relieved condition. Second, in order to investigate and compare the J. Warren et al. RM 36 83_

15" International Plansee Seminar, Eds. G. Kneringer, P. Rddhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 effect of heat treatment on the structure and properties, the target disks were exposed to an extensive recrystallization vacuum heat treatment at 1900°C for 3h. Tensile specimens were then machined (EDM) from these target disks, tensile tested and compared. Specimens were removed from the target disk centers (see Fig. 1) far removed from the influence of the W-Re track located at the outer edge.

The microstructures of each alloy composition were studied using optical, SEM and TEM methods. EDS was used to acquire chemical composition data for the precipitates observed (both the SEM and TEM used in this investigation had this capability). In Fig. 2a and b, SEM photomicrographs of the TZM and Alloy-2 compositions respectively, exposed to the stress relieving heat treatment, are shown and compared. The microstructures shown are views which are normal to the through-thickness of the disk. Average grain sizes, measured using the Hilliard circle technique, were found to be similar at 22|um for the TZM alloy and 27|im for Alloy-2. Conventional, stress relieved TZM wrought product, subjected to extensive forging or hot extrusion processing will have grain sizes finer than the ones measured here. Also, the relatively high temperature stress relieving temperature used in this study results in recovery and partial recrystallization of the matrix.

In the stress relieved TZM composition two distinct precipitates were observed. SEM analysis revealed large Mo and Ti carbide precipitates, Mo2(Ti)C, located both in grain interiors and along grain boundaries while TEM analysis revealed an extremely fine (<50nm) Zr oxide, ZrO2, precipitate array located within the grain interiors. Based on the size and large separation distances of the carbide precipitates revealed in the TZM microstructure it is reasonable to assume that they do not contribute to strengthening in the classical fashion. An example of the two precipitate morphologies can be seen in the TEM micrograph shown in Fig. 3 where the small arrows highlight the extremely fine ZrO2 precipitate array and the large arrows locate a typical grain boundary carbide. 84 RM36 J. Warren et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Figure 2. Microstructure of A) TZM and B) Alloy-2 in stress-relieved condition.

By comparison, the stress relieved Alloy-2 composition possessed two precipitate types. Large oxide precipitates of Hf and Zr located within the grains and a much finer distribution of Hf, Zr and Mo-rich precipitates inside the grains with a size distribution ranging from approximately 30nm to about J. Warren et al.RM 36

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1[im with a separation distance which appears to be significantly smaller than the corresponding ZrO2 precipitate spacing observed in the TZM alloy. EDS spectra taken from these particles over the range of sizes observed were nearly identical suggesting that the extremely small precipitates are nuclei of the larger precipitates. Diffraction patterns were taken from these precipitates but, at the time of this writing, the analysis could not accurately determine their structure. Diffraction patterns typical of the MC structure were not observed. EDS spectra did not reveal the presence of oxygen hence the form is likely that of a carbide. A TEM micrograph of the material, Fig. 4, shows how these (carbide) precipitates, especially the extremely fine ones, are arranged in the matrix. The (carbide) precipitates are highlighted by the arrows. Also, the grain boundary shown in Fig. 4 is free of large precipitates and typifies the grain boundaries in this composition. Grain boundary precipitates were not observed in this composition. The (carbide) precipitates can be seen in the SEM photomicrograph shown in Fig. 2b (see arrows). As is evident from the figure, the precipitate density appears to vary significantly from grain to grain.

Higher magnification TEM micrographs are shown in Figures 5 and 6 and illustrate the extent of retained hot work in the TZM and Alloy-2 matrices due to forging and subsequent stress relieving, respectively. Both figures show a typical dislocation arrangement. Close examination of Fig. 5 also shows clearly that the extremely fine ZrO2 precipitates interact and impede dislocation slip generated during forging. The comparison (Fig. 6A and B) also reveals that Alloy-2 appears to contain a significantly higher number of dislocations than the TZM composition with many of the Alloy-2 grains containing a refined cell structure (6B) with some grains possessing a roughly equiaxed cell morphology. 86 RM36 J. Warren et al. 15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

1 micron

Figure 3. Large arrow indicates typical grain boundary carbide in stress-relieved TZM and the small arrows show typical matrix precipitate morphology. J. Warren et al. RM36 87

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1 micron

Figure 4. Arrows highlight precipitates in stress-relieved Alloy-2.

Recrystallizing each alloy results in significant changes in the microstructure. As expected the average grain sizes increased. For example, the TZM alloy increased to 48|um while for the Alloy-2 composition a value of 44^m was measured which indicates that the grain growth kinetics are similar between the two alloy systems. TEM analysis of each alloy also indicated that the ZrO2 in the stress relieved TZM and the Hf, Zr, and Mo-(carbide) precipitates in the stress relieved Alloy-2 coarsened to some extent as a result of the heat 88 RM36 J. Warren et al. 15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

treatment with a corresponding decrease in the separation distance between particles.

0.5 microns Figure 5. Arrow highlights dislocation interactions with matrix precipitates. J. Warren et al. RM36 89

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0.5 micron

Figure 6. Alloy 2: A) dislocation density B) refined cell structure 90 RM36 J. Warren et al.

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Tensile Testing Procedure

Pin-loaded tensile specimens per ASTM E21-92 were cut from both the stress relieved and the recrystallized TZM and Alloy-2 forgings as shown in Fig. 7. Gage extension was measured using a direct contact extensometer. To maintain a constant tensile strain-rate in the gage, the output signal from the extensometer was used to control the cross-head velocity using a PC interface. All tests were conducted with a constant gage strain-rate of 8.3x10"5/sec. After reaching approximately 1% strain, the test was switched to a constant cross-head velocity test of 0.02mm/sec (8.3x10"4in./sec). All tensile tests were conducted to rupture.

Figure 7. A typical tensile rupture.

Tensile tests were conducted in a two-zone furnace "clam-shell" type furnace purged with a constant flow of high pressure nitrogen gas for all test temperatures up to and including 1000°C. Argon gas flow was then used at the highest test temperature, 1200°C. Tension tests were conducted at 400°, 600°, 800°, 1000° and 1200°C. Prior to load application, specimens were soaked for 30 minutes at each temperature. Outputs from the extensometer and load cell were sent to a data logging PC. Engineering stress-strain curves were then determined from the recorded data. Three tensile tests (replicates) were conducted at each test temperature listed above. From the stress-strain plots the 0.002 offset yield strength was established for each alloy and associated heat treatment. The average strength is reported. J. Warren et al.RM 36

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Results and Discussion

The tensile testing results of the two compositions exposed to the stress relieving heat treatment are shown in Fig. 8 where the average offset yield strength at 0.002 (0.2%) engineering strain is plotted against the test temperature range of 400 -1200°C. As is evident in the Fig. 8, the yield strength of Alloy-2 is approximately 100MPa higher than the values measured in the TZM material over the entire test temperature range. From a qualitative standpoint, a higher yield strength in Alloy-2 is expected based on the observation of the refined dislocation substructure retained after stress relieving as shown in Fig. 5. A comparison of average ultimate tensile strength (UTS) data, Fig. 9, clearly indicates a higher strain hardening coefficient associated with Alloy-2 resulting in a greater dislocation density and corresponding flow stress for a given strain increment. Based on this observation, it is reasonable to assume that the precipitate density in Alloy-2 is higher than TZM. Examination of Fig. 8 also shows that the yielding is somewhat insensitive to the test temperature and explanations in terms of precipitation strengthening and dynamic strain aging (DSA) are considered. For example, each alloy system relies primarily on the dispersion of hard incoherent precipitates to strengthen the matrix by both stabilizing a finer substructure during the forging operation and by blocking dislocations during plastic deformation in the classical Orowan fashion. In the case of the TZM composition fabricated for this study, the precipitates consist primarily of the very large grain boundary Mo2(Ti)C type and extremely fine ZrO2 precipitates observed in the molybdenum matrix. In this analysis it is assumed that the principle strengthening of the TZM matrix is associated with the Zr oxide precipitate array. By comparison, particle strengthening in Alloy-2 is derived from the Hf, Zr, and Mo-(carbide) precipitates. Hence the observed differences in 0.2% offset tensile yield strength observed between the two alloy compositions can be analyzed based on the incremental tensile strength increase, Aa, associated with the presence (or absence) of matrix precipitates and the associated dislocation blocking mechanism. 92 RM36 J. Warren et al.

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320 300 280 260 240 220 200 180 160 140 en c CD 120 CM O 100 q ö 80 60 40 Stress-relieved 1550°C, 20 30 min., Hydrogen 0 400 600 800 1000 1200 temperature (°C) Figure 8. Yield strength of TZM and Alloy 2 after 1550°C heat treatment.

Both long-range and short-range interactions contribute to strengthening. The long range interactions arise from the internal elastic strain fields associated with the presence of the precipitates and the resultant interaction of dislocations with the distorted matrix. Short range interactions however, contribute to the majority of the strengthening seen in these alloys [3]. The tensile yield stress increment associated with short range strengthening has been calculated by Orowan and is approximately equal to: ACT ~ 4 Gb/L where G is the temperature dependent shear modulus of the Mo-matrix, b is the Burgers vector of the Mo-matrix (2.73X10"10 m) and L is the particle to particle separation distance. A quantitative analysis of the density of extremely fine precipitates present in each alloy composition was not made for this study. If however, the assumption is made that the extremely fine precipitates in each alloy strengthen the Mo-matrix in a similar fashion the higher strength associated with Alloy-2 can, in part, be attributed to the presence of the dispersion of the somewhat coarser Hf, Zr, and Mo-(carbide) precipitates observed under the SEM (see arrows in Fig. 2b). For example, the estimated particle to particle separation distance of these precipitates in Alloy-2 is about 4j.im which, based on the Orowan relationship, corresponds J. Warren et al. RM36 93

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to an incremental strength increase of about 34MPa at 400°C. Hence, Orowan strengthening, in conjunction with higher resistance to recrystallization, likely accounts for the greater offset yield strength observed with the stress-relieved Alloy-2 composition.

380

360

240

220 Stress-relieved 1550°C, 30 rnin., Hydrogen 200 400 600 800 1000 1200 temperature (°C)

Figure 9. Ultimate strength after 1550°C heat treatment 94 RM36 J. Warren et al.

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150 - -. - ... ! , 140 130 120 'S" 110 D. 100 H W 90 CD T2M __———"^ -isr~~^~*r~ - - ' ' ; c 80 i t 70 neer i

ng i 60 CD CM 50 O 40 0. 0 30 Recrystallization 20 -• 1900°C, 3h, 10 vacuum

0 , , •••• ; | 400 600 800 1000 1200 temperature (°C)

Figure 10. Yield Strength after 1900X heat treatment

The shear modulus of Mo at 400°C is about 94% of the value at 25°C while at 1200°C it reduces to approximately 83%. With the exception of the matrix shear modulus, the remaining physical parameters of the Orowan relationship are independent of temperature. Based on this analysis, if the Orowan mechanism was the only mechanism active in each alloy, a relatively small (-12%) reduction in strength would be observed between the two test temperature extremes. Counteracting the reduction in strength associated with increasing temperature is the anomalous hardening behavior observed in dispersion strengthened BCC alloys due to the interaction of dislocations with solutes (i.e. DSA) [2]. The DSA strengthening mechanism is likely active in the stress relieved compositions and clearly apparent in the recrystaiiized alloy compositions, Fig. 10, where the average 0.2% yield stress is plotted over the test temperature range. As is shown in Fig. 10, the yield stress of both alloys increases gradually above about 1000°C. Based on the classical interpretation of DSA, the drag stress imposed on mobile dislocations increases as the mobility of solutes within the lattice increases resulting in a gradual increase in strength. A number of investigations have shown that the J. Warren et al. RM 36 95^

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activation enthalpy of DSA is the same as interstitial diffusion [2]. The temperature range where DSA is active in Mo-based alloys is between about 800° and 1200°C and is readily observed in Mo-based alloys that retain a significant fraction of alloying elements in solution, i.e. quenched from a recrystallization heat treatment. DSA is also observed in recrystallized TZM that is furnace cooled which suggests that the precipitation kinetics associated with TZM are sufficiently slow enough to retain a reasonable fraction of solutes in the matrix. Hence, it is reasonable to assume that DSA is active in both recrystallized alloy compositions and can explain the increase in strength shown in Fig. 10 above 1000°C.

As expected, comparison of Fig. 10 to Fig. 8 shows that recrystallizing the stress relieved alloys reduces the yield strength of each composition dramatically. TZM exhibits approximately a 70MPa average reduction while the Alloy-2 composition drops nearly 150MPa; a difference of 80MPa. Grain growth, particle coarsening and a reduction in the dislocation density associated with hot working are the principle factors that contribute to the drop in strength. Of particular interest is the larger drop in strength observed in Alloy-2 likely a result of a significantly greater reduction in the dislocation density than that which occurs in the TZM composition.

The role of the tantalum addition and its contribution to strengthening, if any, was not investigated in this report. No precipitates containing tantalum were observed and, based on EDS analysis of the matrix, it appears to be present entirely in solution

Conclusions

In summary, the strength observed in the Alloy-2 material has been attributed to the presence of a fine distribution of Hf, Zr, and Mo-(carbide) precipitates that strengthen the material in the classical Orowan fashion. The strength of this alloy is significantly higher than TZM after thermal exposures to temperatures as high as 1550°C yet possess similar forging and machining characteristics. Modifications to this novel alloy are planned and include an analysis designed to study the effect varying the weight fraction of tantalum has on strength. 96 RM 36 J. Warren et al.

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References

1. US Patent No. 5222116; Metallic Alloy for X-ray Target (1992) 2. E. Pink, A Review of Deformation Mechanisms in Particle Strengthened Molybdenum Alloys, Refractory Metals and Hard Materials, Dec. 1986, pp. 192-99. 3. A. Lou et al, High Temperature Tensile Properties of Molybdenum and a Molybdenum-0.5% Hafnium Carbide Alloy, Scripta Metallurgica et Materialia, Vol. 29, 1993, pp. 729-732. AT0200472

M.I. Karpov et al. RM 44 9

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Nanolaminate - bulk multilayered Nb-Cu composite: Technology, Structure, Properties

M.I. Karpov*, V.I.Vnukov*, N.V. Medved*, K.G.Volkov*, I.I.Khodoss**

*Solid State Physics Institute, Chernogolovka Moscow distr., 142432 Russia **lnstitute of Microelectronics Technology, Chernogolovka Moscow distr., 142432 Russia

Summary:

The Cu-Nb multilayered composite as a ribbon 50 mm in width and 0,35 mm in thickness was produced by the following procedures constituting a technological cycle: integration of a packet of specific number of layers (16 layers 0,35 mm in thickness both for Cu and Nb), rolling of the packet in vacuum at a temperature of 750-800 °C, cold rolling in air down to the thickness equal to that of one initial layer constituting the composite. The structure of the composites was examined by optical and electron microscopy, X-ray diffraction at all the stages. The hardness was measured in the course of cold rolling for each of 3 technological cycles. Upon completion of 3 cycles a nanolaminate was produced consisting of 32768 layers of 11 nm in thickness, the ribbon thickness being 0,35 mm. A change in the hardness as a function of true deformation has an ordinary parabolic behaviour in the first and second cycles. A significant growth of hardness and the change over from the parabolic to the linear behaviour of the dependence is observable at a thickness of the layers less than 200 nm in the 3-rd cycle. The hardness amounting to 350 HB was observed at layers thickness of 11 nm. Noticeable changes of the hardness and halfwidth of X-ray peaks of Nb and Cu occur after annealing at temperatures above 400 °C. After annealing at 1000 °C copper grains of about 200 - 600 nm in size form which comprise niobium inclusions of about 1-10 nm in size. The hardness drops down to 135 HB.

Keywords:

Nb-Cu, composite, layer, technology, rolling, structure, hardness 98 RM44 M.I. Karpov et al,

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1. Introduction:

An interest in nanokrystalline materials has been growing in recent years not only in connection with an essential difference between their fundamental physicochemical properties and those of materials of the ordinary structure (1). Still greater number of researchers attempt to produce on the base of nanocrystalline materials novel materials of specific mechanical, magnetic, or electrical properties. So, the authors of ref. (2) examined a possibility of producing a high-strength multilayered Mo-W nanolaminate using the CVD technique. The material produced by them was 50 |um thick and was composed of 4 nm thick molybdenum and tungsten layers. The hardness and ultimate strength of that material were 15 times as high as the analogous characteristics of an alloy of the corresponding composition. The authors of (3-5) produced Fe-Ag and Fe-Cu nanolaminates, the layers being 20 nm in thickness, through the cycle: diffusion welding-pressing-rolling repeated 2-3 times. Their aim was to study magnetoresistance on them. They observed the magnetoresistive effect analogous to that on a CVD-produced Fe-Cu nanolaminate. The result suggests that purely metallurgic procedures are quit prospective for producing bulk nanolaminates. Their efficiency is far beyond that of the CVD technology. The aim of this work is to study a possibility of producing a metal nanolaminate by an alternative metallurgic method based on compacting by hot rolling in vacuum. This method is still more efficient, compared to (3-5). Moreover, the packet being compacted stays at an elevated temperature much shorter time period. This offers ample scope for producing nanolaminates of pure metals whose mutual solubility is quite big. We took a Cu-Nb composite for the object of our investigation as it possesses interesting electric properties. From the standpoint of materials science it is the case of complete mutual insolubility of copper and niobium.

2. Experimental technique:

The nanolaminate was produced by the following procedure, constituting a technological cycle, repeated three times: integration of a packet of 32 layers (16 both for Cu and Nb in the first cycle), hot rolling of the packet in vacuum, cold rolling in air down to the thickness equal to that of the initial layer. The initial layers measuring 50x100 mm in area, 0,35 mm in thickness were integrated to a 11,2 mm thick packet and fixed at the angles with rivets. Before the integration the layers surfaces were worked with sandpaper • I. Karpov et al. RM 44 99^

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in order to remove oxides and produce roughness. The samples were rolled in vacuum at 750-800 °C (heating time being 15 min) in two passes, the sum reduction being 40 %. After that the samples were rolled in vacuum in air at room temperature, the final thickness of the rolled packet was 0,35 mm. In the second and third cycle the packets were made up of the layers rolled in the previous cycle. The produced nanolaminates were annealed in vacuum at temperatures of 400, 600, 750, and 1000 °Cfor 1 hour. The structure of the composites was examined at all the production stages and after the anneals. The plane of the section for metallographic tests and the plane of the foils for transmission electron microscopy were taken normal to the rolling direction. The structure of the sections was examined in a metallographic microscope Neophot and scanning electron microscope Jeol JSM-25S. The foils were examined in an electron microscope Jeol JEM- 2000FX. The X-ray spectra were taken from the surface of metallographic sections prior to their etching in a diffractometr Siemens B-500 using a Cu Ka radiation in the reflection mode from an 1x1 mm of section. The obtained spectra were processed by a specific Siemens program packet and JCP file. The samples were measured for hardness in the process of cold rolling in all the cycles and after the anneals. Down to a thickens of 1 mm the measurement was performed in a Rockwell instrument on HRB scale (100 kg load, 1,588mm dia ball). With smaller thickness the Vickers hardness was measured (5 kg load). The obtained data were adjusted to the Brinnel hardness - HB. 3. Results:

The results of the structural studies of the composite after rolling are illustrated in fig.1. The quantitative data on the layers thickness and their nonuniformity at different production stages are listed in table 1. It can be seen that the layers nonuniformity in thickness grows with number of layers and with the degree of deformation. So, after the rolling in the first cycle down to a thickness of 1 mm, a mean layers thickness being 0,0311 mm, ultimate changes come to 0,02 mm on either side (64 % of a mean value). After the rolling down to 0,7 mm (second cycle) at a mean layers thickness of about 680 nm the ultimate change towards an increasing reaches 1400 nm (206 %) and 600 nm (88 %) towards its decrease. Upon completion of rolling in the third cycle, a mean layers thickness being 11 nm, the corresponding changes come to 24 nm (218 %) and 7 nm (64 %). The layers become wavy herewith. 100 RM44 M.I. Karpovet al.

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But the layers retain their laminarity even at a mean thickness of 11 nm (fig. 1c), they have fine clear boundaries and not intermix. The hardness measurement data are illustrated in fig. 2. A change in the composite hardness as a

100 x 1000

Fig.1

The structure of composite after rolling: 7 a - first cycle, thickness of the composite -1 mm; b - second cycle, thickness of the composite - 0,7 mm; c - third cycle, thickness of the composite - 0,35 mm. M.I. Karpov et al. RM44 101

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Table 1.

Layers thickness at different nanolaminate production stages. Ultimate changes are given for some values.

After Cold Cold Cold Cu-Nb Initial rolling in rolling rolling rolling vacuum down to down to down to 1 mm 0,7 mm 0,35 mm thickness thickness thickness

1 -st cycle 0,35 mm 0,21 mm 0,0311 mm 0,0218 mm 0,0109 mm + 0,02 mm -0,02 mm

2-nd cycle 10900 nm 6540 nm 971 nm 680 nm 340 nm +1400 nm - 600 nm

3-rd cycle 340 nm 204 nm 31 nm 22 nm 11 nm +24 nm -7 nm

function of true (logarithmic) deformation at rolling has a near-parabolic behavior in the first and second cycles. Late in the second cycle, however, when the layers thickness changes from 680 nm to 340 nm, a change from parabolic to the linear behavior of the dependence is observed. It becomes near-linear in the third cycle, the hardness coming up to 350 HB late in the cycle, which is characteristic of heat treated medium-carbon steel. It should also be mentioned that the hardness grows appreciably (from 55 to 135 HB) after sample rolling in vacuum in the 3-rd cycle (a mean layers thickness being 204 nm), compared to the corresponding value in the second cycle (6540 nm). 102 RM44 1.1. Karpov et al.

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Fig. 3 illustrated the data on hardness measurement as a function of the annealing

i ' i ' i 350 — 204nm-11nm 350 300 •210000nm-ipg00nm 250 §200 200 &150 150 I 100 100 50 50 Deformation, e Fig. 2

The hardness as a function of true (logarithmic) deformation - e by rolling in first, second and third cycles.

0 200 400 600 800 1000 i i i . i . i 350- •— - 350 \ 300- -300 "Pi \ I •

100- -100 0 200 400 600 800 1000 Annealing temperature, °C

Fig. 3

The hardness of the Cu-Nb nanolaminate as a function of annealing temperature. It is seen that 400 °C annealing decreases the nanolaminate hardness only insignificantly. At higher temperatures the hardness drops nearly-linear M.I. Karpovet al. RM44 103

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. • with the temperature, and after annealing at 1000 °C its value comes to 135 HRB. A change in the halfwidth of X-ray (211) Nb and (220) Cu peaks as a function of the annealing temperature is illustrated in fig. 4. The both curves are near- parallel throughout the temperature ranges. The halfwidth changes most significantly within 400-600 °C.

200 400 600 800 1000 i . i , i , i , i •1,3 —•—Nb(211) •1,2 -•- Cu (220) 1,1 •1,0 •0,9 •0,8 •0,7 0,6 0,5 •0,4 0,3 •0,2 0 200 400 600 800 1000 Annealing temperature, C

Fig. 4

The halfwidth of X-ray (211) Nb and (220) Cu peaks as a function of annealing temperature.

Fig. 5 illustrated the results of the electron-microscope study of the nanolaminate after annealing at 600 and 1000 °C. Laminarity is seen to be disturbed already after the 600 °C anneal (fig 5a, b). A greatest number of the survived layers are 50 - 60 nm in thickness but there are layers less than 2 nm thick. The boundaries between the layers are diffuse. The are also structural regions comprising newly-formed grains of 60-70 nm in size. Inside these grains finely divided inclusions of the second phase are seen, the particle size ranging from 1 to 10 nm. After the 1000 °C anneal the structure is no longer laminar (fig.5c, d). It constitutes extended copper grains measuring 200 - 600 nm across. Some grains scarcely contain second-phase inclusions, the others contain a large number of particles from 1 to 10 nm in size. 104 RM44 M.I. Karpov et al.

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Fig. 5

The structure of the nanolaminate after annealing at 600 °C (a, b) and 1000 °C (c, d). M.I. Karpovetal. RM 44 105^

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4. Discussion

The data of table 1 and figure 1 suggest that the goal sought in this work is accomplished: nanolaminate has been produced consisting 32768 layers of copper and niobium, a mean layer thickness being 11 nm. The occurrence of a layer thickness nonuniformity can, possibly, be remedied or decreased by changing the laminate production process parameters. Noteworthy is a change in the hardness upon rolling. Ordinary materials exhibit parabolic behavior of hardness as a function of deformation, when hardness growth slows down with the increasing deformation. This is e result of dynamic recovery in the process of deformation: dislocations climb to other glade planes under the action of growing internal stresses as a result of which dislocation density growth slows down and the internal stresses relax. The deviation of this curve towards the linear hardness growth with the deformation signifies the absence of the dynamic recovery: the relationship between the dislocations which passed the sample and those stuck in it remains constant. The change to this new composite quality occurs at a layer thickness of 680 - 340 nm. This finding is also supported by the fact that even before rolling in the third cycle the hardness was 135 HB that is nearly twice as high as the corresponding values in the first and second cycles. Also, the hardness of the nanolaminate drops to the same level upon annealing at 1000 °C following which the grain size also ranges from 200 to 600 nm. Quite unusual come about changes in the hardness, halfwidth of X-ray peaks of Nb and Cu and in the composite structure upon annealing. A temperature of 400 °C corresponds to the onset of recristallisation in copper. After the 600 °C anneal copper will be completely recristallized. In this temperature range Nb layers demonstrate only the occurrence of dislocation recovery, niobium starts recrystallizing at 1100 °C. The date on change in the hardness and halfwidth of the X-ray peaks upon annealing indicate that the dynamics of structure and internal-stress alterations both in Nb and Cu layers is governed by the changes occurring in Cu layers. These and the electron microscopic data, illustrated in fig. 5, will allow an understanding of how structure alterations proceed in nanolaminate upon annealing. Within RT - 400 °C all the structure alterations in nanolaminate are related to the processes of dislocation restructurization in copper and niobium layers. Within 400 -600 °C the diffusive mobility of copper atoms suffices for recrystallization to develop in copper layers. New copper grins form in the structure regions where copper layers were thick. The other process is progressing concurrently: copper atoms penetrate into niobium layers diffusing across the 106 RM44 M.I. Karpov et al.

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boundaries of old grains, dislocation boundaries, and other structural imperfections in niobium.

As a result, niobium layers "swell", their thickness grows while that of copper layers decreases. The result of this process is clearly seen in fig 5a. Niobium layers do not represent niobium any longer, they constitute a copper matrix with inclusions of Nb particles (fig. 5b). The size of niobium particles is dictated by the structural elements of niobium layers in the post-annealing state: dislocation cells, subgrains (if at all one can speak about such elements in 11 nm thick layers) or they analogs. Judging from the fact that the particles size is 1-10 nm, these structural elements were of the same sizes. The copper layers formed at their boundaries established the conditions for stress relaxation in them so that the internal stresses in niobium relax the same as in copper. On further increase in the annealing temperature due to development of recrystallization in the copper matrix there forms a structure involving copper grains relatively free of niobium inclusions and copper grains comprising a large number of niobium inclusions (fig 5c, d). The size of these particles remains constant, i. e. 1-10 nm, as they cannot coagulate in the copper matrix because niobium atoms are incapable of diffusion in copper once it is defectless. It should be mentioned that even after annealing at 1000 °C the grain size in the composite is 200-600 nm, and it is still a nanometrical material.

5. Conclusion

Nb-Cu nanolaminate has been produced as a 0,35 mm thick ribbon. The nanolaminate consists of 32768 layers of a mean thickness 11 nm. Hardening upon rolling the composite becomes anomalously strong at layers thickness less than 340 nm due to the absence of dynamic recovery in the process of deformation. The restructurization of nanolaminate upon annealing is associated with the occurrence of two basic processes: recrystallization in the copper layers and diffusive penetration of copper into the niobium layers. As a result there forms a structure, consisting of copper grains relatively free of niobium inclusions and those comprising a significant number of niobium particles measuring 1 - 10 nm . M.I. Karpov et al. RM 44 107

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The authors are grateful to G. E. Abrosimova, senjor scientist for the ISSP RAS for conducting X-ray diffraction analysis of the samples and to A. S. Aronin, leading scientist for the ISSP RAS for his participation in discussing the results of the electron microscopic studies.

This work is supported by Russian Foundation for Basic Research Project 9902-17411

References:

1. H. Gleiter "Nanokrystalline Materials", Progr. Mater. Sei. 33 (1989) 223. 2. D. P. Adams, M. Vill, J. Bilello and S. M. Yalisove "Controlling strength and toughness of multilayer films: A new multiscalar approach", J. Appl. Phys., 1993, v. 17, N 2, p. 1015-1021 3. P. H. Shingu, K. Yasuna, K.N. Ishihara, A. Otsuki and M. Tarauchi, Mater. Sei. Forum, p. 235-238, (1996), 35 4. K. Yasuna, M. Tarauchi, A. Otsuki, K.N. Ishihara and P.H. Shingu, J. of Appl. Phys. 82 (1997), p.2435 5. B. Huang, K.N. Ishihara, P.H. Shingu "Bulk nano-scale Fe/Cu multilayers produced by repeated pressing-rolling and their magnetoresistance", J. Of Mat. Sei. Let., 19,2000,1763-1765. AT0200473 tO8 RM76 P. Harmat et al.

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A Novel Method for Shape Analysis: Deformation of Bubbles During Wire Drawing in Doped Tungsten

P. Harmat1, T. Grösz2, L. Rosta3, L. Bartha1

1MFA Res. Inst. for Techn. Phys. and Mat. Sciences, POB 49 Budapest H-1525 Hungary institute of Chemistry, Chemical Research Center, POB 17 Budapest H-1525 Hungary 3KFKI SZFKI Res. Inst. for Solid State Phys. and Optics, Budapest H-1121 Hungary

Summary:

A novel technique has been developed for monitoring shape and size of microscopic pores, bubbles, second phase particles in deformed PM materials. The anisotropic small angle neutron scattering (ASANS) measurement provides direct visualisation of the shape of second phase objects after rolling, swaging, wire drawing. Also in case of mixture of different objects e. g. uniformly elongated bubbles and spherical ones they can be separated and their morphological parameters like relative number density, diameter, aspect ratio can be obtained from the quantitative analysis of ASANS data. Rods and wires from K-AI-Si doped tungsten containing residual porosity and K filled bubbles were studied from 6 mm to 0.2 mm in diameter. The increase of the average aspect ratio (~1/d) was found to be much slower than expected from the usual theory (~1/d3). Instead of "constant volume" assumption, the "constant length" seems to be reliable. The ASANS investigation revealed also the occurrence of a small amount of spherical bubbles after several steps of wire drawing.

Keywords: doped tungsten, bubbles, deformation, wire drawing, small angle scattering

1. Introduction

Single phase materials subjected to sintering contains powder particles or agglomerates of the solid material (metal, oxide) surrounded by gas or vacuum. The sintered body after sintering is usually close to the theoretical density and the residual porosity consists of gas filled bubbles or empty P. Harmat et al. RM 76 109

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pores. From the point of view of small angle neutron scattering (SANS), both the initial single phase solid material (powder) and the sintered body should be regarded as two-phase systems having sharp interfaces between the solid material and vacuum [1].

2. The method

SANS method can also be applied to determine the shape or size of pores during several steps of thermomechanical processing after sintering, like rolling, swaging and wire drawing. In case of wire drawing the definitely oriented macroscopic strain being parallel to the wire axis results in a uniform orientation of deformed pores in the body. In consequence of this, a significant anisotropy in the two dimensional scattering distribution of SANS intensity can be observed.

Fig. 1 Schematic plot of SANS intensity arising from a large set of equally elongated ellipsoids with identical orientation. The aspect ratio of ellipsis contour corresponds to the average aspect ratio of elongated ellipsoids.

The application of anisotropic SANS (ASANS) is relevant in investigation of the elongation of pores and bubbles. The scattering distribution on the detector plane can be derived from the basic theory of small angle scattering [2]. We have concluded that in case of an ellipsoid shaped particle the isointensity contours of the map representation of the 2D scattering intensity 110 RM76 P. Harmat et al.

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are ellipsis and they have the same aspect ratio as the particle itself. In our wire samples, the bubbles have always elongation parallel to wire axis. Due to identical orientation of bubbles the superposition of scattering intensities of individual bubbles reserves this anisotropy. Applying neutron beam perpendicular to the wire axis, the whole ensemble of elongated bubbles reveals this anisotropy and the aspect ratio of ellipsis contour corresponds to the average aspect ratio of elongated ellipsoids.

3. Material and experiments

We investigated K-AI-Si doped tungsten alloy produced for incandescent filament in conventional light sources produced by sintering. Only small (50- 500 nm) K filled bubbles and larger (1000-5000 nm) empty pores remain in the bulk after sintering.

Theory of Moon and Koo (1971): 1) Deformation during wire drawing: 2) Morphological changes after wire drawing due to annealing: - identical deformation rate for wire and bubbles Aspect Ratio A=ai/ag A =A *(d /d )3 bubble 0 ' 0 wire - neither "breaking-up-into-row-of bubbles" norspheroidization occurs A<8.9

doped tunsten direction wire of wire 6 drawing A >= 8.9 o \J A » 8.9 bubbles bubbles and pores t t t t f t and pores 0 after drawing o o o o o o Fig. 2 Summary of the usual theory made by Moon and Koo [4].

The subsequent steps of mechanical working at 900-1300 K result in more and more elongated bubbles [3]. It is assumed by Moon and Koo [4] that during the mechanical working, the amount of deformation of microscopic P. Harmat et al. RM76 111

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bubbles and pores is identical to that of the rod/wire and these elongated ellipsoid shaped bubbles return to their spherical shape or break up into smaller spherical pieces by spheroidization during annealing. The statistical study of SEM and TEM micrographs led to the conclusion that the average number of bubbles in a single row after heat treatment is essentially smaller than expected [5].

0.4 mm i n A ~*~~ * annealed 04mmAlilt1650K

'3S E .2 0.6 mm

ns o 2 a 1 mm '5

to _c 'to 3 mm O o o

6mm 6 mm annealed

Heat treatment

Fig. 3 Processing map of ASANS intensity maps measured on samples of doped tungsten from selected stages of thermomechanical processing. One step to right corresponds to an annealing, while one step upwards corresponds to a uniaxial mechanical working. The measurement was made in LLB CEA Saclay, France using neutrons with wave length of 12 A and detector sample distance of 4.8 m. 112 RM76 P. Harmat et al.

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The whole process of fabrication between 6 mm down to 0.4 mm in rod/wire diameter has been monitored on selected samples. At this smallest size, extra samples (partly annealed and fully annealed) was prepared for comparison. Each individual two dimensional scattering distribution maps are arranged into a processing map.

The first sample of 6 mm diameter rod underwent swaging process, its ASANS map shows a slight anisotropy. The second one with the same diameter was annealed, it shows a very small anisotropy. The further samples were subjected to the subsequent cumulative deformation process by swaging and wire drawing, they show an increasing anisotropy which corresponds to more and more elongated ellipsoids. The 0.4 mm diameter sample was then annealed at 1650 K for 10 s, its 2D map shows a large anisotropy decrease. The anisotropy vanishes completely after full annealing of the sample at 2700 K for 300 s.

4. Quantitative analysis of ASANS intensities

Spheres with minimal elongation: Ellipsoid parameters from fitting parameters from fitting Wirediam= 6 3 2 1 0.6 0.4 mm Wire diam = mm

./ \ 10000 - 10000. \ Volume •

=*1000,

2 : 22 s> a^^ ^\ Length g 100, ^" • - Diameter • para i

3> 10. i 2 i Aspect ratio ; A 5 ^' A A A-—-

ULI /^/ ! 0.1- •—-• Rel. number cone.

0.1 Reciprocal wire diameter [mm' ] Reciprocal wire diameter [mm" 4a 4b P. Harm at et al. RM 76 113

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Fig 4a and 4b Results of model parametric fitting ofASANS data on Fig. 3. Parameters of elongated ellipsoids are plotted on Fig 4a and that of nearly spheres on Fig 4b. The length and diameter data from fitting are absolute values in nm. The dimensionless aspect ratio increases from 1 to 14. Average bubbles volumes are calculated from the length and diameter data and plotted in arbitrary units.

The analysis of scattering maps reveals further details. The circular shape in the central region at samples with 1 mm in diameter and below suggested us that among the majority of very elongated ellipsoids, spheres are also present in the sample and the scattering distribution corresponds to their overall superposition. The reduced anisotropy of 0.4 mm sample after 1650 K annealing shows that the number density of ellipsoids decreases because many ellipsoids undergo breaking up during the incomplete annealing. The perfect vanishing of the anisotropy on 0.4 mm sample after 2700 K annealing reveals that all the ellipsoids disappeared by the break-up process. 114 RM76 P. Harmat et al.

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0.4 mm crowded 0.4 mm «1650K o.4mm ^si-

0.6 mm

o 0) m CO O) '35 3 mm (A V U o re o 6 mm 6 mm annealed re o 0> (0)

Heat treatment

Fig. 5 Simulated processing map of ASANS intensity maps. Results from parametric fitting on Fig 4 were used to simulate ASANS intensities for each calculated 2D maps.

A simple quantitative model has been developed to describe the observed behavior. We supposed that there are two populations of bubbles in the wire: the elongated ellipsoids with uniform aspect ratio and the sphere-like ellipsoids with an aspect ratio slightly greater than or equal to 1. The shape of overall superposition of scattering intensity depends on the relative number densities and sizes of these two populations. We used G. Pepy's PXY program (CEA Saclay, France, [6]). Figure 4 shows the result of parametric fitting to the elongated ellipsoids and the spheres with moderate elongation. P. Harm at et al. RM76 115

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Data from this fitting can be used for simulation of scattering intensities. Fig 5 shows such a totally simulated processing map having a good similarity with that on Fig 3.

5. Discussion

Wire diameter [mm] 3 2 1 0.6 0.4

0.1 1 Reciprocal of wire diameter [mm1]

Fig. 6 Theoretical and experimental dependence of the ellipsoids' aspect ratio on the reciprocal of wire/rod diameter. The slope of fitting to experimental data is essentially different from the theoretically expected value. Also the aspect ratio of the population of almost spherical bubbles is plotted showing a tendency of increase.

The aspect ratio values are plotted in function of reciprocal sample diameter. The almost linear dependence (A ~ 1/d) is drastically different from the value calculated from the theoretical expectation (A~1/d3). Even the moderate enhancement of the aspect ratio of almost spherical bubbles can be evaluated. They seem to be produced as perfect spheres during wire drawing 116 RM76 P. Harmat et al.

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but they become also more and more elongated during the subsequent processing.

Theory of Moon and Koo (1971): From SANS experiments (2000): -3 -1 A~d A-d

VOLUME of bubbles remains constant LENGTH of bubbles remains constant Identical elongation rate for wire and bubbles Volume of bubbles decreases (constant ratio of wire & bubble diameter) Fast elongation rate Slower elongation rate

Fig. 7 Schematic comparison of results from our SANS measurement with the usual theory.

The huge decrease of the average volume of the elongated ellipsoids indicates that one of the assumptions in the theoretical description for the deformation is missing: the constant volume of bubbles during wire drawing is not valid. A constant average length of the elongated bubbles seems to be a better assumption. P. Harmat et al. RM 76 117

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6. Acknowledgement The work was supported by the National Scientific Research Fund (OTKA), contract No.T 32730.

7. References

[1] F. Rustichelli: Material. Science and Technology 11 (1993) pp. 118-141 [2] G. Porod: in: Small Angle X-ray Scattering, Eds. O. Glatter, O. Kratky, Academic Press, London, 1982 [3] The Metallurgy of Doped/Non-Sag Tungsten, Eds Pink and Bartha, Elsevier Sei. Publ., London (1989) [4] D. M. Moon and R. C. Koo: Metall. Trans. 2 (1971) pp. 2115-2122 [5] O. Horacsek, M. Menyhärd and J. Läbär: High Temp. Materials and Processes 14 (1995) pp. 207-213 [6] G. Pepy: PXY fitting program for SANS data. Private communication (1999) JJ8 RM89 H.W. King et al.

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RESIDUAL STRESS IN TIN COATINGS ON Ti-6AI-4V

H.W. King, S. Gursan I. Kim D.R. Nagy S. Ferguson, M. Yildiz School of Advanced Materials Liburdi Engineering Ltd. Mechanical Engineering Kum-Oh National Univ. of Tech. 400 Highway No. 6 University of Victoria Kumi, Kyung Bug, 730-701 Dundas, ON, L9H 7K4 Victoria, BC V8W 3P6 KOREA CANADA CANADA

ABSTRACT

Residual stresses in TiN coatings deposited by a reactive ion process on Ti-6A1-4V sheet were determined over the temperature range from 20-700 °C, using the sin2\]/ X-ray diffraction method. Calculations based on the relative thermal expansion coefficients of the coating and substrate showed that the magnitude of the thermal stress at room temperature was only -0.04 GPa, so that the observed compressive stress of the order of -2 GPa could be attributed almost entirely to microstresses generated during the deposition of the TiN coating. The observed stress remained approximately constant after heating for 1 h periods at temperatures up to 500 °C and then reduced irreversibly to -0.6 GPa over the temperature range from 500 to 700 °C. The reduction in compressive stress was accompanied by a decrease in peak breadth and an increase the perfection of the [111] texture of the TiN coating, confirming that the stress relief was associated with recovery processes involving point defects and dislocations. A simultaneous decrease in the d-spacings of planes parallel to the surface of the coating over the temperature range from 500 - 700°C is attributed to the relief of the 833 Poisson generated strain, rather than a loss of nitrogen from the TiN structure.

INTRODUCTION

The effectiveness of TiN coatings for improving the erosion and corrosion resistance of compressor blades, and the wear resistance of cutting tools, is strongly dependent on the magnitude and nature of residual elastic stresses generated within the coating. While a moderate compressive stress is beneficial for increasing coating adhesion, a tensile stress can cause coating failure by crack formation and loss of adhesion (1-4). Two types of residual stress are observed in deposited coatings:

1. Macroscopic thermal stresses that originate from the thermal expansion coefficients of the coating and substrate. These stresses, which are uniform within the grains of the coating, can be either tensile or compressive in nature and are fully reversible on thermal cycling. H.W. King et al, RM 89 119

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2. Microscopic growth stresses originate from defects in the coating caused by the impact of randomly arriving particles and hence are not uniform throughout the coating. These stresses are usually compressive in nature, which is favourable for coating integrity, but they can be partially relieved, if not completely eliminated, by heating to a sufficiently high temperature. The temperature dependence of residual stresses can be determined by high temperature X-ray diffraction, while room temperature stresses can also be monitored after thermal treatments, to provide a reference for the evaluation of non-reversible effects. The incidence and relief of the non-uniform growth microstresses in coatings can be determined from measurements of the breadth of X-ray diffraction profiles (4,5), while the intensity ratios of diffraction peaks can provide information on textured microstructures (5,6). As the thermal component of residual stress is expected to be insignificant in TiN coatings deposited on Ti-6A1-4V, due to the close similarity between the thermal expansion coefficients of the coating and substrate (2,7), this system provides a unique opportunity for investigating the relief of the growth stress in PVD coatings as a function of controlled thermal treatments.

EXPERIMENTAL

The sin2\]/ X-ray diffraction method for determining residual stresses is based on measurements of elastic strain in terms of the variation of the interplanar spacings, d^, of a selected diffraction plane inclined at an angle i|/ with respect to the sample surface. Using the interplanar spacings, dn, of planes that lie parallel to the sample surface as a reference, the strains sv in planes inclined to the surface can be expressed in terms of Ad/dn, where Ad = dn-dv. The biaxial residual stress, cs^, in a plane parallel to 2 the surface of a coating can then be determined from linear slopes fitted to plots of Ad/dn versus sin i|/, in accordance with the relationship (2,8):

2 sv = Ad/dn = V2S2 a

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A 10 |am TiN coating was deposited on to both sides of a 1 mm thick Ti-6A1-4V sheet at a temperature of 475 °C, using the Liburdi Engineering reactive ion coating method, which is described elsewhere (12,13). X-ray specimens 6 mm wide and —30 mm long were cut from the coated sheet with a slow cut diamond saw. High temperature X-ray measurements were made in air using a Buehler HDK 2.3 high temperature furnace mounted on a Scintag XDS 2000 X-ray diffractometer, operated in the 8:6 configuration. The coated specimens were clamped firmly to the current electrodes of the furnace by a pair of stainless steel clips, and were heated by the passage of a dc current through the Ti-6A1-4V substrate, as described in a previous publication (14). The temperature of the coating was measured with a Pt/Pt-10%Rh (type S) thermocouple welded to the under surface of the specimen and was controlled to ± 1 °C using a Micristar PID controller. The X-ray source was operated at 45 kV and 40 mA, using a Cu target X-ray tube. Since resolution was not an important factor for recording the broadened diffraction peaks, the intensity of the diffracted beam was maximized by using a 3 mm receiving slit with no Soller slits. The selected diffraction peak was step scanned over an angular range of at least 10 °28, using a step width of 0.3° and a dwell time of 20 s, which gave a minimum of 30 data points per diffraction profile. The peak positions of the broadened profiles were determined with a profile fitting program based on a Pearson VII function. To investigate the temperature dependence of residual stress, the peaks were first scanned at room temperature (20 °C), using tilt angles of 0-40° in increments of 10°, and these scans were then repeated at temperatures of 100, 200, 300, 400, 500, 600 and 700 °C. Allowing 10 minutes to equilibrate at each temperature, and 10 minutes to perform the scan at each \.\J angle, the sample was held for approximately one hour at each temperature. The step scan data obtained at each temperature were also used to determine the breadth of the selected diffraction profile in terms of the full width at half maximum (FWHM), to investigate the relief of microstresses. After each high temperature scan, the sample was returned to room temperature and re-scanned over the range of tilt angles, \\i, to detect any permanent changes in residual stress or changes in the dn-spacings of planes parallel to the surface, i.e. for \\) = 0. The relative intensities of the low angle 111 and 200 TiN peaks were also measured after each thermal treatment, to monitor the state of the texture of the coating.

RESULTS AND DISCUSSION

The low angle section of the room temperature diffraction pattern of the TiN/Ti-6Al-4V sample is shown in Figure 1. The intensity ratio (I111/I200) of the 111 and 200 diffraction peaks was found to be 12.6, which is significantly greater than the intensity ratio of 0.72 observed for randomly oriented powder samples of TiN (15), indicating that the present PVD coating exhibits a strong [111] texture. The high angle section of the diffraction pattern in the insert in Figure 1 shows that the highest angle 511 peak at 141.3 °29 cannot be used for determining residual stress, because it is overlapped by a diffraction peak from the underlying Ti-6A1-4V substrate. The next highest angle peak, i.e. the 422 peak at 125.7 °26 was thus selected for the sin2i|/ determination of residual stress, because it is not overlapped by any of the peaks of the substrate pattern. H.W. King et al. RM89 121

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1520_ 136DJ 00 laooj 111 Q_ 1O4D_ o 1152.198 CPS h SIT Ti-6-4 || I f\ 560 \ !L i 124.1) 127.0 130.0 133.0 136.0 139.0 142.0 400_ 240_ 200 91.419 CPS 36.IVD 38.0 40.0 42.0 44.0 46.0 Deg. Diffraction Angle, 26

Figure I. Diffraction pattern of the TiN coating on Ti-6A1-4V.

0.1 0.2 0.3 0.4 sin \|/

2 Figure 2. Temperature dependence of the strain Ad/dn versus sin v|/.

The strain Ad/dn determined at various tilt angles \\i is plotted against sin^y in Figure 2, for the temperatures between room temperature and 700 °C. There is significant overlap between the plots at temperatures up to 500 °C, but the compressive strains at 600 and 700 °C are distinctly smaller, in contrast to previous observations on TiN coatings on 316 stainless steel where the sign of the measured strain changed from negative to positive at the deposition temperature (16,17). The strain plots in Figure 2 exhibit the "serpent's tail" effect, which is suggested to be the result of anisotropic X-ray elastic constants (18). The consequent error in the determination of residual stress by fitting linear slopes to the observed strain vs. sin2\]/ plots in Figure 2 was estimated to be ± 0.2 GPa. 122 RM89 H.W. King et al.

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100 200 300 400 500 600 700 800 Temperature (Celsius)

Figure 3. Temperature dependence of residual stress (o^,) and peak breadth (FWHM) in the TiN coating. 2 The residual stresses calculated from linear slopes fitted to the plots of Ad/dn versus sin v|' plots are shown as functions of the experimental temperatures in Figure 3. The dashed line at the centre of the figure represents the thermal stress ax in the TiN coating calculated using the relationship (1):

T O"T = -VTiN)] («TiN " O-Ti-6-4) ( D - T) equation (2) where the Poisson's ratio VXJN of TiN is taken as 0.2 (10), axiN is thermal expansion coefficient of TiN = 8.5 x 10"6 K"1 (2), axi-6-4 is the thermal expansion of the substrate = 8.6 x 10"6 K"1 (7) and Tp is the deposition temperature. Due to the close values of the expansion coefficients, the room temperature thermal stress is very small (-0.04 Gpa) which means that, in the absence of significant compressive growth stresses, the TiN coating on Ti-6A1-4V would be mechanically unstable. Further, since all values of the thermal residual stress are very much smaller than the experimental error of ±0.2 GPa for the sin-\\i determinations of the total residual stress, the observed trends in residual stress can thus be attributed to the growth stresses in the TiN coating. The magnitude of the total room temperature stress, at -2 MPa, is much smaller than the values of 4-8 GPa reported for TiN coatings on austenitic stainless steel (1,2,16-18). One reason for this difference is that while the stress observed in the present system is composed almost entirely of the growth stress, the stress observed in the TiN/stainless steel system has a significant thermal component of -3 GPa (17). The results in Figure 3 show that the growth stress remains approximately constant at -2 GPa on heating for one hour at temperatures up to 500 °C and then decreases sharply to -0.6 GPa over the temperature range from 500-700 °C. Room temperature H.W. King et al. RM89 123

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scans after thermal treatments showed that the decrease in stress observed on heating above 500 °C was not reversible on cooling to room temperature. The breadths of the 422 diffraction peak at different treatment temperatures, expressed in terms of the full width at half the maximum intensity (FWHM) of profiles recorded at \\i = 0, are also included in Figure 3. These results show that the FWHM remained approximately constant at just below 3.0 at temperatures up to 500 °C and then decreased to just less than 2.0 when the temperature was raised from 500 to 700 °C. The reduction in peak breadth is specifically associated with the recovery of microstresses due to the removal of point defects and dislocations (1-3,5). It is also significant to note that the temperature at which this process is initiated coincides precisely with the temperature at which the residual stress is reduced, confirming that both processes are associated with the relief of the growth microstresses in the TiN coating.

200 400 Temperature (Celsius)

Figure 4. Changes in texture coefficient and dn-spacings after thermal treatments.

The texture in the TiN coating was estimated from the relative intensities of the low angle 111 and 200 diffraction peaks using the coefficient (Im/Pm) I (l/2( Ini/I°lll + I20o/I°200)) suggested by Rickerby et al. (5), where the 1° values refer to the respective peak intensities in the ICDD powder diffraction file for a randomly oriented powder sample of TiN (15). As shown in Figure 4, the texture coefficient determined at room temperature after each thermal treatment fluctuates about 1.875 after treatments at temperatures up to 600 °C and then showed a distinct increase up to 1.935 when the treatment temperature was raised to 700°C. This observation is consistent with the observed relief of growth microstresses by the removal of point defects and dislocations, which would also cause an increase in the crystalline perfection of the texture. Since the [111] texture in TiN coatings is columnar in nature (6), a reduction of the dislocation density in the grain boundaries by thermally stimulated dislocation climb will also contribute to an enhanced texture.

The results presented in Figure 4 show that there is no significant change in the dn-spacings after thermal treatments up to 400 °C, but a distinct decrease in dn is observed after the higher temperature treatments at 500 to 700°C. Decreases in dn observed after heating at temperatures of 800-900 °C have been associated with a loss of nitrogen from the TiN structure (4,17), but this effect is not considered to 224 RM89 H.W. King et al.

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be the cause of the reduction in dn observed at lower temperatures. Since the decrease in dn occurs over the same temperature range as the recovery of the compressive growth stress, as shown by the results in Figure 2, it is thus attributed to an associated relief of the tensile £33 stress generated normal to the coating surface by the Poisson effect, as indicated by the second term in equation 1. The possibility of nitrogen loss from TiN coatings on Ti-6A1-4V substrates at higher treatment temperatures will be investigated in a further extension of this work.

SUMMARY OF OBSERVATIONS AND CONCLUSIONS

1. The TiN films on Ti-6A1-4V have a pronounced [111] texture, which is known to encourage columnar growth. 2. The room temperature thermal stress in TiN coatings on Ti-6A1-4V (at -0.04 GPa) is an insignificant component of the total compressive stress of-2 GPa, so that all observed trends in total stress can be attributed to changes in growth stress. 3. On heating for periods of lh between 500 and 700 °C the growth stress is irreversibly reduced to a lower level of-0.06 GPa, due to the removal of point defects and dislocations generated during the deposition of the coating. As this stress is little more than 10 x the magnitude of the detrimental tensile thermal stress generated at temperatures above the deposition temperature, it is concluded that TiN coatings on the Ti-6A1-4V substrate should not be used at temperatures above 500 °C for extended periods of time. 4. The [111] columnar texture of the coating is enhanced after the thermal treatments at 600-700 °C, in association with the recovery of microstresses by the removal of dislocations. This observation confirms that there is no breakdown of the TiN crystal structure at temperatures up to 700 °C.

5. A reduction in the dn-spacings of planes parallel to the coating surface after thermal treatments above 500 °C is attributed to the relief of the Poisson generated S33 normal strain, rather than a loss of nitrogen from the TiN structure.

ACKNOWLEDGEMENTS

This work was supported in part by an individual operating grant from the Natural Sciences and Engineering Research Council of Canada.

REFERENCES

1. H. Oettel and R. Wiedemann, Surf. Coat. Techno!., 76/77, (1995), 265. 2. D.S. Rickerby, SJ. Bull, A.M. Jones, F.L. Cullen and B.A. Bellamy, Surf. Coat. Technol., 39/40, (1989), 397. H.W. King et al. RM 89 125_

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3. A.J. Perry, M. Jagner, P.F. Woerner, W.D. Sproul and P.J. Rudnik, Surf. Coat. Techno!., 43/44, (1990), 234. 4. K. Xu and J. He, Surf. Coat. Technol, 70, (1994), 115. 5. D.S. Rickerby, A.M. Jones and B.A. Bellamy, Surf. Coat. Techno!., 37, (1989), 111. 6. A.S. Korhonen, Vacuum, 45 (10/11), (1994), 1031. 7. American Society for Metals, Metals Handbook -10'1' edition, (1990) Vol. 2, p.620. 8. H. Fujiwara, T. Abe and K. Tanaka, Eds., "Residual Stresses - III Science and Technology", Elsevier Applied Science, (1992), p.12. 9. A.J. Perry, Thin Solid Films, 193/194, (1990), 463. 10. A.J. Perry, J. Vac. Sei. Techno!., A8 (3), (1990), 1351-1358. H.A. Reuss, Z ang. Math. Mech., 9, (1929), 49. 12. V.R. Parameswaran, J.P. Imarigeon andD. Nagy, Surf. Coat. Technol., 52, (1992), 251. 13. D.R. Nagy, V.R. Parameswaran, J.D. McLeod and J.P. Imarigeon, in Proc. Propulsion and Energetics Panel (PEP) Symposium, Rotterdam, Netherlands, April, 1994, p.27-1. 14. H.W. King, J.D. Brown, T.A. Caughlin and D.R. Nagy, Adv. X-RayAnal, 40, CD-ROM International Center for Diffraction Data (ICDD), (1997/98), p.509. 15. W. Wong-Ng, H. McMurdie, B. Paretzkin, C. Hubbard, A. Dragoo, NBS, ICDD PDF crystal data for TIN, Gaitherburg, MD, USA 16. H.W. King, T.A. Caughlin and D.R. Nagy, Proc. 13th. International Plansee Seminar, Vol. 3, G. Kneringer, P. Rodhammer and P. Wilhartitz, Eds., (1997), p.270. 17. H.W. King, T.A. Caughlin and D.R. Nagy, J. Adv. Materials, 33(1), (2001), 63. 18. L. Chollet and A.J. Perry, Thin Solid Films, 123, (1985), 223. 126 RM98 EP. Schalunov et al.

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Anwendung der hocheffizienten dispersionsgehärteten Werkstoffe auf Pulverkupferbasis in den Teilen von Motoren und Kraftanlagen der Transportmittel

E.P. Schalunov, V.A. Dovydenkov*, V.S. Simonov*

Wissenschaftlich-technische Firma TECHMA G.m.b.H., P.O.Box 27, 428015 Tscheboksary, Russland * Werk der Komposit- und Metallkeramik-Werkstoffe KERMET AG, 424003 Joschkar-Ola, Russland

Summary: Three types (CO/94: Cu-AI-C-O; C3/04: Cu-AI-Ti-C-0 and CO/97: Cu-AI-C-O) of Oxide and Carbide Dispersion Strengthened Materials based on copper powder (QCDS-Copper) of DISCOM® logo accordingly for guiding sleeves and valve seats for modern petrol and diesel engines and also for current- removing plates of pantographs of high-speed electric locomotives, developed by Scientific and Technological Company TECHMA Ltd. and produced by mechanical alloying in attritors by Composite and Ceramic Materials Works KERMET JSC are considered. It is significant that materials CO/94 and CO/97 with the same system and the same technology of producing but with different compounds are situated on opposite poles of gamma of materials with Cu-AI-C-O system and they indicate the range of physical and mechanical properties of the gamma. For example, materials' hardness changes from 92 to 15% of values of pure copper's analogous characteristics. Above-mentioned and other types of OCDS-Copper have their temperature of recrystallization from 830 to 1000°C, great high-temperature strength and excellent wear-resistance coefficient under conditions of sliding electric contact which provides products of them working under extreme conditions with the resource substantially higher than the resource of products from traditional materials. All three mentioned types of OCDS-Copper and also more than 10 other types have been producing and delivering in the form of semi-finished OCDS- Copper items (round bars, plates, pipes etc.) and ready-made articles both Russian and foreign customers for more than 10 years.

Keywords:

Dispersion Strengthened Copper, Oxide and Carbide Dispersion Strengthened Copper, Mechanical Alloying, Guiding Sleeve of Engine, Valve Seat of Engine, Current-Removing Plate of Pantograph. E.P. Schalunov et al. RM 98 127

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1. Einleitung:

Die Entwicklung der Kommunikationssysteme auf dem neuen Niveau ist eine der Hauptbedingungen für die Sicherung des stabilen Funktionierens und des Wachstums von der Wirtschaft der Länder im 21. Jahrhundert. Der wichtigste Bestandteil der Kommunikationssysteme ist und wird das Verkehressystem sein. Die Grundforderungen an dieses System im 21. Jahrhundert werden noch strenger, weil von diesem System noch höhere Geschwindigkeiten des Fahrgast- und Güterumschlages, noch höhere Zuverlässigkeit und Sicherheit der Benutzung von Verkehrsmitteln, noch höhere Wirtschaftlichkeit der Transporte bei der gleichzeitigen Senkung der negativen Folgen des Verkehrsmittelbetriebes auf die Umwelt verlangt wird. Diese Forderungen bedingen die Notwendigkeit der Schaffung auch der Verkehrsmittel der neuen Generation selbst. Schon jetzt gibt es ziemlich große Zahl der Muster von Kraftfahrzeugen, Eisenbahnzügen, Schiffen und Flugapparaten, die den obenangegebenen Forderungen in gewissem Grade gerecht werden. Es sei aber sicher zu betonen, daß Oberflächen- und Wasserverkehrsmittel schneller geworden sind, und Flugmaschienen zudem größere Steighöhe haben. Wachsende taktisch-technische Daten der Verkehrsmittel bedürfen der Anwendung in ihnen der vor allem dynamischeren Motoren und Antriebe. Die Steigerung der Leistungs- und Geschwindigkeitsdaten der Motoren und Antriebe wird unvermeidlich mit der Erhöhung der Betriebstemperaturen in ihren Arbeitsorganen begleitet. Daher tritt in den Vordergrund die Aufgabe der Suche oder Entwicklung der Werkstoffe, die fähig sind, die Betriebstemperaturen im Laufe von ganzer Nutzungsdauer der Motoren und Antriebe auszuhalten. Es handelt sich also vor allem um hitzebeständige und warmfeste Werkstoffe. Diese Werkstoffe müssen auch verschleißfest und korrosionsbeständig unter Bedingungen der hohen Temperaturen und aggressiven Medien sein. Die Werkstoffe für elektrische Antriebe müssen auch verschleißfest unter Bedingungen des elektrischen Gleitkontaktes und des Lichtbogenbrennens. Für Wärmeabfuhr und Vorbeugung der Überhitzung der Motoren und Antriebe müssen die Werkstoffe für sie über gute Wärmeleitfähigkeit, und folglich, über gute elektrische Leitfähigkeit (was besonders wichtig für elektrische Antriebe ist) verfügen. 128 RM 98 E.P. Schalunov et al.

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Neben dieser sowie anderen technischen Sonderforderungen müssen die Werkstoffe für Motoren und Antriebe, unter Berücksichtigung deren Serien- produktion im großen Umfang, gute Fertigungsgerechtheit und annehmbare Kosten haben. Unter Berücksichtigung der immer höheren Anforderungen an den Umweltschutz dürfen diese Materialien in ihrer Zusammensetzung chemische Elemente und Verbindungen, deren Anwendung aus ökologischen Gründen verboten ist, nicht enthalten. Bereits bei der flüchtigen Gegenüberstellung von erforderlichen Eigens- chaften der Werkstoffe stellt sich ihre offensichtliche Widersprüchlichkeit heraus. Beispielsweise haben die Werkstoffe mit hoher elektrischer- und Wärmeleitfähigkeit, in der Regel, nicht hohe Festigkeitswerte, besonders bei erhöhten Temperaturen, und ungenügende Verschleißfestigkeit, besondes unter Bedingungen des elektrischen Gleitkontaktes. Wie bekannt, gibt es in der Natur kein solches chemisches Element, das alle obenangeführten Eigenschaften hätte. In meisten Fällen ist es unmöglich, alle erforderlichen Eigenschaften in einem Material, das durch Verfahren der herkömmlichen Metallurgie gewonnen wird, zu vereinen. Zugleich ermöglichen Errungenschaften und Erfolge letzter Jahrzehnte auf dem Gebiet der Festkörperphysik und der Metallkunde ganz bewußt innerhalb bestimmter Schranken in einem Material verschiedene widersprüchliche Eigenschaften zu synthetisieren. Besonders erfolgreich wird dieses Problem mit Verwendung der Pulvermetallurgieverfahren gelöst. Pulvermetallurgische Verfahren ermögli- chen es, in einem Material unterschiedliche chemische Elemente und Verbin- dungen zu verwenden, wobei jedes von ihnen bestimmte Gesamtheit der Eigenschaten hat und das synthetisierende Material additiv beeinflußt. Dabei kann man Ausgangs-komponenten verwenden, indem man ihren Einfluß aufeinander, d.h. chemische Wechselwirkung zwischen ihnen ausschließt. Als solche Komponenten können Metalle auftreten, die sich miteinander weder in festem noch in flüssigem Zustand vermischen sowie Metalle und Nichtmetalle, die über begrenzte gegenseitige Löslichkeit verfügen (Karbide, Oxide, Nitride, Boride und andere Verbindungen). Für Gewinnung wärmeleitender Werkstoffe, die bei erhöhten Temperaturen arbeiten, wurden in der Regel schwerschmelzbare Materialien (W, Ni, Ti, Cr u.a.) oder deren Verbindungen (Karbide, Oxide u.a.) verwendet, die porige Grundlage bildeten (Skelettkörper), deren Poren mit dem reinen chemischen Element (z.B. Cu u.a.) gefüllt wurden [1]. E.P. Schalunov et al. RM 98 129^

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Aber bei großen Belastungen und erhöhten Temperaturen erweichte sich der Tränkestoff, und die Basis, indem sie Belastungen nicht aushielt, wurde deformiert oder gebrochen. Es wurden auch Versuche gemacht, ein chemisches Element, z.B. Kupfer oder seine Legierungen als die Basis und darin einführende schwerschmelzbare keramische Teilchen (Karbide, Oxide, Boride, Nitride u.a.) als härtende Phasen zu verwenden. Dabei war die Abmessung dieser Teilchen (Pulver) nicht weniger als 0,5... 1,0 (im. Die Erweichungstemperatur solcher Kompositwerkstoffe wurde wirklich höher als bei herkömmlichen Legierungen auf Basis dieses chemischen Elementes. Aber diese Erhöhung war doch ungenügend, damit diese Werkstoffe bei Temperaturen über 800°C betrieben werden könnten. Zugleich ist aus der Theorie der Dispersionshärtung bekannt, daß die Betriebstemperaturen bei Abmessungen der härtenden Phase 0,01...0,05 |im und der Entfernung zweischen Teilchen dieser Phase 0,1...0,5 um bis zu 0,90...0,95 der Schmelztemperatur von der Werkstoffmatrix erhöht werden können [2]. Die Teilchen der härtenden Phase solcher Abmessungen zu schaffen, ermöglicht das Verfahren der „inneren Oxidation" sowie das Verfahren des „Reaktionsmahlens", das in Hochenergie-Kugelmühlen, z.B., in Attritoren verwirklicht wird. Davon ausgehend, kann man die Erweichungstemperatur der dispersionsge- härteten Werkstoffe auf Kupferbasis bis zu 1000°C erhöhen, was völlig ausreichend für Anwendung solcher Werkstoffe in vielen Organen der Motoren und Antriebe verschiedener Transportmittel ist. Letzte Jahrzehnte haben bestätigt, daß solche Werkstoffe wirklich geschaffen sind und in der Wirtschaft sogar angewendet werden. Es handelt sich in erster Linie um dispersionsgehärtete Werkstoffe Cu + AI2O3, die sowohl durch „innere Oxidation" (z.B. GlidCop® AL-15, AL-25 und AL-60) als auch durch „Reaktionsmahlen" gewonnen werden [3, 4, 5]. Aber wegen der Notwendigkeit der Verwendung bei der Fertigung dieser Werkstoffe von mehrstufigen Redoxoperationen im Medium der Spezialgase, darunter im Wasserstoffmedium, ist der technologische Prozess der Fertigung von solchen Werkstoffen ziemlich arbeitsaufwendig und ihe Kosten übersteigen wesentlich die Kosten der herkömmlichen Legierungen ähnlicher Anwendung. In diesem Zusammenhang haben die angegebenen Werkstoffe bis jetzt fest jene Stellung nicht eingenommen, die sie in der Wirtschaft von ihren einzigartigen Charakteristiken ausgehend hätten einnehmen müssen. 130 RM 98 E.P. Schalunov et al.

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Außerdem wird dieses metallische System nicht in allen Fällen den Betriebsforderungen der Erzeugnisse aus ihnen gerecht. Unten wird gezeigt sein, daß die Senkung der Kosten von Werkstoffen solcher Klasse mit gleichzeitiger Modernisierung ihrer chemischen Zusammensetzung für konkrete Betriebsbedingungen erlaubt, doch, diesen Werkstoffen eine breitere Anwendung in verschiedenen Technikbereichen, darunter in Motoren und Antrieben der Transportmittel zu sichern.

2. Durch Oxide und Karbide dispersionsgehärtete Werkstoffe auf Pulverkupferbasis der Handelsmarke DISCOM®:

DISCOM® ist die Handelsmarke neuer Klasse der speziell von der Wissenschaftlich-technischen Firma TECHMA G.m.b.H. (Tscheboksary, Russland) für Maschinen- und Motorenbau, Elektrotechnik und anderen Gebieten der Technik entwickelten durch Oxide und Karbide dispersiondge- härteten Werkstoffe auf Pulverkupferbasis (Oxide und Carbide Dispersion Strenthened Copper: OCDS-Copper). OCDS-Copper wird durch Bearbeitung der Rohpulverkomposition im Attritor, Kompaktrieren der im Attritor gewonnenen Granulate in Bolzen und durch nachfolgendes Heißstrangpressen dieser Bolzen in Stangen, Rohre und andere Profile hergestellt [6,7,8]. Die absolute Abwesenheit irgendwelcher Schutz- und Redoxatmosphären bei der Herstellung von OCDS-Copper ist die Hauptbesonderheit der ausgearbeiteten Technologie: der ganze Fertigungsprozess wird in der Luft ausgeführt. Der Luftsauerstoff oxydiert im Laufe des Fertigungsprozesses Legierungselemente bis zur Bildung der Uitradispersionsteilchen (0,02...0,03 jam) von Oxiden (AI2O3, TiO2, Cr2O3 u.a.) dieser Elemente. Der in der Rohpulverkomposition vorhandene Kohlenstoff bildet mit Legierungselementen Karbide (AI4C3, TiC, Cr3C2, Cr7C3, VC u.a.), deren Größe mit der von obenerwähnten Oxiden vergleichbar ist. Darüber hinaus reduziert der Kohlenstoff das Kupfer aus seinen Oxiden, was der Matrix der Werkstoffe gute elektrische und thermische Leitfähigkeit verleiht. Diese in der Matrix gleichmäßig verteilen Oxide und Karbide gewährleisten für OCDS- Copper hohe Festigkeit, besonders bei hohen Temperaturen. Dadurch, daß im OCDS-Copper härtende Phasen (Oxide und Karbide) verschiedenartig sind, ist deren Koagulation bei hohen Temperaturen fast unbemerklich. Das, wie bekannt, verleiht solchen Werkstoffen hohe Glühbeständigkeit und Warmfestigkeit. E.P. Schalunov et al. RM 98 131

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Oxide und Karbide mit dem Restkohlenstoff, der in der Matrix in der Ultradispersionsform vorhanden ist, gewährleisten für OCDS-Copper hohe Verschleißfestigkeit. Gemäß dem obenangeführten Verfahren der Gewinnung kann man diese Werkstoffe auch als „innen-oxidierte und innen-reduzierte" nennen. Als oxide-und karbidebildende Elemente können Metalle von III, IV, V und VI Gruppen des Periodensystems von D. Mendeleev verwendet werden. Aber die breiteste ATwendung haben Zusätze von Aluminium, Titan, Chrom, Vanadium gefunden, indem folgende Systeme gebildet wurden: Cu-Al-C-O, Cu-Ti-C-O, Cu-Cr-C-O, Cu-V-C-O, Cu-Al-Ti-C-O, Cu-AI-Cr-C-0 und andere. Ausführlicher kann man sich mit chemischen Zusammensetzungen der von der Firma TECHMA entwickelten Werkstoffe und Technologie ihrer Gewinnung in russischen (Nr. Nr. 2103103, 2103134, 21031135, 2104139, 2113529, 2116370, 2117063, 2118393 u.a.) sowie ausländischen, z.B. österreichischem (Nr. 400580) Patenten der Firma bekannt machen. Dank relativ einfacher Technologie und nicht hohem Arbeitsaufwand haben diese Werkstoffe die Kosten, die 1,5...2 Mal niedriger als der bei Werkstoffen GlidCop® sind. Verschiedene Typen von OCDS-Copper als heißstranggepresste oder durch Ziehen kalibrierte Halbzeuge mit verschiedenem Profil des Querschnittes werden von der Firma TECHMA selbst sowie nach ihrer Lizenz in einigen spezialisierten, nach den Projekten der Firma in verschiedenen Städten Russlands gegründeten Betrieben hergestellt. Unter ihnen sei vor allem das Werk der Komposit- und Metallkeramik-Werkstoffe KERMET AG in der Stadt Joschkar-Ola hervorzuheben. In meisten Fällen wird OCDS-Copper als bereits Fertigerzeugnisse geliefert. Für ihre Fertigung werden alle nötigen Verfahren der mechanischen Bearbeitung (Drehen, Fräsen, Bohren, Gewindewalzen, Schleifen u.a.) und des Gesenkschmiedens verwendet.

