Ceramic Materials and Components for Engines
Edited by Jurgen G. Heinrich and Fritz Aldinger
@WILEY-VCH
Further Titles of Interest
J. Bill, F. Wakai, F. Aldinger (Eds.) Precursor-Derived Ceramics ISBN 3-527-298 14-2 R. Riedel (Ed.) Handbook of Ceramic Hard Material ISBN 3-527-29972-6 G. Miiller (Ed.)
Ceramics - Processing, Reliability, Tribology and Wear ISBN 3-527-30194-1
Ceramic Materials and Components Ior Engines Edited by Jurgen G. Heinrich and Fritz Aldinger
Deutsche Keramische Gesellschaft
Weinheim - New York - Chichester Brisbane - Singapore Toronto
-
Prof. Dr. Jiirgen G. Heinrich TU Clausthal Institut fur Nichtmetallische Werkstoffe Professur fur Ingenieurkeramik ZehntnerstraBe 2a 38678 Clausthal-Zellerfeld Germany
Prof. Dr. Fritz Aldinger Max-Planck-Institut fiir Metallforschung Heisenbergstrak 5 70569 Stuttgart Germany
71hInternational Symposium “Ceramic Materials and Components for Engines” Applications in Energy, Transportation and Environment Systems Organizer: Deutsche Keramische Gesellschaft
This book was carefully produced. Nevertheless, editors, authors and publisher do not warrant the information contained therein to be free of errors. Readers are advised to keep in mind that statements, data, illustrations, procedural details or other items may inadvertently be inaccurate.
Cover: The cover picture shows the high-performance brake of the new Mercedes-Benz CL 55 AMG “F1 Limited Edition”. The brake discs are made of a carbon fiber-reinforced ceramic and are the f i s t CMC material that is introduced to series production in the automobile industry. The picture is courtesy of DaimlerChrysler Communications, Stuttgart (Germany).
Library of Congress Card No.: applied for A catalogue record for this book is available from the British Library. Die Deutsche Bibliothek - CIP Cataloguing-in-Publication-Data A catalogue record for this publication is available from Die Deutsche Bibliothek ISBN 3-527-30416-9
0 WILEY-VCH Verlag GmbH, D-69469 Weinheim (Federal Republic of Germany). 2001 Printed on acid-free paper. All rights reserved (including those of translation in other languages). No part of this book may be reproduced in any form by photoprinting, microfilm, or any other means - nor transmitted or translated into machine language without written permission from the publishers. Registered names, trademarks, etc. used in this book, even when not specifically marked as such, are not to be considered unprotected by law. Printing: betz-druck gmbh, 64291 Darmstadt. Bookbinding: Wilhelm Osswald & Co., 67433 Neustadt. Printed in the Federal Republic of Germany.
Foreword With the 7'h International Symposium Ceramic Materials and Components for Engines, June 19-21, Goslar, the German Ceramic Society was hosting this international conference for the second time since 1986. Since that time a lot has changed in regard to the use of ceramic components in engines. A number of parts have proven their suitability through large-scale experiments for seriel application in very impressive ways, especially in Japan, the USA and in Germany. Despite this fact, neither ceramic gas turbines in automobiles nor ceramic components in reciprocating engines, as for example valves, turbocharger motors, precombustion chambers or portliners, were successful in mass production. While the reliability of these components could be proven by seriel application, there is still a lack of economical quality assurance concepts. The ceramic components are still too costly compared to metallic components. Therefore the main emphasis of research for these conventional products is the reducement of costly starting materials and the development of processing techniques at a good price. For the past few years a new generation of ceramic components, for the use in energy, transportation and environment systems, has been developed. The efforts are more and more system oriented in this field. The only possibility to manage this complex issue in the future will be interdisciplinary cooperation. Chemists, physicists, material scientists, process engineers, mechanical engineers and engine manufacturers will have to cooperate in a more intensive way than ever before. This is the only way to successfully develop the complicated systems by using ceramic materials. The R&D activities are still concentrating on gas turbines and reciprocating engines, but also on brakes, bearings, fuel cells, batteries, filters, membranes, sensors and actuators as well as on shaping and cutting tools for low expense machining of ceramic components. As a result, the range of materials, which are actually in discussion for these applications, have considerably increased. Besides presentations on the systems mentioned, this conference offered discussions on important issues such as performance, reliability, design, modelling and simulation, expense effective manufacturing as well as material design and process development. This book summarizes the scientific papers of the conference. Especially in the plenary lectures some of the most fascinating new applications of ceramic materials in energy, transportation and environment systems are presented They are followed by special articles on the main topics of the conference. The symposium led to a lot of contacts between collegues from industry universities and research institutes. The proceedings shall lead to new ideas for interdisciplinary activities in the future.
Goslar, June 2000
Jurgen G . Heinrich Fritz Aldinger
Contents
1.Systems. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1 New Technologies in the Light of Materials . . . . . . . . . . . . . . . . . . . . . 3 The Promise of SOFC Power Generation Technology . . . . . . . . . . . . . . . . . . 7 Ceramic Matrix Composites for Disk Brakes and Their Manufacturing Technologies . . . . . . . . 13 Ceramic Cutting Tools . . . . . . . . . . . . . . . . . . . . . . . . . . . . 21 Practical Use of Ceramic Components and Ceramic Engines . . . . . . . . . . . . . . . . 27 Beta Ceramic in Zebrae and NAS Batteries . . . . . . . . . . . . . . . . . . . . . 33 Oxygen Sensors for Lean Combustion Engines . . . . . . . . . . . . . . . . . . . . 39 Ceramic Gas Turbine CGT302" Development Summary . . . . . . . . . . . . . . . . 45 Preparation of Planar SOFC-Components Via Tape Casting of Aqueous Systems. Lamination and Cofiring . 51 Glasses from the System RO.R,O,.SiO, as Sealants of High Chromium Steel Components in the Planar SOFC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57 Lightweight and Wear Resistant CMC Brakes . . . . . . . . . . . . . . . . . . . . 63 Development of Ceramic Sheathed Type Thermocouple with High Heat Resistance and High Durability . . 69 Advances in Hot Gas Filtration Technique . . . . . . . . . . . . . . . . . . . . . 73 Nano-Scaled Ceramic Membranes for the Filtration of Fine and Sticky Dust . . . . . . . . . . . 79 Production and Characterization of TiCN-Based Materials for Cutting Tool Applications . . . . . . 85 Sensitivity Characterization to Flamable Gas . . . . . . . . . . . . . . . . . . . . . 91
.
.
I1 Performance / Reliability . . . . . . . . . . . . . . . . . . . . . . . . . Evaluation of Mechanical Reliability of S1.N. Nozzles after Exposure in an Industrial Gas Turbine . . . Friction and Wear of Advanced Ceramics . . . . . . . . . . . . . . . . . . . . . Lifetime Prediction for Silicon Nitride . . . . . . . . . . . . . . . . . . . . . . Standardising Measurement and Test Methods for Advanced Technical Ceramics . . . . . . . . Investigations on the Stable Crack Growth of Indentation Cracks . . . . . . . . . . . . . . Prediction of Thermal Shock Resistance of Components Using the Indentation-Quench Test . . . . . Ceramic Components for Metal Forming Tools . . . . . . . . . . . . . . . . . . . Effect of Grain Boundary Composition on High-Temperature Mechanical Properties of Hot-Pressed Silicon Carbide Sintered with Yttria . . . . . . . . . . . . . . . . . . . . . . Thermal Shock Properties of SiALON Ceramics . . . . . . . . . . . . . . . . . . . Corrosion of Nonoxide Silicon-Based Ceramics in a Gas Turbine Environment . . . . . . . . . Mechanical Properties and Wear Behaviour of Differently Machined Silicon Nitride and Silicon Carbide Ceramic Surfaces . . . . . . . . . . . . . . . . . . . . . . . . . . . . Design of Wear Resistant Polycrystalline Alumina . . . . . . . . . . . . . . . . . . Crack Growth of Ceramic Materials in Sliding Contact . . . . . . . . . . . . . . . . . Damage Detection in Tetragonal Zirconia Polycrystals (TZP) by Impedance Spectroscopy . . . . . . Reliability and Reproducibility of Silicon Nitride Valves: Experiences of a Field Test . . . . . . . Self-Mated Tribological Properties of Plasma Sprayed Chromium Carbide Coating . . . . . . . . Role of Grain Size in Scratch Damage Resistance in Zirconias and Silicon Nitrides . . . . . . . . Ceramic Coatings with SolidLubricant Ability for Engine Applications. . . . . . . . . . . . Tribological Behavior of Silicon Nitride/Steel Contacts under Lubricated Conditions . . . . . . . Optimization of the Brazilian Disc Test for Ceramic Materials . . . . . . . . . . . . . . .
VI
95 97 103 109 115 121 127 133 139 147 153 157 163 169 175 181 187 193 199 205 211
The Impulse Excitation Technique for Rapid Assessment of the Temperature Dependence of Structural Properties of Silicon Nitride and Zirconium Oxide Ceramics . . . . . . . . . . . . . . . 217 VAMAS Round Robin Testing of High Temperature Flexural Strength . . . . . . . . . . . . 223 Effect of High Voltage Screening Method on Titania Ceramics with Different Surface Finishing . . . . 229 In Situ Observation of Tension and Cyclic Fatigue Damage in Hi-NiCALON Fiber/SiC Composite . . . 233 Multi-Axial Strength Data for A1.0.. and MgO-ZQ-Ceramics . . . . . . . . . . . . . . 239
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I11 Design. Modelling and Simulation. . . . . . . . . . . . . . . . . . . Fatigue Design and Testing of Ceramic Intake and Exhaust Valves . . . . . . . . . Design and Testing of a Prototype SiSiC heat Exchanger for Coal Combustion Power Stations DesignandTestingofCeramicComponentsforIndustrialGasTurbines . . . . . . . An Investigation on Paste Flow in a Press-Moulded Ceramic Dome . . . . . . . . . Design Standard for Advanced Ceramic Materials and Components . . . . . . . . . Static and Cyclic Stress-Lifetime Curves of Ceramics . . . . . . . . . . . . . Lifetime Prediction of Ceramic Thermal Barrier Coatings Based on Lifetime Analyses of Close to Reality Tests . . . . . . . . . . . . . . . . . . . . . . . . . .
. . . 245 . . . . 247
. . . . . . . . . . . .
255 261 267 273 279
. . .
285
. . . .
. . . .
CreepofaSiliconNitrideunderVariousSpecimen/LoadingConfigurations. . . . . . . . . . 291 Mathematical Model of Microhardness of Plasma Sprayed Chromium Oxide Coating . . . . . . . 299 Lifetime Prediction Model for Plasma-Sprayed Thermal Barrier Coatings Based on a Micromechanical 305 Approach . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Analytical Design and Experimental Verification Methods of Ceramic Radial Turbine Rotors 311 for a Gas Turbine . . . . . . . . . . . . . . . . . . . . . . . . . . . . The Evolution of Damage in Ceramic Materials for Gas Turbine Applications under Complex Load 319 Conditions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Mechanical Failure of Electroceramic Components . . . . . . . . . . . . . . . . . . 325 The Energy and The Power Time Dependence on the Ultrasonic Welding Process - A Weibull Statistical 333 Based Model . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
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IV Cost Effective Manufacturing . . . . . . . . . . . . . . . . . . . . . . . Advances in Brazing of Ceramic Materials for Engines . . . . . . . . . . . . . . . . . Stereolithographyfor Ceramic Part Manufacturing . . . . . . . . . . . . . . . . . . Application of the Mold SDM Process to the Fabrication of Ceramic Parts for a Micro Gas Turbine Engine Potential of the Hydrolysis Assisted Solidification Process for Wet Forming of Si.N. Ceramics . . . . Green Machining of Aluminium Oxide Ceramics . . . . . . . . . . . . . . . . . . . Laser-Assisted Turning of Silicon-Nitride Ceramics . . . . . . . . . . . . . . . . . . Ceramic Engineering with Preforms for Locally Reinforced Light Metal Components . . . . . . . Laser Beam Welding of Alumina - a New Successful Technology . . . . . . . . . . . . . SRBSN Material Development for Automotive Applications . . . . . . . . . . . . . . . Process Design for High Performance Grinding of Advanced Ceramics in Mass Production . . . . . New Ceramic Excellence for Complex Machining of Engine Materials . . . . . . . . . . . . Process Strategies for Grinding of Advanced Ceramic Cutting Tools . . . . . . . . . . . . . Ultrasonic Assisted Face Grinding and Cross-Peripheral Grinding of Ceramics . . . . . . . . .
.
V Material Design and Process Development . . . . . . . . . . . . . . . . . . . Advancing in Mechanical Properties of Silicon Nitride: The Roles of Starting Powders and Processing . . Materials Design of Composite Materials . Compatibility of Self-Damage Monitoring and Strengthening . The Design of Composition and Mechanical Properties of a-P SiAlON Ceramics Densified with Higher Atomic Number Rare Earths . . . . . . . . . . . . . . . . . . . . . . Grain-Boundary Phase Control of Silicon Nitride Materials . . . . . . . . . . . . . . . Sintering and Microstructure of Silicon Nitride with Magnesia and Cerium Additives . . . . . . . Characterisationof Multi-Cation Stabilised Alpha-SiALONMaterials . . . . . . . . . . . .
345 347 353 359 365 371 377 383 387 393 399 405 411 417 423
425 43 1 435 439 443 447
VII
Grain Boundaries of SiC.Si0. Composite . . . . . . . . . . . . . . . . . . . Millimeter Wave Sintering of Ceramics . . . . . . . . . . . . . . . . . . . . Internally Cooled Monolithic Silicon Nitride Aerospace Components . . . . . . . . . . Preparation and High Temperature Strength of Gd.Al. 0 J M g 0 Composites. . . . . . . . . Characterization of In Situ Sic-BN Composites . . . . . . . . . . . . . . . . . Coating Experiments on Carbon Fibers Using a Continuous Liquid Coating Process . . . . . . Effect of A1 Component on Mechanical Properties in Al-Penetrated Alumina . . . . . . . . Cavitation Creep in the Next Generation Silicon Nitride . . . . . . . . . . . . . . . Thermal Conductivity and Phonon Scattering Mechanisms of p-Si3N4Ceramics . . . . . . . A High Thermal Conductive p-Silicon Nitride Substrate for Power Modules . . . . . . . . Characterisation of the Pore Structure of Biomorphic Cellular Silicon Carbide Derived from Wood by Mercury Porosimetry . . . . . . . . . . . . . . . . . . . . . . . . TBC Consisting of New Metal-Glass Composites . . . . . . . . . . . . . . . . . Aspects on Sintering of EB-PVD TBCs . . . . . . . . . . . . . . . . . . . . Crack Propagation in a Thermal Barrier Coating System . . . . . . . . . . . . . . Porosity Graded Silicon Carbide Evaporator Tubes for Gasturbines with Premix Burners . . . . Porous Ceramics Functional Cavities for System Innovation . . . . . . . . . . . . . Feasibility Studies on Applying In Situ Single Crystal Oxide Ceramic Eutectic Composites in Non-Cooled High Efficiency Turbine System . . . . . . . . . . . . . . . . Synthesis and Property Tailoring of Reaction-Based Composites: The RBAO and the 3A Process . Properties of Silicon NitrideKarbide Nano/Microcomposites- Role of S i c Nanoinclusions and Grain Boundary Chemistry . . . . . . . . . . . . . . . . . . . . . . . Layered Si. Nd(SiAlON+TiN) Composites with Self-Diagnostic Ability . . . . . . . . . . Multilayer C/SIC Composites for Automotive Brake Systems . . . . . . . . . . . . . Sinter Additive Optimization in Processing of Aluminum Nitride for Heat Exchanger Components . New Opportunity for Bimodal Microstructure Control in Silicon Nitride . . . . . . . . . Design of SiCN - Precursors for Various Applications . . . . . . . . . . . . . . . Gelatin Casting and Starch Consolidation of Alumina Ceramics . . . . . . . . . . . . Liquid-Phase Sintered Silicon Carbide Based Ceramics with AlN.Y,O, and AlN.L%O. Additives . Slip Casting of ATZ Ceramics . . . . . . . . . . . . . . . . . . . . . . . The Preparation for Sintered Body of CeO. Based Complex Oxide in Low Temperature Solid Oxide Fuel Cells Using Colloidal Surface Chemistry . . . . . . . . . . . . . . . . . Processing and Properties of Tic-Ni3AlComposites . . . . . . . . . . . . . . . . Ceramic Joints Between S i c Bodies: Microstructure. Composition. and Joining Strength . . . . Fatigue Behaviour of Ceramics Stressed Near Fatigue Limit under Rotary Bending . . . . . . Laser Cutting and Joining of 2D-Reinforced CMC . . . . . . . . . . . . . . . . Analysis of Compounding and Injection Moulding Process of Ceramic Powders . . . . . . . Use of the SOL-Gel Method in the Extrusion of Alumina Ceramics and ATZ Ceramics . . . . . Advanced Hot-Pressed Ceramic Matrix Composites (CMC) in Sicw.Sic. C,,/Si.N. Systems . . . High Temperature Si3N4-BNComposite . . . . . . . . . . . . . . . . . . . . Processing of SICN-Fibres Prepared from Polycarbosilazanes . . . . . . . . . . . . . Design of Grain Boundary Phases in Silicon Nitride - Silicon Carbide Nano-Composites . . . . Yb-Si-A1-0-N Glasses and Glass-Ceramics as Grain-Boundary Phases for Silicon Nitride Materials .
Index.
VIII
. .
453
. . 457 . . 463 . . 469
. . . .
471 477 . . 483 . . 487 . . 495 . . 499 3
. . . .
. .
505 513 . 517 . 523 . . 531 . . 537 543 549
. . 553
. . . . .
. . .
559 565 571 . 577 . 581 . . 587 . . 593 . . 599
. . 605
. . . .
. . . .
611 617 621 627 . . 631 . . 637 . . 641 . . 647 . . 653
. .
657
. . 661
. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
665
I.
Systems
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NEW TECHNOLOGIES IN THE LIGHT OF MATERIALS Prof. Dr.-Ing. Heinrich A. Flegel DaimlerChrysler AG D-70546 Stuttgart, Germany
ABSTRACT The industrialized nations' prosperity and living standards depend to a major extent on the rapid transfer of technical innovations into marketable products. Alongside classical materials - metals, plastics, ceramics, and glass - tailor made combinations like composite materials play an increasingly important role. In the transportation sector materials are expected to render a major contribution towards solving the target conflicts which arise from contradictory customer requirements: high performance and reliability, and minimum weight and low costs. The rapid transfer into new products can only be achieved, if the development of materials and the pertinent production processes and system development are closely coordinated and synchronized by a simultaneous engineering approach.
Sustainability of transport and other sectors is now a major goal for the EU and progress is required. Since fossil fuels will not be available forever in limitless quantities, alternatives will have to be found to meet the future energy demand. PEM (proton exchange membrane) fuel cells are expected to be a suitable alternative for future mobile applications. As shown in fig. 2, the electrical energy can be generated from hydrogen, methanol or hydrocarbons. Hydrogen and methanol can be generated from fossil as well as from renewable primary energies. Therefore, both items increase the security of supply and help to reduce carbon dioxide production. Key success factor is an affordable polymer membrane, which is coated on both sides with a catalyst containing platinum [ 11.
INTRODUCTION It is evident, mobility is a basic need of mankind. Efficient, effective and flexible transport systems are essential for economic activity and quality of life. Sustainable mobility is a megatrend of the future and therefore a major driving force for the development of new technologies. The awareness that resources are limited is growing. Therefore, also political pressure towards resource conservation and waste reduction has significantly increased. Fig. 1 shows the prevailing trends in the EU and USA, two very important segments of the world markets.
-
Awarenessthat resourcesare Limited is growing in aUape smnger than in the US.-"berefon, also political pressure towards resource CMlserVationand waste reduction bas s i g n i E d y i n m w d
Ewopean end-of-live legislationfor vehicles almost finalized conrtatoneE are fixed: myding quota 95% in 2015, free of
rnal@zzzg
the amouot of new materials in the r) eahance use
uroducts will come into the w l i t i d focus brncycling and renewablek t e r i a l s
Curreatly new move towards more product related envinnunentallegislation (Integrated Roduct Policy, IPPt
Fig. 2:
Fuel Cells in Mobile Applications
Very often in mobile applications, progress in the development of new, affordable materials is crucial to the economical success in volume production.
AUTOMOTIVE LIGHTWEIGHT ENGINEERING Fuel efficiency regulations and the voluntary commitment of the European Automobile Manufacturers Association (ACEA) to reduce fuel consumption by 25% over the period 1995 to 2008 are the main drivers for automotive lightweight engineering. Fig. 3 shows the amount of fuel saved per 100 km by 100 kg weight reduction, specified for cars powered by internal combustion and Diesel engines.
.r)End-of-Liivc legislationnot in ptace yet, but Likely in the htture
=rl)No strong legislationfor free of charge iake back d goods, c m t infrashucm supports dismaattiag indusby Fig. 1:
Resources and Environment
3
1990:
I
5 1200 kg
It is very important to keep the complete process chain in mind: from the raw material to the finished part. Since lightweight design doesn't mean lightweight at any cost.
2010: c lo00 krr
Fig. 3:
has excellent die casting properties, it can be used for highly integrated die castings. Plastics are widely used for bumpers and interior parts like dashboard, heading and rear-window shelf.
Automotive Lightweight Design
As a secondary effect, a lighter car needs a less powerful engine to achieve the same performance and the brakes as well as the load-bearing components do not need to be as strong. When these effects are also taken into account, the bonus offered by lightweight engineering is seen to increase to about 0.6 liters of fuel saved over the same distance for the same weight reduction. The average car weight will presumably be reduced by about 20% within the next 10 years. With regard to the distribution of materials, the percentage of steel will be steadily decreasing in favor of lighter materials like aluminum, magnesium and plastics. However, there has also been considerable transformation and improvement with regard to the individual steel quality used. There is a strong trend towards higher and high-tensile strength steel qualities. It is obvious that lightweight engineering is the major driving force for the development of new materials in the transportation sector. However, it can only be completely successful, if new materials, new design principles and suitable economic high volume manufacturing technologies are combined.
Carbon fiber-reinforced plastics (CFRP) are well known as structural materials from formula 1 racecars. CFRPs have the potential to significantly reduce weight. These laminates do not corrode and are almost immune to material fatigue. They show excellent crash behavior because of their high energy absorption ability. Advaotage * Weight saving * High strength and stiffness * Excellent crashbehavior becauseof high energy absorption Disadvantages
costs
* No experience in volume production
Fig. 5 :
Carbon Fiber Reinforced Plastics
However, their use is still limited to small scale feasibility studies and high performance applications (fig 5).
Fig. 6:
Textile Technologies
For automotive applications cost effective manufacturing technologies for volume production have to be developed [2]. The necessary development steps are shown in fig. 6. Fig. 4:
Trend for Body Concepts in Car Design
As shown in fig. 4, multi material design is the trend of the future for body concepts in car design. Because of high tensile strength - which is good for deformation energy absorption - steel is used in crash relevant sectors of the car body. Extruded aluminum profiles are used for frame structures; with aluminum die castings, high functional integration can be achieved. Magnesium
4
High temperature applications: titanium aluminides have the potential to substitute the heavy and expensive nickel based alloys in aircraft engines. The turbine blades should be capable of withstanding temperatures of more than 1200" C. Success factors are their low density and high creep limit. Compared to ceramics ductility increases with increasing temperature[3].
aeroengine
automotive powertrain
PW4086Engine
Turb0charge.r
clhnrst Range:300-450 kN)
VdEnginc
I
I
commnents
- Low Pressure Turbiie
- Blades High Pressure Compress Blades - or HPC-Castings
Engine pans
-
Fig. 7:
Titanium and Titanium Aluminides in Engines
Potential applications in automotive engineering are exhaust valves, piston rods and turbine wheels (fig. 7). A weight saving potential for these rapidly moving inertial masses of up to 50% was proven.
CERAMICS IN AUTOMOTIVE AND AEROSPACE APPLICATIONS
Fig. 9:
Ceramic High Performance Brake (CBrake)
The ceramic high performance brake or C-brake is a brake disc made of fiber reinforced ceramics. The benefits compared to a typical grey cast iron brake disc are: 60% weight reduction, no corrosion, high performance and comfort, and lifetime corresponding to the useful life of the car (fig 9). The C-brake will be introduced in the market in autumn 2000 in a special edition of the Mercedes S class coup6 by AMG. Fiber reinforced ceramics for hot spacecraft structures (fig. 10): C/SiC is an excellent material for lightweight and heat resistant structures [ 5 ] .
Ceramics are stable at high temperatures, resistant to chemical attack, lighter than steel but more brittle than metals or polymers. Fiber reinforcement is a well known method for reducing brittleness. The benefits of ceramics are concisely listed in fig.8. The favored automotive applications are carbon pistons, brake discs and tribo-coatings. The all-ceramic engine that was enthusiastically promoted almost 15 years ago has faded from the scene. Potential Components Bnke discs Piston (carbon) Tribo-coatings Fig. 10:
Benefits Weight reduction High temperature resistat Low wear Adjustable friction behav Long lifetime Fig. 8:
Ceramics in Automotive Applications
C/C-Sic composite materials offer great advantages for new lightweight and wear resistant brakes [4]. These CMC materials have demonstrated their high potential for the application in disc brakes for high speed trains and cars.
Aerospace Components from Fiber Reinforced Ceramics
It enlarges the application range of fiber composites by opening up the high temperature field of up to 1600"C. By using this material, the spacecraft industry can manufacture heat shields and engine components with longer lifetime, lower operating cost and reduced weight; e.g. replacing the metal by C/SiC in the expansion nozzle of the ARIANE upper stage engine would allow 60% weight savings. First ground combustion tests have been successfully performed just recently. Smart materials, either piezoceramics or shape memory alloys, are able to act as sensors or actuators (fig. 1 1).
5
can a be injected more precisely, which leads to a more quiet engine running without the diesel knock.The high precision injection leads to cleaner combustion and lower emissions.
Advanced funclionalites by
- PiezoCeramics - Shape Memory Alloys(SMA)
CONCLUSIONS Electrically kqeramics) or thfmyUy SMA) mQ$ledshape basts for the appllcatlon ok smart materials as Sensor materials (sensing conditions) Actuator materials (performing actions)
changes
areL
Fig. 11:
Smart Materials - Function Concepts
Piezo-active structures are based on structurally integrated piezoelectric actuators to enable active shape control, vibration damping and precision positioning. Aerospace structures usually have to realize high strength and stiffness at minimum weight, and are therefore made of CFRPs. Piezoceramics are seen as most promising materials for low fiequency vibration damping and shape control of lightweight aerospace structures [6]. Due to piezo-driven control flaps in the rotor blades, the rotor induced vibrations inside the cabin as well as the rotor noise during approach could be r duced significantly (fig. 12).
To sum up briefly, four major statements can be made: 0 There is an increasing demand for lightweight materials and the suitable manufacturing technologies in transportation. 0 There is a fierce competition among different materials. 0 Ceramics are preferably used in high performance applications. 0 Application in volume production only at reduced costs. 0 The key issues of structural ceramics are - reliability - interface to surrounding, non-ceramic components - costs
ACKNOWLEDGEMENTS Partial funding of these research activities by the German Ministry of Education, Science, Research and Technology (BMBF) is gratefully acknowledged.