3. Anwendung von OCDS-Copper der Handelsmarke D1SCOM® in Ventilführungshülsen und Ventilsitzringen der Benzin- und Dieselmotoren:

Der Mechanismus der Ladungswechselsteuerung ist einer der besonders verantwortlichen Organe des Motores, da er bestimmte Reihenfolge und Solldauer des Verlaufs von Prozessen des Einlaßes vom Treibsstoff und des Auslaßes der Produkte seiner Verbrennung (Gase) im Arbeitszyklus des Motores genau sichert. 132 RM98 E.P. Schalunov et al.

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Auf der Abb.1 ist die typische Konstruktion vom Mecha- nismus der Ladungswechsel- steuerung mit dem Einlaß- ventil 1 und dem Auslaßventil 2, die in Automobil-, Trak- toren-und anderen Typen der Motoren angewendet werden, dargestellt. Der Ventil ist der heißeste Teil des Motores. So kann die Temperatur im Zentrum des Auslaßventilkopfes der Moto- ren mit Zwangszündung 927...1002°C, und der Diesel- motoren 7O2...9O2°C erreichen. Bei Ventilen ohne inneren Abb.1. Antrieb von zwei Ventilreihen Kühlungshohlraum werden durch zwei Nockenwellen: 70...80% der Wärme durch 1-Einlaßventil; 2-Auslaßventil; Tragfläche des Kopfes zum 3-Ventilführungshülse; Ventilschaft und von ihm zur 4-Ventilsitzringe; 5-Nockenwellen; Führungs-hülse abgeleitet. 6-Fedem. Auf solche Weise kann die Temperatur des Sitzringes in der Zone des Kontaktes mit dem Ventilkopf in leistunsfähigen Hochgeschwindigkeits-motoren über 900°C betragen. Die Temperatur sinkt nach der Höhe der Ventilführunghülse in Richtung der Entfernung vom Ventilkopf, aber deren maximale Wert kann fast 700°C erreichen. Bei der hin- und hergehenden Bewegung des Ventils wird die Hülse dem mechanischen Verschleiß ausgesetzt und der Sitzring ist zudem unter hoher Stoßbelastung der Quetschung. Zudem ist hinzuzufügen, daß alle drei Teile unter Bedingungen eines aggressiven Mediums (der sich mit Geschwindigkeit von 400...600 m/s bewegenden, bis zu 727...1202°C erhitzten Gase) arbeiten. Ventile werden aus speziellen warmfesten Stählen hergestellt und für deren Verschleißverhütung, in der Regel, verschiedenen Verfahren der thermischen und thermochemischen Bearbeitung ausgesetzt. Ventilführungshülsen und Ventilsitzringe werden je nach Leistungs- und Geschwindigkeitschakteristiken der Motoren aus grauem Perlit- und gefeintem Roheisen, austenitischen Ventilstählen und Sinternwerstoffen auf E.P. Schalunov et al. RM 98 133

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Eisenbasis, metallurgischen Materialien und Sintern Werkstoff en auf Kupferbasis sowie Verbundwerkstoffen hergestellt. Die früher mit Erfolg angewendeten graues Roheisen Gh 1051 (3,65% C; 2,65%Si; 0,55%Mn; 0,45...0,75%P; S<0,15%; Fe-Rest) für Ventilführungs- hülsen in Motoren, z.B. der Kraftfahrzeuge FIAT und das Spezialroheisen mit Härte 207...255 HB (3,0...3,5%C; 1,8...2,5%Si; 0,6...1,2%Ni; 0,5...0,9%Mn; 0,25...0,55%Cr; Fe-Rest) für Einlaß- und Auslaßsitzringe der Ventil in Moto- ren, z.B. der russischen leistungsfähigen Motorräder URAL, werden zur Zeit immer weniger angewendet. An ihre Stelle treten härtere Spezialroheisensorten mit Trostit-Gefüge und Karbid-Einschiüßen. Aber sie zeichnen sich durch komplizierte chemische Zusammensetzung und niedrigere Fertigungsgerechtheit aus. Zum Beispiel, in Einlaß- und Auslaßven-tilsitzringen der Motoren russischer leistungsfähiger LKWs KAMAZ wird Spezialroheisen mit Härte von 420...490 HB angewendet, das folgende chemische Zusammensetzung hat: 2,4...2,9%C; 3,0...3,5%Cr; 1,5...2,0%Si; 0,5...1,5%V; 0,8...1,5%Mo; 1,2...2,0%Cu; 0,5...0,7%Ni; 0,5...1,0%Mn; 0,18...0,35%P; S<0,1%; Ca<0,02%; Fe-Rest. Einfache und billige Sintern-Stähle für Ventilführungshülsen (z.B. 1,5%C; 3,0%Cu; 0,3...0,6%S; Fe-Rest) mit Härte von 60...110HB (porige mit Tränken mit Maschienenöl) und Härte von 60...130HB (mit Sulfidierung) sind wegen der Erhöhung der Leistungs- und Geschwindigkeitscharakteristiken der Motoren praktisch bereits in die Vergangenheit weggegangen. Gut haben sich Sintern-Stähle der Firma BLEISTAHL für Sitzringe des Einlaßventils (FSN 335: 0,5...1,0%C; 2,5...3,5%Ni; 2,0...3,5%Pb; Fe-Rest) mit Härte 280...320HB und für Sitzringe des Auslaßventils (Como 7SH: 0,3...0,8%C; 1,0...2,0%Mo; 0,5...1,5Pb; 6,0...7,0%Co; 1,0...2,0%Ni; Fe- Rest) mit Härte von 320...360HB gezeigt. Aber diese Werkstoffe enthalten wesentliche Menge von Blei, dessen Anwendung wegen seiner Umweltunfreundlichkeit verboten ist. Aus demgleichen Grund wird der Sinternwerkstoff (0,5...2,0%C; 2...4%Cr; 0,2...0,4%Mo; 0,2...0,4%V; 0,2...3,0%CaF2 oder BaF2, oder MnS, oder 0,1...0,4%S; Fe-Rest, Tränken mit 1,0...2,0%Pb) ausgewechselt, der von der Fa. TOYOTA für Ventilsitzringe entwickelt wurde. An seine Stelle traten auf den japanischen Automobilmarkt verschiedene verwickeltlegierte Sinternstähle (z.B. 0,3...1,5%C; 4,0...8,0%Co; 0,5...1,5%Cr; 4,0...8,0%Mo; 1,0...3,0%Ni; 0,2...0,6%Ca; Fe-Rest). Der Bedarf des japanischen Automobilmarktes an Ventilführungshülsen beträgt ca. 10Mio Stck. jährlich. Bereits jetzt wird 30% dieses Bedarfes mit Ventilführungshülsen aus dem Verbundwerkstoff auf Eisenbasis Fe+C+ 134 RM 98 E.P. Schalunov et al.

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Zusätze keramischer Teilchen gedeckt. Wenn der Verschleiß der Roheisenhülsen ungefähr 50 \xm beträgt, so macht der Verschleiß der Hülsen aus diesem Werkstoff 10...20 jj.m aus. Gut hat sich in Ventilführungshülsen der Hochgeschwindigkeitsmotoren LADA-SAMARA das verwickeltlegierte Messing bewährt, das 56,5...59,5%Cu; 0,002...0,05%Ca; 1,0...2,0%Si; 2,0...4,0%Mn; 0,5...2,5%Ni; 0,25...1,75%Pb; 0,01...0,5%AI; Fe<0,35%; Einschüße < 1,0%; Zn-Rest (sein ausländisches Analogon ist CuZn 40 AI2 PTL 2101) enthält. Seine Härte beträgt 150... 190 HB. Aber dieses Messing darf man als Werkstoff der Zukunft wegen seiner beschränkten Warmfestigkeit nicht nennen. Zadum enthält es Blei. Moralisch veralten ebenso andere, früher mit Erfolg in Ventilführungshülsen angewendeten Legierungen auf Kupferbasis. Dazu gehört z.B. die von den Firmen Motoren- und Turbinen Union Friedrichshafen G.m.b.H. und PORSCHE AG angewendete Legierung CuNi2SiF65 (1,6...2,5%Ni; 0,5...0,8%Si; Mn<0,8%; Cr<0,04%; Cu-Rest) mit Härte von 180...220HB sowie Legierungen Thermo-Hedul FS-15 (10,0...12,0%Zn; 0,12...0,25%Si; 0,30...0,50%Mn; 0,60..0,90%Te; S<0,03%; Fe<0,10%; Ni<0,20%; Cu-Rest) mit Härte von 130... 170 HB und Aeterna VL 22 (Zn<0,5%; Si<2,0%; Al<1,4%; Pb<0,7%; Cu<58%) mit Härte von 160...190HB. Aus dem Obenangeführten ist ersehen, daß Bronzen und Messinge niedrigere Härte als spezielle Roheisen und spezielle Sinternstähle haben. Dafür aber ist niedriger der Verschleiß des Reibungspaares in Führungshülsen aus Bronze und Messing. Aber das betrifft nur den Betrieb dieser Kupferlegierungen in Motoren mit niedriger Leistung und/oder niedriger Geschwindigkeit. In leistungsfähigen- und Hochgeschwindigkeitsmotoren bei Forcierensbe- triebsweisen, wenn die Arbeitstemperatur in ihnen große Werte erreicht, ist der Betrieb dieser Legierungen wegen ihrer niedringen Erweichungstem- peratur (nicht mehr als 400...550°C) unmöglich. Trotztdem verläßt die Konstrukteure moderner Benzin- und Dieselmotoren die Idee nicht, Kupferwerkstoffe anzuwenden, die über hohe Erweichungstemperatur (über 800°C), hohe Härte und Festigkeit bei diesen Temperaturen, bessere als bei Roheisen und Stählen Wärmeleitfähigkeit und Verschleißfestigkeit des Reibenspaares unter Bedingungen aggressiver Medien verfügen. In wissenschaftlich-technischer Firma TECHMA G.m.b.H. ist die ganze Palette dispersionsgehärteter Werkstoffe auf Kupferpulverbasis (OCDS-Copper) der Handelsmarke DISCOM® speziell für Ventilführungs- hülsen und Ventilsitzringe der Benzin- und Dieselmotoren entwickelt (Tab.1). E.P. Schalunov et al. RM98 135

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Tabelle 1: Grundeingeschaften der heißstranggepreßten Haizeuge aus den dispersionsgehärteten Werkstoffen auf Pulverkupferbasis (OCDS- Copper) der Handelsmarke DISCOM® für Motorenbau.

Typ von OCDS-Copper Nr. Technische Charakteristiken CO/94 C3/03 C3/04 1. Anteil der Dispersoiden, Gew.% -6,5 -6,9 -10,0 2. Dichte, g/snr3 8,10 8,22 8,00 3. Schmelztemperatur, °C 1080 1080 1078 4. Erweichungstemperatur, °C 1000 1000 950 5. Spezifische Wärmeaufnahmefähigkeit, J/kg -°C: - bei 25 °C 411 411 420 - bei 100 °C 421 423 437 - bei 200 °C 422 428 445 - bei 300 °C 443 446 463 - bei 400 °C 466 469 488 6. Wärmeleitfähigkeit in Preßrichtung, W/m-°C: - bei25°C 65,0 55,0 50,0 - bei 100 °C 64,3 54,7 49,7 - bei 200 °C 63,6 54,3 47,9 - bei 300 °C 60,8 51,6 46,7 - bei 400 °C 56,4 47,5 43,2 7. Wärmeleitfähigkeit in Querpreßrichtung, W/m-°C: - bei 25 °C 53,4 41,8 36,8 - bei 100 °C 53,0 41,3 36,4 - bei 200 °C 51,7 40,3 35,4 48,0 37,8 33,3 - bei 300 °C 43,0 33,9 29,7 - bei 400 °C 136 RM98 E.P. Schalunov et al.

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Fortsetzung der Tab.1 8. Linearer termischer Ausdehnungskoef- fizient, x1061/°C: - bei 20...100 17,0 15,0 17,0 - bei 100...200 18,8 19,1 18,7 - bei 200...300 22,5 22,6 20,3 - bei 300...400 24,7 25,4 20,1 - bei 400...500 23,8 21,7 21,0 - bei 500...600 22,4 22,6 22,1 - bei 600...700 22,6 23,7 23,3 - bei20...300°C 19,6 19,3 19,1 9. Brinellhärte HB 5/750/30 230 258 270 10. Zugfestigkeit, MPa: - bei25°C 770 931 966 - bei 200 °C 328 562 616 - bei 400 °C 212 320 327 - bei 500 °C — — — - bei 600 °C 131 196 204 11. Relative Dehnung, %: - bei25°C 2,0 2,3 1,5 - bei 200 °C 2,5 7,3 1,6 - bei 400 °C 1,1 3,4 1,9 - bei 500 °C — — — - bei 600 °C 1,0 2,8 1,7 12. Relative Einschnürung, %: - bei25°C 4,3 20,9 3,0 - bei 200 °C 7,2 38,5 4,4 - bei 400 °C 4,5 25,4 6,9 - bei 500 °C — — — - bei 600°C 3,3 18,2 7,3 13. Druckfestigkeit in Preßrichtung, MPa 1022 1092 890 14. Druckfestigkeit in Querpreßrichtung, MPa 1062 1117 1065 E.P. Schalunov et al. RM98 137

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Fortsetzung der Tab.1 15. Relatives Stauchen in Preßrichtung, % 30 24 15 16. Relatives Stauchen in Querpreßrichtung, 28 15 10 % 17. E-Modul, MPa 110000 92000 — 18. Schubmodul, MPa 39600 43000 43500 19. Schubfestigkeit, MPa 529 563 478 20. RAD. Bruchfestigkeit, MPa — 690 590

Aus der Tab.1 ist zu ersehen, daß alle drei Typen von OCDS-Copper die Erweichungstemperatur von 950...1000°C und hohe Werte der Härte bei hohen Temperaturen haben, was sie für hitzefeste Werstoffe zu halten ermöglicht. Auf der Abb.2 sind Diagramme der Abhängigkeit der Härte von der Prüfungstemperatur („heiße" Härte) für OCDS-Copper und herkömmliche warmfeste Aluminiumbronzen angeführt.

240 OCDS-Copper DISCOM ®: 1 1 r n«w 200 N JM 160 Standa rtbronv.en : — (:UAHOFe3Mn2 120 \ C:UAIIO \

40 A

0 100 200 300 400 500 600 700 800 Prüfungstemperatur, °C Abb. 2. Vickershärte verschiedener Kupferwerkstoffe je nach der Temperatur ihrer Prüfungen. Aus der Abb.2 ist zu ersehen, daß beide Typen von OCDS-Copper im ganzen Temperaturbereich der Prüfungen höhere „heiße" Härte als die der warmfesten Bronzen hat. Im Zusammenhang damit kann man die Werkstoffe C 0/94 und C 3/04 für warmfeste halten. Die Warmfestigkeit und die Hitzebeständigkeit der angegebenen Werkstoffe sind auf das Gefüge des Werkstoffes (Abb.3) zurückzuführen, das, wie entsprechende Prüfungen gezeigt haben, aus Kupfer, a-Ti, ultradispersen Teilchen y-AI2O3 (für C 0/94) oder TiC und y-AI2O3 (für C 3/04) sowie restlichem Kohlenstoff besteht. 138 RM98 E.P. Schalunov et al.

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Abb.3. Gefüge von OCDS-Copper C 0/94: bei der 400- facher Vergrößerung (a) und bei der 20.000-facher Vergrößerung (Abdruck) (b).

Stereologische Analyse des Abdruckes hat gezeigt, daß die Durchschnitts- große von Teilchen der härtenden Phasen bei C 0/94 ca. 0,034 |um und bei C 3/04 - ca. 0,026 jim beträgt. Durchstrahlungsaufnahme einer Folie (Abb.4,a) demonstriert, daß der Werkstoff C 0/94 die subkörnige Struktur hat, was von der Abwesenheit der Erweichung vom legierten Kupfer beim Heißstrangpressen (870...900°C) seiner Granulien in die Stange zeugt. Das Elektronenbeugungsdiagramm (Abb.4,b) das von dieser Folie aufgenommen wurde, zeugt davon, daß die Subkörnige Struktur hohen Grad der Polykristallisierbarkeit (Ringe auf dem Elektronenbeugungsdiagramm) hat. Dritter vom Zentrum, dünner, ganzer Ring bestätigt, daß in der Struktur von

C 0/94 ultradisperse Phase AI2O3 tatsächlich vorhanden ist. E.P. Schalunov et al. RM98 139

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Abb. 4. Durchstrahlungsaufnahme (a) von OCDS-Copper C 0/94 bei der 57.000-facher Vergrößerung und das Elektronenbeugungsdiagramm (b), das von ihm aufgenommen wurde.

Obenbeschriebene Besonderheiten der Werkstoffstruktur sowie Vorhanden- sein in diesem Werkstoff des feindispersen freien Kohlenstoffes in Menge von 0,69...0,73 Gew.% zusammen mit obenangeführten physikalisch-mechanischen Eigenschaften mußten dem Werkstoff auch gute Betriebseigens- chaften sichern. In der Tab.2 sind Ergebnisse der Verschleißprüfungen vom Reibungspaar „Führungshülse-Ventil" angeführt, die auf dem Automobilwerk Auto-VAZ (LADA) AG durchgeführt worden sind. 140 RM98 E.P. Schalunov et al.

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Tabelle 2: Gewichtsverschleiß der Ventilführungshülsen aus dem OCDS- Copper C 0/94 und grauem Roheisen Gh 1051 bei deren Arbeit im Satz mit dem ionitrierten Ventil

Werkstoff Nr. des Gewichtsverschleiß, mg der Ventil- Reibungs- Der Ventilführungshülse Des Ventils führungshülse paares 1 0,30 0,10 OCDS-Copper 2 0,15 0,15 CO/94 3 2,65 0,85 4 0,15 0 Durchschnittswert: 0,81 0,28 graues 1 4,8 1,2 Roheisen 2 0,65 0,75 Gh1051 3 0,95 0,35 Durchschnittswert: 2,13 0,77

Ventilführungshülsen wurden aus dem OCDS-Copper C 0/94 und grauem Roheisen Gh 1051 gefertigt, das in Motoren FIAT angewendet worden war. Als Gegenkörper wurden Auslaßventile mit dem ionitrierten Ventilischaft und der Ausgangsrauhigkeit von Ra = 0,92... 1,25 |im verwendet. Schmiermitteldosierung erfolgte wie in einem realen Motor durch serienmäßige Ölabstreifkappe. Die Temperatur des Schmiermittels wurde auf dem Niveau von 121+0,5°C gehalten. Der Ventilshub war 11 mm. Die Freguenz der hin- und hergehenden Bewegung des Ventils betrag 1500 Arbeitsspiele in der Minute. Die Prüfung eines Reibungspaares dauerte 4 Stunden. Aus der Tab. 2 ist zu ersehen, daß der Gewichtsverschleiß der Ventil- führungshülsen aus dem OCDS-Copper C 0/94 2,63-fach niedriger als deren aus grauem Roheisen ist. Der Verschleiß des Ventils selbst wurde 2,75-fach niedriger. Weitere vergleichende Prüfstanderprobungen wurden im Allrussischen Fors- chungsinstitut des Motorenbaus (Sankt-Petersburg) durchgeführt. Geprüft wurden Ventilführungshülsen, die aus dem OCDS-Copper C 0/94, C 3/04 und C 3/03 sowie aus der Bronze CuNi2Si F65 DIN 17666, die die Anwendung in Motoren der Fa. Motoren- und Turbienen Union G.m.b.H. und PORSCHE AG gefungen hatte, gefertigt wurden. Die Ventile wurden aus dem Verchromten E.P. Schalunov et al. RM98 141

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Ventilstahl gefertigt. Die Ergebnisse dieser Prüfstanderprobungen sind in der Tab.3 angeführt.

Tabelle 3: Die Vergleichsdaten der Verschleißfestigkeit und Gratbestendigkeit und Zerreißfestigkeit der Ventilführungshülsen aus OCDS-Copper und der Bronze CuNi2Si F65 und der Kopplung „Ventilführungshülse-Ventilschaft"

Material Bezei- Mittelkennziffer OCDS-Copper Bronze chnung CO/94 C3/04 C3/03 CuNi2SiF65 Härte, MPa HV 2540 2600 2650 2250 Spezifischer Verschleiß: 3 - der Hülse, mm /km JH 0,024 0,036 0,057 0,403 - des Ventils, mg/km Jv 0,028 0,073 0,057 0,036 Koeffizient der Verschleißfestigkeit:

- der Hülse KH 17 11 7 1 - der Kopplung KH-V 21 5 3 1 Reibungswert fr 0,020 0,018 0,030 0,008 Kritische Belastung, kg PKR 145 134 110 88 Spezifische Einklem- mungsbelastung, PF 663 468 368 169 kg/sm2 Koeffizient der Gratbe- K 2,54 2,09 1,65 1,00 ständigkeit F Koeffizient der 7,3 3,2 2,2 1,0 Lebensdauer KL

Die in der Tab.3 angeführten Parametern und Charakteristiken sind nach folgenden Formeln berechnet. Koeffizient der Verschleißfestigkeit für den Werkstoff der Ventilführungsshülse K„ und Kopplung "Ventilführungshülse-Ventil" K„_V entschprechend:

KH = -^ und KH_y =KHX. — , JH J v wo Jfr spezifischer Verschleiß der Ventilführungshülse aus der Bronze (Etalon); 142 RM 98 E.P. Schalunov et al.

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jy- spezifischer Verschleiß des Ventils, der in Kopplung mit Ventilführungshülse aus der Bronze arbeitet. Koeffizient der Gratbeständigkeit:

wo P^K,Py- entsprechend kritische Belastung und spezifische Einklemmungs- belastung für Kopplung des Ventils mit Ventilführungshülsen aus der Bronze. Koeffizient der Lebensdauer bei der Annahme des gleichen Einflußes von Verschleißfestigkeit und Gratbeständigkeit auf Lebensdauer:

Aus der Tab.3 folgt, daß alle drei Typen von OCDS-Copper nach tribotechnischen Charakteristiken wesentlich die Bronze CuNi2SiF65 übersteigen, wobei die besten Ergebnisse der Werlstoff CO/94 hat, der den niedrigsten Verschleiß nicht nur der Ventilführungshülse sondern auch dem Ventil, d.h. ganzen Kopplung sichert. Aufgrund durchgeführter obenangegebener und anderer Prüfstanderprobun- gen sowie Naturerprobungen auf über 10 Automobilwerken in verschiedenen Typen der Benzin- und Dieselmotoren als der Werkstoff für Ventilführungs- hülsen wurde OCDS-Copper CO/94 und für Ventilsitzringe-OCDS-Copper C 3/04 empfohlen. Obenangegebene Typen von OCDS-Copper, und zwar C 0/94 und C 3/04 werden für Fertigung aus ihnen entsprechend der Ventilführungs- hülsen und Sitzringe der Benzin- und Dieselmotoren angewendet. Diese Werkstoffe werden hergestellt nach der Lizenz der Firma TECHMA im Werk KERMET als Stangen und Rohre, aus denen dann durch Verfahren der mechanischen Bearbeitung in der Firma TECHMA Ventilführungshülsen und Sitzringe gefertigt werden. Sie werden auf den Montagefließband des russischen MOTOREN WERKes Zavolshje AG zur Komplettierung mit ihnen der 16-Ventil Benzin- und Dieselmotoren für LKWs und PKWs geliefert. Diese Erzeugnisse aus OCDS-Copper der Typen C 0/94 und C 3/03 werden schon seit langem in allen Rennwagen (LKWs und PKWs) der führenden russischen Automobilwerke (LADA-SAMARA, UAZ, GAZ, KAMAZ u.a.) angewendet. Mit Ventilführungshülsen aus dem Werkstoff C 0/94 werden im Dieselwerk DIESELPROM AG in der Stadt Tscheboksary die Motoren MTU 8V396TC4 des deutschen Konzernes Motoren- und Turbinen Union G.m.b.H. Friedrichshafen während ihrer Reparatur komplettiert. E.P. Schalunov et al. RM 98 143

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Ventilführungshülsen und Sitzringe aus den obenangegebenen Typen von OCDS-Copper werden von einigen russischen Kfz-Reparaturwerkstätten bei der Reparatur der Motoren ausländischer Autos angewendet. Obenangegebene Werkstoffe wurden von einigen deutschen, österreichis- chen und koreanischen Firmen untersucht, die das Interesse an Durchführung der Prüfstand- und Naturerprobungen der Erzeugnisse aus diesen Werkstoffen ausgesprochen haben. Heißstranggepreßte Halbfabrikate werden gemäß Technischen Forderungen (TF 7960-001-13092819-99) der Fa. TECHMA, die vom Staatskomitee der Standartisierung der Russischen Föderation (GOST RU) verabschiedet und registriert wurden.

4. Anwendung von OCDS-Copper der Handelsmarke DISCOM® in stromabnehmenden Leisten der Scherenstromabnehmer der Hochge- schwindigkeitszüge:

Die Stromabnahme von Kontaktleitungen für den Antrieb des Elektrotransports (Eisenbahnzüge, U-Bahnzüge, Obuse, Straßenbahnen u.a.) erfolgt immer unter Bedingungen des elektrischen Gleitkontaktes. Wegen der Bewegung zeichnet sich dieser Kontakt durch Unbeständigkeit aus, die die Möglichkeit der Erscheinung von Funken und Lichtbogen bedingt. Das Funken vergrößert den Verschleiß des Kontaktpaares „stromabnehmende Leiste - Kontaktleitung" und der Lichtbogen ruft Verdampfung des Werkstoffes hervor, was die Rauhigkeit der Kontakte erhöht. Das Eis auf Kontaktleitungen ruft unstetige Bewegung hervor und vergrößert Funken. Der Regen erhöht auch den Verschleiß des Kontaktpaares. Unter diesen Bedingungen erhöhen sich die Forderungen an den Werkstoff des stromabnehmenden Bauelementes vom Stromabnehmer des Elektro- transports, das die ununterbrochene Stromabnahme bei jeder Wetterlage sichern muß. Funken und Lichtbogen, dessen Leistung sich mit der Erhöhung des abnehmenden Stromes erhöht, sowie mechanisches Reiben im Kontaktpaar bedingen wesentliche Erhöhung der Temperatur des stromabnehmenden Bauelementes im Vergleich zur Umwelttemperatur, wobei die Temperaturer- höhnung größer je nach der Vergrößerung von der Stärke des abnehmenden Stromes sowie der Bewegungsgeschwindigkeit des Transportmittels wird. Im Zusammenhang damit ist es nötig, daß die stromabnehmenden Bauelemente erhöhte Lichtbogenfestigkeit, Verschleißfestigkeit, hohe Festigkeitseigens- chaften bei erhöhten Temperaturen sowie hohe Antifriktionseigenschaften 144 RM 98 E.P. Schalunov et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 haben. Für Vorbeugung der Stromverluste muß der Werkstoff dieser Bauelemente auch gute elektrische Leitfähigkeit haben. In elektrischen Transportmitteln, die sich mit niedrigen Geschwindigkeiten (30...80 km/h) bewegen und elektrischen Strom der Stärke von 70...150 A verbrauchen (Straßenbahnen, Obuse) fanden und finden bis jetzt Anwendung in vielen Ländern (Russland, USA, Deutschland, Schweden u.a.) Kohlen- und Graphitstromabnehmer (Gleitschune). Ihre Härte ist nicht sehr hoch, trotzdem macht ihr Verschleiß nur 2,0...2,5 sm3/1000km beim Verschleiß der Kontaktleitung bis zu 1,5 sm3/1000 km aus. Obusstromabnehmer, die aus Graphitkompositionen gefertigt werden, haben bei der optimalen Wetterlage den Lauf bis zu 250 km. Beim Regen oder Schneefall erniedrigt sich ihre Lebensdauer. In einigen Fällen fanden Anwendung in stromabnehmenden Bauelementen der Obuse Kompositionen aus Eisengraphit Fe-3...5%C, die mit Metallpolymeren (Cu+Kapron, Cu-Cd+Kapron) gatränkt sind. Die Lebensdauer der Stromabnehmer aus solchen Kompositionen erhöhte sich bis zu 2700...3600 km. Die Werkstoffe auf Kupfergraphitbasis, z.B. 82%Cu-5%C-5%Ni-8%Sn, dessen Härte 48...49 HB und IACS = 50% ist und 70% Cu-5%C-15%Fe - 5%Pb, dessen Härte 54...55 HB und IACS = 39% beträgt, fanden breite Anwendung in stromabnehmenden Leisten der Scherenstromabnehmer von elektrischen Zügen, die sich mit der Geschwindigkeit bis zu 90 km/h beim Speisestrom bis zu 400A bewegen. Dabei betrug die Lebensdauer der Arbeit solcher Leisten bis zum Wechsel nicht mehr als 10.000 km. Mit der Erhöhung der Bewegungsgeschwindigkeiten der elektrischen Züge bis über 100 km/h und Erhöhung der Erhitzugstemperatur von der Kontaktzone des Stromabnehmers bis zu 400°C hat man begonnen in Eisenbahnen vieler Länder die härtere (150...160 HB)und die mit höheren elektrischen Leitfähigkeit (IACS = 72...75%) Bronze CuCriZr anzuwenden. Die Leisten aus dieser Bronze ordnet man an den Rändern von der aus dem Graphit hergestellten Leiste an. Die Graphitleiste dient als Quelle der Trockenschmierung für Bronzeleisten beim Gleiten auf ihnen der elektrischen Kontaktleitung. Als die Geschwindigkeit der Elektrozüge sich bis über 160 km/h und die Temperatur in der Kontaktzone des Stromabnehmers bis zu 500-600°C erhöhte, wurde die Anwendung von CuCriZr wegen ihrer Erweichung bei solchen Temperatuten verboten. In vielen Fällen hat man begonnen, Werkstoffe auf Eisenbasis anzuwenden. Es wurde z.B. ein Werkstoff angeboten, der aus der im voraus zusammengesetzten Pulvermischung (2... 14% Cr, 2...7% MeS, 1% P, Fe - E.P. Schalunov et al. RM 98 145

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Rest) und dann aus dem festen Sintem-Gerippe besteht. Das gewonnene Gerippe wird im Vakuumofen durch Blei oder Bleilegierung getränkt. Dank solcher komplizierten Struktur hält dieser Werkstoff tatsächlich Temperaturen von 500...600°C aus. Aber er verfügt über sehr niedrige elektrische Leitfähigkeit (IACS<10%), was große Stromverluste und das zusätzliche Selbsterhitzen der stromabnehmenden Leiste bedeutet. Das führt seinerseits zum erhöhten Verschleiß dieser Leiste. Die angegebene Aufgabe wurde mit Hilfe des speziell von der wissenschaftlich-technischen Firma TECHMA G.m.b.H. entwickelten disper- sionsgehärteten Werkstoffes auf Kupferpulverbasis Cu-AI-C-0 (OCDS- Copper) C 0/97 der Handelsmarke DISCOM® gelöst. Die physikalisch-technischen Grundeigenschaften dieses Werkstoffes sind in derTab.4 angeführt.

Tabelle 4: Vorläufiges Produktblatt für OCDS-Copper C 0/97 der Handelsmarke DISCOM® Kriterien Daten 1. Bezeichnung OCDS-Copper (Cu AI2O3 C) DISCOM® C 0/97 TF 1479-002-13092819-01 2. Matrix des a-Cu(AI) Werkstoffes 3. Härtendene Phasen im AI2O3> C Werkstoff 4. Eigenschaften Spezifisches Gewicht bei 20 °C 8,69±0,09 g/snr Schmelzpunkt ca. 1083°C Erweichungstemperatur min. 830 °C Wärmeleitfähigkeit bei 20 °C 335±5 W/mx°K Ausdehnungskoeffizient (20- 150 °C) 16,6±0,1 um/mx°C Elektrische Leitfähigkeit bei 20 90:2 % IACS °C 3 146 RM 98 E.P. Schalunov et al.

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Fortsetzung der Tab. 4 Zugfestigkeit bei 20 °C ca. 425 N/mm2 bei 500 °C ca. 120 N/mm2 0,2-Dehngrenze bei 20 °C ca. 365 N/mm2 bei 500 °C ca. 105 N/mm2 Bruchdehnung A5 bei 20 °C 22±2% bei 500 °C 20±2% Zugfestigkeit nach 1h Glühen 2 bei 800°C in N2 ca. 385 N/mm 2 in H2 ca. 390 N/mm 0,2-Dehngrenze nach 1h Glühen 2 bei 800°C in N2 ca. 330 N/mm 2 in H2 ca. 350 N/mm Bruchdehnung A5 nach 1h Glühen bei 800°C in N2 20,8±2% in H2 19,6±2% Rockwellhärte bei 20 °C ca. 71,5 HRB Rockwellhärte nach 1h Glühen bei 800°C in N2 ca. 63,2 HRB in H2 ca. 65,2 HRB Brinellhärte bei 20 °C ca. HB 125 bei 500 °C ca. HB 70 Brinellhärte nach 1h Glühen bei 800°C ca. HB 115 Vickershärte bei 20 °C ca. 135HV30 Druckfestigkeit bei 20 °C ca. 1900 N/mm2 Stauchung zur Zerstörung bei 65±2% 20 °C 5. Anwen- Anlaßbeständiger Kupferwerkstoff mit hoher dungsbereich elektrischer Leitfähigkeit für elektrotechnische Anwendungen (z.B. Stromgleitkontakte, Instrument für Schweißen u.s.w.)

Anmerkung: In der Tabelle sind Ergebnisse der von wissenschaftlich- technischen Firma TECHMA G.m.b.H., Max-Planck Institut und IFAM durchgeführten Prüfungen verallgemeinert. E.P. Schalunov et al. RM 98 147

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Wie aus der Analyse der Tab.4 hervorgeht, verfügt der entwickelte Werkstoff über gute elektrische Leitfägigkeit (IACS=86...92%), hohe Erweichungstem- peratur (über 830°C), ausreichende Härte (ca. 125HB) und Prastizität (A5 = 20...24%). Die Endgefüge dieses Werkstoffes stellt an sich die Kupfermatrix mit gleichmäßig verteilten in ihr ultradispersnen Teilchen (0,024...0,036 (am) von schwerschmelzbaren Verbindungen Aluminiums (y-AI2O3) sowie Teilchen vom restlichen Kohlenstoff dar. Die Produktion dieses Werkstoffes und fertiger mechanischbearbeiteter stromabnehmender Leisten aus ihm mit verschiedenem Profil des Querschnittes und verschiedener Länge ist nach der Lizenz der Firma TECHMA auf dem Werk KERMET AG organisiert. Beim Zusammenbau der Abnehmer-Vorrichtung des Elektrozuges wird zweischen zwei Leisten aus OCDS-Copper eine Leiste aus Graphit aufgestellt (Abb. 5). ,10» tiillu- Stromabnehmende Leisten der ~ Scherenstromabnehmer der Elektrozüge von drei Typenmaßen, die aus dem obenangegebenen Werkstoff gefertig sind, werden auf Hochgeschwindigkeitseisenbah nstrecken (über 160 km/h) der italienischen Eisenbahn angewendet. Der Werkstoff OCDS-Copper CO/97 der Handelsmarke DISCOM® wird gemäß Technischen Forderungen (TF 1479-002-13092819-01) der Fa. TECHMA, die vom Abb.5. Stromabnehmer der Staatskomitte der Standar- Elektrozüge mit hoher (>160 tisierung der Russischen km/h) Bewegungs- Föderation (GOST RU) geschwindigkeit: verabschiedet sind, hergestellt. 1-stromabnehmende Leisten; 2-Graphitleiste. 148 RM 98 E,P. Schalunov et al.

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5. Schlußfolgerungen:

Zwei obenangeführte Beispiele effektiver Anwendung der dispersionsge- härteten Werkstoffe auf Kupferpulverbasis (OCDS-Copper) der Handels- marke DISCOM® schöpfen alle Anwendungsgebiete von diesen Werkstoffen nicht aus. Die Werkstoffe, die in vorliegender Arbeit präsentiert sind, befinden sich auf entgegengesetzten Polen aller möglichen Palette von Werkstoffen dieser Klasse. Wahrlich, wenn die Werkstoffe für Ventilführungshülsen und der Motoren (C 0/94 und C 3/04) Härte von 230...270HB und entsprechend elektrische Leitfähigkeit und entsprechend Wärmleitfähigkeit) von ca. 15...19% von analoger Charakteristiken des Kupfers haben, so hat der Werkstoff für Stromabnehmende Leisten der Scherenstromabnehmer von Hochgeschwindigkeitszügen die Härte von ca. 125HB bei der elektrischen- und Wärmeleitfähigkeit von 86...92% von analoger Charakteristiken des Kupfers. Zwieschen diesen Werkstoffen liegt breite Palette der Werkstoffe (C 0/70, C 0/77, C 0/98, C 0/99, C 1/15, C 1/20, C 1/24, C 1/54, C 1/56, C 2/05, C 3/06, C 3/07, C 4/01, C 4/12 u.a.), die in Schweißtechnik, Maschienenbau, Gerätebau und anderen Technikbereichen weit verbreitet sind. Unter ihnen gibt es auch Werkstoffe, die aufgrund durchgeführter Untersuchungen und Prüfungen auch für Anwendung in modernen Benzin- und Dieselmotoren empfohlen werden können. Es geht z.B. um die Möglichkeit des Wechsels vom Werkstoff GLYCO-40 in der Lagerbuchse für Ölpumpe, von der Bronze CuSu8F46 in der Buchse der Kipphebelachse u.a. Gemeinsames, was verschiedene Typen von OCDS-Copper vereinigt, sind ihre hohen Warmfestigkeit, Hitzebeständigkeit und Verschleißfestigkeit, hohe Fertigungsgerechtheit und nicht hohe Produktionskosten im Vergleich zu nicht nur Produktionskosten anderer dispersionsgehärteter Werkstoffe auf Kupferbasis (z.B. Glidcop®) sondern auch sogar im Vergleich zu herkömmlichen Bronzen der Systeme Cu-Cr-Zr, Cu-Ni-Ti-Be, Cu-Co-Be und anderen.

Danksagung: Die Autoren vorliegender Arbeit sprechen ihnen Dank aus Herrn Professor Dr.G. Jangg (TU Wien) für wertvolle Ratschläge bei der Entwicklung der Werkstoffe dieser Klasse, Herrn W. Pleus und Herrn Dr.-Ing. H. Figge E.P. Schalunov et al. RM 98 149

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(Pleuco G.m.b.H., Carl Pleus und Söhne) für Organisation und Durchführung einiger Untersuchungen und Prüfungen der Werkstoffe C 0/94, C 3/03 und C 3/04, der Abteilung EGL2 (Herr Stenzel) der Firma PORSCHE AG für Untersuchung chemischer Endzusammensetzung und der Härte von Ventilführungshülsen aus dem Werkstoffe 0/94 des Rennwagens VAZ-LADA Gamma 2 nach einer langwieriger Autorallye sowie der Firma RÖTECH G.m.b.H. (Herr S. Wendland) für Einsatz der stromabnehmenden Leisten aus OCDS-Copper C 0/97 in Elektrozügen mit höher (> 160 km/h) Bewegungsgeschwindigkeit.

Literatur:

1. F. Eisenkolb: Fortschritte der Pulvermetallurgie (Akademie-Verlag, Berlin, 1963) 2. E. Orowan: Inst, of Metals, Rep. Ser. (1948) 3. GLIDCOP® - Aluminiumoxideverstärkter Kupferdispersion-werkstoff. Prospekt der Fa. SCM Metal Products, Inc., 1994 4. G. Korb, G. Jangg, F. Kutner, M. Seirafi: Metall, Heft 11 (1981), pp. 1126- 1130 5. L. Berglin and O. Sjöden: Scand. J. Metallurgy 12 (1983), pp. 184-188 6. ÖP Ns 400580, 1996 (E.P. Shalunov u.a.) 7. V. Dovydenkov, E. Shalunov: Proc. 1998 PM World Congress, Vol.1, pp. 372-376 (EPMA, Granada, 1998) 8. E. Shalunov, J. Lipatov, D. Golubyatnikov: Proc. Int. Conf. Deformation and Fracture in Structural PM-Materials, Vol. 1, pp. 365-373 (ed. L. Pariiak et al, IMR SAS Kosice, Piestany, 1999) AT0200474

JT50 RM 102 V. Reglero et al.

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Tungsten and Optics

V. Reglero*, I. Ruiz Urien+, T. Velasco*, I. Santos+, J. Rodrigo*, J. Zarauz+, J. L. Gasent*, J. Alamo*, R. Chato*.

* Astronomy and Space Science Group, Instituto de Ciencia de los Materiales. Universidad de Valencia. Apdo. 22085. E-46071 Valencia, Spain + SENER, Avda. Zugazarte 56, E-48930 Las Arenas, Vizcaya, Spain

Abstract

The ESA INTErnational Gamma Ray Astrophysical Laboratory (INTEGRAL) satellite is the largest Gamma-Ray Astrophysics mission foreseen for this decade. Since its selection in 1994, three very innovative technologies were identified to achieve the INTEGRAL scientific goals: new generation Solid State Detectors, Active Veto Systems and Spatial Signal Multiplexing Systems. The high energy instruments of the mission, IBIS, SPI and JEM-X, have been developed by using the above referred technologies.

The Grupo de Astronomia y Ciencias del Espacio (GACE) of the University of Valencia was chosen as the responsible to design and develop the Spatial Signal Multiplexing Systems (Coded Masks) for the three instruments, appointing the Spanish company SENER as the Prime Industrial Contractor. During the early design phase (1994-96) different materials were considered to modulate the high energy signals. Tungsten was selected as the optimum material for implementing the codes due to both its high density and large atomic number that provide the required stopping power. Codes follow HURA and MURA patterns, being the total mass of W used of around 280 kg.

The scope of this paper is to summarise the milestones of the development programme carried out from 1994 until now, once the Flight Models have been manufactured, tested, delivered and mounted on the INTEGRAL satellite at Alenia and CNES (Oct/Nov 2000). We pay special attention to describing the complex Masks Themoelastic Designs, involving W coupled to composites and Ti. Mask Assemblies stability was tested in a wide range of temperatures (-80°C +40°C) and mechanical stress (12g) at INTA facilities for Flight Qualification. V. Reglero et al. RM 102 151

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Keywords: Tungsten, Densimet, Optics, High Energy Astrophysics, INTEGRAL, imaging. 1. Introduction.

1.1. INTEGRAL Mission

The ESA INTErnational Gamma Ray Astrophysical Laboratory (INTEGRAL) was selected in 1994 as a medium size scientific mission in ESA Horizon 2000 Programme. INTEGRAL main scientific goal is to perform fine spectroscopy and imaging of celestial gamma-ray sources in the energy range from 15 keV to 10 MeV. Launch foreseen date is April 2002.

The instrumentation on board INTEGRAL was designed to use three very innovative technologies: Massive Active Shielding (Veto Systems), Solid State Position Sensitive Detector Planes and Spatial Signal Multiplexing Systems.

INTEGRAL payload consists of two main gamma-ray instruments, the Spectrometer (SPI) and the Imager (IBIS). Each of them has spectral and angular resolution capabilities, but they have been optimised in order to complement each other and to achieve overall excellent performances: AE/E=500 at 1MeV and point source location accuracy 1 arc min. The two main instruments are complemented with two monitors: the X-Ray Monitor (JEM-X), in the 3 to 35keV energy range and the Optical Monitor (OMC) between 500nm and 850nm. See Figure 1. The three high energy instruments (SPI, IBIS and JEM-X) share a common principle of operation: they are all Coded Mask Telescopes.

All the instruments have benefited from a wide collaboration encompassing scientific institutes in the ESA member states, as well as USA, Russia, Czech Republic and Poland. 152 RM 102 V. Reglero et al.

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Imaser IBIS X-Ray Monitor JEM-X Optical Monitor OMC Spectrometer SPI

itiiilitP

Figure 1. Instrumentation on board INTEGRAL

1.2. Coded Masks

Imaging gamma rays is a difficult task. Photons can not be focussed by conventional mirrors or lenses. The solution is to use a pinhole camera, one of the less sophisticated optical devices. The radiation passes through a hole in an opaque plate and the image is formed in the detector. It is the oldest photograph camera from Daguerre in the XIX century. See Figure 2a.

However, gamma-ray photons from distant objects are rare (few per hour), and many holes, following an identifiable pattern (code), are needed in the opaque plate in order to get enough number of photons to produce images, this is the principle of the Spatial Signal Multiplexing. See Figure 2b. The image on the detector is thus the convolution product of the source (sky) and the code pattern. Hundreds/thousands of images coming through the holes are then superimposed. The original source sky picture must be reconstructed (deconvolved) by software techniques.

Figure 2a. Pinhole Camera Figure 2b. Signal Multiplexing V. Reglero et al. RM 102 153

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2. Coded Masks on INTEGRAL

The Astronomy and Space Science Group (GACE) of the University of Valencia have been the responsible group for the design, manufacturing, qualification, scientific calibration and delivery of the Coded Masks on board INTEGRAL since the instruments selection in 1995. Main contractor for the Masks activities was the Spanish company SENER.

The coded aperture geometry of each Mask was chosen according to the geometrical properties of the detector plane, instrument focal length and scientific requirements. Their development has been accomplished in close connection with the detector plane designers of each instrument. For further explanations see (1).

Tungsten was the baseline material selected for implementing the codes of the masks, due to both its high density and large atomic number which provide the required stopping power (opacity) at the different energy ranges. Main drivers for the design of the INTEGRAL Coded Masks have been the use of large amounts of Tungsten of different thickness together with both composites and Ti elements providing support and interfaces for the W Codes with the Satellite Payload Module.

2.1. SPI Coded Mask.

The SPI Coded Mask is a circular array of 127 hexagon pixels 60mm side to side, 63 opaque and 64 transparent to gamma ray within the operational energy range (20 keV to 8 MeV). The HURA pattern is inscribed into a 780mm diameter circle. The code has been selected according to the detector geometry (19 hexagonal Germanium detectors), instrument focal length (1.7m) and field of view (35 deg).

Tungsten was the material selected to implement the Mask Code. Densimet 18 alloy was preferred instead of pure Tungsten due to its thermoelastic properties. The code thickness is the largest of INTEGRAL Masks (30mm) with an opacity of 95% at 1MeV and a total mass of 109 kg. Code plates have been machined by Electrodischarge Wire Cutting (see paper by Reglero et ai in this Seminar). The code is supported by a carbon fibre structure 154 RM 102 V. Reglero et al.

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surrounded by a Ti ring that provides the interface with the Satellite. See Figure 3. For further descriptions, see (2).

The main design goal of the Mask was to minimize the amount of passive materials in the open pixels. Large amount of passive materials could jeopardise the detector performances at low energies. The use of composite materials guarantees the required stiffness and strength for the whole mask assembly with a minimum loss of transparency. 6 kg of Carbon Fibre support the code under Qualification, Launch and Operation conditions (acceleration 18g; temperature -55°C to +40°C), being the total weight of the Mask 131 kg, including interface layers.

Figure 3. SPI FM Coded Mask during Dimensional Control at INTA

2.2. IBIS Coded Mask.

The IBIS Mask is a MURA code (53x53) four times repeated. Pixel size and shape have been selected according to detector geometries. Pixel size is 11.2x11.2mm with a total coded area of 1064x1064mm2. The IBIS Field of View up to zero response is 29 deg, the angular resolution 12 arc min and the point source location capability 1 arc min. The instrument energy range covers from 15keV to 10 MeV. V. Reglero et al. RM 102 155

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Tungsten was the material selected to implement the Mask Code. As in SPI Coded Mask, Densimet 18 alloy was preferred instead of pure Tungsten. The code has been machined in four 16mm thickness plates having a total mass of 164 kg. This thickness provides the required opacities in the instrument energy range (70% at 1.5 MeV). Code plates have been machined by Electrodischarge Wire Cutting (see paper by Reglero et al in this Seminar).

The code is supported by a carbon fibre structure surrounded by a composite frame that provides the interface with the Satellite. The use of composite materials guarantees the required stiffness and strength for the whole mask assembly with a minimum loss of transparency, while thermoelastic deformations are also minimised. 16 kg of carbon fibre support the total weight of the Mask (197 kg) under Qualification, Launch and Operation conditions (acceleration 12g; temperature -60 to +40°C). See Figure 4. For further descriptions, see (3).

Figure 4. IBIS FM Coded Mask during Vibration Test at INTA

2.3. JEM-X Coded Mask

The JEM-X instrument is the INTEGRAL X-Ray Monitor (two units) designed to enlarge the mission wavelength coverage down to the 3 keV range. The 156 RM 102 V. Regiere et al.

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JEM-X Mask code is a HURA pattern formed by 24247 hexagonal pixels 3.3mm side to side (25% open and 75% opaque). The total coded area has a diameter of 535mm. The pattern is cut in a 0.5mm thick pure Tungsten plate, providing the required opacity in the JEM-X energy range. The instrument Field of View is 13 deg and their angular resolution is 3 arc min.

The design driver was to define a system able to survive launch and operational conditions (acceleration: 12g, 1st axial eigenfrequency > 60Hz, temperature: -60°C to +35°C) maintaining the code elements location and minimising the use of support structures obscuring the Field of View. The code was manufactured by Electrodischarge Wire Machining (see paper by Reglern et al in this Seminar).

The interface between the Satellite and the Masks is provided by a Ti forged ring. An innovative element is the presence of 44 Cu-Be elements (Pretensioning System) attaching the coded plate to the Ring and reducing dramatically the need of additional support structures. Two Ti Strongbacks (Exoskeleton) are located on both sides of the Code membrane to increase its stiffness and strength. They have been designed to minimize the occlusion of open pixels (less than 2%). See Figure 5. For further descriptions, see (4).

Figure 5. JEM-X FM1 Coded Mask V. Reglero et al. RM 102 157^

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3. Coded Masks Development.

3.1. Model Philosophy.

Three models for each Coded Mask have been developed between 1997 and 2000:

• Development Model (DM): Technology and manufacturing demonstrator.

• Structural and Thermal Model / Qualification Model (STM/QM): Built in flight standards and tested at full Qualification at Subsystem and Satellite level.

• Flight Models (FM): Tested at Acceptance Level and accepted by ESA (IBIS & JEM-X Masks) and CNES (SPI Mask). Mounted on the INTEGRAL Payload and SPI instrument in October 2000

3.2. Manufacturing and Assembly

More than 10 Spanish companies have taken part in the Masks manufacturing and assembly. Main activities can be summarised as follows:

• Tungsten Codes Machining (see paper by Reglero et al in this Seminar) • Metallic parts machining, including the JEM-X and SPI Ti forged rings, JEM-X Strongbacks and Pretensioning System (Cu-Be elements), inserts, etc • Composites parts manufacturing, including IBIS & SPI sandwich panels (honeycomb Nomex and two Carbon Fibre layers) and IBIS Carbon Fibre frame. • Joining between composites and metallic parts. • Codes integration on Mask Assembly. Three specific tools were developed for each Mask, in order to locate the code elements within the hard positioning requirements (better than 0.1mm over surfaces up to 1 m2) 158 RM 102 V. Reglero et al.

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• Coded Masks storage and transport: Three specific transport containers were manufactured in order to store and transport the Masks, minimising the induced stresses.

3.3. Testing.

The Masks have been largely tested in order to verify their capability to survive storage, launch and operation environments during INTEGRAL foreseen life time (5 years).

The Coded Masks have passed both Qualification and Acceptance test campaigns at Subsystem level (at INTA) and at Satellite levels (at Alenia, Italy and CNES, France).

Main tests performed on the Masks were:

• Dimensional Control • Mass properties • Vibration Test (Sinus and Random Vibrations, 12g input acceleration) • Thermal Vacuum and Thermal Cycling (-60°C to +45°C) • Electromagnetic Compatibility

The Masks have also passed through scientific test in order to determine their response to the X and Gamma radiation.

3.4. Integration.

The Flight Models of the Coded Masks were delivered to ESA and CNES in October/November 2000. Figures 5a and 5b show the FM Masks integrated on the Satellite (IBIS and JEM-X) and SPI instrument. V. Reglero et al. RM 102 159

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Figure 5a. IBIS & JEM-X FM Masks Figure 5b. SPI FM Mask after its after their integration on Satellite FM integration on SPI instrument at at Aienia, Turin (October 2000) CNES, Toulouse (November 2000)

4. Conclusions.

Imaging is the driver of the ESA INTEGRAL Mission. High Energy Astronomy research requires accurate point source location to perform multiwaveiength studies of the cosmic gamma-ray emitters. New technologies have been developed (1995 to 2000 by GACE) to achieve this scientific goal, the use of large Spatial Signal Multiplexing Systems (Masks).

The Optical Systems based on the use of Coded Masks together with Solid State Pixelated Detector Planes provide a Point Source Location Capability of 1 arc min, that is 3600 times better than that of the last NASA CGRO mission (1990-98). 160 RM 102 V. Reglern et al.

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Tungsten was the selected as the optimum material for the new High Energy Optics based on binary signal multiplexing. The four Masks (SPI, IBIS and two JEM-X) have been delivered to ESA and mounted on the Satellite in October 2000 (Alenia/CNES) waiting four launch foreseen by the middle of 2002

The four INTEGRAL Masks are by far the more sophisticated and larger "optical systems" for High Energy Astrophysics ever built. Furthermore, they represent an innovative application for the W as raw material for Optics.

Acknowledgements.

We would like to thank more than 14 European industries that have participated in the Mask development, with special mention to our Prime Contractor SENER, test facilities people at INTA, our Tungsten supplier PLANSEE and its representative in SPAIN Technalloy.

References.

(1)Reglero V. Sanchez F., Perez F., Rodrigo J.M., Collado V., Peris J., Donate M., Velasco T., Navarro A,, Sabau D., Reina M., Sanchez A., Eiriz E., Urteaga J.M., Sebastian P., Mayo L, Beiztiegui R., INTEGRAL Signal Multiplexing, 1999, Astrophysical Letters and Communications, Vol 38, p 381. (2) Rodrigo J, Velasco T, Zarauz J, Sanchez F, 2000, SPI Mask General Requirements Specification, IN-SP-UV-SPE-002. (3) Gasent J.L, Rodrigo J, Zarauz J, Reglero V, 2000, IBIS Mask General Requirements Specification, IN-IB-UV-SPE-002. (4) Velasco T, Rodrigo J, Zarauz J, Reglero V, 2000, JEM-X Mask General Requirements Specification, IN-JX-UV-SPE-001. AT0200475 V. Reglero et al. RM 103 161

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High Precision Tungsten Cutting for Optics

V. Reglero*, I. Ruiz Urien+, T. Velasco*, I. Santos4", J. Rodrigo*, J. Zarauz+, J. L. Gasent*, J. Älamo*, R. Chato*, G. Clemente#, C. Sanz-Tudanca#, J.L.Lopez*.

* Astronomy and Space Science Group, instituto de Ciencia de los Materiales. Universidad de Valencia. Apdo. 22085. E-46071 Valencia, Spain + SENER, Avda. Zugazarte 56, E-48930 Las Arenas, Vizcaya, Spain # Mecanizados Gines S.A., P.l.de Bayas C/Orön-parc 4, 09200 Miranda de Ebro, Spain Abstract

Coded Masks provide the Signal Multiplexing tools for the High Energy Instrumentation of the INTEmational Gamma Ray Astrophysics Laboratory (INTEGRAL) ESA Project addressed to fine imaging of celestial gamma ray sources. Milestones on its development are described in the paper by Reglero et al in this Seminar.

The aim of this paper is to summarise the results obtained during the INTEGRAL Masks development programme on implementing the HURA and MURA Codes on Tungsten plates of different thickness. Hard scientific requirements on pixels size and location tolerances (tenths of microns over large areas -1m2- and thickness from 0.5mm to 60mm) required the set up of a dedicated programme for testing cutting technologies: Laser, Photochemical Milling, Spark Machining and Electro Discharge Wire Cutting.

After a very intensive test campaign the Wire Cutting process was selected as the optimum technology for code manufacturing . Accuracies achieved on the code cutting fulfil scientific requirements. In fact, they are 5 times better than required. Pixel size and centroids location accuracies of 0.01mm over a 1m2 area have been obtained for the 10,000 pixels on IBIS, 100 pixels on SPI and 24000 pixels on JEM-X Masks.

Comparative results among different cutting technologies are also discussed. J_62 RM 103 V. Reglero et al.

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Keywords: Tungsten, Densimet, High Energy Astrophysics, imaging, INTEGRAL, Electrodischarge Wire Machining

1. The Coded Masks on INTEGRAL Satellite

The Astronomy and Space Science Group (GACE) of the University of Valencia was the responsible of the design and developement the Coded Aperture Systems (Coded Masks) for the ESA Mission INTEGRAL. The three high energy instruments on Satellite (Spectrometer SPI, Imager IBIS and X- Ray Monitor JEM-X) use Coded Masks to become imagers in the gamma ray domain. An overview of the INTEGRAL Coded Masks operation principle and performances is presented in Reglero et al in this Seminar.

The main function of the Coded Masks is to perform spatial multiplexing of the incident radiation. The Masks are basically matrixes of opaque and transparent pixels, following a bidimensional pattern, each of the holes acting as an individual pinhole camera. The image on the detector plane is therefore the convolution of the sky image and the code pattern. Each Mask use a different code pattern, optimised for the scientific goals of the instrument. The performance of the Masks acting as Spatial Multiplexing Systems depend to a great extent on the accuracy of the code elements position and size.

Tungsten (pure or alloyed) was the baseline material chosen to implement the codes because of both its high density and large atomic number which provides the required stopping power (opacity) in the gamma and X-ray energy ranges.

The main parameters for the Mask Code Patterns are:

IBIS Coded Mask:

• Material: Densimet 18 • Code pattern: MURA 53x53, four times repeated (symmetry 180°). See Figure 1a. • Pixel Size: Square, 11.2x11.2mm2 • Total Nr of Pixels: 95x95; 50% transparent / 50% opaque • Total Coded Area: 1064x1064mm2 • Code Thickness: 16mm V. Reglero et al. RM 103 163

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SPI Coded Mask:

• Material: Densimet 18 • Code pattern: HURA 127 elements; 64 transparent, 63 opaque (symmetry 120°). See Figure 1b. • Pixel Size: Hexagon, 60mm side to side • Total Coded Area: 780mm diameter • Code Thickness: 30mm

JEM-X Coded Mask:

• Material: pure Tungsten • Code pattern: HURA 22501 elements (symmetry 120°) • Pixel Size: Hexagon, 3.3mm side to side • Total Coded Area: 535mm diameter • Total Nr of Pixels: -23500; 25% transparent / 75% opaque • Code Thickness: 0.5mm

-JUv.-v--uU---J.JL

mjji ILÜ GliJ SI: !;•

nirOU'-C-J^^CJ^C1^

'"CMTJ'TU"!" "i n niinn ULULLil. •i • ! ; ä OJl

Figure 1a. IBIS Mask Code pattern Figure 1b. SPI Mask Coded pattern

For further information see (1)

Strong requirements were defined for the code parameters at the beginning of the project (1994) in order to fulfil the instruments scientific goals (angular resolution, source location accuracy and sensitivity). These requirements 164 RM 103 V. Reglero et al.

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define the performance of the Masks as Optical Systems, and refer mainly to two aspects:

• Position and size of the code elements, defined by the code manufacturing process and code integration into the Mask Assembly • Transparency/Opacity of the hole/opaque pixels in the instrument energy range. Deviations from its theoretical values (0/1) induced by the presence of support structures need to be characterizes in detail.

Tables I, II and III summarise the pixel position and size scientific requirements (deviations from theoretical values) altogether with the those imposed to the codes cutting process itself, before their integration in the Mask Assembly. Error budgets are quite demanding in order to achieve the required final precision for the full Mask Assembly.

The IBIS Mask Code was machined in four individual W 16mm thickness plates, with hard requirements in pixel position (+0.075mm) and size (0.1mm).

Mask Requirements Code Manuf. Requirements Position Size Position Size Y +0.15mm 0.1mm +0.075mm 0.1mm Z +0.15mm 0.1mm +0.075mm 0.1mm

Table I. IBIS Mask scientific and manufacturing requirements (pixel position and size)

The SPI Mask Code was machined in individual W 30mm thickness blocks (1, 2 or 17 hexagons each one). The manufacturing requirement (+0.025mm) refers to these blocks dimensions.

Mask Requirements Code Manuf. Requirements Total Position Total Size Total in plane +0.15mm +0.025mm

Table II. SPI Mask scientific and manufacturing requirements (pixel position and size)

The JEM-X Mask Code hexagons are machined in a very thin (0.5mm) Tungsten sheet. Centroid pixel position (+0.06mm) and pixel size (+0.01 mm) requirements are quite demanding for more than 5400 hexagonal holes. V. Reglero et al. RM 103 165

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Mask Requirements Code Manuf. Requirements Position Size Position Size Y +0.2mm 0.01mm +0.06mm 0.01mm Z +0.2mm 0.01mm +0.06mm 0.01mm

Table III. JEM-X Mask scientific and manufacturing requirements (pixel position and size)

Figure 2. JEM-X Mask code detail: "island"

2. Cutting Technologies.

One of the milestones of Coded Masks development carried out from 1995 to 2000 was to define the cutting optimum technology for the Code machining. The main constraint for the optimal cutting technology selection was the extremely high accuracies needed for the pixel position, size and geometry. As it is shown in Tables I, II and III, the manufacturing accuracies should be better than 0.1mm over Tungsten plates larger than 500mm, having several thousands of pixels. The standard machining processes for hard materials like Tungsten were not prepared to obtain these accuracies in such large areas.

Another challenging problem was the machining of the JEM-X Code on a 535mm diameter plate of only 0.5mm thickness. The code of JEM-X (-5400 hexagonal holes 3.3mm side to side) includes "islands" (opaque hexagons surrounded by transparent holes) joined to the plate by 0.4mm thickness ribs (see figure 2). The selected machining process must avoid the presence of 166 RM 103 V. Reglern et al.

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microcracks coming from the manufacturing process in a very fragile coded plate.

In all cases, after the manufacturing and assembly of the Coded Masks, they had to pass through a hard Qualification Test campaign, in order to prove their capability to survive INTEGRAL launch and operation environments, maintaining their scientific performances. These tests include Vibration tests (more than 12g inputs) and Thermal Vacuum and Cycling (from -60°C to 45°C).

Technologies tested were: LASER cutting, Photochemical Milling, Electrodischarge Milling and Electrodischarge Wire Machining. Main advantages and disadvantages of each method are summarised in the following paragraphs.

2.1. LASER Cutting.

BENEFITS: • Low cost • Good accuracy, regular hexagon shapes

PENALTIES • Material damage: cracks due to overheating • Material mechanical and thermal properties affected • Recasting

Figure 3. LASER cutting. JEM-X Code sample, hexagons with 0.4mm ribs. On the left, cracks on the ribs can be observed. On the centre and right pictures rests of materials along the ribs are clearly shown. Microphotographs from ICMUV facilities. V. Reglern et al. RM 103 167

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2.2. Photochemical Milling.

BENEFITS: • Low cost • Material properties not affected

PENALTIES: • Poor accuracy in pixel dimension and shape • W high resistance to chemical milling • Photoresin: difficult to resist the acid required for Tungsten.

Figure 4. Photochemical Milling. JEM-X Code sample, hexagons with 0.4mm ribs. A poor geometrical finishing is clearly seen in the three pictures. Inhomogeneities and shapes Microphotographs from ICMUV facilities.

2.3. Electrodischarge Milling.

BENEFITS: • Good accuracy, regular hexagon shapes • Minimum material damage

PENALTIES: • Large uncertainty in pixel position • Quick wearing of electrodes • Several electrodes required to comply accuracy requirements • Expensive 168 RM 103 V. Reglern et al.

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2.4. Electrodischarge Wire Machining (EDWM).

BENEFITS: • Splendid accuracy of hexagon shapes and positioning • Minimum material damage • Flexibility to improve pixel accuracy (multiple passes)

PENALTIES: • Expensive • Time machine consuming

Figure 5. Electrodischarge Wire Machining. JEM-X Code samples, hexagons with 0.4mm ribs. Good finishing in the pixel wall and pixel regular shapes are clearly seenn. V. Reglero et al. RM 103 169

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The EDWM was the technology selected to implement the Codes of the three INTEGRAL Masks. More information about the EDWM can be found in (2). The Code plates for the Coded Masks Development Models (DM), Qualification Models (QM) and Flight Models (FM) have been manufactured by this method, with increasing successful results in the Spanish company Mecanizados Gines.

Main parameters of the wire cutting are the following:

• Wire Diam: 0.25mm • Wire Material: Cu (Zn coating) • Voltage:-80/-100V • Cutting speed: 2.2 to 2.5mm/min • Dielectric: Distilled water • Nr. of cutting hours (one JEM-X Code): -500

Figure 6. Cutting of the JEM-X code by Electrodischarge wire machining. More than 5000 hexagonal holes 3.3mm side to side. 170 RM 103 V. Reglero et al.

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3. FM Masks Results on Pixel Size and Position.

The results on pixel size and position obtained in the implementation of the INTEGRAL FM Masks codes are presented hereafter, compared with both scientific and manufacturing requirements.

3.1. IBIS Mask Results.

a) Plates Manufacturing (all pixels measured)

Required Std Dev(a) Max Dev Pixel Position Y +0.075mm 0.010mm 0.030mm Z +0.075mm 0.015mm 0.040mm Pixel Size Y 0.1mm 0.010mm 0.060mm Z 0.1mm 0.010mm 0.060mm

Table IV. IBIS Mask pixel position and size results, after code machining.

b) Mask Assembly (16 control pixels measured)

Required Std Dev(a) Max Dev Pixel Position Y +0.150mm 0.030mm 0.040mm Z +0.150mm 0.030mm 0.050mm

Table V. IBIS Mask pixel position and size results, after Mask assembly.

Detailed IBIS Mask results can be found in (3) and (5). Main conclusions can be summarised as follows:

• The results obtained are 5 times better than required. • Stability of the IBIS Mask as an optical system: no significant deviations have been found before and after environmental tests at INTA V. Reglero et al. RM 103 171

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Figure 7. IBIS Code plate cutting by Electrodischarge Wire Machining.

3.2. SPI Masks Results.

a) Plates Manufacturing (all pixels measured)

All Code blocks (1, 2 & 17 hexagons) are within manufacturing tolerances (+0.025mm).

b) Mask Assembly (all pixels measured)

Required Std Dev(a) Max Dev Total in plane +0.150mm 0.020mm 0.080mm Pixel Position

Table VI. SPI Mask pixel position and size results, after Mask assembly.

Detailed results of SPI Mask scientific performances can be found in (4) and (6). Main conclusions can be summarised as follows:

• The results obtained are 7 times better than the scientific requirements • Stability of SPI Mask as an optical system: no significant deviations found before and after environmental tests at INTA 172 RM 103 V. Reglero et al.

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Figure 8. Dimensional Control of the SPI FM Coded Mask

3.3. JEM-X Mask Results.

a) Plates Manufacturing (200+200 pixels measured)

Required Std Dev(cr) Max Dev Pixel Position Y +0.060mm 0.015mm 0.030mm Z +0.060mm 0.015mm 0.030mm Pixel Size Y 0.01mm 0.005mm 0.010mm Z 0.01mm 0.005mm 0.010mm

Table VII. JEM-X Mask pixel position and size results, after code machining,

b) Mask Assembly (100+100 pixels measured)

Required Std Dev(c) Max Dev Pixel Position Y +0.2mm 0.025mm 0.060mm Z +0.2mm 0.025mm 0.050mm

Table VIII. IBIS Mask pixel position and size results, after Mask assembly V. Reglero et al. RM 103 173

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Detailed JEM-X Mask results can be found in (7) and (8). Main conclusions can be summarised as follows:

• The results obtained are 8 times better than the scientific requirements • Stability of JEM-X FM Masks as optical systems: no significant deviations found before and after tests

Figure 9. Dimensional Control of the JEM-X FM code plates after their machining.

4. Conclusions.

An extensive programme on testing technologies for high precision Tungsten cutting in large areas was carried out from 1995 to 1998 by the Spanish team involved in the INTEGRAL ESA Mission (GACE, SENER, INTA and Mecanizados Gines): LASER, Photochemical Milling, Spark Machining and Electrodischarge Wire Machining. The EDWM was selected as the optimum technology.

Qualification and Flight Models for the four Mask Assemblies (SPI, IBIS and JEM-X) demonstrate the high accuracy achieved in the code cutting and geometry. Code elements size and shape fulfil scientific requirements. In fact they are 5 times better than specified. _174 RM 103 V. Reglero et al.

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Qualification and Acceptance Tests carried out at INTA proved the stability of the code positioning under extreme conditions (12g input acceleration; -60°C to +45°C), exceeding requirements by a factor 8. EDWM has demonstrated its capability to produce fine Tungsten plate cutting to implement new Optical Systems for High Energy Astrophysics.

Acknowledgements.

We would like to thank more than 14 European industries that have participated in the Masks development, with special mention to our Prime Contractor SENER, test facilities team at INTA, our Tungsten supplier PLANSEE and its representative in SPAIN Technalloy.

References.

(1) Reglero V, Sanchez F, Perez F, Sanchez F, M, Collado V, Peris J, Donate M, Velasco T, Navarro A, Sabau D, Reina M, Sanchez A, Eiriz E, Urteaga J.M, Sebastian P, Mayo L, Beiztiegui R, INTEGRAL Signal Multiplexing, 1999, Astrophysical Letters and Comm, Vol 38, p 381. (2) Fuller, John E, Nontraditional Machining Processes. Electrical Discharge Machining, 1984, E.J.Weller Ed, Society of Manufacturing Engineers (3) Reglero V, Sanchez F, Rodrigo J, Velasco T, Gasent J.L, Chato R, Alamo J, Burgos J.A, Suso J, Blay P, Martinez S, Donate M, Reina M, Sabau D, Ruiz-Urien I, Santos I, Zarauz J, Vazquez J, IBIS Mask Pre-Calibration Matrix, 2001, ESA SP series, in press. (4) Sanchez F, Reglero V, Chato R, Rodrigo J, Velasco T, Gasent J.L, Alamo J, Burgos J.A, Suso J, Blay P, Martinez S, Donate M, Determination of SPI Coded Mask Transmission Matrix, 2001, ESA SP series, in press. (5) Velasco T, Gasent J.L, Zarauz J, Reglero V, IBIS Mask FM Dimensional Control Report, 2000, IN-IB-INT-RPT-012 (6) Velasco T, Reglero V, Zarauz J, Sanchez F, SPI Mask FM Dimensional Control Report, 2000, IN-SP-INT-RPT-021 (7) Velasco T, Reglero V, Zarauz , JEM-X Mask FM1 Alignment and Positioning Check Report, 2000, IN-JX-INT-RPT-015 (8) Velasco T, Reglero V, Zarauz , JEM-X Mask FM2 Alignment and Positioning Check Report, 2000, IN-JX-INT-RPT-021 AT0200476 R.M. German et al. GT4 175

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Unique Opportunities in Powder Injection Molding of Refractory and Hard Materials

Randall M. German Brush Chair Professor in Materials Director, Center for Innovative Sintered Products P/M Lab, 147 Research Building West The Pennsylvania State University University Park, PA 16802-6809 USA

Summary:

Powder injection molding (PIM) is a relatively new manufacturing process for the creation of complicated net-shapes outside the range usually possible via powder metallurgy technologies. This new process is now in production at more than 550 sites around the world. Although a small industry, PIM will soon pass SI billion dollars (USA) in annual sales. This presentation overviews the PIM process, some of the new developments and some of the successes that have occurred with both refractory metals and hard metals. Example applications are seen in medical and dental devices, industrial components, wristwatches, jet engines, firearms, automotive components, and even hand tools. To help establish the novel growth opportunities, PIM is compared to other fabrication routes to better understand the design features arising with this new approach, providing a compelling case for substantial opportunities in the refractory and hard materials. Illustrations are provided of several components in production. New opportunities abound for the technology, since it eliminates the shape complexity barrier associated with die compaction and the cost of machining associated with complicated or dimensionally precise components. Further, a relative cost advantage exists for refractory and hard materials because PIM can use the same powders at the same prices as employed in alternative processes. Future successes will occur by early identification of candidate materials and designs. Early examples include tungsten heavy alloy components now reaching production rates of six million per month.