REFERENCES
Piezo-driven control flap of helicopter rotor blades
Fig. 12:
BdtS 50 % noise reduction duringappmch
90 % reduction of rotor-inducedv i i o n s inside the cabin
Smart Materials - Helicopter Rotor Blades
The specifications for a piezo-controlled high pressure injection valve for a common rail diesel injection system are shown in fig 13.
(1) K. Kordesch, G. Simader, Fuel Cells and their Applications; VCH, Weinheim (1996) (2) J. Brandt, K. Drechsler, F. Strachauer, Kostengun-
stige und groherienfahige Herstellung von Verbundwerkstoffen mit textilen Faserstrukturen, in Kunststoffe im Automobilbau, VDI Verlag, Dusseldorf (1998) 227 (3) N. Eberhardt, A. Lorich, R. Jtirg, H. Kestler, W. Knabl, W. Kiick, H. Baur, R. Joos, H. Clemens, Pulvermetallurgische Herstellung und Charakterisierung einer intermetallischen Ti-463 Al4(Cr,Nb,Ta,B)-Legierung, Z. Metallkunde 89 (1998) 772 (4) W. Krenkel, R. Renz, B. Heidenreich, Lightweight
and Wear Resistant CMC Brakes, Proceedings of 7" International Symposium Ceramic Materials and Components for Engines (2000) ( 5 ) W. Schafer, H. Knabe, W. D. Vogel, Fiber Rein-
forced Ceramics for Hot Spacecraft Structures, ICCE 7, Denver (2000) (6) U. Herold-Schmidt, E. Floeth, B. Last, W. Schafer, H.W. Zaglauer, Qualifications of Smart Composites for the Use in Aerospace Applications, SPIE Conference on Smart Structures and Materials, San Diego (1997) Fig. 13:
Smart Materials - Function Concepts
Because of the higher switching speed of the piezocontrolled valves compared to solenoid valves the fuel
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THE PROMISE OF SOFC POWER GENERATION TECHNOLOGY D.S. Schmidt* and R.A. George SOFC Power Generation Siemens Westinghouse Power Corporation 1310 Beulah Road Pittsburgh, PA 15235 USA
ABSTRACT The tubular solid oxide fuel cell (SOFC) technology developed by Siemens Westinghouse is rapidly approaching commercial readiness. Cell technology has progressed from 36cm active length cells producing 30W/cell in the mid-1980s to our commercial size (150cm active length) cell producing 160W/cell in the late 1990s. Further cell technology advances currently under development related to power output enhancement and cell cost reduction will also be discussed. Our SOFC stack and power systems technology has also advanced quite substantially from a 3kW unit operated on H2/C0 in the late 1980s to a lOOkW combined heat power unit operated on pipeline natural gas in the late 1990s. In addition, the first pressurized SOFC/gas turbine power system rated at 220kW will be in operation in the spring of 2000. An overview of the power system demonstrations will also be presented.
INTRODUCTION The Siemens-Westinghouse Power Corporation (SWPC) tubular SOFC design is currently being evaluated in prototype generators, but commercializationwill require a large reduction in cell and generator costs. This focus has been applied throughout the entire system, but one of the greatest potentials for cost reduction lies in the development in low cost cell manufacture, a function of both the processing and raw materials. Current State of Technology Currently, SWPC has two operating field units: one lOOkW unit in operation and one 22OkW-unit undergoing site acceptance testing. The lOOkW unit has been operating for 13,000 hrs as of mid-June, 2000 on natural gas. Located in the Netherlands, it has been delivering 1lOkW AC (130kW DC) as well as hot water to the district heating system. The efficiency of this atmospheric pressure unit is 46%.
A second unit has been constructed for southern California Edison (SCE). The SCE unit will produce 220kW continuously by utilizing a combined cycle with 3 atm operating pressure. When combined with a microturbine generator (MTG), the SOFC is utilized as a high-efficiency burner, delivering clean exhaust gas (no NO,/SO,) to the gas turbine. SOFCs will deliver higher
powers at higher pressures, and in combination with a turbine, overall efficiencies are expected to be at or above 60% with essentially no emissions of CO, NO,, SO,, and very low levels of CO1 emissions due to the high efficiency. Both units utilize the SWPC seal-less tubular design operating near 1000°C. This design allows for small footprints (the prototype SCE unit has a foot print of approximately 25m2) as well as eliminates the problems associated with seals between air and fuel at 1000°C found in planar technologies. As with any new technology, the fxst successful prototype SOFC involved large amounts of developmental and first-time engineering and processing costs. One of the main aspects of achieving commercial success is the technical development of the materials and processes necessary to achieve lower SOFC cell costs. Cell costs comprised 55% of the overall unit costs for the first prototype, and need to compromise no more that 25% of the overall mature product costs. Costs in this high-temperature system such as piping, insulation, recuperators, gas turbines, valves, power electronics, and controls also need to decrease. The cell costs represent the biggest potential for generator cost reduction as production increases due to the ability to directly affect the processes by which this high-tech device is manufactured. Therefore, this paper will focus on the reduction of cell costs through manufacturing improvements.
TECHNOLOGY Basic Materials As commercialization efforts progress, the basic concept remains the SWPC tubular cell. The cell consists of four basic parts: the cathode, the electrolyte, the anode, and a cathode interconnection. The backbone of the tube is the cathode, which operates in an air environment and supports the layers of interconnection, electrolyte, and anode, which are placed on it. The interconnection provides an external connection to the cathode for generator assembly, described below. The basic requirements of the SOFC parts are shown in Table 1.
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conductive
in
Because the entire exposed surface w ill be covered, and a direct connection to the cathode is desired through the LaCr03 stripe previously deposited, a sacrificial material is used to mask the interconnection, which is removed after the deposition. The key to the tubular design is the gas-tight nature of the tube, imparted by the electrolyte and the interconnection. These two surfaces, to maintain this integrity, must overlap as shown in Figure 1. The mask on the interconnection is made to be 7.5mm in width, and leaving lmm at each end. This allows the growth of the EVD electrolyte layer over the edge of the densified interconnection, ensuring a gas-tight seal.
both air and fuel
An important part of the material requirements is the matching of the thermal expansion coefficient to YSZ (10.0~10-~ d m m “C)“). Without such a match, the internal stress buildup at the material interfaces will result in cell damage as it is brought from room temperature to its operating point of 1000°C.
Cell Construction Currently, to form a single cell, doped LaMnO3 is mixed with cellulose binder and water to form an extrudable paste. This paste is first extruded from a piston extruder to form a rounded closed end with a radius of 11.6mm and wall thickness of 2.2mm. When the end is formed, the extrusion is continued to form the cylindrical portion of the support tube with dimensions of 181cm in length and 2.2cm in outer diameter. This results in a hollow single body with one end closed. (See Figure 1.) The binder is burned out from the tube and the powder is fired at 1500°C to form the porous gas-diffusion electrode with the mechanical strength to support the subsequently applied layers. Onto one side of this is plasma-sprayed doped-LaCr03 in a stripe that is l00p.m thick, 1lmm wide, and 150cm This plasma-sprayed stripe is the in length. interconnection which will allow an external conductive path to the cathode. Because it will reside in both the reducing fuel atmosphere and oxidizing air atmosphere, it must be a fully dense body to prevent gas leakage. The next step is the application of the electrolyte to the cell. This is performed by electrolyte vapor deposition (EVD)of YSZ to the surface of the tube. This process involves the use of YC13 and ZIC4 at two torr and 1300°C on the outside of the tube while fluxing H 2 0 and 0 2 through the tube wall. This process is covered in detail el~ewhere‘~,~’, but involves first the closing of pores at the surface of the cathode by CVD of the chlorides in contact with the oxygen. Once the pores have been closed, the mixed conductivity of the CVD YSZ allows the passage of oxygen ions to the outer surface where they can continue to react with the chlorides, resulting in a perfectly dense and uniform 40pm layer over the entire exposed surface of the tube.
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Figure 1: Two-part figure of SOFC cell showing construction and orientation. Closed end not shown in this feure. The last step is the application of the anode. The ideal material for this is porous nickel (Ni), but Ni has a thermal expansion coefficient of 13.46~ 10amm/mm 0C(4)which would result in separation when in contact with YSZ. However, ‘fixing’ Ni particles in a YSZ mounting material or cermet has been successful for S W C and At S W C , this ‘fixing’ of the anode is performed by first dipping the tube into a Ni slurry, placing a porous layer of Ni on the electrolyte surface. The interconnection is again masked, and EVD is again performed on the tube. Growth of the YSZ occurs around the Ni particles, enveloping and ‘ f i n g ’ them in a l00p.m thick porous matrix to the electrolyte surface, providing intimate electrical contact and mitigating thermal expansion mismatch. Estimated conductivities of the anode and subsequentlyapplied layers are shown in Table 2. The last step in producing a working SOFC is ensuring good electrical contact with the interconnection surface. Ni plating has proven effective, providing a
very thin Ni layer with intimate contact, to which additional nickel bonds may be made. Cell part
Electrical Role
Resistivity
Thickness
at 1OOO"C
Cathode Electrolyte Anode Interconnection
Porous gas-difision electrode for oxygen reduction Fully dense gas barrier, passing 0'ions only Porous gas-diffision electrode for fuel oxidation Fully dense gas barrier, electrical conduction only
0.0 12 Q c m
2.2mm
10~cm
40p
-48 x 10*Q-m(Ni)
1oOMm
1Q c m
lOOp
Bundle Construction Cells, once made, must be arranged for efficient space usage and convenient electrical contact. Multiple cells are formed into three-by-eight (twenty-four cell) bundles. Electrical connections are made by the application of Ni felts to the NiNSZ anode and Niplated IC. Electrically, three parallel path are provided, and the voltage across any bundle is increased by 8x the cell voltage. (See Figure 2) Fortyeight bundles are joined to form a complete power unit such as the 1lOkW atmospheric or 220kW pressurized units, which both uses 1152 cells.
$300/kW. In the past, SOFC development and cell construction has largely been performed by highly labor-intensive processes. With the small number of cells produced per year, material throughput was low, resulting in high material costs for SOFC production as well. The combination of these two factors led to an actual cell cost that was well beyond a marketable value. This must change for commercialization.
Figure 3: Cell potential at 300mA/cm2, 83% FU over 34,000 hrs. with 11% H20/89%H2 atmosphere. Manufacturing Advancement Figure 4 shows the past, current, and projected future cost per cell. Table 3 describes the processing path to achieve these goals. Three actions are primarily required volume increase, automation, and vertical integration of the manufacturing process. The specialized ceramics required for SOFC operation can be costly due to the purities required. However, by negotiation of long-term contracts with large volume requirements, significant price reduction (lO-SOo/o) can immediately be achieved.
Figure 2: Actual 8x3 bundle of SOFC cells. 7-
Another strong indicator of cell performance is voltage versus time for a constant current density and fuel utilization (expressed as a percentage). This is shown in Figure 3. This graph indicates a very low degradation rate of O.lo/dlOOOhr. At this rate, after 40,OOOhr operation, the cell would be expected to have lost only 4.0% of its power, or about 6.3W.
.--A-L----
These graphs represent the highly reliable prototype developed.
MANUFACTURING/TECHNICAL ADVANCEMENT For commercialization of the technology, the cost of the product must be lowered to a level of below $2000/kW, of which cell costs must be less than
Figure 4: Actual and projected costs for a cell and power system.
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Labor intensive processes must be highly defined and automated for quality control as well as cost reduction purposes. By increasing volume while labor remains constant or reduces, the labor costs per cell can be dramatically reduced by as much 95% if high volume levels are reached.
I
volume
automation 25%
However, the largest cost savings will be obtained by performing previously subcontracted processing inhouse. The first step in this movement up the value chain can be seen in the progression of air electrode processing. In the past, an outside vendor was contracted to deliver to SWPC a complete and processable cathode support tube. This type of tube cost approximately $1600/tube. Recently, SWPC has brought in-house the ability to extrude and form the cathode support tube starting with LaMn03 powder. This ability has been primarily responsible for drop in cost shown in Figure 5. The next step, due to weight, is the in-house production of the cathode powder, as is shown in future processing (Table 3). The cathode represents 92wt% of the completed SOFC cell.
Figure 5: Relative contributions of increased volume, automation, and in-house process control.
A second major advantage to in-house production is the ability to optimize the overall process due to our control over all the processing steps. Changes and process improvements can be quickly implemented, with careful control and monitoring of results, allowing quality feedback not only within the individual step but between steps at the facility. While feedback with an outside vendor is possible, it is difficult to reach a communication level comparable to that developed internally. The relative role of all these cost-cutting methods are shown in Figure 5. Process Cells/var Cathode
I Past I 1000 Received tube
Current 10,Ooo
Receive powder Produce tube
Interconnection
Received powder Plasma sprayed
I Efectrofyte I Received I I powder
Mask EVD Remove mask
Future lO0,OOO
-
Receive raw materials Produce powder Produce tube
om,po3msm,sn,an,wgn,sm,momosm,mowo
Receive powder
Make powder
Figure 6: Operating History of EDBELSAM
Automated plasma spray
Automated plasma spray
The role of manufacturing improvements is to achieve a reliable product at acceptable cost through the use of the tools of volume increases, automation, and inhouse processing.
I Receivepowder I Makepowders I I Mask EVD Remove mask
Automated application
I I
SYSTEM STATUS Currently, SWPC has two generators in operation: the 220kW SCE pressurized SOFC/gas turbine (PSFOFC/GT) unit and the lOOkW EDBELSAM combined heat and power (CHP) atmospheric unit. Figure 6 shows an operating graph for the EDBELSAM unit over its lifetime.
Ni-plating
10
Plated single cells
Plate small batches of cells
Fully automated plating system
EDB/ELSAM 100 kW Delivered in 1997 to EDBELSAM (a Dutch/Danish utility consortium) this unit has provided electricity and district heating to a substation in Westervoort, Netherlands. Through 2000, it was the largest operating SOFC system, producing between 1051lOkW AC to the utility grid and 65kW to the hot water district heating system with an average electrical efficiency of 46%. There are two stages of operating shown the in the graph, which are termed "Build 1" and "Build 2." Build 1 was the originally delivered unit, started in November 1997. The unit operated unattended for 3700 hours at 42% efficiency. This was significantly below the analytical predicted efficiency of 47%. Temperature and localized voltage anomalies were observed, and air leakage to the fuel side of the cell was suspected due to measured fuel compositions at the cell exit. Due to these observations, the unit was shut down for inspection and repair on June 26, 1998, despite the fact that no degradation of the terminal voltages was observed, except for a sulfur-poisoning incident that was corrected by replacement of the desulfurizing reagent. The unit was returned to Siemens Westinghouse Power Corporation for evaluation. It was determined that there existed a failure of the baffle boards at the open end of the cells. Since this allowed air leakage into the fuel feed, the fuel side of the cell was partially oxidized, especially as the cell was cooled. Also, it was noted that there was a partial loss of nickel contact between some of the cell within bundles. The unit was rebuilt with improved baffle boards and cell bundles, and sent back for restart in March 1999. Almost immediately, significant improvement was observed, with efficiency increasing from 42% for Build 1 to 46% for Build 2, very close to theoretically calculated efficiency predictions. DC efficiency is actually about 53%, with the disparity between the two due parasitic loss of system loads (7-8kW) and inverter efficiency (92.5%). On June 1, 2000, Build 2 reached one year operation with >99% availability. Overall, the unit will reach two complete years operation in November 2000.
efficiency will be approximately 57%. shows operating parameters.
Table 4: 0 eratin arameters Cell Current CellVolta e Pressure Ratio SOFCDC ower SOFC ossAC ower Gas turbine AC power System net AC power Efficiencv (net AC/LHV)
Table 4
267am s 0.610 V 187 kW 176 kW 47 kW 220 kW 57Y0
CONCLUSION Siemens Westinghouse Power Corporation has shown technical viability of the SOFC in both a lOOkW atmospheric and 220kW pressurized system in combination with a microturbine generator. The biggest hurdle currently facing commercialization of this technology is the development of an economically acceptable system. Key to the lowering of capital costs is the development of a low-cost cell production process. The three major categories available for lowering cell costs are increased volume, automation of production system, and process control from raw materials through finished product. Siemens expects to reduce overall costs significantly through these three areas, making possible the complete commercialization of the SOFC power systems.
REFERENCES Shackleford, J.F., Introduction to Materials Science for Engineers. 2"d Ed.. Macmillian, New York, 1988, p. 372. Pal, U.B., Solid State Zonics, 52 (1992) 227. Isenberg, A. O., Solid State Zonics, 3/4 (198 1) 43 1. Hodgeman, C.D., Weast, R. C., Selby, S. M., Handbook of Chemistw and Physics: 42"d Ed., Chemical Rubber Company, Cleveland, 1960, p.2242. Primdahl, S., and Mogensen, M., J. Appl. Electrochem., 30 (2000) 247. Steel, B. C. H., Solid State Zonics, 86-88 (1996) 1223.
SCE 220kW A 220kW unit recently completed construction for Southern California Edison (SCE). This unit utilized a high pressure SOFC (PSOFC) at the bottoming cycle in combination with a microturbine generator (MTG) from Northern Research and Engineering Corporation. After completing lOOhrs successful operation at SWPC in April 2000, it was shipped and installed at Irvine, California. The PSOFC operates at 3-atm pressure, with the pressure provided by the compressor portion of the MTG. With approximately 20% of the electrical power provided by the MTG, the overall unit
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CERAMIC MATRIX COMPOSITES FOR DISK BRAKES AND THEIR MANUFACTURING TECHNOLOGIES R. Gadow University of Stuttgart Institute for Manufacturing Technologies of Ceramic Components and Composites Allmandring 5b D-70569 Stuttgart GERMANY
ABSTRACT Thermally stable and corrosion resistant light weight components are a central challenge in modem automotive engineering. The competition in the automotive industry, especially for heavy, high performance, luxury and sports cars, demands retardation performance, comfort and all weather braking ability for new disk brake systems. The required mass reduction with simultaneously improved performance and durability in modem trucks and high speed trains requires disk materials with life time corrosion and wear resistance. The reinforcement by short, chopped and endless carbon fibers results in fracture toughened ceramic matrix composite (CMC) properties with appropriate friction and reliable mechanical properties in comparison with conventional materials. Reaction bonding is a competitive technology for manufacturing of near net shape formed components like brake disks. A review is given on chemical processing, manufacturing and design as well as first application results of these new refractory CMC components in brake technology are shown.
fast, reliable and reproducible compound manufacturing methods fast forming process with serial manufacturing capacity net shape sinteringkhemical transformation to CMC minimized grinding and finishing effort short production cycles high temperature and corrosion resistance combined with CMC fracture toughness steady friction coefficient for antilocking brake system under severe road vehicle conditions. Different types of chemical processing and manufacturing of components for brake disks can be used. With a background in aircraft systems chemical vapour impregnated carbodcarbon composites (CFC) were the first successful materials, followed by pitch and precursor impregnated cheaper CFC disks and rotors. As they cannot fulfil the main requirements for low cost passenger car application, the recent development is focussed on silicon ceramic composite with SiSiC or other reaction bonded matrices.
c
to Final Statc CMC,
Iclcansurface, n u r ncr shop
TECHNICAL AND ECONOMIC REQUIREMENTS FOR CMC BRAKE SYSTEM COMPONENTS Due to their limited corrosion resistance CFC materials are not suitable for long term operation temperatures above 500 "C in atmosphere('.*', so that their application will be limited to aircrafts and racing cars, where pure performance overcomes cost. Low melting and softening temperatures limit the application of coated light metal alloy components and MMCs to temperatures below loo0 OC. Only CMC provide sufficient strength and corrosion resistance up to higher temperatures. The application of CMC materials, e.g. as friction materials in brake systems of passenger cars and trucks, depends on the fulfillment of the following industrial system requirements (3): 0 low cost raw materials and additives
Fiber Preform., GrcmCompoct
1 II PCS!PLI"
Xomal h n n c
RRSC
RB-S6
ZC-Lonp F i h
Shon Fiber
Tecbnksl Process Variant
fig. 1
manufacturing cycle times for CFC and CMC composites('2,
For the production of carbon fiber reinforced S i c ceramics different methods are used. Usually 2D C-fibre woven fabrics with phenolic resin as C-precursor or temporary binder are used for reinforcement. Prepregs are laminated to a fibrous preform and the Sic-matrix is formed by chemical vapour infiltration (CVI)(u6' or by infiltration of organosilicon polymers which are pyrolized to ceramics up to about 1600°C under inert gas
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atmosphere (LPI)(4*7*8'. Due to its relatively short production cycles (fig. 1) and consequently cost effectiveness the reaction bonding by silicon melt impregnation (LSI) and reaction in porous carbonaceous preforms is a promising technology for industrial applications (%I2).
PHYSICOCHEMICAL BASICS FOR REACTION BONDED SILICON CERAMIC COMPOSITES Reaction bonding is a process where a strong ceramic body is formed from a porous compact or prepreg by an in-situ chemical rea~tion"~'.For the fabrication of Sic-composites the liquid metal impregnation method is used for reaction bonding of fiber skeletons in originally porous carbon containing matrices. The use of the heterogeneous chemical reaction between silicon powder containing compacts and nitrogen spender phases results in a fine porous or nearly dense reaction bonded silicon nitride matrix composite. Reaction bonding is a cost effective and technically advantageous method to obtain dimensionally stable and precise CMC components in manufacturing of serial products as they are used in the automotive industry. The advantages of the S i c route are fast impregnation and chemical transformation, but a quite sophisticated metal melt technology is needed. The Si3N4 route allows convenient reacting with gaseous species but is more time consuming and leads to less dense products. The silicon penetration rate in carbonaceous preforms is high due to the strong capillary flow effect and the low wetting angle of liquid silicon on solid carbon(I4'.The initial wetting angle is changed to values in the range of 30"...40"during melt infiltration with excess silicon mainly by two processes: The dissolution of solid carbon in liquid metal and the formation of a continuous solid Sic-layer at the interface. The capillary flow effect of silicon is given by Hagen-Poiseuille's equation. For more complex models used by Fitzer and Gadow the pore utilization is taken into account.(4.1 2.15.I6.I7) In pore systems and networks there are various pore radii and a pore shrinkage effect is observed due to the chemical reaction between carbon and silicon during melt infiltration. The combined effect of mass transfer from the liquid reactant Si to the solid surface, its diffusion through the primary formed solid S i c layer and its chemical reaction (first order) with the solid reactant carbon was studied by Fitzer and Gadow using reaction models for non catalyzed heterogeneous chemical reactions between fluid metal and a shrinking core of carbon containing ~olid"~'. It was shown that the process is controlled by silicon diffusion through the Sic-layer resulting in the known parabolic correlation between reaction time and layer thickness. In any case of infiltration by silicon in carbon-carbon preforms the open porosity content and
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consequently the pore radii distribution of the preform are the critical value'"' to control and optimize the reaction. A non optimized porosity and non optimized carbon filler/fiber/SiC powder ratio cause problems like incomplete infiltration or inhibition of reaction, e.g. in the presence of non desired byproducts like SiO which effect the wetting properties. For RBSN-matrices the fiber damage by mechanical degradation in contact with sharpedged Si powders and the drawback of mandatory post impregnation by metals or liquid precursors must be mentioned.
FUNDAMENTALS FOR CERAMIC COMPOSITES
fig. 2
interacting factors determining the composite
proper tie^"^' The general aim in CMC development is the improvement of the thermomechanical behavior of the matrix ceramic by incorporation of high tenacity refractory fibers'24* 25'. The performance of such composites strongly depends on the fibedmatrix interaction like adhesion, chemical and mechanical bonding at the interface as well as on the fiber volume content (fig. 2) and the fiber orientation, distribution and geometrical arrangement (fig. 3).
fig. 3 fiber arrangement variations For short and chopped fiber CMC the critical fiber length has to be reached or exceeded to provide the desired stress transfer from matrix to fiber'**'. The reinforcement mechanism of brittle fibers in brittle
matrices is also dependent on the fiber content. If this content is lower than the critical one the composite fails by spontaneous brittle fracture (fig 4). The fiber volume content to obtain a damage tolerant mechanical behavior under load must exceed a critical value (- 20...30 vol. %). Short fiber structures are used for cost effective products and need special compounding technologies. Unidirectionally and multidirectionally designed CMC show improved mechanical properties, but are not gas and fluid tight or fully oxidation resistant. A recently developed novel concept consists of a short fiber reinforced layer with a high ceramic matrix content at the surface of the component and strong endless fibers in the core (fig. 3 and fig. 10). singulary
multiple
.......fnctun .............
fnctun
oBF
............................................
b
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.....
.......
................... ........... ,.-”’
u)
........
‘BC=‘BM
‘BF
bending strain E
4F=o
4‘F
+F=l
fiber volume content c$F
fig. 4 stress strain behavior of composites with brittle matrix(26’ To control the relatively harsh conditions during reaction bonding, protective coatings for the carbon fibers are necessary in most cases. These coatings can be produced e.g. by chemical vapor deposition. Refractories like TIC, TiN, S i c and pyrolytic carbon can form a temporary diffusion barrier“3’, but because of cost arguments other processes like the cheaper liquid phase fiber coating”’ or the completely different approach by reducing the chemical activity of the alloyed silicon melt are to be preferred(”, 23’.
MANUFACTURING TECHNOLOGIES FOR REACTION BONDED CERAMICS The manufacture of fiber reinforced preforms in general involves two stages. Incorporation of fibers into the unconsolidated matrix material, followed by a consolidation of the matrix.(”’. The highest priority for reaction bonding processes is the manufacturing of a well designed preform porosity, which is obtained during pyrolysis of carbon containing resins, and an advantageous fiber arrangement. In general two kinds of preform manufacturing technologies are used which are followed by the high temperature treatment: conventional ceramic manufacturing methods and modified compounding techniques of the plastics industry.
Conventional ceramic manufacturing methods (Slip Casting, Mixing, Granulation): In the case of slip casting a slurry is poured into a porous mould which is covered by the fiber skeleton. The mould absorbs the liquid carrier, but the densification is hindered by the fibers and a high porosity is the result. The matrix deposition can also occur on the fiber skeleton which acts as a porous medium itself resulting in a porous preform, which is compressed by hot-pressing. Another possibility to obtain a fibrous preform is to fill pellets consisting of fibers, powders and binder in a mould. For manufacturing of these pellets, either mixing technologies (e.g. kneader with extrusion screw) or build up granulation processes are used(20’.Depending on the binder system the pellets are either cold- or warmpressed to manufacture the preform. If carbonaceous precursors are used as binder a reaction bonding (RB) takes place after pyrolysis converting the carbon preform by a chemical reaction with liquid silicon into a SiCceramic (LSI-liquid silicon infiltration). In order to obtain a high matrix content n reinfiltratioddensification cycles can be used (21). In the case of using organometallic precursors as binder, the pyrolysis is followed by a crystallization of the glassy matrix without RB-process (LPI-liquid polymer infiltration). By using silicon powder based porous preforms a reaction bonding to RBSN under nitrogen atmosphere is possible.
Modified compounding techniques adapted from polymer manufacturing: Long fiber and woven fabrics reinforced preforms are usually manufactured with these technologies mainly used for thermoplastics and thermosetting resins. The matrix consolidation is either performed in one step with incorporation of fibers into the matrix (Autoclave Process, RTM) or in a separate step. Autoclave and RTM pressure forming processes: These processes allow the production of large sized, thin walled articles in medium production quantities. For an autoclave process heated and pressurized forms are used. After coating the form with resin mixture, the fiber reinforcement and, if appropriate, the core material is applied as an insert. The counterform, which is also coated with resin mixture, is used to form an airtight seal. Alternatively the counterform can be a removable foil. In the resin transfer moulding (RTM) process, two-component or multicomponent rigid moulds are employed, which are opened and closed by hydraulic mechanisms. The fiber reinforcement is cut to shape and placed in the mould. The mold is closed and the resin mixture is injected from the bottom under pressure. The application of vacuum can assist the impregnation state. The injected resin must be of low viscosity and highly reactive.
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Filament winding and slurry infiltration of fiber rovings: The equipment for filament winding consists of a rotating winding mandrel and an oscillating fiber guide, the drives of which can be both regulated. Rovings are generally used for reinforcement, although woven tapes can also be employed. The rovings are fed through an impregnation bath, containing a mixture of resin and powder fillers. The impregnated fibers are then wound onto a core under constant tension and with exact geometrical pattern, governed by the ratio of rotation speed of the winding mandrel and the desired winding pattern. This technology is a highly reproducible process, but only suitable for the manufacture of hollow articles. For the production of long fiber reinforced sheets, the laminates on the winding mandrel are cut into pieces, laminated to two-axial structures (mostly O', 90') and reimpregnated with a resin paste in an autoclave. Another possibility is to use a solution, containing metal compounds ( e g metal alkoxide, acetate or halide), which is reacted to a sol. This sol-gel technique has been used mainly to produce oxide and glass ceramic matrices. The modified filament winding process is theprecursor route for manufacturing long fiber reinforced structures'22' Sheet Moulding Compound (SMC) Technology: The SMC (Sheet Moulding Compound) is transformed from its liquid, fiber and powder filler ingredients into a sheet product that is usually about 4 mm thick and 1SO0 mm wide. The SMC process is illustrated in fig 5.
simple and cost effective incorporation of short and long fibers in one combined process. Before molding the SMC sheets have to be cut into pieces of predetermined size and shape. The cut pieces are subsequently laminated and assembled into a charge pattern. The combination of short and long fiber reinforced sheets opens up the possibility to realize a great spectrum of suitable structures for different applications. The charge is finally placed in a preheated mold and compressed. A near-net-shape forming process is possible due to the high mold filling capacity of the prepregs and the precision of the mold.
High temperature treatment and thermal processing The fiber-matrix incorporation and consolidation step is always followed by a pyrolysis resulting in a porous preform due to mass loss and the resin shrinkage. As mentioned before, it depends on the used binder and resin system, if the ceramic matrix is formed by a ceramizatiodcrystallization or a reaction bonding process. An overview of different manufacturing routes is shown in fig 6.
Raw Materials
Compoundirg lncorporahonand Alament of Fibers Consolnlabon of the Matm (Polymemabon, Curing, Thermosethng)
High TemperatureTreatment
fig. 6
.
V
compaction roller
fig. 5 manufacturing process for SMC -prepregs"O' The SMC process creates a resin paste and carbon fiber 'sandwich' which is subsequently sent through a series of compaction rollers where the carbon fibers are wetted with the resin paste and trapped excess air is squeezed out of the sheet. In contrast to e.g. the precursor route the amount of resin is very high, and so there is no reimpregnation necessary. Endless fibers can be obtained as well by SMC technology if no cutter blade is used. Consequently SMC is the only compounding technology which allows a
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manufacturing scheme for reinforced carbonand S i c ceramics
CMC BRAKE MATERIALS CHARACTERIZATION The highest mechanical strength values for RBSic composites (UD: flexural strength 600 MPa) were obtained by use of C V B SiC fibers and S i c coated carbon fibers(''). 2D-composites with woven fabrics are reported to reach 150-250 The mechanical properties of most industrial brake disks must exceed the flexural strength of high quality cast iron (-100 MPa). A comparative summary of friction materials properties is shown in table 1. The key properties for these components are stable and predictable friction behavior and wear resistance. Surface modified CFC disk materials show a steady increase of the friction coefficient p with relatively low initial values in comparison to S i c matrix CMC (fig 7). The carbon and
-
graphite dominated friction values are even inferior to cast iron (GG25) at room temperature. Ferrosilicon alloyed CMC show the most promising friction behavior. The friction behavior depends on the obtained surface roughness and homogeneity, which have been optimized by appropriate machining and finishing operations during manufact~ring‘~~’.
mechanical properties are limited. Examples for short fiber reinforced components are shown in fig 9.
fig. 8
0
5m
l
a
w
)
I
~
2
[
m
2
5
0
3
3
m
Mmerelstmkes
fig. 7
comparison of SiSiC-, FeSi75 alloyed-SicCMC with other brake disk materials(23’ tab 1
friction materials data(23’
)
C/C materials for racing cars (courtesy of SEP and Ferrari)
As mentioned earlier, the SMC technology offers a wide range of possibilities to combine various fiber length and filler contents. Short and reproducible production cycles as well as a net shape forming process, due to the high mould filling capacity of the prepreg sheets, can be realized by the SMC compounding process. The fulfillment of the automotive industry requirements concerning cost effectiveness and reproducibility opens up the possibility for a serial manufacturing method for ceramic brake disks.
COMPONENT DESIGN IN AUTOMOTIVE AND AIRCRAFT BRAKE SYSTEMS I The first introduced S i c CMC brake disks were solid bulk components with 280 mm diameter‘3’. High standard CFC disks and rotors for aircraft and racing car use (fig 8) were made of joined elements up to diameters of 600 mm. The actual commercially available RB-Sic- and carbon composites especially show excellent mechanical and thermophysical properties due to the use of long fibers (tab 1). But on the other side long carbon fibers reinforcements show limited corrosion resistance, because the fibers are attacked not only on the component surface but the corrosion proceeds also inside of the component due to the high fiber length. By using short fibers the isolated fiber strands provide improved overall oxidation behavior, but on the other hand the
fig. 9
I short fiber reinforced brake disks for passenger cars
Realizing combined short- and UD-fiber structures with short fibers at the functional surface of the component could be a promising way to combine both the damage tolerance of long fibers and the corrosion resistance of insulated short fibers with relatively high ceramic matrix volume content (fig 10). A technical design for a 330 mm diameter SIC CMC brake disk with internal ventilation was realized by joining two equishaped semidisks by metal melt infiltration with subsequent reaction bonding (fig 11). The granulate or SMC based semi-monocoques were
17
formed by axial warm pressing in a precision tool. They feature torsional locking by a set of radial keys and slots. The reinforcement design is fitted to the orientation of the mechanically and thermally induced main stresses. The geometry with joined key and slot with a cone shaped design in radial orientation performs an excellent distorsion safety during manufacturing and practical use. The web distance is optimized with reference to the '/z brake pad size providing an advantageous mechanical brake force distribution.
5) 6) 7) 8) 9) 10)
ndomly distributed short fibers
11) 12)
~~
13)
14)
fig. 10 structural design of high performance multilayer CMC friction material
15) 16) 17) 18)
19)
fig. 11 technical design for brake disks with internal ventilation, silicon joined
20)
REFERENCES E. Fitzer: Potentialstudie C/C-Verbundkorper als Werkstoff im Primkkeislauf einer HTR-Anlage, Institut fur chemische Technik, Universitat Karlsruhe D. W. McKee et. al.: Carbon 22,285 f (1984) Gadow, R.; Kienzle, A.,,Processing and Manufacturing of C-Fibre Reinforced SiCComposites for Disk Brakes",Proc. 61h Int. Symp. On Ceramic Mat. And Components for Engines, Arita, Japan, K. Niihara et al. eds., ISBN 49980630-0-6, pp. 412-418 (1997) Fitzer, W. Fritz and R. Gadow, Chem. Ing. Tech. 57, pp. 737-746 (1985)
18
21)
22)
23)
R. Naslain, P. Hagenmuller, F. Christin et al.,J. : Adv. Comp. Mat. 2, pp.1084-1097 (1980) R. Gadow, Fortschrittsber. d. Dtsch. Keram. Ges., Bd. 4 (1989) 5-40 R. Gadow, A. Kienzle; Proc. TAE/TAW coll. ,,Modeme Werkstoffe". pp. 23-24.1997 Esslingen T. Haug and R. Ostertag, in ,,techn. keram. Werkst." chapt. 4.4.1.1, DKG ed. (1995) C. W. Forrest et al.: Special Chemics 5, pp. 99-123 (1970) W. B. Hillig et al.: GEC Tech. Inform. Ser. 74 RD 182 (1974) E. Fitzer and R. Gadow, Am. Ceram. SOC.Bull., 65, 2, pp.326-335 (1986) Fitzer, E Fritz, W.; Gadow, R.: Proc. Int. Symp. Ceram. Comp. Engine, (1983), KTK Scientific Publishers, Tokyo (1983), pp. 505-5 18 E. Fitzer, W. Fritz, R. Gadow: Possibilities for fiber reinforcement of silicon carbide, Advanced Ceramics, S. Somiya ed., KTK Scientific Publishing Company, Tokyo, pp. 81-129 (1987) Li, J.-G.; Hausner, H.: Reactive wetting in the liquidkolid-carbon system, J. Am. Soc.79 [4], pp. 873-880 ( 1996) P. Godard et. al.: J. Appl. Polym. Sci., Vol. 18, pp. 1477-1491 (1974) H. Strohmeier: Dissertation, University of Karlsruhe (1 98 1) R. Gadow: Dissertation, University of Karlsruhe (1986) Singh, M.; Berendt D. R.: Reactive melt infiltration of silicon-molybdenum alloys in microporous carbon preforms, Mat. Sci. Eng., A194, pp. 193200 (1995) D.C. Phillips: Fiber reinforced ceramics, Handbook of Composites, Vol. 4, ed. by A. Kelly and S.T. Mileiko, Elsevier Science Publishers B.V., 1993, ISBN 0 444 864474 Gadow, R; Speicher, M.: Herstellung faserverstiirkter, reaktionsgebundener Siliziumkarbid-Keramiken unter Verwendung intermetallischer Siliziumlegierungen, Mat.-wiss. u. Werkstfftech. 30, No. 8, WILEY-VCH Verlag, pp. 480-486 (1999) M. Nader, et. al.: Herstellung von endlos- sowie schnittfaserverstiirkten C/SiC-Keramiken, Verbundwerkstoffe und Werkstoffverbunde, K. Friedrich ed., pp. 179-184 (1997), ISBN 3-88355250-X P. Greil, M. Seibold, Advanced Composite Material, M.D. Sachs ed., Ceramic Transactions, Vol. 19 (1990) R. Gadow, M. Speicher: Manufacturing and CMC component development for Brake Disk in Automotive Applications, 23d Annual Cocoa Beach Conference & Exposition: B, USA, pp. 55 1558, Transactions of the ACerS, (1999), ISSN 0 196-6219
24) K. M. Prewo: Fiber-Reinforced Ceramics: New Opportunities for Composite Ceramics, Ceramic Bulletin, Vol. 68, No. 2, pp.395-442 (1989 25) R. J. Kerans: The Role of Fiber-Matrix Interface in Ceramic Composites, Ceramic Bulletin, Vol. 68, NO. 2, pp.429-442 (1989) 26) J. Schlichting, Verbundwerkstoffe, Lexika-Verlag Grafenau 1978, ISBN 3-88146-149-3 27) M. H. Van de Voorde, M. R. Nedele: C M C s Research and the Future Potential of CMCc in Industry, 20th Annual Conference on Composites Advanced Ceramics, Materials and Structures: B, Ceramic Engineering and Science Proceedings 4, pp. 3-21 (1996)
28) A. Kelly: Strong Solids, Clarendon Press, Oxford (1966), ISBN 3 528 07703 4 29) U. Papenburg: Faserversmkte keramische Werkstoffe (CMC), Keramische Werkstoffe, DKG ed. (1994) 30) SMC/BMC - Design for Success!, European Alliance for SMC, WDW Werbedruck Winter (1997)
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CERAMIC CUTTING TOOLS G. Brandt
AB Sandvik Coromant, R&D, Materials and Processes, SE-126 80 Stockholm, Sweden
ABSTRACT
they have in some cases also been successfully applied in milling of cast-iron and heat resistant alloys.
Ceramic cutting tools can, successhlly applied, increase the metal removal rate by several times over that obtained with conventional tool materials. Ceramic tools are either based on alumina or silicon nitride and applications in metal cutting are determined by the specific material properties of the tool material in question. An understanding of the wear mechanisms is a prerequisite for further development of these materials. Applications and development trends are discussed.
INTRODUCTION Ceramic materials have always had a potential for being used as cutting tool materials in terms of their high hot hardness and chemical stability. In spite of these advantages they only represent up to 5 % of the indexable cutting tool market. This is mainly dependent on their inferior strength, fracture toughness and thermal shock resistance compared to high-speed steels and cemented carbides. During this century there has been a rapid development of new tool materials and machines for chip-forming within metal cutting, which in case of steel machining has led to, on average, a doubling of productivity every tenth year. Ceramic inserts are not suited for steel machining to any large extent, but ceramic coatings on cemented carbide inserts have no doubt been the major prerequisite for the dramatically increased cutting speed capability of this category of materials. Nevertheless, ceramic inserts based on alumina and silicon nitride, are to-day successfully applied in machining a variety of metals within the engineering industry, including cast iron, hardened steel and heat resistant alloys. Metal removal rates are significantly higher than when using conventional coated or uncoated cemented carbide tools.
METAL CUTTING Metal cutting is an operation where a tool produces thin layers of metal, chips, during the relative motion between the workpiece and the tool. The most common operation is turning, which means that the workpiece is rotating against a stationary tool. This is also the most frequent operation for ceramic cutting tools although
Fig 1. Wear types in metal cutting During the cutting process the geometry of the tool will change as a result of wear to the cutting edge. Toollife, to day is often counted in minutes, and ending when the cutting edge no longer produces acceptable components. Tool life determinant parameters can vary, especially for ceramics it is important to avoid excessive wear, which can cause fracture and unpredictable toollives. Tool wear is the result of different load factors acting on the cutting edge. Metal cutting generates a lot of heat and the thermal load is considerable. Cutting edge temperatures for ceramic tools often exceed 1000 C. Mechanical loads include both static and dynamic components. Due to the high pressures and temperatures along the interface between workpiece material and tool material, diffusion and chemical reactions normally have a significant role in the wear process. Inclusions in the workpiece material can also influence the wear process considerably, hard particles cause abrasive wear where as softer particles can built up protecting layers decreasing chemical interactions. The exposed area of the workpiece material is large, depending on the high cutting speed, even if the inclusion content is very low, millions of particles will pass the cutting edge every minute. Flank wear, which develops on the clearance face and crater, wear on the rake chip face. Notch wear at the extremities of cut is normally observed when machining heat resistant alloys. The mechanisms for a certain type of wear may vary depending on the actual tool and workpiece material combination; these mechanisms will be discussed in more detail later on within this paper.
21
TOOL MATERIALS The various tool materials that are used by the metalworking industry to day are shown in figure 2, arranged according to their hardness and toughness. The figure reflects also to some extent the general application areas since the cutting speed capability increases as we move fiom right to left in the figure at the same time, feed, depth of cut and intermittent cutting capability decreases. The figure also shows the tool material development during this century.
. , ftennrerse Rupture Stre-
Wmm*
Fig 2. The various tool material properties The ceramic materials that are used in metal cutting to day are either based on alumina or silicon nitride. The merit of alumina is mainly the high chemical stability and wear resistance where as silicon nitride shows excellent thermal shock behaviour due to a low thermal expansion coefficient. The various commercial materials available to day will be also discussed more in detail. Alumina-zirconia materials were developed around 1975' and have a much better strength, toughness and thermal shock resistance than pure alumina. The improvements are related to a martensitic transformation of zirconia causing both compressive stresses and microcracking ahead of a propagating crack. The Grades for cast zirconia content ranges from 3-10 YO. iron machining normally have a content around 5 % whereas higher contents are used for machining steel. Alumina-TiC-(TiN)' was developed around 1970 and has a better hot hardness and better thermal conductivity than pure alumina. The hot hardness makes this material suitable for machining of chilled cast iron and hardened steel. A high resistance to attrition wear makes it also suitable for fine machining of cast iron. Alumina-silicon carbide whisker materials were introduced around 1985. A whisker is a single crystal fibre characterised by being very high strength. A 25% addition of silicon carbide whiskers to alumina increases simultaneously strength, toughness3and thermal shock resistance4,to much higher levels than can be arrived when using particle reinforcement. Silicon carbide whisker reinforced alumina is the only alumina-basedmaterial that can with stand the use of coolant during machining without fi-acturing.
22
Silicon nitride was synthesised already during the last century but had not found engineering applications until the last decade. Silicon nitride was one of the primary candidates for high-temperature applications in diesel engines for which very large research efforts were devoted. In approximately 1985 several producers introduced silicon nitride as a cutting tool material. Silicon nitride cannot be consolidated with solid state sintering with different oxide additives such as alumina, yttria, magnesia and zirconia, used to form an oxynitride liquid together with silica, which is always present on the silicon nitride grains, and silicon nitride itself. During sintering the a-Si3N4raw material is dissolved in the liquid and p-Si3N4 is precipitated. The resulting microstructure consists of crystalline silicon nitride grains with an intergranular'phase,which is normally glass. The morphology of the silicon nitride grains 5, and the amount and composition6of the intergranular phase both have a significant influence on the properties of the sintered material. Modem silicon nitride materials for metal cutting have a high proportion of elongated prismatic p-Si3N4grains,often referred to as self reinforced or in situ reinforced. This special structure leads to improved strength and toughness. The dimensions of the silicon nitride grains classify them as whiskers (aspect ratio>3). Sialon is produced using the same raw materials as for silicon nitride but with additions of alumina and aluminium nitride, which are used to form a solid solution' S&Al,ON,-,.where z54,2. The microstructure is similar to silicon nitride with elongated p-Sialon grains embedded in an intergranular glassy phase. The main advantage of Sialon is the increased chemical stability with increasing z-value. However, it's strength, toughness and also thermal shock resistance decrease with increasing z-value, which are why z-values larger than 2 are not recommended for metal cutting. Commercial sialon tools may also contain a-sialon' (same crystal structure as a-Si3N4)with the general formula Me, (SiAl)12(ON)16 where x52 and Me often is yttrium. The a-Sialon increases hardness' but at the same time fracture toughness decreases which may be due to the equiaxed morphology of the a-grains.
WEAR MECHANISMS A good understanding of operating tool wear mechanisms is essential to the further development of ceramic cutting tool materials. The wear mechanisms commonly observed with ceramic tools will be discussed within this chapter. Abrasive wear is very common for most tool materials and caused mainly by the rubbing action of hard particles in the workpiece material. Abrasive wear is normally experiencedas flank wear. Ceramic materials have a high hot hardness and not many particles found in workpiece materials have a higher hardness at high temperatures. Abrasive wear can
anyhow occur due to the lower temperature in the workpiece, compared to the flank face of the tool. Often abrasive wear is caused by the tool material itself, since pulled-out hard grains can cause abrasion on the tool face. When machining cast iron with alumina-Tic tools, the flank wear depends on the grain sue of the TiCgrain", since abraded Tic-grains are the major cause for flank wear.
observed when machining Inconel 7 18 with sialon tools, which probably leads to reduced wear rates. Silicon nitride normally shows high wear rates when used in turning of nodular cast iron due to diffisiodsolution wear. Wear rates are reduced considerably in milling'*, or interrupted c ~ t t i n goperations '~ depending on the formation of protective coatings. The access to oxygen fkom the air is of conclusive significance for the formation of these coating^'^.
Inconel 718
TiN
Fig 3. Abrasive wear When machining cast iron with silicon nitride tools it is generally observed that the wear resistance increases with decreasing amounts of the glassy phase ' I . Chemical reactions during cutting can lead to lowering of the glass transition temperature, above which the viscosity of the glass decreases significantly. This can lead to intergranular flow and pull out of silicon nitride ,-12,13 ,which then can be the source of abrasive wear. Abrasion, by hard particles in the workpiece '4 is, however, often claimed to be the dominant wear mechanism of silicon nitride tools. Solutioddifision wear is a tribochemical wear caused at first hand by the high temperatures generated during metal cutting. It means that the tool material constituents are dissolved in the workpiece metal, through the fiesh interfaces that are continuously produced. An estimation of the magnitude of this wear mechanism can be obtained using solubility data '' for tool constituents and the work material in question. Figure 4 shows such calculated wear rates for steel machining.
Fig 5 . Chemical reactions Formation of coatings due to adherence of plastically defonnable inclusions in the workpiece on the tool surface is sometimes observed. Alumina inclusions in steel accelerate wear on cemented carbide tools due to abrasion, but can have the opposite effect in alumina based tools, due to the higher cutting edge temperatures2'. When turning cast iron without cast skin with silicon nitride tools a layer rich of silica is always observed on the tool surface2'. Such layers will prevent a direct contact between the silicon nitride grains and the iron decreasing difhsiodsolution wear.
Flank face of Sialon after machining grey cast iron
Fig 6. Formation of coatings
Fig 4. Calculated solution wear rates It is clear from these data that silicon nitride and silicon carbide are not suitable for steel machining where as alumina is practically insoluble in steel. Chemical reactions leading to formation of a new compound in the interface is not uncommon with ceramic materials due to the very high temperatures generated in metal cutting ( 1600' C have been measured has been la). The formation of TiN and alumina
Adhesion or attrition wear is commonly observed as an excessive localised wear at the depth-of-cut line when machining Ni-based heat resistant alloys. The deformation-hardening tendency of the workpiece material together with a saw-toothed chip produces intermittent conditions of seizure of workpiece material with pull out and small fiwtures of tool material. Plastic deformation takes place as a result of the combined action of high stresses and temperatures on the cutting edge. High hot hardness is critical to avoid this mechanism and gross plastic deformation is rare with ceramic materials. However, discrete plastic deformation of the outermost layer of the tool is commonly observed
23
on alumina based tools, as this is the main wear mezhanism when machining steel with alusnina based ceimics”.
R8k.W
Of-
- -*
MchhhS
s1..I (8s2641) a450 d m i n
Fig 7.Plastic deformation Thermal cracking is frequently observed especially with alumina based materialsz3,but can occur also in sialon Sometimes these cracks are the major cause for tool failure.
Won l n c o d 718 4M)m h i n
-
At203 TK: steel ss 2541
46omlmin
Fig 8. Thermal cracking Using thermoelastic theories, various thermal shock resistance parameters have been developed2*.Under transient thermal shock conditions, the maximum allowable temperature difference a body can be subjected to is proportional to R (fracture initiation resistance). R’=a,(1-v)/aE Where 6,is tensile strength, v is Poisson ratio, a is thermal expansion coefficient and E is Young’s modulus. A relative ranking using this parameter is shown in figure 9.
Fig 9. Thermal shock parameters Silicon nitride and sialon show a very good thermal shock resistance primarily due to their low thermal
24
expansion coefficients. Alumina based materials generally have a poor thermal shock resistance due to a high thermal expansion coefficient and relatively low thermal conductiv&. Silicon carbide whisker additions reduce thermal expansion and increase strength, why thermal shock resistance is improved. It should be noted that care must be taken using this or other theoretical parameters for the purpose of tool material selection since the parameters were developed for monolithic materials. All of the tool materials listed in the table are composites and actual measurementsof the critical temperature difference can be quite different. An indentation thermal shock test? and with boiling water as coolant gave the results according to figure 9 (ATJ The fact that alumina-Sic, shows a much better performance than can be expected from calculations based on strength, thermal conductivity and thermal expansion coefficient has been attributed to the interaction between microcracks in the matrix and the SiC-whiskerswhich prevent the coalescence of the cracks into critical flaws4. In spite of the low value of both the theoretical and experimental ATc for alumina-zirconia it can withstand thermal cycling better than alumina-Tic because crack extensions are smaller, often not leading to catastrophic tool failure. This phenomenon is probably also related to microcracking during zirconia transformation. Fracture is a failure mode that should be avoided by using appropriate machining parameters, since it gives an unpredictable tool life and can also damage tool holder and work piece. Since ceramics are brittle materials they are sensitive to defects and flaws. Much of the development of ceramic materials has been aimed at controlling these defects by proper processing. The Weibull modulus “m” is one parameter; this describes the homogeneity of a material. Figure 10 shows the tool life of some commercial ceramic tool materials in an interrupted turning operation26,plotted in a Weibull diagram where the slope of the lines give the Weibull modulus m. It is shown that the two particle reinforced alumina tools show a very unpredictable tool life. If fracture is to be avoided tool life is determined by the worst result for the tool material in question and only a hction of the true potential can be utilised. In contrast the whisker reinforced alumina tool and the silicon nitride show a very predictable tool life and with Weibull modulus being similar to cemented carbides. These latter materials also contain defects but the whiskers or elongated grains have eliminated the negative effect of these probably by reducing stress concentrations.
with coolant to minimise part temperature. Sialons are used for rough machining at lower cutting speeds, but higher feeds than whisker ceramics, while the latter can be used for high speed finishing operations. The introduction of ceramic tools in this application meant a 10-fold increase in metal removal rate as a result of development of quite new materials.
su
T l b
li.d
Fig 10. Tool life in interrupted cutting
0
M.M!nmh*.
apa;almrw
1
3
%
-
r o o r o o * & & & U(nrlk*q
APPLICATIONS OF CERAMIC CUTTING TOOLS Cast iron machining is the most common application for ceramic cutting tools. The chips produced in grey cast iron are relatively short due to the graphite flakes and temperatures are lower then when machining steel. Silicon nitride is the preferred material in milling and rough turning operations due to their high strength and thermal shock resistance. They can also be used with coolant. Aluminazirconia and alumina-Tic have a higher wear resistance than silicon nitride and are preferably used in finish turning operations. They are normally used without coolant and interruptions must be avoided. Alumina-Tic is the preferred material if surface finish of the workpiece is the tool life determining criterion due to its high resistance to attrition wear, at the secondary cutting edge. Nodular cast iron is being increasingly used in the automotive industry, as a replacement for steel. Nodular iron is less successfully machined with ceramic tools. Cutting edge temperatures is much higher which puts high demands on chemical resistance and thermal shock resistance. Alumina and Tic/" coated silicon nitride tools can in some applications give a satisfactory tool life, but adhesion of the coating can be a problem due to large differences in thermal expansion coefficient. Hardpart turning has increased significantly during recent years. Hard steel or iron parts generally have hardness between 48-64 HRC and were previously machined by grinding. The operation generates high temperatures and cutting forces. This is why aluminaTic is the preferred tool material. Occasionally also SIC whisker reinforced alumina and silicon nitride cutting tools are used in special applications. Heat resistant alloys were only 15 years ago machined with carbide tools at cutting speeds of 50 d m i n . Superalloys are extremely difficult to machine since they retain strength at high temperatures and generate very high cutting temperatures. Both Sialon and whisker reinforced ceramic are used to day, normally
Fig 11. Machining of Inconel 7 18. Steel machining is by far the most common chip forming operation. Long continuo chips lead to high temperatures and chemical reactions with the tool material. Ceramic coatings on cemented carbide are the most successful solution to these demands. Both alumina-zirconia and.alumina-Tic are used in high speed machining of low alloy and medium carbon steels at high cutting speeds, but only in continuos operations with a minimum of intermittence due to their susceptibilityto mechanical and thermal hctures. Unfortunately, ceramics with high resistance to these failure mechanisms are unsuited, because of excessive crater and flank wear as a result of solutioddifhsion.