Keywords:

Powder injection molding, design, net-shaping, new processes, innovative processing, sintering

1. Introduction:

The concept of mixing a metal powder with a thermoplastic binder to enable shaping by injection molding prior to sintering has been around since the 1920's. Today, powder injection molding (PIM) is 176 GT4 R.M. German et al.

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gaining widespread popularity due to the combination of net-shaping and performance [1], The plastic molding process is ideal for fabrication of large quantities of the same shape, so PIM is simply building from the enormous technological base created by plastics. However, for most metallic systems the powder cost is ten times that of a casting material. Thus, for steel and stainless steel, PIM can only be cost competitive for very complicated and low weight structures where machining proves difficult. However, for the refractory and hard materials, this material cost disadvantage is less evident. Consequently, many new products and shapes are possible, in large production quantities, without a cost penalty via PIM. Accordingly, several applications have been identified and moved into pilot or full production. A large body of literature is available detailing the processing, properties, economics, and applications [see literature reviews in references 1 and 2]. Example systems covered by the prior publications include tungsten, tungsten heavy alloys, tungsten-copper, rhenium, molybdenum, molybdenum-copper, molybdenum disilicide, niobium-base alloys, titanium, titanium alloys, titanium aluminides, and several other important refractory systems. It is important to recognize that PIM provides a double cost benefit - * no powder cost penalty since these materials are already formed from small powders * reduced losses and easy recycling increase the process advantage with costly materials.

Thus, it is appropriate to consider what might be some targets for future conversions to PIM. This paper identifies the current systems and further illustrates some emerging opportunities in military, aerospace, sporting, and other devices including cellular telephones and computer disk drives. Today, we recognize PIM as one of many shaping technologies addressing net-shape production of engineered components. Unlike casting (limited to metals), it is applicable to a range of materials. One of the future gains will be in mixed powder systems, where ceramic and metal powders can be tailored for properties by adjusting particle size and composition prior to molding. This will provide a differentiation in performance with respect to most fusion techniques. Further, when properly executed, PIM is financially and technologically competitive with alternative processes for net-shape production. It provides good dimensional control and is more cost effect than machining and grinding, Thus, PIM has emerged as a technique for large quantity net-shape production of discrete, and complex engineered components with dimensional tolerances that are not tight. It builds a unique technology based on a combination of powder metallurgy and plastic injection molding.

Powder injection molding has been available in various forms for several years, yet significant success was delayed until the 1990's. A few factors contributed to the long incubation period, including basic scientific understanding and a need for the full infrastructure to mature (furnaces, binders, debinding equipment, mixers, and molders). Today, the technology has passed through most conceptual barriers and is widely accepted in several application areas. The cause for the long incubation period included the following factors:

* poor financial performance; profitability was nonexistent * basic process knowledge and problem solution protocols were missing * early variants depended on highly proprietary technologies without production volume raw materials were costly and in limited supply * trained workers were missing and no instructional programs existed R.M. German et al. GT4 177

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* early equipment was built from scratch and only later was equipment customized for PIM * designers were uncomfortable because of uncertain design rules and material properties * the field lacked needed case histories of success with PIM * early equipment and operations lacked required process controls * growth in microelectronics allowed manufacturing instrumentation at low costs * without an infrastructure there was excessive cost to covering the range of technologies * optimization was impossible, since the scientific basis for PIM was missing.

Today the successes in PIM trace to several recognized benefits, including net-shaping and the avoidance of machining and other steps. Further, the products are sintered to full density, providing high performance. A wide range of compositions are now available, essentially all of the known engineering metals and ceramics have been demonstrated, if not put into production. An important aspect it the product cost, especially for complicated shapes. In part the low cost comes from the large existing base in plastic molding. Further, new software in molding and sintering simulations enable faster optimization - indeed growth is highly accelerated by the fact that PIM is piggybacking on the large field of plastic injection molding, but using powders and sintering cycles well known in powder metallurgy.

Production today ranges from a few to millions. As long as a fine, sinterable powder is available PIM is successful from a technological view. Financial considerations might not be as favorable, but technology is done and over 500 firms now practice the technology. Parts in production range from 0.03 g up to 17 kg, with production volumes as low as 10,000 per year up to 72 million per year. Most recently the technology has undergone simplification, with the emergence of a three-step process. The fist step is the formation of a powder-polymer mixture known as a feedstock. There are now over 10 commercial feedstock suppliers and the bulk of the newer entries are moving to premixed feedstock during the evaluation and start-up phase. Molding takes place in classic plastic injection molders, so many of the actors have come from a plastics background. Finally, sintering is being performed in furnaces well known to powder metallurgy - both batch and continuous furnaces. To assist the plastic molders, a few toll-sintering operations have been established to provide early support, in the same manner as many metallurgical operations rely on outside heat-treating. The key decision these days is in the binder, recognizing the debinding process is largely dictated by the binder attributes. Wax- polymer mixtures are most common, with over 55% of the industry relying on these simple systems. The polymer is simply extracted by slow heating, the same way cemented carbides are dewaxed early in a sintering cycle. Some operations separate the debinding and sintering steps, but others have been successful using combined thermal treatments.

As the process becomes simplified and knowledge is widely distributed, sales for PIM products have accelerated and now are approaching $ 1 billion worldwide. Today, over 6200 people are employed in PIM, with 37% of the products for internal use and 63% for external customers. About 53% of the sales are in North America. Curiously, most of the industry is still relatively small, in the pilot stage; hence, 5% of the firms control 42% of the production capacity, 63% of the sales, and 81% of the profits. Those operations that have reached full production average 85 employees, 12 molders, 5 furnaces, 2 178 GT4 R.M. German et al.

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mixers, and product 100 tons of parts per year worth $14.4 million. The average weight is near 5 g and the average sale price is near $0.80 each, making this a high profit field for may applications. As is evident, these are small components, but with high value. Continued expansion at a 20% or faster rate is expected for the next several years. Some of the reasons for this expansion lie in the shapes, performance, and cost benefits. These are best seen via the illustrations that follow.

2. Applications in Computers and Microelectronics:

One of the most exciting opportunities is associated with the explosive growth in computers and microelectronics. PIM has participated in many of the recent surges in consumption, first in the hard disk drive, and later in cellular telephones. Indeed, today some of the largest uses of PIM are associated with computers, microelectronics, and other electronic devices - as weights, heat dissipators, electromagnetic shields, magnets, motor components, vibrators, and balancing devices where a few examples are shown in Figure 1. In the cellular telephone, some of the simple devices have reached production volumes of several million parts per month with similar production quantities associated with computer hard disk drive weights, latches, counterbalances, and magnets. The magnets are typically formed from martensitic or ferritic stainless steels, but there is considerable us of tungsten in many of the computer and cellular telephone devices.

Figure 1. Example hard disk drive components, courtesy of Advanced Materials Technologies. R.M. German et al. GT4 179

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3. Applications in Military Hardware

Tungsten penetrators are a favorite for anti-armor applications. PIM has been used in the synthesis of new alloys and is at that demonstration level. Early production of tungsten ammunition took place in the late 1980's, with successful testing in anti-aircraft and other applications by the early 1990's. Most of these were tungsten heavy alloys. Today, more complicated geometries and alloys are being fielded for tungsten military rounds, such as shown in Figure 2. An interesting aspect of PIM is the ability to put features into the shape without machining, and in this case a functional gradient in tungsten content exists in the wall thickness.

Recent success has been in PIM tungsten-nylon compositions for frangible ammunition or practice bullets to replace lead. Note these do not require sintering and are used in the as-molded condition. Further, several compositions tailored to ballistic are now making progress in the market place, including some success with tin as the cement for tungsten. Initially such compositions were designed for use around nuclear facilities, but now the variants on the tungsten-tin compositions are being applied to a wide range of military, and soon hunting applications. Another derivative technology is based on shape charge and explosively formed projectiles. The first applications are in oil well drilling technologies using tungsten-lead compositions.

Figure 2. A military projectile with functional gradient tungsten content in the wall thickness as a possible high volume PIM product.

4. Applications in Petrochemical Processing

PIM has been in use for selected shapes and typically in caps for drilling teeth off and on for about 10 years. An example drilling tooth is illustrated in Figure 3. Most of these shapes are formed from cemented carbide compositions. However, an opportunity exists in generating gradient microstructures via progressive molding of differing compositions. This represents a significant new production area for PIM.

Figure 3. An example tooth for drilling, which is usually coated by cemented carbide, but could be fabricated as a functionally graded structure by molding layers over each other. JJ30 GT4 R.M. German et al.

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5. Applications in Microelectronic Packaging

Demonstration compositions for microelectronic packaging have been around since the late 1980's, and patents on the concept date to the early 1980's. Some high volume and exciting opportunities exist in these areas, and heat dissipation, electrical conductivity, and low thermal fatigue failure (thermal expansion matching) are clear dictates. Most of the early successes have been based on W-Cu compositions, where copper provides the thermal conductivity and the interconnected tungsten skeleton provides the low thermal expansion coefficient. Although very successful in commercial uses, these materials are still heavy and expensive. Use of Mo-Cu compositions help lower the density, but future portable computer systems are pushing for new lower density and lower cost materials. Hence, PIM of W-Cu and Mo-Cu will likely be associated with stationary devices, and composites in aluminum and copper or even aluminum nitride will be more typical in portable devices. Even so, the applications in computers, workstations, cellular telephone base stations, digital light projectors, and related applications are enormous. Today, most of the successful products are delivering thermal conductivities over 200 W/m/K. Pictured in Figure 4 is an new hybrid microelectronic package, where the tungsten- copper base (microstructure shown to the right) is sinter bonded to a kovar glass-sealing housing.

0 an) " \; * _ -

Figure 4. An advanced microelectronic package consisting of a W-Cu base for heat dissipation and kovar side wall for glass sealing. The composite microstructure of the sintered W-Cu base is illustrated on the right (package courtesy of Advanced Materials Technologies, micrograph courtesy of John Johnson).

6. Applications in Sporting Equipment

Much notice has been directed to the possible use of heavy metal weights in sporting equipment - golf clubs, darts, arrows, and even ammunition. Indeed, one of the first successful applications for PIM was R.M. German et al. GT4 181

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in running shoes with a Ti-Fe composite that provided strength and durability with low density. Today, many sporting items are under investigation. Golf club heads and weights for gold clubs are some of the most exciting products. Most of the tungsten applications are geared to selective counter balances, control of torque, and relocation of the center of gravity during ball strike. An example where a higher density alloy is used to control torque is illustrated in Figure 5, with a low-density center and high density (black) outer region. Production quantities for some of the weights (most not by PIM) have reached 30 million per year. However, there is much frustration for the industry in working on these applications. Truly golf clubs are a fashion industry. Each year a new style is needed, so the product lifetime is often short, while the time to become qualified as a vendor might be long. This requires a new thinking in the PIM industry, where short-intense production is the key concept. One must then always be on the lookout for the next items to enter production; otherwise a large production facility sits idle. Sporting equipment production by PIM is not for the faint of heart. Characteristically, the sporting industry will not enter production itself, since it is easier to move between vendors as fashion changes.

Figure 5. A controlled density distribution golf club head that allows torque control by selective placement of higher density weights on the heel and toe (courtesy of Never Compromise).

7. Titanium and Titanium Alloys:

There is a love affair with titanium and this projects into PIM as well. Many parts are in production from this refractory material, and the mechanical properties are becoming attractive. For example, pure titanium PIM parts are now in the 500 to 630 MPa tensile strength range with 3 to 20% elongation, depending on the powder purity. Alloys such as Ti-6A1-4V are also in production, delivering strengths near 880 MPa and 12% elongation. Applications for titanium PIM are diverse, including golf clubs, automotive items, decorative hardware, and surgical tools, watchcases, watchbands, toys, and industrial components. 182 GT4 R.M. German et al.

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8. Decorative Items:

Watches, jewelry, and watchbands are common now by PIM. The materials range from gold to cemented carbides, and include colored zirconia, titanium, and stainless steels. Other decorative items range from cosmetic cases, wine stoppers, and tokens. Figure 6 is an example of a scratch resistant PIM WC-Co watchcase.

Figure 6. A PIM WC-Co watchcase, courtesy ofRado.

9. Industrial and Metalcutting Tools:

Tools that see abrasion and wear are now common by PIM. These included many cemented carbides, used for drills, cutting inserts, sand blast nozzles, and other high stress, wear, or abrasion situations. Figure 7 shows example components formed by PIM from WC-Co for double sided cutting inserts with chip breakers and internally cooled drills. The former are now estimated at 1% of all cutting inserts by PIM and the latter have achieved a record level of tolerance control - five micrometers dimensional R.M. German et al. GT4 183

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scatter. Many other opportunities exist for net-shape production of complicated cutting tools via PIM, with some interesting possibilities in functional gradients, selective placements of hard phases, and even new shapes.

Figure 7. Examples of PIM WC-Co tools for metal cutting and drilling.

10. Some Future Products:

The future is bright, with many interesting refractory metal and hard material products in the pipeline. Several options were mentioned already, with special focus on shape complexity and cost reduction. Others in the lab include new rocket nozzles, porous electrodes, advanced lighting designs, microelectronic packages, thermal management materials, and porous structures. Many options are under discussion, including new welding electrodes, lighting devices, oil well drilling tips, microelectronic packages, hand tools, hardware, sporting devices, and even conformal refractory metal capacitors. Besides the new shapes, some new materials and functionality will be possible via PIM. These include the construction of hollow components and even the fabrication of controlled porosity or foamed devices. Gradient microstructures are being formulated using two-color molding, where two different materials are injected to form a single device. Most interesting is the use of plastic inserts in the die cavity, allowing burnout of the plastic during sintering to form complicated internal cavities in the structure - a route very useful for rocket nozzles, sand blast nozzles, spraying nozzles, and other fluid control structures. Thus, the future is bright, the options far reaching, and the major barrier is time to chase all of the opportunities. In that regard, 100 companies now sponsor Perm State for its 184 GT4 R.M. German et al.

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Center for Innovative Sintered Products to leverage time and funds to seek out exciting new opportunities.

11. Simple Design Rules:

Success in powder injection molding starts by identification of an application where the production volumes justify the tooling and set-up costs. Figure 8 shows the basic process and key steps with the illustration of one valve component at each stage. Generally, if a shape is simple and can be produced by die compaction, then the PIM process will not be competitive. On the other hand, if die compaction is unable to deliver the shape complexity and sophistication in features, tolerances, and design, then PIM becomes attractive. Selection of candidates pivots on these attributes,

• tolerances - must justify PIM and be compatible with typical production levels • production - quantities over 5000 per year are needed to amortize tooling, and production volumes up to nearly 100,000,000 per year occur with items such as cell telephone vibration weights from tungsten • performance - PIM tends to be better suited when higher performance is desired to better justify the cost • shape complexity - the shape must be beyond the capabilities of die compaction.

When such a situation arises, then PIM production becomes viable. To enhance success with this new form of powder metallurgy, we can go further and isolate some design features that encourage success and ease processing difficulties. The first consideration is to start thinking "plastics" versus "powder metallurgy" by looking at the component in terms of die layout, flow patterns, ejection, gating, and venting [1]. Further, to minimize mass, cost, and processing times, wall thickness should be minimized whenever possible. This suggests enthusiastic use of cores, indents, holes, and other features such as undercuts and even hollow components. Powder cost is a major factor in refractory metal and hard material components, so small tool changes that propagate mass reductions over many components are highly valuable. Another consideration in the component design stage is to consider where parting lines, ejectors, and gates can be placed on the component. These surface blemishes need to be located in non-critical regions to avoid the expense of removal after molding. One new option is to use porous tooling where air pressure is used for ejection after molding. Likewise, ejection from tooling is enhanced by incorporation of a slight draft or taper. Otherwise, special and expensive features are needed to open the tooling for part extraction.

With experience, a designer will learn that other features will greatly assist the production process. Examples are to be sure to include a flat surface on the component to allow for staging during debinding and sintering. Otherwise, special conformal setters need to be fabricated out of ceramic to hold each part. Not that such an option is difficult, but it adds considerable expense onto the production cost. Also, the debinding and sintering stresses acting on the component depend on section thickness. Very thick and very massive components are in production, up to 17 kg and 125 mm thick, but cycle times in molding, debinding, and sintering are slow to accommodate heat transfer and mass flow from R.M. German et a). GT4 185

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thick sections (for example the 17 kg component has a molding cycle time of 5 min, while more typical to PIM is 15 s). Accordingly, mass reduction and uniform wall thickness pay off in faster processing and more uniform material response with better tolerances, less warpage or distortion or cracking. Finally, there is the perpetual issue of tolerances. In general, all of the newer PIM operations with closed-loop feedback control systems and good feedstock homogeneity can sustain tolerances on dimensions near 0.1%. Tighter tolerances are possible, but require more attention and result in higher costs. Recent benchmarks in the industry are tolerances measured in the micrometer range on certain critical dimensions. But this is not typical and not inexpensive. New production operations focused on optical components require such attributes, but for most applications these are neither typical nor required.

12. Acknowledgements:

This brief introduction and overview was assisted by several companies and individuals, with special thanks to my friends Julian Thomas (Center for Innovative Sintered Products at Perm State), Robert Cornwall (Innovative Material Solutions), Kay Leong Lim (Advanced Materials Technologies), Zak Fang (Smith International), Ruven Porat (Iscar), Brian Pond (Never Compromise), and John Johnson (Howmet).

13. References:

1. R. M. German and A. Bose: Injection Molding of Metals and Ceramics, Metal Powder Industries Federation, Princeton, NJ, 1997. 2. R. M. German and R. G. Cornwall: Powder Injection Molding in the Year 2000 an Industry and Market Report, Innovative Material Solutions, State College, PA, 2000. en

Figure 8. A view of the powder injection molding (PIM) which starts by mixing a metal poweder and plastic to form feedstock, which is subsequently granulated and feed into a molding machine. In molding the thermoplastic mixture is heated and rammed into a cold die. The die is constructed to provide shape to the feedstock. On cooling, the shaped mixture is ejected and then the polymer is extracted by heat or other means, a process known as debinding. Finally, the powder structure is sintered, bringing the powder system to full density. These illustrations of the equipment and process steps for the production of a valve body are kindly provided by Advanced Materials Technologies.

CD 3

2- AT0200477 JA Shields et al. GT 8 187

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Fracture Toughness of 6.4 mm (0.25 inch) Arc-Cast Molybdenum and Molybdenum-TZM Plate at Room Temperature and 300°C

J.A. Shields, Jr.*; P. Lipetzky+; A.J. Mueller*

*CSM Industries, Inc., Cleveland, Ohio, USA ""Pittsburgh Materials Technology, Inc., Jefferson Hills, Pennsylvania, USA #Bechtel Bettis Inc., West Mifflin, Pennsylvania, USA

Summary:

The fracture toughness of 6.4 mm (0.25 inch) low carbon arc-cast (LCAC) molybdenum and arc-cast molybdenum-TZM alloy plate were measured at room temperature and 300°C using compact tension specimens. The effect of crack plane orientation (longitudinal vs. transverse) and annealing practice (stress- relieved vs. recrystallized) were evaluated. Depending upon the test temperature either a standard K|C or a J-integral analysis was used to obtain the toughness value. At room temperature, regardless of alloy, orientation, or microstructure, fracture toughness values between 15 and 22 MPa mA (14 and 1/2 20 ksi in ) were measured. These K,c values were consistent with measurements by other authors. Increasing temperature improves the toughness, due to the fact that one takes advantage of the ductile-brittle transition behavior of molybdenum. At 300°C, the fracture toughness of recrystallized LCAC and arc-cast TZM molybdenum were also similar at approximately 64 MPa m72 (58 ksi in72). In the stress-relieved condition, however, the toughness of arc-cast TZM (91 MPa m'/2 / 83 ksi in7*) was higher than that of the LCAC molybdenum (74 MPa m7* / 67 ksi in72).

Keywords: molybdenum, TZM, refractory metals, fracture toughness, K|Cl J Integral 188 GT_8 J.A. Shields et al.

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1. Introduction:

Background: Brittle failure analysis has become a common consideration in the design of many structures. As a result, fracture toughness properties necessary to conduct brittle fracture analysis of many common materials such as iron, aluminum and titanium alloys have been well characterized. Relatively little effort has been devoted to the characterization of the fracture toughness of molybdenum and molybdenum alloys. Because molybdenum's ductile-to-brittle transition temperature (DBTT) is typically within approximately 100°C of room temperature, fracture - critical components have not typically used molybdenum in this temperature range. Electronic heat sinks, for example, rely on the thermal expansion and conductivity of the material, but do not include fracture considerations in component design. Likewise, the high-temperature applications for molybdenum and its alloys (heat treating furnaces, glass-melting furnace components, e.g.), concern themselves more with long-term creep behavior than with fracture resistance.

In order, for molybdenum and molybdenum alloys to be considered for applications that must exhibit fracture resistance, it is necessary to characterize the toughness of the material. The temperature range between ambient and 1200°C forms the region of most interest with respect to toughness. It is in this region that components such as forging dies, metalworking tooling, and furnace structural parts must operate, and in this range where additional characterization of fracture toughness must be performed.

Prior Work: A few authors have investigated the toughness of molybdenum and molybdenum alloys. Measurements of toughness values were performed on pre-cracked bend bars of extruded arc-cast molybdenum as early as 1963 (1). These measurements used thick cross-sections, and calculated the energy of fracture. Other early work investigated sheet molybdenum and TZM alloy (2), employing thin sheet specimens that frequently displayed plasticity. The state- of-the-art at this time did not permit an appropriate stress-analysis solution for such tests. Furthermore the authors were unable to fatigue pre-crack the pure molybdenum samples without fracture, and relied upon Electro-Discharge Machining (EDM) to produce notches in the specimens. More recently the fracture toughness of molybdenum rod was measured (3). These investigators solved the pre-cracking problem by using compression - compression fatigue, and obtained K,c values for the specimen geometries with which they worked. J.A. Shields et al. GT8 189_

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In the preceding work, toughness measurements were obtained at ambient temperatures. Russian investigators have measured the toughness of molybdenum to temperatures of about 500°C (4,5), as have German investigators (6). However, there is no data on the fracture toughness of molybdenum or molybdenum alloys covering temperatures above 500°C, and very little well-characterized data reported between ambient and 500°C.

The results of measurements performed in the preceding references are summarized in Table 1. The toughness numbers vary widely probably due to limitations of the tests, specimen geometries employed, and analytical techniques available at the time of the individual investigations. The most reliable values are probably those of Romine (1), Chen, et. al. (3), and Rödig, et. al. (6).

Current Tests: The testing and analytical tools available to the experimenter today are much improved over those of 20-30 years ago. Analytical techniques like the J-lntegral approach (7) have allowed extension of the art into the realm of tests displaying plasticity. The pre-cracking techniques of Chen, et. al. (3) and Rödig, et. al. (6) permit the introduction of sharp notches into less ductile materials like molybdenum. Furthermore, there is increased interest in the fracture behavior of molybdenum and its alloys, particularly at elevated temperatures where the materials are on the upper shelf of their transition curve.

For these reasons, we have set out to measure the fracture toughness of molybdenum and TZM plate at temperatures up to 1100°C using improved techniques. This paper covers the initial measurements made at ambient temperature and at 300°C. 190 GT8 J.A. Shields et al.

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Table 1. K|C measurements of molybdenum and molybdenum alloys at ambient temperatures.

HRe Product Specimen Crack K,c, f Form Type Dir. Material Mi era structure MPa m1'2 1 Extrusion, Mid-rad. NB ? A/CMo Extruded 17.6a 1 Extrusion, OD NB ? A/CMo Extruded 20.6D 2 Sheet CN Trans. A/CMo SR 35 - 65C Sheet CN Trans. A/CMo RX 27 - 33° 2 Sheet CC Trans. A/CTZM SR 40d 3 Forged bar CT Radial P/M Mo : Forged 8 + 5% 3 Forged bar CT Radial P/M TZM Forged 19 ±5% 3 I Forged bar SC Radial P/M Mo Forged 17e 3 Forged bar SC Radial P/M TZM Forged 17 3 Swaged bar SC Radial P/M Mo Swaged 17 ±11%r 3 j Swaged bar SC Radial P/M Mo RX 18 ± 10%T 3 Swaged bar sc Radial P/M TZM RX 31 + 7%' 4 ! Sheet CT Trans. Mo-1%Zr RX 15-17 5 Sheet ? IRolling Mo ? 15-60 5 Sheet ? In-plane Mo ? 10-40 6 Bar DCT Radial P/M TZM SR 15-17 Key: NB = notched bend; CN = center notched; CC = center cracked; CT = compact tension; SC = surface crack; DCT = dia. compact tension; ? = not reported; SR = stress relieved; RX= recrystallized 3 C Average of 5 tests, calculated from Gic; "Average of 11 tests, calculated from Gic; KO value. Significant plasticity, Kj calculated at deviation of elastic line from linearity. eSpecimen displayed sub-critical crack growth, calc at initiation of crack growth. 'Specimen displayed sub-critical crack growth, calc using the deepest point of semi-elliptical crack

2. Material Description:

The alloys used in this investigation were consolidated from powder precursors using the Press-Sinter-Melt (PSM) process developed in the 1950s (8). In this process, powder is fed from a hopper enclosed in the vacuum melting furnace, into a pressing die where it is pressed into a disc and fed down into the melting chamber. Repeating this process allows an electrode to be built up which is melted in vacuo by an AC arc and solidified in a water-cooled copper crucible. J.A. Shields et al. GT8 191

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The cast ingot is then extruded at elevated temperature to break up the large grain structure associated with arc melting. Analytical samples are obtained after extrusion, and Table 2 summarizes the chemical analysis of the two materials. Interstitial elements (C, N, and O) were determined by LECO inert gas fusion and trace elements determined by DC spark/optical emission spectroscopy. The rectangular extrusion is then sectioned and recrystallized before hot rolling to final dimensions.

Table 2. Chemical analysis of molybdenum plates in weight percent. Material c o N Ti Zr Fe Ni Si Mo Ingot 50839 .005 <.001 <.001 - - .003 <.001 <.001 ASTM B386 £.010 £.0015 £.002 - ~ £.010 £.002 £.010 (365) TZM Ingot 61756 .021 .0008 .0009 .48 .104 <.001 <.001 <.001 ASTM B386 .01-.03 £.0030 <.OO2 .40-.55 .06-.12 £.010 £.002 £.010 (363)

The material used in this investigation was rolled to 6.4 mm (0.25 inch) thick, and subjected to either a final stress-relief or a final recrystallization anneal. The rolling direction was perpendicular to the extrusion direction. Following the final heat treatment, the surface was etched to remove the resulting decarburized zone (~ 25jxm, 1 mil). Table 3 summarizes the average transverse tensile properties measured on each plate. The microstructures of the various material heat treatment combinations are shown in Figure 1. The tensile

Table 3. Mean transverse tensile properties of molybdenum plates at 25°C.

O-y, Elongation Material Structure No. of Tests MPa (ksi) MPa (ksi) Al/lo, % Mo Ingot 50839 SR 2 712(103) 778 (113) 21.0 Mo Ingot 50839 RX 3 377 (55) 498 (72) 52.8 TZM Ingot 61756 SR 2 846(123) 965 (140) 15.3 TZM Ingot 61756 RX 1 476 (69) 623 (90) 50.9 192 GT8 J.A. Shields et al.

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Longitudinal Transverse

•a »^»»^

CO to

4 V N ^f1 "^ S

O s. a:

r -^ * t 1 "^ co CO ^-=--— N #

co 8 P -a < |

Ü a:a?

Figure 1. Representative Optical Photomicrographs of LCAC and arc-cast TZM molybdenum; Murakami's etch. | 1 150>m J.A. Shields et al. GT8 193

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properties and microstructures are typical of the two materials in the conditions obtained. Additionally, the mechanical properties exceed ASTM B386 minimum requirements for yield strength and ductility.

3. Experimental Procedures:

Sample Machining: Square compact tension (CT) specimens were machined from the 6.4 mm (0.25 inch) thick plate of LCAC molybdenum and arc-cast TZM molybdenum using EDM to provide both longitudinal and transverse notch orientations. The specimen design is shown in Figure 2. The notch width made using a 127 urn (5 mil) diameter EDM wire was consistently between 177 and 254 urn (7 and 10 mils) in width and 10 mm (0.400 inches) past the load line of the sample. The surface finish was according to ASTM E-399.

05,08

10,16

Figure 2. Sketch of 6.4 mm (0.25 inch) thick compact tension specimen design (dimensions are shown in mm).

Cyclic Pre-crackinq: A cyclic compressive load was used for initiating the pre- crack (3). Benefits of this method include pre-crack stability, symmetry and low risk of fracture for brittle specimens. Specimens were cycled 2500 times and visually inspected before the maximum load was increased. Atypical loading sequence was 2500 cycles (20 to 40Hz, R=0.1) at 455 kg (1000 lbs) maximum load, followed by 2500 cycles at 909 kg (2000 lbs) maximum load etc. Cycling was continued in this manner until a crack was observed from some point at the notch root. If a crack was visible and growing, no further increase in load was 194 GT8 J.A. Shields et al.

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applied until crack advance stopped. Using this method, specimens were pre- cracked at maximum loads less than 4545 kg (10,000 lbs) and total cycle count below 40,000.

Test Method: Room temperature toughness testing followed ASTM E-399 (7) guidelines wherever possible. Even though the pre-cracking method described above does not meet E-399 requirements, the testing yielded meaningful results by following the other guidelines as highlighted below:

• Loading rate 0.55 to 2.75 MPa my7sec (30 to 150 ksi iny7min). • Pre-crack length must be a minimum of 1.2 mm (50 mils) past the notch root. • Pre-crack must be within 10° of the notch plane. • Surface crack lengths between 85% and 115% of the internal average. • Maximum load at failure no more than 110% ofthat used for K|C calculation.

From a sample thickness standpoint, E-399 requires the specimen thickness to 2 be greater than 2.5 times (K/ ay) . Only one of the room temperature specimens, as noted in the results, showed enough plasticity that the required specimen thickness should have exceeded the 6.4mm (0.25 inch) plate thickness. As temperature and toughness increase, and yield strength decreases, it would not be unusual for certain materials to require very thick (75 mm / 3 inch) samples for valid K tests using E-399. Because this is impractical for most cases of testing, other toughness test methods have been developed. The ASTM E-1820 (7) procedure applies to materials tested under these conditions. A quantity known as J is defined based on crack length, sample geometry and plastic energy absorbed during testing. Like the quantity K described above, J is also a measure of the energy required to create an additional unit of crack advance.

The J test induces either unstable crack extension or stable crack extension (tearing) in the sample. Instability results in a single-point value of toughness, while stable tearing results in a continuous fracture toughness versus crack- extension relationship (R-curve) from which significant point values may be determined. This method requires continuous measurement of load versus load-line displacement crack mouth opening as tearing occurs. In the work presented here the load line opening is not measured; rather crack opening is measured on the edge of the sample as in ASTM E-399. The J method can be J.A. Shields et al. GT8 195_

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used to determine toughness with a single sample if high-resolution crack opening measurements can be made using a compliance unloading technique to determine in-situ crack lengths. An example of the load versus crack opening displacement measurements and derived R curve that were used to calculate the J values is shown in Figure 3. The values of K|C are then calculated from the measured J values for comparison to the room temperature values. A full description of these test methods and surrounding calculations is determined in ASTM E-1820 (7). Although not all ASTM validity requirements were strictly met, the values obtained in this investigation are a good first approximation of the fracture toughness of molybdenum alloys.

4. Results and Discussion:

The room temperature and 300°C fracture toughness test results are shown in Table 4. As can be seen from the listed data, toughness measurements of the duplicate specimens showed good agreement. At both test temperatures, the orientation of the crack did not appear to have a significant effect on the toughness. This is most likely due to the similarity between the longitudinal and transverse microstructures in each heat treatment condition, Figure 1. The similarity in microstructure is developed by the recrystallization of the billet following extrusion and then further working the billet (rolling) perpendicular to the extrusion direction.

At room temperature, regardless of alloy, microstructure, and crack orientation, the fracture toughness values were all between 15 and 22 MPa m/2 (14 and 20 ksi in1/2) with the LCAC molybdenum slightly higher (18-22 MPa m72 /16 - 20 ksi in1/2) than the arc-cast TZM molybdenum (15-20 MPa tnv'114 -18 ksi in%). These fracture toughness values compare quite favorably to the literature values shown in Table 1. 196 GT8 J.A. Shields et al.

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Crack Opening Displacement (mils) 5 10 15 20 25 30 35

0 100 200 300 400 500 600 700 800

Crack Opening Displacement (i^m)

Crack Growth (mils) 0 10 20 30 40 50 60 70 80 90 50 £""' "1|"rT •I 250 E 40

200 -i n CD „n Q. 30 150 3 (ps i Z! J 20 100 CO —^ 10 t- 50

0 0.0 0.5 1.0 1.5 2.0 Crack Growth (mm) Figure 3. Typical 300°C unloading and R curve used to calculate the J - integral of 0.25 inch thick LCAC molybdenum and arc-cast TZM plate. J.A. Shields et al. GT8 197

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Table 4. Fracture toughness of 6.4 mm (0.25 inch) thick LCAC molybdenum and arc-cast TZM plate.

Test Tougl mess Material Microstructure Orientation Temperature MPa m% ksi in% Longitudinal 20,20 18, 18 Stress-relieved 20°C Transverse 21,22 19,20 (68°F) Longitudinal 18,21* 16, 19 Recrystallized LCAC Mo Transverse 19, 19 17, 17 Stress-relieved Longitudinal 73,76 66,69 300°C Transverse 69,77 63,70 (572°F) Longitudinal 59,60 54,55 Recrystallized Transverse 66,68 60,62 Longitudinal 19,20 17, 18 Stress-relieved 20°C Transverse 15, 18 14, 16 (68°F) Recrystallized Longitudinal 15, 18 14, 16 Transverse 18, 19 16, 17 Arc-cast TZM Longitudinal 84,99 76,90 Stress-relieved 300°C Transverse 88,93 80,85 (572"F) Longitudinal 58,64 53,58 Recrystallized Transverse 67,67 61,61 * Did not meet the ASTM E-399 thickness requirement.

At 300°C, the fracture toughness of the recrystallized material was similar for both the LCAC and arc-cast TZM molybdenum ranging from 58 to 68 MPa m'/z (53 to 62 ksi in72). The toughness of the stress relieved materials, however, was higher than the recrystallized material for both alloys with the arc-cast TZM showing the higher overall toughness. The stress relieved LCAC molybdenum toughness was approximately 74 MPa m/2 (67 ksi in1/a) while the arc-cast TZM molybdenum toughness was approximately 91 MPa mA (83 ksi in72).

5. Conclusions:

For molybdenum and molybdenum alloys to be considered for applications that must exhibit fracture resistance, it is necessary to characterize the materials with respect to toughness. The temperature range between ambient and 1200°C forms the region of most interest with respect to toughness since it is in this region that components such as forging dies, metalworking tooling, and furnace 198 GT8 J.A. Shields et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

structural parts must operate. In this investigation, the fracture toughness of low carbon arc-cast (LCAC) molybdenum and arc-cast molybdenum-TZM alloy plate were measured at room temperature and 300cC using compact tension specimens that were fatigue pre-cracked.

At room temperature, regardless of alloy, orientation, or microstructure, fracture toughness values between 15 and 22 MPa m/2 (14 and 20 ksi in%) were measured. These K!C values were consistent with measurements by other authors testing bar stock and employing radial crack orientations. At 300°C, the fracture toughness of recrystallized LCAC and arc-cast TZM molybdenum were also similar to one another at approximately 64 MPa n/2 (58 ksi in14). In the stress-relieved condition, however, the arc-cast TZM toughness of (91 MPa or4 / 83 ksi in%) was higher than the LCAC molybdenum (74 MPa m14 / 67 ksi in1/2). While it is often thought that molybdenum is an inherently brittle material, the results obtained here and by other investigators have shown that commercially available molybdenum has sufficient toughness at room temperature and at elevated temperatures to be a useful material for load-bearing structural applications.

6. Acknowledgements:

The authors would like to acknowledge the efforts of D. Priebe of CSM Industries for tensile testing and N. Ferri and T. Delahanty of Pittsburgh Materials Technology, Inc. for fracture toughness testing and metallographic analysis.

7. References:

1. H.E. Romine, "Fracture Toughness at room Temperature of Some Refractory Metals Based on Tungsten, Molybdenum or Columbium Which are Being Considered for Use in the Nozzles of Large Solid Propellant Rockets," NWL Report No. 1873, U. S. Naval Weapons Laboratory, Dahlgren, VA (1963).

2. C.W. Marschall and F.C. Holden, "Fracture Toughness of Refractory Metals and Alloys," in High Temperature Refractory Metals, L. Richardson, et. al., Eds., Gordon-Breach Science Publishers, pp 129-159 (1964). J.A, Shields et al. GTJ3 199

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3. D.L. Chen, B. Weiss, R. Stickler, M. Witwer, and G. Leichtfried, "Fracture Toughness of High Melting Point Materials," Proc. 13th Intl Plansee Seminar, Vol 1, H. Bildstein and R. Eck, Eds., Metallwerk Plansee, Reutte, Austria, pp 621-631 (1993).

4. A. Yu. Koval, A.D. Vasilev, and S.A. Firstov, "Fracture Toughness of Molybdenum Sheet under Brittle-Ductile Transition," Int. J. Refractory Metals and Hard Materials 15, pp 223-226 (1997).

5. M. Danylenko, "Anisotropy of Structure and Fracture Toughness," in Modeling the Mechanical Response of Structural Materials, Eric M. Taleff and Rao K. Mahidhara, Eds., The Minerals, Metals, & Materials Society, Warrendaie, PA, pp 229-235 (1997).

6. M. Rödig, H. Derz, G. Pott, and B. Werner, "Fracture Mechanics Investigations of TZM and Mo5Re," Proc. 14th Int'l Plansee Seminar, Vol 1. G. Kneringer, P. Rödhammer, and P. Wilharitz, Eds., Metallwerk Plansee, Reutte, Austria, pp 781-791 (1997).

7. Annual book of ASTM standards, Vol. 03.01, West Conshohocken, PA (2000).

8. G. Timmons and R. Yingling, "Arc Melting Molybdenum," in The Metal Molybdenum, Julius, J. Harwood, Ed., American Society for Metals, Cleveland, OH, pp 80-108 (1956). AT0200478 200 GT43 I. Smurov et al.

15" International Plansee Seminar, Eds. 6. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

COMPUTER SIMULATION OF ZrO2+8%Y2O3 AND AI2O3 POWDER PARTICLES HEATING UNDER PLASMA SPRAYING

I.Smurov*, V. Hurevich+, A.Gusarov*, S. Kundas+, T. Kashko+

* Ecole Nationale d'lngenieurs de Saint-Etienne, Saint -Etienne, France + Belarusian State University of Informatics and Radioelectronics, Minsk, Belarus

Keywords:

PLASMA SPRAYING, COMPUTER SIMULATION, PLASMA JET, POWDER PARTICLES, HEATING

1. Introduction:

Improvement of existing plasma spraying processes and development of new ones are associated with considerable problems caused by the difficulty of experimental measurement of parameters in a high-temperature plasma jet [1-3]. The use of computer simulation is one of the efficient ways to solve this problem, which provides the possibility to obtain maximum information about the investigated process with the minimum of experiments [3]. At present a special attention is paid to mathematical simulation of the processes of coating plasma spraying [1-4]. The main purposes of this work were optimisation of plasma spraying processes and investigation of influence of different parameters and variables on particles trajectories, final temperature distributions and velocities. Location of the particle, which has collided with the substrate, is determined by its diameter principally. Diameter of the particle depends on the way of the powder preparing and is approximated by Gauss distribution law. Initial conditions such as position of particle injection and initial velocity of the particle strongly influence on final particle position, particle velocity and its temperature at the substrate too. I. Smurov et al. GT 43 201

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2. Mathematical models and program facilities:

Developed particle heating and moving model consists of the following parts [5]: - heat transfer model between the particle surface and the plasma jet; - particles moving model in the plasma jet; - calculation of the temperature distribution in the particle with taking into account evaporation and phase transformation. All of above models interact between each other. For example, particle diameter's decrease caused by evaporation influences upon the trajectory of the particle. Energy flux acts upon the surface of the particle from the plasma jet. There are two ways to get warm for the particle: by radiation and by convection. Convection flux can be represented as following:

p where Nu is Nusselt number; AS- gas thermal conductivity; dp - particle diameter; 7ps - temperature of the particle at its surface; Tg - temperature of surrounding gas. Problem of convective heat transfer for spherical particles in plasma jet has been considered in [6,7,8,9]. The form of the expression for Nu-number depends on ranges of Reynolds's number and can be represented as the following: Pru-Jj; Re<2 (2)

= l.O5Reo-5Pr03; 2 < Re < 500 (3)

Equation (2) is adequate for laminar plasma forming gas flows and equation (3) is more appropriate for turbulent ones. The expression for calculation of radiation flux is built with admission that radiation acts to the surface of the particle. Materials frequently used for particles are transparent for radiation fluxes, and normally radiation acts not only on the surface of the particle. But it is very difficult to take into account this effect. So radiation flux can be written as following:

where sp is normalised emissivity factor; SB - Stefan-Boltzmann constant. 202 GT43 I. Smurov et al.

S" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reurte (2001), Vol. 4

Total heat flux can be written as:

(5) Choice of evaporation model strongly depends on pressure of the surrounded gas. For low pressures it is better to apply Hertz-Knudsen law [10]. Models based on vaporisation with backpressure are the most suitable for rapid surface vaporization at atmospheric pressures [11]. Jump condition for Knudsen layer can be written as following:

(m2 +0.5)e"l2erfc(m)- p,

/ \2 (5) T_ Y -1 m) i—Y-lm T. 1 + /T --J7V- I U 12J 1 where y is the ratio of specific heats (/ =5/3 for monatomic gases); P - gas pressure; Ps- saturated vapour pressure (it can be measured or calculated by Clausius-Clapeyron equation[12]); Ts - temperature of the particle surface; m - dimensionless velocity of vaporisation M = uj\jyRT = m^2/y ; M - Mach number; u - local mean velocity of the flow outside the Knudsen layer. Equation (5) is adequate only if the flow outside the Knudsen layer is subsonic. Mach number can't be greater than 1 during the calculation. For spherical particle, radius decrease as result of evaporation can be rewritten as:

dr. p _ dt (6)

where T and m are from (5); Pp is the mass density of the particle material; M - molecular mass of the particle. The second model represents facilities for calculation of the particle trajectories. Along the plasma jet axis only drag force is taken into account because all other forces are negligible. Drag force can be described as following: I. Smurov et al, GT43 203

15:n International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

DD (7) dt A ' ppdp

where Vg is the plasma forming gas velocity (velocity distributions of unloaded plasma jet based on experimental measurements from literature references are used in the model); pg - mass density of plasma forming gas; CD - drag coefficient. Drag coefficient depends on ranges of Reynolds's number. Authors of [6] and [13] provide the following equations for CD: 24 Cn= — Re < 0.2 D Re

CD =— (l + 0.187Re) 0.2

081 8 CD=— (l +0.11 Re ) 2.0

^^ (-(l + 0.189Ren , nnr. a<0.6"2J 21.0 < Re < 500.0

For solving equation (7) 4-th order Runge-Kutta method has been applied. In order to obtain temperature distributions at any time layer it is necessary to solve a transient thermal conductivity equation for spherical geometry with dT Stefan conditions. The symmetry condition - u in the centre of the particle, and energy flux at the surface (4) have been used as boundary condition. Developed model is non-linear and allows taking into account temperature dependence of the particle material. The difference approximation has the first order accurate in time and the second order accurate in radius. The models of particle moving (acceleration) and heating are developed as part of the global program complex. The program complex contains several models for simulation of different stages of plasma spraying. It provides common interface and access to the database for all the models. During developing of the program complex the advanced technology of the object orientation programming has been used. The structure of the program complex is shown on Fig.1. 204 GT43 1. Smurov et al.

15™ International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The bus is intended for supporting data exchange between all the modules of the complex including the database, the models and the modules of input and output data representation. It supplies the models of queried data from the database or other models. The bus realises control of the monitoring system and transfers data into it from the models. The bus is the kernel of the program complex. The program complex has an opened architecture universal protocol for data exchange. Any other model can be included into the program complex. The database contains thermal depending properties of the usually used for plasma spraying gases and powder materials. It contains temperature and velocities distributions of the unloaded plasma jets as well. Each data unit has its own reference were it comes from. The database is completely reduced to third normal form to exclude excessiveness. The module of data representation is developed for visualisation of output and input results. The input or output data can be represented as ordinary graphic, distributed diagram or by special way. For example, animation of particle moving and heating in plasma jet, diagram of phase changing, temperature and velocity distribution in plasma jet diagrams are available. Some of data units like thermal 2D distributions can be shown by several ways. For surfaces interpolation Delone's method has been used.

MODULE OF DATA Models \ PRESENTATION PARTICLES HEATING ! AND MOVING MODEL

COATING STRUCTURE FORMATION MODEL

STRESS-STRAIN STATE IN SUBSTRATE AND | COATING MODEL

MODULE OF MATH SERVICES: - 2D interpolation (Bise) - 1D interpolation (b-spline)

Fig.1. Structure of program complex Smurov et al. GT43 205 15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3. Simulation results and its discussion:

All the simulation results have been obtained under the following conditions: the initial particle velocity of 5 m/s; the initial particle temperature of 300K; the powder flow rate of 3 g/s; the distance between the powder entrance and the plasma axis of 2mm; the entrance angle of 90°. The thermal conductivity of AI2O3 and ZrO2+8%Y2O3 as function on temperature is taken from [14]. Specific heat capacity of AI2O3 as function on temperature is taken from [14] too, and for ZrO2+8%Y2O3 - from [15]. Measurements of temperature and velocity distributions in unloaded plasma jet have been used from [16] (see Figs.2, 3).

T,K Y.m

-,..-. rffl76:

II;-'AI '"''-1-'

;-.=:-•, C C:2S;

m 0 C.OIä CK 3 0-15 0.075 0.03 0. 05 0:2 X,m

Fig.2 Temperature of unloaded plasma jet 206 GT43 I. Smurov et al.

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

PfjjjjijjjjiflSHffiff m

i n v ,^'g.&J,-/.$;,|.'|-y f ••

:ro £

2:0 c ^

- ~ - _ _

IM, U.tO IJÜ15 UW 0 0/5 0 0' X,m

Fig.3 Velocity of unloaded plasma jet

Plasma spraying of AI2O3 and ZrO2+8%Y2O3. Let's consider Figs.4-7. In these pictures comparative characteristics calculated for two kind of evaporation models are shown. Particle of ZrO2+8%Y2O3 gets higher temperatures (Figs.4-5), because its specific heat is less and it has higher melting point. Temperature gradient is greater in ZrO2+8%Y2O3 because of less thermal conductivity. The particle temperature calculated by the Hertz-Knudsen law (vacuum condition) is the less than the temperature calculated by the model with the back pressure (atmospheric condition). Such particle temperature behaviour is determined by the evaporation front dynamics (Fig.6). The evaporation into the vacuum goes more quickly and the mass flow from the particle surface is greater. Energy flow to the particle surface doesn't depend on evaporation, therefore bulk of energy spend on evaporation and less one remains for particle heating. For real cases the vapour ejected from the particle can influence on the convection heat I. Smurov et al. GT43 207

15m International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

4000-

3000-

' 2000-

1000-

I I i i I I i i 20 40 60 80 20 40 60 80 Distance, mm Distance, mm

-0.4 15-

14-

\ ' I ' I 20 40 80 20 40 60 80 Distance, mm Distance, mm

transfer, but this effect is not taken into account in the developed models. For the model with back pressure the evaporation starts at the boiling point and stops when the ambient pressure reaches the saturated vapour pressure. For evaporation according to the Hertz-Knudsen law the evaporation starts at the melting point. Greater velocity of AI2O3 particle is determined by its less mass density. 208 GT43 I. Smurov et al.

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Studying of particle diameter influence on the main spraying characteristics. Spraying simulation results for particle diameters of 20, 40 and 80 urn for AI2O3 calculated by the model with back pressure are shown in Figs.8-11.

4000 —, 40 _,

3000 _

S 2000

1000 10 _

[ I 40 20 40 60 80 Distance, mm Distance, mm Fig.8 Particle surface temperature Fig.9 Melting front dynamics

300 —,

i '—i 20 40 Distance, mm Distance, mm Fig. 10 Particle velocity Fig.l 1 Particle trajectory I. Smurov et al. GT 43 209

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The particle of 20|jm and 40um are completely melted (see fig.8). The particle of 80pm is partly melted than it gets stiff with cold. The melting front becomes stationary and the solidification front appears, which moves toward melting one. This tendency is illustrated in fig.9. It is desirable to use completely melted particles for plasma spraying, which are not completely evaporated. Evaporation intensity strongly depends on the surface temperature. So in the order to satisfy all the above requirements it is better to use 40 urn particles. Particle velocity approximately linearly increases when the particle diameter decreases. This behaviour takes place because the drag force rises to the second power of the particle diameter and the particle inertia rises to the third power.

4. Conclusions:

1) Mathematical model for simulation of powder particle heating and acceleration in a plasma jet with the particle evaporation and diameter reduction is developed. Choice of evaporation model strongly depends on pressure of surrounded gas. For low pressures conditions Hertz-Knudsen law is applied. For atmospheric pressures developed model is based on vaporisation with backpressure. 2) An original software with a database for material properties has been developed for practical realisation of models. The software allows us to conduct "start-to-finish" simulation of plasma spraying with data transferred from one model to the other, and independent calculation of separate models, using the intermediate calculation results stored in the database. 3) With the software developed, the process of stabilised zirconium oxide and aluminium oxide coating plasma spraying was simulated. Results of model investigation are showed about reasonable simulation results. For example, velocity of particles strongly depends of its diameter, the evaporation to the vacuum goes more quickly and the mass flow from the particle surface is greater as atmospheric pressure, acceleration and melting of the particle with diameter 20, 40 urn more intensive than 80 urn, ets. 4) Next stage of models improvement is experimental verification and transformation for 3D plasma jet simulation.

References

1. V.V.Kudinov, P.Yu.Pekshev, V.E.Belashchenko et al. The plasma deposition of coatings., Moskow (1990) 210 GT43 I, Smurov et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

2. M.Boulos, P.Fauchais, F.Pfender Advances in Thermal Spraying, Aachen, (1993) 3. S.P.Kundas, A.P.Dostanko, A.llyuschenko at al. Computer simulation of plasma spraying processes, Minsk (1998) 4. M.Vardelle, A. Vardelle, P.Fauchais. AlChE Journal, Vol.29, No.2, (1983) 236-243. 5. Program facilities for integrated simulation of coating plasma spraying / S. Kundas, V. Gurevich, S. Levashkevich, I. Smurov, M. Ignatiev// Proc. of III Intern. Conf. "Plasma Physics and Plasma Technology". Vol. 2. Minsk, 2000. P. 612-615. 6. Deposition of plasma coating // V.V.Kudinov, P.Yu.Pekshev, V.E.Belaschenko etal. Moscow: Nauka, 1990. -406 p. (in Russian). 7. I.S. Burov To calculation of heat transfer of material dispensed particles with plasma / Physic and Chime of Material Treatment, 1979.- Ne4.-c.43- 48 8. V.V. Kudinov Plasma Coating.- Moscow: Nauka, 1977. - 184 p. (in Russian). 9. X. Chen Particle heating in a thermal plasma / Proc. of VIII Intern, symp. on plasma chem. Tokyo, 1987.-Vol.1.- Pp.4-12 10. Modeling of thermo-physical processes of impulse laser influence on metals / A.A.Uglov, I.Yu.Smurov, A.M.Lashin, A.G.Guskov. - Moscow: Nauka, 1991.-228 p. (in Russian) 11. Charles J. Knight. Rapid Surface Vaporization With Back Pressure/ AAIA JOURNAL, MAY 1979 Vol.17 No. 5 12. V. Gusarov, A.G. Gnedovets, I. Smurov. //Gas dynamics of laser ablation: Influence of ambient atmosphere. // J. Appl. Phys. 2000. V. 88. P. 4352- 4364. 13. Y.Yang, C. Coddet, M. Imbert Numerical modelling and simulation of particle behaviours in an HVOF supersonic flow/ Proceedings of 15th International Thermal Spray Conference,25-29 May 1998, Nice, France, 1998.-P. 431-438 14.R.E. Krgiganovski, Z.U. Shtern Thermal Properties Nonmetal Materials.- M.:Nauka,1981.-246p. 15.R. Brand, L. Pawlowski, G. Nener High Temperatures-High Pressure, 1986.-Vol.18.-p.65-67. 16. M. Vardelle, A. Vardelle, P. Fauchais AlChE Journal (Vol. 29, N. 2),1983.- p.236-243 AT0200479 A. Schintlmeister et a!. GT 47 21J_

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

PROTECTION OF MOLYBDENUM AND MOLYBDENUM - ALLOY THIN FILMS AGAINST LOW - TEMPERATURE OXIDADTION

A. SCHINTLMEISTER<+, P. WILHARTITZ*

* Plansee AG, A-6600 Reutte, Tyrol, Austria + Institut für Experimentalphysik, Karl Franzens Universität Graz, A-8010 Graz, Austria

Summary:

Molybdenum is an interesting functional material for thin film applications like electrical connections in thin film transistor flat panel displays (TFT-FPD) or thin film solar cells (CIS). Despite its advantageous properties, the use of pure Molybdenum has been restricted mainly because of its low stability against corrosion in moist air even at room temperature. In this paper we will show how thin films (< 0,5 urn) of Mo can be protected against low - temperature oxidation either by alloying with greater to or equal 10 wt% Rhenium or by coating with a thin layer of MoO2.

Keywords: Molybdenum, Molybdenum - alloys, thin films, low temperature oxidation, corrosion protection

1. Introduction:

The high melting temperature, the high electrical conductivity and a thermal expansion coefficient similar to silicon and glasses make Molybdenum to an interesting functional material for thin film applications in e. g. integrated circuits (IC), thin film transistor flat panel displays (TFT-LCD) or thin film solar cells (CIS). Despite its advantageous properties, the use of pure Molybdenum has been restricted mainly because of its low stability against corrosion in moist air even at room temperature. 212 GJ_47 A. Schintlmeister et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The decisive parameters and the mechanisms of the low temperature oxidation of industrial Mo surfaces have been described elsewhere (1). Due to their improved corrosion behaviour Mo/Cr and Mo/W alloys are commonly used in state of the art applications. But the main disadvantages of these alloys are their higher electrical resistance and their poorer wet chemical etchability as compared to pure Molybdenum. In this paper we will show how thin films (< 0,5 urn) of Molybdenum can be protected against low - temperature oxidation (without affecting the properties of pure Molybdenum significantly) either by alloying with lower to or equal 10 wt% Rhenium or by coating with thin layers of MoO2.

2. Experimental:

All samples were manufactured by sputter deposition of approximately 250 nm of high purity Molybdenum (4N) on alkali free glass plates with a thickness of 1mm and dimensions of up to 350 mm x 450 mm or on commercial 5" Silicon wafers. In the case of the alloyed samples the Molybdenum to Rhenium ratio was varied between 5 wt% and 41 wt% by the use of Mo/Re alloy sputtering targets. The protective layers of MoO2, which had a thickness in the range from 40 nm to 100 nm, were prepared by successive oxidation of the pure Mo films in oxygen at 250°C to 325°C and reduction in hydrogen at410°C. The corrosion behaviour of the samples was investigated in moist air and deionized water. The oxidation experiments in moist air were carried out in a climatic test chamber at temperatures from 25°C to 70°C and a relative humidity of 98%r.h.. The testing procedure with deionized water included slightly rinsing or ultrasonic cleaning of the samples. The etching rates were determined at room temperature with an etchant containing a mixture of phosporic, sulfuric and acetic acid. Classical and modern surface analytical tools such as Auger electron spectroscopy (AES), X-ray photoelectron spectroscopy (XPS), scanning electron microscopy (SEM) and AES-depth profiling have been used to characterize the samples. Changes in the electrical conductivity have been determined by the four-point probe method. A. Schintlmeister et al. GT47 213

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhamrner and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3. Results:

3.1 Corrosion behaviour:

Severe corrosive attack is observed on pure Molybden-um films after storage in moist air (see Fig.1 and Fig.2). By alloying with Rhenium the thickness of the corrosion films Fig.1: Samples after 72h of oxidation on air in a decreases almost linearly climatic test chamber at 40°C and 98%r.h. with increasing Re contents (Fig. 2). No signs of oxidation are detectable on Mo41Re and pure Molybdenum surfaces if the latter ones are coated with thin layers of MoO2 (Fig.1). Additionally the MoO2 layers show a protective effect and remain stable in deionized water (Fig.3).

200

pure Mo Mo5Re MoiORe Mo15Re Mo20Re

Fig.2: Extent of corrosive attack for different Mo-Re alloys as determined by XPS and AES - depth profiling after oxidation on air at 40°C/98%r.h./168h 214 GT47 A. Schintlmeister et al.

151"1 International Plansee Seminar, Eds. G. Kneringer, P. Rödhamrner and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3.2 Electrical Resistivity: O pure Mo Alloying of Molybdenum thin ,] films results in an increase of „.. the electrical resistivity (Fig.4). |" Comparing equal com- |" positions of the alloys, this 1 ] effect is for Chromium and s , Rhenium much stronger than for Tungsten. In the case of the Mo/Re alloys, a linear relation can be observed c. , D •»-.,, „„ ,,,, ~ »,„ «.•

... Fig.3: Resistivity of pure Mo and MoO2/Mo thin between the resistivity and the fi|ms after treatment with deionized water Re content. In contrast to alloying, the resistivity of pure Molybdenum thin films is slightly lowered by coating with thin layers of MoO2 (Fig.3 and Fig. 4).

£ 25 1CJ -

> 15

'35 S> io

_J pure Mo Mo5Re MoiORe Mo20Re Mo41Re Mo50W Mo60Cr MoO2/Mo

Fig.4: Resistivity of several Mo and Mo-alloy thin films after sputter-deposition A. Schintlmeister et al. GT47 215

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3.3 Etchability:

Alloying with Rhenium has no significant effect on the etching rate of Molybdenum in a mixture of phosporic, sulfuric and acetic acid. The etching process is strongly slowered by alloying with Tungsten and strongly catalyzed by coating with MoO2 (Fig.5).

350

pure Mo Mo5Re MoiORe Mo20Re Mo41Re Mo50W MoO2/Mo

Fig.5: Etching rates of Mo and Mo - alloy thin films 4. Conclusion:

The properties of Molybdenum for thin film applications can be improved effectively either by alloying with lower to or equal 10 wt% Rhenium or by coating with thin layers of MoO2.

5. References:

(1) A. Schintlmeister, HP. Martinz, P. Wilhartitz, F.P. Netzer: Proceedings of the 1998 Powder Metallurgy World Congress & Exhibition, Superhard / Refractory Metals, 526 AT0200480 216 GT48 P. Ramminger et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Application of High Temperature X-Ray Diffraction as a tool for material characterisation and product optimisation

P. Ramminger*, R. Tessadri*, R. Krismer**, P. Wilhartitz**

* Institute of Mineralogy and Petrography, University of Innsbruck, Austria ** Plansee AG., Reutte, Austria

Summary

X-ray Diffraction (XRD) combined with insitu high temperature experiments is a very powerful tool for the characterisation of structural parameters of industrial materials. Although this method has been known for some decades, new instrumental developments such as flexible diffracting geometries (Theta-theta movement), high temperature XRD chambers with homogeneous temperature distribution and parallel beam optics provide much better application to a wide range of specimen. Parallel to the progress in instrumental developments, data processing and evaluation of XRD data also underwent significant changes due to the development of new mathematical algorithms and new software packages which allow to gather more precise sample dependent information.

In addition to the qualitative identification of crystalline phases, XRD patterns contain a lot of additional information such as crystallite size distribution, texture, lattice parameters and residual stress. Insitu thermal observations with the high temperature XRD equipment provide additional unique information such as temperature-dependent phase transformations, the thermal changes of structural parameters and changes of the real structure parameters. Furthermore, insitu cycled thermal exposure simulates relevant changes in the physical properties similar to industrial processes. For optimal data quality, the calibration of the high temperature XRD chamber and the setup routines have to be established prior to gathering all sample- related information available. Industrial-relevant examples for the use of this method will be shown. P. Ramminger et al. GT48 217_

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Keywords

X-ray Diffraction, high temperature; parallel beam optics, crystallite size, microstrain

Introduction

X-ray Diffraction (XRD) is a traditional tool for industrial applications (steel- and cement industries) and has been mainly applied for qualitative and semi- quantitative analysis. During recent years the extraction of advanced sample- dependent parameters and the behaviour of industrial samples during non- ambient conditions has become of widespread interest, not only for fundamental research but also for manufacturing processes and quality control.

Major changes in the construction of X-ray equipment were made; i.e. theta - theta movement of the goniometer and changes of the X-ray beam path like parallel beam optics which provides flexibility of sample properties, pin-hole and glass fibre collimators, position sensitive - and 2-D detection systems for large 2-theta coverage and fast data acquisition. For simulations of physical operations or phase reaction studies, non-ambient chambers have been developed for insitu sample characterisation. Construction of new high temperature (HT) and Low temperature (LT) chambers with good temperature homogeneity provide the basis for accurate measurements under non-ambient conditions.

Software developments like advanced profile description methods (Direct Convolution Approach, Fundamental Parameter Approach) implemented into user friendly profile fitting- and Rietveld structure refinement software gathers most of the sample related information accessible. Additional specific software provides the evaluation of mechanical and/or thermal-induced impacts due to residual stress of a sample.

The aim of this contribution is to provide information about the setup and calibration procedure of a modern high temperature X-ray Diffraction system as well as the measurement and evaluation of HT-data. The procedure described in this paper is aimed to industrial research on thin film hard coatings and shows the capabilities of modern XRD analysis in respect to regular quality control and industrial research. 218 GT48 P. Ramminger et al.

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1. Experimental setup

For the characterisation of hard coatings thin films, a standard laboratory X- ray diffractometer (Siemens D5005) equipped with parallel beam optics and a high temperature XRD-chamber (Paar HTK-1200) which can be used for temperatures up to 1200°C. In combination to the HT-appliance, a high vacuum (up to 5 E-6 mbar) and gas supply station for reaction- and inert gases has been installed to provide the flexibility of different well defined gas atmospheres in the chamber in order to prevent uncontrolled reactions such as oxidation and/or reduction.

Powder diffractometer Siemens D5005 8 / 8 goniometer, Scintillation Counter with attached high temperature Chamber (Paar HTK-1200)

X-ray tube /settings fine focus Cu anode; 40kV 40mA Primary Goebel mirror parallel beam optics 2nd generation, 2° Soller slit Secondary Soller slit parallel beam optics 0.4 mm Primary slits 1; 0.6 mm Secondary slit 2 mm

Table 1: Experimental setup used for HT-characterisation

The experimental setup described above has been chosen for maximum resolution and flexibility in respect to the sample material for the following reasons:

• In order to achieve high intensity and fast data acquisition the X-ray optics and additionally attached equipments have to be perfectly aligned and adjusted. This is important, because the installation of the HTK-1200 into the beam path reduces the intensity of the diffracted X- ray by approximately 40% due to the X-ray windows of the chamber. The window material used in this study is made of a flexible capton - aluminium composite and withstands internal temperatures in the chamber of up to 1000°C. For temperatures up to 1200°C a combination of thin capton - graphite layers is used. P. Ramminger et al. GT48 219

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The disadvantage of the second setup is the low oxidation resistance of graphite at temperatures above 500°C. In order to avoid graphite erosion, inert gases such as helium or argon have to be constantly supplied during application.

• The use of gas-atmospheres in the chamber, may result in significant intensity losses due to the absorbance of heavy gas molecules. For example the use of argon reduces the diffracted X-ray intensity by approx. 95%. Depending on the chemistry of the sample alternatives like high vacuum (5 E-6 mbar) or helium should be considered.

2. Calibration of the high temperature chamber

The HT-calibration of the XRD chamber is necessary because of the uncertainties of the heat transfer between the heating coil, the thermocouple and the chamber itself. Additional factors such as the oxidation of the coil and/or the thermocouple, the gas pressure in the chamber, the gas atmosphere, etc. influence the precision of the temperature measurement. To guarantee a thermal equilibrium between the HT-chamber and the sample as well as a homogeneous temperature distribution through out the sample, low heating-rates should be preferred in general. Consideration of the thermal response of the applied thermocouple is also important for the evaluation of the calibration results. The S-type Pt/Pt90- Rh10 thermocouple used in this study responds linear in the temperature range from approx. 350°C up to 1200°C.

Three different kinds of calibrations are commonly used: [1.] measurement of the linear expansion of well characterised substances, [2.] temperature- dependent polymorphic phase transformations. [3.] observation of well- defined melting points. 220 GT48 P. Ramminger et al.

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2.1 Temperature calibration due to linear expansion of cubic materials

This calibration method is based on the isotropic thermal expansion of the lattice parameter of cubic structured materials during temperature treatment. The calibration experiments of this study were carried using a platinum disc (Merck Pt 99,997%, diameter 17mm, 0.1 mm thickness), because of the well known linear expansion coefficient of platinum and its good high temperature and oxidation resistance (1). During the high resolution measurements temperatures where increased in 100°C steps, ranging from 100°C to 1000°C (heating-rate 0.1 °C/min). The results of the obtained lattice spacing of platinum fit well to the reference values of Touloukian et al. (1) are shown in figure 1. The thermal expansion of the lattice of platinum shows a linear thermal behaviour of the chamber during firing.

3.9650

3.9600

3.9550 Calibration Experiment 2% R2 = 0.9997 JS 3.9S00 i 3.9450 in g [Angstöm ) 3.9400 n^ Reference datasets by Touloukian y^- R2 = 0.9988 3.9350 lattic e spa c

3.9300

3.9250

3.9200 100 200 300 400 500 600 700 800 900 1000 1100 Temperature [°C]

Figure 1: Comparison of the lattice parameter of Platinum: reference values (triangular) and calibration experiment (circles) on Pt 99. P. Ramminger et al. GT48 221

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2.2 Temperature calibration due to polymorphic phase transformations

This calibration method is based on isochemical phase transformation of well known substances such as SiO2 (Alpha -Beta transformation at 573°C) or BaCO3 (803°C) at 1 bar. These calibration experiments show the displacive structural change from the LT- to the HT-modification (2).

Quarz at 573"C; t A

Breakdown of (1 0 2) indices at 572° + 2°C

650°C

A Intensity 572°C A Temperature t-r2 ,Theta soo'c

Figure 2: Alpha - Beta Quartz transition experiment: The breakdown of the (1 0 2) of the trigonal Quartz modification was detected at 572°C + 2"C, shown as surface-plot (Intensity - 2Theta - temperature) and as XRD-overlay plot (Intensity vs. 2Theta).

2.3 Temperature calibration due to melting reactions

This calibration method is based on the melting-point determination of well characterised substances such as Au at 1064°C, NaCI at 804°C (2). The exact point of melting is described by the breakdown of Bragg-reflections. This calibration has not been carried out in this work. 222 GT48 P. Ramminger et al.

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3. Application of HT-XRD for investigations on chemical reactions

The use of HT-XRD in industrial laboratories is traditionally used for studying iso- or allochemical phase reactions. This has important implications for studying certain questions of reaction kinetics, such as thus a reactions proceed continuously or discontinuously, or thus a reaction proceeds directly form reactant to products or forms stable or meta-stable intermediate product- phases. The final goal is to establish the stability field in dependence of temperature (T), sample chemistry (X) and fugacity (F) of the gas phase present in the chamber.

An example of a reaction involving solid and gaseous phases is the reaction of Zr + O2 => ZrO2 Baddeleyite

£00 Li n i

•> I 600°C

•

r 1 520°C

i 450°C o 1 i

2-Thela - Scale

Figure 3: Oxidation experiment of an zirconium-sheet shown as overlay-plot ranging from 450°C - 600°C, temp-steps 10°C; The formation of Baddeleyite was established at 520°C ± 10°C in this experiment. P. Ramminger et al. GT48 223^

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4. Determination of real-structure and crystallographic data

The study of chemical reactions and phase transformations as shown in figure 3 is capable to provide a lot more information than transformation- temperatures since comprehensive treatment of XRD-data provides real- structure information (crystallite size, microstrain, residual stress) and information about the crystallographic parameters of the specimen during HT- treatment (i.e. observation of the recrystallisation process during heating and/or cooling). The extraction of lattice parameters (i.e. figure 1; determination of the thermal induced expansion of platinum) from XRD-patterns can be carried out easily using various software packages. Even more sophisticated information (i.e. structural refinement of atomic coordinates) can be calculated from the XRD- data-sets by using Rietveld-based algorithms (3). However, in order to determine real-structure parameters like size of the crystallites and the microstrain of a crystalline sample, a close description of the obtained peak shapes and the instrumental contribution of the observed pattern is necessary. Although there are a lot of software packages, which are capable to describe the sample dependent information sufficiently, the new profile fitting software module TOPAS (Bruker-AXS) is an exciting development for industrial research, since it is based on a powerful mathematical algorithm and has the advantage of an easy to learn and intuitive handling.

4.1. Theoretical background

The enduring development of X-ray diffractometers allows to evaluate the sample specific information on a very high level of precision. To maintain this grade of performance, new advanced mathematical methods are necessary to describe each observed peak shape as close as possible. These methods define an instrumental function which is used to derive the instrumental contribution to the peak profile of an unknown substance. Based on the determination of the instrumental contribution, sample contributions like crystallite size, strain, microstrain can be extracted. 224 GT48 P. Ramminger et al,

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4.2 Parametrisation of the instrumental contribution

The description of the instrumental function is important regardless of the profile fitting methods applied. This instrumental contribution can be calculated in the case of Bragg-Brentano geometries and fundamental convolution based approach for the individual instrument (4). The basis for this approach is the strict separation of the instrumental contribution, the contribution caused by the X-ray source and the sample. This is performed by using a pool of well defined aberrational functions for the description of the instrument. These aberrational functions are included in Fundamental Parameter Approach (FPA) based Software modules like TOPAS. Due to the widespread use of new X-ray geometries (parallel beam optics as used in this study), this convolution based profile fitting method has to be adapted. In order to derive a unique instrumental function for non-Bragg-Brentano geometries, CeO2, LaB6 and Y2O3 were used as experimental standards, because they do not contribute any crystallite size, strain and microstrain to the observed peak profile. Thus the experimentally gathered XRD-pattern reflect the contribution caused by the instrument and the X-ray source. In a second step the aberrational functions provided by TOPAS are convoluted over the peak profile of these idealized specimen profiles to obtain the instrumental function. In addition, the obtained function strongly depends on the X-ray optics and wavelength characteristics used and has to be adopted whenever one of the components is changed. P. Ramminger et al. GT48 225

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4.3. Preparation of reference samples

For determination of the instrumental contributions, well described substances are needed which are nearly perfectly crystallized. Such an ideal substance does not introduce any additional contributions to the instrumental profile function such as domain size broadening caused by crystallite size, strain, microstrain and lattice defects to the peak profile. Three substances, cerium oxide (CeO2), Yttrium Oxide (Y2O3) and lanthanum hexaboride (LaB6) which meet the requirements for a accurate determination of the instrumental breadth were used in the present study .

substance Formula approx. crystallite size approx. crystallite size approx. strain [ran] [nm] Lorenzian Value Gaussian Value

Lanthanum LaB6 -1000 -300 -0,001 Hexaboride

Cerium Oxide CeO2 -1000 -1000-2000 -0,001

Yttrium Oxide Y2O3 -1000 -400 -0,001

Table 2: Reference substances used for parametrisation of the instrumental function

In order to minimize microstrain and the number of lattice defects (crystallite size), Cerium Oxide and Yttrium Oxide were annealed at 1600°C for 7 days in a platinum (Pt95Pd5) crucible using a muffel furnace. To suppress the introduction of strain as a result of cooling after the heat treatment, a cooling ramp of TC/min was applied to decrease temperature from 1600°C to 300°C. After reaching a temperature of 300°C, the crucible was removed from the furnace and stored in a desiccator for further cooling-down to room temperature. As a result of the heat treatment, Y2O3 and CeO2 sintered to form lumps. This required subsequent grinding in a agate-mortar with minimum pressure applied, to prevent the introduction of additional strain. 226 GT48 P. Ramminger et al.

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4.4 Evaluation of the parametrisation procedure

0 J3?91 33.15OSS 55 85 105000 100 000 1 95 000 : S0Ü0O 1 65 000 / 80 000 75 000 70 000- t 65 000 v 60 000 A ! 55 0CO SOOOO i ' 1 \ — ~ -/—V ,1 45 MO 1 / \ «tOOO- | j \ t 35 000 30 000- / 1 2 1\ 25 00O 20 000- 15 030 J j V / 10 000 \ 5 000 y e J \ / \ v 30 •. 31 2 33 34

Figure 4: The XRD-pattern show the profile description of LaB6 (left peak), CeO2 (middle peak) and Y2O3 (right peak) by the obtained instrumental function. The close fit to the observed peak profile is documented by the small difference plot (x-axes).

As shown in figure 4 (detail-B) the LaB6 -reflex at 30.44 °2Theta; (1) marks the observed peak profile. The calculated fit-profile (2) indicates a severely misfit due to a wrong instrumental function. This clearly shows that a correct parametrisation is necessary for a full describtion of the peak-profile, when using non-Bragg-Brentano geometries and the fundamental parameter approach. P. Ramminger et al. GT48 227

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4.5 Application of the instrumental function for the evaluation of real-structure parameters

2-Theta - Scale

Figure 5: This HT-experiment shows the recrystaliisation of hard coatings during temperature treatment. The decrease of peak-broadening with increasing temperature indicates the changes of real-structure parameters with temperature, (for example detail A: increase of crystallite size ranging from 6 nm at 300°C up to 22 nm at 800°C).

5. Conclusion

Using modern X-ray powder diffractometers with high Temperature equipment, unique information can be achieved by insitu thermal observation. This combination enables the determination of thermal changes in crystallographic expansion, real structure parameters and phase kinetics. Furthermore, insitu cycled thermal exposure simulates relevant changes in physical properties during the production and application.

To achieve high quality results in routine analysis, an exact temperature calibration of the HT-chamber has to be performed. Furthermore, the knowledge of the precise instrumental contributions is fundamental for the extraction of crystallographic and real-structure data. 228 GT 48 P. Ramminger et al.

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This work has shown, that the new possibilities of HT-diffraction are interesting improvements for industrial research. The integration of this method into the analytical routine of process- and quality control, makes it necessary to adopt certain aspects for quick and easy data-evaluation. Nevertheless, the results prove that the experimental setup used is capable of resolving slight structural differences, which allows to apply this method of characterisation in a wide range of industrial and scientific tasks.

References:

(1) Touloukian et al.: "Thermal expansion metallic elements and alloys" Thermophysical properties of matter, IFI/Plenum, New York - Washington, 1975, pp. 1348

(2) Kern A.: "Hochtemperatur Rietveld Analysen: Möglichkeiten und Grenzen", (Heidelberger Geowiss. Abh., Heidelberg, 1998), pp. 57-102

(3) Young, R.A.: "Introduction to the Rietveld Method", (IUCR Book Series, Oxford University Press, 1993) pp. 289

(4) Cheary, R.W. & Coelho, A.A., "A fundamental parameters approach of X-ray line profile fitting, (Journal of Applied Crystallography 1992), Vol 25, pp. 109-121 AT0200481 A. Bock et al. HM 2 229

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PRODUCTION AND CHARACTERIZATION OF ULTRAFINE WC POWDERS Andreas Bock and Burghard Zeiler Wolfram Bergbau- und Hüttengesellschaft mbH NfG KG, St. Martin, Austria

Abstract:

The conventional Calzination-Reduction-Carbuhzation (CRC) process offers the potential to manufacture commercial WC powders with median SEM grain sizes below 0.5 urn (ultrafine grades). Strict process control and a high degree of automation have led to increased powder uniformity and high lot-to- lot reliability. The high potential and flexibility of the CRC process is shown by the development of tailor made WC powders with regard to subsequent alloy manufacturing. R&D powders have successfully been processed in full-scale units exhibiting SEM grain sizes of 0.15 - 0.20 urn. This paper discusses the powder characteristics of various ultrafine WC powder grades, stemming from the conventional CRC process. Analytical characterization include also field- emission SEM, bright field TEM, EDX analysis and XRD-line broadening.

Keywords: tungsten carbide, ultrafine, CRC-process, characterization

1. Introduction:

The outstanding success of ultrafine cemented carbides -exhibiting median WC grain sizes between 0.1-0.5 urn in the sintered alloys-, is based on their unique combination of material properties [1]. Hardness, wear resistance against attrition and abrasion, alloy strength / rigidity and resistance to chemical wear are optimized in order to meet the demands in many different fields of application, e.g. pcb drills, extruded parts for metal drilling, wood machining, wear parts or cutting blades [2].

Today, the hard metal industry has gained knowledge to utilize the potential inherent to ultrafine alloys. Individual powder producers and alloy manufacturers have specialized on this type of alloy. In close contact to each other, both have developed carbide powders grades tailor made for further alloy processing. 230 HM 2 A. Bock et al.

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Today's demands for ultrafine carbide manufacture are driven by the necessity to meet the desired alloy properties and to enable proper powder characteristics for alloy processing:

• Optimal combination of hardness, toughness and strength linked with low residual porosity, high alloy purity and strict avoidance of carbide heterogeneities constitute the basis. • Manifold demands arise from alloy processing itself, e.g.: uniform particle dispersion during WC+Co powder milling, suitable sinter activity (as indicated by alloy densification, densification rate in every stage of solid- and liquid phase sintering, uniform diffusion of binder in the ultrafine carbide matrix) and controlled carbide grain growth behavior (as determined by carbide structure, doping respectively doping procedure of Cr, V, Ta and carbon balance throughout processing).

It is important to notice that these trends also contain several conflicting demands: e.g. high sintering activity versus low growth tendency in case of low binder alloys, or increasing hardness combined with low shrinkage.

The aim of this paper is to discuss the intrinsic powder characteristics of standard ultrafine WC powder grades and the potential and flexibility of the CRC process in order to develop tailor made products.

2. CRC Process (Calzination-Reduction-Carburization):

The most challenging task for powder producers is to develop carbide grades, which are tailor-made for the respective conditions during further alloy processing. Thus ultrafine carbides decisively vary with regard to their SEM particle size distribution, defect structure of the WC grains or distribution of grain growth inhibitors (e.g.: chromium and/or vanadium). Among others, such powder properties constitute the sintering activity of the carbide and have to be built in already during the reaction path "tungsten oxide to ultrafine carbide powder". This is achieved by control of powder structure and powder dispersion, even in the first reaction steps, grain size and size distribution, monitoring of trace elements, appropriate addition of growth inhibitors and strict avoidance of impurities [3-8].

Automated or semi-automated processes and stringent control of processing conditions (Fig.1) constitute the basis for a high lot-to-lot reliability and to A. Bock et al. HM 2 231

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Figure 1: Modern CRC process based on a high degree of automation enabling high lot-to-lot reliability

The constitution of ultrafine carbides is determined by the powder agglomeration, the particle size, grain size and the crystal defect structure (Fig. 2). Furthermore powders have to show a high chemical consistency, e.g. local carbon balance, absence of subcarbides, uniform distribution of additives.

FESEM image 0.1

*•.«

defect structure •,

Figure 2: FESEM and bright-field TEM image showing the physical constitution of ultrafine carbides 232 HM 2 A. Bock et al.

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3. Experimental Setup and Powder Characterization:

The experimental setup of this study can be summarized by Figure 3, which shows a plot of BET specific surface area versus FSSS grain size of ultrafine carbides investigated. Carbides cover a broad range between 1.5 to 2.8 m2/g BET surface area.

• Carbides designated CRC 025 and CRC 020 indicate ultrafine standard grades. In this case tailor-made carbides were selected in order to evaluate their intrinsic properties. • Special applications, e.g. wear parts for extreme abrasive conditions, push a trend towards even finer carbides, as characterized by specific surface area (carbides up to 2,8 m2/g).

0,70

0,65

s ° CRC02 S= 0,60 WC06 5 I « w CRC02 .= 0,55 WC05 re re k_ &_ (O 1 V) > 0,50 0,45 |

0,40 i 1,00 1,50 2,00 2,50 3,00 3,50 BET specific surface area [m2/g]

Figure 3: BET specific surface area versus FSSS grain size of ultrafine carbides investigated

Conventional powder characterization still is a powerful tool in order to evaluate the consistency of an individual carbide lot. However previous A. Bock et al. HM 2 233

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studies [9] already indicated the limitations regarding determination of FSSS grain size or grain size distribution measurements due to inherent principles (e.g. detection limit in the ultrafine particle range; assumptions used for calculation of size distribution).

Thus the BET surface area measurement is gaining more importance for routine analysis and production control. BET grain size can be calculated by:

c BET grainsize [|am] = 15.6 * BET surface area [m2/g]

Based on measurement of minimum 300 particles FESEM particle size is given by [9]:

,_,_„,_.„ ,. . . r , sum of diameter fd(i)-d(j)] FESEM particle size [|j.m] = 1~~^ ^ total number of particles

FESEM particle size distribution was calculated by:

sum of diameter [d(i)-d(j)] Frequency[%] = - sum of particle diameter of all measured particles

Bright field TEM and EDX analyses were carried out with a Philips CM20.

XRD-line broadening was derived from measurement of the crystallographic planes [100] and [001]. A Lanthanhexaboride standard (NIST SRM660) was used for measurement of the instrumental width.

4. Standard Ultrafine Powder Grades

Recent developments in the area of standard ultrafine grades were first of all driven by the need to control the sinter-activity and grain growth tendency of ultrafine carbides. Table 1 shows the basic powder characteristics of CRC025 and CRC020, both in the case of the undoped or the chromium-vanadium doped version. 234 HM2 A. Bock et al. 15' International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

CRC025 CRC020 Designation: WC 06 WC05 FSSS grain size [jjrn] powder as supplied 0.63 0.64 0.55 0.56 BET specific surface area fm2/gl 1.71 1.52 2.03 1.95 BET calculated grain size Turn] 0.22 0.25 0.19 0.20 total carbon \%] 6.12 6.17 6.09 6.19 free carbon [%] 0.03 0.03 0.03 0.02 VC [%1 undoped 0.1-0.5 undoped 0.3-1.0 Cr3C2 [%1 undoped 0.1-0.5 undoped 0.3-1.0 A! typical <' 5 ppm Ca typical •i 2 ppm Co typical •i 5 ppm Fe typical 20 ppm Mo typical<20 ppm Ni typical<; 5 ppm O 900 ppm 1030 ppm 1080 ppm 1450 ppm S typical <= 5 ppm Si typical <10 ppm

Table 1: Standard ultrafine WC powder grades CRC025 and CRC020

Field emission secondary electron microscopy (FESEM) has gained increased importance in order to characterize the particle morphology and true particle size (see Figure 4). For CRC 025 the FESEM particle size is ~0.30um, for CRC 020 the FESEM particle size is ~0.25um.

The comparison of FESEM particle size with the calculated BET grain size (Figure 4) shows a close correlation; the FESEM particle size being slightly higher. This difference can be explained by the polygonal morphology of tungsten carbides and surface roughness.

FESEM particle size distribution provides additional, valuable information, as shown in Figure 5. Although the BET grain size is 0.20 urn in both powders designated CRC020, the frequency of fines in the range between 0.1-0.2 urn can vary significantly. A higher frequency of fines can be advantageous in order to increase the sinter activity/densification rate (at low sintering temperatures during solid state sintering) or to enable slightly higher hardness at maximum grain growth inhibition. A. Bock et al. HM 2 235

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CRC025 1 CRC020 WC06 doped WC05 doped BET surface = 1.52 nfVg BET surface = 1.95 nrVg BET calc. grain size = 0.25 urn BET calc. grain size = 0.20 urn FESEM particle size = 0.30 urn FESEM particle size = 0.25 um Figure 4: FESEM imaging and comparison of FESEM particle size with BET grain size in case of CRC025 and CRC020

25%

20%

.5 15%

>> u a 10%

0% 0,01 0,1 1 10 FESEM particle diameter [pm] Figure 5: Variation in FESEM particle size distribution for CRC020 showing distinct variations in amount of fines (0.1-0.2 urn] 236 HM 2 A. Bock et al.

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Typical defect structures proved to be twins or stacking faults; only a minor amount of dislocations were detected (Figure 6). XRD-line broadening provides an additional method to detect grain defect structures: • CRC020: XRD line broadening of [001]: 0.20-0.22 ° • CRC025: XRD line broadening of [001]: 0.17-0.20 ° However it is still not clarified, if such low energy grain boundaries do have a significant effect on grain growth tendency during subsequent alloy sintering.

0.2

Figure 6: Typical occurrence of twinning or stacking faults observed by TEM imaging

In case of predoped carbide powders it was postulated that Chromium can be uniformly distributed at the WC surface if Chromium is added already during powder processing. Long lasting experiences showed a more uniform grain growth restriction during alloy processing.

EDX analysis (see Figure 7) gave clear evidence, that Chromium is preferably present at WC/WC grain boundaries even in the interior of polycrystalline particles or in necks between carbide particles. No Chromium was detected in the interior part of carbide grains by EDX. A. Bock et al. HM 2 237

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WC grain

-AH;

W Cr L ):

X3S73 5i7M.-SS WCV K0 EDX analysis Grain boundary

Figure 7: Results of EDX analysis indicating that Chromium is present at WC/WC grain boundaries and necks

The agglomeration and structure of carbide particles (mono-polycrystalline) plays an important role for WC + Co powder dispersion and the growth behavior during sintering. CRC process is successfully used to alternate the powder structure.

An example is presented for CRC020, both powders showing comparable median BET grain size and FESEM particle size (see Figure 8). However TEM investigations showed clear differences: The first carbide consists mainly of polycrystalline carbides, whereas the second carbide typically shows a high particle dispersion. 238 HM 2 A. Bock et al.

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CRC020 doped 1 um CRC020 doped 2 BET surface area = 1.95m /g BET surface = 1.95 m2/g FESEM part, size = 0.25 urn FESEM part, size = 0.26 urn

0.2 urn

Figure 8: Variation of the carbide structure (poly- or monocrystalline) in the case of CRC020

Furthermore TEM investigations (presented in Figure 9) revealed that the first carbide (left) showed frequent occurrence of stacking faults; areas with low amount of defect structures were more often detected within the second carbide (right). A. Bock et al. HM 2 239

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0.1 um Figure 9: Variation in the amount of defect structures present in CRC020 carbides showed in Figure 8

5. Development Grades CRC015

Special applications, for example in the area of wear parts, set the demand for Carbide grades with BET specific surface areas higher than 2.4 m2/g. Carbides exhibiting calculated BET grain sizes of 0.15 urn were produced in full scale units via the conventional CRC process (see Table 2).

Designation: CRC015 CRC015 FSSS grain size [urn] powder as supplied 0.49 0.49 BET specific surface area [m2/gl 2.43 2.83 BET calculated grain size fl-iml 0.16 0.14 total carbon [%] 6.11 6.15 free carbon [%l 0.01 0.05 vc r%i 0.3-0.6 0.3-0.6 Cr3C2 [%l 0.3-1.0 0.3-1.0 0 1810 ppm 2800 ppm

Table 2: Powder characteristics of carbide grades CRC015 240 HM 2 A. Bock et al.

15- International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The comparison between CRC 020 and CRC 015 (Figure 10) clearly shows the decrease in carbide particle size. FESEM particle size distribution measurement reveals a median particle size of 0.18um and a very uniform size distribution.

%' CRC020 1 pm CRC015 2 2 BET surface = 1.95 m /g BET surface area = 2.83 m /g BET calc. grain size = 0.20 urn BET calc. grain size = 0.14 |jm

0,1 10 FESEM diameter [[im] Figure 10: FESEM image and FESEM particle size distribution for CRC015 compared with CRC020 A. Bock et al. HM 2 241

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Again TEM was used to characterize the intrinsic carbide structure of CRC015 (Figure 11). However the limits of TEM became obvious:

• The structure of WC/WC grain boundaries in the interior of particles and agglomerates was not characterized so far. High resolution TEM might provide a possible solution. However an interpretation would be needed, in order to determine the effective carbide grain size for subsequent alloy production. The question is, which kind of grain boundaries will be penetrated by the cobalt binder phase during sintering thereby reducing the effective grain size, and which grains are not separated and recrystallize during sintering. • Local inhomogeneities, which are relevant for carbide growth processes, can also not be characterized in an appropriate way; e.g. carbon, inhibitor- distribution.

0.2 urn Figure 12: Bright field TEM image of grade CRC015

A sintering test with inhibited grain growth might provide an additional tool to characterize the potential inherent to a carbide grade. 10wt% Co alloys were produced by conventional liquid-phase sintering at 1420 C. Alloys were 242 HM 2 A. Bock et al.

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doped with rather high amounts of chromium- and vanadiumcarbide. The maximum coercivity of the alloys was measured and plotted versus BET specific surface area (Figure 13). The increase in maximum coercivity is directly linked to the BET surface area.

600

580

"ÖT 560 O > 540

•5 520 0 500 u | 480 1 460 ra E 440 420 400 ' 1,00 1,50 2,00 2,50 3,00 BET specific surface area [rrvVg] of carbide Figure 13: Maximum coercivity [Oe] of WC/10wt%Co alloys doped with VC (>1wt%) and Cr3C2 (>0.5wt%)

BSE imaging of polished sections, etched by hydrofluoric acid, again shows the decrease in carbide grain size from CRC 025 to CRC 015. The estimated sintered carbide grain size of CRC015 is comparable to the FESEM particle size of the powder (0.15-0.20um). The mean linear intercept length was so far not characterized but for geometrical reasons must be clearly below 0.15 urn. A. Bock et al. HM 2 243^

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&\*M*:*. Plit ^,®K^ CRC025 CRC020 CRC01S5 max. coercivity = 443 Oe max. coercivity = 465 Oe max. coercivity = 556 Oe

Figure 14: Microstructures of WC/10wt%Co alloys doped with VC (>1wt%) and Cr3C2 (>0.5wt%) - BSE image of HF-etched polished sections

6. Summary

Ultrafine carbides are still of major interest to the alloy manufacturers. At a BET surface area between 1.5 to 2.9 m2/g the FESEM particle size is in the range of 0.18 to 0.30 urn.

• The FESEM particle size can be varied in order to increase the amount of fines (fraction of 0.1-0.2 urn) at a given median BET grain size. • The structure of particles can be decisively changed: monocrystalline or polycrystalline built-up and variations of the internal defect structure were analyzed byTEM imaging. • Chromium as an important grain growth inhibitor can be uniformly distributed at carbide/carbide grain boundaries, as was shown by EDX analysis.

Sintering tests with inhibited grain growth were used to characterize the potential inherent to CRC powders. The finest grades CRC015 showed a maximum coercivity of 560 Oe (WC-10wt%Co).

The potential and flexibility of the CRC process combined with it's high lot-to lot reliability are successful tools in order to meet the demands of ultrafine hard metal manufacture. 244 HM 2 A. Bock et al.

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7. Acknowledgement:

The national fund "Forschungsförderungsfonds für die gewerbliche Wirtschaft" (Vienna, Austria), is gratefully acknowledged for the financial support of this work.

8. References:

[1] Spriggs G.E., Int. J. Refract. Met. & Hard Mater., Vol. 13, 1995, 241-255. [2] LJ. Prakash, Int. J. Refract. Met. & Hard Mater., Vol. 13, 1995, 257-264. [3] Schubert, W.D., Lassner, E., Int. J. Refract. Met. & Hard Mater., Vol. 10, 1991, pp. 133-41 [4] Schubert, W.D., Lassner, E., Int. J. Refract. Met. & Hard Mater., Vol. 10, 1991, pp. 171-83 [5] Schubert, W.D., Bock, A., Lux, B., Int. J. Refract. Met. & Hard Mater., Vol. 13, 1995, pp. 281-96 [6] Lassner, E., Int. J. Refract. Met. & Hard Mater., Vol. 8, 1989, pp. 185-8 [7] Bock, A., et al., Proc. of the 14th Int. Plansee Sem. 1997, HM62, Vol.4, 342-362. [8] A.Bock, W.D.Schubert and B.Lux, Powder Metallurgy Int. 24, 20-26, 1992. [9] B.Zeiler, Proc. of the 14th int. Plansee Sem., 1997, HM13, Vol. 4, 265- 276. AT0200482 B. Roebuck et al. HM 84 245

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Hardmetals - Microstructural Design, Testing and Property Maps

B Roebuck, M G Gee and R Morrell

NPL Materials Centre National Physical Laboratory, Teddington, Middlesex, TW11 OLW, UK

Summary

The production of WC/Co hardmetals and their analogues is often considered to be a mature technology. In recent years however there has been a considerable amount of stimulating research where the concepts of microstructural design have been used to produce alternatives to the conventional two-phase structure. The potential of microstructural design is reviewed, particularly the possibilities of performance improvement via changes in size, shape and distribution of the phases. For example, ultrafine grained materials are key technology drivers and pose severe measurement challenges.

Hardmetal technology is also well served by established and standardised laboratory bench-marking property tests such as hardness, magnetic properties, density etc. Recently there has been extensive activity in a number of performance-related testing areas such as wear, fracture and fatigue. These are also reviewed, bearing in mind their ability to discriminate between the effects of differences in microstructural design.

Finally, the concept of property mapping is introduced as a tool for providing a framework for optimising properties. Their utility in correlating performance properties and their relationships with microstructural parameters is evaluated. Their potential for quantifying the differences in properties that unconventional microstructures might offer is also discussed.

KEY WORDS: Microstructure, Design, Property Maps, Fatigue, Wear. HM 84B. Roebuck et al.

15' International Plansee Seminar, Eids. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

1 Introduction

Hardmetals, cermets and other sintered hard materials are used in a very diverse range of applications. The use of hardmetals is a mature technology. Industry is currently well served by a range of baseline established standards [1] which, if properly followed with good attention to correct quality procedures, will ensure consistent products. However, even with these standards it is essential to have a full understanding of the effects of test method variables. For example, process dependent grinding stresses introduced during testpiece preparation can change apparent bend strength by up to 100% [2]. Consequently, laboratory benchmark testing is vital, in particular for:

• Characterising structural effects • Quality control • Comparing materials • Quantification of properties

Industry and supporting research organisations now use a wide variety of tests, some of which are standardised and of longstanding pedigree, for example, density measurements. Others, that are thought to be relevant and fairly well understood have some limitations, such as bend testing (transverse rupture tests). Finally, there are research or in-house methods, and it is significant that many of the properties which have a strong influence on the performance of hardmetal products, such as corrosion, fatigue, impact, wear and high temperature strength and toughness, are often measured, but not always by standard test methods.

Key microstructural parameters self-evidently control material properties and thus influence the interpretation of the results of test methods. New hardmetals based on WC/Co are evaluated by industry on a continuous basis, particularly to examine new grades of powders, powder mixes and compositions. New materials need to be benchmarked against conventional materials for basic properties such as hardness, wear resistance and toughness. This task is lengthy and frequently difficult because of the lack of good data from validated measurement methods by which to make the comparisons. For this purpose two types of property map are discussed: one where the property is plotted against a microstructural feature such as WC grain size or Co binder phase content and one where different properties, such as hardness and toughness are mapped against each other. B. Roebuck et al. . HM 84 247

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In summary, the paper comprises sections on

• Microstructural Design • Recent developments in testing • Property maps and underlines the importance of baseline understanding for the interpretation of new materials.

2 Microstructural Design

WC/Co hardmetals are generally manufactured to comprise two phases, WC and Co, where Co is an alloy of Co-W-C and WC is stoichiometric. WC/Co is a pseudo-binary section in a three component system (W, C and Co). If the processing conditions are such that the carbon content is either too high or too low then other phases are present in the structure, i.e. graphite when the C content is high and eta-phase (a mixed Co, W carbide) when the C content is low. The structure of two phase WC/Co hardmetals is conventionally defined by these key parameters:

• Grain size of the WC phase • Co binder mean free path • Volume (or wt) fraction and composition of the Co binder-phase

There are no standards for the definition of feature size in hardmetal structures. In their absence the number average linear intercept is used both for the WC grain size and the Co binder-phase mean free path, as measured on 2-D polished and etched sections.

The current most important driver in microstructural design is the reduction in WC grain size. Powders in the nanometre size range are available. The technological challenge is to maintain this small crystal size in the sintered product. Property maps (see section 4) indicate that there are substantial gains to be made in abrasion resistance and hardness from a reduction in the grain size. It is imperative to develop suitable methods of characterisation in order to validate the model predictions. However, there is considerable difficulty in measuring the grain size when the crystals are sub 0.1 |j,m. Co mean free path and distribution play a crucial role in providing adequate toughness in conventional WC/Co hardmetals and some consideration should be given to this parameter when developing ultrafine grained structures. 248 HM 84 B. Roebuck et al.

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Nano WC 10 nm 99.5- NanoCc Fine WC 95- = 7°- CO "o 40- D. CD E 10-

1 - Ultrafine Cc Fine Co Ultrafine WC

0.1 0.01 0.1 Intercept size

Fig 1 Co and WC measured and calculated intercept distributions in ultrafine grained structures.

In conventional hardmetals the Co intercept (mean free path) distribution is linked to the WC intercept distribution [3]. This is shown in Fig 1 where a plot is given of a measured WC intercept distribution in a 10%Co ultrafine grained hardmetal plotted as number probability against log intercept size compared with a calculated Co intercept distribution. It can be seen that as the mean value of the WC intercept distribution decreases it is not possible to maintain similarity in Co mean free paths since the Co intercept values at the lower end of the distribution are not physically realistic when the calculated size is less than atomic dimensions. Even at the scale of 10 nm the properties of Co are unlikely to be similar to those of larger domains. A change in the morphological relations between the WC and Co can thus be predicted and this will clearly impinge on the properties in ways which have yet to be determined. There are added implications for the use of magnetic techniques, such as coercivity, for the quality control of structure and further work is needed to evaluate binder phase distribution as well as the measurement of WC grain size in these nano and ultrafine grained materials. B. Roebuck et al. HM 84 249

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Stoichiometric Dislocation Layered Gradient Solid or Solid Solution Structures Structures Solution (W,X)C

Fig 2 Schematic diagram of intrinsic microstructurai parameters.

Many of the key strengthening and toughening methods common to physical metallurgy, such as solid solution strengthening, phase transformations, precipitation hardening, etc, have been used in the development of WC/Co materials. These can be divided into intnnsic and extrinsic structure-related mechanisms. 250 HM 84 B. Roebuck et al.

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Intrinsic factors relevant to the WC phase [Fig 2] can include:

Solid solution - e.g. Mo, which can partition between binder and hard phase. Orientation - the hardness of individual crystals can vary by a factor of 2, dependent on orientation. Dislocation - recent high pressure studies [4] indicate that structures beneficial mechanical properties may accrue from changes in the internal dislocation structure (twins, stacking faults, tangles). Nanoreinforcement - the principles adopted by the ceramic community (additions of nano-sized carbides, oxides, nitrides etc) to increase strength. Either at boundaries or intragranular. Precipitation - a hardening mechanism not yet evaluated in WC. Multiphase or - perhaps taking advantage of epitaxial growth layered or core/rim between similar crystal structures (TiB2/WC etc). structures within individual crystals

Intrinsic factors applicable to the binder phase could be:

Solid solution - alloying by W and C well understood. Other elemental additions (Ni, Ru, Fe, Cr, Mo and combinations) are fairly well understood but there is more scope for tailoring properties to performance (corrosion, fatigue, etc). Precipitation - Co3W in low C WC/Co hardmetals; science reasonably well understood. Phase homogeneity - fec/hep; difficult to control independently of other and transformations parameters. Martensitic structures in Fe-based binders. Multiphase - Ni alloy gamma/gamma prime (y/y') analogues; structures some science developed but Co-Ni(X) y/y' structures less well understood. Intermetallics - Aluminides investigated but other types could merit evaluation. B. Roebuck et al. HM 84 251

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Clearly there are many possibilities either individually or in combination. Most of these mechanisms predict higher strengths to different degrees [5] but less is known of the potential changes to toughness. At this stage in the evolution of predictive composite microstructural design it is still necessary not only to measure the toughness of structures with altered states but also to be able to compare them with baseline properties. For this purpose materials with single phase WC and a Co(W,C) alloy intermediate between the graphite and eta- phase fields should be used.

Extrinsic mechanisms follow from the principles of composite engineering, either on the level of individual phases or with more macroscopic multi-material morphologies. At the level of the individual phases, e.g. WC and binder, the two most important factors are shape and distribution, although in principle it is difficult to independently change the relevant parameters for each phase as there are space filling constraints on the binder phase.

The unconstrained shape of single crystal WC is a truncated trigonal prism of aspect ratio dependent on the chemistry of the growth conditions. Two dimensional sections observed in conventional micrographs are typical of triangular (basal plane) or prismatic sections through these shapes. It has been shown in both WC/Co and WC/Ni systems that either more platelike crystals [6,7] or more rounded shapes [8] can be developed that result in shifts in the hardness/toughness map either through constraints on the crack morphology in finer grained materials or through changes in the binder-phase size distribution. This would appear to be a fruitful area for further research, as is demonstrated by other papers in this 15th Plansee Seminar [9,10].

With regard to size distribution there are two key aspects:

• changes in feature distribution width (either WC, binder or both) in a monotonic sense

• size mixes of definite modality (bimodal, multimodal, etc) [11,12]; again either for WC and/or binder. 252 HM 84 B. Roebuck et al. 5^ International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

o o

A) Participate B) Laminate C) Fibrous macroscopic bi or multimodal oriented

D) Fibrous E) Particulate F) Hybrid random multiscale cored

G) Hybrid H} Graded I) Particulate layered microscopic

Fig 3 Schematic diagram - composite architectures.

In many cases what appear to be bimodal structures to the eye are difficult to quantify as such using image analysis and the measured distribution can be monomodal, but of wide dispersion. Clearly the issue here is the development of agreed methods of characterising the feature distribution parameters - for both carbide and binder, particularly for the definition of uniformity of size [11,13]. A further aspect of distribution is that of clustering. In principle this can be assessed through measurements of nearest neighbour distances or perhaps B. Roebuck et al. HM 84 253

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through the more traditional parameter of contiguity in hardmetals. However, this parameter may not be easy to manipulate independently of changes in other features, but it should not be overlooked as it may play a significant part in controlling strength and fracture [14].

Other types of extrinsic factors affecting microstructural design are multimaterial composite architectures (Fig 3). These can work at many levels, from scales of 5-10 grains up to dimensions of millimetres. The possibilities are many, and are limited only by the imagination and practical feasibility, but could include the classic particulate, fibrous and laminated structures familiar to polymer composite technologies. In this case the matrix/reinforcement couples can be metal/hardmetal; hardmetal/metal or hardmetal/hardmetal. As well as bimaterial types, multimaterial contributions are theoretically possible. Good examples of the particulate type are the Fang et al [15] metal/ hardmetal and Newcomer [16] hardmetal/hardmetal structures. Since one can use powder metallurgical techniques to vary the microstructural features on both a microscopic and macroscopic scale over extremely wide ranges the scope for new research is very considerable. These principles have been used recently by Berns et al to develop ferrous-based multi-materials [17] with enhanced hardness/toughness characteristics. The driving force as always is to move the hardness/toughness combination away from the relations that define conventional materials, either by increasing toughness at equivalent hardness or increasing hardness at equivalent toughness [18]. Hybrid and graded materials fall also within this class and here measurement techniques are needed for assessing the spatial variation in mechanical behaviour as well as interpreting the overall response of the structure. In these materials the global mechanical response will be greatly influenced by the anisotropy of local properties, from the level of the WC crystals upwards in scale. Understanding this anisotropy in behaviour and its scale dependence is therefore a fundamental requirement in the development of materials with macroscopic composite architectures.

3 Developments in Testing

The technologies for testing hardmetals continuously evolve. Two test techniques are selectively reviewed to illustrate the development of test methods, the interpretation of test results and their relevance to property maps.

• S-N Fatigue • Abrasive Wear 254 HM 84 B. Roebuck et al.

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This choice follows the necessity to develop test methods that are cost effective and have wide applicability in the selection of suitable materials. Abrasive wear is discussed because a single expression can be derived relating abrasion to hardness that is useful in property mapping. By way of contrast S-N fatigue is discussed because it provides an example of mechanical behaviour that is less straightforward in this respect.

3.1 Fatigue tests

Many different kinds of tests method have been used to investigate the fatigue behaviour of hardmetals including fatigue crack growth, compression, tensile, S-N and contact fatigue modes. However, the most developed method in recent years has been that applied by the Erlangen group [19] using an S-N approach to establish a fatigue life for different hardmetals. However, their apparatus is rather specialised and not obviously amenable to wider use. Consequently at NPL we have investigated the use of V-notched testpieces to obtain S-N data through R > 0 bend tests using a straightforward, more commonly available, bend jig.

Rectangular testpieces are diamond ground and then V-notched [2] at the centre of one of the faces to a nominal depth of 1 mm using a 0.5 mm radiused diamond resin-bonded wheel. The notched testpieces are annealed at 800 °C for 1 h in a vacuum to relax residual stresses. Notched S-N fatigue bend tests are performed in a test jig with a span of 30 mm.

Typical results for two hardmetals are shown in Fig 4 and they confirm correspondence with results from the Erlangen group underpinning the typical S-N type behaviour of hardmetals and the ability to discriminate. However, a key issue is how many tests need to be performed to ascribe a given level of confidence for lifetime prediction. The results of a study using randomly selected data sets of size NR from the total population NT is shown in Fig 5 where it can be seen that a minimum of about 20 tests is needed to confer an uncertainty of about 15% on the slope of the fitted parameters to a linear correlation between maximum stress and cycles to failure Nf. Consequently the derivation of unique constitutive expressions describing S-N behaviour that will effectively discriminate between different materials will not be straight-forward because of the inherently larger scatter in data in fatigue. B. Roebuck et al. HM 84 255

15" Internaüonal Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3200 3000- A NPL11M,R=0.1,f=10Hz • NPL6F, R=0.1,f=10hz 2800- Erlangen 6Co+cubics, R=-1, f=22Hz

"E 2600 -

z 2400 - u> 2200 - I 2000- I 1800- E * 1600- 1400- 1200- "1— ITTTTI—i i iuiii|—i 11""i|—i i """|— 10000.1 1 10 100 1000 10000 100000 1000000 1E7 Cycles to failure

Fig 4 Typical notched bend fatigue test results.

1 35- \ —A— Slope jo 30- \ o 25- V •ia t 5 20- "o \A CD ID Ö \4 sp § 10- Ü 5-

0- 10 20 30 40 50 Sample size

Fig 5 Uncertainties in the S-N slope relation parameter. 256 HM 84 B. Roebuck et al.

15lh International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3.2.2 Wear

Assessment of wear resistance is a very important issue from the point of view of end use. However, different wear mechanisms may be appropriate, of which abrasion and erosion are the two most relevant.

The literature on wear behaviour generally fall into two schools

• where there is believed to be a monotonic dependence on hardness [20,21]

• or where the wear behaviour is thought to be dominated by mechanism changes, influenced either by random events in binder and carbide at low values of hardness and bulk deformation mechanisms at higher hardness [22,23].

Other mechanisms caused by oxidative or chemical effects, can also influence the interpretation of test results. Abrasion and erosion test methods are being evaluated at NPL both for their repeatability and reproducibility in order to quantify uncertainty. Quantifying uncertainties gives added confidence in the ability to discriminate materials. Also quantitative relationships are needed between wear parameters and microstructures to underpin the use of property maps. Some of these issues are illustrated with Fig 6 which shows results of the ASTM B611 abrasion test at NPL compared with published data by the University of Witwatersrand group [23]. The results allow an expression to be developed that is applicable to typical grades of commercial WC/Co currently available from HV800-HV2200 such that:

A = aexp[-bH] (1) where A is abrasion resistance in vol loss (cm3), H is hardness (HV30) and a and b are constants. It was found that the repeatability of measurement was less than size of the symbols on the graph. The agreement between the NPL and the University of Witwatersrand group [23] indicates that the reproducibility of the test is also excellent. This single expression is clearly useful in underpinning the development of property maps. B. Roebuck et al. HM 84 257

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

1 - ^4>^ = :\^,d = L::::::::::|:::::::::::::j ::::::::,! ,::::::!::::::::::::::

; ->J- "_S£^k^ t

• :^:::4::::::::::|::::::::::j::::::::::[:::::.:::::::|::::::::::::

o &g> ; ; \ ; 0.1 - o >

< 0.01 - O NPL NPL linear fit • * Witwatersrand group I j A = aexp(-bH) I 6\" :

400 600 800 1000 1200 1400 1600 1800 2000 2200 2400 Hardness HV30

Fig 6 ASTM B611 Abrasion results

4 Property Maps

A structural characterisation exercise was conducted on a set of WC/Co hardmetals with a range of grain sizes and Co contents from 6-25 wt% [24], The measurements of structure were then analysed for comparison with model predictions from an NDE measurement of magnetic coercivity. As well as coercivity the property characterisation exercise consisted of a measurement of hardness, toughness (by the Palmqvist method) and abrasion resistance. The measurements were used to construct the "property maps". In principle, these maps allow the effects of differences in structure and properties to be more easily compared and can provide a benchmark for the evaluation of new materials.

Two types of property map were considered:

• Structural dependence • Property comparisons 258 HM 84 B, Roebuck et al.

15:n International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

In the first type the property of interest is plotted against a microstructural feature that controls that property. Two microstructural features are discussed, WC grain size, d, and wt% cobalt content, W. It is recognised that other parameters are important, such as the composition of the cobalt binder phase, the cobalt mean free path or the size distribution of WC grains. However, for this paper the concept of "property maps" is developed assuming a composition approximately in the centre of the two-phase WC/Co region and that there is a conventional WC grain size distribution with the arithmetic mean number intercept, d, as the parameter characterising WC grain size.

Four properties were evaluated to compare with microstructure:

Coercivity, K Abrasion Resistance, A Hardness, H Palmqvist Toughness, T from which eight property maps based on microstructure can be constructed; K vs d, K vs W, H vs d, H vs W, A vs d, A us W, T vs d and T vs W.

Data from measurements on the baseline WC/Co materials [24] were used to generate equations relating each of the four properties to d and W. These equations were physically-based where possible.

In the second type of property map each of the three mechanical properties were compared against each other in pairs to give three maps:

T vs H AreH A vs T

4.1 Microstructure property maps

Grain size is usually related to coercivity using an inverse expression.

K = a + b(1/d) (2) where a and b are constants, K is the coercivity in kA m"1 and d is the WC arithmetic mean linear intercept size in urn. Initial work found values of 1.79 for a and 10.9 for b for WC/6wt% Co hardmetals with values of K less than about 15 kA m"1. This expression did not include a parameter for the variation of Co content and this is needed for comprehensive property mapping of typical hardmetals. Additional work [24] allowed expression (2) to be modified B. Roebuck et al. HM 84 259

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to give the variation of coercivity, K, with grain size, d for 6 < Co < 25wt% as follows:

K = (ci + dn * wt% Co) + (c2 + d2 * wt% Co) * (1 Id) (3) where Ci = 1.44 di = 0.04 c2 = 12.47 d2 = -0.37 with K in kA m"1 and d in |jm.

Expression (3) can be plotted as a property map, Fig 7, bearing in mind the constraint that it has not been extensively validated for high values of coercivity equivalent to WC grain sizes (arithmetic linear intercept) less than about 0.5 urn. Additional work is needed to evaluate finer grained materials. A property map using Co wt% content as the variable can also be plotted from expression (3). This is also shown in Fig 7 for values of linear intercept WC grain size of 5, 2, 1, 0.75 and 0.5 urn, respectively.

It is also necessary to be aware that coercivity also varies with carbon content and W content of the binder-phase. It has been assumed that the compositions are approximately in the centre of the two phase field. Further work would be needed to refine expression (3) to allow for this effect. It is also possible that microstructural instabilities in low C hardmetals or residual stresses could affect measurements of coercivity through changes in internal strain and composition. These effects contribute added uncertainty to the correlation between grain size and coercivity and need to be systematically examined [25]. it is clear from inspecting the property maps in Fig 7 that the coercivity measurements are fairly insensitive to changes in intercept size as the mean value increases beyond about 3 urn.

Hardness can be related to the inverse square root of the intercept size using the Hall-Petch expression [25].

H = e + fd"05 (4) 260 HM 84 B. Roebuck et al.

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24 WC'Co Property map 22- Cccwil>'' Co ccnlent 5.2,1.0.75.0.5 ^m WC line WC/Co Properly Map 20- Coercivity / WC grain size 5,10,15,20,25% Co 18- Üiii 11 • I \ 16-

14-

12-

10-

2 3 4 5 6 10 15 WC linear intercept um Cobalt wt%

Fig 7 K us d and K vs W property maps.

WC/Co Property Map 2200- 0.5 . WC/Co Property Map i Hardness/ Co content Hardness / WC grain size \ 5.2.1.0.75,0.5 urn WC 6,10,15.20.25% Co 2000- 0.75 A 1800- 1 •

> 1600- ^\ « 2 * B 1400- \ \\\ 5 \ \\\ r 5 * 1200- 10 15~'~ 1000-

800-

12 3 4 5 10 15 WC linear intercept [im Cobalt wt%

Fig 8 H us d and H us W property maps. B. Roebuck et al. HM 84 261

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 where e and f are constants, H is the hardness, HV30, and d is the WC arithmetic mean linear intercept size in urn. Further analysis allowed an expression to be written for the variation of hardness, H, with intercept size, d, for 6 < Co < 25 wt% as follows:

H = (888 - 9.9' wt% Cc) + (229-532'exp(^(wt%Co-6y6.7) (5)

Expression (5) can be plotted as a property map, Fig 8, bearing in mind it is not validated extensively for ultrafine grained materials. There is some evidence that the Hall-Petch expression is inaccurate for the finer grained materials. Further work is needed to assess materials with WC intercept size less than about 0.3 urn.

A property map using Co wt% content as the variable can also be plotted from expression (5): This is also shown in Fig 8 for different values of the WC linear intercept grain size.

Abrasion resistance, A, is dependent on hardness, H, and it was found [20] that:

log A = g + hH where g and h are constants, A is in cm3 and H is the HV30 value,

however H = e + fd"05

thus log A = i + kd"°5 (6) where i and k are constants, and d is in urn. Analysis of the data on the baseline hardmetals [24] allowed an expression to be written for the variation of abrasion resistance, A, with grain size, d, for 6 < Co < 25 wt% as follows:

2 75 a133 %C a0022 (%CO lo9leA = (0.19-0.016'%Co)-' -° - ' r ' » (7) Vd

Toughness can be related to the Co mean free path [26]. The Co mean free path is linearly related to WC intercept size for materials with conventional WC 262 HM 84 B. Roebuck et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

grain size distributions [3]. Therefore the following expression was used to evaluate the Palmqvist toughness results for property mapping:

T = m + nd (8) where m and n are constants. Analysis allowed an expression to be written for the variation in Palmqvist toughness with intercept size as a property map:

T = 8.8 + (10A (-0.33 + 0.088*% Co) )*d (9)

This is probably the least well developed expression in this group because of the difficulty of obtaining extensive sets of valid toughness data for WC/Co.

4.2 Comparative property maps

The expressions developed in section 4.1 for the microstructure property maps can be used to derive comparative property maps. Two examples are shown in Fig 9 (T vs H) and Fig 10 (A vs H). Other combinations are possible, but these are good examples because, industrially, hardness is used quite frequently as the property with which other properties are compared. Also the maps can be plotted for different representative Co contents and for different WC grain sizes.

The first comparative map, Fig 9, shows the relationship between Palmqvist toughness and hardness. The second comparative map, Fig 10, shows the relation between abrasion (vol loss) and hardness. There is an equal correspondence between the properties, and changing either Co content or grain size does not show significant deviation from the overall trend. There is some uncertainty at the high toughness values however due to the difficulty in making accurate measurements of toughness at values greater than about 20 MN m"3/2. B. Roebuck et al. HM 84 263

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

WC/Co Property Map 45- Palmqvist toughness / Hardness

40- 6,10,15% Co

35-

30- v> CD c -C 25-

o 20- > cr 15-

10-

5 800 1000 1200 1400 1600 1800 2000

Hardness HV30

Fig 9 Ti/sH property map.

0.1-

WC/Co Property Map Abrasion / Toughness 0.01 -. 6,10,15% Co

—•—6 -A-10 -T-15

1E-3- 6 9 12 15 18 21 24 27 30 33 36 39 42 45 48 -1.5 Palmqvist toughness MN m " 264 HM 84 B. Roebuck et al.

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

5 Summary

There is an increasing trend in the last 10y to investigate unconventional hardmetal structures, for example, with changes in the shape and distribution of the hard and binder phases, with the physical metallurgy of the different phases and with the architecture used to construct the macroscopic appearance, i.e. graded and particulate structures. Thoroughly well characterised property maps for conventional hardmetal structures are vital in order to fully quantify the changes in properties associated with these unconventional structures.

There are an immense number of structural variants to investigate. Thus, there are clear opportunities for the use of microstructural modelling approaches to accelerate the development process. It is vital to understand that many of these innovative structures will be anisotropic in their properties. It is crucial therefore that more understanding is developed of the effects of this anisotropy, both on a macro and a microscopic level. Modelling studies will be especially needed to develop relations between feature distributions, shape and chemistry and to relate predictions from these studies to mechanical property data such as strength, hardness and toughness.

6 Acknowledgements

The technical work at NPL reported in this paper was performed as part of a programme of underpinning Materials Measurement research financed by the UK Department of Trade and Industry. Members of the British Hardmetal Association Research Group are thanked for the supply of materials and the results of in-house characterisation tests and comments on the development of property maps. Constraint on space has not allowed full reference to all the relevant literature. The global hardmetal research and business community is therefore fully acknowledged. B. Roebuck et al. HM 84 265

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7 References

1. B Roebuck, M G Gee, E G Bennett and R Morrell. Mechanical Tests for Hardmetals. NPL Good Practice Guide No 20. August 1999.

2. B Roebuck. 14th Int. Plansee Seminar, Reutte, Austria, May 1997, V2, 352-365.

3. B Roebuck and E G Bennett. Metallography, 1986, 19, No 1, 27-77.

4. I I Timofeyeva, M A Vasilkovskaya and S M Radic. Science of Sintering, 1999,31(1), 17-22.

5. L Prakash. 13th Plansee Seminar, Reutte, Austria, 1993, V2, 80-109.

6. M Kobayashi and S Yamaya. 10th Plansee Seminar, 1981, V1, 597-611.

7. A V Shatov, S A Firstov and I V Shatova. Mater. Sei. and Eng., 1998, A242, 7-14.

8. P Klar, F Kiefer, K Stjernberg and J J Oakes. PM2TEC99, Int. Conf. on Powder Metallurgy, Vancouver, MPIF, 1999.

9. K Kitamura, M Kobayashi and K Hayashi. 15th Plansee Seminar, Reutte, Austria, 2001, V2, 279-293.

10. A Schon, W D Schubert and B Lux. 15th Plansee Seminar, Reutte, Austria, 2001.

11. B Roebuck. Metal Powder Report, April 1999, 20-24.

12. V N Glushkov, L I Klyachko, V A Fal'kovskii, O N Eiduk and N M Lukashova. Russian J. of Non Ferrous Metals, 1998, 39(8), 49-51.

13. J M Missiaen, A Lackner and L Schmid. 15th Plansee Seminar, Reutte, Austria, V2, 252-266.

14. A Laptev. PM World Congress 1994, Editions de Physique, 1994, V1, 103-106. 266 HM 84 B. Roebuck et al,

15!" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

15. Z Fang, G Lockwood and A Griffe Metall. Trans. A, Dec 1999, 30A, 3231-3238.

16. K J A Brookes. Metal Powder Report, April 1995, 14-17.

17. H Berns, A Melander, D Weichert, N Asnafi, C Broekmann and A Groß-Weege. Computational Mater. Sei., 1998, 11, 166-180.

18. R Raj and L R Thomson. Acta Metall. Mater., 42, 1994, 4135-4142.

19. U Schleinkofer, H G Sockel, P Schlund, P Kindermann, R Schulte, J Werner, K Gorting and W Keinrich, Euro PM96, Advances in Hard Materials Production, Stockholm, May 1996, 239-246.

20. M G Gee, L Fang and B Roebuck. Euro PM99, Turin, 8-10 Nov 1999, 319-326.

21. J Y Sheikh-Ahmad and J A Bailey. Wear, 225-229, 1999, 256-266.

22. V A Pugsley and C Allen. Wear, 225-229, 1999, 1017-1024.

23. D G F O'Quigley, S Luyckx and M N James. Int. J. of Refractory Metals and Hard Materials, 1997, 15, 73-79.

24. B Roebuck, E G Bennett, W P Byrne and M G Gee. Characterisation of Baseline Hardmetals using Property Maps. NPL Report CMMT(A)172, February 1999.

25. B Roebuck. Terminology, Testing, Properties, Imaging and Models for Fine Grained Hardmetals. Int. J. Refractory Metals and Hard Materials, 13,1995,265-279.

26. L S Sigl and H E Exner. Mater. Sei. and Eng., A108, 1989, 121-129. K. Friedrichs HM 88 267

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Transition From HSS Tools to Solid Carbide Tools in the Field of Shaft Tools and the Development of New Types of Carbides by

Dr. Konrad Friedrichs Konrad Friedrichs carbide factory Kulmbach, Bavaria, Germany

Summary:

The development trends for shaft tools made of solid carbide are presented. The reasons for the rapid acceptance of solid carbide tools in areas formerly reserved for high speed steel tools are explained. The development of tougher types of carbide and more modern coating procedures as well as the use of high speed cutting machines with higher efficiency is examined especially. Modern carbide types with 10% and 12% cobalt and higher toughness are described. The influence of reducing the average WC grain size is emphasized. Finally, advanced extrusion processes with screw extrusion presses are explained. They cover a large market segment of the carbide rod production today. The possibilities of finer shaping close to the form of the final product are described.

Keywords:

Extrusion presses, presses for multiple rods, ultrafine grain sizes 268 HM 88 K. Friedrichs

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Introduction:

In recent years, there has been a strong increase in the use of shaft tools made of solid carbide worldwide. This relates to shaft millers and drills with and without internal cooling ducts as well as reamers and tap drills. While the solid metal tools initially increased strongly in the diameter range from 0.1 mm and 12 mm, an increase in the diameter range from 16 mm to 32 mm can also be observed today. The increased use of solid carbide shaft tools actually began only 20 years after the boom in turnplates and the related turnplate holders.

There are several reasons to change rapidly to the use of carbide shaft tools

a) Tougher types of carbide have been developed, which nevertheless have the same hardness as before. Through their higher toughness, fields which were previously reserved for HSS have been conquered.

b) Economical coating processes which increase the lifetime and range of possible applications of solid carbide tools have been brought onto the market.

c) Manufacturing machines have been developed, which have the required performance, speed and stiffness necessary for the use of carbide tools. These machines were previously not available in a large selection. These new machines are forcing the replacement of the earlier high speed steel tools in favor of carbide tools.

d) New areas of use have been added for carbide tools, which were not covered adequately by high speed steel tools. For example, machining of fiber-reinforced plastic Machining of titanium alloys Machining of stainless steel Machining of hardened steel K. Friedrichs HM 88 269^

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e) In many cases, cutting work can be performed more quickly and more economically than with high speed steel tools. This leads to a reduction of the investment with the same production quantity. The plants become more flexible through the higher manufacturing speed. The quality of the products produced also rises because the e-modulus of the solid carbide tools is higher than that of HSS tools. f) A further important reason for the sharp increase in solid carbide shaft tools is the provision of adequate production capacity. It is estimated that there was a production capacity of approx. 13,000 to/year for carbide rods in the year 2000. The Kulmbach region in northern Bavaria accounted for 900 to/year alone. The expansion of the worldwide production capacity is not yet at an end and is increasing, so the use of HSS tools will be on the decline in the foreseeable future. 2!70 HM 88 K. Friedrichs

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2.0 New Developments in the Field of Tough Types of Carbides for Shaft Tools

For years, carbide manufacturers have made efforts to raise the toughness of carbide grades without reducing their high, often desired hardness and resistance to wear. For many years, there were two main types on the market, namely the K types and the P types, whereby the K types were used mainly for cast iron machining and the P types for steel machining. Through the development of coating technologies, the main applications of the carbide types shifted more and more into the K area, meaning into types with very low TaC content and no TiC content. With the coated types of carbide, steel was then machined successfully as well. A so-called 6% was dominant in the K area for a long time, meaning a carbide type with about 6% cobalt content, an average WC grain size of 1.5 urn down to 1.0 urn sometimes. In addition, they were dosed with VC and TaC in low quantities (approx. 0.2% VC and 1% TaC) to avoid grain growth during the sintering process. Every manufacturer of carbide had this type (see Figure No. 1) in its program in the 70s and 80s. K. Friedrichs HM 88 271

15lh International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Figure No. 1: grain size structure of a 6 %er

A variation of this was a 6% with a then very fine WC grain of approx. 0.7 (am to raise the hardness and resistance to wear. This type was used primarily in the electronic drill and the reamer area.

The size of the TaC crystals (> 2um), which led to early cutting break-outs and a reduction of the toughness was extremely annoying with these K types. The so-called RAMET types from an American manufacturer at the end of the 70s and the beginning of the 80s were the first successful types in the K area with Cr3C2 in place of TaC. 272 HM 88 K. Friedrichs

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At the end of the 80s and the beginning of the 90s, all important carbide manufacturers began to develop a Ktype for rod manufacturing. It contained 10% cobalt and had an average grain size of 0.65um to 0.8um.

CR3C2 and some VC were used as doping material. See Figure No, 2.

Figure No. 2: grain size structure of a 10 %er

This 10% type has high hardness and resistance to wear and thereby exhibits considerable toughness as well. This 10% also belongs to the main standard carbide types at almost all important rod manufacturers. K. Friedrichs HM 88 273

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The meaning of the Cr3C2 and VC content was recognized more and more. Variations of the proportion and the quantities of Cr3C2 and VC exhibited significant influence on hardness and toughness.

In the mid-90s, 12% came onto the market. The tools of a Japanese manufacturer were distinguished because of their sharply increased resistance to wear and toughness as well (Kobelco). These carbide types had a significantly higher lifetime in a series of applications with high loading on the tool (diesink milling, machining of hardened steel, machining of stainless steel, milling of titanium alloys, application in aluminum alloys, drilling in fiberglass reinforced plastics). Figure No. 3 shows the grain structure of the 12% at 1500 times magnification. This 12% distinguishes itself through a further reduction in the grain size of the WC. It amounts to 0.5um. Furthermore, the proportion of Cr3C2 and VC was increased again in comparison to the 10%. Further variations were also made in the proportion of the two materials. This type is gaining more and more in importance and will therefore replace the 10% types increasingly. 274 HM 88 K. Friedrichs

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Figure No. 3: grain size structure of a 12 %er with 0.5 micron WC

The grain structures of a 10% and a 12% are to be seen in comparison in Figures No. 4 and No. 5 respectively at 10,000 times magnification. K. Friedrichs HM 88 275

15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

K40UF * 10% Co « 0,65

SEM micrograph 1 10 000 .

Figure No. 4: 10 %er with magnification 10.000 times 276 HM 88 K. Friedrichs

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Figure No. 5: 12 %er with magnification 10.000 times

The transition from an average grain size 0.65u.m to an average grain size 0.5u.m can be seen clearly.

A further development is therefore the transition to types with an average WC grain size of 0.2um, which has been taking place recently. K. Friedrichs HM 88 277

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While the term ultrafine is used for types with 0.5um grain size, the expression superfine is used to distinguish the 0.2um types. In Figure No.6, the grain structure of a superfine carbide can be seen at 10,000 times magnification.

i . "• V. »»*• t '~

'r T v> "^ ?' "*" * »

Figure No. 6: superfine carbide grade with magnification 10.000 times

At the present time, these types are used primarily in the field of electronic drills and PC drills. Hardnesses of 2000 Vickers (93.2 Rockwell) and bending stiffnesses of 4000N/mm2 are normal here. The enormous toughnesses of these materials with considerable hardness are to be emphasized. At the present time, these types of carbide are manufactured with approx. 8% and 9% cobalt. The Cr3C2 and VC contents amount to 1.2 % together. The development of these types will also be followed logically in types with higher Co content (approx. 12%) to achieve further increases in toughness, compared to the old 12% with 0.5|am. The hardness should retain its high level more or less. 278 HM 88 K. Friedrichs

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A summarized presentation of the most important K types is given in Figure No. 7, whereby the 12% and the types with 0.2um average grain size are to be regarded as the latest development of two carbide types for rods at the present time.

Konrad Friedrichs KG HARTMETALLFABRIK KULMBACH

Typical carbide grade recipies in the K-grades

6% Cobalt 1 % TaNbC 0,15% VC 92,85 % we

8% Cobalt 1 % TaNbC 0,2 % VC 90,8 % we 0,7|am

10% 10% Cobalt 0,3 % CraC2 0,2 % VC 89,5 % we

12% 12% Cobalt 0,9 % CraC2 0,4 % VC 86,7 % we 0,5[jm 9% 9% Cobalt 0,63 % O3C2 0,55 % VC 89,82 % we

TlV

Figure No. 7: summarize of important k grades

The most important physical and chemical properties of the types described can be seen in Figure 8. K. Friedrichs HM 88 279

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Konrad Friedrichs KG

WC Co va DeraHy, HareIros Mocpefc Coercive Transverae P«xjsily, c ain Oosiificdtion, rard mefcrJ I ISO HR4, W30, xjhjration fieldsJreng* ISO 450 ittuclure, ISO 513 3349 ISO ISO 4K a HC, slrengiK ISO U99 J

medium/] j. KFK 10/20 92,5 6,5 1 U.8 92,0 1440 12 18 2500 £02 00,00 K10-K20

KFK10F 92,5 4,5 1 14,8 92,7 1810 12 24 3100 £02 00 j 00 tn Fine K01K10

KFK20F 90,5 8 1.5 14,4 92,2 1710 15 23 320O £02 00 00 t« 1 fire K10-K30

KFK40UF 89,5 10 0,5 14,5 91,7 1410 IS 21 3600 £02 00 00 fre - K30-K40

KFK44UF 84,7 12 1,3 14,1 92,1 1460 19 24 3800 £02 00 00 f™ ' " K4O-K5O 1 1 KF K 45 UF 85,7 13 1,3 14,1 91,5 1400 21 24 3900 £02 oo; oo fee - K4O-K5O

KFK55SF 1 89,8 9 1,2 14,3 93,2 2000 14 39 4000 £02 00 00 fir» - -

Figure No. 8: physical and chemical properties of important k grades

The high bending rupture value and the hardness value at the same time of the carbide type with 0.2um average grain size of the WC are remarkable. These types of carbide metal are getting closer and closer to HSS because of their toughness, even if the distance is still relatively large. It is important to note that, in spite of higher hardness, a large bending rupture strength has also been achieved. A schematic representation of the relationship between hardness and toughness is shown in Figure No. 9. 280 HM 88 K. Friedrichs

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Figure No. 9: relationship between hardness and toughness

Every one of the three carbide types can be varied with respect to its relationship between hardness and toughness by, for example, changing the proportion of Cr3C2 and VC or the total quantity of C43C2 + VC. This means that sub-types for special applications can easily be manufactured from one type of carbide. A change in the grain size can also lead to a different carbide with all other parameters fixed as can be seen in Figure No. 9. We will not have a long wait until interesting developments, especially in this field of the variations, are available. K. Friedrichs HM 88 281^

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3.0 Production Methods of Carbide Rods Which Are Used for Cutting Tools

The methods of cutting rectangular rods from cold isostatic pressed blocks and making them into round rods through grinding processes has long been out-of-date. Rods are pressed today, for example, from granulate between upper and lower stamps and matrices horizontally or vertically, whereby the rod lengths are very limited. Lengths which correspond to the standard dimensions of catalog tools are produced, whereby the shrinkage during sintering must be taken into account.

Another, very frequently used procedure today is the pressing of round rods of carbide powder in thick-walled, well-closed polyurethane tubes with water pressure from the outside. The pressure then amounts to approx. 2000 bar.

The most frequently used process is the extrusion process to press out rods of all cross-sections. The length of the rods is almost unlimited thereby and depends only on the type of extruder used. A distinction is made between piston extruders and screw extruders. Screw extruders can work continuously, given the availability of enough mass to be pressed out and, if necessary, produce rods kilometers long if the required cutting and stacking fixtures are available behind the presses. Piston extruders can work without interruption only as long as there is enough material contained in the cylinder in which the piston works. The cross-sections of the rods to be pressed depend upon the plastification agent used. Plastification agents which permit the extrusion of rods with a maximum diameter of 34 mm (sintered diameter) today have been developed by the author. The problem with all plastification agents is that, although they are very necessary with extrusion pressing, they interfere very much with the sintering process later. 282 HM 88 K. Friedrichs

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These plastifiers must be removed before the sintering, at best before the rods reach the sintering oven. The cheapest and simplest process at the present time is a drying process in simple air drying cabinets, during which the majority of the plastification agent is dried out. The rest remaining permits rapid, problem-free sintering without the risk of cracking. Screw extruders have been developed especially for the production of rods with internal cooling ducts. They allow a very exact and precisely-defined pitch of the spiral cooling ducts for drill bits with small diameters and relatively large lengths. Figures No. 10 and No. 11 show a modern screw press with cutting fixture and transport table as well as line camera and monitor to control and adjust the angle of the pitch, based on a patented process. K. Friedrichs HM 88 283

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Figure No. 10: Screw extruder for spiral coolant hole rods 284 HM 88 K. Friedrichs

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Figure No.11: patented procedure to produce helical twisted rods

This type of press can also be used very economically for the production of rods with parallel cooling ducts as well as solid rods without cooling ducts. The plastification agent, which is to be dried out before the sintering, ensures that rods which can be subjected very easily to further shaping steps after the drying process and before the sintering are produced.

As an example, increasing numbers of rods are manufactured, which are already prefluted. Prefluting before sintering, which is very cheap, helps to avoid high costs for the grinding of flutes after sintering. Cost savings of more than 30% in the production of ground tools are no rarity.

Figures No.12 and No. 13 show prefluted, unsintered tool blanks with remarkably sharp contours of the later cutting edges. K. Friedrichs HM 88 285

15:h International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

/ I

Figure No. 12: prefluted rods on sintering plates

Figure No. 13: prefluted rods with clean edges 286 HM 88 K. Friedrichs

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No break-outs may be allowed to occur during hand shaping. A further growing program of sintered metal use is the field of paper knives and wood knives. Soft paraffin is used primarily as the plastifier for the extrusion of small cross-sections in screw extruders. The pressed molded parts are then stepped, although the cross-sections are small. With these small cross-sections, paraffin removal takes place in the sintering oven with no problem.

A typical press for the extrusion of rods of all kinds with low cross-sections in shown in Figure No. 14. Presses of this kind have automatic cutting and transport fixtures for the rods and graphite carriers.

Figure No. 14: automatic press for rods with small cross-sections

Rods for the manufacturing of electronic drill bits have also experienced an upturn in recent years. K. Friedrichs HM 88 287

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Especially in the growth field of mobile telephones and personal computers, so many soldered points with the required drillings are necessary that the demand for such drill bits has risen sharply here. These are drill bits with small diameters. Diameters of 0.1 mm are no rarity thereby. For a long time, drills were manufactured from carbide rods with a blank diameter of 3.25 mm. Nowadays, more and more bits which have a steel shaft with a diameter of 3.175 mm, into which a carbide pin of about 1.4 mm diameter is shrunk, are manufactured today. This pin is then ground down to the necessary outside diameter and given the spiral slots and front teeth.

The development of carbide types with finer and finer grain sizes was described in the previous chapter. The production of rods with a blank diameter of 1.6 mm, for example, has been rationalized increasingly. For this purpose, presses which do not always extrude only one rod at a time, but rather several rods simultaneously have been developed. With the small diameters, this is necessary to obtain reasonably economic extrusion capacity. The next figure (Figure No. 15) shows a press which extrudes and cuts 18 rods at a time and places them on the graphite carrier. 288 HM 88 K. Friedrichs

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Figure No. 15: press for multiple rods in the same press cycle

The growing market for PC drills can be covered adequately only in this way. K. Friedrichs HM 88 289

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4.0 Outlook

The triumph of carbide has not come to an end. With increasing toughness and at least the same hardness, more and more areas which were previously reserved for HSS are being captured. Special development trends are the use of finer and finer WC types, whereby variations of the cobalt content are made to influence the toughness. Increasing attention is also devoted to the content of Cr3C2 and VC and the proportion of the two to each other. The pressing process and tools are being rationalized further and further, whereby attention must be paid especially to screw extrusion processes with automation aids. The use of multiple nozzles will be forced. 290 HMJ38 K. Friedrichs

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5.0 Bibliography

Friedrichs, K. A new technology for the production of carbide rods with helical twisted coolant holes. Proceeding of the 13th International Plansee Seminar, Reutte (1993) Vol. 2, HM48.

Friedrichs, K. Extruder twisting processes for internal cooling channels in carbide rods. Proceeding of The Powder Metallurgy World Congress PM'94, Paris, June 6-9, 1994, Vol. 1, page 153 ff, ref. 140A21A.

Friedrichs, K. A new plastification agent for the extrusion moulding of sintered carbide rods with cooling channel borings. Proceedings of the 14th International Plansee Seminar, Reutte (1997) Vol. 2, HM57 AT0200483 J. Zinyana et al. HM 89 291

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EFFECT OF COMPOSITION ON THE (V,W)C CONSTITUTION IN V-W-C-Co ALLOYS.

J. Zinyana1, S. Broccardo1, S. Hamar-Thibault3, LA Cornish1, MJ Witcomb2, C. H. Allibert3 and S. Luyckx'

1 School of Process and Materials Engineering and 2Electron Microscope Unit, University of the Witwatersrand, Johannesburg, South Africa 3LTPCM (UMR 5614 CNRS-INPG), INP Grenoble, France

SUMMARY:

Fifteen W-V-C-Co alloys varying in composition were sintered at 1450°C for 5 hours and quenched. The sintered alloys consisted of (V,W)CX particles in a Co-rich matrix, with a VC- or WC-rich eutectic present in most alloys. The (V,W)CX particles were analysed by XRD, EDS and EPMA and their microhardness was related to their composition and structure.

1. INTRODUCTION:

In recent years much work has been done to assess the potential of cemented carbides where a substantial amount of the tungsten carbide of conventional WC-Co alloys has been replaced by vanadium carbide (i.e. WC-(V,W)Cx-Co alloys, with x always < 1 ) [1]. It has been found that the materials are suitable for applications in abrasive and corrosive environments on account of their hardness and corrosion resistance being superior to those of conventional WC-Co of equal cobalt content [1,2].

However, these materials, similarly to conventional WC-Co alloys, present the problem of the possible precipitation of graphite or eta phase when the carbon content is above or below the stoichiometric amount. Thus the optimization of the materials has required a systematic investigation to determine the composition range at which no precipitation of graphite or eta phase occurs [3]. The result of the investigation was that the composition range suitable for the formation of 3-phase WC-(V,W)Cx-Co alloys is a narrow strip located along the plane defined by the relationship

C/(V+W) = 1 - 0.33 (V/(V+W)) for 0 < V/(V+W) = 0.72 (1) with graphite or eta phase expected to precipitate at compositions corresponding to higher or lower values of C/(V+W) [3]. 292 HM 89 J. Zinyana et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The (V,W)CX phase has been found [3] to vary in composition with varying the overall composition of the alloy. Therefore the optimization of the materials requires also the determination of the (V,W)CX composition which corresponds to optimum properties. Although for the purpose of optimizing the materials one would need to study only (V,W)CX compositions within the strip defined by equation (1), the study of (V,W)CX has been extended beyond these compositions. The work is in progress, but the results available to date are summarised in this paper.

2. METHOD:

Fifteen WC-(V,W)Cx-Co alloys have been studied. They are part of a set of alloys prepared for the investigation reported in reference [3]. They were prepared by mixing in a ball mill for up to 48 hours VCX (V8C7), WC, Co and, in seven cases, C powders, bringing the mixtures to 1450°C, keeping them at temperature for 5 hours and quenching them at the highest possible cooling rate.

The cobalt content of the alloys was kept constant at 50 at.%. The high Co content was selected to allow the formation of each phase without impingement from other phases. The carbon content was intended to be constant at 27.8 at.% in seven of the alloys (high- carbon or HC alloys) and at 22.8 at.% in eight of the alloys (low-carbon or LC alloys), so that the (V+W)/C ratio would be 0.8 and 1.2 respectively and the substitution of V with W (in VCX) was systematically increased.

The alloys were examined by scanning electron microscopy (SEM) at 20 kV, including backscattered electron microscopy (BSE), and analysed in South Africa [4] by energy dispersive X-rays (EDS) for 100 s using pure metal standards, and in France [3] by electron probe microanalysis (EPMA). The EDS system available could not detect carbon but was used to determine V/(V+W) ratios. Each V/(V+W) ratio reported below represents the average of measurements from at least 5 randomly chosen particles. Carbon contents were determined by EPMA, each reported value being the average of 10 measurements. The samples were also analyzed by X-ray diffraction (XRD), using Cu radiation despite the large Co content, to determine the phases present and the lattice parameter of (V,W)CX in each alloy. The (V,W)CX peak intensity ratios were also measured to deduce possible variations in structure or composition.

The microhardness of (V,W)CX was measured at 50g. The values reported below represent the average over measurements from at least 3 randomly chosen particles. J. Zinyana et al. HM 89 293

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3. RESULTS:

The XRD analyses results showed that all samples contained (V,W)CX and Co, but only the W-rich samples contained also WC. In addition, all the HC (high carbon) alloys contained graphite. The HC alloys contained both fee and hep Co, while the LC (low carbon) alloys contained only fee Co. No residual VCX was found in any of the alloys.

SEM and EDS examinations established that the alloys consisted of large round particles of (V, W)CX and fine WC grains (if present) in a continuous Co-rich matrix. Most of the alloys exhibited also a VC-rich or WC-rich eutectic and the graphite in the HC alloys precipitated in the form of flakes. The (V,W)CX particles were up to 100 \xm in size, i.e. large enough for Vickers microhardness measurements at 50 g load (HV0.05).

Table I summarises the V/(V+W) ratios obtained from EDS and EMPA analyses of the (V,W)CX particles in each alloy. The agreement between EDS and EMPA results was excellent for most alloys. The standard deviation was generally low, which indicates that a good degree of homogeneity had been achieved in most cases. Table I lists also the (V,W)CX formulae deduced from the analyses' results. Table I also indicates whether WC or graphite where present in the alloy.

Figures 1 and 2 show that the lattice parameter of (V,W)CX decreases with increasing V/(V+W) ratio (i.e. with decreasing W content) and C/(V+W) ratio respectively. The V/(V+W) data in Fig. 1 are those obtained from EDS analyses and the C/(V+W) data in Fig.2 are those obtained from the EMPA results which were in agreement with the EDS results. The lattice parameters in Figs. 1 and 2 are the average values obtained from the d- values of the three most intense XRD peaks of (V,W)CX, assuming (V,W)CX to be cubic in all cases. Figs 1 and 2 show that for equal V/(V+W) and C/(V+W) ratios the lattice parameters of (V,W)CX in C-rich (HC) alloys are consistently larger than in C-poor (LC) alloys.

Figures 3 and 4 show that the microhardness of (V,W)CX varies with both the V/(V+W) and the C/(V+W) ratios. The scatter in the microhardness results is large (mean standard deviation: ± 150 HV0.05) because the measurements were carried out on random cross sections of (V,W)CX particles (thus there may be an anisotropy effect) and because it is difficult to measure accurately the diagonals of 50 g indentations, although the diagonals were measured on SEM micrographs at high magnification. Despite the scatter, the results show that the microhardness appears to reach a maximum value at V/(V+W) between 0.8 and 0.88. Both Figs 3 and 4 show that the microhardness of LC alloys was consistently higher than the microhardness of HC alloys of equal V/(V+W) and C/(V+W) ratios. 294 HM 89 J. Zinyana et al.