Fig 12. Machining of steel with Sialon
DEVELOPMENT TRENDS Undoubtedly, the strive for cutting tool materials with higher speed capability will continue, at an escalating speed. High speed machining is an important area, only in the last five years the maximum spindle speed has increased f?om 10.000 to 80.000 rpm. Another factor which will favour the use of ceramic tools is the trend to eliminate cutting fluids, both for environmental and for economical reasons. The cost of cutting fluids is at present 3 times the cost of that for cutting tools.
25
The share of ceramic tools in different workpiece materials and development trends are shown in figure 13.
Fig 13. Ceramic share and development trends for indexable inserts At present ceramic cutting tools expand their application only in 2 areas, namely hard part turning and heat resistant alloy machining. They represent at the moment only about 5 YOof the total machining operations this is why the impact on the overall consumption of ceramic inserts is low. In cast iron machining the competition from newly developed cemented carbides, with thick ceramic coatings is strong, with the share of ceramic inserts being fairly constant, although the development of more wear resistant silicon nitrides has been extensive. In steel, only aluminazirconia and alumina-Tic have a potential for growth, but in such a case their toughness and thermal shock resistance have to be very much improved. It therefore seems that a substantial growth of ceramic inserts will not take place until completely new materials have been developed with, at the same time, high strength, toughness, thermal shock and wear resistance. None of the existing materials currently fulfil these requirements to the full extent.
REFERENCES N.Claussen, J. Am. Cer. SOC.59 (1976)49 M. Furukawa et al, Nippon Tungsten Review 18 (1985) 16 P.F.Becher and G.C.Wei, Am. Ceram. SOC.Bull. 64 (2) 298 T.N.Tiegs and P.F Becher, J. Am. Ceram. SOC. 70(1987) C-109 P.Sajgalik et al, J. Am. Ceram. SOC. 78(1995) 2619 A.J.Pyzik et al, Mat. Res. SOC.Symp. Proc. 287( 1993) 4 11 K.H.Jack, J. Mater. Sci. ll(1976) 1135 T.Ekstrom and M.Nygren, J.Am.Ceram.Soc.75 (1992) 259 T.Ekstrom, J.Hard.Mater.4 (1993) 77 10) Y.Katsumura, Tribology Transactions, 36(1993) 43 11) H.Tanaka, Trans. Mater. Res. SOC.Jpn. 14A(1994) 54 1
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12) H.K.Tonshoff et al, Mat. -wiss. u. Werkstofftech. 26(1995) 255 13) R.F.Silva et al, Wear, 148 (1991)68 14) P.J. Mehrotra, SME Technical Paper MR94-173. 15) B.M.Kramer and N.P.Suh, J. of Engineering for Industry 102(1980)303 16) J.F.Huet and J.F.Kramer, Proc. of the loth North American Metalworking Research Conference, ed. R.S.Hahn. ASME Dearborn, MI, (1983) 297 17) G.Brandt et al, J.Eur.Ceram. SOC.6 (1990) 273 18) B.Dedenka, Fortschr.-Ber. VDI Reihe 2 Nr 249. Diisseldorf: VDI-Verlag 1992. 19) J.Lauscher, Drehen mit SiliciumnitridSchneidkeramik-Verschleiss-vorgiinge und mechanismen-, Dissertation, TH Aachen, 8 dec 1988. 20) G.Brandt and M. M h s ,Wear, 118(1987)99 2 1) G.Brandt et al, S h e engineering 3( 1987)221 22) G.Brandt, Wear 112(1986)39 23) G.Brandt, Surface Engineering 2(1986)121 24) D.P.H.Hasselman, Ceramurgia International, 4( 1978)147 25) T.Andemson and D.J.Rowcliffe,J.Am.Ceram.Soc. 79(1996)1509 26) W.Konig and KGerschwiler, VDI-Z, 131(1989)52
PRACTICAL USE OF CERAMIC COMPONENTS AND CERAMIC ENGINES Hide0 Kawamura I s m Ceramics Research Institute Co., Ltd. Japan
Abstract In opder to improve the air pollution m the Tokyo metropolitan aty, DPF which was developed by Isuzu ceramics Research h h t e Co., Ltd. are tnstalled on vehicles using in Tokyo, Yokohama and Kawasaki Cities. This DPF is collsisted of laminated Ceramics(SiC)fiber for fihxationandmetalnetheatersfcrqenemtionby burning down mated PM. The system is very tough m the markets using light oil contained sulfur above SOOPPM, because Sic fiberhas hlgh s i l x q t h a n d ~ m t h e exhaust gas. To keeping nice enviromnd on the earth by reductngco2,ceramicsturbocompoundenginesaresilldied. The tfiamos shzture coI1sis&d of cgdmics(Si3N.,)parts is very e W v e to reduce heat rejection hm combdon chambers and the recovering turbine produce shaft work h m t h e exhaust gas enhancsdthe enblpymthe cadmics bcompolnndenghaes. However, as very big exhaust gas eneqg is retnaining still after the turbine,we studied to use the chemical system recovaingtheenerpyhmtheexhaustgastoshaftwork As the resuh of fimshmg the study, we will be realized the engine with the high hnnal efficiency of about 65% or above whichis the dream value for engine engineers.
Introduction In order to improve e n .
~11prOblemsOnthe
earth,we should make effints m pur&-mg the air which has been polluted by Nitrogen Oxidati~Ox),Particulate Matea(pM)andmUchCadmdeOxidation(C03e~ hm vehlcles as soon as possiile. Althougk It’s seem hat this prcblems has big scale and it can’t solve except using very big deal and political powers between ahnost Countries, sud.1 p o w d person as the Governor of T&yo meiropolitan aty Mr. Ishibara appealed to the dizm of Tokyo to h e diesel vehlcles out of the city last y a , ihenxpnthe ripple exteded all of Japan m an instant and public opinion will make the policy realize. The diesel
particulate fikx(DPF) which we have beenpursuingto
develop by using cedamics fibeas with the Bureau of lib- ProQctiOn m Tokyo Metropolitan Government has been dmm into the limelight and components made of ceramics is just about chngmg the -ne of big cities. On the other hand, Japan Government must give up to promote the development of nuclear genaating systems l-ecedyblxaLKe of the accident happened in an atomic file1 plant m Japan. Resuhmg m the decisio~CNG (compressed ~n a~t d *)has been mthe spotlightdue to a snail quantity ofexhaust pollution when it is used for llltanal combustion engine, as well as the Ceramics turbo ampound engine which we have developed with Japan MlTI d Japan &as aSSOciation has became a promising engine system in the fhue.
Air Pollution and DPF We can see m a y medical papers studied on the relation between lung diseases and died partiailate matter &om 4-5 years ago. b f m , DPF having honey comb shape made of cmdiright was mvdgated on the pehmmmanddudnhtytouseitpmhcally.A s t h e d t s , we could know that thistype ofDPFwere broken Sometimes by thermal stress when PM is M o u t onthe &rafter PM is W-ated and collected on the filter adce, as well as the same type of DPF made of Sic was c o n c t u c t e d ~ w ~.~tothetestrasult,mspiteofthehi~~ materials, Sic DPF also were brokenbythethend siress at the #on times because of the micro size ofholes jnducedmthewallsto~PMarereduclngthe~ AmmhgIy, we started to develq the DPF laminated Sic tibe€S.”
Fig 1 shows the outer view of DPF developed by us. The DPF is COrlSiSted of lmmatmg seats like blanket sandwiched bymetal~~andthemerseatis~Ceramics tibe€S at I.andomand dmse seats are bent like bellows. The filterislaminatedaseatwithsnalldiametertibe€Sandaseat
27
M Carbon Graphite
20
0 a43em
065-11
043-06
11-21
21-7
7rm<
Length of c.rbm Graphite ( II m)
Fig3 Filtrating Efficiency of Each Graphite Size Investigated by Tokyo Governor in the use of Ceramics Fiber DPF
- - -
100
-
250;
Fig.1 Photo of Filter for DPF
Fig 2 shows the photoscaph of ('xoss sectional view when PM was fihmted in the filter pded up &a fibas at random, We can see PM like black gabs which are caught on the ceramics Gbers and ceramics fibas are shown as Whaeliraesmthephotograpk A s t h e n d t o. f offihmtmgdonbytheuse ~ ofhigh speed camera, snail size of PM are caught on the Ceramics iibm and the PM grain inaease the Size by gahermg much particle of PM. b f m , the DPF can &ate the smallsize PM bavmgthel@ ofbelow0.4 p m Dense Mesh Filter
Law Mesh Filter
a
Cleaned Exhaust Gas
Exhaust Gas
h .
h .
Filtrating PM in the Depth of Filter Fig.2 Cross Sectional View of Fiber Filter after Filtrating Particulate
Fig 3 shows the i i h t m g efficiency mvdgated on every length of PM. I pmume that his highef6cienqof fihatmg is due to 7igZlg flowing of exhaust &as topass thKnlghthe
CetIlmicsfiberiilterandthechanCeofimpingingthePM grain to the fibas is incaeaSing. Fig 4 shows the fihmng efficiencywhenfiltmngtimepassed Accmhngtothe
28
50 0
5
10
15
20
25
30
35
40
45
F h a t i n g limeshin)
Fig4 FiltratingEfficiency of PM in the use of Ceramics Fiber DPF
Many kdofdumbihy tests were condllcted First, I carriedout to select the fiber mordeato see throughthe Sic. Fig 5 shows thephoQqph ofmicro slrudm afkr exposing test m the air of 1OoO"C fix 5WB. Sic fiber has excellent dwatnhyathightemperahneairconditionIIherefore,we
cxmtrolledthetLmpameof fibmas keeprngbelow lOoo0C. As the lhiclaless O f l h d x n l deposit m the iilter difkron the positions of the iilter, teqemme clislrjMon become very difkent because the combustian of the c l e p l t exblded witfiout~lifdleheate€ofthewirenetdoesnotmakethe depcmtigniteatthesamethaes. Fig6 shows a schemaofthe filtexshuctwe. Thewire net is very efFed to control the &npemtm of combustrng PM depwt clue to ignition it on all spots. As the DPF coukl guambxithe chaabilhy fbr long thnes,the colltrol system of DPF was decided time schedule shown m Fig 7. Two ~werepleparedasusingthe*bytumsaudaoesi& ofiilterisused fix fihdionwhen another *is used fix regenedm
d t h e DPF as well as almostbigcitiesare also preparing the same lllllolcIpal ordinance. bwe must produce the DPF systems ofabout400 thousand sets m a year h2001.
Ceramics Turbo Compound Engines
YXlHr at lOOOt:
Apowdmethodofre
Having Durability and Reliability
t
Mechmcal Loss
Mechanical Loss
Fig& Schema of DPF Structure Laminstad SE Fiber Mat Like Blanket and Metal Nets
Heat Reji
(23.0%) Ceramics LH.R Endne
Ftlvrl Flltcr2 b
--
Fig8
Comparison of Heat Balance between Conventional Cooled Engine and LH.R Engine
Fig7 Schema of DPF System
We supplied the DPF systems of about 150 sets to big cities m Japan included Tokyo, Yokohama, Kawasaki and othas. Tokyo metropolitan city has jhished to install the systems on the many kind of &cles and they have already c20dinddle dlddltyandthe advantage ofthe system on cicybuses becausetfiebuseswere usedabout 100 thwsand k m and they were no e x p e r b ~on troubles. Tokyo metropolhan city intend to installthe syslans on the vehicles ofabout 1500mthisyearandtheyarep~par&jamutllcIpal opdinance to putdiesel truck usas&anobli&ationto
29
passage used tlae beat insulation WLIqmIEm made of &a lnateds. The aleqgy recovaing t t n b k is laested of turbo c.&ager and the rotor fix genemtm as shown m Fig 10 which is namedbdo chaqjerandpemb~. Whenthe turbine efiiciemy was dievedby 80% and the caqmsor efficiency was improved about 79!%we can achieved the e5ciency of 44% m the d turbo c o m p o d engiue used C.N.G. However, as we could not succeeded m dle ldizallm of the talge which was atlout 5004 how to US steam power which was produced by ranainedsrhaustgas enaKywas .studied Insteadofusing the fidl scale of Ranloine cycle after the bdo cbarger and genemmr (T.C.G) water is sent to a beat exchangerindled atta T.C.Gand boiled steam is imwhxi totheeidlaxxof the T.C.G.As the gas volume is incnrlsedfordriving hntrine, the shaft power dthe e&iency of T.C.G imrease.As the result, the efficiency of the turbo campcxmd engiee was improvedtoby about46-47%. On the other hand, the diesel combustion of C.N.G in an &, which enhance temperature and pressure in cylmder due to the thermos structure had been great troubles. As CH, has small number of moleculaq the time of resolving the molecular and reactingwith oxygen are very short and its very
I
Fig 11 The Combusbon Process of CNG Diesel Engine which CNG Wm, Low Pressure by 0.3 to 0.5Mpa is Supplied
18 16 14
r"
u
e
12 10
1
In v)
E a U
m
I
I
I
I
I
I
0
2
-60 -30 0 Crank Angle [ded Cylinder Pressure Rise (Sharp Combustion)
-90 Exhaust Valve
In Original Engine Steam Tubnc
18
~~~~~
Fig.9 Schema of a Ceramics Turbo Compound Engine for CNG
-a
-
16 14
I
-P
12 10
%
T 2
8
800
6
600
4
400
2
200
: B
0
0
:
0
UI
? n
Y
0
0
Y
-90
-60
-30
0
Crank Angle [degl Cylinder Pressure Rise (Smoothly Combustion)
\
.
Cedtnic Turbine Blade Reinforcing Tube Made of Si3N4 . Fig10 The Structure of The Energy Recovering
Turbo Charger and Generator
30
In New Engine
Fig 12
Conparison of Cylinder Pressure Rise between Original and New Engine
30
easy to 0cc1.1~knoclnng in the hgh temperature and pressure air condition Therefore, we have developed new combustion chamber system which is consisted by pre-combustion chamber with control valve, the lugh rate of exhaust gas reCirculation (E.G.R) and learn air fuel mixture system as the gas fuel of 80% is introduced to the cylinder for homogeneous mixture and c o m p r e d C.N.G of about 20%is injected to the pre- chamber Fig 11shows the schema of combustionprogress in the C.N.G engine. The combustion has been improved very smmthly by the use of the new combustionchamber systems. E'lg 12 shows comparison of the results of pressure rise in cylinder, which is investgated in the new combustion chamber and the oxi& type of
Then, we stated to develop w c convertas for -er CH, and s e p l a t m to extract cq fiom tJle exhaust gas. Fig14 shows the shu~hae of vessel for r e f a CH+ which is COIlSiStBd of heat ex&mger and refhmercoated Ru,PC Ni Rhand CsO, .gains on the fins. Fig.15 shows the result of that the efficiency of refCH, lIlBeased with the enhancmg the tenpmwe of Cahlytlc converter and the e f E c i e n c y ~ b y a b o u 75% t at 750°C.
nn
High Temperature Exhaust Gas Inlet
H4
co
I
4.
Ht
couhmonsystem.
Ceramics Turbo Compound Engine Used Reformed C.N.G Fuel We had made areview on the cerinnics turbo compomd thermal efficienq Although T.C.G system was used for recovering the exhaust gas awgy m the ceramicsturbocompoundengille,theexhaustgasawgyhas remainedabout32%(Fig.8). Inordertou~etheranareed enagyinexhaustgas,dlechanicalrecoveringsystemwae studied. Fig13 shows the schema of the &cs turbo cmpounclengiae a d d e d r e f v CH, system.When we can refom and cq to 2 c o and 2H, m the atalytlc convatgathigh~d~thecaldcvalueof
Fig.14 The Structure o f Vessel for Reforming CH4
engine on the
Catalyst Co&d by
@?
Osm
sso
am
850
700
Tmmpntum of R.fomns 0..
CC)
Is0
100
Fig15 Reforming Rate of CH4 when Temperature or Reforming
~lsm~asshownmthefollowingequations.
Gas is Changed(Experimenta1 Date)
Ca+CI-&-+2CO+2H2 2CO+2H,+20z-+2CQ+2H,0+250,580KcalDkmoI CH,+2q+CQ+2HzO+191,290KcavZlanol
(1) (2)
(3)
I Thermal Efficiency of 65%and Coz Emission of 1/2willbe Achieved in the Engine ~ C a v r h r
Fig. 13 Ceramics Turbo Compound Engine with COz Separator from Exhaust Gas and Catalyst for Reforming CH, and C 0 2 to Hz and CO
Next we ~edto prochacethe sepa~atarof C q . As we could not look fbr the suitable sepznator m the conventional system, many materials were mdgated for lm43lng with the separator and we f d out hit carbon gt.aphae haviug much micro pore deposaed ph-c had the excellent
performanceinFig.16 shows the results of hit separahg mte of C02 fi-om exhaust gas iuaeaxd withinmasing the additional mte of phosphoric acid As we have takenthe brightly views m the cuamics turbo coIllpollDd engine used reformed C.N.G &l, producing the enginesystansh started, w h o s e p j = t ~sponsor by shrp &Ocean Foundation. When we will have finished to study the engine system, we will be reali2edthe thermal efficiency of 65% or above which is a dream value for engine engineas. (Fig.17)
31
100
290
-0“
0
80
m c
70
a m
.s
recovering the ahaust gis eeergy by the use of T.C.G.
60
F1-q
1 50
E
II]
to reco~athe exhaust energy, the chernid
the eaimation tile thermal ef6ciw willbe imploved about 65% OT more.
recovering system has
40
30 0
5
10
15
20
25
30
35
be
studied According to
AddiQonal Rate of Phosphonc Acid (Z) Fig.16 Separating Rate o f COzwhen Carbon Graphite Deposited Phosphoric Acid is used for Separating Materials
ued Cell
G/T. S/T Combined Thormal Power
Nuclear Steam
Turbine
1kW
10
1MW
100
10
100
1000
Power Fig.17 Thermal Efficiency on Many Kind of Engines
on the other hand, we inteed ill use dle ceramics turbo compound mginepl-;acticayl fbrthe c q p e d o n m to supply very cheap electric power m the atyaniq because the engine systemhas such e x c e l l e n t k h as compt, simple, low NOx aaissioDsand high efficiencyas well as,we intehd to apply all of the techwlogies for the vehicle needing dean exhaust emissions a d high theamal efficiency which is sponsor by Japzln MlTI in order to achieve the target of the truck with NO.l level of fuel conslrmption aud low anissionSmthe world Conclusion The DPF made of Sic fiber will be used fbr the d i d vehicles used m Tokyo mtropohu city. We eshmted that the e o n volume is about 1300 thousand-3300 thousand kits per years. Thz movanent obhgatmg the installation of DPF is exteding hTokyo to all of Japan The prachcal use ofthe d c s engine is progressing for the o-n system to sqply the electric power to the office m the city area When we fbkhed to conform the
32
References 1) T. Sakagdu et aL ‘~velopmentof High Durabday Diesel Patticulate Filter by using Sic Fiber” SAE paper No. 1995Lo1-0463 2) D. W Dickey “The E€kt of Insulated Cornbushan S ~ o n ~ I n j e d i o n D i ~ ~ P ~ , Ermssians and Gmbustion” SAE paper No. 89029 H Kawannna and H Matsuoka “Low Heat Fkyction En+ w - h Thema Struchae” SAE papa No.950978 H Kawauuna, A Egashmo and S. Selayama “Combustion and (hmbushon Chamber For a Low HeatRejection&&” SAEpaperNo. W506 H Ka“Ceramics Engine far Eeagy Saving” 6th Sympmium of Cemmics Mateds and Components forIWmA~itaKqrNoteLectme. A Higashut0 H sasalo: a d H Ka“Compression Ignition Comtnmon m a l+dmhmd and Heat insulated bgme Using a Homogeneous N a n d Gas Mixture” SAE paperNo. 2oooO14330
BETA CERAMIC IN ZEBRA@A N D NAS BATTERIES Cord-H. Dustmann
CH-6821 Stabio, Switzerland
ABSTRACT Batteries as electro-chemical energy storage devices a~ characterised by the electrode material and the electrolyte. BETA batteries use V-Al203 ceramic fbr the electrolyte as the key component in combination with sodium (Na) for the negative electrode and NiCl2 (ZEBRA@battery) or sulphur (NAS battery) for the positive electrode. The V-cerarnic is conductive for sodium ions, an insulator for electrons and inert for the electrode material. Batteries of this type are in production in Switzerland (ZEBRA@battery) and in Japan (NAS battery). ZEBRA@batteries are designed for a power to energy ratio of 2 for electric vehicles or hybrid vehicles. The cells have an open circuit voltage of 2,58V and a capacity of 32Ah. NAS batteries are designed for load levelling with a power to energy ratio of 0,125, The cells have an open circuit voltage of 2,08V and a typical capacity d 632Ah. In Europe and Japan there is production experience d more than 500 000 electrolyte tubes which provide the confidence for the next phase of industrialisation with millions of p” tubes produced annually.
INTRODUCTION 600 million vehicles that are in operation worldwide obtain their energy from crude oil. The transition to regenerative energies for mobility requires the availability of storage devices for electricity and Hz. The infrastructure for electricity already exists and the batteries are available now as descibed in this paper. 1/3 of the primary energy consumed worldwide is converted to electricity (14 TWh annually) for instantanious use because there is very little storage capacity. Power plant capacity is necessary for peak power demand which leaves large investments in equipment unused in low load periods. NAS batteries are developed for load levelling in order to reduce this problem. Both fields represent an enormous demand for a low cost and effective device for the storage of electricity. Electrochemical solutions are an important option.
ELECTROCHEMICAL STORAGE
ENERGX
All electrochemical energy storage devices use anode and cathode materials that should be light and offer as much as possible free energy out of their reaction. This reaction is controlled by the electrolyte which separates the electrodes ( fig.1 ). Charger
Load
+
+
N1
2NaCl-2Na
NICC
Fig. 1. ZEBRA@Electrochemical reaction The requirements on the electrolyte material are the following: It has to be stable long term in contact with both electrode materials 0 It has to be a low resistance conductor for the related ions and a high resistance insulator for electrons It must not release hazardous material or gases even under worst case abuse, short circuit or crash conditions It must be available in large quantities for low cost It has to fit into the battery recycling process without a cost burden p”-A1203 ceramic does meet these requirements for sodium batteries which have a theoretical specific energy of 790 Whkg. This is very high compared to the corresponding values of the well known systems Pb/H2S04/Pb0 with 161 Wh/kg and Cd/KOH/NiOOH with 208 Whkg. The energy content of P-batteries is determined by the volume of the cathode whereas the power is determined by the surface of the electrolyte Two different types of these &batteries are in use: 1. ZEBRA@ batteries are designed for automotive applications in electric and certain types of hybrid electric vehicles that require a power to energy ratio of about 2 (1,2)
33
2. NAS batteries are designed for load levelling applications optimized for high energy content with a power to energy ratio of 0,125 (3,4)
ZEBRA@CELL DESIGN 9 8
mBRA@ BA'ITERY SYSTEM The ZEBRA@Battery system (fig2) is divided in the subgroups: p"-ceramic electrolyte, cells, battery box with thermal insulation and peripheral equipment.
Air Cooling
cdls &to 16 baltefyunitsvl parsld
Fig. 2. ZEBRA@BatterySystem
P"CERAMICELECTROLYTE
The cell case which is connected to the negative pole is made out of steel by deep drawing. It is laser welded to the outer nickel component of the TCB seal. The inner nickel component of the TCB seal is laser welded to the current collector which is the positive pole (fig. 4) in contact with the positive electrode, a mixture of Ni powder and salt (NaCI) impregnated with a molten salt (NaAlc1.1). This liquid electrolyte has several functions:
The p"ceramic electrolyte (fig3) is the key component of the battery. The battery perfiormance and the reliability is dependent on it. The fact that the ceramic electrolyte is a perfect insulator for electrons results in a 100% Ah efficiency for the battery. Since the first report on the extraordinary high conductivity for sodium p-alumina in 1967 (5) intensive R&D work has prepared maturity for commercialization (6).
Fig. 5 . Liquid electrolyt reaction Fig. 3. p"- alumina Electrolyt The processes used for production of these tubes are: Calcination of boehmite produces a mixture of y- 6and 6-alumiNAS. This calcined powder is mixed with sodium carbonate and lithium hydroxide to produce a coarse powder, which is then calcinated at -1200°C to produce p"-alumina granules. Upon dispersing this p"alumina in water and milling to a particle size of 1-2 Pm The slip is then spray-dryed to produce a fi-ee-flowing powder. This powder is isostatically pressed into the desired monolith tubular shape and sintered in a batch furnace inside a spinel crucible to prevent soda loss. Finally the tube is cut to length and sealed to an aalumina ring with a sealing glass. After final inspection this component is delivered to the cell assembly line.
34
1. Sodium ion conductivity between the inner surface of the p"-electrolyte and the NiC12 inside the cathode in order to utilize its total volume 2. Overcharge protection following the overcharge reaction in fig.5. This protects the ceramic due to the availability of additional sodium and stops the charge current by the higher open circuit voltage 3. Overdischarge protection following the overdischarge reaction in fig.5. 4. Short circuit of the cell in the case of a ceramic fiacture by the formation of metallic aluminium following the overdischarge reaction of fig.5. This makes the battery tolerant against cell failures up to about 5%. 5. Contribution to battery safety. In case of a heavy accident there must not be any additional danger from the battery. This was demonstrated in a test in which the battery was crashed against a pole with 50 km/h (fig.6). As soon as the ceramic breaks the liquid electrolyte reacts with the liquid sodium and the resulting NaCl and A1 passivates
the cathode. This reaction produces only 213 of thermal load compared to the normal discharge reaction (7). A complete review on battery safety is given in (8).
The cells are assembled in a double walled vacuum insulated box. The volume between the inner and the outer wall is filled with an insulation material with the very low heat conductivity of only 0,006 W/mK. This material also cames the atmospheric pressure so that the rectangular shape of the box can be made using 0,5 mm stainless steel sheet metal. An ohmic heater and an air cooling fan are operated by the battery controller BMI in order to keep the internal operating temperature in the limits of 270°C and 350°C. The outside surface of the box is only 5 - 10°C above ambient. The main contactors and the charger are integral parts of the battery system.