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The first set of results in Table II reports the approximate relative intensities of the three most intense XRD peaks of (V,W)CX as obtained from a first set of analyses. The second set reports results from a second set of analyses on the same samples, carried out after rotating the samples.

4. DISCUSSION AND CONCLUSIONS:

The results reported above confirm that the mixed carbide (V,W)CX can exists in a Co-based matrix in a range of compositions.

The monotonic behaviour of the results in Fig. 1 (which agree with earlier results by Rudy et al [5]) and in Fig. 2 suggest that (V,W)CX is cubic at all compositions, as it was assumed when calculating the lattice parameters. The differences in relative peak intensities at different compositions (Table II) initially suggested possible variations in crystal structure since some VCX compounds of different crystal structure (e.g. V6C5 and V8C7) have equal d- values but different relative peak intensities. However, by repeating the XRD analyses after rotating the samples it was found that the peak intensity ratios changed in the same sample. The differences in peak intensity ratios, therefore, appear to be due to the (V,W)CX particles having grown with some preferential orientations.

The microhardness of (V,W)CX was found to depend on both the V/(V+W) and the C/(V+W) ratios. A comparison between the results obtained for (V,W)CX containing very small amounts of W (e.g. Alloy 5) and the results reported in the literature for VCX show that the microhardness of (V,W)CX is closer (statistically equal) to the microhardness of orderered VCX than to the microhardness of disordered VCX [6], although the alloys had been quenched. This can be explained in terms of the high diffusion rate of carbon atoms [7]-

The highest microhardness of (V,W)CX appears to be achieved at a V/(V+W) ratio between 0.8 and 0.88. At equal V/(V+W) or C/(V+W) ratios the microhardness of (V,W)CX in low- C alloys has been found to be consistently higher than in high-C alloys while the lattice parameter in low-C alloys has been found to be consistently smaller than in high-C alloys.

ACKNOWLEDGEMENTS:

The authors wish to acknowledge the financial support of the CNRS (France) and the NRF (South Africa) under the French-SA Scientific Co-operation Agreement as well as the continual support of Vanitec and of the MRSP of the University of the Witwatersrand. J. Zinyana et al. HM 89 295

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REFERENCES:

1. S Luyckx, DJ Whitefield, MJ Witcomb and LA Cornish, Advances in Powder Metallurgy & Particulate Materials - 1998, Metal Powder Industries Federation, Voll, 85-91. 2. S Broccardo, Progress Report: "Work done on the corrosion of WC-VC-Co alloys", University of the Witwatersrand, Johannesburg, South Africa, September 2000. 3. EG Obbard, S Luyckx, S Hamar-Thibault and CH Allibert, 7th Int. Conf. Science on the Science of Hard Materials, Mexico, 5-9 March, 2001 and Int. J. of Refractory Metals and Hard Materials, in press. 4. J Zinyana, MSc Dissertation, University of the Witwatersrand, Johannesburg, South Africa, December 2000. 5. E Rudy, F Benesovsky and E Rudy, Mh. Chem., 93, 693, 1962. 6. VN Lipatnikov, W Lengauer, P Ettmayer, E Keil, G Groboth and E Kny, J. of Alloys and Compunds, 261, 192, 1997. 7. J. Billingham, S Bell and MH Lewis, Acta Cryst, A 28, 602, 1972. 296 HM 89 J. Zinyana et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

TABLE I : V/(V+W) RATIOS IN (V,W)CX PARTICLES AND (V,W)CX FORMULAE

V/(V+W) V/(V+W) from from (V,W)CX FORMULAE OTHER PHASES EDS EMPA PRESENT Alloy 1 0.91 0.77 (Vo.76. Wo.23, COo.Ol) Co.77 binder, graphite Alloy 2 0.74 0.74 (Vo.73. Wo.26. Con.oi) Co.76 binder, graphite, WC Alloy 3 0.80 0.79 (Vo.76. Wo.21. CO0.03) Co.81 binder, graphite Alloy 4 0.85 0.87 (V0.86. Wo.13. COo.Ol) C0.86 binder, graphite Alloy 5 0.90 0.96 (Vo.95. Wo.04. COo.Ol) Co.87 binder, graphite Alloy 6 0.95 0.77 (V0.74, Wo.22. CO0.04) Co.85 binder, graphite Alloy 7 0.98 0.98 (V0.97. W0.02 .Coo.oi) C0.88 binder, graphite

Alloy 8 0.65 0.70 (V0.68, Wo.30. CO0.02) Co.71 binder, WC Alloy 9 0.71 0.74 (V0.74. Wo.26. CO0.01) Co.73 binder, WC Alloy 10 0.78 0.76 (Vo.72, Wo.23. CO0.05) Co.75 binder Alloy 11 0.83 N/A - binder Alloy 12 0.87 0.86 (Vo.85, Wo.14, COo.Ol) Co.76 binder Alloy 13 0.90 0.89 (V0.86. Wo. io, C00.04) C0.79 binder Alloy 14 0.93 N/A - binder Alloy 15 0.96 0.96 (Vo.%. Wo.04, CO0.00) Co.84 binder

N/A = Not analyzed.

TABLE II: COMPARISON OF RELATIVE MAIN XRD PEAKS INTENSITIES OF (V,W)CX OBTAINED BEFORE AND AFTER ROTATING THE SAMPLES

1st set of analyses 2nd set of analyses

Alloy 2 2.42, 2.IO9 1.493 2.42X 2.IO4 1.492 Alloy 3 2.42, 2.10s 1.48s 2.42, 2.10s 1.484 Alloy 4 2.09, 1.489 2.427 2.09x 2.429 1.483

Alloy 5 2.41, 2.099 1.48s 2.41, 2.097 1.482 Alloy 6 2.08x 1.476 2.404 2.08, 2.404 1.47,

Alloy 7 2.40, 2.08, 1.476 2.08, 2.408 1.473 J. Zinyana et al. HM 89 297

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4.21

0.7 0.8 0.9 V/(V+W)

Fig. 1. Variation of the lattice parameter of (V, W)CX with the V/( V+W) ratio in (V,W)CX. HC stands for "high carbon V-W-C-Co alloys" i.e. alloys where the intended carbon content was 27.8 at% and LC stands for "low carbon V-W-C-Co alloys" i.e. alloys where the intended carbon content was 22.8 at%.

4.22 € ~ 4.21 4 2 i ii - • HC £ £ 4.19 a o) 4.18 ssLC S< 4.17 5 — 4.16 5 4.15 0.7 0.75 0.8 0.85 0.9 C/(V+W)

Fig. 2. Variation of the lattice parameter of (V,W)CX with the C/(V+W) ratio in (V,W)CX. HC and LC have the same meaning as in Fig. 1. 298 HM 89 J. Zinyana et al.

International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3100 V) (fl 3000 y i I ' ! .§ 2900 * HC re 2800 \x s^\ y N a LC t i 2 2700 N f J 'E 2600 1 ;/ .. -L " 2500 /[ 0.64 0.74 0.84 0.94 V/(V+W)

Fig.3. Variation of the microhardness (HV0.05) of (V,W)CX with the V/(V+W) ratio in (V,W)CX HC and LC have the meaning as in Fig. 1.

3200 i

• HC ALC

0.71 0.76 0.81 0.86 C/(V+W)

Fig. 4. Variation of the microhardness (HV0.05) of (V,W)CX with the C/(V+W) ratio in (V,W)CX. HC and LC have the same meaning as in Fig. 1. /|. Sommer et al. HM 12 299

15" International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

ON THE FORMATION OF VERY LARGE WC CRYSTALS DURING SINTERING OF ULTRAFINE WC-Co ALLOYS

Michael Sommer, Wolf-Dieter Schubert, Erich Zobetz*, Peter Warbichler*

Institute for Chemical Technology of Inorganic Materials "Institute for Mineralogy, Crystallography and Structural Chemistry Vienna University of Technology, A-1060 Vienna, Austria

** Research Institute for Electron Microscopy A-8010 Graz, Steyrergasse 17/111, Austria

Abstract:

During liquid phase sintering of very fine-grained WC powders (average particle size in the range of 100 to 200 nm) very large WC crystals form with sizes of >20 urn. Such grains typically exhibit the shape of plates. On tracing back the formation of these crystals during sintering it was observed that already on heating to 1150°C small platelets were present in the still porous sintering body. At 1250°C, i.e. before eutectic liquid formed, they had grown to a plate length of up to 7 urn with aspect ratios higher than 1:10.

This early rapid WC crystal growth was observed preferentially in the outer areas of the sintering body where the sample is in closer contact with the sintering atmosphere. SEM imaging and TEM studies reveal the presence of twin boundaries in the plate-like WC crystals, rendering the possibility of a defect-assisted growth of the WC. Intermediate formation of r|-carbides (M12C, M6C) and their subsequent transformation might be an explanation for this early occurrence of extraordinarily large grains.

Keywords: WC grain growth; abnormal grain growth; WC platelets; defect- assisted growth; growth twins. 300 HM 12 M. Sommer et al,

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1 Introduction

The current trend in the hardmetal industry to finer and finer-grained alloys has put high demands on the manufacturing process, for both powders and alloys. Microstructures can be obtained today with WC intercept sizes of less than 100 nm, originating from WC powders obtained by conventional as well as alternative routes [1,2].

One of the crucial aspects in fine-grained hardmetal manufacture is the strong tendency of the very fine WC grains (in the order of 50-200 nm) to coarsen, due to their high interface energies (increased chemical potential) as well as differences in individual grain sizes (originating from the grain size distribution in the green compact), constituting the driving force for the growth process.

To control grain growth during sintering, so-called grain growth inhibitors are commonly added, such as VC or Cr3C2, which retard the growth process. In this regard it is important to avoid any local WC grain growth during sintering, since large grains (> 4 urn) in an otherwise ultrafine-grained matrix can cause material breakdown during application. The smaller the average WC grain size the more critical is this aspect. This has been demonstrated using numerical calculations based on an Ostwald ripening process, in particular for grain sizes smaller than 0.5 urn [3].

In industrial practice, local WC grain growth can be a vexing problem, since its origin is often far from clear, because of the manifold interactions between powder and processing parameters.

On sintering of a -150 nm WC powder grade without growth inhibitor it was observed that extraordinary large WC grains already form during the heating- up period of the sintering cycle, well before the eutectic liquid is formed [4], During subsequent liquid phase sintering very large WC crystals were then obtained with sizes of >20 urn. Such grains typically exhibit the shape of plates.

An investigation was carried out to trace back the origin of early growth to shed more light on this formation process, and to find out under which conditions such large WC plates are formed. /l. Sommer et al. HM 12 301

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2 Experimental

2.1 Starting Materials

High purity starting materials (WC, Co, carbon black) were used for the preparation of alloys. A SEM image of the WC powder is presented in Fig. 1. It originates from the direct carburization process [2], and represents the finest WC powder grade currently commercially available.

0.5 Fig. 1: Field-emission scanning electron micrograph of the WC powder used for the present study; average SEM-size: ~120 nm; calculated BET-particle size: 100 nm.

2.2 Alloy Preparation and Microscopic Investigations

Powder mixtures (90 wt% WC, 10 wt% Co) were prepared either by ball milling (72 hours; weight ratio of milling media: powder 6:1; 60% of the critical rotation speed) or by blending the powders in a TURBULA mixer (1 hour mixing time), subsequently granulated and pressed into compacts at 200 MPa. Sintering was carried out in an industrial GCA vacuum sintering furnace at 1150 to 1350°C. The heating rate was 10°C/min and the total pressure during isothermal sintering <0.01 mbar. The sintered samples were cut and embedded in Bakelite, ground and mirror polished with diamond suspensions. SEM imaging was performed on etched sections (Murakami solution; -5-8 minutes) as well as on fracture surfaces using a JEOL 6400 electron microscope. 302 HM 12 1. Sommer et al.

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Selected areas of the samples showing WC plate formation were marked on an optical microscope and prepared for TEM imaging (Philips CM20 STEM; 200 kV).

3 Results

Fig. 2 shows a comparison between two WC-10 wt% Co alloys, prepared by ball milling or by blending of the powders in the Turbola mixer, respectively; and sintered at 1350°C for 1 hour.

Mixed

Fig. 2: Microstructures of two WC-10 wt% Co alloys, sintered at 1350°C for 1 hour; alloy prepared by mixing of the powders by mixing in a Turbula mixer (left) and, alternatively, by ball milling (right).

In both cases a bimodal WC grain size distribution was formed on sintering, as one would expect for the case of a very fine WC powder grade without additions of growth inhibitors. However, in case of the ball-milled material, an additional, exaggerated local grain growth occurred, with the shape of the large grains in plate form. Obviously, the powder was "activated" during the ball milling process. Tracing back the formation of the very large WC grains to temperatures below the eutectic melt formation (<1275/80°C) [5], several large WC grains were i\. Sommeretal. HM 12 303

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observed in the ball-milled alloy with sizes of up to 4 urn. These grains could be easily detected in the optical microscope, as shining stars in an otherwise dull matrix. In the case of the mixed material only a few significantly larger grains than those of the matrix were present (Fig. 3), with sizes of up to 2 urn. In both cases, all larger WC grains (<0.5 urn) were already facetted, indicating that solution/reprecipitation processes had occurred already on heating in the solid state.

Mixed ball-milled *••

Fig. 3: Microstructures of the two alloys, shown in Fig.2; but heated to only 1250°C (holding time: 10 minutes); both alloys still exhibit residual porosity.

A closer inspection of the ball-milled alloy by means of the electron microscope revealed an interesting microstructural feature (Fig. 4). In the outer, surface-near areas of the sample, where the compact was in direct contact with the slightly decarburizing sintering atmosphere, very thin WC plates were grown, with plate sizes of up to 7 urn and a plate thickness between 100 and 500 nm. In the inner part of the sample, on the other hand, only now and then were such plates observed, as well as some larger grains which had grown more isotropically. On subsequent liquid phase sintering of such heat-treated alloys large WC plates were observed, which on etching showed indications of twinning (Fig. 5). 304 HM 12 1. Sommer et al.

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ball-milled (inner sample area) ball-milled (outer sample area)

Fig. 4: Microstructures of the ball-milled alloy heated to 1250°C and held for 10 minutes; section of the inner part of the sample (left) and close to the sample surface (right). ball milled

Fig. 5: Microstructures of the alloy shown above, but subsequently liquid phase sintered at 135CTC for 1 hour; section close to the sample surface; note that the etching with Murakami solution revealed the presence of twin boundaries. /l. Sommer et al. HM 12 305

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The formation of WC plates with indications of twinning can be observed even better on the sample surfaces (sintering skins; Fig. 6), where it also becomes obvious that surface ledges were formed on twinning, facilitating the rapid growth.

Fig. 6: SEM images of the sintering skin of the ball-milled alloy shown in Fig. 4; demonstrating the formation of growth twins; alloy heated to 1250°C and held for 10 min.

On reproducing this sintering experiment it was observed that the area of plate formation was not always the same within the ball-milled alloy, and obviously did depend on the local gross carbon content within the alloy at the moment of plate formation. It did occur only when the gross carbon content of the alloy was high (i.e. free carbon was present) but the sintering conditions were decarburizing. In the case of very low carbon contents (r\ phases formed all over the sample) no plates were observed. No large WC plates were observed under any of these conditions in the Turbula-mixed alloy.

In the ball-milled alloy, however, WC plates were detected even at 1150°C, with plate sizes of up to 3 urn (Fig. 7). 306 HM 12 M. Sommer et al.

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Fig. 7: WC plate formed in the ball-milled alloy on heating to 1150°C; SEM image of a fracture surface, showing the plate attached to an agglomerate of very fine WC particles.

4 Discussion

On sintering of very fine-grained hardmetals (so-called ultrafine or nanophase grades) significant grain growth occurs during heating to the isothermal hold. This early, non-isothermal contribution to the growth process is a peculiarity of the fine structures, and reflects their high degree of "metastability" as well as a sufficiently high atomic mobility to support the growth process. Even at 1000°C, i.e. at least 275°C below the eutectic melt formation, WC grain growth occurs in such alloys, and concurrent to this growth, low-energy prismatic interfaces are formed [6-8].

Grain growth during liquid phase sintering has been phenomenologically treated as an Ostwald ripening process [3,9,10]. Smaller grains dissolve due to their higher dissolution potential (increased chemical potential), while coarser ones grow by material reprecipitation, thereby reducing the interface area of the system. In the case of the growth of faceted grains in a liquid matrix with atomically flat surfaces (as for faceted WC grains) it is, however, generally assumed that further growth is possible only by two-dimensional surface nucleation or a defect assisted growth process [11,12]. In the case of the former the growth rate is expected to increase abruptly at a critical supersaturation rather than to linearly increase with the driving force arising from differences in individual grain sizes. So, the largest grains (with sizes larger than the threshold value) M. Sommer et al. HM 12 307

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will grow at a significantly higher rate than the finer (but still "stable") ones, whereas the smallest grains will disappear at the expense of the larger, leading to a bimodal grain size distribution as demonstrated in Fig. 2 for the case of a fine-grained WC-Co alloy. In this growth mode it is the coarse grains of the grain size distribution that cause the bimodal growth process. Such a growth mode can be further promoted, simply by adding small amounts of larger WC grains, as demonstrated in Fig. 8 for the case of additions of a 1 urn WC powder grade to the ball-milled alloy. From this comparison it can be seen that significant grain growth had already occurred at 1250°C, whereas the coarse grains obviously had acted as seed crystals for growth. The growth occurred isotropically, and no indications were found for the formation of WC plates.

ball milled (normal) ball milled (with addition)

(iiiliiSiilii! Illililiiiiiiiil

Fig. 8: Microstructure of the ball-milled alloy without (left image) and with addition of 2 wt% of a 1 urn WC powder grade (right image); heated to 1250°C and held for 10 minutes. The present work has also shown, however, under certain conditions of sintering, that significant anisotropic WC grain growth occurred and large, thin trigonal plates were formed rather than equiaxed trigonal prisms. These plates can be traced back to about 1150°C, where they grow at high growth rates at a stage where significant residual porosity still prevails in the compact. This porosity (free space) seems to be a necessary prerequisite for 308 HM 12 M. Sommer et al.

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the growth of such large grains in the solid state leading to plate sizes of up to 7 urn (Fig. 4). TEM imaging of such plate-containing microstructures revealed that the thin plates always contained growth twins, which, obviously, facilitated the rapid growth process (Fig. 9). Their low plate thickness (in the range of 100 to 500 nm) further indicates that the origin for the high growth rates cannot be traced back to the portion of the largest WC grains within the alloy but in the defect itself; i.e. surface ledges are formed through twin formation and a defect- assisted growth can be assumed. In this case also finer WC grains can grow at high growth rates even at low supersaturations, originating from the differences in grain sizes. The formation of twins can have quite different origins. Twins might be present already in the starting WC powder (as growth twins), or might result from intensive powder milling in attritors or ball mills (introduction of defects) and subsequent recrystallization on heating (recrystallization twins). However, twinning might be caused as well by chemical reactions which take place during sintering and which might result in the formation of growth twins and subsequent plate formation. A strong indication for this formation mode is given by the present investigation, which has shown enhanced twinning as a result of powder milling.

0 2 |jm AS

Fig. 9: TEM bright-field images of the ball-milled alloy shown in Fig.4, indicating the presence of growth twins in the large WC plates; note that even the smallest WC grains are facetted; heated to 1250°C and held for 10 minutes. However, twinning occurred only in certain parts of the samples (and not all over the alloys, as one would expect in case of recrystallization twinning), and, obviously, as a result of the gross carbon content. In this regard it is important to consider that on ball milling a significant oxygen pick-up occurs. M, Sommer et al. HM 12 309

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This oxygen (present at the WC and Co powder particle surfaces in the form of the respective oxides) is released later during sintering in the temperature range of 500-1000°C as CO2 and CO, causing local carbon deficiencies, where -q-carbides (Co3W3C, Co6W6C) can be formed. Subsequent transformation of these carbides to form WC and Co, according to

C03W3C + 2C -> 3 Co + 3 WC under carburizing conditions (by carbon diffusion from carbon-rich areas) might result in the in-situ formation of WC and the formation of growth twins. An important prerequisite for this formation mechanism is the assumption of a non-equilibrium during early sintering (cobalt spreading onto the WC surfaces) and the build-up of carbon activity differences during sintering. These conditions will be influenced by the "local microchemistry" of the WC- Co compact (presence of off-stoichiometric areas; carbon nests; graphite particles) as well by the sintering atmosphere. This is more or less decarburizing, depending on temperature and residual atmosphere. In 1981 it was reported that twinned WC crystals form on the carburization of milled tungsten powders in the presence of an iron binder metal in the temperature range of 900 to 1200°C [13]. Only recently it was then demonstrated that WC plates with growth twins are formed "in-situ" during sintering of a Co3W3C + graphite (or carbon black) powder mix [14]. Thus these reactions can be seen to proceed locally in a fine-grained WC powder matrix with local fluctuations in carbon content. Since only a few WC twins form in this case (as compared with a pure reaction sintering process [15]), only a few, but very large WC plates are formed, which rapidly grow at the expense of the surrounding still very fine WC matrix. It is easy to demonstrate that the presence of small amounts of ri-carbides does promote the formation of large WC plates in a fine-grained WC matrix. The result of such an experiment is shown in Fig. 10 (left image) for the case of small additions of Co3W3C (produced from elementary powder mixes) and the respective equimolar amounts of carbon black to a ball-milled WC-Co powder batch. In such an alloy WC plates are formed in isolated, nest-like plate structures, rather than as individual plates. This defect-assisted anisotropic growth mode can be easily differentiated from the case of a rapid isotopic growth mode, which can be forced by additions of a small amount of larger WC grains, as shown earlier in Fig. 9. 310_ HM 12 M. Sommeretal.

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ball milled (addition of Co3W3C and graphite)

10 um ^ 5 1 um

Fig. 10: Microstructure of the ball-milled alloy; 2wt% Co3W3C/graphite added to the powder mixture; heated to 1250°C and held for 10 minutes; plates formed close to a small graphite particle.

Considerations on twin formation

WC crystallizes in the non-centrosymmetric space group P-6m2 (No. 187) with one formula unit per unit cell [16]. The cell dimensions are a = 2.906 A and c = 2.837 A with cla = .976. The atoms are in the following special positions: W in 1 a: 0,0,0; C in 1 of: 1/3,2/3,1/2. Fig. 11 shows a projection along the c axis. The atoms are arranged in alternate close packed layers of W and C in (0001) planes. Close packed layers are 36 nets composed of equilateral triangles. For such layers to be stacked the nodes of one net must center alternate triangles of nets above and below. The stacking symbols A,B,C and cc,ß,y, where the Roman letters represent W atoms and the Greek letters represent C atoms, relate the positions of the nodes of the nets of the hexagonal cell. For WC the stacking repeat sequence is Ay.

Each W atom in WC is surrounded by a trigonal prism (Fig. 11), W by six C at 2.197 A. It might be noted that the structure is its own anti structure so there is also each C atom coordinated by six W atoms. Each atom is further coordinated by eight atoms of the same kind at 2.837 A (2 x) and 2.906 A 1. Sommer et al. HM 12 311

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(6 x). These distances are comparable with the nearest neighbor distance in pure tungsten, 2.735 A.

A, a Fig. 11: Unit cell of the WC structure projected along [0001]. W, open circles; C, black circle. A,B,C and a,ß,y denote the stacking sites for close packed layers. Each atom in the WC structure is coordinated by six atoms of the other kind (indicated by three pairs of double lines).

The structure of WC contains three kinds of dense plane forms, namely the pinacoid {0001}, and the two crystallographically different trigonal prisms {10-10} and {-1010}, which seem to determine the morphological properties. WC crystals occur preferentially as trigonal prismatic plates, but it is not known yet, whether they are bound by {0001} and {10-10} or by {0001} and {-1010}. Due to the fact, that for characteristic X-rays the differences between the intensities of (hkl) and (-h-k-l) reflections are rather small with respect to the large absorption effects, the only way to determine the absolute structure (and thus to differentiate between {10-10} and {-1010}) would be the use of synchrotron radiation with wavelengths near the L-absorption edge of W {X= 1.21545 A).

It is often observed (Fig. 12a) that WC crystals occur as contact twins (twinning by merohedry) with [01-10] as the twin axis (rotation through 180°) and (0001) as the composition plane. Twinning of WC might be due to stacking faults. Above and below the composition plane WC adopts the NiAs structure (Fig. 12b), which can be coded as AyAß. In this structure W is surrounded by six equidistant C atoms situated at the comers of a distorted octahedron. The layers of trigonal prisms consisting of C atoms are the same as in the WC structure, but now adjacent layers are staggered so that trigonal prisms share only edges (and not faces) between the layers. The contact region might consist of alternating small parallel crystals referred to as twin lamellae. Each lamella is related to its adjoining neighbor by the appropriate twin law. Contact twins with two more or less large twin domains occur, if the number of stacking faults (composition planes) is odd. 312 HM 12 L Sommer et al.

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(0001)

NiAs-lypc

(0110) (1010) (1010) (0110) (1100)

Fig. 12: a) Twinning by merohedry in WC. b) WC structure with one stacking fault projected on (11-20). W, open circles;C, black circles. The broken line indicates the composition plane parallel (0001). Thin dotted lines outline the coordination polyhedra of W.

If the number of stacking faults (composition planes) is even (Fig. 13a) the two domains have the same orientation (Fig. 13b). Initially polytypism has been defined as a special case of polymorphism in which the various polytypes are derived from the different ways of stacking structurally and chemically equivalent layers.

(0001)

WC-typc

NiAs-type

(0110) (1100) (1010) WC-typc

Fig. 13: a) Two domains of WC with the same orientation, maybe due to an even number of stacking faults, b) WC structure with two stacking faults projected on (11-20). W, open circles; C, black circles. The broken lines indicate composition planes parallel (0001). Thin dotted lines outline the coordination polyhedra of W.

However, in the case of silicates it became clear that the relative stability of stacking variants is a function of the composition of the constituent layers. M. Sommer et al. HM 12 313

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This has been reinforced by the discovery that polytype stability in SiC is affected by trace impurities [17]. Impurities or enrichment of elements therefore might be the reason for stacking faults and twinning in WC.

5 Conclusion

• On sintering of very fine-grained WC powders (particle sizes in the range of 100 to 500 nm) significant local grain growth already occurs on heating to the isothermal hold. • This growth occurs both isotropically as well as in an anisotropic growth mode, where large trigonal WC plates are formed. • While isotropic growth can be related to coarse WC grains within the sintering body acting as seed crystals for the local growth, a defect- assisted growth has to be assumed in case of the plate formation. • The plates contain growth twins, which promote the grain growth by forming surface ledges. • The trigonal plate containing growth twins can be regarded as a growth shape (habitus) of the hexagonal WC.

6 References

1. W.D. Schubert, A.Bock and B. Lux: Advances in Powder Metallurgy & Particulate Materials (Ch.L. Rose, M.H. Thibodeau, eds.), MPIF, Princeton, New Jersey, Vol.3 (1999) pp. 10-23. 2. Y. Yamamoto, A. Matsumoto, Y. Doi: Proc. 14th Int. Plansee Seminar, Vol.2, pp. 596-608 (ed. G. Kneringer et al, Plansee AG, Reutte (1997). 3. K. Hayashi, N. Matsuoka: Proc. 14th Int. Plansee Seminar, Vol.2, pp. 609-621 (ed. G. Kneringer et al, Plansee AG, Reutte); 1997. 4. Ch. Walch: Densification and Grain Growth of Ultrafine WC during Hardmetal Sintering, Diploma Thesis, Vienna University of Technology, Austria (1994). 5. L. Akesson: Sandvik Coromant, Lab. Report No.2510, Stockholm (1978). 6. W.D. Schubert, A. Bock, B. Lux: Int. J. of Refractory & Hard Materials 13 (1995) pp. 281-296. 7. T. Taniuchi, K. Okada, T. Tanase: Proc. 14th Int. Plansee Seminar, Vol.2, pp. 644-657 (ed. G. Kneringer et al, Plansee AG, Reutte);1997. 314 HM 12 M. Sommer et al.

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8. K. Okada, M. Hisanaga, T. Tanase: Proceedings of PM2000, Kyoto, to be published. 9. C. Wagner, Z. Elektrochemie 65 (1961) 581. 10. Fischmeister, G. Grimval: Sintering and Related Phenomena, ed. G.C. Kuczynski, Plenum Press, New York (1980) pp. 119-46. 11. Y.J. Park, N.M. Hwang, D.Y. Yoon: Metall. Trans. A, Vol.27 (1996) pp. 2809-2819. 12. C.H. Allibert: Int. J. of Refractory & Hard Materials 19 (2001) pp. 53-61. 13. M. Kobayashi, S. Yamaya: Proc. 10th Int. Plansee Seminar, Vol.1, pp. 597-612 (ed. H. M. Ortner, Plansee AG, Reutte); 1981. 14. S. Kinoshita, T. Saito, M. Kobayashi, K. Hayashi: J. of the Japan Society of Powder and Powder Metallurgy, 47, No.5 (1999) p. 526. 15. M. Kobayashi, et. al: European Patent EP 0759 480 A; 1995. 16. E. Parthe, V. Sadagopan: Monatshefte für Chemie, Vol. 93 (1962) pp. 263-270,. 17. N.W. Jepps, T.F. Page: J. Cryst. Growth and Characterisation (1983) pp. 259-307. AT0200484 A.A. Ryzhkin et al. HM 33 315

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Wear Resistance and Electronic Structure of Cutting Tool Materials on a Basis Carbides of Tungsten and Titanium

A.A. Ryzhkin, V.V. llyasov, A.V. Lyulko

Don State Technical University, 344010 Rostov-on-Don, Russia

Summary:

The decision of problems of durability, in particular of shock resistance and resistance to wear of tool materials, has allowed to formulate a set of the requirements to them, and also to establish dependence between physical properties and characteristics of wear. However for understanding of a nature of processes, determining, for example, tribological property of the cutting tool it is necessary to consider interaction of atoms among themselves in a crystal, carried out valence electrons. We spend theoretical study of the physical properties of cutting tool materials of system W-Ti-C in a wide interval of changes of concentration of titanium. Are calculated total and partial local density of electronic condition for each atom in a hard solutions W(1-x)Ti(x)C. Within the framework of one approach comparison of electronic structure of considered(examined) firm solutions in comparison to binary analogues is spent. Calculation partial local charges of valence bandof system W(0.83)Ti(0.17)C has shown, that there is the carry charges to W (0,33e) and carbon (0.29e), and Ti acts in a role of the donor of electronic configurations W and carbon. The analysis of the given accounts of electronic structure and absolute termo-e.m.s. for system W(1-x)Ti(x)C to predict optimum concentration of Ti in the given system, determined by area of the minimum meaning of absolute termo-e.m.s., that proves to be true by x-ray researches. 316 HM 33 A.A. Ryzhkin et al.

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Keywords: Cutting tool, W(1-x)Ti(x)C, durability, tribological property, wear, electronic structure, termo-e.m.s.

System researches of the last years, in a combination to the analysis received in work [1] the formulas, have allowed to define the basic ways of decrease the wear and management of the wear resistance at friction in conditions of cutting: • Maintenance of a TD-condition of a tool material with the minimum meaning of density saved entropy; • Localization of thermodynamic processes in thin-film structures; • Application of cutting materials with high meaning of critical density of entropy; • Maintenance of a mode of self-organizing; • Development of criteria of optimization and analytical dependences for a choice of an optimum mode; • Diagnostics of wear during cutting. The significant interest represents consideration of the first direction, the realization of which assumes the decision of problems macro- and microlevel. In particular, at a microlevel a major problem, on our sight, is perfection of structure of a tool material. As in quality of structural - sensitive parameter size termo-e.m.s. can serve, which is defined by character of distribution of density electronic condition (DOS) in a vicinity of a level Fermi, the establishment of correlation communication between parameters of the wear resistance and termo-electrical of the characteristics of a tool material represents scientific and practical interest [1-3]. Carried out in [3] the analysis A.A. Ryzhkin et al. HM 33 317

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and estimated accounts of the wear resistance have shown, that lifted in [4] the concept of the termo-e.m.s. a parameter of the wear resistance is objective. As is known [5], termo-e.m.s. is caused by three reasons: by dependence of a level Fermi from temperature; diffusion of electron and enthusiasm fonons of electrons. For understanding of a nature of physical processes determining the tribology properties of a material of the cutting tool, it is necessary to consider interaction of atoms among themselves in a crystal carried out valence electrons and described by the equation of Shredinger. Size termo-e.m.s. are defined by character of distribution DOS in a vicinity of a level Fermi. Less size termo-e.m.s. is experimentally established, that than., the above the wear resistance of a material. Thus, absolute meaning termo- e.m.s. can act by the important indicator in search of structure of new cutting tool materials, acting in a role «bridge» between their electronic structure and the wear resistance. One of the main tasks of a science about friction and deterioration is the development of analytical dependences for a settlement estimation of size of deterioration (or intensity of wear process) with the account probably of greater number of the influencing factors. Let's consider the most investigated kind of wear process - abrasive, which agrees by the data Vierrege G. [6], is shown with uniform intensity in all range of speeds (temperatures) cutting. Therefore we shall establish interrelation between termo-electrica! by the characteristics of a tool material and relative wear resistance at abrasive wear process, using classical dependence Hrushev M., the hardness of tool materials of materials can be found on correlation dependence. For definition 318 HM 33 A.A. Ryzhkin et al.

15!" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 expressions allow on data of account complete and local partial of density of electronic condition and experimental meanings termo-e.m.s. to receive estimations of wear resistance of a material of a product at abrasive wear process. The electronic structure carbide of systems WC, W^T^C, TiC paid off by a method local coherent of potential within the framework of the theory repeated scattering under the circuit, described in work [9]. The theoretical study of the specified properties of tool materials of system W - Ti - C is spent in a wide interval of changes of concentration titanium, considering them as firm solutions of the carbide tungsten W^Ti x C (x = 0.. 0.6) with structure of a lattice as NaCI. Are calculated complete and partial of density of electronic condition for each atom in a firm solution. Within the framework of one approach comparison of electronic structure of considered firm solutions in comparison to binary analogues is spent. Calculation partial charges of electrons in the top part of top valence band of firm solutions of the carbide tungsten and titanium (W, Ti) C has shown, that to increase of concentration the carbide tungsten in system there is the change of numbers of filling. At replacement of atoms of the tungsten by titanium, there is the downturn of statistical weight electronic sp3 - configuration of atoms of tungsten. Similarly, using experimental meanings termo-e.m.s. of the materials, calculate the hardness, relative wear resistance and energy of communication, which are submitted in tabl. The analysis of data table shows, that the given model reflects the tendencies of increase the wear resistance of tool materials in a this line, with reduction termo-e.m.s. Observable increase relative wear resistance of tool materials on a basis carbide of systems correlates with increase of energy of communication in a tool material.The increase of energy of communication A.A. Ryzhkin et al. HM 33 319

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 can be explained hybridization of 2p-condition of carbon and 5d-condition of tungsten, doped of 3p- and 3d-condition of titanium, that indirectly proves to be true by increase of hardness in the this line.

Table. Settlement the wear resistance tool materials for temperature of cutting t = 600° C

Material EF, s ,|dV/deg H, ' OTH Ry GPa Ry/un.cell

we 0.900 -23.0 23.57 10.7 330.45 BK8 0.920 -20.3 25.54 11.1 358.13 (W - 8Co) T5K10 0.913 - 19.4 30.68 11.5 430.06 (W-5TiC-10Co) T15K6 0.918 - 16.6 34.40 11.9 482.29 (W-15TiC-6Co) T30K4 0.930 -11.9 39.38 12.3 552.21 (W-30TiC-4Co) T60K6 0.970 0.96 57.65 13.6 808.83 (W-60TiC-6Co)

Thus, with reduction of meanings absolute termo-e.m.s. As the characteristics of electronic structure of a tool material is increased the wear resistance. Received settlement data can basically explain marked many by the authors presence of communication between termo-e.m.s. and intensity of wear process hard-alloy T5K10, T15K6 and BK8 on automatic transfer lines. 320 HM 33 A.A. Ryzhkin et al.

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Correlation connection between intensity of wear process and size termo- e.m.s. In conditions of friction and cutting, established for firm alloys, is kept and for other groups of tool materials, for example, oxide-metal material and high-speed steels. Thus, despite essential differences of structure of the high-speed steels from firm alloys and for them character of correlation communication between the wear resistance and absolute termo-e.m.s. is kept. The settlement sizes of energy of communication of carbides can basically be aplicable in future for an estimation of intensity of wear process and for other kinds of wear process, in particular, adhesive wear and diffusion wear.

References

1. Ryzhkin A.A. 14-th Inter.Plansee Seminar' 97. May 12 - 16, 1997, Reute/Tirol/Austria. Proc. Vol. 2. PP. 300 - 314. 2. Ryzhkin A.A. Triboelectrical of the phenomenon and wear of tool materials // Reliability and efficiency machine and tool systems: Rostov-on-Don, 1998.-P. 9-51 (in Rus). 3. Solonenko V.G., Zarezhkii G.A. Wear resistance of cutting tools: DGTU and KGTU, 1998. -104 P (in Rus.). 4. Yu.V.llyasov, A.A. Ryzkin, V.V.IIyasov, I.Ya. :lnterConf. DF&PM. Slovac Republic, Piesany.-1999.-PP.216-218. 5. Blatt F. J.,Shreder P.A., Foilz K.L., Greig D. Termoelectromotive force of metals. Translation with angl. Under edit. D.K. Belashenko. M.: Metallurgy, 1980.-248P. 6. Vierrege G. Zerspanung der Eisenwerkstoffe.- Verlag Stahleisen.-2. Auflage.- Düsseldorf, 1970 - P. 81- 83. A.A. Ryzhkin et al. HM_33 321

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7. Fridel J. The Physics of Metals I..Edited by J.M.Ziman (Cambridge University Press, Cambridge, 1969). 8. Ryzhkin A.A., llyasov V.V. and llyasov Yu.V. Estimation abrasive wear of tool materials / Friction and wear, DGTU: 2000.- P.13-22 (in Rus.). 9. Nikiforov I.Ya.,llyasov V.V. and Safontseva N.Yu. Electron energy structure of nonstoichiometric cubic boron nitride//J.Phys. Condens.Matter. 1995. V.7. P. 6035-6044.

Решение проблем долговечности(срока службы), в особенности ударопрочности и прочности материалов инструмента, разрешило формулировать набор требований к ним, и также устанавливать зависимость между физическими свойствами и характеристиками износа. Однако для понимания характера(природы) процессов, определения, например, трибологические свойства режущего инструмента необходимо рассмотреть взаимодействие атомов между собой в кристаллическом, выполненном валентными электронами. Мы тратим(проводим) теоретическое исследование физических свойств материалов режущего инструмента системы W-Ti-C в широком{...} Мы тратим(проводим) теоретическое исследование физических свойств материалов режущего инструмента системы W-Ti-C в широком интервале изменений(замен) концентрации(обогащения) титана. Рассчитаны общее количество и частичная локальная плотность электронного условия(состояния) для каждого атома в жестком Указателе времени Вт (1-х решений(растворов) (х) С. В пределах каркаса одного сравнения подхода(подводящего канала в экструзионной головке) электронной структуры рассматриваемых (исследованных) твердых решений(растворов) по сравнению с бинарными аналогами потрачен. Вычисление частичные локальные обвинения валентности bandof системный Указатель времени Вт (0.83, который ((0.17) С показал, что имеются обвинения переноса к Вт (0,33e и копировальной бумаге (0.29е), и действиям Указателя времени в роли донора электронных конфигураций Вт и копировальная бумага. Анализ данной отчетности электронной структуры и абсолютного termo-e.m.s. Для системного Указателя времени Вт (1-х (х) С, чтобы предсказать оптимальную концентрацию(обогащение) Указателя времени в данной системе, определенной областью значения минимума абсолютного termo-e.m.s., который, доказывается, {...} Анализ данной отчетности электронной структуры и абсолютного termo-e.m.s. Для системного Указателя времени Вт (1-х (х) С, чтобы предсказать оптимальную концентрацию(обогащение) Указателя времени в данной системе, определенной областью значения минимума абсолютного termo-e.m.s., который, доказывается, истинный рентгеновскими исследованиями. 322 HM 35 H. Schön et al.

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WC PLATELET-CONTAINING HARDMETALS

A. Schön, W.-D. Schubert, B. Lux

Institute for Chemical Technology of Inorganic Materials Vienna University of Technology, A-1060 Vienna, Austria

Since Toshiba Tungaloy introduced their Disk Reinforced Cemented Carbide to the scientific community, WC platelet-containing hardmetals have been a topic of major interest.

The present work provides a literature review on the formation and growth of the trigonal-shaped WC platelets. The review establishes two main mechanisms of plate formation. Based on this knowledge a new processing route has been elaborated at Vienna to obtain WC plate-containing hardmetals. This route is based on the recrystallization of very-fine grained WC in the presence of small amounts of TiC or Ti(C,N) [0.2-0.7 wt% TiC.N]. Alloys based on this processing method show a high degree of plate formation with aspect ratios of up to 1:10, as well as promising hardness/ toughness combinations compared with conventional, plate-free materials.

Keywords: hardmetal, cemented carbide, WC platelets, growth twins, abnormal grain growth

1. Introduction

Platelet- and fiber-reinforced ceramics showing enhanced fracture toughness appear to be attractive materials for structural applications. Commonly, the platelets and fibers are added during P/M processing but can also be formed "in-situ" during the sintering process. The latter route has been demonstrated to be successful in case of WC platelet-reinforced hardmetals [1-3], based on a reaction sintering process, and only recently such materials were introduced into the cutting tool market [4].

The present work focuses on this interesting material development. It aims to construct a new alternative route for obtaining WC platelets in hardmetals. H. Schön et al. HM 35 323

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This route originates in a comprehensive literature review and constitutes a simple and straightforward approach based on well-introduced raw materials and common processing techniques.

2. Literature

The story of WC platelets in hardmetals can be traced back to the mid-sixties, when du Pont de Nemours and Company filed a series of patent applications on WC plate formation [5-7]. Since then, several further studies were published/patented, based on different principles of plate formation, leading to the commercializing of platelet-containing alloys in 1998/99 [4]. The following list of investigations gives a short chronological review on the most important work on WC plate formation, which acted as a basis for constructing a new route for obtaining WC plate-dominated microstructures.

1964-72: G. W. Meadows (du Pont de Nemours and Company) was granted several patents for the manufacture of plate-like WC-based hardmetals. It is claimed that the plates are formed during sintering of colloidal WC powders in the presence of an iron binder metal matrix. Such alloys are said to exhibit superior combinations of strength, hardness and toughness. The platelets can be oriented by hot pressing or hot extrusion. [6,7]. 1968: L. Pons (University of Caen, France) presents a paper on the plastic properties and grain shape of WC [8]. In one of his chapters he describes the formation of twinned, plate-like WC crystals on recrystallizing WC with and without binder by heating to 2000°C. 1975: M.K. Brun et a! (Pennsylvania State University) report their results on precipitation studies in the system WC-TiC [9]. They observe the formation of plate-like WC crystals out of supersaturated solid-solution powders on sintering the powders with 10 wt% Co at 1450°C. 1980-1981: M. Kobayashi et al (Toshiba Tungaloy) publish their work on twinned WC grains in cemented carbides [10]. Flaky W powders (produced by powder milling) form plate-like WC during carburization in the presence of ferrous metals. 1989: K. Kobori et al (Toshiba Tungaloy, Tokio University) describe the formation of sheet-like WC in WC-(W,Ti,Ta)C-Co alloys [11]. The sheets form on sintering of (WC/TiC/TaC) powders supersaturated in tungsten. 1992: S. Masahiro (Toshiba Tungaloy) describes a procedure for obtaining plate-like WC on sintering of hardmetals containing 3-40% of cubic carbides, by using fine-grained WC powders (<0.5um) [12]. 324 HM 35 H. Schön et al.

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1994-2000: M. Kobayashi et al (Toshiba Tungaloy) are succeeding in producing plate-reinforced hardmetals on a large scale. The patented process [1] is based on the formation of a W-Co-C powder precursor sub- stoichiometric in carbon (related to the final composition WC-Co), which is subsequently used for hardmetal sintering ("reaction sintering"). The plates can be oriented by the use of proper starting materials and processing conditions [13,14]. Several papers demonstrate the advantage of these materials over conventional, plate-free alloys. Commercialization took place in 1998/99. 1997-98: K. Rödiger et al. (Widia GmbH) patent the formation of WC platelets on reaction sintering of (W, Co, C)-powder mixes in a microwave device [15]. 1998: A.V. Shatov publishes a work on the shape of WC in hardmetals [16]. He observes plate-like WC in the system WC-Ni in the presence of small amounts of TiC. The formation of plates is explained by a face-specific adsorption of titanium on the growing WC.

From this review two main principles of WC plate formation can be elaborated:

• Plate formation through recrystallization of fine-grained WC on heating above a certain critical temperature (which depends on the grain size of the starting WC and the amount/composition of the metallic binder). In this case, the WC platelets are formed by a dissolution/reprecipitation process ("recrystallization") [5,8,12,16]. • Formation of WC plates through nucleation and growth of the carbide in various chemical environments, in most cases in a solid or liquid iron group metal binder matrix. In this case, the WC forms anew during annealing/precipitation/sintering via the transformation of a tungsten- containing source, such as, for example, a W-supersaturated cubic mixed crystal carbide (W,Ti,Ta,Nb)(C,N), W, W2C, Co7W6, ri-phases (e.g. C03W3C) orK-phases (e.g. C0W3C) [1,9,10,11,15].

In both principles the plate-like morphology of the WC can be regarded as a growth shape of the hexagonal WC rather than the equilibrium shape. This conclusion is confirmed by the observation that during rapid growth of the WC, as for example, during eutectic precipitation of the WC in case of very high Co grades (no WC seed crystals present; locally high melt supersaturations) very thin WC plates (lamellas) are always formed, rather H. Schön et al. HM 35 325

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than isotropically grown equiaxed trigonal prisms, which are considered to be the equilibrium shape of the WC [17]).

It can be further concluded that additions of titanium [9,16] support (stabilize) the plate formation both during recrystallization as well as "new formation" (reaction sintering).

3. The Meadows/Shatov route

Recrystallization of very fine-grained WC powders as described by G. Meadows in 1964 [5] constitutes a straightforward way of WC plate formation. By contrast with the manufacture of hardmetal in the 1960's powders are commercially available today with average WC grain sizes in the range of 100 to 200 nm. On sintering of such materials at comparatively high temperatures (up to 1500°C), WC plate formation occurs, as demonstrated in Fig.1 in the case of the finest WC powder grade currently commercially available. However, this kind of growth occurred only in case of heavily milled WC powders, where the milling process had obviously activated the recrystallization mode. In addition, recrystallization was very non-uniform and a portion of very large WC grains with sizes greater than 10 urn always occurred in the microstructures. Further experiments then indicated that additions of titanium carbide or carbonitride can force the recrystallization of the WC in the plate shape, as described earlier by A. Shatov for the case of WC-Ni alloys [16].

Based on these preliminary investigations a new route of hardmetal manufacture was established, which is based on the recrystallization of fine- grained WC powders (so-called ultrafine grades) in the presence of small amounts of TiC or Ti (C,N). With reference to the earlier investigations by G. Meadows and A. Shatov this route was named the Meadows/Shatov route. 326 HM 35 H. Schön et al.

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10 (jm

Fig.1: Light optical micrograph of a WC-10 wt% Co alloy, ball-milled for 14 days and sintered at 1500°C for 20 minutes; alloy based on the WC powder shown in Fig.2a; Murakami etching; 1000 x.

4. Starting Materials and Experimental Procedure

Several ultrafine WC powder grades were taken into consideration for the recrystallization experiments. This included WC powders produced by conventional processing, as well as by alternative routes [18-20]. Fig. 2 presents SEM-images of the finest and the "coarsest" WC powder used.

Fig.2: SEM images of the finest (2a) and "coarsest" (2b) WC powder used for recrystallization experiments; magnification: 20,000 x. H. Schön et al. HM 35 327

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High purity Co (0.95 |jm; H.C.Starck), TiC powder (3.6 pm; Treibacher Chemische Werke) and Ti(C,N) powder were used for alloy preparation.

Powder mixes (12-14 wt% Co) were prepared by ball milling (90 hours; weight ratio of milling media: powder 6:1; 60 % of the critical rotation speed) or attritor milling (2 hours; 700 rpm; ball size: 1.2 mm; disk stirrer), subsequently granulated and pressed into compacts at 200 MPa. Liquid phase sintering was carried out in an industrial GCA sintering furnace at 1350-1500°C for 20 to 180 minutes. The heating rate was selected between 3 and 10°C/min and the pressure during isothermal sintering was <0.01 mbar. Hardness and indentation toughness measurements were carried out on selected alloys as described in [21], using an indenter load of 50kgf.

5. Results and Discussion

Several parameters turned out to influence the degree and mode of recrystallization: • the sintering temperature and sintering time • the average grain size of the starting WC powder grade • the alloy gross carbon content • the amount of cobalt • the amount of TiC or Ti(C,N) added (-0.2 wt% -» 0.7 wt%) • the residual atmosphere/ heating rate

Only the most crucial parameters are discussed in the following.

Sintering temperature/time In general, high sintering temperatures were necessary to achieve a high degree of WC plate formation. This is demonstrated in Fig.3 for the case of a WC-14 wt% Co alloy that was sintered under three different sintering conditions. At 1400°C, the matrix of the alloy still remained fine, but a few WC plates with plate sizes of up to 3um were already formed. At 1500°C and 20 minutes, most of the WC grains are already in plate form, and the microstructure is dominated by a large number of small WC plates with aspect ratios of up to 1:10. After 3 hours at 1500°C, almost a complete recrystallization occurred, with WC plate sizes of up to 7 urn. 328 HM 35 H. Schön et al.

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K , *„<**** 1\ ** ^ * •"•' *> #=» ^

FV~.,t. -? V

I« **.

«*^: J r \ - -^ "t 1**'^ \ ; *-„<** v' 1400°C 60 min 1500°C20 min 1500°C 180min Fig.3: SEM images of WC-0.7TiC-14Co alloys, sintered to different degrees of recrystallization; based on the WC powder shown in Fig. 2a; 4000 x.

Average grain size of the starting WC powder The smaller the grain size of the WC powder the stronger is the tendency to recrystallize. "Coarser" WC powders (-0.35 |jm; Fig.2b) need stronger recrystallization conditions to form a high degree of platelets. Thus, the sintering cycle has to be adapted to the WC powder particle size. Nevertheless, almost all grains can be forced into the plate shape (Fig.4).

10|jm

Fig.4: SEM image of a WC-0.35TiC-14Co alloy; sinterhiped under industrial conditions; 2000 x. H. Schön et al. HM 35 329

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Alloy gross carbon content It is well known that the alloy gross carbon content significantly co-determines the dissolution/reprecipitation processes in hardmetals and, therefore, also the degree of recrystallization. High carbon alloys promote grain growth, and, consequently, the degree of recrystallization. However, recrystallization occurs less uniformly and the plates that are formed show a somewhat lower aspect ratio (Fig.5). Low carbon alloys restrict both the overall as well as the local grain growth.

> 10 pm ,' V V

~y i n K ^ >

\ '"., J- v. /' Fig.5: SEM images of a high carbon (left) and low carbon (right) WC-0.3TiC-12Co alloy; sintered at 1500°C for 20 min; based on the WC powder shown in Fig. 2a; 2000 x.

Aspect ratios of up to 1:10 can be obtained in such low carbon variants under proper sintering conditions (Fig.6).

Fig.6: Microstructure of a low carbon WC-0.4TiC-14Co alloy; sintered at 1500X for 20 min; 4000 x. 330 HM 35 H. Schön et al.

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Local WC plate growth On recrystallization of fine-grained WC powders a certain plate size distribution is always formed. The finer the WC powder, and the broader the starting WC particle size distribution, the stronger is the tendency for the local growth of very large WC plates. These are detrimental to properties and can act as a fracture source during stressing. An extreme for such a local giant plate growth is demonstrated in Fig.7. The formation of such plates can be traced back to the early stages of sintering (~1250°C) and the occurrence of chemical and/or physical heterogeneities, leading to a defect-assisted growth of the WC by formation of growth twins [22]. Such heterogeneities (which can be already present in the WC powder grade or form during intensive powder milling) have to be strictly avoided by choosing proper processing conditions.

•o 7~^ ^ Fig.7: Giant WC grain growth obtained in a WC-0.35TiC-14Co alloy; sintered at 1500°Cfor one hour; such large WC plates can already form during the heating-up of the sintering cycle (still in the solid state) as a result of local excess carbon.

However, appropriate powder raw materials as well as adjusted sintering conditions result in alloys with a high degree of microstructural uniformity and WC plates with plate sizes in the range of 2-5um (Fig.8). H. Schön et al. HM 35 331

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Fig.8: BSE image of a WC-0.35TiC-14Co alloy sintered under optimized processing conditions; note that the largest plates are in the order of 7|jm; random plate distribution.

The twinned trigonal WC plate as the growth shape of the WC The trigonal WC plates can be considered as the growth shape of the hexagonal WC. TEM imaging reveals that the plates contain growth twins and stacking faults that facilitate crystal growth by providing surface ledges for growth (Fig.9). Growth occurs preferentially on the prismatic planes (1010), whereas the basal plane (0001) remains as the dominating crystal face.

Such crystal defects are also easily visible on the sintering surfaces (skins) of the WC plate-containing alloys as demonstrated in Fig. 10. 332 HM 35 H. Schön et ai.

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1 um

Fig.9: TEM image of a WC-Co alloy produced by the "Meadows/Shatov" route; note the presence of growth domains (twins; stacking faults) which obviously facilitate the growth along the prismatic planes.

\im

Fig.10: SEM image of the sintering skin of a WC-0.4TiC-14Co alloy; sintered under industrial sintering conditions; 4000 x. H. Schön et al. HM 35 333

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Hardness to toughness relationship of WC platelet-containing alloys Selected plate-containing alloys were used for hardness (HV50) and indentation toughness measurements. The results are summarized in Table 1 together with those obtained on two commercial, plate-free WC-Co alloys. From this comparison it can be seen that in the hardness range of HV 1300 to HV1400 promising toughness improvements were obtained for the plate-containing alloys, in particular in the case of optimized plate microstructures. SEM investigations of the crack paths give evidence for the various toughening mechanisms that work in brittle materials, in particular crack deflection and cleavage, but also bridging, microcracking, and debonding (Fig.11). However, it cannot be excluded that residual stresses, originating from the cooling due to the different thermal expansion coefficients of the WC and binder (and being different in plate-containing alloys as compared with plate-free materials), interact on the different mechanisms, and, hence, also influence the crack resistance of the respective alloys.

Table 1: Hardness and indentation toughness parameters of WC plate-containing hardmetals, compared with commercial WC-Co grades. recr stalh Co WC powder y - £ crack- Palmqvist- Klc desig. content est. SEM za.' HV50 length toughness [MN r Jn/, • r 7 conditions r , „,, 7 3/2 7 3 [wt%] sizefrm] r1500°ci &m] [N/mm] mm "] commercial, plate-free alloys Std 1 WC-Co T370 343 143Ö 12.3 Std2 WC-Co 1297 203 2416 15.5 platelet-containing development grades 1 12 0.12 20 min 1356 214 2292 15.5 2 12 0.12 20 min 1387 244 2010 14.6 3 14 0.12 20 min 1323 184 2666 16.5 4 14 0.15 60 min 1385 254 1931 14.3 5 14 0.2 180 min 1382 229 2142 15.1 6 14 0.2 20 min 1361 228 2151 15.0 7 14 0.35 180 min 1329 147 3337 18.5 334 HM 35 H. Schön et al.

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Fig.11: SEM images of crack paths in a WC-0.35TiC-14Co alloy, showing different toughening mechanisms.

Conclusion

WC platelet-containing hardmetais can be successfully produced by the sintering of very-fine grained WC powders in the presence of small amounts of TiC or Ti(C,N) at comparatively high sintering temperatures. Both, uniform and non-uniform platelet microstructures can be formed, depending on the starting WC powders and subsequent processing conditions. Optimized platelet-containing alloys show superior hardness-toughness combinations compared with platelet-free commercial alloys. The formation of very large WC plates (sizes >10 urn) has to be avoided, since they lower the crack resistance and can act as fracture origins on loading. H. Schön et al. HM 35 335

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7. Acknowledgements

The present work was supported financially within the framework of a group project "Concepts for New Microstructural Design of Hardmetals by using Advanced Powders and Processing Techniques" in which the following companies participated: H.C. Starck, Kennametal, Mitsubishi Materials Inc., Sumitomo Electric Ind., Teledyne Metalworking Products, United Hardmetal and Wolfram Bergbau- und Hüttenges. The authors wish to thank Dr. Warbichler, Research Institute for Electron Microscopy, Graz, Austria for conducting TEM investigations, Prof. E. Lassner, Graz, for helpful discussions, and Dr. Y. Yamamoto (Allied Materials) for supplying us with WC powders produced via the direct carburization process.

8. References

1. M. Kobayashi, K. Kitamura, S. Kinoshita: EP 0 759 480 A1; European Patent Office, 1995. 2. M. Kobayashi: Machinery and Tools, Japan (written in Japanese), pp. 71- 76, February 1997. 3. T. Saito, Y. Taniguchi, M. Kobayashi, K. Kobori: Proc. Intern. Conf. on Powder Metallurgy and Particulate Materials, Las Vegas (ed. J.J. Oakes et al), Vol.1, 1-149, MPIF, Princeton, NJ, 1998. 4. Tungaloy, Product information: TD900 Grade Series, Toshiba Tungaloy, 1375 East Irving Park Road, Itasca, IL 60143, 1999. 5. G. Meadows: U.S. Patent Application 418,808; United States Patent Office, 1964. 6. G. Meadows: U.S. Patent 3,451,791; 1969. 7. G. Meadows: U.S. Patent 3,647,401; 1972. 8. L. Pons: Plastic anisotropy in WC; anisotropy in single crystal refractory compounds (ed: F.V. Vahldiek, J. Mersol); Vol 2; 393-444; Plenum Press; New York, 1968. 9. M.K. Brun, R.R. Neurgaonkar, V.S. Stubican; Precipitation studies in the system WC-TiC; J. Am. Cer. Soc; [58 No.9-10]; 392-95; 1975. 10. M. Kobayashi, S. Yamaya: Investigation of cemented tungsten carbides with twinned WC grains; Proceedings of the 10th Plansee Seminar, Vol. 1; pp. 597-611 (ed. H. Ortner, Metallwerk Plansee, Reutte,1981). 336 HM 35 H. Schön et al.

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11. K. Kobori, M. Ueki, T. Saito, H. Suzuki: Journal of the Japan Society of Powder and Powder Metallurgy, 36, 379-384, 1989. 12. S. Masahiro et. al.: JP 5339659; Japanese Patent Office, 1993. 13. S. Kinoshita, T. Saito, M. Kobayashi, K. Hayashi, J. of the Japan Society of Powder and Powder Metallurgy, 47, No.5, 526-33, 1999. 14. S. Kinoshita, T. Saito, M. Kobayashi, K. Hayashi, J. of the Japan Society of Powder and Powder Metallurgy, 48, No.1, 51-60, 2000. 15. K. Rödiger, K. Dreyer, M. Willert-Porada, T. Gerdes: International Patent WO 98/40525; Patent Cooperation Treaty (PCT) 1998. 16 A. V. Shatov ; S.A. Firstov: Material Science and Engineering; [A242]; 7; 1998. 17. H.E. Exner: Intern. Metals Reviews, No.4, 149-173, 1979. 18. B. Zeiler: Proc. of the 14th Int. Plansee Seminar, Vol.4, pp.265-276 (ed. G. Kneringer et al, Plansee AG, Reutte) 1997. 19. Y. Yamamoto, A. Matsumoto, Y. Doi: Proc. of the 14th Int. Plansee Seminar, Vol.2, pp.596-608 (ed. G. Kneringer et al, Plansee AG, Reutte) 1997. 20. D.F. Carroll: Proc. of the 14th Int. Plansee Seminar, Vol.2, pp.168-182 (ed. G. Kneringer et al, Plansee AG, Reutte) 1997. 21. W.D. Schubert, H. Neumeister, G. Kinger, B. Lux: Intern. J. of Refractory Metals & Hard Materials 16, pp. 133-142, 1998. 22. M. Sommer, W.D. Schubert, W. Zobetz, P. Warbichler: Proc. of the 15th Int. Plansee Seminar (ed. G. Kneringer et al, Plansee AG, Reutte), 2001, to be published. R. Frykholm et al. HM 44 337

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The effect of binder volume fraction on gradient sintering of cemented carbides

Robert Frykholm1, Malin Ekroth2, Bo Jansson3, John Agren2, Hans-Olof Andren1

1 Dept. of Experimental Physics, Chalmers University of Technology and Göteborg University, Göteborg, Sweden 2 Dept. of Materials Science and Engineering, Royal Institute of Technology, Stockholm, Sweden 3 AB Sandvik Coromant, Stockholm, Sweden

Summary:

To increase the cutting performance of cemented carbide tools, it is common to use a high temperature CVD process to coat them with thin wear resistant layers. During the process cracks are unavoidably introduced in the coating. To prevent crack propagation it is of interest to create a tough surface zone in the substrate, enriched in WC and binder phase. A way to create such a zone is to sinter a nitrogen containing cemented carbide in a nitrogen free atmosphere. This formation of gradient structures has been extensively studied using microscopy and simulations, and it has been shown that the process is driven by diffusion in the binder phase. However, the diffusion paths are partly blocked by the dispersed particles. This effect can formally be handled by considering effective diffusivities by introducing a so-called labyrinth factor, A. In prior work it has been assumed that A = f2, where / is the volume fraction of the binder. The validity of this assumption has been studied by simulations and experimental analysis of gradient sintered WC- Ti(C,N)-Co cemented carbides containing 5.0, 6.7, 10.0 and 20.0 vol.% binder phase. It was found that by using the labyrinth factor / instead of f2, a better correspondence between experiments and simulations can be achieved. 338 HM 44 R. Frykholm et al.

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Keywords:

EPMA; Microstructure; Diffusion; Thermodynamics; Computer simulation

1. Introduction

Cemented carbides, or hardmetals, are wear resistant hard materials, often in the form of cutting tool inserts, used for example for cutting, drilling or turning of metals. The basic cemented carbide consists of hard WC grains embedded in a Co-rich binder phase, but usually there is also a second hard phase present. This phase is a carbide or carbonitride containing, for example, one or several of the elements Ti, Ta and Nb. Addition of this second hard phase can make the material harder, but also more brittle. Cemented carbide tools are usually coated with a wear resistant layer to increase their performance further. The coating is performed using a high temperature chemical vapour deposition process (CVD). Because of a difference in thermal expansion coefficient in coating and bulk material, cracks are unavoidably introduced in the coating. The brittleness of the cemented carbide may then cause a problem since cracks in the coating might easily propagate into the base material and cause failure. A way to get around this problem is to use so called gradient sintering. The material in this study is a gradient sintered cemented carbide cutting tool insert, based on WC-Ti(C,N)-Co. Gradient sintering can be performed with N- containing cemented carbides. The material is then sintered in an N-free atmosphere, leading to an outward diffusion of N. The gradient in N-activity created by this diffusion will lead to an inward diffusion of Ti. In this way a surface zone free of the hard carbonitride phase, and enriched in Co and WC, is created. This tougher surface zone makes the cutting tool insert more suitable for coating since it makes it difficult for cracks to propagate. For the development of new gradient materials it is of great interest to be able to predict the formation of surface zones. A study of the influence of cobalt content on the gradient zone thickness was made by Schwarzkopf [1]. In the present study the software DICTRA has been used to simulate the formation of gradients [2, 3]. A model for diffusion in a continuous matrix with dispersed phases has been used [4]. In this case the matrix is liquid cobalt and the dispersed phases are the carbides and carbonitrides. Due to the presence of the dispersed phases the diffusion is reduced in the matrix. In earlier studies R. Frykholm et al. HM 44 339

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it has been assumed that a so-called labyrinth factor f2, where / is the volume fraction of the matrix, can be used in the simulations to reduce the diffusion coefficient matrix of the binder. The purpose of the present study is to investigate if this assumption gives a good description of the process. Since the formation of surface zones has been shown to be controlled by diffusion, a description of the labyrinth factor, as correct as possible, is crucial for correct simulation results. Therefore, the influence of the binder content on the gradient zone thickness has been studied using electron probe microanalysis and simulations.

2. Material and experimental

2.1 Material

The cemented carbide alloys used in this work were based on a mixture of WC, (Ti,W)C, Ti(C,N) powders and metallic Co powder. A set of powders were mixed with the aim of obtaining 5.0, 6.7, 10.0 and 20.0 vol.% binder phase and a constant ratio between volume fractions of cubic carbonitride phase and WC in the sintered state. Further, the N and C contents were controlled in order to obtain the same nitrogen and carbon activity at the sintering temperature for the four alloys. A pressing agent, which is removed at sintering, was added to the raw material powders. The mixture was ball milled with cemented carbide milling bodies in a liquid based on ethanol and subsequently spray dried. The sintering process contained several stages, including dewaxing, densification and gradient formation. For the gradient formation the samples were kept at 1450°C during 2 h in a nitrogen free atmosphere consisting mainly of Ar and CO. The chemical compositions of the sintered materials are given in Table 1. Full details about the mixing of the powders and the sintering process can be found in [2] 340 HM 44 R. Frykholm et al.

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Table 1. Chemical composition of the studied alloys (mass%).

Alloy Co Ti N C W 1 3.32 5.96 0.39 6.52 balance 2 4.53 5.64 0.40 6.40 3 6.81 5.47 0.38 6.24 " 4 13.85 4.99 0.33 5.71

2.2 Experimental

To study the elemental distribution, electron probe microanalysis (EPMA) was used. The analysis was performed on a Jeol JXA8900R combined WD/ED microanalyser operated at 15 kV. The electron beam was set to perform line- sweeps over 100 |j,m parallel to the insert surface, and was moved perpendicular to the surface with 1 jam between every sweep. For each sweep an average composition was calculated. In this way the cutting insert was scanned down to 500 \im from the original surface.

3. Simulations

3.1 The model

As already mentioned, a model for long-range diffusion occurring in a continuous matrix with dispersed phases has been used [4]. This diffusion model is available in the DICTRA software [5], It is assumed that all diffusion occurs in a matrix, which in this case is the liquid binder phase. The calculations are divided into two steps, one diffusion step, and one equilibrium step, where during the diffusion step gradients in activities drive the diffusion, and during the equilibrium step local equilibrium compositions are calculated. The Thermo-Calc [6] package is included in DICTRA and is used for all thermodynamic calculations. Due to the presence of a dispersed phases (the carbides and carbonitrides) the diffusion is reduced in the matrix (the liquid binder phase). It has earlier R. Frykholm et al. HM 44 341

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been assumed that a so-called labyrinth factor, A(f) where / is the volume fraction of the matrix, can be used in the simulations to reduce the diffusion coefficient matrix [4]. This is done by simply multiplying the diffusivity with the labyrinth factor. In earlier simulations the factor f2 has been used. Since typical binder fractions are 10-15 vol.% this leads to a rather drastic reduction in diffusivity. In this work the labyrinth factor / has also been tried. For simpler systems it is possible to calculate the labyrinth factor exactly [7], but for the complicated structure of a cemented carbide during liquid phase sintering this is not possible. Because of the low solubility of nitrogen in the matrix phase (liquid), an unrealistically high nitrogen supersaturation may be created in the matrix phase during the diffusion steps. This leads to numerical problems. The way to handle this is described elsewhere [2].

3.2 Thermodynamic and kinetic basis

A thermodynamic database [8] for carbon- and nitrogen- containing cemented carbides has been used for the Thermo-Calc calculations and DICTRA simulations. Model parameters for short-range order in the carbonitride phase were included in the database [9]. For the DICTRA simulations a kinetic description of the liquid is also needed. Lacking more detailed information it was assumed that all elements, i.e., Co, Ti, W, C, and N, have the same mobility in the liquid binder. A temperature dependent Arrhenius expression was used:

M= RT ^ni' M\

The activation energy, ß, was assumed to be 65,000 J/mol [10,11,12]. A value of the frequency factor, M°, was chosen in order to give approximately the experimentally measured gradient zone width. From the mobility it is possible to get the chemical diffusivity. In this case we get a diffusion coefficient matrix where diffusion of one element also depends on the concentration gradients of other elements. The coefficients of the matrix are given by: 342 HM 44 R. Frykholm et al.

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(2)

where xk is the mole fraction of element k, Mi is the mobility of element i, which here is assumed to be the same for all elements, /a-, is the chemical potential, and n is an arbitrary chosen reference element. To calculate the derivative of the chemical potential, access to thermodynamical data is necessary. Here this is solved by using the software Thermo-Calc. To get the effective diffusivity of the binder, equation 2 is multiplied with the labyrinth factor:

When using the labyrinth factor / at the sintering temperature of 1450°C this yields diffusion coefficients in the range D » 1 • 10"8 m2/s, which is quite reasonable. The flux of any species is then given by the multicomponent extension of Fick's first law:

dc, n" J

hieff QZ (4) where Jk is the flux of species k in the direction of the z-axis and dc^ldz is the concentration gradient of species j. It should be emphasised that eq. 4 yields the flux in the binder. In order to obtain the flux referring to a cross section of the material the flux given by eq. 4 must be multiplied with the volume fraction of the binder.

4. Results and discussion

Figure 1 shows experimental Co concentration profiles. The results were obtained using EPMA. Note that these are not profiles for the binder phase volume fraction, but only for the Co content. The binder phase also contains some W. The W content is determined by the C activity, which during R. Frykholm et al. HM 44 343

"IS-" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

sintering is almost uniform throughout the substrate. This results in a close to constant W content in the binder of approximately 3 at.% after cooling. Of the other elements in the cemented carbide virtually nothing will be left in the binder phase after cooling. This means that there will be a slightly larger volume of binder phase than that corresponding to the Co content, but this increment of volume due to W content is spatially independent.

Alloy

50 100 150 200 250 300 Distance

Figure 1. Co concentration profiles obtained with EPMA.

In figure 2 the same data as in figure 1 has been used, but the length scale has for each curve been normalised with respect to the corresponding zone thickness, and the Co bulk content has been set to zero, i.e. the bulk content of Co has been subtracted from the total Co content. As can be seen, all four curves match each other. It seems as if the Co content drops more steeply at the zone border for the materials with a high total Co content, but this is only an effect of the more compacted length scale due to thicker zones. 344 HM 44 R. Frykholm et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The amount of binder phase in the material does not affect the absolute variations around the bulk value. The only effect of a high Co content is that the larger binder volume allows for higher diffusion fluxes, and thus faster gradient growth. For all four alloys, the drop in Co content from the Co enriched region outside the zone border, to the slightly Co depleted region inside the zone border, is the same and approximately 15 at.%. This of course also means that the relative variations in phase distribution around the zone border will be larger with a small total binder phase content. Calculations give the result that this is what to expect. If the gradient formation is considered only as a replacement of carbonitride phase with binder phase, and one also allows for some of the W from the carbonitride phase to form WC to maintain local equilibrium, the difference in Co content when crossing the zone border should be approximately 17 at.%.

Alloy

3 4 5 Relative distance

Figure 2. Experimental Co concentration profiles with bulk value set to zero and length scale normalised with respect to the corresponding zone thickness. R. Frykholm et al. HM 44 345

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Figure 3 shows a simulation of the Co concentration profile for the material with 10 vol.% binder phase (alloy 3). The simulated profile shows the same type of appearance as the experimental profile. The largest Co content is found just outside the zone border, and inside the border there is a depletion. The difference in Co content between the enriched and the depleted region is 13 at.%, which is approximately the same as found experimentally.

25- i i i i i i i

20-

15-

O 10-

5- -

0- I I I : I II 0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 Relative distance

Figure 3. Simulation of Co concentration profile for alloy 3.

Figure 4 shows simulation results and experimental values of the gradient zone thickness vs. the cobalt content in the alloys. Due to natural variations in the cemented carbide it is not possible to exactly measure the gradient zone thickness. It is reasonable to believe that there might be an error of ± 3 urn present in the thickness measurements. There is also a variation in zone thickness due to two-dimensional diffusion near corners. Since the simulations are based on one-dimensional diffusion, thickness measurements were performed as far away from corners as possible. The dashed line in the 346 HM 44 R. Frykholm et al.

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figure shows the behaviour using the labyrinth factor f1, and the solid line using /. For the calculation of the highest cobalt content alloy the grid has been changed to be denser at larger distance. The results from the simulations are compared with experimental data. The agreement between simulations and experiments are satisfactory using the labyrinth factor/, confirming the hypothesis that the gradient thickness varies approximately linearly with the volume fraction cobalt in the cemented carbide. This indicates that the increase in diffusion distance, due to the curved diffusion paths around dispersed particles, is not as severe as previously thought.

120 gradient sintering 2h, 1723 K + simulation A.=f2 iooH x simulation A=f A experimental values / 3 80 I 60

20-

0 0 2 4 6 8 10 12 14 16 18 20 cobalt (mass%)

Figure 4. Gradient zone thickness vs. the cobalt content in the alloys. The dashed line shows the behaviour using f2, and the solid line using /. Triangles are experimental values. R. Frykholm et al. HM 44 347

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Extrapolation of experimental data will give a negative Co content for zero gradient. Still it is reasonable to expect this type of behaviour. As the Co content decreases the diffusion paths become longer, and for some critical value there will no longer be a continuous binder phase, and it will no longer be possible for a gradient to be formed. As the Co content decreases towards this value there will be a drastic drop in the zone thickness vs. cobalt content graph, and thus extrapolation of the values beyond this point is of no physical interest. To compare with earlier research in this area figure 5 shows experimental data from Schwarzkopf [1]. In the figure a growth rate constant, K, is plotted vs. the Co content. The constant is defined by

x2=K-t (5) where x is the gradient zone thickness and t is the sintering time. Schwarzkopf assumed a linear dependence between K and the Co content, i.e. he assumed that the gradient thickness varied with the square root of the Co content. If the data is extrapolated using this assumption, the x-axis is intersected at about 6 vol.% Co, which means that for alloys with a Co content lower than 6 vol.% there will be no gradient formation. This is a rather high value, and it is not reasonable since the present study shows that a total binder content of 5 vol.% is quite sufficient to create a gradient in the material. Also, figure 4 shows that there should be a linear dependence between zone thickness and Co content. When analysing Schwarzkopf's data further it is found that also here a linear dependence between zone thickness and Co content is the most correct assumption. In figure 6, data from the sintering performed at 1490°C is instead plotted in the same way as the data from the present study, with zone thickness vs. Co content. When plotted this way, the data shows a linear behaviour and extrapolates, as expected, to a slightly negative Co content for zero gradient formation. 348 HM 44 R. Frykholm et al.

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'« 6-101-3

£ 5-10"9

UJ 4-1Q"9 2T

& 3-10"9 o Si ! 2-10^ UJ

1.10* ,N0! G 5: X 2 4 6 8 10 12 14 16 18 20 Co - GEHALT Vu? Vol. %

Figure 5. Experimental data by Schwarzkopf, after [1].

i 0 2 4 6 8 10 12 14 16 18 20 cobalt (mass%)

Figure 6. Experimental data by Schwarzkopf replotted. The data are for the material sintered at 1490°C R. Frykholm et al. HM 44 349^

15" International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

5. Summary and conclusions

• The binder phase fraction of the alloys only affects the thickness of the surface gradient zone. Since all diffusion occurs in the liquid binder phase, the binder content determines the rate of gradient formation. A higher binder content allows for faster gradient formation because of higher diffusion fluxes in the material. • The absolute variations in Co content, and the overall gradient structure is independent on the amount of binder phase. • Due to presence of a dispersed phase (carbides and carbonitrides), the diffusion paths are partly blocked by the dispersed particles, which grow or dissolve as a result of diffusion in the binder. This effect can formally be handled by considering effective diffusivities. The diffusivity of the binder is multiplied with a so-called labyrinth factor, X, which is dependent on the volume fraction, /, of binder phase. • The labyrinth factor A(f) = f gives a better agreement between experiments and simulations than the factor f2 previously used. It can therefore be concluded that the gradient zone thickness and binder content show a linear relationship. • Earlier assumption that the gradient zone thickness is proportional to the square root of Co content has been shown to be wrong. A consequence of the earlier assumption would be that no gradient should be formed in materials with less than 6 vol.% Co. For the present study, however, gradient materials with a total binder content of 5 vol.% were manufactured. If the gradient zone thickness is instead assumed to be directly proportional to the Co content, extrapolation of data leads to zero gradient growth for approximately zero Co content, as expected.

Acknowledgements

Ulf Rolander and Per Gustafson at AB Sandvik Coromant are thanked for valuable discussions. AB Sandvik Coromant is gratefully acknowledged for financial support. 350 HM 44 R. Frykholm et al.

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References

1. Schwarzkopf, M., Kinetik der Bildungen von Mischkarbidfreien Randzonen auf Hartmetallen, Ph. D. Thesis, Montanuniversität Leoben (1987). 2. Ekroth, M., Frykholm, R., Lindholm, M., Andren, H.-O., Ägren, J., Acta Mater., 48, (2000), pp. 2177-85. 3. Frykholm, R., Ekroth, M., Jansson, B., Andren, H.-O., Ägren, J, "Effect of Cubic Phase Composition on Gradient Zone Formation in Cemented Carbides", to be presented at 7tn Int. Conf. on the Science of Hard Materials, Ixtapa, Mexico, March 5-9, 2001, to be published in Int. Journal of Refractory Metals and Hard Materials. 4. Engström, A., Höglund, L., Agren, J., Metall. Mater. Trans. A, 25A, (1994), pp 1127-34. 5. Andersson, J.-O., Höglund, L., Jönsson, B., Agren, J., Fundamentals and Applications of Ternary Diffusion, Purdy, G. R. ed., Pergamon Press, New York, (1990), pp 153-63. 6. Sundman, B., Jansson, B., Andersson, J.-O., CALPHAD, 9, (2), (1985), pp. 153-90. 7 McCullough, R. L., Composite Sei. and Tech., 22, (1985), pp 3-21 8. Frisk, K., Dumitrescu, L., Ekroth, M., Jansson, B., Kruse, O., Sundman, B., "Development of a Database for Cemented Carbides: Thermodynamic Modelling and Experiments", Research Report IM-2000- 029, Swedish Institute for Metals Research, Stockholm, (2000). 9. Ekroth, M., Frisk, K., Jansson, B., Dumitrescu, L.F.S, Metall. Mater. Trans. B, 31B, (2000), pp. 615-9. 10. Ono, Y., Yagi, T., Trans. Iron Steel Inst. Japan, 11, no. 4, (1971), pp. 275-79. 11. Faber, T.E., Introduction to the Theory of Liquid Metals, London, (1972), p. 159. 12. Kubicek, P., Trans. Iron Steel Inst. Japan, 22, no. 5, (1982), pp. 391-5. AT0200485 W.D. Münz et al. HM 96 351

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Microstructure, Composition and Performance of PVD Coatings Designed for Successful Dry High Speed Milling

Münz Wolf-Dieter (a'b), Lembke Mirkka I.(a), Lewis D. Brian (a), Smith lain J (b) a) Sheffield Hallam University, Materials Research Institute, Sheffield, S1 1WB, UK b) Bodycote SHU Coatings Ltd, Sheffield, S1 4QF, UK Summary: Dry high speed machining (HSM), particularly dry high speed milling, demands hard coatings, which exhibit high toughness, high oxidation resistance, a limited amount of residual stress and excellent adhesion to the cemented carbide (CC) substrate. These requirements are met by TiAICrYN coatings grown by the combined cathodic arc/ unbalanced magnetron deposition method. Fully sufficient adhesion is achieved by ion implantation of Cr into the CC prior deposition. Residual stress is controlled by an Y - free base layer; high oxidation resistance is provided by an Y - containing 3|im thick hard coating with 29 GPa hardness and a residual stress well below -7 GPa. Under the influence of temperatures above 800°C, Y segregates along the columns of TiAIN and plugs the in/ outdiffusion of elements. A top layer of Y - containing oxynitride reduces the friction against the work piece material (0.9 to 0.65). Cutting tools coated as such may be used for dry milling up to 25k rpm in steels HRC > 60.

1. Keywords High speed cutting, PVD, TiAIN hard coating, Yttrium, Oxidation resistance

2. Introduction Dry high speed machining is gaining increasing importance in modern production technology. Dry high speed milling of dies and moulds represents the current state of the art. Cutting speeds between 15k-25k rpm have been reported to machine steels in the hardness range of 55-62HRC. In this application cemented carbide cutting tools must be protected with hard oxidation resistant coatings as localised temperatures are typically up to a 1000°C. In this temperature range tungsten carbide but also PVD - TiN, PVD - TiCN coating material, suffer severe oxidation. Since the introduction of 352 HM 96 W.D. Münz et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

PVD - TiAIN, temperatures up to 800°C can be tolerated [1]. To increase the peak temperature still further, the addition of 1-2at% Y to TiAIN based coatings allows peak temperatures up to 950°C [2].

E TiAIN C '(0 A 4 (D TiAICrYN CD TiAICrYN +topcoat

500 600 700 800 900 1000 Temperature [°C]

Figure 1: Continuous thermogravimetric measurement. Heating rate was with 50°C/min to 400°C and with 1°C/min to 1000°C.

Thermogravimetric data shown in Figure 1 confirm the dramatic difference in the oxidation behaviour of tungsten carbide when compared with TiAICrYN. [2], [3]. Temperature measurements on the cutting tool indicate indeed the occurrence of cutting temperatures greater than 900°C (Table 1). These experiments have been evaluated by milling various steels with two flute 8mm diameter solid carbide end mills. The coatings described in this paper have been deposited by the combined cathodic steered arc/ unbalanced magnetron deposition method (Arc Bond Sputtering: ABS - Technology) [4], [5]. W.D. Münz et al. HM 96 353

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Workpiece Cutting Temperature Material DIN Hardness Conditions (after 5 minutes) (HRC) EN24 Vc =385 m/min, Ax C0.40Mn0.60Ni 1.6582 38 =3.8 mm, Radial < 700 °C 1.5Crl.lMo0.3 Feed=l mm, Feed =0.1 mm/rev, Spindle 16 000 rpm. P20 Vc =385 m/min, CO.32MnO.80Cr Ax= 3.8 mm, 1.60Mo0.40 1.2311 48 Radial Feed= 0.4 <700 °C mm, Feed= 0.1 mm/rev, Spindle 16 000 ipm. Vc= 510 m/min, Ax= 3.8, Radial Feed= 0.4 ~810°C mm, Feed= 0.1 mm/rev, Spindle 20 000 rpm. H13 Vc=385 m/min, C0.35Sil.0Cr5.2 1.2344 53 Ax= 3.8 mm, 5Mol.5Wl.25V Radial Feed= 0.4 ~810°C 0.30 mm, Feed= 0.1 mm/rev, Spindle 16 000 rpm. A2 Vc =385 m/min, Ax= 3.8 mm, C1.0Mn0.65Cr5 1.2363 58 Radial Feed= 0.4 ~ 880 °C .OMol.OVO.3 mm, Feed= 0.1 mm/rev, Spindle 16 000 rpm. Vc= 510 m/min, Ax= 3.8, Radial Feed= 0.4 mm, Feed= 0.1 ~910°C mm/rev, Spindle 20 000 rpm.

Table 1: Overview of temperatures at the cutting edge of coated cemented carbide ball nosed end mills. 354 HM 96 W.D. Münz et al.

15;h International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3. The substrate/ coating interface In dry high speed milling the tools are subjected to enormous interrupted reciprocating mechanical impacts and under these conditions adhesion is of paramount importance. Previous work has shown that the adhesion measured by scratch testing (critical load : Lc) can be directly related to the microchemistry of the interface resulting in localised epitaxial growth of the deposited TiAIN hard coating [6], [7]. The absolute value of the critical load Lc can be controlled by a dedicated low energy ion implantation process, carried out in a steered arc discharge, as process step of the ABS technology, during the metal ion etching procedure of the substrate surface (20min duration) prior to coating [4], [5]. The implantation of 1.2keV Cr++ ions prior to deposition of a TiAIN coating led to the highest scratch test values (Lc = 135N), when compared with a pre-treatment of 1.2keV Ar (l_c = 35N), Nb (Lc = 90N) and V (l_c = 115N) ions. Additionally a decrease of adhesion was observed, when the accelerating voltage during ion implantation was only 0.6keV (Cr; Lc = 60N). Figure 2 reflects the life-time of TiAICrYN coated cemented carbide end mills machining HRC 58 A2 die steel (cutting speed 15k rpm). In these experiments two-flute end mills fabricated from identical cemented carbide material were pre-treated with various ion species and then coated with 3|um TiAICrYN using identical deposition parameters.

40 Crl200V_2 Crl200V 1 135 c

•=30

J 25 Crl200V5min Cr600V 2 Arl200V 20 20 40 60 80 100 120 140 LcinN Figure 2: Life time in dry high speed machining versus critical load values in scratch adhesion testing W.D. Münz et al. HM 96 355

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Figure 3: Local epitaxial growth of TiAIN on tungsten carbide after Cr ion bombardment at Ub=-1200V (a) Selected area diffraction pattern (SADP) of TiAIN grain (b) superimposition of the TiAIN pattern and tungsten carbide pattern (c) SADP of tungsten carbide grain.

Selected area electron diffraction patterns (Figure 3) explain the excellent adhesion of TiAIN based coatings deposited on cemented carbide after a 1.2keV Cr++pre-treatment. Figure 3a) shows the electron diffraction pattern of fee TiAIN, whereas Figure 3c) represents the diffraction pattern of hexagonal tungsten carbide. The superimposed patterns of fee TiAIN and hexagonal tungsten carbide show perfect matching of the lattices at the interface between coating and substrate, which is a direct indication of localised epitaxial growth. 4. Successful Layer Architecture TiAICrYN based coatings with nanolayered Y incorporation have been proven particularly successful in dry high speed milling operations. However, these coatings exhibit relatively high compressive residual stresses up to - 7GPa. To retain sufficient adhesion a stress reduced 250nm thick TiAICrN base layer is employed to grade the stress between substrate interface and TiAICrYN coating [8]. In addition, a 400nm thick special toplayer of has been developed to achieve a reduced friction coefficient (0.9 to 0.65) of the wear resistant ceramic hard coating against steel as work piece material. This toplayer consists of a TiAIYN/CrN nanoscaled multilayer, which is embedded in a TiAICrY - oxynitride. The oxygen content is continuously increased during the deposition of the outermost 200nm of the coating by continuous replacing in the reactive sputter deposition process nitrogen with medical dry air [9]. Figure 4 shows a schematic diagram (a) and a TEM cross section (b) of the coating. Detailed description of the PVD deposition process is given in [4], [6], [5], [9]. This coating has been made commercially available as Supercote 11 by Bodycote SHU Coatings Ltd., Sheffield. 356 HM 96 W.D. Münz et al.

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

TiAICrYN + Topcoat

Base Layer Ion Implanted Zone a) Substrate

Figure 4: TiAICrYN with topcoat as a) schematic and b) cross-sectional bright field image W.D. Münz et al. HM 96 357

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5. Diffusion of Substrate Elements at High Operating Temperature Previous work on stainless steel substrates revealed a rapid diffusion of substrate material elements through the Y free coating above 800°C [2]. In extreme cases (1h at 950°C) cavities started to form at the interface between the stainless steel substrate and the hard coating, which leads automatically to a loss of adhesion (Figure 5). The outdiffusing substrate elements form complex mixed oxides at the coating surface containing both substrate and coating elements. These disintegrating effects can be completely avoided by incorporation of 1at% Y [2]. The incorporation of Y during deposition led to a fine grained interrupted columnar growth structure, which increases the diffusion path for in- and outward diffusion of substrate elements. Furthermore, it has already been reported that Y diffuses preferentially to the column boundaries after heat treatment for 1h at 950°C, thus further blocking this diffusion path [2]. In case of cemented carbide as substrate material, we have carried out extensive investigations using STEM and EDX mapping on cross sectional samples to localise the influence of Y. Distribution maps of the coating and substrate elements taken after heat treatment 1h at 900°C are shown in Figure 6. Whereas all elements such as AI, Ti and Cr are uniformly distributed in the coating, there is clear evidence of out-diffusion of Co from the cemented carbide substrate into the base layer. In parallel, we found a very uniform distribution of Y in the TiAICrYN coating. Figure 6 is a clear indication that Co diffuses mainly along column boundaries in the base layer and is blocked in the Y containing layer section. At the higher magnification (Figure 7) the above mentioned Y segregation taking place in the Y containing part of the coating becomes evident, too. 358 HM 96 W.D. Münz et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

TiAlCrYN

TiAICrYN+Ox on CC after 1h at900°C Figure 6: Elemental distribution maps of TiAlCrYN with topcoat deposited onto cemented carbide after 1h at 900°C.

Figure 7: YK map of TiAlCrYN with topcoat. 6. Oxidation Behaviour of TiAlCrYN with Topcoat The progress in improvement of oxidation resistance is shown in Figure 1. The oxidation resistance of TiN is substantially improved by the addition of AI and incorporation Y. The TG analyses indicate, that the oxidation behaviour of the layer described in Figure 4 shows an even improved performance when compared with the toplayer free TiAlCrYN. STEM - EDX point analysis (Figure 8) explains the important role of Cr in this context. The interlayer consisting of the nanoscale multilayer TiAIYN/CrN seems to be fully intact after the heat treatment in air. We observed in this region a decrease in N and the expected increase in O content stemming from the exchange of the reactive gases during the deposition process and indicating the formation of TiAICrY - oxynitride. In the outermost region we found discrete particles of TiO2 (rutile), AI2O3 (corundum) and Cr2O3 (eskolite). In the diagram shown here AI2O3 is predominant. XPS mapping of the surface confirmed the existence of the above mentioned discrete particles [8]. W.D. Münz et al. HM 96 359

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TiAICrYN+Ox after 1h at 900°C

!- Y x 10 -o—N (nol in at%) - -O {not in at%)

2000 2100 2200 2300 2400 distance from substrate/ coating interface [nm] Figure 8: EDX point analysis through the topcoat region of TiAICrYN after heat treatment for1hat900°C. 7. Results in Dry High Speed Milling Finally, the performance of the here described TiAICrYN with the Cr containing topcoat (Supercote 11) is compared with commercially available Y free TiAIN coatings deposited by cathodic arc evaporation (supplier I and II) and an increasingly popular sandwich coating TiN/TiAIN (-30 layers), also deposited cathodic arc [10]. All tools were manufactured by the same tool supplier using cemented carbide material stemming from an identical production lot. In the tests 8mm two flute ball nosed end mills were used to mill A2 (HRC 58) die steel. Figure 9 summarises the results of end of life tests 1 for two different cutting speeds, namely v0 = 385 m min" (16krpm) and vc = 500 m min"1 (20krpm). The tests at 16krpm illustrate the superiority of Supercote 11 against arc deposited TiAIN and the multilayer system TiN/TiAIN. In the 20krpm test Supercote 11 is compared with Supercote 02 a coating with a similar composition to Supercote 11 however, without the incorporation of Y. At these high cutting temperatures (see table 1, ~800°C) the influence of Y becomes again clearly evident. One at% Y doubles the lifetime of TiAIN 360 HM 96 W.D. Münz et al.

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vc=500m/min 20krpm

Sputter Arc TiAIN II Arc TiAIN I Arc TIN/ Sputter Sputter TiAiCrYN TiAIN TiAiCrYN TiAiCrYN Supercote Supercote Supercote 11 2 11

Figure 9: Milling test results from various coating grown on cemented carbide ball nosed end mills.

Acknowledgements The investigations have been partially financed by the DTI - Link Surface Engineering Project GK/ K 76351 and by the Brite Euram Project Hidam (No: BRPR CT 98-0753). Particular gratitude is expressed to Hydra Tool International Pic Sheffield, and their managing director Mr A.P. Deeming, who have supported the project by supplying the cemented carbide test tools. We also like to acknowledge the Department of Materials at the University of Oxford for the use of the STEM HB501 microscope. The encouraging support of Prof. J.M. Titchmarsh was extremely helpful and is highly appreciated.

References 1. W.-D. Münz, Titan aluminium nitride films: A new alternative to 777V coatings. Journal of Vacuum Science and Technology A, 1986. 4: p. 2717. 2. L.A. Donohue, I.J. Smith, W.-D. Münz, I. Petrov and J.E. Greene, Microstructure and oxidation resistance of TiAiCrYN layers grown by combined steered-arc/ unbalanced magnetron sputter deposition. Surface and Coatings Technology, 1997. 94-95: p. 226-231. 3. M.I. Lembke, D.B. Lewis and W.-D. Münz, Localised oxidation defects in TiAIN/CrN superlattice structured hard coatings grown by cathodic arc/ unbalanced magnetron deposition on various substrate materials. Surface and Coatings Technology, 2000. 125: p. 263-268. W.D. Münz et al. HMJ36 361_

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4. W.-D. Münz, D. Schulze and F.J.M. Hauzer, A new method of hard coatings: ABS (arc bond sputtering). Surface and Coatings Technology, 1992. 50: p. 169-178. 5. W.-D. Münz, C. Schönjahn, H. Paritong and I.J. Smith, Hard Coatings Deposited by Combined Cathodic Arc Evaporation and Magnetron Sputtering (Arc Bond Sputtering: ABS). Le Vide, 2000. 3/4(297): p. 205. 6. C. Schönjahn, LA Donohue, D.B. Lewis, W.-D. Münz, R.D. Twesten and I. Petrov, Enhanced adhesion through local epitaxy of transition- metal nitride coatings on ferhtic steel promoted by metal ion etching in a combined cathodic arc/ unbalanced magnetron deposition system. Jornal of Vacuum Science and Technology A, 2000.18(4): p. 1718. 7. C. Schönjahn, Surface treatment in a cathodic arc plasma - key step for interface engineering, in Materials Research Institute. 2001, Sheffield Hallam University: Sheffield. 8. D.B. Lewis, L.A. Donohue, M. Lembke, W.-D. Münz, R.K. Jr., V. Valvoda and C.J. Blomfield, The influence of the yttrium content on the structure and properties of TiAICrYN PVD hard coatings. Surface and Coatings Technology, 1999. 114: p. 187-199. 9. M.I. Lembke, D.B. Lewis, W.-D. Münz and J.M. Titchmarsh, Significance of Y and Cr in TiAIN hard coatings for dry high speed cutting. Surface Engineering, 2001.17. 10.E. Zoestbergen, N.J.M. Carvalho and J.T.M. DeHosson, Stress state of TiN/TiAINPVD Multilayers. Surface Engineering, 2001. 17(1): p. 29-34. AT0200486 362 GT43 I. Smurov et al.

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Neue Oxid-Composite Beschichtung für schwierige Zerspanungsaufgaben

Hartmut Westphal, Henk van den Berg, Volkmar Sottke, Ralf Tabersky

WIDIAGmbH, Essen (Germany)

Summary

The changes in today's metal working technology are driven by increasing cutting speeds, heavy/hard machining and an enormous amount by changes in workpiece materials. These applications are asking for more tailor made cutting tool solutions. Together with the well established multi component coating technology a new approach of composite coatings is giving solutions for the tough demands of the cutting tool market. In this paper is presented composite coatings of AI2O3/ZrO2/TiOx made by CVD. The coating is like High Performance Oxide Ceramics for cutting applications. The coating is used in combination with MT CVD coatings and different carbide substrates. The CVD coating has optimum stress for cutting applications, low friction and very high thermal isolation. The outstanding performance of this coating is demonstrated in different applications.

1. Einleitung

Um im Fahrzeugbau Gewichte zu verringern oder im Maschinen- und Anla- genbau Aggregate mit höherer Leistungsfähigkeit zu bauen, werden zuneh- mend höherfeste metallische Werkstoffe eingesetzt, die meistens eine un- günstigere Zerspanbarkeit aufweisen. Gleichzeitig besteht die permanente Forderung nach effektiveren Bearbeitungsverfahren, was für Spanungsopera- tionen mit geometrisch bestimmter Schneide bedeutet, schnell, mit großem Zeitspanvolumen, möglichst trocken und dabei prozesssicher zu arbeiten. Nach dem CVD-Verfahren erzeugte Mehrkomponentenschichten bieten zur Lösung solcher Aufgaben ein interessantes Leistungspotenzial [1]. Den An- forderungen der Zerspanungstechnik Rechnung tragend und unter Anwen- dung der CVD-Mehrkomponenten-Beschichtung wurde auch in Anlehnung an Erfahrungen aus der Hochleistungsschneidkeramik eine neue Oxid- I. Smurov et ai. GT43 363

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Composite Beschichtung zur Lösung schwieriger Bearbeitungsaufgaben und für HSC-Anwendungen entwickelt.

2. Aufbau und Herstellung der Oxid-Composite Beschichtung

Diese wird nach dem Niederdruck CVD-Prozess erzeugt und ist mit MT-CVD Schichten kombiniert. Die aus 3 Komponenten bestehende Oxidschicht wird aus einem aus AICI3, ZrCI4, TiCI4, CO2, H2, N2 und Ar bestehenden Gasge- misch kondensiert. AI2O3 liegt danach als K-Phase und ZrO2 in Abhängigkeit von der Beschichtungstemperatur >990°C als Phasengemisch tetrago- nal/monoklin und bei Beschichtungstemperaturen <960°C monoklin vor. Da das ZrO2 in diesem Schichtsystem nicht stabilisiert ist, ist davon auszugehen, dass sich tetragonales ZrO2 teilweise mit Volumenzuwachs in monoklines ZrO2 umwandelt. Die aus AI2O3, ZrO2 und TiOx bestehende Composite Schicht erinnert auch hinsichtlich ihrer Zusammensetzung mit 5 bis 15% ZrO2 sowie 0,5 bis 5% TiOx und Rest AI2O3 an Hochleistungsschneidkeramiken ähnlicher Zusammensetzung. In Abb. 1 ist die Composite Schicht in Kombi- nation mit MT-CVD TiCN dargestellt. In der dargestellten Ausprägung ist dieses Schichtsystem auf allen gängigen Substrathartmetallen anwendbar.

Kalottenschliff Bruchgefüge

Composite

Al2O3/ZrO2/TiOv

Ti(C,N) (MTCVD) -*•

Abb. 1: Kalottenschliff und Bruchgefüge einer TiN-TiCN(MT)-AI203/Zr02/Ti0x- Schicht 364 GT43 I, Smurov et al,

15'° International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Aus dem Schliffbild ist ersichtlich, dass die ZrO2-Teilchen in die AI2O3 Matrix eingelagert sind. TiOx ist im Schliff und im Bruchgefügebild nicht sichtbar, es ist z. B. mittels EDX-Analyse gut nachweisbar, bisher aber als Phase nicht eindeutig identifiziert [2]. Wegen seiner nanodispersiven Verteilung hat TiOx eine positive Auswirkung auf das Schichtwachstum und die Korngröße der beiden anderen Oxide.

3. Eigenspannungen der Oxid-Composite Beschichtung

Für beschichtete Schneidkörper ist die Einstellung von Druckeigenspannun- gen in der Hartstoffbeschichtung wünschenswert, weil Druckeigenspannun- gen der Rissentstehung und der Rissausbreitung entgegenwirken und deshalb vorzugsweise beim Zerspanen mit Schnittunterbrechungen höhere Standzeiten erwartet werden können.

Abb. 2 zeigt Fräs-Standzeitvergleiche von Proben mit Zugeigenspannungen und mit Druckeigenspannungen in der Hartstoffschicht. Die Beschichtung ist identisch mit der in Abb. 1 dargestellten, Substrathartmetall und Geometrie sind ebenfalls identisch. Im AI2O3 der Composite Schicht wurden die Eigen- spannungen röntgenographisch nach dem sin2\|/-Verfahren [3] an jeweils 3 unterschiedlichen Gitterebenen bestimmt und der Mittelwert gebildet. Die für die unterschiedlichen Gitterebenen erhaltenen Spannungswerte differieren im Mittel um ca. 20%. Das Vorzeichen - (minus) steht für Druckeigenspannun- gen.

Noch deutlicher als im Standzeitdiagramm sichtbar, zeigt sich der Unter- schied zwischen den Schneiden mit Zug- und Druckspannungen in der Aus- bildung des Verschleißes. In Abb. 3 sind solche Schneiden mit Sicht auf die Spanfläche abgebildet. Während die Schneide mit Zugspannungen in der Beschichtung eine große Ausplatzung über die Nebenschneide zeigt, ist die mit Druckspannungen normal verschlissen.

Die Änderung des Spannungszustandes von Zug- in Druckspannungen wurde durch Mikrostrahlen mit Korund erzeugt. Weiterhin zeigte die Untersu- chung der Eigenspannungszustände in der unter der Oxidschicht befindlichen TiCN-Zwischenschicht eine deutliche Verkleinerung der Zugspannungen und die in der WC-Phase des Substrathartmetalls gemessenen Werte eine Erhö- hung der Druckeigenspannung im Vergleich zu den nicht gestrahlten Proben. H. Westphal et al. HM 97 365

15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

D =-3800Mpa D =+350MPa

7000 6000 1 5000 »4000 + I 3000-

^ 2000 1000 0 vc=200m/min

Abb. 2: Einfluss der Eigenspannungen in der Beschichtung auf die Standzeit beim Fräsen von GGG70

Abb. 3: Verschleißbilder auf der Spanfläche nach Standzeitende aus Frästest gem. Abb. 2 - links Beschichtung mit Zugspannung, - rechts Beschichtung mit Druckspannung. 366 HM 97 H. Westphal et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

4. Anwendungen der Oxid-Composite Beschichtung

Zur Erhöhung der Lebensdauer von Schmiedegesenken werden diese mit Stellitaufschweißungen gepanzert. Sicher wäre diese Technologie schon wei- ter verbreitet, wenn die meist 2 bis 3 mm starke Stellitpanzerung nicht so schwierig zu bearbeiten wäre. Wie die Abb. 4 verdeutlicht, sind die im For- men- und Gesenkbau zum HSC-Fräsen herkömmlich beschichteten Hartme- tallschneidkörper für die Bearbeitung von Stellitpanzerungen eher ungeeig- net. Sie versagen schon bei relativ geringen Schnittgeschwindigkeiten wegen Aufbauschneidenbildung und thermischer Überlastung. Der in der Composite Beschichtung enthaltene ZrO2-Anteil trägt offensichtlich erheblich zum Erfolg bei dieser Fräsbearbeitung bei.

VB max • K10+Composite [mm] -K10-TiAlN Fräsen auf Stellit21EHL 2,5 RDMW 0802MO

vc (siehe links) 1,5 ap = 0,3 mm ae = 4 mm fz = 0,4 mm Fräsweg = 2000 mm

80 100 200 300 400 500 600 Vc [m/min]

Abb. 4: Verschleißverhalten bei der Bearbeitung aufgeschweißter Stellitpanzerungen

Bekanntlich hat ZrO2 einen geringen Reibwert zu Bestandteilen des Stellit und eine sehr geringe Wärmeleitfähigkeit, die den Schneidkeil zusätzlich vor thermischer Überlastung schützt [4]. Um den Nachweis zu führen, dass dieser Schneidstoff sowohl zum Schruppen als auch zum Schlichten von Stellit- panzerungen geeignet ist, wurde über einen Vorschubbereich fz von 0,2 bis 0,6 mm getestet. Die Ergebnisse sind in Abb. 5 zusammengefasst dargestellt. H. Westphal et al. HM 97 367

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Material: Stellit21EHL Werkzeug: Torusfräser D = 16 mm, r = 4 mm Schneidstoff: Widia W6-42 Typ RDMW0802MO Beschichtung: A12O3/ZrO2/TiOx KSS: Druckluft 2 bar Parameter: ap = 0,3 mm, ae = 4 mm Fräsweg: s = 2000 mm

3 0,21 I o.o 100 200 300 400 500 Schnittgeschwindigkeit vc [m/min] Abb. 5: Verschleißverhalten bei der Bearbeitung aufgeschweißter Stellit- panzerungen in Abhängigkeit von der Schnittgeschwindigkeit und des Zahnvorschubs 368 HM 97 H. Westphal et al.

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Wie in Abb. 5 darstellt, gibt es um 400m/min Schnittgeschwindigkeit einen nur unwesentlichen Einfluss des Zahnvorschubs auf die Standzeit. Die ganze Un- tersuchung wurde im Einzahnversuch durchgeführt.

Nach der für den verwendeten Stellit angegebenen Richtanalyse enthält das Material 0,15 - 0,30 % C / 30 - 33 % Cr / 4,5 - 5,5 % Mo / 3,0 - 4,0 % Ni / Rest Co.

In Praxiseinsätzen konnten die hier ermittelten Schnittwerte gut reproduziert werden. Für die Gesenkhersteller bedeutet dieses Ergebnis, dass für die Be- arbeitung der Stellitpanzerungen ähnliche oder gleiche Schnittwerte anwendbar sind, wie sie für den Grundkörper aus Werkzeugstahl verwendet werden, wenn der neuentwickelte Schneidstoff aus feinkörnigem K10 Substrathartme- tall mit der in Abb. 1 dargestellten Beschichtung eingesetzt wird.

In einer weiteren Anwendung wurde der beim Stellitfräsen erfolgreiche Schneidstoff mit dem in Abb. 6 abgebildeten Fräser für die Zylinderkopfbear- beitung von Truck-Dieselmotoren verwendet. Wie beim Stellitfräsen ergibt sich auch hier eine deutliche Leistungsverbesserung bei höherer Schnittge- schwindigkeit (Abb. 7).

Abb. 6: Planfräser M750 Hexacut (Werksfoto WIDIA) H. Westphal et al. HM 97 369

IS1' International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. WiJdner, Plansee Holding AG, Reutte (2001), Vol. 4

Standzeit pro Schneide in min

M750 (Hexacut) 60°

HNGX-MM

vc = 400 m/min ap = 2,0 - 4,0 mm Planfräser 75° fz = 0,14 mm

Z = 20 SPKN 1204 EDR SPFR 1204 EDR

K10+TiN-TiCN- K10+TiN-TiCN-AI2O3 K10+TiN-TiCN-AI2O3 AI2O3/ZrO2/TiOx Abb. 7: Planfräsen Zylinderkopf GG25 mit Sinterhaut (trocken)

Die mit Abstand umfangreichste Anwendung findet die Oxid-Composite Be- schichtung beim Außenfräsen von Kurbelwellen [5]. Diese Bearbeitung ist immer trocken, in Abhängigkeit vom Material der Kurbelwellen mit Schnittge- schwindigkeiten zwischen 250 und 350 m/min bei der Zapfenbearbeitung (Hub- und Mittellager) und beim Topping, sowie mit 180 bis 250 m/min an den Wangen. Der Zahnvorschub (hmax) wird mit 0,20 bis 0,30 mm eingestellt. Problematisch ist die Gußhaut bei Kurbelwellen aus GGG-Material, ungleiche Aufmaße und Schmiedekruste sind typisch für Stahlkurbelwellen. In Abb. 8 ist ein solches Fräswerkzeug zur Lagerzapfen-, Unterstich- und Wangenbearbei- tung zu sehen. 370 HM 97 H. Westphal et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Abb. 8: Fräswerkzeug zum HSC-Außenfräsen von Kurbelwellen (Werksfoto WIDIA HEINLEIN)

Anzahl der bearbeiteten Zapfen

10000 7

9000

8000 WSP- Geometrie: spezial

7000 Schnittbedinqunqen:

6000

vc = 320 m/min 5000

4000 ap = 2,0-2,9 mm

3000 hma>i = 0,24 mm 2000 Trockenbearbeitung 1000

0 I TiN/Ti(C,N)/AI 2O3-Schicht TiN/Ti(C,N)/AI 2O3-ZrO 2-TiO ^-Schicht Substrat: K10- Substrat: K10- Feinstkorn-HM Feinstkorn-HM Abb. 9: Standzeit beim Hochgeschwindigkeitsaußenfräsen von Kurbel- wellen aus GGG70 (Unterstich Zapfen)

Die außerordentliche Verschleißbeständigkeit der Oxid-Composite Beschich- tung verdeutlicht der Standzeitvergleich in Abb. 9. Die Verbesserung wurde in H. Westphal et al. HM 97 371

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diesem Falle ausschließlich durch die Änderung der oxidischen Deckschicht von reinem AI2O3 auf AI203/Zr02/Ti0x erzielt.

Die in diesem Beitrag beschriebene Oxid-Composite Beschichtung wird zur Zeit auch für Drehoperationen mit gutem Erfolg in der Schwerzerspanung er- probt.

5. Literaturverzeichnis

[1] H. Westphal, U. König, H. van den Berg, R. Tabersky, V. Sottke: Neue Hartstoffbeschichtungen auf Basis der Carbonitride von Titan und Zirkon; 14. Int. Plansee-Seminar 1997, Vol. 3, C7, S. 55-62

[2] H. van den Berg, H. Westphal, R. Tabersky, V. Sottke und U. König: New Multi-Component and Multi-Phase Coatings by CVD and PVD; EURO PM'99 Advances in Hard Materials Production, Turin, 8.- 10.11.1999

[3] K. Kamachi, T. Ito and T. Ymamoto: Comparison of residual stress in cemented carbide tips coated with TiN by the CVD and PVD processes and their effect on failure resistance; Surfacing Journal International 1986, Vol. 1, No. 3, S. 82-86

[4] H. Westphal, H. van den Berg, V. Sottke: Neuere CVD-beschichtete Hartmetalle für die Metallbearbeitung; Spanende Fertigung, 3. Auflage, Vulkan-Verlag Essen, erscheint zur EMO 2001

[5] H. Westphal, H. van den Berg, V. Sottke; mav Innovationsforum Werk- zeugtechnik 10, 2000, S. 74-75 AT0200487 372 HM 101 R. Haubner et al.

15'" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

LOW-PRESSURE c-BN DEPOSITION - IS A CVD PROCESS POSSIBLE ?

Roland Haubner, Xinhe Tang Institute for Chemical Technology of Inorganic Materials, TU-Vienna, Getreidemarkt 9/161, A-1060 Vienna, Austria

Summary Since the low-pressure diamond deposition was discovered in 1982 there is a high interest to find a similar process for the c-BN synthesis. During the last years many researchers propagated success, but many of them failed in analytical problems characterizing the BN coatings. A new paper by S. Matsumoto and W. Zhang shows c-BN deposition using a DC Plasma Jet and characterization by various analytical methods. Our experiments were carried out with a simple CVD reactor without plasma activation. Various BN layers were deposited and characterized. X-ray diffraction, IR-spectroscopy and SIMS indicate that BN-coatings containing c-BN were deposited. However a final verification of c-BN crystallites by TEM investigations was not possible till now.

Key words: c-BN, Physical Vapor Deposition (PVD), Chemical Vapor Deposition (CVD)

1. Introduction Boron and nitrogen are neighbors of carbon in the periodic table and therefore BN phases are isoelectric to the corresponding carbon phases. Because of structural similarities between h-BN and graphite Wentorf [1] successfully tried the high-pressure high-temperature (HP-HT) synthesis of c-BN in 1957, equally to the diamond synthesis. The second hardest material - after diamond - was bom. After the first high-pressure experiments the B-N phase diagram was designed [2] and after some modifications the c-BN was described as metastable phase at room temperature [3]. Contrary to this opinion it was reported in 1988 that c-BN is the stable phase and many experiments R. Haubner et al. HM 101 373

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

confirmed this result, but exact thermodynamic data are still not available [4,5,6]. Therefore conversion experiments of c-BN into h-BN were carried out under normal and high pressure. Such experiments are helpful to understand the BN phase diagram. The transformation at normal pressure was investigated during DTA-analysis and exhibits the stability of c-BN under standard conditions [7]. Influences of c-BN grain size and purity (oxide content) on the conversion temperature become obvious. The lowest conversion temperature observed was at about 900°C for 1.5 urn c-BN powders containing boron oxide. Based on SEM images of various conversion stages it could be shown that the reaction can be explained by two mechanism: « Solid state mechanism: The c-BN reacts to the intermediate rhombohedral BN phase (r-BN) which is at least transferred into the hexagonal BN modification. • Gas phase mechanism: This is possible because h-BN shows a vapor pressure (1 mbar at 1300°C [8,9]). According to this vapor transport the reaction (phase transformation) is not reversible. High pressure transformation of h-BN into c-BN (at 6.5 GPa) and the reverse transformation of c-BN into h-BN (from 0.6 to 2.1 GPa) were investigated in a Li3N-BN catalyst system [10]. These experiments showed that the phase boundary line in the BN phase diagram should be at about 500°C (extrapolated to normal pressure) [10].

Because HP-HT synthesis for c-BN and diamond work under similar conditions, it was considered that the low-pressure synthesis of c-BN should be possible equally to that of diamond. But a lot of differences between c-BN and diamond make a low-pressure deposition of c-BN rather difficult: • c-BN is the stable phase - diamond is metastable (at standard conditions) • c-BN consists of two elements - diamond is only carbon • selective etching of h-BN is complicate - graphite can be etched easily by atomic hydrogen • stabilization of c-BN surface is difficult - diamond surface is stabilized by hydrogen • complex precursors are required to deposit c-BN - for diamond methane is convenient • analytical characterization of c-BN is complex - diamond can be identified easily by Raman. 374 HM 101 R. Haubner et al.

15° International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Deposition of nano-cBN films is possible with ion-assisted CVD and ion- assisted PVD techniques. The amount of grain boundaries in such materials is rather high and therefore these materials should be called nano-cBN instead of "pure c-BN".

The PVD methods for nano-cBN deposition are: • ion-beam-assisted deposition (IBAD) [11] • mass selected ion beam deposition (IBD) [12] • ion plating [13] • RF- or magnetron sputtering [14] and • laser deposition [15].

Most methods for nano-cBN deposition generate ions, which are accelerated and hit the substrate surface. Parameters for the substrate bias (ion energy, ion mass, etc. [16]) are mostly similar and therefore an equal growth mechanism was considered for the different methods. By starting the deposition process commonly an oriented h-BN layer is deposited onto the substrate surface. On this interlayer the c-BN nucleates and forms a layer [17]. Because of the deposition conditions and the ion- bombardment - which are necessary for the c-BN deposition - crystal defects and secondary nucleation occur and the single crystalline areas are very small (nm range) [18].

To describe the nano-cBN deposition four models have been proposed: • the compressive stress model [19,20] • the subplantation model [21,22] • the selective sputter model [23] and • the momentum transfer model [24].

Plasma CVD methods can be applied to activate the gas phase. Using conventional thermal CVD various BN modifications but no c-BN or nano-cBN are formed. Typical plasma CVD deposition methods are: • ECR plasma CVD [25,26]; . ICP CVD [27]; • Bias enhanced Plasma CVD [28] and • DC Plasma Jet [29]. R. Haubner et al. HM 101 375

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Selective etching of h-BN compared to c-BN is one possibility to increase the c-BN grain size in the layers. For a successful CVD-synthesis of c-BN a reaction system has to be found where c-BN is deposited and h-BN and amorphous BN formation can be prevented or be removed by selective etching during the deposition. In order to find a substance which allows selective etching of h-BN several compounds (e.g. atomic hydrogen, chlorine, mixtures of chlorine with hydrogen, BF3) were tested. From these substances BF3 showed the best selectivity for etching the h-BN compared to c-BN [30,31]. The disadvantage of BF3 - in contrast to atomic hydrogen in the carbon system - is the smaller etching rate of h-BN (about 10 times lower than in the carbon / atomic hydrogen system). The consequence is that only at very low BN growth rates the selective etching becomes relevant. Matsumoto et al. [29] tried a combination of bias assisted plasma jet and the reactive gas mixture Ar-N2-BF3-H2 for selective etching of h-BN [32]. Using a d.c. bias voltage the nano-cBN grain size in the layers could be increased from 7 nm (-150 V) to 12 nm (-80 V) [33]. By optimizing the process nano- cBN coatings thicker than 20 urn and c-BN grain size up to 100 nm could be deposited.

The simple chemical way ? Several attempts to grow c-BN at normal pressure were performed. Similar to the high-pressure high-temperature synthesis melts were used in the temperature range between 600 and 800°C at normal pressure [34]. Various compounds were mixed with fine grained c-BN powder (for seeding) and then heated up to allow grain growth. The stability of c-BN in chemical active media depends on the properties of the reacting agent. Reductive media like lithium metal or lithium-boride generally leads to a strong degradation of c-BN. In this case there is no uniform reaction with c-BN, and different phases occur. The reductive dissolution of c-BN first led to the formation of boron or boron-rich boron nitride B50N2, which further reacts with excess of lithium to lithium-borides. It could be shown that the system Li3BN2/c-BN is interacting at elevated temperatures leading to a severe change in the morphology of the c-BN seed crystals. In all cases a formation of well faceted surfaces could be detected. From this system the growth of c-BN is most likely [34]. Actually a spontaneous nucleation of c-BN from a degradation of Li3BN2 will not occur due to the fact that the less dense phase will nucleate first 376 HM 101 R. Haubner et al.

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according to Ostwald-Volmer rule, which is a general thumb rule for the kinetic behavior of reactions [35]. Thus, seeding with c-BN can overcome this problem. Hence, this approach seems to be well suited to develop a low- pressure synthesis of c-BN.

2. Characterization of BN products Characterization of BN-phases by spectroscopic methods (e.g. IR and Raman) is difficult due to the complexity of BN-structures and atomic bonding conditions. For example the x-ray diffraction lines of c-BN correspond with those of Cu, Ni and many other cubic phases. Elemental composition must be known or measured to be sure that no other phases are present. Some data are summarized below.

Infrared spectroscopy (IR) or Fourier transform infrared spectroscopy (FTIR) are often used to characterize BN products. To identify c-BN the Transverse Optical mode (TO) at 1065 cm"1 and the Logitudinal Optical mode (LO) at 1340 cm"1 were described [36]. Investigating commercial c-BN only one IR-peak between 1050 and 1100 cm"1 is observed commonly. If pure BN mixtures with B:N ratio of 1:1 are analyzed it will be easy to distinguish between h-BN and c-BN. However, if the chemical composition of the sample is unknown many artifacts can occur and a clear statement will often not be possible. More complex is the situation for h-BN, t-BN and a-BN because all of them show peaks between 780 and 1370 cm"1, which makes it impossible to differ between these phases. An IR-spectrum of c-BN and the peak positions of other compounds are shown in Fig. 1.

Raman spectroscopy is also useful to differ between h-BN and c-BN. As described above impurities as well as non-stoichiometrical mixtures can result in miss-interpretations. In case of c-BN two characteristic peaks are observed; h-BN shows only one peak. The two peaks for c-BN are described by several authors as TO-mode at 1055-1057 cm"1 and as LO-mode at 1305-1306 cm"1 [37-39]. The crystaliinity seems to be important for the Raman peaks of c-BN because in case of very fine grained c-BN (mainly in PVD layers) these peaks were not observed [40]. For h-BN the typical peak is located at 1367 cm"1 [38]. R. Haubner et al. HM 101 377

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Raman spectrum of c-BN and the peak positions of other compounds are shown in Fig. 2.

600 800 1000 1200 1400 1600 wavenumber (cnr1)

Fig. 1: IR spectrum of a nano-cBN coating and peak positions of other compounds in the relevant region are marked.

0 C

1600 1400 1200 1000 800 wavenumber (cnr1) Fig. 2: Raman spectrum of a HP-HT c-BN powder. Peak positions of other compounds in the relevant region are marked. 378 HM 101 R. Haubner et al.

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X-ray diffraction is another possibility to distinguish between c-BN and h-BN. In that case again the elemental composition of the sample is important because many cubic substances pretend c-BN (e.g. Cu, Ni,...) (Fig. 3). The peak position and its intensity are mainly influenced by grain size, stress within the layers and impurities.

(111)

(200)

(220) c " (311) 'S , 1 Cu Ni c-BN 30 40 50 60 70 90 2G (calculated diagram, Cu Ka) Fig. 3: X-ray diffractograms of c-BN, copper and nickel. The peak positions are very similar.

Methods for measuring the composition of the layers as weil as the B : N ratio are important. Mainly SIMS and Auger spectroscopy are used for such characterizations but for this methods it is difficult to get high quality standards [41].

Summing up, it may be said that for the characterization of c-BN in unknown samples (mainly PVD and Plasma-CVD deposits), the results of only one analytical method is not adequate for a statement. Measurements of the elemental composition in combination with IR and/or Raman and/or X-ray diffraction are necessary to ensure that c-BN is present. R. Haubner et al. HM 101 379

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3. Low-pressure BN and BCN deposition For the deposition experiments Tris(dimetylamino)borane was used as single source precursor which was decomposed in a cold wall reactor.

3.1 Experimental procedure In a cold-wall reactor (Fig. 4) substrates were heated up by high-frequency [42]. Argon at normal pressure (1 bar) was used as protecting gas and for carrying the precursor. The carrier gas (regulated by a mass flow controller) is saturated with the Tris(dimetylamino)borane precursor [43] according to its vapor pressure [44] in the evaporator/cooler unit. The carrier gas transports the precursor to the hot substrate, where deposition occurs. WC-Co hardmetal plates containing 6 wt.% Co, Mo and Si were used as substrates. Gravimetric analysis of the deposited layers reveals deposition rates ([mg/cm2*h]) and an estimation for the layer thickness (assumption: density of the obtained layer is similar to h-BN and graphite: 2.2 g/cm3).

Sample

C Fig. 4: Apparatus for deposition of BCN-layers by inductive heating of the samples TC temperature control, C gas cleaning R rotation pump (needed for evacuation/flushing) 380 HM 101 R. Haubner et al.

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The morphologies of the BCN-layers were determined by Secondary Electron Microscopy (SEM JEOL 6400). Chemical composition was measured by EPMA (electron probe micro analysis) integrated to the SEM. Such microprobe measurements give only uncertain hints about the composition because the elements B, C, N and O can not be detected by the system. Micro-Raman-spectroscopy (ISA Ramanor U1000, 488 nm Ar-ion Laser with 100 mW power, back-scattering geometry) was carried out to determine the film quality and phase composition. Also infrared-spectroscopy (Biorad FTS 175 with microscope) was used to characterize the deposits obtained. X-ray diffraction (CuKa radiation) was used to investigate the phase composition of the layers and products formed be reactions with the substrates. SEM, Raman and XRD examinations were normally carried out in the center of the samples.

3.2 BN layers deposited from Tris(dimethylamino)borane Deposition experiments were carried out in a wide parameter range [45,46], but only in a very small region XRD and IR peaks of c-BN could be observed. Typical parameters for such deposits are: hardmetal substrates 750 - 800°C, 100 - 400 seem Ar, 2.0*10"2 mmol/min Tris(dimetylamino)borane flux, normal pressure 6 - 8 h deposition. The layer growth rate is mostly very low and therefore the layer thickness was below 1 urn. Some SEM pictures are shown in Fig. 5.

•'•,'.. -** J*£fV* ft • « ' , ' - *

100 urn M 10 urn Fig. 5: BN deposits on hardmetal substrates (deposition time 7 h, 2.0*10"2 mmol/min precursor, WC-Co substrates, 100 seem Ar/min) R. Haubner et al. HM 101 381^

15" International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The IR-spectrum of such layers clearly show the peak for c-BN at 1088.3 cm"1 and additionally a peak at 1551.5 cm'1 is observed corresponding to sp2 carbon. A peak for h-BN at 1380 cm"1 was not observed but a small peak at 814.35 cm"1 which can be caused by the two h-BN peaks 780 cm"1 and 828 cm1. X-ray diffraction of such layers show a clear peak at 29 = 43.2° corresponding with c-BN (111). Also peaks of WC and Co could be observed and a small peak for h-BN (29 = 40.7°) (Fig. 6). The elemental composition measured by SIMS and Auger-spectroscopy shows a B : N ratio of about 1 : 1 and additional carbon. Other impurities like Co and oxygen are not observed in the layers. TEM investigations were also carried out but no clear reflexes of c-BN could be observed by electron diffraction [47].

4. Conclusions In the field of c-BN deposition one goal might be the coating of tools, which could be used for the machining of iron based materials. Such coatings offer the possibility to use a superhard material instead of the conventional carbide and oxide layers. From the present point of few the outlook for industrial applications of nano- cBN coatings is quite good, but some problems with the deposition process and the substrate materials still have to be solved.

4.1.1 Properties and applications of nano-cBN coatings Matsumoto and co-workers showed that the deposition of thick nano-cBN layers is possible and the adhesion on Si substrates is acceptable [32,33]. The one major question is whether this process is also suitable for hardmetal (WC-Co) substrates or not and if there will be any problems with the Co binder phase and the BF3 in the gas phase? 382 HM 101 R. Haubner et al.

15'° International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

IR-spectrum 2,5 T c-BN C (sp*) (1065), (1530-1590) , 1088,3 2 .. ^551,5 h-BN (780)(828)

1 ..

0,5 -.

600 800 1000 1200 1400 1600 1800 2000 wavenumber crrv1

XRD-diagram WC

I I 45 40 35 50 26 CuKa Fig. 6: IR spectrum and XRD diagram from layers deposited by decomposition of Tris(dimetylamino)borane

The stress within the c-BN layers which is caused by the high energetic ions is a problem, because many of the nano-cBN coatings delaminate immediately after or during deposition. Several investigations show ways to reduce the stress within the layers (e.g. buffer layers [25,48], regulation of the ion energy [27], or ion-induced stress relaxation [49]). R. Haubner et al. HM 101 383_

15™ International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

4.1.2 A CVD process for c-BN deposition? The present results are promising for the future that c-BN deposition is possible by a simple CVD process. From the parameters which were found further experiments are necessary to verify and optimize c-BN growth. It would be too early to talk about the industrial realization of such a CVD process and its applications.

ACKNOWLEDGMENT The authors gratefully acknowledge financial support of the present work by the Austrian Science Foundation (FWF) (Project P12055-CHE). This research project was carried out under the auspices of the trinational "D-A-CH" German, Austrian and Swiss cooperation on the "Synthesis of Superhard Materials". Parts of the present work were supported by NEDO (New Energy and Industrial Technology Development Organization) and JFCC (Japan Fine Ceramic Center) as a part of the FCT (Frontier Carbon Technology) project promoted by AIST, MITI; Japan.

References

[I] R.H. Wentorf Jr.: J. Chem.Phys. 26 (1957) 956 [2] R.H. Wentorf: J. Phys. Chem. 63 (1959) 1934-1941 [3] F.R. Corrigan, F.P. Bundy J. Chem. Phys.: 63 (1975) 3812-3820 [4] G. Will, G. Nover, S. von der Gönna: J. Solid State Chem. 154 (2000) 280-285 [5] V.L. Solozhenko: High Pressure Res. 13 (1995) 199 [6] O. Fukunaga: Diamond and Related Materials 9 (2000) 7-12 [7] H. Sachdev, R. Haubner, H. Nöth, B. Lux: Diamond and Related Materials 6 (1997) 286-292 [8] R. Lorenz, J. Woolcock: Z.Anorg.Chemie 176 (1928) 283-304 [9] J. Campbell: J. Electrochem. Soc. 96 (1949) 318-333 [10] J. von der Gönna, HJ. Meurer, G. Nover, T. Peun, D. Schönbohm, G. Will: Mater. Lett. 33 (1998) 321-326 [II] W. Kulisch, S. Reinke: Diamond Films and Technology 7 (1997) 105-138 384 HM 101 R. Haubner et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

[12] H. Hofsäss, C. Ronning, U. Griesmeier, M. Gross, S. Reinke, M. Kuhr, J. Zweck, R. Fischer: Nuclear Instruments and Methods in Physics Research B 106 (1995) 153-158 [13] H. Saitoh, W.A. Yarbrough: Diamond and Related Materials 1 (1992) 137-146 [14] K.H. Seidel, K. Reichelt, W. Schaal, H. Dimigen: Thin Solid Films 151 (1987)243-249 [15] G. Reisse, S. Weissmantel: Thin Solid Films 355-356 (1999) 105-111 [16] S. Reinke, M. Kuhr, W. Kulisch: Diamond and Related Materials 3 (1994) 341-345 [17] J. Widany, F. Weich, Th. Köhler, D. Porezag, Th. Frauenheim: Diamond and Related Materials 5 (1996) 1031-1041 [18] D.V. Shtansky, O. Tsuda, Y. Ikuhara, T. Yoshida: Acta mater. 48: (2000) 3745-3759 [19] D.R. McKenzie, W.D. McFall, W.G. Sainty, CA. Davis, R.E. Collins: Diamond and Related Materials 2 (1993) 970-976 [20] D.R. McKenzie: J. Vac. Sei. Technol. B11 (1993) 1928-1935 [21] W. Dworschak, K. Jung, H. Ehrhardt: Thin Solid Films 254 (1995) 65-74 [22] S. Uhlmann, T. Frauenheim, U. Stephan: Phys.Rev. B51 (1995)4541 [23] S. Reinke, M. Kuhr, W. Kulisch, R. Kassing: Diamond and Related Materials 4 (1995) 272 [24] D.J. Kester, R. Messier: J. Appl. Phys. 72 (1992) 504 [25] M. Okamoto, H. Yokoyama, Y. Osaka: Jpn. J. Appl. Phys. 29 (1990) 930-933 [26] A. Weber, U. Bringmann, R. Nikulski, C.P. Klages: Diamond and Related Materials 2 (1993) 201 [27] W. Kulisch, R. Freudenstein, A. Klett, M.F. Plass: Thin Solid Films 377- 378 (2000) 170-176 [28] A. Chayahara, H. Yokoyama, T. Imura, Y. Osaka: Jpn. J. Appl. Phys. 26 (1987) L1435-L1436 [29] S. Matsumoto, N. Nishida, K. Akashi, K. Sugai: J. Materials Science 31 (1996)713-720 [30] W. Kalss: Doctoral Thesis (1998) Vienna University of Technology [31] H. Sachdev, M. Strauss: Diamond and Related Materials 9 (2000) 614-619 [32] W.J. Zhang, S. Matsumoto: Chemical Physics Letters 330 (2000) 243-248 R. Haubner et al. HM 101 385^

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[33] W.J. Zhang, S. Matsumoto: Appl. Phys. A 71 (2000) 469-472 [34] H. Sachdev, X. Tang, R. Haubner, B. Lux: Refractory Metals & Hard Materials 16 (1998) 243-251 [35] A.F. Hollemann, E. Wiberg: "Lehrbuch der anorganischen Chemie" Walterde Gruyter& Co. (1995), Berlin, 105th edn. [36] P.J. Gielisse, S.S. Mitra, J.J.N. Plendl, R.D. Griffis, L.C. Mansur, R. Marshall, E.A. Pascoe: Phys. Rev. 155 (1967) 1039-1046 [37] O. Brafman, G. Lengyel, S.S. Mitra, P.J. Gielisse, J.N. Plendl, L.C. Mansur: Solid State Commun. 6 (1968) 523-526 [38] P.V. Huong: Diamond and Related Materials 1 (1991) 33-41 [39] H. Herchen, M.A. Cappelli : Phys. Rev. B47 (1993) 14193-14199 [40] T. Ikeda, Y. Kawate, Y. Hirai: J. Vac. Sei. Technol. A8 (1990) 3168-3174 [41] T. Kolber, K. Piplits, K. Nowikow, X. Tang, R. Haubner H. Hutter: Thin Solid Films (in press) [42] X. Tang, R. Haubner, B. Lux, A. Zechmann, E.Hengge: Journal de physique IV Colloque C5 (1995),Vol.5, 777-784 [43] S. Bohr, Doctoral Thesis, TU Vienna, 1995 [44] Gmelins Handbuch der anorganischen Chemie, B, Borverbindungen, Springer Verlag, 22 (1975), 4-40 [45] K. Aigner, X. Tang, R. Haubner, B. Lux: Proc. of ADC/FCT CONFERENCE 1999, Editors: M. Yoshikwa, Y. Koga, Y. Tzeng, C.-P. Klages, K. Miyoshi, pp. 301-306, Aug. 31- Sep. 3, 1999, AIST-Tsukuba Research Center, TSUKUBA, JAPAN [46] K. Aigner, X. Tang, R. Haubner, B. Lux: New Diamond and Frontier Carbon Technology 10 (4) (2000) 181-189 [47] P. Pongratz (TU-Vienna), S. Weissmantel (Hochschule Mittweida) private communication [48] Y.K. Yap, T. Aoyama, Y. Wada, M. Yoshimura, Y. Mori, T. Sasaki Diamond and Related Materials 9 (2000) 592-595 [49] H.-G. Boyen, P. Widmayer, D. Schwertberger, N. Deyneka, P. Ziemann: Applied Physics Letters 76 (2000) 709-711 386 HM 46 G. Korb et al.

15'"1 International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. WilrJner, Piansee Holding AG, Reutte (2001), Vol. 4

Gradient structures in SiAION's for improved cutting performance

Georg. Korb*, Friedrich Ivo Bulic*, Pavol Sajgalik** and Zoltan Lences**

*) Austrian Research Centers Seibersdorf, A-2444, Seibersdorf,Austria **) Institute of Inorganic Chemistry, SAS (Slovak Academy of Sciences); SK-84236 Bratislava, Slovakia

Summary

Typically at the present time used SiAION's for cutting purposes consist of ß-type, which has a hardness of about 1500 HV and a significant toughness. The authors have made available a type of SiAION which has a relatively tough core of a'/ß'-SiAION with a rim of nearly complete a'-SiAION. The hardness and microstructure have been illustrated. A true gradient in concentrations and microstructure can be observed from the inner core continuously to the outer region. It could be imagined, that an insert with such a structure together with the more tough core can be used as a cutting tool with high performance.

Keywords

Silicon Nitride, SiAION, Ceramic, cutting tools, cutting FGM, Functionally Graded Materials ,wear resistance, inserts, Si3N4, G. Korb et al. HM 46 387

15:h International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

1 Introduction

The well known types of Si3N4 based ceramic cutting materials are based on a composition of Si3N4, AIN, AI2O3, Y2O3 in a complex inner Silicon Nitride- Oxide- called "SiAION "

Typical SiAION's have a composition of about

Si3N4 to 83,77 wt% AIN to 5,54 wt% AI2O3 to 8,79 wt% Y2O3 to 1,09 wt%

RNrSAIN

glass SUä 2'FI SIM liH^H roolon A!N flJK

ct'-SiAION plsfie

The system between the constituents is given in Fig. 1

Fig. 1 Janecke prism of oxynitride systems of silicon, aluminum and R

(rare-earth cation) a'-SiAION plane lies on an inclined plane spanned

by Si3N4-RN:3AIN line and the Si3N4-AIN:AI2O3 line

At present commercially used cutting tools consist of ß-SiAION, this type does have sufficient toughness for machining purposes. The a-SiAION is much 388 HM 46 G. Korb et al.

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more brittle but has an outstanding hardness about HV 2000; and could be expected to keep sharp edges for a longer period of time and to have lower wear. At present, generally nitrogen ceramics (or called nitrides) have generated great interest as a group of attractive new ceramic materials due to their unique combinations of higher fracture toughness and better thermal shock resistance compared with the other structural ceramics [2]. Since Si3N4 does not occur naturally, it has been synthesized by several different processes. The major technical barrier for structural application was the densification of Si3N4: it is virtually impossible to fabricate it into a fully dense body by classical processing technology, e.g. solid state sintering, because the volume diffusivity in Si3N4 is not large enough to offset the densification retardation effects of surface diffusion and volatilization phenomena. The covalent bond is believed to be the reason for such a low volume diffusivity. To allow easier sintering during hot-pressing and to tailor the product properties to use, additives have been introduced into Si3N4 powder.

1.1 Silicon Nitride [2] [3] Silicon nitride exists in two phases a and ß both of which have hexagonal crystal structures. The a-phase has a unit cell approximately twice as large as that of the ß-phase. The dense body of ß-Si3N4 is e.g. fabricated by hot-pressing a-Si3N4 powder with a suitable densification additive (e.g. Y2O3, MgO, or AI2O3) at 30 MPa in a graphite mold at 1700 - 1800°C. The Si3N4 powder is generally covered with a thin layer of SiO2. The purpose of the additives is to react with this SiO2 and a small amount of Si3N4 at the high hot-pressing temperature to form oxynitride liquid in which a-Si3N4 dissolves and from which ß-Si3N4 is precipitated. Si3N4 based ceramics are densified by liquid phase sintering and consist of two phases such as crystalline Si3N4and an intergranular bonding phase. The intergranular phase is a glass or partially devitrified glass phase based on SiO2 and sintering aids - However, the intergranular glassy phases in Si3N4 tends to soften at elevated temperature and thus leads to reduced strength. Several commercial products in various compositions have been introduced. For instance, to improve hardness and abrasive wear resistance, a composite consists of a matrix phase of 70% (92% Si3N4. 6% Y2O3, and 2% AI2O3) and a dispersed phase of 30% TiC, orTiN. Another composition consists of Si3N4 with additions of AI2O3 and Y2O3 but, do not contain TiC orTiN. G. Korb et al. HM 46 389

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1.2 SiAlON [2] [3] In the early 1970's it, was found that aluminum and oxygen could be substituted for silicon and nitrogen, respectively, in the Si3N4 crystal structure to form an Si(6-z)Alz0zN(8.2) solid solution -, usually referred to by the acronym "SiAlON". To produce solid solution ß-SiAION, a mixture of ß-Si3N4(77%). AI2O3 (13%). Y2O3(10%), and Metal Nitride is used as the starting material. The powder mixture is preformed by cold isostatic pressing, and subsequently sintered at 1800°C for 1 hour. During sintering, the mixture for ß-SiAION produces a larger volume of lower-viscosity liquid than in the synthesis of Y2O3-stabilized Si3N4- Therefore ß-SiALON can be fully densified by pressureless sintering. Rapid cooling from processing temperature produces a microstructure of ß-SiALON grains with an intergranular glassy phase. If Y2O3 is the sintering aid, a portion of this glassy phase can be converted to crystalline yttrium-aluminum-garnet (YAG) by heat treatment or slow cooling. However, most sintered SiAlON s contain some residual glass phase, particularly at grain-boundary triple points. Like the Si3N4 sintering aid systems, the properties of SiAION's are dependent on the type and amount of sintering aids employed the processing route followed during part fabrication. In general, silicon nitride cutting tools have better thermal shock resistance but lower chemical inertness than alumina tools have. To improve the chemical stability AI2O3 can be added in ß-SiAION.

To improve fracture toughness ZrO2 can be added in ß-SiALON, however, the addition of only ZrO2 causes an undesirable chemical reaction without densification. Therefore, AI2O3 and AIN are further added to produce a dense material consisting of ß-SiAION and ZrO2 [2].

1.3 Graded SiAION's [1]

First attempts have been made to achieve gradients in 1996/97 by Limin Chen

[1] where the outer layers consist of a-SiAION . This was achieved by a special

"powder bed" sintering process. 390 HM 46 G. Korb et al.

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2 Experimental As a starting point a designed a + ß composition has been prepared in quadratic shaped specimens. Atypical starting composition was e.g.

Si3N4 AI2O3 AIN Y2O3

79,02 wt% 8,00 wt% 8,24 wt% 4,25 wt%

The material had a porosity of about 50%. By a special chemical method (at present proprietary) the continuos changing phase from a'/ß' (in the center) to a' in the rim has been introduced.

Fig. 2. piece of graded SiAION after grinding off most outside layers (barrier, impurities) roughness etc.

After full densification a density of 3,2 g/cm3 has been achieved.

After grinding off the most outside impurities and barrier layer, a typical size which is useful for cutting inserts has been obtained. Fig. 2 (12 mm by 12 mm by 4 mm). G. Korb et al. HM 46 391

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3 Investigations and results

3.1 Microstructural Investigations The specimens have been partially been broken and also cutted and grinded through the center to make visible the core-rim structure.

Fig. 2, Fig. 3 and Fig. 4 are demonstrating the gradient structure achieved. The different "grey"-graduation is clearly visible. The high resolution in Fig. 5 is demonstrating the typical microstructure of the SiAION, a free Si3N4 -SiAION matrix material with glassy phase.

Fig. 3 Gradient structure from center to the outside of the specimen

Fig. 4 Different microstructures in the FGM material, higher magnification 392 HM 46 G. Korb et al.

15° International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Fig. 5 Microstructure of the rim and core (magnification: 1000x) 3.2 Chemical composition A typical line scan for Yttrium in the SEM shows a gradient from the inner to the outer regions which is typically for the compositional change in such FGM materials (Fig. 6)

--. ,-*

i= 16,48 GPa != 20,94 GPa

Fig. 6 EDS-line scan for Yttrium on the FGM sample plus different hardness

indentation (HV1) G. Korb et al. HM 46 393

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3.3 Hardness Hardness as HVi (1 kg indentation load) from the center to the outer region step by step is given in Fig.7 that means an outer hardness of about 2000 HV was achieved in comparison to 1600 in the center.

21 FGM-4C

s 20 Eä m a 19 a E3 a a B B 18 B B

i.i. I.I. I i 1 2 3 distance / mm

Fig. 7Hardness HV-, of the graded a'/ß' SiAION vs. distance from the surface 394 HM 46 G. Korb et al.

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3.4 Crystallographic Structures XRD patters of the surface gave the confirmed result, that the outside region of the specimen consist of mainly a'-SiAION. The X-ray pattern is given in Fig. 8 where clearly both phases can be identified but of course the a is much stronger. A rough but specific calculation out of the peak intensities gives a ratio of a'/(a'+ß') of about 87% for the a-modification on the surface.

i?nn FGM-4C "innrr o a-Si3N4 °- ß-Si N 800 3 4

cps

400 O O AO ?on o ?0 40 50 70 2 0

Fig. 8 X-Ray pattern of outer zone of FGM SiAION G. Korb et at. tM46 395^

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4 Discussion

To achieve substantial change in microstructure and achieve sufficient toughness in a SiAION-structure one has two possibilities: first to change the outer region to a'-SiAION and to keep ß or (a + ß) core or second to change the a-structure in that way that elongated grains can cause the desired thougness property, (see also contribution of Mr. Li, Jang Tao, this seminar, paper No. HM 48) [4].

Unfortunately the wear at cutting operations does not consist of one type of wear. Resistivity against e.g. pure mechanical wear can be influenced by increase of hardness but chemical wear, which is a type, when e.g. chemical diffusion from chip to insert or vice versa takes place, can only be influenced by changing the chemical composition of the cutting insert. Both possibilities are feasible with the SiAION but a wide field of scientific work and a great challenge is given. 396HM 46

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Literature

[1] Limin, Ch., G. Groboth and Kny E.: "Newly developed ceramic cutting tool materials with functionally graded microstructures", Proceedings of the 5th European Conference and Advanced Materials, Processes and Applications, Materials, Functionality & Design, Vol. 2, Polymer and ceramics, EUROMAT 97, Maastrich, Netherlands (1997)

[2] Sukyoung;K.: "Advanced Ceramics as Cutting Tool Materials, Part II- State of the Art and Developments", Journal of the Canadian Ceramic Society Vol.61, No.1. Feb.1992

[3] Ekström, T: and Nygren, M.: "SiAION Ceramics";J.Am.Ceram.Soc.75 [2] 259-76(1992)

[4] J. T. Li, B. Djuricic, G. Korb, and C.C. Ge:

"a-SiAION — new opportunities via new microstructure" this seminar, HM

Acknowledgement: The work for this publication was supported by the European Union , Fifth Framework Program Contract Nr: G5RD-CT2000 00221 AT0200488 N. Frage et al. HM 49 397

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The Effect of Cr2O3 Additions on the Pressureless Sintering of B4C.

N. Frage, M. Bizaoui and M. P. Dariel

Department of Materials Engineering, Ben-Gurion University of the Negev, Beer-Sheva, Israel

Summary: The effect of Cr2O3 and CrB2 additions on the sintering behaviour of B4C was studied in the 2050 to 2130°C temperature range. According to thermodynamic predictions, chromium oxide is reduced by carbon from the B4C phase, thereby generating carbon deficient carbide. CrB2 is formed in situ in the course of the interaction between B4C and Cr2O3. By sintering at 2130°C for 1h, a mixture of B4C and 15wt% Cr2O3 is transformed into a 94% dense two-phase composite that consists of sub-stoichiometric boron carbide and CrB2 and displays a bending strength of 170 MPa. The significantly improved sintering behaviour of this composite ceramic is attributed to the enhanced mass transport processes that take place in sub-stoichiometric B4C. The relatively low value of the bending strength reflects, probably, the influence of the cracks generated as a result of thermal stresses that arise at the CrB2/B4C interface during the cooling of the sintered samples.