BA'ITERY LIFE AND RECYCLING
Fig. 6. Battery safety 212 crash test
=BRA'
BATTERY DESIGN
ZEBRA@cellscan be connected in parallel and in series. The series connected cells get never out af balance because of the 100% Ah efficiency and the parallel connected cells balance themself automatically by redistribution of the current due to the internal resistance of the cells. The smallest battery has about 105 cells and the largest one has 528 cells at present. Fig.7 shows the standard battery Z5C with 216 cells.
25-278
Type
ML-32 Capacity 32 Rated Energy 17.8 Open Circuit Voltage 0-15% DOD V 278.6 557 Max. discharge current A 224 112 ML3 I 2 1 6 Cell Type/N" of cells 195 Weight with BMi kg Specific energy without BMI Wh/kg 94 Energy density without BMI Wh/l 148 Specific power W/kg 169 Power density W/I 265 Peak power kW 32 80% DOD. 213 OCV. 30s, 335%
"C W
-40 to +70 < 120
at 270'C internal temperature
Fig. 7 Z5 standard battery with main data
APPLICATIONS This battery system is used in the DaimlerChrysler VITO as a van for citylogistic and postal services as well as in different bus types. Fig. 8 shows the example of a small city bus which is available in an electric, a hybid and a diesel version.
25-557
ML-64 Ah 64 kWh 17.8
Ambient temperature Thermal loss
The average lifetime of a battery is 1000 cycles OF 100% charge and discharge and 10 years. 9 years and 1500 cycles have been demonstrated. Ageing effects known are only a rise of the internal resistance due to slow cathode morphology change and ceramic cracks due to small non detected production faults. Up to 5% cell failures are tolerated. For recycling the cells are sold to steel production plants where the contained Ni and steel is utilized in the stainless steel production process in which the salt and ceramic content contribute to the slag.
Hybrid
Electric
Capacity
10+1 Seating 29 standing
10+1 Seating 29 standing
Drive
Brushless DC
Brushless DC
50kW
50kW
Type of Battery
2xz5c
5xz5c
Energy Content
35.6 kWh
89 kWh
Battery Voltage (OW
279 V
279 V
Weight of Battery
400 kg
1000 kg
Maximum Speed E
50 kmlh
60 kmlh
Drive power
Fig. 8. Autodromo bus 35
The electric bus offers zero emission transportation for up to 150 km per day with overnight charge and up to 350 km with interim charging during the day.
table I).Due to the specially designed safety tube the amount of sodium available for the reaction with sulphur is limited in the event of a break in the p”alumina tube. The battery safety was demonstrated for stationary use.
NAS BATTERY SYSTEM The other battery system that utilizes b”-mamic for the electrolyte is the NAS system developed by NGK Insulators in Japan in cooperation with Tokyo Electric Power Company (4). The basic reaction is 2 Na + 3 S e Na& The main differences to the ZEBRA@systemare that the cathode is sulphur instead of NiCh and the absence of the liquid electrolyte. The subgroups of the NAS battery are the ci..p..”.electrolyte, the cells the battery box and the ACDC converter with the transformer for the integration into the 6,6 kW, 3 0 utility network for load levelling.
NAS CELL DESIGN The NAS cells (fig. 9) are designed for very high energy content in stationary applications.
Fig. 10. NAS
PI-
alumina Electrolyt
b*- Alumina-TubeDimension Type Outer Diameter (mm)
ZEBRA
NAS
ML3
T5
+ 33
058
Length (mm)
220
477
Thickness (mm)
1.25
1.7
ACC. Production 1998 130.000
70.000
Table 1
NAS BATTERY APPLICATIONS 384 NAS cells are connected in parallel and in series for a 50 kW, 400 kWh module (fig. 11).
Fig. 9.Nas Cell
Fig. 11 50 kW modul NAS Battery
The cell case which is the positive pole is made out d deep drawn aluminium and inside the electrolyte is the sodium electrode which is the negative pole. The ceramic electrolyte has impressive dimensions (fig. 10, 36
For a 6 MW/48 MWh load levelling power plant 120 of these modules are assembled in 6 blocks. All together 20 plants with 15,7 MW/125,6 MWh total are commissioned until March 2000. This system is designed for an 8 h charge over night, 4 h interval, 8 h discharge during daytime peak power and 4 h interval
before the next charge starts. For this mode of operation cooling is not necessary.
CONCLUSION Batteries using p"-alumina for the electrolyte are under development since 30 years for NAS and 20 Years for ZEBRA (NaNiClz). Both systems have now reached maturity for industrialization and first production experience has been gained. Within the next years they have to prove their competitiveness to other options for electric energy storage.
REPERENCES (1) J. Coetzer: ,,A New High-Density Battery System" J. Power Sources 18:377 - 380 (1986) (2) J. Coetzer, J.L. Sudworth: ,,A Second Generation Sodium/Nickel Chloride ZEBRA@Cell" Electric Vehicle Symposium 13, Osaka, Oct. 1996 (3) J.L. Sudworth, A.R. Tilley: ,,The Sodium/Sulfur Battery" London, Chapman and Hall (1 985) (4) Communication from NGK Insulators, Ltd., Nagoya, Japan (5) Y.F.J. Yao, J.T. Kummer: J. Inorg. Nucl. Chem. 29 (1967) p. 2453 (6) J.L. Sudworth, P. Barrow, W. Dong, B. Dunn, G.C. Farrington, J.O. Thomas: ,,Toward Commercialisationof the Beta-Alumina Family d Ionic Conductors" MRS Bulletin March 2000 (7) A. van Zyl, C.-H. Dustmann: ,,Safety Aspects of ZEBRA High Energy Batteries" Proceedings d the WEVA Conference for Electric Vehicle Research, Development and Operation, Nov. I995 in Paris (8) D. Trickett: ,,Current Status of Health and Safety Issues of SodiudMetal Chloride (ZEBRA) Batteries, NREL/TP-46025553, November I998 (9) E. Kodema, A. Okuno, F. Kiuchi, Y. Kurashima, T. Mima, K. Mori:"Development of Compact Sodium Sulfur Battery", Proc. of Electrical Energy Storage Applications and Technologies, Chester, June 16 - 18, 1998.
37
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OXYGEN SENSORS FOR LEAN COMBUSTION ENGINES E. Ivers-Tiffbe', K. H. Hardtl', W. MenesMou' and J. Riegel** (*) Universitat Karlsruhe(TH), Karlsruhe, Germany (**) Robert Bosch GmbH, Stuttgart
ABSTRACT Oxygen sensors are used in automotive applications to control the air-fuel ratio in order to reduce emissions and fuel consumption. The three way catalyst system (TWC) represents the most effective system for the emission control at this time. In TWC systems, the airfuel ratio is kept at the stoichiometric point lambda=] and is controlled by potentiometric zirconia sensors (Nernst type sensors). New control strategies with linear lambda control at lambda=l, for direct injection engines and other lean bum engines operating with air excess (lambda>1) need alternative sensor concepts. Hence, current limiting electrochemical pumping cells (amperometric sensors) based on zirconia have been developed for engine control applications. The paper will give an overview about potentiometric and amperometric sensors based on zirconia and will present new research works in resistive type oxygen sensors based on semiconducting metal oxides as a future option.
INTRODUCTION The lambda (1)closed loop control together with the lambda sensor and the three-way catalyst (TWC) represent today's most effective concept for the reduction of toxic emissions in spark ignition (SI) engines. Since 1976, when Bosch started with the world's first ZrO, based oxygen sensor to go into operation in vehicle exhaust emission control systems, worldwide a few hundred millions of lambda sensors have been produced. To meet the new exhaust emission requirements LEV (low emission vehicle), EU4 and ULEV (ultra low emission vehicle) and fkture regulations like SULEV (super ultra low emission vehicle) the two-stage controller using a conventional thimble type ZrO, oxygen sensor seems to be at its limit. Today the main emissions appear during warm up phase and, especially with aged catalysts, at maximum deviation from h = 1 due to the oscillation of the conventional two-stage controller behavior. Therefore new emission control strategies have been developed and were applied in vehicles during the last few years [ 1,2,3,4]. One strategy keeps the two-stage controller but starts the closed-loop mode during warm-up phase. The engine is preferentially slightly lean driven or uses a secondary air pump to enable fast light off of the catalyst. This requires fast light off capability of the upstream sensor (115 s resp. 4 0 s) which needs to be
much better than the capability of the conventional thimble type oxygen sensor. Another strategy uses the linear h-control to keep the conversion rate of aged catalysts high by reducing the deviation from the ideal h = 1 point. Together with lean warm-up and secondary air pump concepts these strategies require a linear wide range lambda sensor with fast light off capability and high accuracy. Offering less fuel consumption lean bum engines and recently gasoline direct injection (GDI) engines were developed which need linear lambda control outside h=l. Current lean burn and GDI systems don't operate exclusively at lean condition but also at h=l for better acceleration and emission conversation due to threeway-catalysts. If NOx-storage catalysts are used under lean condition, cyclic rich phases have to guarantee the NOx-regeneration. Wide range sensors can improve the quality by closed loop control during this phases. While simple limiting current oxygen sensors only can operate in lean exhaust gas wide range lambda sensors have a measuring range from h=0.7 to air. The functional principles of such universal oxygen sensors are already described [5, 6, 7, 8, 91 and Bosch started with series production of its wide range lambda sensor LSU in 1998. The ZrO, planar technology is the basis for the universal wide range oxygen sensor with high performance for lean bum applications further miniaturization and a platform for the next generation of exhaust gas sensors for NOx, HC, etc. [8,9]. Resistive type oxygen sensors based on semiconducting metal oxides are an alternative to the above-mentioned sensors because they have also a good oxygen sensitivity and can be manufactured using a low-cost screen printing technique. A resistive gas sensor for internal combustion engines was proposed by Logothetis [lo] for the first time. He suggested semiconducting TiO,. However its decisive disadvantage is its insufficient chemical stability. A further disadvantage of TiO, as well as of many other semiconducting oxides is its high temperature dependence of the conductivity [l I]. For this reason, a family of semiconducting oxides has been investigated where the temperature dependence is low, for example cO,-m,o ~ 1 s~,MP,J~,.,o, [ 131, ~ a ~ e , . , ~ ~ . ~ 0 , [14]. In this paper we present a thick film sensor based on Sr(Fe,Ti)O, whose temperature dependence is suppressed by an adequate adding of iron [15]. Such sensors give new aspects for applications of resistive type sensors in lean exhaust gas.
39
PLANAR POTENTIOMETRIC OXYGEN SENSORS The planar technology of ZrO, oxygen sensors is quite similar to the multi-layer-technology of electronic circuits (MCM - multi-chip-modules). All operating layers are arranged in plane, consecutive surfaces. This technique allows the integration of complex functions for example heating and sensing layers, gaps, cavities and channels with determined porosity in a compact 3dimensional monolithic design.
-
the conventional heated ZrO, oxygen sensor in planar technology [9, 171. Fig. 2 illustrates a simplified version of the layer structure.
Porousprotective layer
L Outer e w d e Sensor sheet Inner electrode
a -
THICKFILM TECHNOLOGY OF PLANAR ZIRCONIA OXYGEN SENSORS
,.
"
'1 Airdudsheet Insulation layer
In a tape casting process [ 161 ZrOzceramic green sheets are manufactured by adding an organic binder phase to the Zirconia powder (fig.1). These green tapes are punched into separate sheets. On each side of the sheets in a screen-printing process (thick-film) individual function groups (e.g. Pt heater, Alz03 insulation layers, leads, electrodes, electrochemical cells and porous ceramic layers) are attached to arrange the desired layer structure. These layers can consist of various inorganic materials. By punching and drilling channels and holes (e.g. reference air duct, contact holes) can be produced. The layout for several individual sensing elements in a parallel arrangement is printed on one substrate sheet. A number of sheets can be stacked and laminated together to form complex composite structures. After separation the sensor elements are sintered.
2%
' . Heater .
.
9
.
- \ Heatersheet .
Fig. 2. Thick film technology of planar zirconia Nernst sensor LSF (Bosch) [9, 171 Fig. 3 shows a cross cut of the sensing cell [9, 171. Compared to the conventional thimble type zirconia oxygen sensor the planar sensor LSF operates as a ZrO, solid electrolyte galvanic oxygen-concentration cell with the advantages - fast light off capability (-10 s) and reduced electrical power consumption (5 - 7 W) on account of an integrated heater and lower thermal mass - small element size (-59x4~1 mm') and reduced weight - basic technique for new functions, complex planar designs (e.g. universal, HC and NOx-sensors) and M e r miniaturization. ZrO~-Ceramic Electrodes EEtl Porous protective layer Insulation
Sinteri
Fig. 1. Process scheme for manufacturing of planar zirconia oxygen sensors [171. Mechanical and thermal shock resistance, long term phase stability and high ionic conductivity are demanded to ensure adequate thermal and mechanical stability of the ceramic elements and to fulfill the requirements for automotive application over a life time of 15 years and 150000 miles. The properties of the different materials e.g. tape casting performance, thermal coefficient of expansion and cofiring capability of ZrO,, A1,0, and Pt layers must be accurately adapted ~71.
PLANAR LAMBDA=l SENSOR LSF The planar lambda=l sensor LSF4 is the realization of
0.98 1.0 1.02A NormalizedAIF ratio
Fig. 3. Cross cut of planar lambda=l sensor LSF (Bosch) Because of the temperature dependence in the rich region and the very low voltages and flat curve at the lean region, this galvanic lambda=l sensor only can be used near the stoichiometric point with high accuracy.
continuous h=l-control to lean bum control, diesel control, CNG engines, burners and measuring devices (Lambdameter LA3). Due to its fast response time cylinder balancing has been demonstrated to be feasible. This wide range oxygen sensor LSU4 is in series production for linear lambda control and lean burn control since 1998.
understood on a defect chemical basis and published in numerous papers e.g. [24].
0
-I
-2x=0935
RESISTIVE OXYGEN SENSORS The principle of operation of resistive sensors is based on the dependence of the electrical conductivity of a metal oxide on the oxygen partial pressure in the ambient gas atmosphere at elevated temperatures. Electrical conductivity (J of semiconductingoxides can be expressed generally by
- 5 ~ " ' I ' " ' I ' ~ ' " ' " ' I " ' -10 -5 0 5 lO!aPO, 1 Pa)
~~~
Fig. 8. Conductivity versus oxygen partial pressure and temperature for Sr(Ti,-,Fe,)O, ceramics. The first term, the activation energy for conduction E,, the Boltzmann constant k, and the temperature T, describes the temperature dependence of the electrical conductivity. Oxygen partial pressure dependence is described by the second term, including m as a constant which depends on the dominant type of bulk defects and corresponds to the sensor sensitivity. Figure 7 arranges some semiconducting oxides according to temperature dependence (EA) and sensitivity (m). 24
0 112
1I6
118
Fig. 7. Sensitivity (m) and temperature dependence (Ed of some semiconductingoxygen sensor materials. The ternary metal oxide strontium titanate is a promising material for resistive high temperature (T > 700 "C) oxygen sensors. SrTiO, is able to tolerate large levels of dopants without phase transformations, and the perovskite crystal structure is stable in a large temperature (T < 1200 "C) range [22, 231. This makes an adjustment of the electronic properties by doping the parent structure oxide with alternative metal ions possible and renders some interesting combinations of sensor properties. Figure 8 shows the electric conductivity of SrTi,-,Fe,O, bulk ceramics for x = 0.01 and x = 0.35 as a function of the oxygen partial pressure PO, and temperature T. The electric behavior at high temperatures is well
42
The PO,-dependence of (J can be divided into 3 ranges of the PO,, see figure 8. In the case of very small PO, (< lo4 Pa) the material is a n-type conductor. The conductivity decreases with increasing PO, (m < 0) and shows a strong temperature dependence for all iron contents x (EA > 1 eV). A conductivity minimum at mean values of PO, Pa - 1O-I pa), that almost shows partial pressure independence, indicates traces of ionic conductivity for x = 0.01. For an application of SrTi,,Fe,O, as a temperature independent oxygen sensor for lean bum engines, the partial pressure range PO, > 10' Pa (m > 0) is of interest. In this p-type range the isothermal data lines are very much closer together than in the n-type range. This can be explained. In metal oxides at a fixed oxygen partial pressure of the ambient atmosphere oxygen leaves the lattice with increasing temperature. This leads to an increase of electrons (n) or due to the generation-recombination (characterized by the band gap E,) to a decreasing hole concentration @J).The latter effect is opposite to the thermal promotion of the dominating charge carriers (p?), which leads to an increasing conductivity. Equation 3 describes the temperature dependence of conductivity (characterized by E,), that is determined in the p-type range by the reduction enthalpy of the oxygen vacancies AHRd and the energy of generation-recombination of electronic charge carriers E, [25].
L
Since the effective band gap E, of SrTi,~,Fe,O,depends strongly on the iron content x the temperature dependence of conductivity can be adjusted by the iron content x [26, 271. In the case of x=0.35, the temperature dependence of conductivity can even be eliminated (E, = 0) consequently, this iron content is the material base for an ideal resistive oxygen sensor for a use in a lean-bum engine @0,=103 Pa - 2.104 Pa). The
single-cell limiting current oxygen sensor is essentially used in exclusively lean exhaust gas applications [8].
PLANAR LIMITING CURRENT OXYGEN SENSORS For lean burn applications (GDI, diesel engines, CNG engines, gas burner) the Nernst measuring principle with insufficient lambda-range is replaced by an amperometric principle [ 18, 191. Principle of Function
MftusiDn Wr
I---it---siF--------1
cathode &+4e--..2P
Arode 2P-Q+48-
A single cell type universal oxygen sensor with a distinctive measuring range in the rich region can be realized by applying the anode to the reference air. Such a sensor is in series production from a Japanese manufacturer as a conventional thimble type universal oxygen sensor [21]. But the measuring range in rich exhaust gas is limited due to the limited gas supply from the reference air side. Further disadvantages of this sensor type are high power consumption and limitation of the light off time due to the thimble type construction.
DUAL CELL PLANAR UNIVERSAL WIDE RANGE OXYGEN SENSOR LSU
I L =axrpI.Dc&-co,
04:Cfdfusbn me(iiciant d 0 2 cq:"""bno14
0 Ic-----X----.CI
Fig. 4. Principle of limiting current oxygen sensor At elevated temperatures (>600 "C) an active transport (pumping) of oxygen ions can take place through the solid electrolyte ZrO, ceramic. By applying of an external pumping voltage (Up) on the electrodes oxygen ions (02?are pumped from the cathode to the anode (fig.4). The electrical current corresponds to the oxygen ion current and therefore to the current of oxygen molecules diffusing to the cathode in the exhaust gas. With increasing pumping voltage (Up) the current increases according to the internal cell resistance. A diffusion barrier in front of the cathode impedes the flow of oxygen molecules to the electrode, the oxygen concentration at the cathode decreases till zero and results in a current saturation (IL) beyond a certain pumping voltage threshold (fig.5). The resulting limiting current is roughly proportional to the exhaust gas oxygen concentration (coz):
I, =4FD--c,, Q L
(1)
(F = Faraday's constant, D = diffusion coefficient, Q = effective diffusion cross section, L = effective diffusion length).
-
3.0 20
8 i
a"
1.0
0 -1 0
-2.0
071.0 13 1.6 1.9 22 5 NotmalizedMF ratio
Fig. 6. Planar wide range oxygen sensor LSU4 (Bosch) Depending on the polarity of the applied voltage on the inner and outer pumping electrode, oxygen can be pumped out or into the gap. The sensing cell, consisting of a sensing electrode inside the gap and a reference electrode in a reference air duct, is a Nemst type solid electrolyte cell measuring the lambda value in the gap. By regulating the pumping voltage by an electronic closed loop control circuit a constant lambda value of the gas in the gap is maintained, corresponding to a Nernst voltage of the sensing cell of Umf 450 mV. In lean exhaust gas 0, is pumped out of, in rich exhaust gas, 0, is pumped into the gap which leads to a positive or negative pumping current depending on the Lambda value of the exhaust gas. At h=l no oxygen needs to be pumped in any direction and therefore no pumping current is measured. As for the limiting current sensor the pumping current is direct proportional to the oxygen concentration in the lean region (or oxygen requirement in the rich region). The signal characteristic is continuous over the whole lambda range. The sensor needs an operating temperature of -700...800 "C. A special electronic control circuit which is integrated in an ASIC is necessary for operation. Because of the integrated heater this sensor offers fast light off behavior ( 4 5 s). The range of application varies from
-
mA 2
5
E
.p
For the application in the whole lambda range from lean to rich a universal dual cell wide range oxygen sensor is preferred [8, 9, 171. This sensor consists of two cells, a pumping and a sensing cell separated from each other by a porous layer with a gap width of 20 - 50 pm which acts as a diffusion barrier (fig. 6).
1
a" a0 0
0.5
1
1.5
pumpins wage
V
0
8.36
20.9
%
& -
Fig. 5. Characteristic of limiting current oxygen sensor Depending on the porosity of the diffusion barrier the difhsion current is a combination of gas phase and Knudsen diffusion with different dependences on temperature and exhaust gas pressure [20]. This simple
41
slope m of SrTio,6sFeo,3s03 is 1/5 in this oxygen partial pressure range. SrTi,-,Fe,O, is chemically stable in the temperature and partial pressure range which is represented in figure 8. The ceramics (d = 500 pm) respond to a sudden oxygen partial pressure change within a time constant of a few seconds, which is in the same range as slightly acceptor doped SrTiO, [28], where the kinetic behavior is also determined by the diffusion of oxygen vacancies.
T / "C
, o ~11po ~
In the following, it will be shown that the temperature independence and the slope m of the characteristic curve of the ceramics (x = 0.35) can be transferred to thick films of the same material, as shown in figure 9.
1oo]
lop0
*---
*
~
4
porous thick film
thin film
-
yo
970
-
SrTiO, CeO,
0.70
SrFe0.35Ti0.6503
1 .oo
0.90
0.80
1000
tvr
Fig. 10. Response times of a Sr(Tio,6sFeo,3s)03 thick film (d = 15 pm, grain size 0.5 pm) compared with fast thin film sensors (d = 1 pm).
Fig. 9. SrTio,6,Feo,3,03 thick film sensor on a Al,O, substrate with Pt bottom side contacts. 1 square = 1 mm2. For the preparation of SrTio,,,Feo,,,O, thick films a screen printing paste prepared of the same powders as above was printed onto an alumina substrate (96 %). The thick films were dried and fired at 1050 "C. The resulting layer thickness was 15 pm, with grain sizes of 0.5 pm and an open porosity of 30 %. All examined thick films were equipped with Pt bottom side contacts, that were bumed in at 1300 "C. Sensors manufactured that way show very short response times due to their small layer thickness (1 5 pm). The response times of an oxygen thick film sensor (d = 15 pm) are shown in figure 10 as a function of temperature. The fast kinetic behavior (response times) of the thick films were investigated by a method developed by Tragut [29, 301. The total air gas pressure and the corresponding oxygen partial pressures p02 were changed periodically from 2.104 to 4.104 Pa. The modulation frequency f can be varied from f = 0.02 Hz to 2 kHz. The magnitude IAl and phase
Figure 1 1 shows a measurement result of the thick film sensor under lean-bum conditions in a car engine (Daimler-Benz, 2.0 1, 4 cylinder) with preceding fuel injection. The sensor was heated by a printed Pt substrate heater to its working temperature. It was protected only by a metallic housing with bores. It was mounted a few centimeters behind the outlet valve of the engine in the exhaust gas stream. Because of this installation close to the engine the sensor detects only one single cylinder. The aidfuel ratio was changed via the injection duration t, (4.2 to 5.1 ms) stepwise from h=1.4 to h=l.l. The sensor was operated at 5 V via a serial resistor of 3.9 kn. This leads to a sensor signal between 1 V and 2 V (sensor resistance: 1.2 kn at h = 1.4). Considering the delayed outflow process, the sensor can detect the very fast PO,-changes on account of the varied injection times ti. fuel
2 -
- injection cycle
1.4
1
.
2
-
7 I
.
h =1.4
1,42 h z 1,l
'
h =1,1
I
.I .n ! "
0
50
100
150
200
t I ms
Fig. 1 1 . Profile of the sensor signal after change of the fuel injection times ti of a car engine. The dashed line shows the injection process. On account of the delayed outflow process, the sensor signal is time-shifted. Exhaust gas temperature: 820 "C k 5 K; engine speed: 3500 r.p.m.
43
CONCLUSIONS ZrO, based oxygen sensors have been proved for a long time as reliable sensors for engine control systems. For lean burn engines universal wide range oxygen sensors in ZrO,planar technology are preferred in current series applications. Temperature independent resistive type oxygen sensors as presented might be a future alternative for distinctive applications in lean exhaust gases and offer high performance for fast lambda control.
ACKNOWLEDGMENTS We would like to thank the Institut f ir Kolbenmaschinen of the Universitiit Karlsruhe (TH) for the measurements on the car engine, and the Keramikverbund Karlsruhe Stuttgart (KKS) of the State Government of Baden-Wiirttemberg for the financial support.
REFERENCES [ 11 Bosch: Kraftfahrtechnisches Taschenbuch 23, Auflage, Vieweg-Verlag ( 1999). [2] K. Antonius, A. Gamer, S. Garrett, Honda first to
have gasoline engine verified at ULEV exhaust levels, Honda Press Release, Jan. (1995). [3] K. Winkler, M. Kusell, E. Schnaibel, W. Strehlau, U. Gobel, J. Hohne, W. Muller, The Development of an aftertreatment System for Gasoline Direct Injection Passenger Cars, F98T2 18, FISITA World Automotive Congress Proceedings, Sept. 27 Oct. 1 (1998) Paris. [4] M. Kiisell, W. Moser, M. Philipp, Motronic MED7 for Gasoline Direct Injection Engines: Engine Management System and Calibration Procedures, SAE 1999-01-1284, Detroit (1999). [5] S. Sojima and S. Mase, Multi-Layered Zirconia Oxygen Sensor for Lean Bum Engine Application, SAE 850378, Detroit (1985). [6] T. Takeuchi, Oxygen Sensors, Sensors and Actuators, 14, 109-124 (1988). [7] T. Yamada, N. Hayakawa, Y. Kami, T. Kawai, Universal Air-Fuel Ratio Heated Exhaust Gas Oxygen Sensor and Further Applications, SAE 920234, Detroit (1992). [8] H.-M. Wiedenmann, G. Hotzel, H. Neumann, J. Riegel, and H. Weyl, Exhaust Gas Sensors, in Automotive electronics Handbook, edited by Ronald K. Jurgen, McGraw Hill Inc. (1995). [9] J. Riegel, G. Hotzel, H. Neumann, Advanced Electrochemical Exhaust Gas Sensors for Automotive Application, Proc. 44" Meeting of the International Society of Electrochemistry, Berlin, Sept. 5-10 (1993). [ 101 E.M. Logothetis in Ceramic Engineering & Science, Proceedings 1, Westerville, Ohio, 28 1301 (1980). [ l 11 P.T. Moseley, Sensors and Actuators B 6, 149-156 (1992).