Keywords: Boron carbide, chromium diboride, reactive sintering

1. Introduction: The well-documented, outstanding properties of B4C make it a valuable potential material for a variety of applications (1,2). The realization of this potential is hindered, however, by two major drawbacks, namely, the low fracture toughness of B4C and the very high temperature required for its sintering. Nearly full density cannot be achieved by pressureless sintering and can be attained in pure B4C only by hot pressing above 2300°C. Numerous proposals (1-20), which were discussed in a previous paper (21), have been put forward in order to lower the sintering temperature of boron carbide.

The present study was undertaken with the purpose of gaining additional insight into the mechanism whereby Cr2O3 additions improve the sintering behavior of B4C. 398 HM 49 N. Frage et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Boron carbide exists over a wide composition domain that extends from stoichiometric B4C (20 at% C) to a sub-stoichiometric compound with 9 at%C. It was of interest to find out whether, in the absence of an extraneous carbon addition, chromium oxide could be reduced by carbon that originated from boron carbide, thereby altering the composition of the latter. A specific objective of this study was to determine the contribution of the in situ developed deviations from stoichiometry on the sinterability of B4C. In order to provide an appropriate frame of reference, a comparative study was carried out on two sets of samples, namely boron carbide with: (a) additions of Cr2O3, leading to the formation of TiB2 simultaneously with depletion of the carbon content in the carbide phase, (b) additions of CrB2. An analysis of the deviations from stoichiometry due to the reduction of oxides was suggested in (21). The first part of this paper describes the thermodynamic analysis. The results of the analysis led to the experimental study described subsequently.

2. Thermodynamic Considerations A. The effect of Cr2O3 additions on the final composition of the sintered samples. The interaction between B4C and Cr2O3 additives at high temperature leads to the partial decomposition of B4C and to the formation of new phases according to reactions:

2+x 2+x

-D4C H Lr2Ui = Crx>2 H LC>2 H -°4M-x f2)

The relative amounts of B4C and CrB2 in the final product and the composition of the boron carbide depend obviously on the amount of Cr2O3 that had been added and on the importance of reaction (1) as compared to reaction (2). The calculations are straightforward and are detailed in (21).

The weight losses (per 100g of mixture) Am(l)and Am(2) in the course of the interaction according to reactions (1) and (2), respectively, may be expressed as: Am(l) = 0.553 • (Cr2O3,%) and Am(2) = 0.434 • (Cr2O3,%).

The CrB2 content in the final mixtures, calculated according to the stoichiometric coefficients for reactions (1) and (2), are shown in Fig.1. N. Frage et al. HM 49 399

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

40 for reaction (I).

0.969 •(Cr2O3,%) CrB2,Vo = 100-0.553-(O-203,%)

20- for reaction (2)

S 0.969.(0-A,Q/o) 2 10- 100-0.434-(O-203,%)

10 15 20 25 30 35

Cr2O3,wt.%

Fig.1. The CrB2 content in the final mixtures as function of the amount of added Cr2O3.

The carbon content in the residual boron carbide phase, as a function of the amount of Cr2O3 that had been added, is shown in Fig.2. The two curves stand for the two limiting cases, corresponding to reactions (1) and (2), respectively.

Cr2O3, %

Fig.2. The dependence of the carbon content in the boron carbide phase on the amount of Cr2O3 that had been added. 400 HM 49 N. Frage et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

According to Fig.2, up to 41.0 wt.%, Cr2O3, if eq.(1) is valid, or 59.0 wt.%, Cr2O3p if eq.(2) is valid, can be added without any free boron being formed.

B. The effect of the carbon content of the boron carbide phase on the equilibrium parameters For a thermodynamic analysis of Cr2O3 being reduced by carbon that originated in non-stoichiometric boron carbide, it is useful to treat the latter phase as a solid solution of B and C with varying composition. For the simple case, when only CO gas is formed during the interaction, the equilibrium in the Cr-C-B-0 system can be described by the chemical reaction:

4[B] + 3[C] + Cr2O, = 2CrB2 + 3CO (3) where the square brackets denote that boron and carbon are in solid solution B-C within the composition limits, corresponding to the stability range of the boron carbide phase.

The equilibrium constants for reaction (3) can be written as

AVV = 43 (4) aBac where at is the activity of the components in boron carbide, Pco is the partial pressures of the components of the gaseous phase (atm).

According to equations (4), at a given temperature, the composition of the boron carbide determines the composition of the gaseous phase. In order to determine the equilibrium pressure, values of the activities and the temperature dependence of the equilibrium constants K(3) must be known.

The temperature dependence of AG°(3) was derived from the thermodynamic data [22, 23] of the following reactions:

C + O2 = CO2 AG° = -394762 - 0.8367

-O2 + 2Cr = Cr2O3 AG° = -1120266 + 259.8267

CO2+C = 2CO AG° =+169008-172.19T

Cr + 2B = CrB2 AG° = -134122 +14.237 N. Frage et al. HM 49 401

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

For reaction (3), we have: ÄG°(3) = 513391-4888.397 and

+ 58.744 (5) For the determination of the thermodynamic properties of boron carbide within its range of the stability, the only available results are those of Froment et. al., (24,25). These results were used to derive the composition dependence of the activities, as shown in eqs. (6,7). The details of this derivation are given in ref. (21),

53524 2 In aB = In xB - 'c c -0.0081) +0.0943 (6) T 53524 2 In ac - In xc - [(l-xc) - 0.64] + 1.609 (7) T By combining equations (4-7), it became possible to calculate the equilibrium parameters for the reduction of Cr2O3 by carbon originating from the boron carbide phase. The equilibrium compositions of the gaseous phase as function of carbon content in the carbide phase at various temperatures are shown in Fig.3.

As can be seen from Fig.2, the CO partial pressure in the gaseous phase is high enough to ensure reducing of chromium oxide by the carbon that originated from the boron carbide phase.

3.5 ' 1700K • 1300K A 1 QnnTif 1 yuuiv 2.0 1.5 1.0 0.5 o.o 0.10 0.12 0.14 0.16 0.18 0.20 Carbon content, at. fraction

Fig. 3. The equilibrium partial pressure at various temperatures of CO for the reduction of Cr2O3 by carbon originating from B4C. 402 HM 49 N. Frage et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3. Experimental Procedures Boron carbide (Fig. 4) of about 5.9 fim mean particle size (a carbon-to-boron atomic ratio of 3.94), Cr2O3 (2 jam and less particle size) and CrB2 powder of about 3 jam average particle size were used to prepare ceramic B4C samples containing 5-25 wt.% Cr2O3 or CrB2additions.

Powder compacted powder mixtures were heated to 1300-1500°C for two hours in a vacuum furnace to ensure the completion of the reaction between the B4C and the additives. Sintering was carried out for 1h, in a "ASTRO" furnace with graphite heat elements under an argon atmosphere in the 2000- 2130°C temperature range. The density of the sintered samples was measured by the liquid displacement method in distilled water. The relative densities were calculated using the theoretical density values of the mixtures. Flexural strength was measured by the three-point method with a fulcrum spacing of 9 mm in an "lnstron-1186" instrument, at a constant crosshead speed of 0.5 mm/min. For each composition and treatment temperature, five or six identical samples were tested. The reported data points represent the average of the measured values and the error bars stand for the standard deviation.

0 5 10 Particle size,|am Fig.4. Particle size distribution of the boron carbide powder. N. Frage et al. HM 49 403^

15"1 International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

X-ray diffraction ("Rigaku", RINT-2100, Cu-K, line) and diffraction line fitting were applied to identify the phases present in the sample and determine the lattice parameters of the boron carbide phase.

The weight-loss of samples that underwent high temperature anneals was determined at room temperature. Samples that contained Cr2O3 additives lost weight due to the reaction between B4C and the additive and formation of volatile, CO and CO2, reaction products. Even for B4C samples with no Cr2O3 additive, a loss of approximately 4 wt.% was observed and attributed to the evaporation of B2O3 on the surface of the carbide particles or/and to its reaction with some excess carbon present in the carbide. The experimental results regarding to the weight losses were corrected accordingly.

4. Results and Discussion: A. Phase composition The weight losses incurred during the heat treatment of mixtures that had different Cr2O3 content are shown in Fig. 5. The results fall within the range defined by the two calculated curves that correspond to the presence of either CO or CO2 in the gaseous phase. In Fig.3, the 4 % weight-loss of blank boron carbide was taken into account. The composition of the gaseous phase cannot be defined unambiguously. It varies as a function of the time elapsed; it also varies as a function of the location within the sample and is probably dominated by kinetic factors in addition to the thermodynamic factors.

X-ray diffraction patterns of sintered B4C with Cr2O3 additions are in agreement with thermodynamic predictions and the only phases detected were boron carbide and chromium diboride. The measured lattice parameters of the B4C compound provide additional support for the thermodynamic prediction. The dependence of the lattice parameters of boron carbide on the additive fraction of Cr2O3 is shown in Fig.6 (a and b). The additions of Cr2O3 increase both the a and c lattice parameters of B4C. According to Aselage and Emin (26), the lattice parameters of B4C, increase with decreasing carbon content in the carbide. Thus, the formation of sub-stoichiometric boron carbide that was predicted by the thermodynamic analysis was actually observed experimentally. B. Densification No noticeable densification was observed for treatments carried out up to 2000°C. In the 2050-2130°C temperature range, significant densification 404 HM 49 N. Frage et al.

15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

occurred with a strong dependence on the composition of the mixture. The results are shown in Fig. 7 for two sintering temperatures. Both, the Cr2O3 and the CrB2 additives improve the densification. Chromium oxide, however, is significantly more effective than CrB2 in improving the densification of the boron carbide powder. This feature is more remarkable at 2050°C than at 2130°C and may be due to the dissolution of carbon in the nonstoichiometric carbide and the change of the boron carbide stoichiometry during sintering at high temperature in a furnace with graphite heat elements. In principle, a sub-stoichiometric compound with a significantly higher structural defect concentration may display improved mobility of its constituents, leading to higher diffusivities and, thence, to higher sinterability. A multi-step dynamic mass transport process in which the evaporation- condensation mechanism plays a major role can also be considered. Over the width of the carbide phase, the activity of boron increases by an order of magnitude to reach its highest value in equilibrium with pure boron at the boron-rich boundary. Thus the rate of evaporation increases with decreasing carbon content in the carbide phase. Condensation of the boron vapour onto neighbouring boron carbide particles gives rise to carbon concentration gradients and associated carbon atom fluxes. The resulting chemical diffusion may lead to the enhanced densification that was observed.

Cr2O3, %

Fig. 5. Weight loss of mixtures of various Cr2O3 content. The full lines standfor the calculated values, the open circles for the experimental data. N. Frage et al. HM 49 405

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The optimal amount of additives was 15% Cr2O3. The relative density of such mixture reaches 94-95% of TD after sintering at 2130°C.

5.620

5.600-^ 5 10 15 20 25 Cr2O3, %

12.10-

12.08-

12.06-

20 ' 25 Cr2O3, %

Fig.6. Cell dimensions of B4C after its interaction with Cr2O3. 406 HM 49 N. Frage et al.

15!h International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

0 —o— Cr,O32130 C 90- y u —n— Cr,03 2050°C ' <. —t-

—Y—CrB22130°C

—•— CrB2 2050°C

/ Relativ e Density , %

O N - J o o / 10 15 20 25 Additives, %

Fig. 7. Relative density of the B4C samples with Cr2O3 and CrB2 additives, sintered at temperature range 2050-2130°C

C. Microstructure The optical micrographs of the sintered and polished ceramic samples containing of 15% Cr2O3, before and after electrolytic etching (1% KOH, 5V, 45s) are shown in Figs. 8 and 9. The bright particles in Fig.6 were identified as CrB2 particles, the dark spots are either residual pores or voids formed when CrB2 particles were torn out from the free surface, in the course of the specimen preparation. The relatively easy removal of the CrB2 particles may be due to the substantial residual thermal stresses that arise at the interface between the CrB2and B4C particles as a result of the significant difference in their respective thermal expansion coefficients. The thermal expansion coefficient (TEC) of the boron carbide phase is 4.5-10"6K"1, whereas that for 1 the CrB2 phase is significantly higher (about 11.58-10' K" ) (22). The level of the stresses at the interface, ar, can be estimated (eq.8) according to (27)

(8) 0.5(1

where E , a , and v are the elastic moduli, thermal expansion coefficients and Poisson coefficients, respectively, of the components. N. Frage et al. HM 49 407

15!h International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The Young modulus and the thermal expansion coefficient for B4C are 445GPa and 0.19, the respective values for CrB2 are 211 GPa and 0.25 (22). For a A71 value of 2000°, the calculated residual stress is about 3.3GPa. Such level of residual stresses may easily lead to crack formation during cooling of the samples after sintering.

Almost all CrB2 particles were removed after electrolytic etching (Fig.9), the grain boundaries between the B4C particles however, are clearly visible. The average size of the boron carbide grains, derived by image analysis, is about 40)am. According to these results, significant coarsening took place during sintering.

Fig.8. Microstructure of the sintered B4C with 15% Cr2O3 additive (x400)

Fig.9. Sintered boron carbide with 15% Cr2O3 additive after etching. 408 HM 49 N. Frage et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

D. Mechanical properties The measured bending strength of the sintered samples is shown in Fig. 10. The bending strength of the samples that were sintered at the lower temperature was about 50-80 MPa. As shown in Fig.8, the bending strength of samples that were obtained from a mixture of B4C with 15% chromium oxide reaches 170 MPa. This relatively low value reflects, probably, the influence of the weak CrB2/B4C interfaces. This feature is more remarkable for the samples with 25% additives and leads to a considerable decrease of the bending strength. T " Cr,O 150- 3 * CrB2

T T bß 120- c 1

C 1 _ ] | I 90- II

0 5 10 15 20 25 Additives, wt.% Fig. 10. Bending strength of the sintered (2130°C) boron carbide as a function of the amount of additives.

4. Summaries 1. Thermodynamic analysis of the B4C+Cr203 system and its experimental verification indicate that chromium oxide is reduced by carbon that originated in the carbide phase. The resulting product is a mixture of sub-stoichiometric boron carbide with in situ formed CrB2. 2. The addition of Cr2O3 and CrB2 improves the sintering behaviour of B4C. Thus the addition of 15 wt.% Cr2O3 to B4C powder allows reaching 94% relative density after sintering for 1 h at 2130 °C. This density level is higher than that attained with an equivalent amount of CrB2 addition. 3. The striking improvement of the sintering behaviour of B4C by Cr2O3 N. Frage et al. HM 49 409

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additions is attributed to the formation of sub-stoichiometric boron carbide. 4. Metallographic examination of the compact sintered at 2130°C revealed a two-phase microstructure with 40 ^m size B4C grains. The in situ formed CrB2 has a weak interface with B4C grains due to the difference between their thermal expansion coefficients. 5. The maximum bending strength was reached for samples with 15% of Cr2O3. The relatively low value of the bending strength (170MPa) and its considerable decrease for samples with 25% additive represent, probably, the influence of crack formation at the weak CrB2/B4C interfaces during cooling of the sintered samples. REFERENCES

1. R. Teile: Materials Science and Technology, Vol 11, pp. 175-258 (ed. R.W. Cahn, P. Hansen, and J. Kramer, VCH, New York 1994) 2. S. Dole, S. Prochazko, and R. Doremus: J. Am. Ceram. Soc. 72 (1989), pp. 958-966 3. M. A. Kuzenkova, P. S. Kislyi, B. L. Grabchuk, and N. I. Bondaruk: J. Less-Common Met. 67 (1979), pp. 217-223 4. J. D. Katz, R. D. Blake, J. J. Petrovich: Met. Powder Rep. 43 (1988), pp. 835-841 5. K. A. Schwetz and G. Vogt: Dtsch. Patent DE 2751998 C2, (1977) 6. J. V. Henney and J. W. S. Jones: Br. Patent Application 2014193 A, (1988) 7. H. Suzuki and T. Hase: in Proc. Conf. Factors in Densification of Oxide and Nonoxide Ceramics, pp.345-349 (ed. S. Somiya and S. Saito, Japan 1979) 8. Z. Zachariev and D. Radev: J. Mater. Sei. Lett. 7 (1988), pp. 695-698. 9. Z. Zachariev and D. Radev: Boron and Refractory Borides, V. I. pp. 464- 467 (ed. Matkovich, Springer-Verlag, Berlin 1977) 10. S. Prochazko, US Patent 4005235, (1977) 11. R. G. Lange, Z. A. Munir, and J. B. Holt: Mater. Sei. Res. Vol. 13 (1980), pp. 311-315 12. Y. Kanno, K. Kawase, and K. Nakano: J. Jap. Ceram. Soc. 95 (1987), pp. 1137-1141 13. J. Kriegesmann, Dtsch. Patent DE 3711871 C2, (1989) 14. B. L. Grabchuk and P. S. Kislyi: Poroshkovaya Metallurgiya 7 (1975), pp.151-155 15. M. Bougoin, F. Thevenot, F. Dubois, and G. Fantozzi: J. Less-Common Met. 114(1985), pp. 257-261 410HIVU9

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16. J. H. Oh, K.K. Orr, C.K. Lee, D.K. Kim, and J.K, Lee: J. Korean Ceram. Soc. 22(1985), pp. 60-65 17. K. Adlassnig: Planseeberrichte for Pulvermetallurgie, Vol 6 (1958), pp. 92-98 18. B. L. Grabchuk and P. S. Kislyi: Poroshkovaya Metallurgiya, Vo! 8 (1974), pp.11-16 19. V. Skorochod, JR. D. Vlajis, and V. D. Krstic: J. Mat. Sei. Let. 15 (1996), pp. 1337-1339 20. H-W Kim, Y-H Koh and H-E Kim: J. Mat. Res. 15 (2000), pp. 2431-2436 21. L.Levin, N.Frage, M.P. Dariel: Met. Trans. 30A (1999), pp. 3201-3210 22. T. Y. Kosolapova: in „Refractory Compounds, Properties, Production and Application" (Metallurgy, Moscow, 1986) 23. JANAF Thermochemical Tables, Handbook of Chemistry and Physics, NBS 37, Washington D.C. (1985) 24. K. Froment, C. Chatillon, and M. Colin: Rev. Int. Hautes Temper. Refract., France, 1991, vol. 27, pp. 141-158 25. K. Froment, C. Chatillon, J. Fouletier, and M. Fouletier: J. of Nucl. Mat. 188(1992), pp. 280-284 26. T. L. Aselage and D. Emin, in „Boron and Refractory Borides", AIP Conference Proceedings 231 (ed. D. Emin, T. L. Aselage, A. C. Switendick, E. B. Morison, and C. L. Beckel, American Institute of Physics, New York 1991, pp. 187-191) 27. L. S. Sigle: Acta Met. 44 (1996), pp. 3599-3609 AT0200489 B. Wollein et al. HM 56 411

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PHASE REACTIONS IN Ti(C,N)/(Ti,W)C, Ti(C,N)/(Ti,Mo)C, (Ti,W)(C,N)/Co AND (Ti,W)(C,N)/Ni DIFFUSION COUPLES

B. Wollein, W. Lengauer

Institute for Chemical Technology of Inorganic Materials, Vienna University of Technology, Getreidemarkt 9/161, A-1060 Vienna, Austria; www.tuwien.ac.at/physmet

Summary:

In order to investigate interactions of (Ti,W)(C,N) and (Ti,Mo)(C,N) with binder metals solid/solid diffusion couples were annealed. These two-dimensional arrangements provide good access to phase reactions occurring upon sintering already in the solid state. It was found in (Ti,W)(C,N)/Co- and (Ti,W)(C,N)/Ni-based couples that the reaction zone is thinner in contact with Co than with Ni. It was also observed that the reaction rate with both Co and Ni is lower if nitrogen is added to the hard phases. Beside a thickness variation of the diffusion zones a change in the microstructure was found. At the interface of nitrogen-free hard phases in contact with Co elongated microstructural constituents are formed with the main axis perpendicular to the interface, while at the interface of nitrogen-containing hard phases these elongated microstructural constituents were found in contact with Ni. Also phase reactions and the diffusion behaviour between the different hard phases were studied by means of solid/solid diffusion couples of the type Ti(C,N)/(Ti,W)C and Ti(C,N)/(Ti,Mo)C with various concentrations. After 1008h inside the Ti(C,N) phase of the Ti(C,N)/(Ti,W)C diffusion couple W-rich rims were found around Ti-rich cores. The thickness of these rims did not change with time. If W is exchanged by Mo it was observed that the well known core-rim structure is built by diffusion of Mo into the Ti(C,N) grains. But with time the Mo diffuses into the core of Ti(C,N) phase showing that the structure is not stable.

Keywords: hardmetals, binder metals, EPMA, X-ray mapping 412 HM 56 B. Wollein et al.

15": International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

1. Introduction:

The W-C-Co system is one of the most important material systems in the field of hardmetal cutting tools. The interactions in WC-Co or WC-Ni based hardmetals which often contain a number of other hard phases such as TaC, NbC, TiC and TiN, are difficult to investigate because of the small grain size (ca. 1um) of the hard particles embedded in the Co or Ni binder phase. In order to investigate the interactions of hard phases with the binder metal and also the reactions in between the hard phases diffusion couples were made. The diffusion couples should provide a better access to investigate phase reactions at the interfaces (1,2). The first step in our investigations reported here were diffusion couples based on (Ti,W)(C,N) hard phases contacted with different binder metals. To investigate the two main regimes the reactions were performed below and above the melting temperature of the binder metal. Beside the conventionally used binder metals Co and Ni, also mixtures of Fe/Ni/Co were used. Another part of this work consisted of solid/solid reactions in diffusion couples of the type Ti(C,N)/(Ti,W)C and Ti(C,N)/(Ti,Mo)C with various concentrations in order to observe the tendency of mutual solubility of these ternary phases. From contact reaction investigations (3) it is known that increasing nitrogen content influences the diffusion kinetics in such a way that faster diffusion rates are observed. In our work we investigated the change of diffusion kinetics if Mo is exchanged by W.

2. Experimental:

2.1 Sample preparation:

2.1.1 Hard phase/binder metal diffusion couples:

The different powders (WC, TiC, TiN, Mo2C, C, W) were mixed and hot pressed at ca. 2200°C and 620bar in Ar to get almost dense samples. The compositions were (vVo.48Tio.52)C, (Wo.42Tio.52)C and (Wo^TiojeXCo./No.a). The binder metals were Co and Ni sheets and pre-sintered powders of the composition Co/Ni 50/50 and Fe/Ni/Co 30/30/40. For the solid/solid diffusion experiments, both the hard phases and the binder metals, were lapped on a cast iron disk with B4C in order to get plane surfaces to guarantee good contact between the plates during annealing. The samples were annealed in an induction furnace in 200mbar Ar. Firstly, a Co/Ni diffusion couple was prepared by annealing a Co plate contacted to a B. Wollein et al. HM 56 413

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Ni plate at 1300°C for 40h to reach a gradient from 100% Co to 100% Ni. Afterwards the Co/Ni sample was cut, lapped and annealed in contact with a hard phase plate. A sketch of the arrangement is shown in Fig.1 and the annealing conditions are listed in Tab.1.

W weight ca. 2 kg

Fig.1 Arrangement of the contacted plates (solid/solid diffusion couples).

For the solid/liquid hard phase/binder metal diffusion couples the hard phases were hot pressed (see above), cut into small pieces and put into AI2O3 crucibles. Onto these pieces pre-alloyed binder metal powders were pressed. The crucibles were covered with an AI2O3 plate to avoid vaporisation of the binder metals and annealed in a Mo crucible, which served as susceptor, in an induction furnace at 800mbar Ar. The annealing temperatures were slightly above the melting temperature of the various alloys and metals (Tab A).

2.1.2 Hard phase/hard phase diffusion couples:

Onto hot-pressed (see above) Ti(C,N) samples mixed hard phase powders, different in composition, were pressed in Ar at 2200°C and ca. 620bar. The composition of the samples was Ti(Co.7No.3)/(Tio.7Wo.3)C and Ti(C0.8No.2)/(Tio.8Moo.2)C. These samples were annealed at 1bar Ar above ambient pressure at 1500°Cfor504, 1008 and 1512h. 414 HM 56 B. Wollein et al.

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Tab.1 Annealing conditions of the solid/solid and solid/liquid diffusion couples;

hard phase binder metal T[°C] t

(W0.48Tio.52)C Co 1230 69h

(WO.48Tio.52)C Ni 1230 69h

(Wo.24TiO.76)(Co.7NO.3) Co 1230 97h

(Wo.24Ti0.76)(Co.7No.3) Ni 1230 97h

(WO.48Tio.52)C Co/Ni 50/50 1445 5min

(W0.48Ti0.52)C Fe/Ni/Co 30/30/40 1460 5min

(W0.24Tio.76)(Co.7No.3) Co/Ni 50/50 1440 10min

(WO.24Tio.76)(Co.7No.3) Fe/Ni/Co 30/30/40 1475 10min

2.2 Metallography:

The samples were cut with a diamond disk, embedded in cold-setting resin (Araldite®, Bayer) and ground with a 20pm diamond disk. After this procedure a 3um diamond paste was applied on a hard cloth (Struers DP-NAP) for polishing.

2.3 Microprobe investigations:

The microprobe investigations were performed with a CAMECA SX 50 microprobe equipped with five spectrometers and a variety of crystals to make the characteristic radiation of a full range of elements accessible. Chemically analysed carbide and nitride standards were used.

2.3.1 Linescans:

A variety of different conditions was chosen for the beam current, voltage, type of crystal and background position to carry out the measurements (see Tab. 2). To avoid the problem of peak coincidences and reduce the problem of contamination during measurements, C was subtracted from all elements, N from Co, Ni and Ti (4). B. Wollein et al. HM 56 415

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Tab.2 EPMA conditions for the systems (Ti,W)(C,N) + Co and (Ti,W)(C,N) + Ni

Element0 Beam Crystala Peak +bgb -bgb conditions Position [105 sin 0] [105sin©] [kV/nA] [1O5sin0] W 10/100 TAP 27176 820 530 C 10/100 PC1 74125 7000 7000 Ti 10/100 PET 31424 800 600 Mo 10/100 PET 61787 600 1.2 Ni 10/100 TAP 56705 600 1.2 Co 10/100 TAP 62139 600 1.2 N 10/100 PC1 52998 5000 4500 PC1:59.8A, PET:8.742A, TAP: 25.9Ä b Background to be added to (+bg) or subtracted from (-bg) the peak position. c All the measurements were carried out with cold trap and without oxygen jet

2.3.2 X-ray mapping:

For X-ray mapping the beam was fixed and the sample holder moved across a defined area. This provides strictly focused conditions in the spectrometer. The scanning time per point and per element was 50ms. It should be noted that the X-ray mapping was not quantified, meaning that the same colour in different images is not equal to the same concentration. A difficulty in this approach is the low intensity of the X-ray signal. Compared to the BSE signal, the X-ray signal from an element is lower by several orders of magnitude. Quantification of the X-ray mapping would cause a much longer scanning time (e.g.: to neglect matrix effects for an element with around 1wt%, the scanning time would have to be increased to 2000s) (5). 416 HMJ56 B. Wollein et al.

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3. Results:

3.1 Solid/solid hard phase/binder metal diffusion couples:

Comparing the reaction behaviour of the solid hard phase (W0.48Ti0.52)C with different solid binder metals it was found that the reaction zone is 2 1/i times thicker in contact with Ni than with Co (compare Fig.2b and Fig.3b). In addition, elongated microstructural constituents perpendicular to the interface were found (see Fig.3a) if Co is in contact with the hard phase. It was not possible yet to characterise these parts by SEM or EPMA. This phenomenon is in so far interesting as the microstructure changes if nitrogen is added to the hard phase which will be shown later. The reaction zone with Ni builds two zones while in contact with Co only one zone can be identified (Fig.2a, brighter and darker parallel zones and Fig.3a). This observation was confirmed by the EPMA line scans made across the interface of both diffusion couples (Fig.2b, 3b), which also provide quantitative information on the reaction zones. In the line scan of the Ni/hard phase diffusion couple (Fig.2b) it can be seen (13) that Ti, C and W first decrease from right to left slowly to 26at% Ti, 38at% C and 12at% W (12) and then steeply to 4at% Ti, 12at% C and 6at% W. Inverse to this, Ni increases from 13 to 12, then steeply to 11. Comparing these results with the line scan of the Co/hard phase diffusion couple (Fig.3b) the concentrations of Ti, C and W decrease steeply from 42at%, 46at% and 12at% (12) to 2at% Ti, 9at% C and 3at% W (11), while Co increases from Oat% (12) to 86at% (11). If the results of the Ni/hard phase diffusion couple (Fig.2) and the Co/hard phase diffusion couple (Fig.3) are compared it is obvious that 4at% Ti and 6at% W diffused into Ni (Fig.2b, 11) while nearly no Ti and W was found in Co (Fig.3b, 11). The better solubility of Ti in Ni is in agreement with other investigations (6,7) and deviated somewhat from (8), who found 11at% W in Co. Because of the long annealing time Ti was found in Ni up to about 500um away from the interface. B. Wollein et al. HM 56 417

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J ! • ; • 1 U ' ' ' A 12 13 • Ti _ 90.0 : _ . - - W - " ,W2C - - . isjr- 80.0 g ~ ~~ •' "' \ O 7O.0Ü Ü 1 _ 6O.O o - V in - __p Ü 1 50.0 F Ni 1

'• 40.0 ra

A - ) - 30.0 M J 1 I J / 1 20.0 ; i I

- -' - N / 10.0

O.O f '. l 1.1,1 1 . ' 1- T • I" .- r " " 1O0 140 18O 220 260 3OO 340 38O 42O 4SO 5O0 distance / microns

Fig.2 a: Microstructure of the diffusion couple Ni/(Wo.48Tio.52)C etched with FeCb/HCI and Murakami solution; b: Line scan over a distance of 400um across sample Ni/(Wo.48Tio.52)C.

To investigate the influence of nitrogen on the diffusion behaviour diffusion couples with the solid hard phase (W0.24Tio.76)(Co.7N0.3) were annealed in contact with solid binder metals. The most interesting phenomenon found was, that contrary to the samples without nitrogen, in the Ni/(Wo.24Tio.76)(Co.7No.3) diffusion couple the black microstructural constituents perpendicular to the interface were found, while in the Co/(Wo.24Tio.76)(Co.7N0.3) diffusion couple they were absent (see Fig.2a, 3a, 4a and 5a). Comparing the different hard phases with the same binder metal the reaction zone is thinner, both with Ni and Co, if the hard phase contains nitrogen (Figs. 2b, 3b, 4b and 5b) although the reaction time was 30% longer. 418 HM 56 B. Wollein et al. 5" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

J ' 1 • [ ' 1 1 • 1 ' 1 • 1 ' 1 ' u 1OO.O n I2 - . . . Ti N 90.0 - *" ""— G ~ . w - 1 so.og o 70.0^ ü O " g Co 1 60.0 0 _p 50.0F ' \. - - -r 40.0 ra - W "' 30.0 1 •'A - 20.0 •••/V- •'' 10.0 :,f .4 V \ """'_ "1" *~V *~"r . r . i , o.o 25O 285 32O 355 39O 425 4SO 49S 53O 565 GOO d stance / microns

Fig.3 a: Microstructure of the diffusion couple Co/(W0.48Ti0.52)C etched with FeCb/HCI and Murakami solution; b: Line scan over a distance of 350um across sample Co/(W0.48Tio.52)C.

If the nitrogen-containing hard phase is contacted with Ni again two zones within the diffusion zone were built (Fig.4b). Contrary to the reaction of Co with the nitrogen-free hard phase (Fig.3), two zones have built in the reaction zone of Co with the nitrogen-containing sample (Fig.5b). In the line scan of the Ni/(Wo.24Tio.76)(Co.7No.3) sample (Fig.4b) it was found that the nitrogen content decreases down to Oat% in the first zone (13 to 12), while Ti and C decrease slowly to 4at% Ti and 7at% C across the reaction zone (13 to 11). W remains nearly constant in between the reaction zone and decreases within Ni. Analysing the line scan of sample Co/(Wo.24Tio.76)(C0.7N0.3) (Fig.5b) a slight decrease of Ti from 13 to 36at% (12) was found. Contrary to this the N and C content increase from 13 to 11at%N and 42at% C (12), while the W concentration is nearly constant in between 13 and 12. Then W, Ti and C decreases to 3at% W, 1at%Ti and 5at% C (11). The N concentration decreases to 8at% (11) and remains constant at around 7at% in the binder phase (Fig.5b). B. Wollein et al. HM 56 419

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3ONI1 .COR

• i ' 1 ' l_ n 13 o W - o -v. I — " - TSI CO _ 70.0 z d 60.0 ~ Nl r Ni ? / 500C 1 40.0 ra I • •;. 'r d r 3O.O 1 CM O * * ~" "H fj 20.0

y; r* 'V I 10.0 JT. - •"•— AA O.O -,•:.-_..,_.._,.._ •i~ ;• -r — "i— V —i" . 1 . i . r 84 112 140 168 196 224 2S2 230 distance / microns

Fig.4 a: Microstructure of the diffusion couple Ni/(Wo.24Tio.76)(Co.7N0.3) etched with FeCI3/HCI and Murakami solution; b: Line scan over a distance of 280um across sample Co/(Wo.24Tio.76)(C0.7N0.3).

w If the reaction of (Wo.24Tio.76)(Q).7No.3) ^h Ni is compared with that of the Co side it was found that again Ni reacts faster with the nitrogen-containing hard phase (compare Fig.4b, 5b). The most significant difference is the N diffusion into the binder metals. While the N concentration with Ni as the binder metal decreases down to Oat% already in the reaction zone (Fig.4b, 12), 8at% N were found within Co at 11 (Fig.5b). This change of solubility of N was already seen in contact with a binder metal with the composition of 56at% Co and 35at% Ni where the N concentration decreased already to Oat% in within the reaction zone. So already little additions of Ni influence the solubility of N in the binder metal substantially. 420 HM 56 B. Wollein et al.

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1OO.O 12 13 9O.O ,

SO.O_

CO Ö

u Co Ö 50.0

4O.O ^

. > •"* ~~ o • • . 3O.O

CM ^ 2O.O

10.0

\. .. 0.0 O 14 28 5e 7O 84 98 112 12S 14O [stance / microns

Hg.5 a: Microstructure of the diffusion couple Co/(W0.24Tio.76)(Co.7N0.3) etched with FeCI3/HCI and Murakami solution; b: Line scan over a distance of 140um across sample Co/(W0.24Tio.76)(Co.7N0.3);

3.2 Solid/liquid hard phase/binder metal diffusion couples:

In order to investigate the influence of addition of Fe to Co and Ni the reactions of two hard phases, one with and one without nitrogen, with liquid Co/Ni and Fe/Ni/Co binder phases were investigated. The reaction of (W0.48Ti0.52)C with Co/Ni 50/50 is such intensive that after 5 minutes annealing the compact hard phase was dissolved completely (Fig.6a). If only the binder metal is changed to Fe/Ni/Co it was observed (Fig.6c) that the reaction slows down substantially. This reduction in reaction rate was also observed in the nitrogen-containing diffusion couples (Fig.6c, 6d). Interestingly, the reaction of the nitrogen-containing sample in contact with the Fe-containing binder metal (Fig.6b) is much faster than the reaction of the nitrogen-free sample contacted with the same Fe-containing binder metal (Fig.6c). B. Wollein et al. HM 56 421

15;h International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

•:.T :^.JS WF?

CoM

{W0.24TiOJ6)(C0.7N0.3)

Abb. 6a-d Microstructures of solid/liquid hard phase/binder metal diffusion couples;

This is inverse to the results of the solid/solid diffusion experiments, where the reaction zones of the nitrogen-free diffusion couples were, independent of the composition of the hard phase, thicker than the reaction zones in nitrogen-containing samples (compare Figs. 1-5). Another phenomenon is the change in the microstructures of the nitrogen-containing samples by changing the binder metal from Co/Ni to Fe/Ni/Co. In the Fe-containing diffusion couple it is possible to identify the original shape - and also the original interface - of the hard phase (Fig.6b), while with in the Fe-free binder metal the hard phase is dissolved (Fig.6d). Comparing the reaction of the Fe-containing binder metal with a nitrogen-containing (fig.6b) and a nitrogen-free (Fig.6c) hard phase the same change in microstructures was found. The nitrogen-free hard phase (Fig.6c) is dissolved within the binder metal without the possibility to identify the original geometry of the hard phase.

3.3 Hard phase/hard phase diffusion couples:

At 1500°C the reaction rate of the Ti(Co.7No.3)/(Tio.7Wo.3)C diffusion couple is so low that after an annealing time of 1008h it was not possible to identify an diffusion zone by light-optical microscopy. In Fig.7 the SEM images of the 422 HM 56 B. Wollein et al.

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sample Ti(Co.7No,3)/(Tio,7Wo.3)C (a) and Ti(C0.8No.2)/(Tio.8IVlOo.2)C (b) after 1512h are shown. In both of these microstructures it is possible to identify a reaction zone between the two hard phases but it can be seen that the diffusion zone of sample Ti(Co.8N0.2)/(Tio.8Mo0.2)C (Fig.7b) is thicker than in sample Ti(Co.7No.3)/(Tio.7Wo.3)C (Fig.7a). Hence, the Mo-containing couple shows faster reaction than the W-containing diffusion couple.

(TL Mo ,)C 50 urn 50 urn 8 0

Fig.7 A, b: SEM images of the hard phase/hard phase diffusion couples annealed for1512hat1500°C;

In the X-ray maps made across the interface of the diffusion couple Ti(Co.7No.3)/(Tio.7Wo.3)C annealed for 1008h thin W-rich rims can be observed (Fig.8a) In these rims the nitrogen content decreases (Fig.8b). This corresponds to the low chemical affinity between W and N. By continuing the annealing of this diffusion couple it was found that the thickness of these rims remains constant (Fig.8a, 8c) and W diffuses deeper along the grain boundaries into the Ti(C,N) side. Obviously, the core-rim structure remains stable at positions near the interface where near-equilibrium conditions are reached. To find out more about the diffusion processes it seems to be necessary to anneal the diffusion couples even for a longer time to investigate whether the thickness of the W-rich rims remains constant or changes if the equilibrium is reached. B. Wollein et al. HM 56 423

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Fig.8 X-ray maps on sample Ti(CojNo.3)/(Tio.7Wo.3)C; a: W; b: N; c: W; d: N; a, b: 1008h; c, d: 1512h; white: high concentration, black: low concentration of the respective elements; for the colour version see the microstructure section at www.tuwien.ac.at/physmet

In the X-ray maps of the Ti(C0.8N0.2)/(Tio.8lvloo.2)C diffusion couple (Fig.9) it can be seen that after 1008h (Fig.9a, 9b) Mo diffused already into the Ti(C,N) and hence Mo-rich rims were built around the Ti- and N-rich zones (Fig.9a) and is dissolved. Again an enrichment of N inside the Mo-poor parts is clearly visible (see Fig.9b and 9d). This is due to the low chemical affinity between Mo and nitrogen. After 1512h (Fig.9c, 9d) the rims around the Ti- and N-richer parts grew and thus became thicker; contrary to the rims in the W-containing diffusion couples. This diffusion couple should be annealed for a longer time in order to observe whether the core vanishes and the grains become homogeneous. However, already from these investigations it can be 424 HM 56 B. Wollein et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 concluded that in the W-containing sample a stable core-rim structure is formed because of the miscibility gap whereas no such stable miscibility gap is present at the investigated conditions (temperature, composition) in the Ti- Mo-C-N system and hence the grains in diffusion couples of the latter homogenise to a single phase near the interface. In sintering experiments the core-rim structure of hard phase particles was attributed to the presence of a miscibility gap (compare (9)).

Fig.9 X-ray maps on sample Ti(C0.8N0.2)/(Tio.8Moo.2)C; a: Mo; b: N; c: Mo; d: N; a, b: 1008h; c, d: 1512h; white: high concentration, black: low concentration of the respective elements; for the colour version see the microstructure section at www.tuwien.ac.at/physmet B. Wollein et al. HM 56 425

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4. Conclusion:

It was found that both with nitrogen-free and nitrogen-containing hard phases the reaction with Ni is faster than with Co at T<1230°C. If the same hard phases were compared with different binder metals it was observed that in the nitrogen-free hard phase in contact with Ni 4at% Ti are soluble, while in contact with Co no Ti dissolved in the binder metal. Adding nitrogen to the hard phase the most interesting finding was the change in the microstructures. While contacted with the nitrogen-free hard phase Co built black microstructural constituents perpendicular to the interface, while contacted with the nitrogen-containing hard phase Ni built black microstructural constituents perpendicular to the interface. Future attempts will concentrate to characterise these phases. In contact with nitrogen- containing hard phases, ca. 8at% N at the interface with Co were found, where no N diffused into Ni. Ni seems to influence the solubility because in the interdiffusion zone between 100at% Co and 100at% Ni it was observed that no N diffused into a binder metal with 56at% Co and 35at% Ni. Adding Fe to Co and Ni the reaction both with the nitrogen-free and nitrogen- containing hard phase slowed down substantially. Interestingly the Fe- containing binder metal reacts faster with nitrogen-containing hard phases which is inverse to the Fe-free binder metals, independent on the N content. The Ti(C,N) phase contacted with (Ti,W)C showed already after annealing for 1008h W-rich rims surrounding Ti- and N-rich cores. The thickness of these rims did not change even if the diffusion couples were annealed for further 504h. In Ti(C,N) in contact with (Ti,Mo)C even after 1008h a core-rim structure was built by diffusion of Mo into Ti(C,N), but the structure was not stable and the core size decreased.

5. References

1 H. Holleck: „Binäre und Ternäre Carbid- und Nitridsysteme der Über- gangsmetalle", (Geb. Borntraeger Berlin, Stuttgart, 1984) 2 V.V. Popov: Phys. Metall. Metall. 80 (1995), pp. 550-556 3 P. Wally, M. Ueki, P. Ettmayer, J. Jap. Soc. Powder & Powder Metall., 43 (1996), pp. 520-525 4 W. Lengauer, J. Bauer, M. Bohn, H. Wiesenberger and P. Ettmayer, Mikrochim. Acta, 126 (1997), pp. 279-288 5 J. Goldstein, A. D. Romig, D. Newbury, C. Lyman, P. Echlin, C. Fiori, D. Joy and E. Lifshin: in "Scanning Electron Microscopy and X-ray Microanalysis" (2nd ed., Plenum Press, New York, 1992) 426 HM 56 B. Wollein et al.

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6 D. Bandyopadhyay, R.C. Sharma and N. Chakraborti, J. Phase Equilibria, 21 (2000), pp.179-185 7 Y. Du and J.C. Schuster, Z. Metallkunde, 89 (1998), pp. 399-410 8 O. Lavergne, C.H. Allibert, High Temp. - High Pressures, 31 (1999), pp. 347-355 9 P. Lindahl, P. Gustafson, U. Rolander, L. Stahls, H.-O. Andren, Int. J. Ref. Mat. & Hard Mat., 17 (1999), pp. 411-421 B. Wittmann et a). HM 57 427

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WC GRAIN GROWTH AND GRAIN GROWTH INHIBITION IN NICKEL AND IRON BINDER HARDMETALS

Bernhard Wittmann, Wolf-Dieter Schubert, Benno Lux Institute for Chemical Technology of Inorganic Materials Vienna University of Technology, A-1060 Vienna, Austria

Abstract:

WC grain growth and growth inhibition of an 0.6 (jm FSSS WC powder (average SEM size: 0.35 pm) was studied in WC-10 wt% Ni alloys by adding 0-2 wt.% of inhibitor carbides (VC, Cr3C2, TaC, TiC and ZrC).

Alloy gross carbon content turned out to be a crucial factor for WC growth in Ni alloys, even with high inhibitor additions. Coarsening was more pronounced in high carbon alloys, compared with low carbon grades, resulting in a significantly lower hardness. VC proved to be by far the most effective grain growth inhibitor in WC-Ni hardmetals, followed by TaC, Cr3C2, TiC and ZrC. Hardness increased with increasing amount of additive but reached a maximum above which it remained about the same.

Experiments on WC-Fe-(VC) alloys revealed that WC grain growth is strongly restricted in Fe-binder alloys, even without additions of growth inhibitors.

Binder chemistry thus strongly influences both continuous and discontinuous WC grain growth. This chemistry is determined by the nature of the binder matrix (Fe, Co, Ni), the alloy gross carbon content (which determines the composition of the binder matrix) as well as the inhibitor additive.

Keywords: hardmetal, alternative binder, WC grain growth, WC grain growth inhibition. 428 HM 57 B. Wittmann et al.

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1 Introduction For the production of hardmetals tungsten carbide (WC) is the basic and most widely used hard compound, whereas cobalt was found to be the optimal binder metal for most applications. Other iron group metals (nickel and iron) are employed only to a limited extent for specialized applications, e.g. where hot hardness and resistance against thermal cracking or corrosion/oxidation resistance are required [1,2]. Compared with the WC-Co system, which has been extensively investigated in the past due to its commercial importance, there is still limited information on alternative binders. Up to now, for example, no systematic study has been performed on WC grain growth and growth inhibition in those systems. Thus it appeared to be worthwhile to carry out a respective investigation, and to compare the results on WC grain grain growth and growth inhibition with those already known from the well investigated WC-Co system [3]. WC grain growth and grain growth inhibition was investigated using a straight WC-10 wt% Ni grade and additions of small amounts of inhibitor carbides (VC, Cr3C2, TaC, TiC, ZrC); as practice in the manufacture of fine-grained WC-Co hardmetals. WC grain coarsening in a WC-10 wt% Fe alloy was additionally investigated.

2 Literature Little has been reported concerning WC grain growth and growth inhibition in WC-Ni alloys. No direct evidence was found in related literature on grain growth in unalloyed WC-Ni. However, the early results of Suzuki et al. [4] demonstrated a strong influence of the carbon content on the hardness of the sintered material. For a WC-10 wt% Ni alloy, the hardness varied between -1060 HV30 for a high carbon alloy and -1280 HV30 for a low carbon alloy, indicating a strong impact of binder chemistry on WC grain coarsening. Excessive grain growth was observed in WC-6 wt% Ni alloys at 1500°C with a plain nickel binder [5]. Both chromium and molybdenum were found to act as grain growth inhibitors [5,6], but no systematic data were available on the inhibiting effect. Alloying additions of V, Mo and Crwere observed to increase the hardness in Fe-Co-Ni-bonded hardmetals, and VC was found to be the most effective grain growth inhibitor, as it is in the WC-Co system [3]. Data for the solubility of VC, Cr3C2, TaC, NbC, TiC and ZrC in nickel is scarce [7]. B. Wittmann et al. HM 57 429

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3 Experimental 3.1 Starting materials High purity starting materials (WC, Ni, Fe) were used for the preparation of alloys. A SEM image of the WC powder used is presented in Fig. 1.

Fig.1: SEM image of the WC powder used for this investigation; grain size: 0.35 |jm (estimated by SEM-image analysis) calculated diameter from BET: 0.2 pm; magn. 20 000 x.

3.2 Alloy preparation (90-x) wt% WC, 10 wt% Ni and x wt% inhibitor carbide were immersed in cyclohexane and either ball milled for 72 hours (steel lining, 60% of the critical rotation speed) or attritor milled for 2 hours (700 rpm, plate stirrer). 2 wt.% wax was added as a pressing aid. Different amounts of carbon black were added to the individual batches in order to obtain the desired gross carbon content of the alloys. After milling, the cyclohexane was evaporated using a Rotavapor. The powder mix was passed through a 0.5 mm sieve and subsequently granulated in a Turbula Mixer for 10 minutes to obtain a smoothly flowing granulate. 11 g of this granulate were pressed in a hydraulic press at 200 MPa to bars of 6 x 5 x 50 mm. 430 HM 57 B. Wittmann et al.

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3.3 Sintering Sintering was carried out in an industrial GCA vacuum furnace for 1 hour at 1480°C (1 hour at 1350°C in case of the iron alloys). The heating rate was 10°C/min in the temperature range to 1250°C, and 3°C/min above. The total pressure was <0.01 mbar. The samples were sintered in graphite boats with ribbed bottoms. In order to achieve alloys with high gross carbon content (graphite precipitation) some of the samples were covered with graphite or carbon black. The boats were covered with a graphite plate in order to create a micro- atmosphere whose carbon potential is dominated mainly by the alloys and not by the degassing of the carbon felt or by small furnace leakages.

3.4 Vickers hardness measurement and microscopic investigations The sintered samples were cut with a diamond saw, embedded in Bakelite® , ground and mirror polished with diamond suspensions. Vickers indentations were carried out on a hardness tester M4U 025 by EMCO (Austria). An indenter load of 30 kgf was used for indentation, as recommended by ISO 3878. SEM imaging was performed on etched sections (Murakami solution; ~5 minutes) using a JEOL 6400 electron microscope.

4 Results Table 1 presents an overview of the sintering experiments. Both two phase (WC+Ni) as well as three phase alloys (WC+Ni+eta and WC+Ni+C, respectively) were investigated, to consider the influence of the alloy gross carbon content on the growth process. Table 1 also summarizes the results of the hardness measurements that give a first indication of the efficiency of the growth inhibition. B. Wittmann et al. HM 57 431

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Table 1: Overview on sintering experiments and HV 30 values of the respective alloys (sintered at 1480°C for 1 hour) containing 2-phase: graphite r|-phases WC-10 Ni precipitation WC-10Ni HV30: 1403 1032 1000

VC 0.2% 1585 1582 1308 0.5% 1574 1649 1464 1.0% 1826 1879 1630 2.0% 1738 1802 1699

Cr3C2 0.2% 1355 1383 1201 0.5% 1480 1490 1405 1.0% 1591 1580 1463 2.0% 1574 1504 1467 TaC 0.2% 1377 1012 935 1.0% 1593 1615 1506 TiC 0.2% 1589 1373 1187 1.0% 1405 1370 1290 ZrC 0.2% 1554 1116 1047 1.0% 1512 1357 1023

4.1 WC grain growth in undoped (inhibitor free) WC-Ni alloys SEM images of an undoped high carbon and low carbon alloy are presented in Figure 2. From this comparison it can be seen that the alloy gross carbon content exerts a considerable influence on WC grain growth. Strong grain growth occurred in the high carbon alloy (HV30:1032); whereas a very non- uniform but significantly finer microstructure was formed in case of the low carbon grade (HV30:1403). 432 HM 57 B. Wittmann et al.

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WC-10 wt% Ni; low carbon alloy WC-10 wt% Ni; high carbon alloy isothermal hold for 1 h at 1480°C 10 urn

Fig. 2: SEM images of microstructures of I.e. (left) and h.c. alloy (right); 2000 x.

This characteristic difference in grain growth already develops during the heating-up of the sintering cycle, as can be demonstrated by interrupted sintering experiments, where the samples were cooled down immediately after heating to 1480°C (Figure 3). On isothermal liquid phase sintering, only limited further WC grain growth occurred, shifting the average grain size to higher values. While, in case of the low carbon alloy a significant part of the fine-grained portion survived the isothermal sintering, most of it disappeared in the high carbon alloy by dissolution of the finer particles and growth of the larger (Ostwald ripening).

WC-10 wt% Ni; low carbon alloy WC-10 wt% Ni; high carbon alloy heated up to 1480°C/1 mm 10 um

!LJ Fig. 3: SEM images of microstructures of I.e. (left) and h.c. alloy (right); 2000 x. B. Wittmann et al. HM 57 433

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4.2 WC grain growth in doped WC-Ni alloys 4.2.1 VC-doped alloys VC proved to be by far the most effective grain growth inhibitor in WC-Ni alloys. The strong growth inhibiting effect can be deduced from Figure 4, which shows a considerable hardness increase with increasing VC additions, corresponding to the decrease in WC grain size (Figure 5). High carbon alloys always showed a significantly lower hardness than their low carbon variants, even with high VC additions.

1 wt% VC turned out to be the most efficient additive, leading to a high hardness, and a high degree of microstructural uniformity. No WC grains larger than 1 urn were detected in this alloy. At higher additions (2 wt% VC), no further increase in hardness was observed in two phase alloys, but nest- like precipitates of (V,W)C occurred.

2100 ,

1900 2-phase eta-phase 1700 C-precipitations O ro 1500

1300

1100

_l 0,5 1 1,5 2,5 VC-addition [wt%] Fig. 4: Hardness [HV30] vs. VC-addition [wt%]; WC-X wt% VC-10 wt% Ni alloys; sintered at 1480°C for 1 hour. 434 HM 57 B. Wittmann et al.

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HV30: 1582 0.5% VC

Fig. 5: Microstructures of VC-doped two-phase WC-Ni alloys; sintered at 1480°C for 1 hour; magnification: 4000 x; at 2.0% VC additions (V,W)C is formed on cooling.

4.2.2 Cr3C2-doped alloys

The influence of Cr3C2-additions on the WC grain growth was less pronounced than that of VC in terms of hardness and grain refinement (Fig. 6). Again, high carbon alloys showed a lower hardness than low carbon alloys (Table 1). Growth inhibition was more pronounced with increasing amounts of inhibitor carbide, but reached a saturation limit, above which neither a higher hardness, nor a finer microstructure was obtained (Fig. 7). B. Wittmann et al. HM 57 435

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Hardness vs. Inhibitor addition

0 0,5 1 1,5 2 2,5 inhibitor addition [wt%]

Fig. 6: Two-phase WC-10 wt% Ni alloys with additions of VC, Cr3C2 and TiC; sintered at 1480°C for 1 hour.

WC-10 wt% Ni-1 wt% Cr3C2 WC-10 wt% Ni-2 wt% Cr3C2 H 1 |jm

Fig. 7: Microstructures of two-phase WC-10wt% Ni alloys containing 1 and 2 wt% Cr3C2, respectively; magnification: 4000 x. 436 HM 57 B. Wittmann et al.

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4.2.3.TaC-doped alloys Small additions (0.2 wt.% TaC) did not significantly alter the grain growth as compared with the undoped alloy, but at 1 wt% TaC a considerable growth inhibition took place. The hardness of this alloy was about the same as that of the 1 wt% Cr3C2 addition (HV30: 1615 [Ta] and 1580 [Cr], respectively). At 1 wt% TaC addition, nest-like precipitation of (W,Ta)C occurred in the microstructure, independently of the gross carbon content.

4.2.4.TiC-doped alloys For TiC a clear growth inhibition was observed, even at a level of 0.2 wt% TiC (Figure 6). This inhibition was more pronounced in the low carbon (eta-phase containing) alloy. This indicates a strong additional effect, due to the tungsten-rich melt formed in such alloys. Even in the case of the 0.2 wt% TiC additions, precipitation of (Ti,W)C occurred on cooling, indicating a low solubility of titanium carbide in nickel. The most striking microstructural feature in TiC-doped alloys is the occurrence of plate-like WC grains (sizes up to 6 urn), which are embedded in a still fine-grained WC matrix (Fig. 8).

Fig.8: Microstructure of a WC-1 wt% TiC-10 wt% Ni alloy; sintered at 1480°C for 1 hour; note the plate morphology of the WC; dot-like precipitations of (W,Ti)C form on cooling; magn. 4000 x. B. Wittmann et al. HM 57 437

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4.2.5 ZrC-doped alloys Compared with TiC, ZrC is an even weaker growth inhibitor, and only works in the low carbon activity range (i.e. in a tungsten-rich, but carbon-poor environment of the growing WC). No growth inhibition at all was observed in high carbon alloys with graphite precipitations. Precipitation of (Zr,W)C occurred even at 0.2 wt% ZrC additions, indicating a low solubility of ZrC in nickel. Interestingly, such precipitates were not only formed in the nickel binder, or at the WC/Ni interface, but also appeared to be incorporated within large WC grains.

4.3 WC grain growth in WC-10 wt% Fe alloys Also, for comparison purposes, two WC-10 wt% Fe alloys were investigated for their grain growth behavior; one alloy without inhibitor additive, the other with 1 wt% VC. The WC powder used for this investigation was the same as that used as the standard grade for the WC-Ni hardmetals. The undoped alloys showed a very fine and uniform microstructure, even in case of the high carbon grade (Figure 9). Overall grain growth was more pronounced in this variant compared with the low carbon alloy, but in both cases the microstructure was fine and very uniform, showing only a few WC grains larger than 1um.

HV30: 1916 HV30: 1812 I—i 1 urn :•» ik..1 ;\'V% * .** *.%' '***' ' «•*••>• ^%>r *r A-ys iv-t: r:>i **, * i * «- .* * - • . n *

Fig. 9: WC-10 wt% Fe alloys: two phase (left), with graphite precipitates (right); magnification: 10 000 x. 438 HM 57 B. Wittmann et al.

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Adding VC to the WC-10 wt% Fe alloy did lead to an additional further growth inhibition but the effect was much less pronounced than in case of nickel- or cobalt-bonded grades, due to the already very strong inhibition through the iron matrix (Figure 10). The lower hardness of this alloy (compare Fig. 9, 10) can be explained by the high degree of residual porosity that occurred in this material.

*•% f ^

o? * - > -<*- * *x t > »'

* \,*;r

v •* 1 urn t *, Fig. 10: Microstructure of a WC-1wt% VC- 10wt% Fe alloy; HV30:1884; sintered at 1350°C for one hour; magnification: 10000 x.

Discussion

5.1 General considerations on WC grain growth and growth inhibition WC grain growth during liquid phase sintering of WC-Co alloys is phenomenologically treated as an Ostwald ripening process [8-10]. Smaller grains dissolve due to their higher dissolution potential (increased chemical potential), while coarser ones grow by material re-precipitation, thereby reducing the interface area of the system. Overall (continuous) WC grain growth proceeds rather slowly in conventional WC-Co hardmetals compared with other carbide/binder systems (VC/Co, NbC/Co, TaC/Co). This is attributed to the very low WC/Co interface energy (<10"2 J/m2) and the relatively high activation energy of the solution- B. Wittmann et al. HM 57 439

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reprecipitation step, indicating an interface controlled growth mechanism [11].

The growth rate of the WC in WC-Co hardmetals is remarkably increased in high carbon alloys (stoichiometric and over-stoichiometric alloys) compared with low carbon alloys (sub-stoichiometric alloys), and can be significantly decreased by additions of growth inhibitors, such as VC, Cr3C2, TaC and NbC. In all cases, the influence of the "additive" (C, W, inhibitor) has been explained by effects at the WC/Co interface. Thus, theories on the mechanism of grain growth inhibition assume either an alteration of the interface energies or an interference of the growth inhibitor with the interfacial dissolution-nucleation-reprecipitation steps [11]. The additives are soluble in the cobalt binder and most apparently segregate at the WC/Co interface during sintering.

Several mechanisms have been suggested: face-specific adsorption (leading to a decrease of the respective interface energies), face-oriented deposition (interface alloying) or blocking of active growth centres of the crystals [10], including a change in edge energy (on considering a two-dimensional nucleation-controlled coarsening process) [12]. Also the retardation of grain growth during sintering in the presence of Cr3C2-additions by a reduction of the tungsten flux through the binder phase was discussed [13].

The results obtained in this study, which extend the knowledge on WC grain growth as described above, to nickel- and iron-based hardmetals, have revealed several new and interesting aspects in WC grain growth and growth inhibition: • First of all, the strong growth inhibition in WC-Fe alloys, which occurred even without additions of growth inhibitors. • The small influence of variations in the gross carbon content on the WC grain growth in WC-Fe alloys. • The absence of any discontinuous WC grain growth in WC-Fe alloys (despite the small average particle size of the starting WC: ~0.35um). • The strong growth in WC-Ni alloys, in particular in inhibitor free high carbon alloys. • The significant growth inhibition in low carbon WC-Ni alloys (tungsten- rich binder environment), in particular in alloys with eta-formation. • The outstanding role of VC as growth inhibitor in WC-Ni alloys, similar to that observed in WC-Co grades [3]. • Effects of TiC on WC shapes in WC-Ni hardmetal. 440 HM 57 B. Wittmann et al.

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The results further demonstrate that observations made on WC-Co alloys in terms of grain growth inhibition through inhibitor carbides (VC, Cr3C2, TaC, TiC, ZrC) can be directly transferred to WC-Ni alloys. For a certain inhibitor additive (e.g. VC) the finest microstructure is obtained when the saturation concentration of the respective compound in nickel is reached at sintering temperature [3]. This relationship is demonstrated in Fig.6 for the WC-10 wt% Ni alloys with additions of VC, Cr3C2 and TiC. Above this limit, no further growth inhibition took place.

5.2. The importance of the growth environment

The chemistry of the solid and liquid binder surrounding the growing WC seems to be a crucial aspect in WC grain growth. This chemistry varies significantly with changes in the gross carbon content or additions of growth inhibitors, and determines the interface chemistry and kinetics. This is shown in Fig. 11 for the case of a high and low carbon WC-Ni alloy at a temperature of1500°C.

Thus, a carbon-rich, tungsten-poor nickel interface (formed on sintering), significantly promotes WC grain growth, whereas a tungsten-rich, carbon- poor nickel interface has a growth inhibiting effect. This significant change in binder chemistry becomes even more obvious if the atomic C:W ratio is considered, which changes from -2:1 (high carbon) to -1:3 (low carbon).

A: 5.4 at% W 10at%C 84.6 at% Ni (binder composition in high carbon alloys)

B: 17.6 at%W 6 at% C 76.4 at% Ni (binder composition in low carbon alloys)

Fig.11: Nickel-rich area of the Ni-W-C phase diagram at 1500°C, according to [15]. B. Wittmann et al. HM 57 441

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Table 2 gives several examples for the large variations in binder chemistry, depending on the nominal composition of the respective alloys. Variations in chemistry are most likely to result in both changes in interface energy (i.e. driving force for growth) as well as changes in activation energies and mass transport of the nucleation/growth process (growth kinetics; resistance against growth).

Table 2: Melt composition of alloys based on different binders and carbon activities; values taken from [14,15]; temperature: 1500°C. alloy high carbon (W+M)/C* low carbon (W+M)/C* at% wt% at% wt% at% binderat binderat binder binder 1500°C 1500°C 1500°C 1500°C WC- 72%WC 19.4% W 5.7% W 0.27 48.0% W 19.6% W 1.5 10 Fe 28% Fe 4.7% C 21.1% C 2.1% C 13.3% C 75.9% Fe 73.2% Fe 49.9% Fe 67.1% Fe WC- 73%WC 22.5% W 7.5% W 0.5 48.3% W 21 % W 1.9 10 Co 27% Co 2.9% C 15% C 1.7% C 11% C 74.6% Co 77.5% Co 50.1% Co 68% Co WC- 73%WC 16.3% W 5.4% W 0.54 41.5% W 17.6% W 2.9 10 Ni 27% Ni 2%C 10% C 0.9% C 6%C 81.7% Ni 84.6% Ni 57.6% Ni 76.4% Ni WC- 72.2%WC 15.1% W 5.2% W 0.76 39.4% W 17.2% W 3.4 10 Ni- 27% Ni 1.9% C 10.1% C 0.85% C 5.7% C 1 TaC*) 0.8% TaC 75.9% Ni 82.2% Ni 54.6% Ni 74.8% Ni 7.1% Ta 2.5% Ta 5.1% Ta 2.3% Ta WC- 70.9%WC 15.3% W 5%W 1.3 39.65% W 16.4% W 4.2 10 Ni- 26.6% Ni 1.9% C 9.5% C 0.86% C 5.5% C 1 VC*) 2.5% VC 76.6% Ni 78.2% Ni 55% Ni 71.4% Ni 6.2% V 7.3% V 4.5% V 6.7% V *) this assumed that the addition is fully dissolved in the nickel at 1500X; The values are estimates, since no information is available on the quaternary systems *M...V, Ta 442 HM 57 B. Wittmann et al.

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Only recently it was postulated that the formation of metal-carbon clusters inhibits liquid phase diffusion of W and C atoms in the Co-rich melt, and that the grain coarsening is impeded by the incorporation of W and C atoms in such a stable cluster form. The strong bonding between the refractory metal (W, V, Cr, Ta) and non-metal (carbon) was considered to be the key for cluster stabilization [16].

The present work demonstrates that the higher the metal-to-carbon ratio (W+V/C, W+Cr/C, W+Ta/C) for a certain alloy (binder), the more effective is the growth inhibition in WC-Co(Ni,Fe) alloys (the upper limit is set by the thermodynamics and reaches its maximum in low carbon alloys at temperature). In this regard it is important to consider that although small amounts of inhibitor carbides are commonly added (0.2 to 1 wt%), large amounts of inhibitor atoms are dissolved in the binder at temperature (see Table 2).

More generally speaking, the nature of the binder (Co, Ni, Fe) contributes to determining the growth behaviour by influencing the metal-to-carbon bond relationship. For example, iron has the highest affinity to carbon, followed by cobalt and nickel. This affinity can also be interpreted as the ability to form stable metal-carbon bonds (also in the liquid phase), the iron masking the carbon, and impeding carbon transport and precipitation by increasing the activation energies of the nucleation and growth processes (i.e. increasing the resistance against growth).

A decrease of the interface energy of the WC/binder interface might also be a possible explanation for the extremely low coarsening rate of the WC in case of the iron binder (through bond formation), compared with the WC-Co and WC-Ni system.

The strong influence of the respective binder system on WC grain growth is demonstrated in Fig. 12 for the case of three high carbon alloys (containing graphite precipitations), originating from the same WC powder (Fig.1) but with different binder matrices (Co, Ni, Fe). No inhibitor carbides were added. B. Wittmann et al. HM 57 443

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all alloys saturated with carbon 1 1 10 pm sintered at: 1480°Cfor1h

»a, WC-10Ni-C * ^ „ f V W HV 30: 1000 #^ n •I * ^ \ K V V eutectic melt formation: X s. ~1342°C 1 "V sintered at: 1400°Cfor 1h 4 i

WC-10CO-C 1 v'J HV30: 1395 eutectic melt formation: i7 Vj £&? $* ^K f x ^ C • ~1275°C EX ^ "*" A sintered at: 1350°Cfor 1h

• ^ i ^ ^ *

<• WC-10Fe-C >•• HV30: 1812

eutectic melt formation: ; i -i ' ~1143°C i Fig. 12: Microstructures of WC-10wt% Co (Ni, Fe) alloys, based on the same WC powder (shown in Fig.1). 444 HM 57 B. Wittmann et al.

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6. Conclusion

The present investigation on WC grain growth in WC-Ni and WC-Fe hardmetals has revealed several new and important aspects which might render a more general understanding of the growth process.

Three main chemical parameters were elaborated to control WC grain growth under the experimental conditions chosen: • The chemical nature of the binder matrix (Fe, Co, Ni), which surrounds the growing (and dissolving) WC grains, • the carbon activity, which prevails during the sintering process in the respective alloy system (and which is determined by the alloy gross carbon content), • as well as the presence of growth inhibitors, such as, e.g. VC or TaC.

WC grain growth occurs to the least extent in iron binder alloys where it is also only slightly influenced by the alloy gross carbon content. By contrast, WC growth is most pronounced in WC-Ni alloys and a strong influence is performed by the carbon potential. Cobalt alloys are intermediate between, showing a moderate overall WC grain growth and a less pronounced influence of the alloy gross carbon content compared with WC-Ni. The strong influence of the carbon activity on overall WC grain growth can be explained by the fact that the carbon activity significantly co-determines the binder composition at temperature (both of the solid and liquid binder), and thus the interface chemistry of the suspended WC, which determines the rate of crystal growth. High carbon alloys exhibit a carbon-rich but tungsten-poor binder, which obviously enhances the WC grain growth; whereas, low carbon alloys exhibit a tungsten-rich and carbon-poor binder, which acts as a growth inhibiting environment compared with the former. The strong influence of the iron-binder metal itself (Fe, Co, Ni) on WC grain growth can be explained as a result of changes in the interface energy of the respective WC/binder interface (i.e. a change in the driving force) as well as in interactions between the binder atoms with the "growth active" carbon atoms (i.e. an increase in activation energy of the interface reaction). The latter hypothesis assumes the formation of stable metal/carbon clusters through bond formation which result in a depletion of "free available" carbon atoms necessary for the further growth of the WC crystals. This strong impact of the binder chemistry on WC grain growth supports the concept of an interface controlled growth mechanism, as assessed earlier from WC growth studies in the WC-Co system. B. Wittmann et al. HM 57 445_

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7. Acknowledgements

8. References

[1] L.Prakash: Proc. of the 13th Int. Plansee Seminar, Reutte, Austria, Vol. 2, eds. Bildstein, H. & Eck, R., pp. 80-109 (1993).

[2] US Patent 6,024,776: Cermets having a binder with improved plasticity; 15/02/2000.

[3] K.Hayashi; Y.Fuke; H.Suzuki: J. Japan Soc. of Powders and Powder Metall., 19,67-71 (1972).

[4] H. Suzuki; T.Yamamoto: J. Japan Soc. of Powders and Powder Metall., 15,68-71 (1968).

[5] A.M.Human; I.T.Northtrop; S.B.Luyckx; M.N.James: J. of Hard Materials, 2, 245-56(1991).

[6] S.Ekemar; L.Lindholm; T.Hartzell: Proc. of the 10th Int. Plansee Seminar, Reutte, Austria, Vol. 1, ed. Ortner, H.M., pp. 477-92 (1981).

[7] H.Holleck; H.KIeykamp: Int. J. of Refractory Metals & Hard Materials, 1, 112-16(1982).

[8] C.Wagner: Z. Elektrochemie, 65, 581 (1961).

[9] H.Fischmeister; G.Grimval: Sintering and Related Phenomena, ed. G.C. Kuczynski, Plenum Press, New York, pp. 119-49 (1980). 446 HM 57 B. Wittmann et al.

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[10] W.D.Schubert; A. Bock; B. Lux: Int. J. of Refractory Metals & Hard Materials 13, 281-296 (1995).

[11] H.E.Exner; H.Fischmeister: Arch. f.d. Eisenhüttenwesen, 37(2), 417-26 (1966).

[12] K.Choi; N.M.Hwang; D.-Y.Kim: Powder Metallurgy, 43(2), 168 (2000).

[13] K.Okada; T. Taniuchi; T.Tanase: Effect of Cr3C2 addition on the sintering of cemented carbide, International Conference on Powder Metallurgy & Particulate Materials, Las Vegas (1998).

[14] P.Gustafson: Metallurgical Transactions A, 18 A, 175-88 (1987).

[15] A.Gabriel; H.Pastor; D.M.Deo; S.Basu; C.H.AIIibert: Proc. of the 11th Int. Plansee Seminar, Reutte, Austria, Vol. 2, eds. Bildstein, H. & Ortner, H.M., pp. 509-25(1985).

[16] R.K.Sadangi; L.E.McCandlish; B.H.Kear; P.Seegopaul: Int. J. of Powder Metallurgy, 35, 27-33 (1999). M. Jilek et al. HM 65 447

15:h International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Hard PVD coatings nc-(Tii_xAlx)N/a-Si3N4 not only for dry machining. Mojmir Jilek, Pavel Holubar, Michal Sima

SHM, Ltd, Novy Malin, Czech Republic

Summary:

Nanocomposite coatings nc-(Tii.xAlx)N/a-Si3N4 prepared by PVD technology based on low voltage arc and using two cylindrical central cathodes, which were made from Ti and AISi, respectively, were studied for various combinations of evaporating cathode current on the both cathodes. All the depositions were done on the coating unit of a production size and the results were used in the industrial production. XRD, SEM-EDS, TEM resp. HRTEM methodes were used for structural analyses. GD OES method was used for composition analysis. The hardness measurements were done by means of the load-depth sensing technique using Vickers diamond indentor at a maximum load of 70 mN. The effect of the substrate bias on the stochiometry, crystallite size and mechanical properties of the deposited films was investigated. The crystallite size increases from 3 nm to more than 100 nm when the bias voltage decreases from - 200 to - 50 V and the titanium content increases from 35 to 90 at. %. The surface roughness decreases with increasing bias. The plastic Vickers hardness was in the range between 30 GPa to 47 GPa. The mechanical properties were also studied after the annealing at 800°C.