-
44
[ 12) K. Park, E.M. Logothetis, J. Electrochem. SOC.: Solid-state Science and Technology 124 (9), 1443-46 (1977). [13] C. Yu, Y. Shimizu, H. Arai, Chemistry Letters, 563-566 (1986). [ 141 P.T. Moseley, D.E. Williams, Polyhedron 8 (13/14), 1615-18 (1989). [15] D.E. Williams, B.C. Tofield, P. McGeehin, European Patent Specification EP0062994 (1982). [16] R. E. Mistler, D. J. Shanefield, R. B. Runk, Foil Casting of Ceramics, in Ceramic Processing before Firing, edited by G. Y. Onada and L. L. Hench, John Wiley and Sons, Inc., New York ,411-448 (1978). [ 171 H. Neumann, G. Hotzel, G. Lindemann, Advanced Planar Oxygen Sensors for Future Emission Control Strategies, SAE 970459, Detroit (1997). [ 181 H. Dietz, Gas-Diffusion-Controlled SolidElectrolyte Oxygen Sensors, Solid State Ionics 6, 175-183 (1982). [ 191 H.-M. Wiedenmann, Characteristics of Oxygen Sensors for Lean Exhaust Gas, VDI-Berichte 578, 129-151, VDI-Verlag, Dusseldorf (1985) [20] K. Saji, Characteristics of Limiting Current Type Oxygen Sensors, J. Electrochem. SOC.: Electrochemical Science and Technology, Vo1.134, No.10, Oct. (1987). [21] K.Mizusawa, K. Katoh, S. Hamaguchi, H. Hayashi, S. Hocho, Development of Air Fuel Ratio Sensor for 1997 Model1 Year LEV Vehicle, SAE 970843, Detroit (1997). [22] R. Moos, T. Bischoff, W. Menesklou, K. H. HZirdtl, Solubility of Lanthanum in Strontium Titanate in Oxygen-Rich Atmospheres, J. Mater. Sci. 32 4247-52 (1997). [23] S. Steinsvik, R. Bugge, J. Gjonnes, J. Tafto, T. Norby, J. Phys. Chem. Solids 58 (6), 969-76 (1997). [24] R. Moos, K.H. HZirdtl, J. Am. Ceram. SOC.80 (lo), 2549-62 (1997). [25] W. Menesklou, H.-J. Schreiner, K.H. Hiirdtl and E. Ivers-Tiffke, Sensors and Actuators B 5912-3, 184-189 (1999). [26] W. Menesklou, H.-J. Schreiner, R. Moos, K. H. HZirdtl, E. Ivers-Tiffke, Sr(Ti,Fe)O,: Materials for temperature independent resistive oxygen sensors, be published in MRS fall meeting proceedings, Materials for Smart Systems 111, Volume 604 (1999). [27] H.-J. Schreiner, Temperaturunabhhgige resistive Sauerstoffsensoren auf der Basis von Sr(Ti,Fe)O,. *, Ph.D. thesis, Karlsruhe (1999). [28] R. Waser, J. Am. Ceram. SOC.74 [8], 1934-40 (1991). [29] Ch. Tragut, Kinetik schneller Sauerstoffsensoren, Fortschr.-Ber. VDI 8 (291), Diisseldorf, Ph.D. thesis (1992). [30] Ch. Tragut, K.H. Hiirdtl, Sensors and Actuators B 4,425-429 (1991). [313 U. Lampe ,J. Gerblinger, H. Meixner, Sensors and Actuators B 7,787-91 (1992).
CERAMIC GAS TURBINE "CGT302" DEVELOPMENT SUMMARY Tetsuo Tatsumi", Isashi Takehara, Yoshihiro Ichikawa Kawasaki Heavy Industries Ltd., Gas 'hrbine Research & Development Center 673-8666 Akashi, Japan
ABSTRACT In Japan, a 300kW Ceramic Gas Turbine (CGT) research and development program was begun in 1988 as a part of "the New Sunshine Project" promoted by the Japanese Ministry of International Trade and Industry. The development target was to demonstrate the thermal efficiency of over 42% and low "Ox emission at 1350 "C of turbine inlet temperature (TIT). Kawasaki Heavy Industries, Ltd. (KHI) has been taking a part of this project and developed a regenerative two-shaft CGT, which was named CGT302. We achieved 42.1% of thermal efficiency and 31.7 ppm NOx emission, and also we conducted 2,100 cumulating operating hours at 1200 "C TIT, which is considered to be reasonable temperature for commercial use. This project was finished March 1999. For the next step development, MITI started "Research and Development on Practical Industrial Co-generation Technology". The project objective is to encourage prompt industrial applications of co-generation technology by confirming its reliability and soundness of Hybrid Gas Turbines (HGT).
INTRODUCTION Efficiency of gas turbines has been improved by utilizing higher TIT. In the field of large sized gas turbines, TIT has been steadily increased owing to progress of high temperature resistant alloys and turbine blade cooling technologies. For small gas turbines, however, it is difficult to manufacture air-cooled blades because of their size limitations. The purpose of using ceramics is to increase TIT without cooling air by utilizing their superior heat resistant characteristics. New Energy and Industrial Technology Development Organization (NEDO) started a CGT project in 1988 to develop three different types of CGT, namely CGT301, CGT302 and CGT303 under a contract with the Agency of Industrial Science and Technology of MITI. Kawasaki Heavy Industries Ltd. proposed the development of CGT302, two-shaft CGT with recuperator, in co-operation with Kyocera Corporation and Sumitomo Precision Products.
DEVELOPMENT TARGET The development target is shown in Table 1. This target was set to 42% ambitiously to have comparable efficiency with that of diesel engines expected at the end of this project.
Table 1 Development target Target
I
Thermal Efficiency
42 %
Turbine Inlet Temperature
1350 "C
I
Output Power
300 kW class Satisfy the national regulation
Exhaust Emission
DEVELOPMENT SCHEDULE The development schedule is shown in Table 2. The development phases were as follows. Phase 1: Basic design (Ceramic gas turbine) Phase 2: Basic MGT (900°C Metal gas turbine) Phase 3: Basic CGT (1200°C ceramic gas turbine) Phase 4: Pilot CGT (1350% ceramic gas turbine) The each phase corresponds to the part load condition of the final target of Phase 4, because the basic designs for all phases are the same.
Table 2 Development schedule
I
I
I
I
I
I
'88 '89 '90 '91 '92 '93 '94 '95 '96 '97 '98 Ceraiic Coiponet!t Fab!icatior! T&ch!-iology Component Technology (Turbine. Compressor, Heat Exchan
Interim AppraisLl Basic Design 900°C MGT 1200°C Basic CGT
1350°C Pilot CGT
45
ENGINE CONFIGURATION AND DESIGN FEATURES The design policy of CGT302 is as follows. To adopt conventional component layout to concentrate the development effort for ceramic utilization avoiding other difficulties. To adopt the metallic plate-fin type recuperator as a heat exchanger instead of a ceramic rotary regenerator which has been commonly used in the CGT for automobile use. To adopt module configuration to facilitate assemble and disassemble at numbers of test operation. CGT302 consist of single stage centrifugal compressor, single stage axial gas generator turbine, single stage axial power turbine, single can combustor, a plate fin recuperator, speed reduction gearbox for output shaft and other auxiliary equipment. Cutaway view of CGT302 is shown in Fig. 1.
Composite-component structure ('hrbine nozzle) The turbine nozzle is a large component consisting of complicated shaped nozzle vanes and an inner and an outer shroud. It is exposed in the highest temperature among the turbine components in which a large thermal stress may occur due to the temperature distribution on it. Accordingly, the segment configuration has been adopted to reduce thermal stress. This method has the merit of easy manufacturing but has demerits of stress concentration at the joining position and gas leakage between the segments. As a countermeasure for these demerits, we developed composite components. This concept is shown in Fig. 2. The fabricating process of the composite components is schematically shown in Fig. 3.
Binding
Impregnation (organwilimn Polymer)
'
.c
Heat Exchanger
Conversion
Repeat
.c Coating (To prevent separauon of fibers)
$. CompositeCompoment
Fig. 2 Turbine nozzle concept Fig. 3 Fabrication process Gas
Generator
/turbine L. Compressor
/
Scroll
Combustor Fig. 1 Cutaway view of CGT302 The main design features of the CGT302 are described below.
We can obtain a composite component that has a monolithic ceramic gas path with high heat resistance and a FRC backup ring with high fracture toughness. We applied this technology to the gas generator turbine (GGT) nozzle and the power turbine (PT) nozzle. By this technology, GGT and PT nozzle segments can be assembled tightly enough to seal gas leakage between the nozzle segments and also the centering function can be kept precise enough to minimize the turbine tip clearance. For example of this composite component technology, the GGT turbine nozzle is shown in Fig. 4.
ADVANCED UTILIZATION OF CERAMICS Si3N4, a candidate ceramic material, is much more brittle than metals and its thermal expansion coefficient is almost 1/4 of that of typical heat resistant alloys. Considering these characteristics, we have developed advanced utilization technologies for ceramics contrary to the conventional component layout. These technologies are as follows. Composite-component structure (Monolithic - FRC* hybrid component) *Fiber Reinforced Composite
Stress-free elastic supporting structure Stress-free independent supporting structure
46
Fig. 4 GGT turbine nozzle (Composite component)
Stress-free structure A stress-free elastic supporting structure and independent supporting structure were developed to protect ceramic components from excessive force derived from collision of components and loss of tightening force due to the thermal expansion difference between ceramic and metallic components and asymmetric deformation or expansion of metallic supporting structures. We did not adopt the conventional structure such as stacking up the ceramic components. We designed to support ceramic components independently from each other and to keep appropriate clearance between them. Accordingly, ceramic components can be displaced freely avoiding interference of each other when the metallic support is deformed asymmetrically. Ceramic coil springs and wave rings are properly utilized in the CGT302. The clearance between each ceramic component is sealed by ceramic elastic components such as coil springs or wave rings. These ceramic elastic components are also used for absorbing difference of thermal expansion between metallic support and ceramic components. Fig. 5 shows the configuration of the stress-free supporting structure. eolL SPRllBS
Fig. 6 Schematic Drawing of DLE ceramic combustor BLISK type turbine rotor There are two types of ceramic turbine configuration. One is "BLISK type that the blades and disk are made in one-piece, and the other is hybrid type, which is made by using a slotted metal disk and inserted ceramic blades. As these two types have advantages and disadvantages, we have to choose the better one considering application requirements. The BLISK type is difficult to make a large size turbine and has to be joined with the metal shaft, but it has better properties than the hybrid type in heat resistance, oxidation resistance, strength-weight ratio and manufacturing cost. We introduced a BLISK type turbine rotor for a GGT that is relatively small. And we introduced a hybrid type for PT rotor that has large diameter and could not be manufactured at the start point of the development. According to the development progress, manufacturing technology was improved and BLISK type PT rotor has become able to be manufactured. We changed PT design to BLISK from hybrid. Manufactured GGT and PT rotor are shown in Fig. 7. The outer diameter of this PT rotor is 192 mm and must be the biggest BLISK ceramic rotor in the world.
Fig. 5 Configuration of stress-free supporting Combustor The combustor is a rather simple shaped component among the ceramic components for CGT, but it is exposed directly in the combustion flame that has extremely large temperature distribution along both the radial and axial direction. The single can configuration we adopted inevitably causes asymmetry of the casing. Accordingly, careful design is required to support it freely from the excessive force caused by the interference of asymmetrical deformation of the combustor casing and associated casings while there is difference in thermal expansion between the metal casing and the ceramic combustor. We designed this component as a simple one-piece cylindrical combustor liner to facilitate a reliable support system and manufacturing. It is supported by the coil springs arranged to push it against the scroll gas inlet port. This configuration is shown in Fig. 6 Also excellent low NOx emission characteristics of dry low emission (DLE) combustion system can be expected by adopting the ceramic liner in which the maximum gas temperature decreases as the ceramic does not need cooling air.
Fig. 7 BLISK tvDe GGT and PT rotor The turbine tip clearance greatly affects turbine efficiency especially in these small size gas turbines. Ceramic BLISK turbine is more suitable than hybrid or conventional metal turbine to minimize the tip clearance, because the coefficient of thermal expansion of ceramics is low to minimize the tip clearance in various operation conditions. BLISK turbine also contributes to obtain better performance with no gas leakage between the blades.
47
ENGINE TEST RESULTS 900 "c Metal GT The MGT is manufactured as a test bed for the CGT and was tested to confirm the performance ability and propriety of the design concept of the CGT. At the MGT design, all ceramic components of the basic design are replaced by metal parts that have the same aerodynamic geometry as the CGT. The efficiency at the initial stage was only 17.3% and 42 kW output at 913°C TIT, and the final stage performance was improved to 23% efficiency and 59 kW output at 899°C TIT. The performance improvement is shown in Fig. 13 compared with the later CGT development progress. The cumulated operation time of MGT was 97 Hr and the number of start stop was 282 times. 1200 "c BasicCGT The high temperature metallic components in MGT were replaced with ceramics for the 1200°C Basic CGT. Ceramic material for Basic CGT was SN252 Si3N4. Before installing ceramic components, we conducted component tests such as the cyclic thermal-shock test for stationary components and the cold and hot spin test for turbine rotors. Before the all-ceramic configuration test, we conducted the confirmation test by installing single or minimum ceramic component step by step. We experienced unexpected obstacles through these tests, but after appropriate countermeasures, we have proved its durability through 40 hours of operation at 1200°C TIT, and also obtained 33% of thermal efficiency and 164 kW at 1190°C TIT as shown in Fig. 13. This R&D result of Metal GT and Basic CGT was evaluated by the neutral committee of industrial technology deliberation, and progressing to Pilot CGT stage was decided.
1350 "c Pilot CGT Fortunately the Basic CGT development was successfully demonstrated and we could step forward to Pilot CGT stage without a wholesale design change. Ceramic material was changed to SN281 for GGT rotor and SN282 for all other ceramic components from SN252. The difference between SN281 and SN282 is HIP process that is applied for SN281 only. And the turbine tip speed at rated condition was reduced to 480 d s e c from 570dsec as a countermeasureagainst FOD. We tested Pilot CGT increasing TIT stepwise by 50°C from 1200 "C at which successful operation was confirmed by the Basic CGT, and in these each step we confirmed stable operation and performance. To operate the engine safely according to increased TIT through these steps, we needed some modifications such as cooling of the joining part of GGT rotor and the clearance of roller bearings in the hot section. After these steps, we succeeded in 30 hrs of 1350°C TIT operation, but the target efficiency was not achieved.
48
To obtain further performance improvement, we redesigned the compressor to increase the airflow rate to get better matching to the turbine, because we found that the turbine nozzle throat area was slightly larger than the optimum value. We also improved the manufacturing accuracy of the turbine nozzle throat area because it had become evident from the engine test and dimensional inspection that the throat area fluctuated in each product and in each throat. Also we redesigned the GGT rotor blade shape to match the reduced rotational speed. But even after these efforts, the efficiency was staying around 40%, still slightly short from the target. To achieve the 42% target efficiency, improvements such as a larger heat exchanger with increased stack layer, an abradable PT shroud to minimize the tip clearance, and a PT rotor with thinner blades for improved turbine efficiency, and an improved compressor, were implemented. After several attempts, we achieved 42.1% at 1396 "C TIT with 322 kW and 41.5% at 1359 "C TIT with 302 kW in test #4-39. The pilot CGT performance improvement is shown in Fig. 8. CGT302 Performance Progress
i
1050
Target bilot CGT)
1
I
I
I
I
1150
1250
13M
1450
Turbine Inlet Temp.(oC)
Fig. 8 Pilot CGT Performance Improvement The performance comparison in the range of power output and efficiency improvement history is shown in Fig. 9 and Fig. 10, respectively.
10 1oo
10'
1o2
^.
*.
-
1o3. .
10' ......
10'
1o6
Fig. 9 CGT302 Efficiency Compared with Conventional GT The achieved efficiencies are comparable or superior to those of the largest gas turbines for power generation that have almost lo00 times the output power of the CGT. The cumulative operating time for the Basic CGT and the Pilot CGT was 334 hrs and the number of startstops was 645.
deterioration, turbine blade resonance and foreign or domestic object damage, but we could not determine what was the real cause of the GGT rotor damage. The endurance test results are as follows.
45, I I I I ............ .'.S'YFS 1i 1~5o a
g 401 Y
6
35
c - I
' m F l 1I
Target=1,000 Hrs (Weekly Start Stop) Cumulative operating time over 1,200 OC = 2,117 Hr 57 Min 1st.Try : 592 Hr, GGT rotor damaged : 519 Hr, DITTO, the combustor support ring survived 2nd.Try 3 r d . T ~ :782 Hr, DITTO, the combustor liner and support ring
@ Pilot CGT
0 Basic CGT
survived
4th. Try : 218 Hr, Combustor liner=1,000 Hr, Stop the test Other : 7 Hr, Check and adjustment Total cumulative operating time = 2,146 Hr 10 Min Longest time of the same assembly = 782 Hr Longest time for the same component (Combustor Liner) = 1,000 Hr Operation time of Heat Exchanger = 2,118 Hr
Year
Fig 10 Efficiency Improvement History Concerning to the NOx emission, we confirmed 9 ppm by the rig test and 31.7 ppm (02=16%) by the engine test at 1350 "c TIT. Measured NOx emission data is shown in Fig. 11 and Fig. 12.
- 95
P l = 8.1 .ta 11 = 717 C Ur = 15.6 J s
ae
E
valve P i l o t
._
:t
RESEARCH AND DEVELOPMENT ON PRACTICAL INDUSTRIAL CO-GENERATION TECHNOLOGY
CCI-901(ReaI Press)
100
H
?052
50X 251 10%
OX 01
TUlLt
40
60
80
Based on the successful results of 300 kW ceramic gas turbine development, MITI and NED0 started "Research and Development on Practical Industrial Co-generation Technology" project. The purpose is to encourage prompt industrial applications of co-generation technology that employs hybrid gas turbines (using both metal and ceramic parts in its high-temperature section) by confirming its reliability and soundness. The development activities are performed through material evaluation tests and long-term operation tests for a hybrid gas turbine of the medium size (8,000 kW class). It is expected that the development can realize low pollution by reducing the emission of C02 with highly efficient use of energy.
OX
t-6-cc)-
d
ffl
00 120
100
A/F kdkg 0.40.350.3 0.25 0.2 0.15 Eguivalsnce r a t i o
Fig. 11 NOx emission in rig test 100 ...............
h
80
-5 60 v
Main Fuel Ratio : 100% Prim. Fuel Ratio : .................................... I
I
I
I
I
.....................................
0 1100 1150 1200 1250 1300 1350 Turbine Inlet Temperature (deg-C)
Fig. 12 NOx emission in engine test
cumulative operating hours by total of three engines. The longest endurance operation hours by the same engine is 782 hours and the high time component is the combustor liner, which was operated to 1,OOO hours. We experienced three times of unexpected termination of the test by the engine damage. We supposed that it was initiated by the GGT rotor failure from the investigation results of damaged engine. There are several suspects
1
Items Output Power Thermal Efficiency Turbine Inlet Temperature Operation period Exhaust gas properties
Targets 8,000 kW class
I
34% or more 1,250oC 4,000 hours Standard regulation values or less
( Research on soundness and reliability )
Kyocera Corp. (Development and evaluation test of the heat-resistant ceramic comuonents ) Tokyo Gas Co., Ltd.
. Osaka Gas Co., Ltd. Toho Gas Co., Ltd.
49
I
FY 1999
2000
2001
2002 Interim Evaluation
v
ODevelopment and Evaluation tests of Ceramic Components @Research on Soundness and Reliability Basic design Detailed design Manufacturing Operation @General Investigation of the System
50
2003 Final Evaluation
b
b b
b
REFERENCES (1) Yamagishi, K., Yamada, Y., Echizenya, Y., Ishiwata, S., 1989 "Current Status of Ceramic Gas Turbine R&D in Japan", ASME, 89-GT-114 (2) Nagamatu, S., Mizuhara, K., Matsuda, Y., Iwanaga, A., Ishiwata, S., "Current Status of Industrial and Automotive Ceramic Gas Turbine R&D in Japan", ASME, 91-GT-10 (3) Shimada, K., Ushijima, H., Yabe, A., Ogiyama, H., Tsutsui Y. "Advanced Ceramic Technology Developed For Industrial 300 kW CGT (Ceramic Gas Turbine) Research And Development Project In Japan", ASME, 93-GT-188 (4) Arakawa, H, Suzuki, T, Saito, K, Tamura, S, Kishi, S, "Research and Development of 300kW class Ceramic Gas Turbine Project in Japan", ASME, 97-GT-87
PREPARATION OF PLANAR SOFC-COMPONENTS VIA TAPE CASTING OF AQUEOUS SYSTEMS, LAMINATION AND COFIRING Bernd Bitterlich*’,Christiane Lutz, Andreas Roosen University of Erlangen-Nuremberg, Department of Materials Science, Glass and Ceramics, D-91054 Erlangen, Germany *) Present address: Technical University of Clausthal, Institute for Nonmetallic Materials, D-38678 Clausthal-Zellerfeld, Germany
ABSTRACT The Solid Oxide Fuel Cell (SOFC) gains a high economical efficiency but its manufacturing costs are too high to compete successfullyat the market. Tape casting via the doctor-blade method is used as a low cost process to produce flat components for the planar SOFC. A decrease of the manufacturing costs could be achieved if the different green sheets would be laminated and then cofired in a single firing step. A further advantage for low cost manufacturing would be the replacement of organic solvents by water in the tape casting process. Organic solvents which are widely used in the industry are flammable, explosive, and a hazard to environment and health. In this paper anode and electrolyte of the SOFC were prepared by tape casting of aqueous slurries. To achieve a highly porous anode graphite was used as a filler which burns out after binder burnout. The anode/ electrolyte composite structure is formed by lamination by the thermo-compression method. The laminates were cofired. Defect-free connections between the porous anode and the dense electrolyte were obtained. The paper discusses the results in detail.
using water as solvent safety considerations during processing can be kept low [9]. But it is more difficult to obtain green sheets of high quality. This is due to the different physical and chemical properties of water compared to organic solvents, which have a strong impact on the drying behavior. This has to be taken into account if defect-free green tapes of the electrolyte and the anode are developed on water based slurries. The lamination of the different tape cast sheets is a fast processing step and can easily introduced in a production line. Energy and time are saved by cofiring of the laminated green sheets. In the common concept of planar SOFCs the electrolyte is the substrate for the electrodes and must have a certain thickness to guarantee the mechanical stability [ 10, 111. Reducing the thickness of the electrolyte would reduce its electrical resistance making it possible to lower the operating temperatures while gaining the same power density [12]. In addition the lower operating temperature would reduce the complex material problems of the SOFC-periphery which cause high costs. This concept is realized in the design proposed by the Research Center of Jiilich, which consists of a thick porous anode supporting a thin 20 pm electrolyte layer and the second electrode (Fig. 1) [13]. Cathode -50 p m
INTRODUCTION Due to its electrochemical energy conversion the solid oxide fuel cell (SOFC) offers the advantages of high efficiency and low pollution in comparison with conventional power stations [l]. But its spreading is still hindered by the high manufacturing costs [2]. The main functional components of the SOFC are the electrolyte and its two electrodes. Fuel (e.g. hydrogen) is oxidized on the anode side while the oxidant (oxygen or air) is reduced on the cathode side. Oxygen ions diffuse through the dense electrolyte via oxygen vacancies in the lattice which results in an electrical voltage. Components of a planar SOFC can be produced by the economic tape casting process [3, 41. Lamination and cofiring of the single green sheets would reduce the production costs drastically. Cofiring of components for the SOFC has been carried out by some authors but on different processing routes [5-71. In this work tape casting on the basis of aqueous slurries was used to fabricate the SOFC components. In the industry organic solvents are widely used. They are flammable, explosive, and a hazard to environment and health [8]. By
Anode
-50pm
a) Electrolyte is substrate
b) Anode is substrate!
Fig. 1: Planar SOFC designs: Substrate is either the electrolyte (a) or the anode (b). The single components must fulfill several requirements: The electrolyte, which consists of fully yttria stabilized zirconia (YSZ), must be impervious to gases. In order to reduce the electrical resistance enabling a lower working temperature of the SOFC, the electrolyte must be very thin (< 20 pm). The anode consists of a Ni/YSZ-Cermet, which is formed in a reducing step, when NiO is transferred to Nickel. The anode has various functions: it is the load bearing substrate for electrolyte and cathode and therefore it must have a sufficient thickness (appr. 1,5 mm)
51
to provide an appropriate mechanical strength. The anode must have a high gas permeability for the firing gases. This means it must have an interconnected POrous structure [12] of a specific dimension. Nickel and YSZ must be homogeneously distributed but the Ni must form a continuous phase to facilitate electronic conductivity.
Lamination, binder burnout and sintering
Because cofiring means to fire different materials at one temperature the sintering shrinkage and its rate should be as similar as possible to achieve defect-free samples [14-161. No diffusion between the layers should occuf. The thermal expansion coefficient of the sintered ceramics must be adapted.
The tape cast sheets were laminated by the thermocumpression method. The anode tapes were moisten with water, that resolves the PVA binder to improve the cohesive strength. Thermal compression was carried out at 1 MPa, 50°C and a holding time of 5 min. Stacks of several anode sheets and one electrolyte sheet at the top were laminated. To prevent camber symmetric stacks (anode-electrolyte-anode)were made as well. The dried laminated samples were slowly heated with 25 K h up to 650°C and presintered at 1200°C for 3 h (heating rate 60 Kh). The final sintering was carried out in another furnace at 1400°C for 3 h with a heating rate of 300 Kh).
EXPERIMENTAL PROCEDURES
Measurement of porosity, permeability and strength
Slurry compositions For the electrolyte a very sinter active 8 m o l I yttria stabilized zirconia powder (Tosoh, Japan, dso = 0,26 pm, ABm = 16 m2/g) was chosen. After several investigations an ammonia salt of a polyacrylic acid as a dispersing agent was used. An acrylic polymer emulsion (Mowilith, Hoechst Perstorp, Sweden) was used as a binder. The binder emulsion was chosen because these kind of binders are ready-to-use and achieve a high density and tensile strength of the green sheets [17]. Furthermore they are successfully used by other authors [18, 19,20,21]. The slurry for the anode consists of a mixture of NiO (Baker, USA) and YSZ powder. This Y203fully stabilized zirconia powder (Unitec, UK,d50 = 0,84 pm, ABET = 7,7 m2/g) has larger grains and a smaller specific surface area than the one which was used for the electrolyte, because a lower sinter activity is needed to o b tain a more porous substrate. To disperse the two powders another polyelectrolyte was used as a dispersant. To achieve a high porosity graphite powder (Timcal KS6, CH, d50 = 3,7 pm, ABET = 19 m2/g)was added. The graphite burns out after the binder burnout, but before sintering occuzs. Dispersion of the graphite in water is very difficult but could be managed in a solution of polyvinylalcohole. PVA (Mowiol, Hoechst, Germany) was also used as the binder. This binder needs a plasticizer (Glycerol, Merck, Germany) to improve flexibility and to lower the Tg. In case of the anode, no high green densities are necessary because porosity is desired.