Keywords: Nanocomposite coatings, nc-(Tii.xAlx)N/a-Si3N4, cathodic arc

1. Introduction: The development of hard and superhard TiAISiN' coatings for machining applications, such as steel turning, drilling and milling begun at SHM company in 1993. It was motivated by the state of the art development of the TH.XAIXN coatings which were originally prepared by magnetron sputtering [1-3]. Several papers reported investigation into the dependence of the properties of these coatings on the Ti:AI ratio (e.g. [4-6]). Upon annealing of Th-xAlxN coatings at 500 °C a decrease of the hardness commences and at > 700 °C it becomes fast [7-9]. The work of Li Shizhi et al. [10] on the preparation of superhard TiSiN films inspired us to try if an addition of silicon 448 HM 65 M. Jilek et al.

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into the (Ti1_xAlx)N coatings can lead to an improvement of their mechanical properties. We have chosen the vacuum arc cathodic evaporation as a pilot technology for the deposition [11] and indexable inserts made of cemented carbide as standard substrates for cutting tools operating at high temperature which are expected in dry cutting at high speed. The initial experiments have shown a pronounced decrease of the crystallite size from several 100 nm in the single phase (Ti1_xAlx)N to about 20 nm for our TiAISiN which Vickers microhardness increased to about 30 GPa. Also the thermal stability of the coatings was significantly improved. In course of further development of the technology the central cathode for a simultaneous evaporation of all three elements Ti, AI and Si was improved by using two segments consisting of Ti and AISi solid solution with eutectic composition of 11.7 wt.% of Si. The ratio of Ti:AISi in the coatings was optimized by means of the control of the movement of the cathodic arc spot on both segments. The optimized 'standard' TiAISiN coating has a crystallite size of about 5 nm. Its Vickers microhardness of about 35 GPa is stable up to 850°C. In a collaboration with the Technical University Munich further information on the nature of these coatings was obtained [12,13]. In particular, the thermal stability against recrystallization could be increased to more than 1000 °C when the crystallite size was decreased to about 3 nm [13,14]. The crystallite size was verified recently by means of HR TEM analysis (made by SECO Tools and SANDVIK Coromant). The presence of amorphous Si3N4 was documented by a combination of X-ray diffraction (XRD) and X-ray photolelectron spectroscopy (XPS) [14]. Thus it was verified that the coatings represent another kind of superhard nanocomposites nc-(Ti1.xAlx)N/a-Si3N4 according to the generic design principle elaborated by S. Veprek et al. [15-18] and recently summarized in a review [19]. A subsequent development of the coatings and their deposition technology and collaboration with other research institutions and with Pramet Tools as an important toolmaker resulted in a further improvement [20-22] of the coatings towards a 'standard' TiAISiN single layer coating with trade mark MARWIN® SI. Furthermore we developed novel 'multilayer' coating on the basis of TiAISiN with a variable Ti content with trade mark MARWIN® MT. Both coatings MARWIN® were modified in 2000. In the present paper we shall describe the applied coating technology and the effect of the variation of the coating parameters on the properties of the nc-(Tii-xAlx)N/a-Si3N4 coatings in order to illustrate the technological possibilities. We shall confirm their nanocomposite structure and high thermal stability. . Jilek et a!. HM 65 449

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2. Experimental:

All coatings were prepared by the vacuum arc cathodic evaporation in a coating equipment of a production size which was designed and built by SHM. It consists of two central cylindrical cathodes made of titanium and AISi - eutectic alloy, respectively. The movement of the cathode spot of the vacuum arc is controlled by magnetic field in such a way that the evaporation from both cathodes occurs controllably. An independant control of both arcs enables us to prepare coatings with different stochiometry using the same electrodes. This design is different from that which we reported in [22]. The substrates are mounted on several fixtures which are placed symmetrically around the cathode and rotating. The typical deposition conditions are as follows: Substrate temperature of about 520 °C to 550 °C, negative bias of - 30 to -200 V, total arc current on the both electrodes of 270 A, current density on the substrates of 12 - 15 mA/cm2 and total pressure of 0.5 Pa (0.005 mbar). Typical coating thickness is about 3 (am. The coatings were investigated as deposited and after annealing to 800 °C for 30 min in 1 at. of pure N2. The hardness measurements were done at a load of 70 mN in order to assure that the indentation depth does not exceed 5 to 10% of the film thickness. All hardness values reported here are average values of at least 5 measurements on polished surface of the coating which was done in order to avoid artefacts due to surface roughness [22]. XRD measurements were performed with the Bragg-Brentano diffractometer working in the quasi parallel-beam mode (a long Soller collimator with the opening 0.4° and a curved graphite monochromator were inserted into the diffracted beam). The angle of incidence of the primary beam was kept constant at 5°, which has reduced the penetration depth of the Cu Ka radiation (X - 1.5418 x 10"10 m) to approximately 0.9 jum. A quantitative depth analysis of chemical composition was done by means of optical spectroscopy with glow discharge excitation (GDOES) on the LECO SDP 750 instrument.

3. Results and discussion:

We used four different values for both, the bias and the arc current ratio on the two cathodes made of titanium and AISi eutectic alloy. This yielded 16 different coatings with different composition, structure and properties as summarized in table I and II. The crystallite size was obtained from the physical broadening of diffraction lines using the Williamson-Hall plot, i.e., as the reciprocal value of the intercept of the linear dependence of physical 450 HM 65 M. Jilek et al.

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broadening on sin 8 with the y-axis. The physical broadening was taken from the diffraction line broadening after the correction for instrumental effects (LaB6 from NIST was used as external standard). One can see that the composition of the coating is influenced mainly by the arc current ratio. The effect of the bias being relatively small. Thus, the somewhat different color of the samples is due to a different morphology.

Table I Composition of the coatings Ti + AI + Si = 100%. Total arc current l(Ti)+l(AISi) = 270 A.

Coating Arc current Bias Crystallite Ti content Al + Si number ratio (V) size (nm) (at.%) content l(Ti)/l(AI+Si) (at.%) 1 0,6 30 34,48 65,52 2 0,6 50 4,8±1 36,92 63,08 3 0,6 100 34,64 _ 65,36 4 0,6 200 3,2±0,4 34,58 65,42 11 1,0 30 55,59 44,41 12 1,0 50 56,71 43,29 13 1,0 100 11+5 55,19 44,81 14 1,0 200 7,9+1,8 55,47 44,53 21 1,75 30 76,62 23,38 22 1,75 50 84,01 15,99 23 1,75 100 18±7 76,69 23,31 24 1,75 200 12±1 77,97 22,03 31 3,3 30 91,35 8,65 32 3,3 50 150±50 89,45 10,55 33 3,3 100 92,64 7,36 34 3,3 200 32±8 94,87 5,13

Table II summarizes the values of the hardness and elastic modulus measured after the deposition and after annealing at 800°C in pure nitrogen for 30 minutes. M. Jilek et al. HM 65 451

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Table II Hardness and elastic modulus of the coatings

Coating As deposited Annealed at 800°C, 30 min number HVp|ast 70 mN Elastic HVpiast 70 mN Elastic (GPa) modulus (GPa) modulus (GPa) (GPa) 1 37 364 42 377 2 36 347 45 393 3 37 350 48 394 4 35 350 34 348 11 46 480 46 504 12 47 461 53 491 13 47 455 - - 14 43 427 44 435 21 44 508 48 552 22 42 517 47 560 23 46 529 53 561 24 47 491 56 519 31 35 532 37 521 32 40 534 40 547 33 37 525 37 538 34 32 510 33 521

From table II it is clear that there is no significant dependence of the hardness and elastic modulus on the bias for the as deposited coatings, but there is a tendency of hardness to increase after the annealing, in particular for coatings with higher fraction of AI + Si (current ratio 0,6). For the coatings with the highest Ti content (current ratio 3,3) the hardness does not change after the annealing. The optimal high hardness is reached for coatings with an Al+Si content in the range of 25 - 45 at.% (current ratio 1,75 - 1). The maximum elastic modulus is reached for coatings with a Al+Si content in the range of 10 - 20 at.% (current ratio 3,3 - 1,75). There is a pronounced dependence of the crystallite size on the Al+Si content and bias: The higher Al+Si content and higher the bias the lower the crystallite size. This can be the reason for the different color of the coatings mentioned above. 452 HM 65 1. Jilek et al.

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Our commercial coating process uses bias of- 200 V. Therefore the effect of the composition of the coatings on the residual stress was studied only for this bias. Figure 1 shows the measured residual stress as obtained from the XRD experiments. Residual stress and the "stress-free" lattice parameters were calculated from the well-known sin2v|; plot [23,24]. The compressive biaxial stress initially increases with decreasing Al+Si content (increasing the current ratio to 1) but decreases again when the Al+Si content decreases below approx. 20 at.% when the highest hardness is reached.

Bias 200V

16 14 12 CL 10 < CD Q 6 -• - Stres s 4 2 f \ 0 L 3,3 1,75 1 0,6

' arc Ti'i arc Al+S

Fig.1 Compressive biaxial stress in the coatings vs. composition (arc current ratio) for Bias - 200V.

From the application point of view it is necessary to study the influence of coating parameters on the resulting surface roughness as shown in Fig. 2. It is seen that higher ion energy has a positive influence on the surface rouhgness. This is due probably to a lower content of macroparticles because of the high ion bombardment. Lower surface roughness in the case of higher Al+Si content is caused by the known fact that the alloys based on AI are evaporated with lower content of macroparticles in the nitrogen atmosphere 1 Jilek et al. HM 65 453

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than titanium. While in the case of titanium the evaporating arc is divided into the cathode spots with current 80 - 120 A in the case of AI alloys this cathode spot current is lower.

Ra (jam)

Fig.2 Surface roughness Ra (urn) vs. the bias and composition of the coatings (arc current ratio).

Figures 3 and 4 show HRTEM micrographs of the nanocrystalline structure of the multilayer coating nc-(Ti-|.xAlx)N/a-Si3N4. These images confirm the nanocrystalline structure and the small crystallite size as determined from the XRD analysis. The multilayer coating consists of the two different layers corresponding to the experimental ones mentioned in the tables I and II with numbers 14 and 34. The layer with higher amount of Ti has crystals with higher size. The sub-multilayer structure of the layer 14 is the result of the substrate rotation. 454 HM 65 M. Jilek et al.

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Fig.3 SAED Ti(AI,Si)N multilayer, FCC, (L=20cm) (by courtesy of Sandvik Coromant)

Fig.4

HRTEM images of multilayer nc-(Tii-xAlx)N/a-Si3N4 coating, (by courtesy of Seco Tools) 1 Jilek et al. HM 65 455_

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4. Conclusions:

Nanocomposite coatings nc-(Ti-|.xAlx)N/a-Si3N4 were prepared by modified PVD process with higher deposition rate. A decrease of the surface roughness with increasing bias voltage was found. There is also a well controllable relation between the biaxial compressive stress and the composition of the coatings. The crystallite size decreases with increasing substrate bias and with increasing fraction of Ti. This enables us to optimize the tribological properties of the coating by an optimum choice of their composition and process parameters. The coatings have also a good thermal stability which, together with their high hardness and low surface roughness makes the nc-(Ti1.xAlx)N/a-Si3N4 nanocomposite coatings very suitable for dry machining.

5. Acknowledgement:

We would like to thank Prof. Dr. Stan Veprek and his colleagues from the Technical University Munich for valuable discussions, Swedisch companies Sandvik Coromant and Seco Tools for their structural analysis, Pramet Tools Sumperk for the cutting tests, Dr. D. Rafaja (Charles University in Prague) for the XRD measurements and Mr.M.Vnoucek (West Bohemian Univ.) for GD OES analysis. This work was supported by the NATO under the SfP Project 972379, by the Czech Ministery of Education and Youth under the KONTAKT Project ME302 and by EC under the Fifth Framework Programme Project No. G1RD-CT- 2000-00222.

References: [1]W.-D. Münz, J. Vac. Sei. Technol. A4 (1986) 2717. [2]O. Knotek, M. Böhmerand T. Leyendecker, J. Vac. Sei. Technol. A4 (1986)2695. [3] H. A. Jehn, S. Hofmann, V.-E. Rückborn and W.-D. Münz, J. Vac. Sei. Technol. A 4 (1986) 2701. [4]Y. Tanaka, T. M. Gür, M. Kelly, S. B. Hagström, T. Ikeda, W. Wakihira and J. Satoh, J. Vac. Sei. Technol. A 10 (1992) 1749. [5]T. Ikeda, H. Satoh, Thin Solid Films 195 (1991) 99. [6] A. Kimura, H. Hasegawa, K. Yamada, T. Suzuki, Surf, and Coat. Technol. 120-121 (1999)438. 456 HM 65 M. Jilek et al.

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[7]D. T. Quinto, G. J. Wolfe and P. C. Jindal, Thin Solid Films, 153 (1987) 19. [8]O. Knotek, R. Elsing, M. Atzorand H. G.Prengel, in: K. C.Ludema (ed.), Proc.Wear of Materials Conference, April 9-13, Denver, Colorado, 1989, pp. 557. [9] E. Vancoille, J. P. Celis, J. R. Roos, Thin Solid Films, 224 (1993) 168. [10] Li-Shizi, Shi Yulong, Peng Hongrui, Plasma Chem. and Plasma Processing 12 (1992) 287 . [11] I. I. Aksenov, V. G. Bren, V. G. Padalka, B. M. Choroschich, J. Tech. Phys., 48 (1978)1165. [12] S. Veprek, P. Nesladek, A. Niederhofer, F. Glatz, M. Jilek and M. Sima, Surf. Coat. Technol. 108-109(1998) 138. [13] A. Niederhofer, P. Nesladek, H.-D. Männling, K. Moto, S. Veprek and M. Jilek, Surf. Coat. Technol. 120-121 (1999) 173. [14] S. Veprek, T. Bolom, P. Nesladek, M.Jilek, Technical University of Munich, private communication (2000), to be published. [15] S. Veprek, S. Reiprich and Li Shizhi, Appl. Phys. Lett. 66 (1995) 2640. [16] S. Veprek and S. Reiprich, Thin Solid Films 268 (1996) 64. [17] S. Veprek, M. Haussmann and S. Reiprich, J. Vac. Sei. Technol. A 14 (1996)46. [18] S. Veprek, M. Haussmann, S. Reiprich, Li Shizhi and J. Dian, Surf. Coatings Technol. 86-87(1996)394. [19] S. Veprek, J. Vac. Sei. Technol. A 17 (1999) 2401. [20] P.Holubar, M.Jilek, M.Sima, Surf. Coat. Technol. 120-121 (1999) 184. [21] P.KIapetek, M.Jilek, M.Sima, EURO PM99, EPMA, Turin, November 1999, p.525. [22] P.Holubar, M.Jilek, M.Sima, Surf. Coat. Technol. 133-134 (2000) 145-151. [23] D. Rafaja, V. Valvoda, R. Kuzel, A.J. Perry, J.R. Treglio, Surf. Coat. Technol., 86-87 (1996) 302-308. [24] D. Rafaja, V. Valvoda, A.J. Perry and J.R. Treglio, Surf. Coat. Technol. 92(1997)135-141. R. Weissenbacher et al. HM 79 457

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Deposition of BCN Layers from B-C-N-H Single Source Precursors

Ronald Weissenbacher, Roland Haubner

Vienna University of Technology, Vienna, Austria Institute for Chemical Technology of Inorganic Materials

Summary:

This work deals with the deposition of BxCyN2 phases by decomposition of a B, C, N containing single source precursor called Tris(dimethylamino)borane in a hot-filament CVD apparatus. In contrast to the deposition of diamond layers the deposition of cubic-BN or cubic-BxCyNz phases with low pressure CVD-methods was not successfully until now. To avoid the problems of filament failure and deposition of Ta into the produced layers, the Ta-filaments were pre-treated by carburization or nithdation. BCN layers were deposited on Si and on WC-Co substrates in a hydrogen atmosphere similar to diamond deposition. Layers deposited on Si showed higher growth-rates than layers deposited on WC-Co at the same deposition conditions. An increase of the filament temperature led to an increasing growth rate. This increase of growth rate was caused by higher BCN decomposition but also by additional Ta incorporation. At filament temperatures higher than 220CPC Ta evaporated much easier from the filament. An increase of the precursor flow rate as well led to an increase in the deposition rate on both types of the investigated substrates. Layers produced at a precursor flow rate of 1,7*10"5 mol/min showed some weak peaks in Raman and IR spectra, but none of the phases that exist in the ternary BCN system could be clearly identified. The morphology of the layers was investigated by SEM.

Keywords: CVD, Coatings, BCN, BN 458 HM 79 R. Weissenbacher et al.

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1. Introduction Boron, nitrogen and carbon are relatively small atoms, this fact results in short bondage lengths and high bonding energies (1). This is the reason for high hardness of the phases that exist in the ternary BCN-system, like diamond and cubic boron nitride called superhard phases, or boron carbide as a hard material (2). Cubic BxCyNz phases, which are expected to have high hardness, can be seen as an intermediate link between diamond and c-BN. The deposition of diamond at low pressure has become an industrial process (3), but no low pressure process is known for the deposition of high quality c-BN crystals in p.m size. Nano-crystalline c-BN can be grown by PVD methods (4) or ion-assisted CVD (5) at low pressure. Cubic BxCyNz phases are formed by shock wave synthesis (6) or by static high pressure and high temperature techniques (7). The lattice constant of c-(BN)xC1.x phases determined by x-ray diffraction is in between those of diamond and c-BN (6,7), therefore these phases can be considered as solid solutions of carbon andBN. Studies on the phase stability of c-BN and h-BN showed that at standard conditions c-BN is the stable phase and h-BN the metastable one (8,9). However the exact transition temperature between c-BN and h-BN (10) as well as exact thermodynamic data are still lacking. In this work the decomposition of Tris(dimethylamino)borane, which is a BCN- containing single source precursor, and deposition of BCN layers in a hot- filament CVD apparatus were investigated. The deposition experiments were done in hydrogen atmosphere similar to the one used for the deposition of diamond layers. The use of a single source precursor was preferred because the B:C:N ratio is fixed.

2. Experimental Setup and Characterization of the Samples 2.1. The Apparatus The deposition of BCN-layers by decomposition of the single source precursor B(N(CH3)2)3 was carried out in a hot-filament CVD apparatus. The reaction chamber was a quartz tube with Viton sealed stainless flanges. As a hot filament a coiled Ta wire (6 mm inner diameter, 4 cm heated length) was used for the activation of the gas phase (11). For the heating of the filament and to keep the filament temperature constant a power constanter was used. The filament wire was fixed between two Mo electrodes and held in place with a hook in the middle, which guaranteed a constant distance between filament R. Weissenbacheretal. HM 79 459

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and substrate above the entire substrate area. Filament temperature was measured by a two-colour pyrometer (IRCON Mirage OR1500).

2.2. Substrates and Chemicals As substrate materials Si-wavers and WC-Co hardmetal inserts (K20 SPGN WC 6% Co) were used. Hydrogen (quality 5.0) was used as reaction gas and argon (quality 5.0) as carrier gas for the precursor. The precursor flow, which was empirically quantified, was adjusted by the flow rate of the carrier gas. As B-C-N containing single source precursor, Tris(dimethylamino)borane was used. A single source precursor was used because of the fixed B:C:N ratio. Literature data for the vapour pressure of Tris(dimethylamino)borane were contradictory (12,13) and therefore the precursor flow rate was empirically investigated for different Ar flows - from 0 to 50 seem - at a evaporation pressure of 15mbar and a temperature of 20°C. Precursor flow rates between 1,01*10"5 and 1,50*10A mol/min Tris(dimethylamino)borane were used for the deposition experiments.

2.3. Analytical Investigations The layer deposition rate was measured by gravimetric analysis before and after the deposition ([mg crrf2h~1]). The morphology of the layers was investigated by scanning electron microscopy (SEM JEOL 6400). X-ray diffraction (CuKa radiation) was carried out to determine the phase composition of the layers. Infrared spectroscopy (Biorad FTS 175 with a microscope) was used to characterize the BN deposits.

2.4. Filament Pre-treatments Using a Ta filament for the deposition of BCN layers two main problems occurred. Primarily at temperatures higher than 2000°C the reaction of the B-C-N containing precursor Tris(dimethylamino)borane with the Ta wire led to the formation of a liquid phase as consequence of the eutectic reaction at 2055°C in the Ta-B system (14). Secondly the BCN-layers were contaminated with Ta, which evaporated from the filament surface. Therefore filament pre-treatments were carried out to increase their lifetimes and to avoid or at least reduce the Ta contamination into the produced layers. As an alternative to the pre-carburization (15,16) of the filament pre-nitridation was investigated. In a nitrogen atmosphere (12 mbar, N2-flow 50 seem) the Ta-filament was heated up to a temperature of 2000°C and heating power was kept constant. Within five minutes an abrupt temperature drop of 500°C (from 2000°C to 1500°C) could be observed (Fig. 1), which started at the 460 HM 79 R. Weissenbacher et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 outer twists of the filament and moved continuously to the middle twist within 60 s (Fig. 2). Because the electrical heating power was kept constant, this effect could not be described by simple changes in the conductivity of the filament (17). XRD measurements showed a change in the lattice constant of Ta shortly after the temperature drop as a result of the dissolution of nitrogen in the Ta lattice (17). After 15 minutes of nitridation time Ta2N could be found as an additional phase, which could be expected from the phase diagram (18). From metallographic investigations a change in surface roughness of the filament before and after the temperature drop could be observed (Fig. 1). Cross sections of the filament showed the growth of a second phase in a concentric circle when nitridation time was increased.

2000- Temperature drop of the filament C1900 measured on one fixed point of the £1800 filament with a two-colour 21700 0 pyrometer (IRCON Mirage OR15). |"1600 01 •"1500

60 80 100 120 140 160 time in s surface morphology before and after temperature drop

' - •• 4.-

Fig. 1: Temperature drop of the filament during nitridation and change in

surface morphology (TFN.= 2000°C, p = 12 mbar, N2-flow 50 seem) R. Weissenbacher et al. HM 79 461

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temperature distribution during if- conversion

0-120 sec. r t

135 sec.

I 150 sec.

165 sec.

Fig. 2: Temperature drop of the filament during nitridation (TV«.= 2000°C, p = 12 mbar, N2-flow 50 seem)

3. Results and Discussion

3.1. Hot Filament Deposition of BCN-Layers Deposition experiments were carried out in order to investigate the influence of the substrate material, the filament temperature and the precursor concentration on the deposition rate and morphology of BCN-layers in hydrogen atmosphere.

3.1.1. Influence of the substrate material on the deposition rate The substrate materials Si and hardmetal (WC-Co) were used for these experiments. The deposition rate at equal deposition parameters (filament temperature of 2000°C, substrate temperature 740°C, precursor flow rate of 462 HM 79 R. Weissenbacher et al.

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0,15 mmol/min, H2-flow of 250 seem and pressure in the reaction chamber of 15 mbar) was significantly higher on Si (0,34 mg*cm"2h"1) than on WC-Co (0,18 mg*cm"2h"1) after 18 h of deposition (Fig. 3). SEM investigations showed that the morphology of the layers was similar on both substrates (Fig. 4). Ball-like structures were found with about 10 (im diameter. XRD measurements showed no additional peaks beside the peaks of the substrate material. This indicated that the deposited layers are of amorphous structure. Raman spectroscopic examinations did not reveal any Raman active vibrations. Also IR spectroscopy showed no characteristic peaks on such BCN layers.

WC-Co Si substrate material

Fig. 3: Deposition rate on different substrate materials (TFi|.= 2000°C, Tsub.= 740°C precursor 0,15 mmol/min, H2-flow 250 seem, 18 h).

3.1.2. Influence of the filament temperature The filament temperature was varied between 1800°C and 2200°C (at a reaction pressure of 15 mbar, a hydrogen flow of 250 seem and a precursor flow rate of 1,5*10"1 mmol/min). Changes in the filament temperature always resulted in a change of the substrate surface temperature, which increased from 620 to 890°C with increasing filament temperature (the distance between filament and substrate was kept constant at 1,4 cm during all the experiments). R. Weissenbacher et al. HM 79 463

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WC-Co

100 urn

H 10 jam

Fig. 4: Morphology of BCN-layers deposited on Si and WC-Co by decomposition of Tris(dimethylamino)borane (TRi.= 2000°C, TSub= 740°C, precursor 0,15 mmol/min, H2-flow 250 seem, 18 h).

Increasing filament temperature resulted in an increase in deposition rate on the Si substrates (Fig. 5). At 1800°C and 2000°C a non-faceted, flat surface containing sperolitic balls was observed. The morphology was similar at both experiments. At 2200°C the layers were flat and only sporadically ball like structures could be observed. On these samples also sharp needles with a size of about 5 |iim were irregularly distributed (Fig. 6) (19). EDX analysis showed, that the layers were contaminated with Ta that was evaporated from the filament at high temperatures (20). The tantalum evaporation from the filament increased with increasing filament temperature. The atomic weight of tantalum is relatively high compared to B-C-N and therefore the higher deposition rate at 2200°C is also a result of Ta-B-C-N co- deposition. 464 HM 79 R. Weissenbacher et al.

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T 0,6 -r

1700 1900 2100 2300 filament temperature °C

Fig. 5: Influence of the filament temperature on the deposition rate on Si (TSub= 620 - 890°C, precursor 0,15 mmol/min, H2-flow 250 seem, 18 h).

filament temperature 1800°C 2000°C 2200°C_

••^.

110 jam Fig. 6: Morphology of BCN-layers deposited on Si substrates at different temperatures (TSub= 620 - 890°C, precursor 0,15 mmol/min, H2-flow 250 seem, 18 h). R. Weissenbacher et al. HM 79 465

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3.1.3. Influence of the precursor flow rate To investigate the influence of the precursor flow rate on the deposition rate experiments at a filament temperature of 2000°C, 15 mbar reaction pressure and a gas flow of 500 seem H2 were carried out. Precursor flow rates of 1,0*10"2, 1,7*10~2 and 2,4*10~2 mmoi/min were used. Deposition experiments with different precursor flow rates were done on Si as well as on WC-Co substrates. Both showed an increase of the deposition rate with increasing precursor flow rate, but deposition rates were higher on Si than on WC-Co as mentioned above (Fig. 7). Differences in the morphology of the layers have also been observed.

0,25 -i

0,01 0,015 0,02 0,025 precursor flow rate mmol/min Fig. 7: Dependence of the deposition rate on the precursor flow rate on o WC-Co and Si substrates (TR|.= 2000°C, Tsub=740 C, 18 h, H2-flow 500 seem).

At a precursor flow rate of 1,0*10'2 mmol/min a weight gain of 0,2 mg was measured after 18 hours of deposition which corresponded to a deposition rate of 0,002 mg*cm"2h'1. This very little weight-difference might as well be caused by uncertainties in gravimetric analysis. Probably no layer was deposited under this conditions. Depositions on WC-Co showed non-faceted layers with spherolitic ball like structures as mentioned above. The surface of the layer was rougher at lower precursor flow rates (Fig. 8). A reason for the finer structure might be a higher nucleation rate at higher precursor flow rate. 466 HM 79 R. Weissenbacher et al.

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On Si the deposited layers looked different from the ones on WC-Co. At 1,7*1CTZ mmol/min also ball like structures could be observed, but for some reason there were holes in the layer and surface roughness was higher than on WC-Co substrates. The layer consisted of different levels of ball like structures, and it seemed that it grew layer by layer (Fig. 8). At 2,4*10"2 mmol/min a continuous layer with relatively rough surface was deposited. The ball like structure observed on WC-Co is not so different from the one on Si, but the layer looked very homogeneous. The differences between the two substrates could be caused by the substrate composition. In case of WC-Co the cobalt could influence the deposition. It is known that Co could diffuse into coatings and react with boron precursors to borides (21). Additionally the influences of Co on the decomposition process and the growth of BCN-layers is not cleared now.

precursor flow rate 1,7*10"5mol/min 2,4*1 CTmol/min

WC-Co

10 jim

Fig. 8: Morphology of BCN-layers deposited on WC-Co and Si substrates at different precursor flow rates. (TFii.= 2000°C, TSub= 740°C, H2-flow 500 seem, 18 h) R. Weissenbacher et al. HM 79 467 15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

3.2. Raman and Infrared Spectroscopy Raman spectroscopic investigations of the layers deposited at a precursor flow rate of 2,4*10"5 mol/min did not reveal any Raman activities (Fig. 9a), no matter they were deposited on Si or on WC-Co. Depositions at a precursor flow rate of 1,7*10~5 mol/min showed some Raman activity on Si as well as on WC-Co (Fig. 9b), although no sharp peaks could be detected. Some phases in the B-C-N system (h-BN, micro-crystalline graphite or B4C) show peaks in this area, but they are not sharp enough to be identified. Further investigation would be necessary to identify the deposited phases. IR measurements showed similar diagrams for all investigated layers (Fig. 10). A typical double peak was observed in all diagrams at 2358 cm"1, a broad peak around 1600 cm"1. Both peaks could not be identified up to now.

For example some typical IR peaks are: h-BN 828, 1380, 1610 [cm'1] c-BN 1065, 1340 [cm"1] B4C 1100, 1590 [cm"']

precursor flow rate 2,4*10"0 mol/min 1,7*10"0 mol/min 8000 - 8000-

6000 - 6000 - _>*

4000 - jj 4000 • |c 2000 - 2000 -

0 800 1000 1200 1400 1600 1800 800 1000 1200 1400 1600 1800 RAMAN-Shift [cm"1] RAMAN-Shift [cm"1] Fig. 9: Typical Raman spectra of layers deposited on WC-Co or Si at precursor flow rates of 1,7*10"5 mol/min and 2,4*10"5 mol/min (TFil.= 2000°C, TSub= 740°C, 18 h) 468 HM 79 R. Weissenbacher et al.

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c o

800 1200 1600 2000 2400 2800 3200 3600 wavenumber [cm"1] Fig. 10: Typical IR spectrum of the deposited BCN-layers

(TF|i.= 2000°C, Tsub= 740°C, 18 h).

4. Conclusions The deposition of superhard c-BN or c-BCN layers as wear resistant coatings on tools is an interesting field of research. Also for industrial production CVD methods would be of great interest. However, the low pressure deposition of cubic-BxCyNz layers with crystals in jim size is not possible until now. There is not enough information available about BxCyNz deposition and therefore some basic research is necessary. In this work the decomposition of Tris(dimethylamino)borane in hydrogen atmosphere was investigated. To activate the gas atmosphere, experiments in a hot-filament apparatus were carried out. The experiments showed, that the layers morphology and the growth rates can be influenced by filament temperature, the precursor flow rate and the substrate material (WC-Co and Si). However, the deposited layers are mainly of amorphous structure because no hint for a crystallinity was observed. Weak peaks were found with Raman and IR spectroscopy, but none of the phases existing in the ternary B-C-N system could be clearly identified. Nevertheless various types of BCN coatings could be deposited. The layers composition, its hardness and other properties will be examined in the next future. Further work with other precursors or other deposition parameters is necessary to investigate more details about reaction conditions and the analysis of the deposited products will be improved. R. Weissenbacher et al. HM 79 469^

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5. Acknowledgements This work was supported by NEDO (New Energy and Industrial Technology Development Organization) and JFCC (Japan Fine Ceramic Center) as a part of the FCT (Frontier Carbon Technology) project promoted by AIST, MITI; Japan.

References

(1) H. Ehrhardt: Surf. Coat. Technol. 74/75 (1995) 29-35 (2) A. Bartl, R. Haubner, B. Lux: Refractory Metals & Hard Materials 14 (1996)145-157 (3) "Low Pressure Synthetic Diamond", B. Dischler and C. Wild, Springer Verlag, Berlin, 1998 (4) T. Yoshida: Diamond and Related Materials 5 (1996) 501-507 (5) M. Kuhr, S. Reinke and W. Kulisch: Surf. Coat. Technol. 74-75 (1995) 806-812 (6) Y. Kakudate, M. Yoshida, S. Usabe, H. Yokoi, S. Fujiwara, M. Kawaguchi, K. Sako and T. Sawai: Trans. Mat. Res. Soc. Jpn. 14B (1993) 1447-1450 (7) A. R. Badzian: Mat. Res. Bull. 16 (1981) 1385-1391 (8) V. L. Solozhenko and V. Ya. Leonidov: Russian J. of Physical Chemistry 62(11) (1988), 1646-1647 (9) H. Sachdev, R. Haubner and B. Lux: Diamond and Related Materials 6 (1997)286-292. (10) G. Will, G. Nover, J. von der Gönna: Proc. of AIRAPT-16 Meeting (Kyoto, Japan aug. 22.-30., 1997) (11) R. Brunsteiner: Ph. D. Dissertation, University of Technology Vienna (1993) (12) Gmelins Handbuch der Anorganischen Chemie, B, Borverbindungen 4 (Bd 22), Springer Verlag, 1975, S. 5-44 (13) W. Kalss: Ph. D. Dissertation, University of Technologie, Vienna (1998) (14) T. B. Massalsky et al, "Binary alloy phase diagrams", ASM International, 1990, Vol. 1 (15) R. Haubner, B. Lux: Diamond and Related Materials 2 (1993) 1277- 1294 (16) S. Okoli, R. Haubner, B. Lux: Surface and Coatings Technology, 47 (1991)585-599 470 HM 79 R. Weissenbacher et al.

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(17) R. Weissenbacher, R. Haubner, K. Aignerand B. Lux: Diamond and Related Materials (submitted) (18) H. Wiesen berger, W. Lengauer, P.Ettmayer: Acta Mater, Vol 46/2, (1998)651-666 (19) W. Kalss, R. Haubner, B. Lux: Refractory Metals & Hard Materials 16(3), (1998)233-241 (20) K. Aigner, X. Tang, R. Haubner, B. Lux New: Diamond Front. Carbon Technol. (2000), 10(4), 181-189. (21) W. Kalss, S. Bohr, R. Haubner, B. Lux, M. Griesser, H. Spicka, M. Grasserbauer, P. Wurzinger: Refractory Metals & Hard Materials 14 (1996)137-144 EUROPEAN FUSION DEVELOPMENT AGREEMENT § §" w CD ! I ITER and the Fusion Reactor: j Status and Challenge to Technology *I K.Lackner, EFDA-Garching based on material provided by H. Bolt 1>, D. Campbell 2>, M. Gasparotto 2>, A. Möslang 3> 1) IPP-Garching, 2) EFDA-Garching, 3) FZK Karlsruhe

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15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 478 RM 1 K. Lackner et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. <

v- 00 CO O o o O5 CO 00 UO a r^ o CO CNJ CO ^ io in o 00 O en CO TJ- T- CN O ^ T- o o ^~ -c- o" Al on" CM CO CD 1.7 0 0.3 3 ITE R

"ST CO (m ) F ? £ i O CM II a> E i. ro S" j— c X CO 'x X Q. F A" CD ro TO ai — A aj 1— CL rv" Ct > CO c V a J a y ^r SEFDA EUROPEAN FUSION DEVELOPMENT AGREEMENT R & D for ITER : the 1-7

Resource Allocation Summary for the Seven Large R&D Projects

(Unit: klUA)

Projects EU Japan RF US* Total LI - Central Solenoid Model Coil 10 61 4 22 91 L2 - Toroidal Field Model Coil 40 0 0 1 41 L3 - Vacuum Vessel Sector 4 19 4 2 29 L4 - Blanket Module 29 14 12 9 64 L5 - Divertor Cassette 13 12 9 21 55 L6 - Blanket Module Remote 3 18 0 o 21 Handling L7 - Divertor Remote Handling 26 3 0 0 29 Total 125 127 29 55 336

' US contributed until July 1999

Status: June 2000 The 1B$ ITER design effort and the 0.4 B$ spent on dedicated component development have produced a solid fundament and are a highly tangible asset of the ITER-project

45«. •vl CO 4^ OD O

EUROPEAN FUSION DEVELOPMENT AGREEMENT

R&D - Divertor Cassette (L-5) (3)

CfC monoblock and W brush armoured vertical target (EU) W and Be armour fast Pure Cu-clad DSCu tube brazing to liner CuCrZr armoured vertical target heat sink (RF) with saddle block CfC and CVD-W armours (JA). ct> K. Lackner et al. RM 1 481

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

C o c 0) • • "ü TJ CD O SZ O LU it Q. H T3 c CO 5 CD 1 o 2 o Ü r- CO c _ CD i9 — CD JK CL 1 CD b 482 RM 1 K. Lackner et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reirtte (2001), Vol. <

Q. O 8 -O en 't/5 SS a> (1) a. T3 o

G CO ! .o co •a a. o o CO in o o z Q.

LU 03

Ü t ü CD 2 QL a) Q. c o O z c cö CD w o •o

CO 0 Q. .Q Q. 00 z o to g CD 0 CO 1 a. D a o E a. o CO O c: O g Ü

t > CO c - o o o Ü

^ k ^ j LL SEFDA F.UROPEAN FUSION DEVELOPMENT AGREEMENT Challenges for ITER & Fusion Power Plant I

(believed) solved on ITER: heating, superconductivity, remote handling;

Physics i f 30 • provide steady state operation (non-inductive current drive, | stellarator) § X • soften technology requirements (heat loads, disruptions, tritium I inventory) -| Technology: material problems i • high heat flux components(divertor target plates, first wall) I > • radiation tolerant, low activation structural and functional materials | (blanket) |

4 CO CO 00

HEFDA EUROPEAN FUSION DEVELOPMENT AGREEMENT First Wall and Divertor: materials in a hostile requirement

first wall: tritium inventory issue: modest flux of high energy •tritium in plasma vessel neutral particles (» 100s eV), must be kept low for safety

chemical erosion 'sputtered, desorbed, 30 evaporated material co- bulk plasma: redeposits with high 3 hydrogen fraction cannot tolerate impurities (< _j W, 10-2 Be, 5x10-2 He)

divertor target piates: high heat flux (without radiation protection: < 60 MW/m2, withr.p.-^10MW/m2) > a impulsive heat loads: ELMs, disruptions

7T CD CD. 01 EUROPEAN FUSION DEVELOPMENT AGREEMENT HEFDA CD CD Main technological Problem for ITER: plasma facing materials GAS PUFFING PFCs Area 2 issues ror ine ÜHJ öivenor isreeis (m ) are IW 300 »Erosion lifetime (chemical 380 sputtering), T-co-deposition (450g T ow limit inside the vessel) and C dust L 10 production (~200kg limit) B 50 Hssues for the Be FW are D 30 •Be dust production (100kg limit) in VT 55 'Beryllium particular on hot surfaces (6kg limit) - Liner 60 Tungsten > Hydrogen production in off normal Carbon-Fibre Composites events - explosion ? «Issues for W clad divertor targets are •W dust production (100 - 300kg limit) causing a radiological hazard in case of a by-pass event »Disruption erosion - melt-layer loss «Plasma compatibility Present preference: 3 different materials

4 00 (Ji GO CD

•Bg!| r~ r^ |-\ A EUROPEAN FUSION DEVELOPMENT AGREEMENT |

Structural and Functional Materials for Fusion Power Plants j CD CD CO requirements: I resilience to neutron flux/fluence s high power density * high availability (interval for replacements) f low activation | activation is only source of radioactive waste I

(reaction product He; T is intermediate product) i good thermomechanical properties up to high temperatures f determine thermodynamic plant efficiency I (chemical compatibility with breeder material) I

> 1'Tt.R can be constructed with existing (conventional) materials a (3 - 10 dpa integrated iiuence) but will test DEIV1O blanket module inserts 1 | — CD EUROPEAN FUSION DEVELOPMENT AGREEMENT 77 SEFDA CD 2. 0) 12.5 MWa/m2 Bestrahlung

Cfi CD

Material for DEMO n—MANET-II o—OPTIFER & First Generation A— EUHOFER-87 v— F82H-mod Power Plants: O EUROFERrai ferritic steels

(remote controlled) recycling of dominant fraction of waste is a primary target

10'1 10D 101 1OZ 103 10" 105 108 Time after Irradiation [y]

1 00 00 00

EUROPEAN 1-USION DEVELOPMENT AGREEMENT

300 EUROFER 97 Heat E 83698 250 compared to previous (F82H-mod):

200 reduction of ductile- CD 'W c brittle transition LU temperature by 50° 150 F82Hmod. Heat 9741: (O,= 124 ppm) .Q (T-L), 1040°C 30' + 750 °C 1h i O w 100 EUROFER: (O,= 10 ppm) 0 (T-L), 980°C27' + 760°C 1.5h 50 £ (T-L), 1050°C 30' + 750°C 3h - m (T-L), 1075°C 30' + 750°C 2h a (L-T), 1075°C 30'+ 750°C 2h D 0 3)

-120 -100 -80 -60 -40 -20 0 20 40 60 80 100 ill O Test Temperature [°C] CD SB O 7C EUROPEAN FUSION DEVELOPMENT AGREEMENT CD SEFDA 2- Radiation effect on DBTT 0)

300 10 and 15 dps- existing neutron sources Neutron irradiation (high flux reactors): 250 T. = 300 °C wrong spectrum - fusion 200 produces 14 MeV neutrons (ratio of displacement to He-production) insufficient fluxes

extrapolation and modelling promise 150 dpa feasiLft forEDRDFER-liB materials

-100 12 3 4 5 6 Damage dose [dpa] > a dedicated test facility (stripping source) for fusion needed: IFMIF (International Fusion Materials Irradiation Facility)

00 -ts. CD O SEFDA EUROPEAN FUSION DEVELOPMENT AGREEMENT Improvements beyond (mainstream) ferritic steels I: ODS

1000

900 p0.2 800

E 700

600 ODS (oxygen dispersion *, strengthened) EUROFER 500 alloys could raise operating

engt h F temperature for blanket by en 400 100° (improve efficiency)

300 MANET II - '>, 200 --O- F82Hmod. Q\ *'•• O OPTIFER V V, EUROFER Heat E83699 as-rec. 100 '•*.. ODS-EUROFER H. 958 as-rec. SSTT - -*- - ODS-EUROFER H. 958 as-rec. 03x18mm 0 0 100 200 300 400 500 600 700 800 900 1000

Test Temperature T [°C]

Could still impact on DEMO / 1st generation power plant designs CD a) o3T F.UROPEAN FUSION DEVELOPMENT AGREEMENT CD SEFDA 2 Improvements beyond ferritic steels II: lower activation materials

10 Dose rate of different materials after irradiation to 10 MWy/m2

1 - Vanadium alloy (V5Ti) 2 — reduced activ martensitic stainless steel

Material plus impurities Pure material Options for 2nd generation power plants (test modules on DEMO)

icrs io~6 1(T4 10"2 10° 102 104 Time after shutdown [years] CD SEFDA EUROPEAN FUSION DEVELOPMENT AGREEMENT Improvements beyond ferritic steels II: reduce activation and enhance thermai efficiency

exemption SiC- SiC- metal ' 1O"5H SiC low activation property after 100 y [Sv/h] low act. ODS recycling steel steel

med, active waste

500 600 700 800 900

operating temperature [°C] CD 03 O EUROPEAN FUSION DEVELOPMENT AGREEMENT 7 CD

s. Fusion Power: if we develop it - will there be a market ? "0 fusion power plant cost MARKAL projections of Fusion Electricity Production in Western Europe estimate in 2100, as function of imposed Constraints on CO2 emission cost estimates for fossile 25 fuels and renewables 20 • solar assumptions on fission Dwind t> 15 B biomass acceptance 3 D fusion B fission 10 detailled demand-supply • gas model Dcoal £ 5 El hydro UJ 7WT\ base 750 650 550 450 minimization of Irdsgraisi CO2 target (ppm) discounted costs for 33 producing required fusion would be a cost-effective supplier CD electrical energy of base-load electricity, if emission constraints are implemented

CD CO CD

EUROPEAN FUSION DEVELOPMENT AGREEMENT | I c Summary and Conclusions j \ c fusion has a high potential, but requires an integrated physics and s technology effort without precedence in non-military R&D I c

CJ the basic physics feasibility demonstration will be concluded with ITER ' (although R&D for efficiency improvement will continue) j The essential technological issues remaining at the start of ITER i operation concern material questions: f "" • first wall components \ • need to be solved for and on ITER (as they involve the complex g plasma environment) I radiation tolerant, low activation materials • development for a long-range target (up to 2nd generation power | plants) f • due to long time-lead for qualification, and strategy to test modules in « preceding device generation, an active development and test f programme is needed § K. Lackner et al. RM 1 495

15» international Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4 1 CD

EUROPEAN FUSION DEVELOPMENT AGREEMENT ÖEFDA 5 ITER Physics Basis - critical issues Steady state cfivertor power loading

V-shaped geometry used in target region favours development of partial detachment straight 86MW oh; V

new V

1 ; [ L—- j o^r-~- • \ >\\ Annular flow vertical target optk -100MW straight -100MW oldV - 100MW new V 130MW new V

0.26 0.28 0.3 0.32 0.34 0.36 FEAT: 3 n [10 ^m" ] geometry variation

Modelling using B2-EIRENE for ITER shows that under partially detached conditions, peal

power load on outer divertor remains below O 10MWm-2 over a range of separatrix densities 7T CD EUROPEAN FUSION DEVELOPMENT AGREEMENT 77 CD SEFDA CD ITER Physics Basis - critical issues

0.25 ' ' ' ' ' t

• Dlli-D ® JET 0.2 - § A ASDEX-U - 1 0.15 -

0.1 - Extrapolation to ITER: : A 0.05 - 10MJ/ELM x, A " A' ITER

i divertor life-time ! 0.01 0.1 10

> a Recent analysis of ELM energy loss indicates that padestal coHisionsiiiy and parailes transport time in the SOL are important

CD l CD CD SEFDA EUROPEAN FUSION DEVELOPMENT AGREEMENT the quest for tolerable ELMs (2) argon seeding 2

Reference Discharge Argon seeded discharge .3000Pulse No.53549 3000PulsaNo:53550 Inner divertor Argon Seeding § 2000 t 11000

18 20 22 24 16 18 20 22 24 Time (s) Time (s)

j - Outer dlvertor . g . Outer dlvertor to I Argon Seeding I 2000 5 2000 I I | s f Mm 1000 I 1000 §• r n'ii ii i i 1 a ^ '' f'( '' f i- ß ... '-• .j^_ 0 i • 1 '\ . '",' " V n J ' V' 16 18 20 22 24 16 18 20 22 24 Time (s) Time (s)

Strong target plate temperature reduction during Argon seeding CD CD 03 FUROPEAN FUSION DEVELOPMENT AGREEMENT 7T Z3 SEFDA CD R&D - Vacuum Vessel (L-3) (2)

33 Sector-B<1/2Secto Sector-A (1fZ Sector)

View of full-scale sector model of ITER vacuum vessel completed in September 1997 with dimensional accuracy of ± 3 mm t CD CD 500 RM 1 K. Lackner et al.

15» International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

O CD -X- -}_* sit e offe r CO Marc h 0 1 M— 0 Expecte d dat e o f z CO 5:

'ai * •K- O en < CD governmen t Statu s withi n -4—» ns : c .0

•X- -X- a -4—• 0 communit y

Ol Statu s withi n

z CD Commercia l offe r S0i c- iQ" z "oo 0 0. 0 sit e Hokkaid o prefectur e prefectur e Darlingto n Naka , Ibarak i Caderach e (F ) d. mura , Aomor i >3 ) Tomatoh , 1 ) JAER I Site , 05 >2 ) Rokkasho - E •0 CO 0 Canad a Europ e Japa n 7T EUROPEAN FUSION DEVELOPMENT AGREEMENT 13 CD

Roadmap, costs, negotiations: ITER costing iTER-JCT has produced a cost estimate in kiUA which establishes the : ; .: • value of individuai procurement packages This estimated determines the r value of each partners contribution klUA* Co ns t ruct io n Cos ts This estimate has established Direct ca pit al 2755 Manage men t & Support 477 the attainment of the goal: R&D During Construct ion -70 FEAT < 0.5 FDR Operation Costs (average per year) Permanent personnel 60 Energy -30 Fuel -8 Preferred option for partner's Maintenance/ imp rove men ts -90 contribution is in kind De c o mm is s io ning 335

*1 kIUA= S,,S,1M- S,„„, 1.39 2M-ü20001.27 9M- ¥,OCX1148M Each partner has carried out a commercial costing of . packages in his currency This costing is the basis for the /•: costs of his contributions and for the ; •; : cost of the project he quotes to his system

o en o

BS9 ^ ET i~\ A i:UR0PEAN FUSION DEVELOPMENT AGREEMENT Replacement of CFC by W for divertor strike zones considered

W in the divertor is practically only eroded by disruptions The issue is how much of the melt layer (~ IOOLUTI) will be lost What kind of longer term distortions of the target surface are to expect

Plasma compatibility, i.e. the W Time, u.s impurity migration into the main plasma is a concern A sacrificial thickness of 10 mm could last only for Promising results from ASDEX-UP suggest about 50-100 disruptions that W should be more seriously Situation for carbon might not be better, if brittle considered destruction plays a role. Longer disruption times can further reduce From Engineering point of view W disruption lifetime. clad targets have reached almost Experimental data regarding formation of melt the same development level as layers (and their properties)are very meagre. 8 CFC CD K. Lackner et al. RM 1 503 15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

\ o § | A i A

< S gi o 5 - A § 8 A h- Z

'S c

i 3 O

Majo r radiu s j 8.1 4 m Mino r radiu s 1 2.8 m a 1

c q < 111 LL.

t O Q LL < UJ Q h- LU Ol o SEFDA •UROPEAN FUSION DEVELOPMENT AGREEMENT

Zusammensetzung von EUROFER 97 und anderen (RAFM) Stählen

Legierung Radiologisch Spezifikation EUROFER 97 OPTIFER la F82Hmod. MANETII Schmelze / Analyse erwünscht EUROFER97 ' E 83699 SI 664 '9741 50804'" A) Cr 8,5-9,5 [9,0] 8,87 9,3 7,7 10,3 c 0,09-0,12 [0,11 ] 0,12 o,iö 0,09" 0,11 Mn " 0,20-0,60(0,40] 0,42 0,50 0.16 0,78 ' P

N2 0,015-0,045 [0,030] 0,018 0,0155 0,006 0,031 o2 < 0,01 0,013 0^0047 (0,01) B) W 1,0- 1,2 [ 1,1] 1,10 0,965 1,94 Ta ftÖ6"-"Ö,Ö9 0.14 0,066 0,02 ' C) Nb < 0,01 ppm < 0,001 (10 ppm) < 0,001 (10 ppm) 0,009 (90 ppm) 1 ppm 0,14 Mo < 1 pprn < 0,005 (50 ppm) < 0,001 (10 ppm) 0,005 (50 ppm) 30 ppm 0,61 Ni <10 ppm < 0,005 (50 ppm) 0,0075 (75 ppm) 0,005 (50 ppm) 200 ppm 0,68 Cu < 10 ppm < 0,005 (50 ppm) 0,021 (210 ppm) 0,035 (350 ppm)" 100 ppm " 0,01 AI < 1 ppm < 0,01 (100 ppm) 0,008 (80 ppm) 0,008 (80 ppm) 30 ppm 0,004 Ti < 200 ppm < 0,01 (100 ppm) 0,008 (80 ppm) 0,007 (70 ppm) 100 ppm ' Si < 400 ppm <Ö,05 (500 ppm) ' 0,06 (600 ppm) " "Ö",Ö6(60Ö"p'pm)"' 1100 ppm 0.19" Co <10ppm < 0,005 (50 ppm) 0,005(50 ppm) 50 ppm 0.005" As *) < 0,005 (50 ppm) 0,0093 (93 ppm)' (< 50 ppm) 0,01 Sn < 30 ppm *) < 0,005 (50 ppm) "0,0005 (50 ppm) (< 20 ppm) 0,001 < 0^005 (50 ppm) 0,0002" Sb *) p Zr *) < 0,005 (50 ppm) < 0,0002(2 ppm) 0,008 3J

A) ~ Grufidzusömnienscuung [ ] Zielwe-te B) ~ Substitutionseleinentc 500 ppm C) - rad:o!ogisch unerwünschEe ßeglei'elemente CD 2- . Mahalingam et al. RM 8 505

151' International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Materials Requirements for Thermal Management in Wireless Communication Mali Mahalingam, Vish Viswanathan Wireless Infrastructure Systems Division, Motorola Inc. 2100 E. Elliott Road, Tempe Az. 85284 mali.mahalinqam @ motorola.com

Wireless communication has been fueling the growth in the microelectronics industry in the last decade. Continued growth in the cellular voice communication and newer applications such as wireless broadband services, wireless internet, and universal wireless connectivity are expected to create increased growth in the wireless segment. RF (Radio Frequency) power amplifiers are key components to enable this growth. RF power amplifier semiconductor components for basestations dissipate large amounts of thermal energy, 1 - 3 kW/cm2 die level heatflux.

After a brief review of thermal management practices in the microelectronics industry, this paper will focus on RF power semiconductor components. A generic metal-ceramic package will be discussed to highlight performance parameters as well as materials, process, and design technologies associated with such packages.

Current and future trends in the RF device technologies will be reviewed. Successful RF LDMOS device will be contrasted with traditional RF Si bipolar device. New developments in GaAs, SiC, and GaN device will be reviewed. LTCC technology holds promise to integrate RF passives leading to higher level of functionality. These trends will be translated into packaging and materials requirements with a focus on thermal management. 506 RM 8 M. Mahalingam et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Materials Requirements for Thermal Management in Wireless Communication

by Mali Mahalingam, Vish Viswanathan Wireless Infrastructure Systems Division SPS, Motorola Inc.

Planscc Seminar May 28 200] «ai<™ou ^ ^ ;;: Mali Mahalingam & Vish Viswanathan IlltfTiSl DIM

Outline Background - Market - Thermal & Thermo-Mechanical Management in electronics

RFPAs - Why an emphsis on Thermal Performance? - LDMOS: Device & Packaging

Materials and Comparison - Assembly stack up, Process Flow and Interfaces - Materials used & Rationale for their use Requirements - Performance & Reliability - Future Technologies: Plastic Pwr., LTCC, Device Summary and Recommendations

„••••••... ßk t .w„-~ Planscc Seminar May 28 2001 "B™"°^ ,, .qgp ...... Mali Mahalingam & Vish Viswanalhan W§>t$l iV M. Mahalingam et al. RM 507

15m International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

The Evolution of the Wireless Consumer Electronics Industry Devices with Ubiquitous Wireless Connectivity

Cellphone

CO PC (Mass produced digital hardware) VCR (Mass produced complex analog hardware Radio and high-tech consumer pricing models)

1900 1920 1940 1960 1980 2000 2020 Courtesy: Lance Wilson

Planscc Seminar May 2S 2001 Mali Mahalingam & Vish Viswanathai

The Wireless World Today

° The market is huge and increasing - 866 million subscribers worldwide (2001 forecast). Growing to 1.6 billion in 2005. - 1 million operational base stations world wide (2000). - 106K operational base stations in the U.S. (2000).

e Wireless is expanding rapidly - PDA's - Laptops, notebook and desktop computers - Wireless Internet appliances (for email/Internet access) - Wireless "smart" appliances (microwaves, dryers, etc.) Coimesv: Lance Wilson

Plans«; Seminar May 2S 2001 Mali Mahalingam & Vish Viswanalha 508 RM8 1. Mahalingam et al.

15™ International Plansee Seminar, Eds. G. Kneringer, P. Rädhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Why Thermal Management in Electronics?

- CMOS: Higher channel T -• slower speed -Bipolar: Higher junction T -*• leaky junction

- Higher speed -* higher packing & power density

- Higher temperature ~* lower reliability

Planscc Seminar May 28 2001 *lonM0^ , : = Mali Mahalincam& Visit Vii.wiircilh.in IIIj?'WI li'lV.?

Electronic Component Reliability Thermal Acceleration Factor, pj = Failure rate at temp T Failure rate at 75°C Digital Bipolar Devices Part AT CC) Stress nT Si 105 -5-7 Transistor

Resistor 75 -22 (carbon composition)

Transformer 60 -33 (coils)

Capacitor 95 -47 75°C lOO-C 125°C 1SO°C (glass)

Reference: MIL HDBK-217B

Plansco Seminar May 2S 200] Mali Mahalingnm & Vish Visnanathan Mahalingam et al. RM8 509

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol.'

Why Manage Thermo-Mechanical Stress?

• In packaging Variety of Materials are used with a wide range of -- Thermal Expansion Coefficient -- Strength • During assembly & test, packages are exposed to large Temperature Ranges

-* Increased potential for failure due to thermally induced stress

Plans« Seminar May 28 2001 «Olmou ^ zz Mali Mahalingam & Vish Visvranalhan \\\Q!V>\ li'fW

Trends In Variables ~ Electronic Package Thermo-Mechanical Integrity INTERFACES: MATERIALS WITH WIDE RANGE OF TCE & STRENGTH ASSEMBLY & TEST: LARGE RANGE OF TEMPERATURE EXPOSURE

TCE (PPM/°C) PROCESS TEMP (X) R&QA TEMP. STRESS (°C) . Ag GLASS t GOLD EUTECTIC — MOLD COMPOUND ( 26) PbSn • (

SILICON (~3)

Planste Seminar May 2S 2001 Mali Maha ingam & Vish Viswanalhan Si 1 ft" F.S 1II A? 510 RM8 1. Mahalingam et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Probability of Material Failure

PROBABILITY

Process Induced Stress

Material Strength

FORCE/AREA OVERLAP OF THE MATERIAL STRENGTH & PROCESS INDUCED STRESS DISTRIBUTION CAUSES FAILURE.

Plans« Seminar May 28 2001 Mali Mahalingam & Visli Viswanathan

Cellular Basestation Forced Air Cooled Rack

IS isLst ltion ( oiilrulki hi**

Basestation «/Antenna

Courtesy; WISP Marketing RFPVs Melil Ceramic

PLanscc Seminar May 2S 2001 Mali Mahalinsam & Vish Viswamuhar M. Mahalingam et al. RM8 511

15th International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Heat Flux vs Operating Temperature

Flux Process Temp (°C) (W/m2) 10 s

7 RFPA Device j* Heat loss from 37 -500 10 Human Body 106 / Nuclear Blast 6 -2000 -10 5 . Flux at 1.5 Km 1O 2 W/m y 104 Sun's Surface - 6000 ~5xlO7 103 2 RF Power w 7 Chip - 150 ~3N 10 10 Temp°C 100 1000 10,000 Challenge: Very high device level heat flux. Small AT to transport the heat.

Plansec Seminar May 28 2001 Mali Mahaüngam & Vish Viswanathan

Why Enhanced Device Packaging Techniques Are Important to Wireless Infrastructure

• Base stations are running high power RF transmitters - Base stations are getting physically smaller (practically the same heat in a smaller box) - Transmitter efficiency is a constant battle - Users and manufacturers don't want active cooling systems if possible

- Efficient Device Thermal Performance is a must

Courtesy: Lance Wilson

Plansee Seminar May 28 2001 MaH Mahalingam & Vish Viswanathan 512 RM. M. Mahalingam et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

LDMOS vs Bipolar - RF Device & Pkg Structures iipolar

Base Emitter

P 11- coHector V=~26V

backside n + subslrlie metal

]> I S^^^^^-z^-E-zrEJ-z:-^ Dielectric ^^\v^\Xt^\^\^^^V^W\\ • Source :it Ground Potential (Heo) ^> ^ • Ver(ic;il current flow 1'kg Metal ^ ^> Flange ^—- ^* Key Packaging Advantages Major Packaging Disadvantages

-- Die bonded directly to metal flange. Collector needs to be electrically isolated. Better Thermal Performance. Need Beo or A1N isolator. Complex pkg. Simpler pkg. Structure. — Topside emitter contact forces higher lead -- Very low source inductance. inductance.

Planscc Seminar May 28 2001 Mali Mahalingam & Vish Viswanathan

Assembly Stack up for RF LDMOS Devices

AIIoy42 LE VDS

DEVICE ASSEMBLY PKOCKSS

- -^"^ CuSIL BRA/I- K ^-> ALUMINA ^-^>^^ -^-^^» WINDOW ,. i, \ ^ * «i - •%•

CuSIL BRA/L_^

^-^ ' PACKAGE FOR DEVICE ASSEMBLY v DIE BOND AT 320 C t CuW FLANt #* WIRE BOND

BRAZING BY PKG. SUPPLIER°AT 820 C LID SEAL 1 DCYRF TEST LAPPING (OPTIONAL) 1 PLATING

A • ••• • ..--.fli-.r;..---: Planscc Sem nar May 28 2001 Mali Mahalingan & Vish ViswaiKithan Mahalingam et al. RM8 513

15™ International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Trends In Variables -- RF Power Package Thermo-Mechanical Integrity • INTERFACES: MATERIALS WITH WIDE RANGE OF TCE & STRENGTH • ASSEMBLY & TEST: LARGE RANGE OF TEMPERATURE EXPOSURE

TCE (PPM/°C) PROCESS TE,Wf°C) R&OA TEMP, STRESS f °C) CuSil - •»- BRAZING

; _ tGOLD — EPOXY(-26) EUTECTIC PbSn — (

(~12)AuSi- EPOXY TEMP. CYCLE THERMAL SHOCK - ALUMINA (-7) (•65<>CTO+150OC) -SILICON (-3)

Planste Seminar Ma}' 28 2001 Mali Mahalingam & Vish Viswanathan

Process Flow

ALLOY42 LEADS CuW FLANGE (HIGHTC, UNIFORM DENSITY) + t NICKEL PLATING (FOR BRAZABLE SURFACE) METALLIZED CERAMIC T WINDOW FRAME ANNEALING (TO BOND Nl TO CUW)

JOINING COMPONENTS WITH CuAg BRAZE @ 800C (CAMBER DUE TO CTE MISMATCH) * SPECIAL NICKEL PLATING (TO IMPROVE DIE ATTACH QUALITY)

GOLD PLATING (TO IMPROVE DIE ATTACH QUALITY RF ELECTRICAL CONTACT)

(ft ,.„-.,..• • Planscc Seminar May 28 2001 Mar&aaLj. ^ -•- W •'•'"••• Mali Mahalingam & Vish Viswanathan illij'-'ISl Ol'W 514 RM M. Mahalingam et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Heat Sink Materials

HEAT SINK MATERIALS ADVANTAGES DISADVANTAGES

X n COPPER TUNGSTEN CTE, E , TC S, IVIACHINING

COPPER MOLY CTE, $, STAMPABLE TC

i npos i o

Me t J AlSiC CTE, WEIGHT TC, FORMING, $

Cu(CuMo)Cu $, TC.STAMPABLE CTE

Laminate d Composit e _Cu(CuW)Cu TC, FORMING S, CTE

FUNCTIONARY GRADED TC, CTE S, FABRICATION _MATERIALS Functiona l Composit e

LCOPPER S$$, TC, STAIVIPABLE CTE, E, Temp stb. Meta l

4£h •- •• •>'•' Plansee Seminar May 2,S 2001 Hir " Mali Mahalingam & Vish Viswanalhan

Camber vs Heat Sink CTE

2 ä 50 > c Preferred Zone (0 to 25 Micro 25 • Ito O '^meter Convex)

Jö ^""^•»«üüü^; >v Low AT

CTE of Heat Sink Material (for Process AT)

Plansec Seminar May 28 2001 Mali Mahalingam & Vish Viswanathan M. Mahalingam et al. RM 515

15* International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Copper

KEY: Somiconduclo A1N • Q Electrically Electrically Insulating Conducting

30 60 90 120 150 180 210 250 260 310 340 370 Thermal Conductivity (W/mK) at RT

Planscc Seminar May 28 2001 Mali Mahalingam & Visll Viswanalhan

Interactions: Flange Vs. System Materials

zu?

Planscc Seminar MnylS 2001 Mali Mahalingam & Visli Viswanalhan fllff-Ti?! SlitS 516 RM 1. Mahalingam et al. 15;h International Plansee Seminar, Eds. G. Kneringer, P. Rörjhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

IR Scan - RF LDMOS Power Device

J Hottest temp. ~ 157°C | Cell to cell temp, varies

i i

45W DC Pwr. Dissipation Heatsink Stage @ ~ 92 °C Theta jh = 1.43 °C/W

Plansce Seminar May 28 2001 Mali Mahnung» & Vish Visu-.innthan

Metalized Ceramic Pkg.-Robust Design Manufacturing, Assembly, Mounting Process Thermo-Mechanical Loading Tr.rasKrDc Temperature Profile During Assembly and \O9d\P Testing

Bolt Down simulated separately

1 Stresses in critical areas (Si, Ceramic) computed. 1 Failure strength of critical materials (Si, Ceramic) measured. 1 Robust designs being developed.

••-tEi- .•• •*••• Plansce Seminar May 28 2001 MOTHMiau» :; " W ' •"' ' ' Mali Mahalingam & Vish Visranathan iliff.'f&UbH 1. Mahalingam et al. RM 517

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Failure Strength Testing of Ceramics Weibull Analyses Weibull Plot Failure Strength of Ceramics 55000

Means and Std Deviations Level Number Mean Std Dev Std Err Mean Weibull Parameter Estimates AINVA/Vias 24 42381. U 5275.28 1076.8 Group Beta BeO/Vias 25 28A35.V 5502.21 1100.4 AIN/AA'ias 9.S60SO

BeOA'ias 5.M.165

Plansec Seminar May 28 2001 Mali Mahalingam & Vish Viswanatha

Non-Destructive Tests Rationale Flatness Mountabilitv, Low Rth, Electrical Groundinq Sonoscan Integrity of Attach Joints - Window Frame - Die Attach Gross Leak Prevent Gross Contaminants into Device

Destructive Tests Rationale Temp Cycle (-65C to +150C) DC/RF Mechanical Integrity of the system Mech. Shock/VVF Integrity under shock loading HTSL (150° C, 1000 Hours) DC/RF Robustness of Device HTSL (300°C, 24 hour, die bond only) Robustness of Flange Plating - Die Attach - Die Shear Die Attach Integrity - Auger Surface contaminants after the test. H3T Drop Test Integrity under shock Humidity Biased TC Elecrochemical Robustness of System (50 cycles, -5° to 70°, 5V) Solderability Device Mountability (8 hour steam age, tin clip) JEDEC Preconditioning {-lOTC's, 150° C, 24 hours) Device Mountability -Build, Solderability Lead Straightening/Pre-Tinning Ease of use by cudstomer - Solder Ingress, Gross Leak, Sonoscan HAST (Highly Accelerated Stress Test) Robustness (130° C, 85% RH, 2 atm, 96 hrs, bias)

Planscc Seminar May 2S 2001 Mali Mahalingam & Vish Viswanathan 518 RM8 M. Mahalingam et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Performance & Reliability Requirements

• CONSISTENTLY HIGH THERMAL CONDUCTIVITY (Kth). • UNIFORMITY IN THE FLANGE (DENSITY, Kth) • CTE OF FLANGE COMPATIBLE WITH PKG. ASSEMBLY. • Ni AND Au PLATING FOR EXCELLENT DIE ATTACH. • PACKAGE ROBUST UNDER THERMAL CYCLING • PACKAGE FLATNESS CONSISTENT WITH MOUNTING REQUIREMENTS. • HUMIDITY AND TEMPERATURE STABILITY. • COST EFFECTIVE SOLUTION: DRIVE DOWN $/W METRIC

Plansee Seminar May 28 2001 Mali Mahalingam & Visli Viswanachan

Plastic Power RF Packages Package Max.Power Dissipation *

PLD-1 8 i 6 Watts IS

PLD-1.5 Hm m 25 Watts

PFP-16 20 Watts

TO-270(PRFP-i) jBH V 35 Watts

TO-272(I>RFi>-2) "JHpiHi Pjfff" 65 Watts

* Tc = 100»C, Maximum Tj = 150°C

Piansee Seminar May 28 2001 Maronen.* Mali Mahalingam & Vish Viswanathaa StiRi M. Mahalingam et al. RM8 519 15* International Plansee Seminar, Eds. G. Kneringer, P. Rodhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

LTCC Packaging Technology LTCC = Low Temperature Cofired Ceramic What is LTCC Technology? Multilayer circuit fabricated by laminating unfired sheets of glass-ceramic containing printed circuit components (capacitors, inductors and resistors) and interconnections, and firing the structure to form a monolithic ceramic circuit. .i;:!: Until .Wjihlr Conceptual LTCC Module: - Passives can be integrated - Hi Q passives can be accomplished - Controlled impedance can be achieved - Chips/Components can be surface mounted

Source: Microwave & RF. JuJy 98, page 45

Planscc Seminar May 28 2001 Mali Mahalingam & Vish Viswanathan

LTCC Packaging Technology LTCCvs HTCC Properties Parameter LTCC HTCC

CTE of dielectric -3-10 -3-7 (ppm/C) Kth of dielectric -2-6 - 15 - 250 (W/mK) Flexural Strength -17-26 - 40 - 50 (Kpsi) Fired Shrinkage (%) (for Dupont951) Std. Hi Density x,y -13 -1 - 0.5 z -15 -10 -5

Plans« Seminar May 28 2001 m"™"" , Mali Mahalingam & Vish Viswanalhan #fS>X3\ »f\ 520 RM8 M. Mahalingam et al.

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

Future DeviceTechnologies For RFPAs LDMOS: Appropriate for high linearity, high efficiency, very high power levels, and low cost up to 3GHz. Best overall solution for frequencies up to 2.2GHz., 2.2-3GHz spectrum being a source of contention. SiC: High voltage capability, but material quality issues and manufacturability concerns consign to niche applications only. GaAs: At 3 GHz and higher, the technology of choice today. Limited high voltage capability is an issue in base station PA design. GaN: Great potential for high voltage, high frequency PA devices. Material limitations today, not commercially viable at this time. May displace GaAs as the high power technology of choice for 3GHz and higher as material properties are improved.

Plansce Seminar May 28 200I Mali Mahalingam & Visl] Viswanatl

Package Cost Per Watt for > 45 Watt RF Transisters

3.5

No r 3

2.5

2 ^^Pout 1.5

Pfijs 0.5 b ®^ 0 0 500 1000 1500 2000

Frequency (MHz)

Planscc Seminar May 28 2001 Mali Mahalingam & Vish Viswanalhan M. Mahalingam et al. RM 521

15" International Plansee Seminar, Eds. G. Kneringer, P. Rödhammer and H. Wildner, Plansee Holding AG, Reutte (2001), Vol. 4

SUMMARY AND CONCLUSIONS > WIRELESS MARKET CONTINUES TO BE A HIGH GROWTH AREA.

• POWER, SPACE AND COST CONSIDERATIONS IN RFPA PACKAGING HAVE NECCESSITATED CONSTANT IMPROVEMENTS IN HEAT SINK SOLUTIONS.

• THE SELECTION OF HEAT SINK MATERIALS IS BASED ON i) COMATIBILITY WITH OTHER COMPONENTS OF THE PKG. AND SYSTEM ii) WEIGHT AND THERMAL REQUIREMENTS OF THE SYSTEM.

• SIMULATIONS & TESTING ARE USED TO SELECT MATERIALS OF PKG. CONSTRUCTION TO ENSURE PERFORMANCE AND RELIABILITY.

•COLLABORATION BETWEEN DEVELOPERS OF HEATSINKS, PACKAGES, DEVICES AND SYSTEMS WILL BE NEEDED TO MEET MARKET DEMANDS.

Planscc Seminar May 28 2001 Mali Mahalingam & Vish Viswanath.in

RECOMMENDATIONS

• INDUSTRY HAS TO PROVIDE COMPETITIVE SOLUTIONS TO THERMAL REQUIREMENTS OF RFPAS FOR WIRELESS COMMUNICATIONS MARKET.

• THIS IS ACHIEVED BY COMPLETE UNDERSTANDING OF DOWN STREAM PROCESS / REQUIREMENTS OF CUSTOMERS AND BY COLLABORATION ACROSS THE "SOLUTION CHAIN".

• DEVELOPMENT AND COMMERCIALIZATION OF HIGH THERMAL CONDUCTIVITY, LOW COST, FUNCTIONAL MATERIAL SYSTEMS WILL CONTINUE TO MAKE AN IMPACT IN THE ELECTRONIC SYSTEMS BUSINESS.

• OFTEN SUCH THERMAL SOLUTIONS ENABLE NEW DEVICE AND SYSTEM TECHNOLOGIES TOWARDS LOW COST AND HIGH PERFORMANCE SOLUTIONS.

Planscc Seminar May 28 200) Mali Mahalingam &. Vish Viswanathan