Slurry preparation and tape casting The slurries were prepared by dispersing the powder in water plus dispersing agent by ball milling for 66 hours. Then the binder emulsion or the PVA solution plus graphite, respectively, was added. The slurries were then homogenized for further 24 hours. After degassing the slurries were tape cast on a tape casting machine with stationary blades on a polypropylene carrier film. The casting width is 25 cm.The cast slurries were dried at room temperature.
52
The open porosity of the sintered anode substrates was investigated by mercury porosimetry (Porosimeter 2000, Car10 Erba Intr., Italy). The permeability was determined using a modified blaine-apparatus (after DIN EN 196, part 6). The evaluation was done considering the equation of Darcy [22]. The strength was measured in a 3-point bending test. The sample geometry of laminated substrates (2 layers) was 11 x 33 mm*.The samples were tested as fned to determine the strength including all defects of the surface. For each type of laminate 20 samples were tested.
Measurement of thermal expansion and shrinkage The thermal expansion of sintered samples was determined in a dilatometer (Netzsch, Germany) with a heating rate of 5 Wmin up to 1400°C in air. Shrinkage during binder burnout and sintering was investigated on laminated green sheets in the same dilatometer. Laminated stacks of about 5 x 5 x 101111113 were put in the dilatometer in such a way that the shrinkage parallel to the layers was measured. Up to 650°C the heating rate was 1 Wmin. Afterwards it was increased to 5 Wmin up to 1400°C.
RESULTS Tape casting and properties of green sheets The slurries of the electrolyte and anode could be tape cast easily without the formation of cracks during drying. The slurries showed the desired strong pseudoplastic rheological behavior. AU sheets could be separated fiom the carrier film without problems. The tapes showed a drying shrinkage in height of about 65 %. The thickness of the electrolyte sheets was about 50 pm. In spite of the low thickness they had a high green strength and very high flexibility. The green density was about 2,8 g / c m 3 (including the binder) which corresponds to 52 % of the theoretical density. Anode tapes of about 500 pm thickness were received. These sheets showed a sufficiently high strength and flexibility. The green density was in the range of 2,4-2,9 g/cm3depending on the slurry composition.
Properties of sintered sheets After sintering the electrolytes had a thickness of 30 to 35 pm. The samples were transparent due to their high density of more than 97 % of theoretical density. The measurements after archimedes' principle were difficult to carry out because of the low mass of the samples due to their small thickness. The microstructure showed only a small quantity of closed porosity (Fig. 2).
anodes prepared by tape casting with organic solvents (TC) [23] and by a coat-mix process (CM) [12] are added in Fig. 4, too.
120-
p"
100-
5 805 F 60t
402
0 15
20
25
30
35
40
a 0,o 45
Open porosity Nol.-%
Fig. 4: Strength and permeability of anode substrates versus open porosity; values of anodes prepared by other methods (TC [23], CM [ 121) for comparison.
Fig. 2: Fracture surface of a sintered electrolyte.
Fig. 4 also shows the strength of the laminated anode structures versus porosity. The weibull modulus of the samples, which were tested as-fired,has a value between 7 and 8. This relatively high value means that the substrates have a small pore size distribution and do not contain great defects in the microstructure.
The sintered anode substrates had a thickness of about 350 pm. Fig. 3 shows the open pore structure of a substrate which contained graphite to increase porosity. Depending on the quantities of binder and graphite different porosities and pore diameters were obtained. Without graphite the substrates exhibit only a low porosity of 15 v01.-% and a small pore diameter of 0,3 pm. Addition of graphite increased the average pore diameter to values of about 1 pm as well as the total porosity. A porosity of 35 v01.-% was achieved by a graphite content of v ~ p h i * $ ( v ~ p h i ~ ~ + v , , , ) of 12 %.
Asymmetric laminates, i.g. anode-electrolyte, had a very strong tendency for camber and bending. If the samples were sinkred without a load, the electrolyte and anode exhibit only insufficient joining or were strongly curled. A higher load during cofiring hindered the shrinkage of the samples and generated a lot of cracks. Thus, symmetric stacks made of layers of anode, electrolyte and
tween porosity and permeability an exponential-curve can be fitted (Fig. 4). For comparison the properties of
Fig. 5: Symmetric cofired laminate (anode/ electrolyte/ anode) showing a crack in the electrolyte.
Properties of cofired samples
53
DISCUSSION Electrolyte
Fig. 6: Good joining between electrolyte and anode after cofiring.
Examination of defect causes In spite of the different binder systems the green sheets have quite a similar binder burnout characteristic.DTAanalysis shows that the binder was completely burned out at temperatures above 450°C. The decomposition of the graphite of the anodes is completed at 600°C. The shrinkage and its rate during heat treatment is quite different for the two components (Fig. 7). The porous remaining anode substrate has a lower shrinkage than ~ . . . . , . . . . , . . . . ,
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The thermal expansion of the sintered samples (Tab. 1) show only small differences. Thus, compared to the difference in sintering shrinkage only small stresses can develop during cooling from the sintering temperature. Table 1: Coefficient of thermal expansion (CTE). /104K' Sample CTE20/1200"c Electrolyte 10,3 Anode 12,9
54
CTE20~1400T /104R' 11,l 13,O
Crackfree green sheets of high flexibility and strength could be tape cast of aqueous slurries. The thickness of the sintered electrolyte was appr. 30 pm. The developed slurry system has the potential to be cast to thinner green sheets of 20 pm thickness because the green sheets show sufficient strength. In case of such thin tapes, more care must be taken to avoid any stretching of the green sheets during handling. In spite of the relatively low green density of max. 52 % of theoretical density the sintered thin sheets are impervious to gases with a very small remaining porosity. This is connected with a high sintering shrinkage. An increased green density would decrease the remaining closed porosity and would lower the sintering shrinkage. This would be advantageous for the cofring process.
Anode Tape casting of the aqueous anode slurries lead to crackfree and flexible green sheets of appr. 500 pm thickness. The required thickness of the anode substrate is 1,5 mm. Thus 4 to 5 green sheets have to be laminated to achieve this thickness after sintering. The cast thickness could be increased thus only 2 sheets would have to be laminated but then the drying time of the cast slurry is increased. The porosity depends strongly on the binder and graphite type and content and can be adjusted in a wide range. The graphite addition leads to an average pore diameter of about 1 pm with a narrow pore size distribution. In further processing of the final component the NiO is reduced to metallic Ni. In this step hardly no shrinkage but an increase in porosity and permeability by 3-4 times will take place. The permeability of the samples is in the range of the reference substrates in the literature indicating that the pore structure is as homogenous as in anodes prepared by other processing routes. The strength is quite low due to the high porosity and the unpolished surEace of the samples, but again comparable with the reference samples and sufficient for the application. The shrinkage of the anode is relatively small, because only a limited densification is desired. Thus it is difficult to increase the shrinkage of the anode to decrease the difference in shrinkage in comparison with the electrolyte.
Cofired samples After sintering a good joining between electrolyte and anode OcCuITed. But the samples showed cracks in the middle of the electrolyte. Thus no crack-free samples of larger size could be produced. Crack free areas were about 10 mm in square. Different sintering shrinkage, different thermal expansion behavior during cooling and defects or inhomogeneities in the green tapes are the main causes for crack development. The coefficient of thermal expansion of the two components do not differ much. The degree of defects was also limited due to an extented deagglomeration treatment and degassing of the slurries. The main contribution for defects is the
sintering shrinkage. Due to the high sintering activity of the fine YSZ powder the electrolyte shrinks much more than the porous remaining anode. From the dilatometer curves the stresses in the electrolyte layer can be calculated (Fig. 8). The stresses at temperatures below sintering were set to zero although there had been some shrinkage differences in this temperature range due to shrinkage during binder burnout. The total shrinkage of the electrolyte is much higher than that of the anode so the electrolyte will finally be under high tension stresses.
- pressure $O-
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CONCLUSION In this work two components of the SOFC, anode and electrolyte, were prepared by aqueous tape casting. The green tapes were crack-free and easy to handle due to their good flexibility and green strength. The sintered sheets fulfilled the required properties: The electrolyte had a thickness of only 30 pm and a high density of more than 97% of theoretical density being impervious to gases. The high open porosity of the anode was achieved by adding graphite which burns out leaving pores. By varying the graphite content the porosity could be adjusted. It was shown that porosity and permeability of the prepared anodes are comparable with samples made by other authors using different techniques. The green sheets were laminated by thenno-compression and using water as a lamination aid. After cofiring the samples had defects due to different sintering shrinkage. During sintering the electrolyte shrinks much more than the porous remaining anode. Thus the electrolyte is under tension stresses leading to cracks. Small crack-free areas demonstrate that a joining by the described technique is possible. To achieve defect-free samples of larger size there are several possibilities: Increasing the green density of the electrolyte green sheets would lower their sintering
shrinkage. Another attempt would be to increase the sintering shrinkage of the anode by using the same very fine zirconia powder as for the electrolyte together with a higher binder or graphite content to gain the same porosity after sintering. A third possibility is to introduce an additional layer between electrolyte and anode with an intermediate porosity and composition to have a graded transition between the components.
REFERENCES [l] N. Q. Minh, Ceramic Fuel Cells, J. Am. Ceram. SOC.76 (1993), 563-588 [2] M. C. Williams, Status and Market Applications for the Solid Oxide Fuel Cell in the U.S., Proc. 3'h European Solid Oxide Fuel Cell Forum, Nantes, France 1998,27-38 [3] R. E. Mistler, D. J. Shanefield, R. B. Runk, Tape Casting of Ceramics, G. Onoda, L. L. Hench (eds.), Ceramic Processing before Firing, John Wiley & Sons Ltd., New York 1978,411-448 [4] A. Roosen, Basic Requirements for Tape Casting of Ceramic Powders, Ceramic Transactions Vol. 1B: Ceramic Powder Science, The American Ceramic Society, Westerville, OH 1988,675-692 [5] M. Maridek, J. Maikk, Cofire Processing of Electrolyte, Anode and Interconnect Materials for Solid Oxide Fuel Cell Fabrication, British Ceramic Proceedings 60 (1999), Vol. 1 [6] T. Kawada, N. Sakai, H. Yokokawa, M. Dokiya, I. Anzai, Fabrication of a Planar Solid Oxide Fuel Cell by Tape-Casting and Co-Firing Method, J. Cerum. SOC.Jpn. 100 (1992), 838-841 [7] S. A. Wallin, S. Wijeyesekera, Y. Chiao, M. J. Semer, Cofired Solid Oxide Fuel Cells for Operation at 800 "C, Electrochem. SOC.Proceedings 13 (1997), 157-167 [8] R E. Mistler, Tape Casting: The Basic Process for Meeting the Needs of the Electronic Industry, Ceram. Bull. 69 (1990), 1022-1026 [9] D. Hotza, P. Greil, Review: Aqueos Tape Casting of Ceramic Powders, Mat. Sci. and Eng. A202 (1995) 206-217 [lo]W. Drenckhahn, H. Greiner, E. Ivers-Tiff&, Materials for Solid-Oxide High-Temperature Fuel Cells, Siemens Power Journal 4/94 [11]H.J. Beie, L. Blum, W. Drenckhahn, H. Greiner, H. Schichl, Development of Planar SOFC at Siemens Status and Prospects, 3d European Solid Oxide Fuel Cell Forum, Nantes 1998, France [12]F. J. Dias, A. Naoumidis, D. Simwonis, D. Stover, H. Thiilen, Properties of NirYSZ Porous Cermets for SOFC Anode SubstratesPrepared by Tape Casting and Coat-Mix Process, AMPT '97, Portugal 1997, 1-5 [13]H.P. Buchkremer, U. Diekmann, L.G.J. de Haart, H. Kabs, D. Stover, I.C. Vinke, Advances in Maufacturing and Operation of Anode Supported SOFC Cells and Stacks, 3'd European SOFC Forum, Nantes, France, P. Stevens (ed), Fuel Cell Forum, (1998) pp. 143-149.
55
[14]C. Hillman, Z. Suo, F.F. Lange, Cracking of Laminates Subjected to Biaxii Tensile Stresses, J. Am. Ceram. SOC. 79 (1996), 2127-2133 [151R. Natarajan, J.P. Dougherty, Material Compatibility and Dielectric Properties of Cefired High and Low Dielectric Constant Ceramic Packages, Electronic Components and Technology Conference 1997,750-754 [16]P.Z. Cai, D.J. Green, G.L. Messing, Constrained Densification of AluminaEirconia Hybrid Laminates, I: Experimental Observations of Processing Defects, J. Am. Ceram. SOC.80 (1997), 1929-1939 [171N. R. Gurak, P. L. Josty, R J. Thompson, Properties and Uses of Synthetic Emulsion Polymers as Binders in Advand CeramicsProcessing, Am. Ceram. SOC.Bull. 66 (1987), 1495-1497 [18]T. Ueyama, N. Kaneko, Effect of Agglomerated Particles on Properties of Ceramic Green Sheets, High Tech Ceramics, P. Vincenzini (ed.), Elsevier Science Publishers B.V., Amsterdam 1987 [19]N. Ushifusa, M. J. Cima, Aqueous Processing of Mullite-Containing Green Sheets, J. Am. Ceram. SOC.74 (1991), 2443-2447 [20]A. Kristoffersson, E. Carlstrom, Tape Casting of Alumina in Water with an Acrylic Latex Binder, J. Eur. Ceram. SOC.17 (1997), 289-297 [21]C. Pagnoux, T. Chartier, M. de F. Granja, F. Doreau, J. M. Ferreira, J. F. Baumard, Aqueous Suspensions for Tape-casting Based on Acrylic Binders, J. Eur. Ceram. SOC.18 (1998), 241-247 [22]H. Heuschkel. G. Heuschkel, K. Muche, ABC Keramik, VEB Dt. Verlag fiir GrundstolTindustrie, 2. Auflage, Leipzig 1990 [23]Ch. Lutz, A. Roosen, D. Simwonis, A. Naoumidis, H. P. Buchkremer, Foliengiekn eines por6sen Anodensubstrats fiir die Hochtemperatur-Brennstoflkelle, Werkstoffwoche (Materialica), Miinchen 1998
56
GLASSES FROM THE SYSTEM RO-R203-Si02 AS SEALANTS OF HIGH CHROMIUM STEEL COMPONENTS IN THE PLANAR SOFC P. Geasee' ,T. Schwickert"",U. Diekmann'*, R Conradt' (*) Institute of Minerals Engineering, Chair of Glass and Ceramic Composites,
Mauerstr.5, D-52064, RWTH, Aachen, Germany (**) Forschungszentrum Juelich GmbH, Central Department of Technology, D-52425, Juelich, Germany
Abstract
Introduction
Planar solid oxide fuel cells (SOFCs) are a promising approach for cost effective electrochemical energy conversion at an operation temperature of 800 "C. They require glass ceramic sealants which have a high thermal expansion coefficient (11.0-12.0 x lo-%-'), high electrical resistance and good thermochemical compatibility and stability with the gases and the fuel cell materials. Different glass types based on the RO-R203-SiOzsystem were melted, crystallized and investigated for thermal expansion, wetting behavior, thermal analysis, long term evaporation stability, and joining properties. Experimental results showed a good adhesion of barium rich glass ceramics to both high chromium steel (18 YO Crz03) and ceramic substrates. SEMEDX was used to investigate the interfaces. Long term evaporation stability of barium rich glass under H2 and HzO atmosphere was observed as a 0.2-0.4 pg/cm2.hmass increase after 2000 h.
Glass or glass ceramics are well known for ceramic, glass or steel joining [1,2,3]. Many of these products are available as commercial products. Their applications depend on sealing temperature, thermal expansion, viscosity and joining ability (wetting and surface tension). Most of the sealing products are produced for low temperature joining and for room temperature application. Figure 1 shows the relation of thermal expansion (a) and glass transformation temperature (Tg) of commercial glass sealants. The goal of a new glass sealing development is also plotted in this picture. Barium calcium aluminosilicate (BCAS) and barium calcium magnesium aluminosilicate (BCMAS) have been studied for high temperature SOFC stack by many researchers. Lahl N. [4] studied BAS, CAS and MAS glasses by using BzO3 as a fluxing agent and using TiOz, ZrOz, Ni and Crz03 as a nucleating agent. She found that barium glass tended to react with chromium steel more than calcium and magnesium glasses. sealing for forsterite, steatite
15
....................................
0.................. Pb-Zn-B-Bi-Ba,Na2Ljo!
Pz05.
..
10
5
0
400
800
1200 Tg ( "C)
Fig. 1 Plot of a of d ~ e r e nglasses t vs. Tg
57
Most of her studies relied on MAS glass with low thermal expansion coefficient (a = 7-9.10%'). Only one BAS glass compositionhad rather high a (13-14 x lo%-'). Doersing and Conradt [5] developed BCAS glasses for high sealing temperatures (1-4 wt. % Al203) by Using Nd24, La203 and Y2O3 as additives. They found that La203 increases a, Tg and TM (according to the ionic radius) more than Nd203 and Y203 respectively. Heilemann and Conradt [6] studied BCAS glasses by using La203, B203, MnO and Zr02 as additives. From their studies, high a glasses but poor joining was found. Schwickert, Geasee et al. [7,8] investigated BCMAS, BS and BCAS glasses The aim of this paper is an adjustment of barium calcium silicate glasses from [7, 81 yielding better joining properties (good sticking, and gas tightness), high thermal expansion, stability under H2 and H 2 0 atmosphere and chemical compatibilitybetween the joining partners.
the solid/liquid interface from heating microscope can be calculated by
where h is the height and r is the radius of the sample. It is also possible to measure the wetting angle directly form the video images of the drop on the solid substrate.
T1
T2
T3 Fig. 1 Three characteristictemperatures of gradual softening
Experimental Table 1 shows the glass compositions investigated, based on a calculation of constitutional compounds. From the calculation, C2BS3, BS, BAS2 and B2S3 were found as major phases. Additives of ZnO lowered viscosity (q) with less effect on thermal expansion than B2O3 [9] while S r O acted as a controlling agent for crystallization [lo]. An effect of these minor oxides was compared to the other barium glasses [111. Nine glass compositions (table 1) were melted from chemical grade raw materials at 1480 "C for 2 h in a Pt crucible in an induction furnace. During soaking at 1480 "C, the glasses were stirred to enhance homogeneity. Glasses were fritted, or cast into bar forms for different investigation purpose. A dilatometer was used for a, Tg and TM determination of both glasses and partially crystallized glasses. The crystallizationprocess was performed by sintering glass powder samples at 900 "C for 10 h.
The melting behavior and wetting angle were studied by heating microscope at a heating rate of 2 Wmin up to complete melting temperature in air. For this purpose, glass powder was pressed in pellet form and placed on steel (no. 1.4740) substrate. Start of sintering temperature el),ball point temperature (T2), complete melting temperature (T3) and the contact angle at T3 were determined (see figure 1). The contact angle (0) of
58
Several joining tests were performed. A quick sticking test of glass pieces (2mm . 1 cm . 1.5 cm) with steel was performed at 900 "C for 2 h. As a simulation of horizontal joining, sandwich samples of two steel plates with glass paste (glass powder mixed with organic) in-between was fired to the same temperature. After firing, a gas tightness test was performed. Then 2 steel plates were broken and the joining behavior was investigated. A simulation of vertical joining of a dummy stack consisting of a conventional housing and a cube of steel were joined with glass paste. The fired dummy stack was also used for gas tightness measurements. Long term evaporation stability was determined from weight loss of glass chips in a tube furnace (flow with 833 mbar N2 + 93 mbar H2 + 74 mbar H20 at 800 "C) after time intervals up to 2000 h. Beside this, thermal gravimetq (TG) was used to confirm the evaporation mechanisms. Crystallization behavior and phase content were investigated by DTA/DSC and XRD. The microstructure of interface layers and of crystallized glasses was studied by light microscope and SEM.
Results and discussion Table 1. Chemic; composition of investigated glasses codes SO2 25 GR50 44 GJ40 35 GJ3 1 34 B57 32 B58 33 B59 32 C6 1 33 C62 31 D70 34 D7 1 30 D72 E73 34
Figure 4 shows a high thermal expansion coefficient of glasses B57 to E73 as measured from glass samples and sintered samples (900 "C, 10 h ) in the range of 20-600 "C. Most of them had hgher a after crystallization except for glass GJ40 because of its content of low a compounds, i.e. diopside (CMS2) and celsian (BAS*) [12,13].
Glass compositions with various additives for lowering viscosity, Tg, TM, surface tension, enhancing a and crystallization are shown in table 1. Figure 3 shows values of Tg and TM. 0 0
0
.... .... 0
-
O r J
0
0 .#
Melting behavior from heating microscope and wetting angles of Werent glasses are shown in figures 5 and 6. Glass GJ40 has a lower ball point temperature (T2) than glass GJ31 by 130 K. This was the reason why glass GJ40 had better sealing behavior at 900 "C than glass GJ3 1. The complete wetting temperature (T3) of glass GJ40 is very high if compared to the other glasses because of the crystallization occurring during heating. According to dilatometric measurement, glass GJ31 also crystallized, but this effect could not be found by heating microscope.
0
0
-
0
T-
' 6
0
o
. . . . . . . . . .
8 *g F2 m2 m2 -8 muQoaF* aFmaFg w codes Fig. 3 Tg and TMof investigated glasses 12
P 1400
-
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1200-
; 1100-
GJ3 1 0
P)
....
+ al
..Q
.......
I
.
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9
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800- ............................. . gGR50
-
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ballpoint(T2)
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0
*
.
0
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...................................*........... I
.
I
.
I
.
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.
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Fig. 4 Thermal expansion of glasses
59
Glass GJ31 had a lower crystallization rate than glass GJ40, which cannot be detected at 2 Wmin in a heating microscope. Glass GFUO had very low ball point temperature ("2) because of its high B2O3 content. This glass started to melt at 743 "C. The wetting angle was detected as 54.6" at 850 "C, which was rather high if compared to the other glasses (25-45"). The effect of its high contact angle may stem from MgO.
40.
60 n
55
%
50
'
30
!
I
'
25
codes Fig. 6 Wetting angle of different glasses at (T3) Glasses B57, B58 and B59 were adjusted from glass GJ31 by small addition of alkali and fluxing oxides. T1, T2, T3, and the wetting angle strongly decreased and yielded better joining. Further development of glass C61 and C62 result in a decrease of T1 and increase of T2 and T3. Glass D70,71 and 72 were adjusted from glass B58. All three glasses had a lower T1. However, glass D72 presented very high T2. This is because of strong crystallization as detected by DTA (see figure 7). The relation of T1, T2 and T3 was found to have a significant effect on the joining properties. Very good joining properties with a good gas tightness was found for glass D7 1. Figure 7 shows an example of DTA signals for glasses B58, D70, D71 and D72. Glass D70 and D72 showed a sharp exothermic peak at 872 "C and 893 "C respectively. The constitutional compound calculation revealed C2BS3, BaO.BZO3 and BASz as major phases of these glasses. An experimental results of the melting behavior, wetting angle and crystallization behavior lead to a conclusion that the glass with good joining properties at a given joining temperature Tj (e.g., 800 or 900 "C ) must have
60
Fig. 7 DTA signal of glasses B58, D70, D71 and D72 start of sintering temperature definitely below Tj, ball point temperature not too far above Tj (m 40-80 K), a low wetting angle at T3 and, a crystallization velocity slow enough so that the spherical shape at T2 can form under a heating rate of 2 Wmin. However, low wetting angle at low temperature does not yet mean a good joining for SOFC at 800900 "C. A quantity is proposed joining rate J. The joining rate J in m/s is calculated from whch shall be
where p = density of glass in g/cm3,g = standard acceleration = 9.8 m/sz, A = joining surface area (cm'), = surface tension (N/m) and q = viscosity of glass in @as. If the density of barium glass at 3.5 g/cm3 and a joining surface area of 1 cmz are used, the value of the joining rate can be calculated as shown in table 2. Table 2. Joining rate calculated form equation (2); q in @as surface joining rate J in m/s at tension (N/ml log q=11.3 I log q=7.6 I log q=4.5 I 1.18E-081 6 E-05 I 7.44 E-04 0.5
I
1
The change of viscosity has a much stronger effect on the joining factor than the change of surface tension. Thermal expansion (a)of ceramic substrate, steel, glass and partially crystal glass are plotted in figure 8. The a of partially crystallized glass is closer to the steel and ceramic substrate than the a of the glass. 14
12
Long term evaporation as determined by a tube furnace transpiration test showed 0.23 pg/cm2.h mass increase at 1500 h for glass GJ3 1, but a mass decrease of 1.5-1.7 pg/cm2-h when B203 was added. This is still acceptable if compared to the high evaporation rate of glass GR50 with high boron content. With the confirmation of on-line weight change measured by thermal gravimetry, the evaporation kinetics of glass GR50 was investigated as a competition of two simultaneous evaporation mechanisms. These are the loss of a matrix component from the surface and the selective loss of a highly volatile compound by diffusion and evaporation.
Summary
5x 10 d
8
I
I
I
I
I
I
I
I
200
400
600
800
1000
temperature ("c)
Fig. 8 Thermal expansion coefficients of glass, partially crystallized glass (sintered at 900 "C, 10 h), ceramic substrate, and steel Figure 9a shows that a good joining to areas reached by the glass through gravity is less critical than joining towards areas above the sealant. As an example, glass GJ31 reached good joining to the lower partner (figure 9a), but a poor result with the upper partner (figure 9b). In this particular case, gas tightness was not reached.
Different glass types base on the RO-R203-Si02 system were developed after the requirements of a high thermal expansion coefficient (11.0 - 12.0 x lo-%-'), high electrical resistance and good thermochemical compatibility with the fuel cell materials. Experimental results from table 3 shows a good adhesion of barium rich glass ceramics to both high chromium steel (18 % Cr2O3) and ceramic substrate. Further adjustment of viscosity, surface tension and crystallization velocity by additive yields gas tight joining. There are several factors influencing the sticking of glass with steel. The most important factors are (1) a proper viscosity, (2) low surface tension and (3) slow crystallization velocity. Long term evaporation stability of barium rich glass under H2 and H20 atmosphere was observed as a 0.2-0.4 pg/cm2.h mass increase at 2000 h (:. 0.2 pg/cm2.h at 1500 h).
Fig. 9 Microstructure of the joining interface between glass and steel, a) excellent, b) poor joining
61
*
Table 3. Physical properties of merent glasses codes properties B58 GJ3 1 845 890 T1 ("C) 862 1128 841 900 1020 900 780 T2 ("C) 1146 1164 1080 1148 1070 914 1112 1137 1153 1175 1250 850 T3 ("C) 41.3 46.6 1 37.3 38.6 I 39.7 39.7 39.5 I 29.0 wetting angle " 56.0 54.6 54.6 44.0 (1146) (1164) (1080) (1165) (1070) (1125) (1180) ( at T in "C) 10.07 10.26 10.71 10.53 11.32 10.77 11.44 10.3 10.56 ax K-' a 654 661 625 622 629 621 719 Tg ("C) 720 681 668 672 685 714 662 658 727 733 769 TM ("c)
I I Pp408
7-Y
+ I I -
10.51 11.52 a x 10-6K-' Evaporation" -0.23 0.69 3.59 -0.05 1.70 (P&n2.h) lo9 gas tightness a : glass sampl :,b:sa iple sintered at 900 "C 10 h, c units of mbml i form s mples sintered for 10 i at 800 C61, C62, D70, and D72 was no surprise.
- I
7
I
References 1. U. Diekmann : High temperature fuel cells- a challenge to joining techniques, DVS-Berichte Band 166, Duesseldorf 1995 2. P.Batfalsky, U.Diekmann, J. GodziembaMaliszewski and T. Koppitz : Joining and sealing of anode-supported planar SOFC stacks using glass and glass ceramics, DVSBerichte Band 184, Duesseldorf 1997 3 I.W. Donald : Preparation, properties and chemistry of glass and glass-ceramic-to-metal seals and coating, Review paper, J. mat. Sci., 28, 1993, p. 2841-2886 4 Nike Lahl : Untersuchungen zur Chemischen Kompatibilitaet Zwischen Glasloten und den Komponenten der HochtemperaturBrennstoffzelle,Berichte des ForchungszentrumsJuelich, Institut h e r Werkstoffe und Verfabren der Energietechnik, 1999 5 H. Dorsing, R. Conradt, Aachen, und U. Diekmann, Jiilich : Alkali and phosphase free solder glasses with high thermal expansion coefficient and high glass transition temperature, 5* Int. Cod. On Joining, Aachen 1998. 6 G.-A. Heilemann, R. Conradt : Enimicklung und Herstellung von Hochdehnenden Glasloten und Laminaren Metall-Glas-Kompositen. Diplomarbeit, ,Lehrstuhl fiir Glas und
62
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I I I I
10.65 0.146
-
I
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I I
10.77 I 11.7
-
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-
4 50 loo 3.10-~ 10 2.10-~ : data from 150 1 h test in tube furnace, d : leak rate in 'C only; therefoi
keramische Verbundwerk-stoffe (R. Conradt), RWTH Aachen 1998 Schwickert, T.; Geasee, P.; Janke, A.; Conradt, R.; Diekmann, U. : Electrically Insulating HighTemperature Joints for Ferritic Chromium, Steel. Proc. of IBSC 2000, Albuquerque (NM), USA, p.116/122 Conference".P. Geasee, T. Schwickert, U. Diekmann and R. Conradt : MO-R203-Si02GlasKeramiklote fiir Hochtemperatur-Brennstoffiellen; 74 GlastechnischeTagung, 29-3 1Mai, 2000, Ulm, Germany. Wolf E. Matthes; Keramische Glasuren : Grunlagen, Eigenschaften, Rezepte, Anwendung., Augustus Verlag Augsburg 1990 10. J.F. MacDowell : Neucleation and crystallization of barium silicate glasses, Proc. Brit. Ceram. SOC.,No. 3,1965, p. 229-40 11. O.V. Mazurin, M.V. Streltsina and T.P. Shvaikoshvaikovskaya : handbook of glass data (Physical sciences data 15, Elsevier Science Publishers B.V., Amsterdam, Netherlands, 1993 12. Y.S. Touloukian : ThermophysicalProperties of High Temperature Solid Materials, Vol. 4, Macmillan, New York 1967 13. Engineering Materials Handbook, Vol. 4, USA; 1991
LIGHTWEIGHT AND WEAR RESISTANT CMC BRAKES W. Krenkel, R. Renz, B. Heidenreich German Aerospace Center (DLR), Institute of Structures and Design 70569 Stuttgart, Germany
Abstract Ceramic matrix composites offer great advantages for new lightweight and wear resistant brakes. C/C-Sic materials with additional surface coatings show in combination with modified pads high and stable coefficients of fiiction and extremely low wear rates, predestining them as lifetime brake disks e.g. for automotive vehicles, motor bikes and for weight-reduced bogies of new trains. To overcome the discrepancy between high material's costs and an acceptable price for the fmal CMC product, new design and manufacture concepts are objectives of the current research.
Introduction Within the national space research programme the DLR started in 1990 first tribological investigations with C/C-Sic composites, manufactured via the Liquid Silicon Infiltration Process (LSI), which have been originally developed as heat shield materials for new thermal protection systems of future space transporters [ 11. First attempts to adapt the mainly carbon fibres and silicon carbide containing material to tribological applications suceeded and showed the high potential of C/C-Sic as a new brake disk material [2]. In comparison to carbodcarbon (C/C) composites [3,4,5], which are the actual braking materials for aircraft and racing cars, the insufficient stability of the coefficient of friction caused by humidity and temperature could be improved essentially. As the substitution of conventional materials results in higher costs in general, CMC brakes will only be accepted, if an economic benefit for customers can be achieved beside all technical advantages like lower unsprung mass, shorter brake distance and higher thermal as well as corrosive stability. Therefore, high wear resistant ceramic brake disks are currently under development which allow an use over the whole life of the vehicle without any exchange or refurbishment. The very promising test results in combination with new developed pad materials lead to actual marketing campaigns of the automotive car producers (e.g. DaimlerChrysler,
Porsche), prognosticating new cars in the near future equipped with lifetime CMC brake disks.
Wear improvement of C/C-Sic brake disks CMC braking materials mainly consist of carbon fibres and matrices of silicon carbide, carbon and silicon. The carbon fibres improve strength, thermal shock resistance and damage tolerance and can vary in type and length, ranging from continuous to short fibres with dimensions of only some few millimeters. Braking against pads of the same material result in surface temperatures which considerably can exceed 1000 "C. This high thermal stability coincides with nearly identical wear rates in both components, i.e. brake disks as well as pads. For example, high energy brakings of 150 kJ lead to wear rates of approximately 170 mm3/MJ of the total brake system [6]. This wear rate may be acceptable in emergency braking systems but is essentially too high for service brakes as used in vehicles. To overcome this restriction, DLR developed a ceramic coating with improved wear resistance which can be deposited on the fiiction surface without any additional manufacturing step (Fig. 1).
CICSIC Substrate
Fig. 1 : SiCralee-coatings for 2D C/C-Sic brake disks
During siliconization, the last processing step of the C/C-Sic composite manufacture, additional silicon as well as carbon are added to the surfaces. By varying the type of carbon material, the amount of shares of the
63
two components and the processing conditions, a layer containing silicon and silicon carbide, permanently and strongly fixed with the C/C-Sic substrate has been achieved by the chemical reaction of carbon and silicon to Sic. The thickness of this SiSiC coating can be adapted to the requirements of the brake and amounts usually between 0.2 to 2 mm. To improve the surface finish and in order to achieve the brake disk’s end contour, the ceramic coating is ground with diamond tools in a concluding step. The coefficient of thermal expansion (CTE) for the ceramic coating is in the range of 3 to 4*104 1/K while the fibre ceramic substrates show normally lower values. Depending on this CTE mismatch, a more or less microcracked surface occurs as a result of the higher contraction of the SiSiC coating during cooling after processing (Fig. 2).
Fig. 2: Top view of the microcracked SiCralee-coating consisting of silicon (white) and silicon carbide (dark)
The most pronounced microcrack pattern in this socalled SiCralee-coating can be observed for twodimensional reinforced substrates whereas isotropic reinforcements lead to nearly crack-free surfaces. Additionally, the width of the microcracks on the outer surface depends on the thickness of the SiCraleecoating. In general, the thicker the coating, the wider the surface cracks. During braking, when the coating is heated up, the cracks get closer as the coating expands more than the substrate. Although microcracked, these SiCraleecoatings show an extremely good adhesion on the substrate’s surface even under thermal shock conditions. The cracks normally run through the total thickness of the coating, but stop at the surface of the C/C-Sic composite and no breakage of the fibres can be observed (Fig. 3). As the braking procedure is complex to simu-
64
Fig. 3: Cross section of a SiCralee-coated C/C-Sic laminate (coating thickness approx. 0.7 mm)
late, original sized brake disks have been fabricated to prove the feasibility of the SiCralee-coating under real tribological conditions. Therefore, brake disks with differing dimensions have been assessed through trials and service tests by the brake system manufacturers (Fig. 4). As a result, SiCralee-coatings increase the wear stability for pads and disks essentially. Using sintermetallic instead of C/C-Sic pads, the wear stability of the SiCralee coated C/C-Sic disk increased considerably. Only 2 mm’/MJ occurred while the pads lost 21 mm3/MJ [ 6 ] . As a consequence, these tribological tests demonstrate the high improvements in wear resistance which are achieved by the ceramic coating. Almost wear-flee brake disks in combination with acceptable wear rates for the sintermetallic pads offer a high potential for lifetime brake disks.
Fig. 4: Automotive brake disks with SiCralee-coatings (diameter from 280-320 mm, mass 1 to 3 kg)
Increase of the transverse thermal conductivity Tribological tests showed the influence of the surface temperature on the tribological behaviour [7]. Extremely high temperatures on the friction surface reduce the coefficient of friction (COF) and increase the wear. One efficient method to limit the temperature and to protect the brake's periphery from high heat radiation is the increase of the C/C-Sic material's transverse thermal conductivity. This can be realized exemplarily by the 0 0
0
use of high heat conductive carbon fibres increase of the angle between the fibres and the friction surface increase of the ceramic content in the C/C-Sic material
While the technical and economical efforts grow considerably for the first two methods, resulting in higher costs for the ceramic brake material, the increase of the silicon and the silicon carbide content within the composite is a more cost-efficient way. Higher ceramic contents can be achieved easily by reducing the fibre content resulting in a higher silicon uptake and S i c formation and consequently a higher density of the C/C-Sic composite material. Therefore, the lower the fibre volume content and the higher the density, the higher the transverse thermal conductivity of the CMC material. Nevertheless, higher densities coincide with a decrease in strength and fiacture toughness. The resulting CMC composites have therefore to be designed according to the individual requirements of the brake system as a compromise between sufficient thermal and mechanical properties. The following figures demonstrate these relationships for some representative C/C-Sic materials. Figure 5 shows the transverse thermal conductivity for short fibre reinforced C/C-Sic composites in dependance on the fibre volume content in the green body stage (carbon fibre reinforced plastic, CFRP). The values vary fiom 23 to 29 W/mK for fibre contents of 55 to 30 Vol %. The tests have been conducted at low temperature (50 "C) with samples of 19 mm in thickness and 25 mm in diameter. All samples have been cut out of plates which were manufactured by hot pressing different compounds of short carbon fibres and polymer precursors. As for most other materials, the thermal conductivity of C/C-Sic decreases with higher temperatures. Figure 6 demonstrates this relationship
30 28 26 24 22
20 20 30 40 50 60 fibre Content (CFRP Stage) [Vol.% ] Fig. 5: Thermal conductivity at 50 "C versus fibre volume content of short fibre reinforced C/C-Sic composites
for different material modifications which have been manufactured by varying the fibre architecture and the processing conditions. The average decrease of transverse conductivity between 300 "C and 900 "C, which represents the range oft high performance brakings, amounts to approximately 30 %. As a consequence, suitable C/C-Sic materials for brake disks must be selected by their high temperature behaviour. However, in contrast to most other materials, C/C-Sic composites remain their mechanical strength level from room temperature up to more than 1000 "C.
35
30
25
20
15
100 250
400
550
700
850 lo00
Temperature ["C] Fig. 6: Thermal conductivity of different C/C-Sic composites versus temperature
65
Figure 7 summarizes the correlationship between the material's density and the transverse thermal conductivity of C/C-Sic. Generally, the lower values correspond with continuous fibre reinforcements (fabrics), whereas the highest densities have been measured particularly with short fibre architectures.
JJ
30 25 20 15 10
5 0 1,so
2,oo
2,20
2940
Density [g/cm'] Fig. 7: Thermal conductivity (at 50 "C) as a hnction of the density of C/C-SIC composites
In tribological tests, C/C-Sic qualities with different thermal conductivities were tested to compare the coefficient of friction. Basis was a standard C/C-Sic composite (C/C-Sic Type I), which was manufactured by stacking up woven fabrics of carbon fibres in the axial direction to a thickness of 8 mm, resulting in high mechanical properties in the radial as well as circumferential direction and in a very regular temperature distribution on the friction surface. Nevertheless, the 0/90° fibre orientation leads to an orthotropic behaviour of the disks, i.e. the thermal conductivity perpendicular to the fricton surface is rather low. To increase the transverse thermal conductivity, improved C/C-Sic qualities have been manufactured: Type 11: Type 111: Type IV:
Orthotropic C/C-Sic with high heat-conductive carbon fibres C/C-Sic with a relevant fibre content in axial direction Orthotropic C/C-Sic with adapted microstructure and higher Sic content
A set of one rotating brake disk pressed against two stationary disks of the same C/C-Sic quality with an outer diameter of 110 mm was tested in a high energy flywheel mass test facility under real conditions [8]. As shown in Figure 8 the coefficient of friction of the C/C-Sic Type I increases with decreasing velocity, reaching the highest frictional values of 0.5 to 0.6 just before braking ends. The
66
braking begins at a sliding velocity of 15 m/s and ends with standstill. High contact temperatures caused by the high energy input are responsible for lower COF especially at the beginning of braking. Due to the orthotropic structure and the low transverse thermal conductivity h of about 8 W/mK, the fiction surface overheat and the coefficient of fiiction decreases to an unacceptable low level of about 0.2 at W, = 145 kl. The decrease of the COF is explained by the creation of friction layers at high temperatures with low COF. Different chemical reactions of the C/C-Sic components (Sic, C and Si) and the sintering of wear particles are responsible for these layers. To obtain a more stable fiiction behaviour it is necessary to reduce the surface temperatures by increasing the transverse heat flux inside the braking disks to use the total heat capacity. Figure 8 compares the high energy friction behaviour of Type I1 to IV composites with the standard C/C-Sic at n = 3000 l/min and WR= 145 kJ. Carbon fibres show their highest values parallel to the fibres direction. If there are relevant proportions of fibres in direction of the brake disks axis the transverse heat flux can be increased essentially. As a result a more stable fiction behaviour can be observed over the total range of velocity for this C/C-Sic quality Type 111. The relation of the maximum and minimum COF is 1.6 compared to the relation of Type I of 3.3. The most economical way to adapt the composite material is to modify the microstructure itself with a simultaneous increase of the Sic content in the matrix (Type IV). The highest stability, however, was observed with C/C-Sic brake disks Type11 of high thermal conductivity achieved by using high heatconductive fibres. The COF remains at high sliding velocities at 0.5 with an increase to a maximum of 0.8 at the end of braking. The relation of hmax to bin is comparable to the composite material Type 111. In summary, these screening tests demonstrate the possibilities of improvements in COF by a material modification.
Type I11 -----Type IV
0
5
10
Awrage sliding speed
15 Ids]
Fig. 8: Coefficient of friction of different C/C-SIC qualities (n = 3000 Vmin, WR= 145 kJ, p = 0,34MPa)
Low cost manufacture of C/C-Sic brake disks The main drawback of C/C-Sic brake disks lies in their current costs which are at least one order of magnitude higher than for grey cast iron disks. In order to reach the optimistic goals of the car manufacturers to introduce CMC brake disks in the near future, cost efficient manufacture routes have to be developed. DLR investigates different designs for solid as well as internally ventilated disks taken benefit of the main advantages of the LSI process [9]: 0
0
0
Applying press techniques for the manufacture of the CFRP body Pre-shaping in the green or carbon material's stage (Fig. 9) In situ joining of elements in the carbodcarbon stage by reaction bonding [ 101 Implementation of SiCralee-coatings into the siliconization step [ 111
Different prototypes have been manufactured and tested by the respective brake system manufacturer or end-user (Fig. 10). The near future will show whether CMC brakes can fit not only the technical but also the economical requirements of such a series product.
CwlrbonfCarbon
CfC-SIC
Fig. 9: Manufacture of internally ventilated C/C-Sic brakes by prefabricated carbodcarbon half-disks [ 121
Fig 10: Prototype of an internally ventilated brake disk of short fibre reinforced C/C-SIC made by press technique
Conclusions C/C-Sic composites demonstrated their feasibility for high performance and lightweight brake systems. SiSiC-coatings (SiCralee) can improve the wear stability of C/C-Sic brake disks essentially, promising lifetime disks in future automotive applications. High transverse thermal conductivities increase the stability of the coeficient of fiiction up to 0.8. Different modifications of the material's composition have been investigated and compared with respect to their tribological behaviour. Using short fibre and applying hot press techniques for shaping CFRP disks promise a cost efficient manufacture route for a series production. Future research work will additionally be necessary to investigate quality assurance methods for these safety-related components to ensure reproducibility as well as reliability of the process.
References Krenkel, W.: Fiber Ceramics for Reentry Vehicle Hot Structures. 40th Int. Astronautical Congress (IAF), Malaga/Spain, 7-13 October 1989 Krenkel, W.: CMC Materials for High Performance Brakes. Proceedings ISATA Conference on Supercars, Aachen, 31 0ct.-4 Nov. 1994 Biihl, H.; Morgenthaler, K.D.; Naumann, E.: CFC-Bauteile im Automobilbau. Faserverbundwerkstoffe, Springer-Verlag, 1996 Sonn, H.W.; Kim, C.G.; Hong, C.S.; Yoon, B.I.: Transient Thermoelastic Analysis of Composite Brake Disks. Journal of Reinforced Plastics and Composites, Vol 14, Dec. 1995 Trefilov, V.I.: Ceramic and Carbon- Matrix Composites. Chapman & Hall, 1995 Krenkel, W.; Renz, R.; Henke,T.: Ultralight and Wear Resistant Ceramic Brakes. EUROMAT 99, Int. Congress on Advanced Materials and Processes, Munchen, 27.-30. September 1999 Heidenreich, B.; Krenkel, W.: Development of C/C-Sic Materials for Friction Applications. EUROMAT 99, Int. Congress on Advanced Materials and Processes, Munchen, 27.-30. September 1999 Renz, R.; Krenkel, W.: C/C-Sic Composites for High Performance Emergency Brake Systems. 9* European Conference on Composite Materials (ECCM-9), Brighton, UK, June 4-7,2000 German patent DE 197 49 462, filed 1997 (10) German patent DE 196 36 223, filed 1996 (11) German patent DE 198 34 018, filed 1998 (12) German patent DE 44 38 455, filed 1994
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DEVELOPMENT OF CERAMIC SHEATHED TYPE THERMOCOUPLE WITH HIGH HEAT RESISTANCE AND HIGH DURABILITY Hideki Kita, Takayuki Suzuki, Tetsuya Isshiki, and Hideo Kawamura Isuzu Ceramics Research Institute Co., Ltd. Kanagawa, Japan ABSTRACT So far, the disposable type of thermocouple has been world-widely used for the temperature measurement of molten metals such as cast iron. Considering the cost performance and resources saving, new type of ceramic sheathed thermocouple has been successfully developed. Tungsten-rhenium alloy wire which has high electromotive force, good linearity and high melting point, was used as sensor for temperature measuring and it was sealed by special ceramics in silicon nitride pipe with ceramic film. In addition, ceramic sleeve with layer structure was arranged around the silicon nitride pipe. It was confirmed that newly developed thermocouple withstood the severe conditions when dipped into molten metal and showed excellent properties about both response and durability.
BACKGROUND AND OBJECTIVES In casting process using molten metal, to measure the temperature of molten metal correctly is the most basic and important item for controlling the quality of the casting. SO far, as the tool for the purpose, various kinds of thermocouple have been used according to the temperature and the nature of the molten metal. Direct temperature measurements on high temperature and strong corrosive molten metal such as cast iron (operation temperature:1723-1773K) and copper (14731573K) are much more difficult than the case of molten aluminum(923-973K). As for molten metals such as the cast iron, copper, frequently intermittent temperature measurement has been carried out using thermocouple in the series of casting processes. In these cases, the thermocouple must withstand the repeated thermal shock and the corrosion when dipping it into high temperature molten metals at a breath without preheating. Also, YOU can imagine easily how severe the work near red-hot furnace is for the worker in the factory. The request to finish temperature measurement work in short time is so strong, and for its purpose, about six seconds are required as the response of the thermocouple. In this way, high response, durability and cost performance are required to the thermocouple for cast iron, however, the hurdle of 4 4 the technology development to meet all of the items is so high, then in the past, such developments were supposed to be given up. Then, a thermocouple having the structure making Pt-Rh wire expose from the tip of the paper pipe, which made durability sacrifice, has been widely used'). The thermocouple is excellent about the response because the wire is exposed, but it is oneso-ca11ed type that needs exchange every time after one to several times uses. Therefore, the
number of the thermocouple consumed in the production factories is proportional to the measurement number of times and becomes an enormous number. As a result, the management expenses which is related with the physical distribution, the conveyance, the disposal in addition to the cost of thermocouple itself became expensive. Moreover, it left a problem from the resources saving viewpoint of consuming a paper pipe, rare metal with enormous quantity in addition to the problem of the displayed temperature is different even if the temperature measurements is done on the same molten metals, due to the differences in performance among the pieces. Then the objective of this study is, considering a situation and a request in the market, to develop the thermocouple which is excellent about the response, the durability, and the cost performance, using ceramics.
EXPERIMENTAL PROCEDURES Concepts of the new thermocou~le Of the new thermocouP1e are The summarized below, corresponding to the market d~mands. ( l ) High Precision (2) Quick response (3) Long durability (4) Little adhesion of molten metal (5)Low cost (6)Resources saving
Basic design and the structural characteristic of the new thermocouple According to the concepts mentioned above, basic design and structures of new thermocouple were discussed, and the outline of the them are mentioned hereinafter. The structure schematic of the thermocouple tip is shown in figure 1.
Fig. 1. The structure schematic of the thermocouple
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The most distinctive characteristic of this thermocouple is that the use of tungsten-rhenium(W-Re) alloy wire and its sealed structure in the silicon nitride ceramic protective pipe. That is, as the sensor for temperature, W-5Rem-26Re wire was used because of its high thermoelectro motive force compared with the conventional platinum-rhodium( Pt-Rh), excellent about linearity and its highest melting point in metal. However, W-Re wire has the fault that it is easily oxidized in the hot air and also that it becomes fragile and easy to be damaged when it is exposed over the temperature of 1373K because of the recrystallization. In order to make use of it in the high temperature oxidation atmosphere possible, it is necessary to make the structure in the way that air doesn't invade into the silicon nitride protective pipe. Then, the space between the W-Re wire and the protective pipe was filled by Ti added reaction sintered silicon nitride(RBSN) which has small expanding characteristic during sintering2) or silicon nitride added glass composite. The oxidation prevention of wire and the support fixation became possible by making the such structure. For the protective pipe, fully densified silicon nitride ceramics was used because of its superior thermal shock resistance, high mechanical strength and fracture toughness and having thermal expansion coefficient almost equal to that of W-Re. However, it was proved that silicon nitride pipe was eroded by the molten metal as a result of a series of basis examinations. Therefore, an effective film for improving the erosion resistance was formed on the surface of protective pipe tip. For the temperature measurement of molten cast iron, Mo/ZrB2 based cermet film was formed with plasma splay process on the surface, and for molten copper, MgO based film was found to be effective for improving the durability. Around the silicon nitride pipe, outer sleeve is arranged. The sleeve is composed of the layer material containing boron nitride(BN), which is difficult for molten metal adhere to. It was supposed that delamination occurred not but immediate destruction in layer materials. On the other hand, for speeding-up the response, assembly was designed in the way that heat would be concentrated on the tip of the thermocouple and it would not flow toward the rear part as much as possible. The reduction of the diameter of protective pipe, and forming the air gap between outer sleeve and the pipe are also effective for the reduction of heat capacity, and heat radiation from the side and to rear part. Moreover, no aerial stagnation in the tip inside in addition to bigger thermo-electromotive force of W-Re wire were also effective to improve the response.
Evaluations The ceramic thermocouple mentioned above was connected with a stainless pipe and a display part and it completed assembly. For testing, cast iron, copper, aluminum molten metal were used. Each metal blocks were fused in the furnace by high frequency induction heating. The temperature were about 1723-1773K, 1473-1573K, 923-973K respectively. After the metal in the furnace reached fixed temperature, thermocouple was dipped in the molten metal, and response, durability,
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precision were evaluated on the conventional thermocouple and the newly developed one, After examination, several kinds of investigation were carried on the sample such as X- ray radiography , SEM observation and EDS analysis.
EXPERIMENTAL RESULTS DISCUSSIONS Results on molten cast iron
AND
The appearance of the thermocouple assembly is shown in figure 2. It is composed of the stainless pipe and the display part which were connected with the ceramic sensor part. The full length is about 1600 millimeters with the standard specification article.
Fig. 2. The appearance of the thermocouple assembly (top and rear view) The result which was examined on molten cast iron of about 1750K using the thermo couple assembly is described below. First, The comparison of the response when measuring the molten cast iron using the conventional thermocouple (disposable type) and the newly developed one, is shown in the figure 3.
I
t
I 0
.+ 0+
+developed TC
1
conventiod TC (disposal type)
1
2
4
6
I 1
8
time (S)
Fig. 3. The comparison of the response As is shown in fig.3, the value stood up rapidly in the early stages, became stable and finally flat with the time. Here, response was defined as the time until it becomes flat after beginning the temperature measurement. As a result, when using a developed thermocouple the
response became about 6 seconds in the case where conventional one shows 5 seconds. The response of the developed one is comparatively quick when it thinks the wire being arranged inside the protective pipe whereas that the tip of the wire is exposed for the disposal type, Also, when using monolithic silicon nitride ceramics or a commercially available molybdenudzirconium oxide(MoIZr0,) cermet, which has been used as protection pipe for molten iron, as the material of the above-mentioned sleeve, the solidified slag, iron and so on became massive and adhered to the tip. And also the damaging by the thermal shock occurred in short time.
I
I
Fig. 6. The appearance of the tip part after 300 times.
Fig. 4.The microstructure of the layer material for the outer sleeve Then, as the solution of such problems it was decided to be composed of the silicon nitride based layer material containing BN, which is difficult for molten metal adhere to. The microstructure of the layer material which becomes a sleeve is shown in figure 4, and the comparison of durability between monolithic and layer structure ceramics are shown in figure 5. Also, the situation of the tip part after 300 times use was shown in figure 6. As is shown in figure 5 and 6, the adhesion quantity of the iron was few, and the delamination among the layers progressed gradually, but that there immediately was not broken, Moreover, the amount of the adhesive slag or iron were reduced.
The temperature measurement result when changing a molten metal temperature, using an conventional thermocouple (the disposable type) and a developed one is shown in figure 7. It should be noted that when using a developed thermocouple, the line becomes smooth compared with the case of using a disposable type thermocouple. This data suggests that the scattering of the temperature measurement result becomes small. After implementing temperature measurement of the repeat, the lifetime of newly developed thermocouple were more than 7OOtimes in the maximum, it became about 400 times in average.
1650
I 0
50
100
cycles (times) Fig. 7. The temperature measurement results 600 ;
Results on molten copper
, 0
----
Monorithic material (cermet)
Layer material (SN/SN * BN)
Fig. 5 . Durability of monolithic and layer structure ceramics
As for molten copper, because the speed of solidification phase boundary is so high compared with that of cast iron, the temperature control with higher precision is required. The thermocouple with high durability is effective for the improvement of the precision in measurement. In case of examination using the silicon nitride protective pipe without a film, W-Re wire inside was broken and became that it didn't work at less than 100 times. According to investigation by Xray radiography, It was proved that a small hole was formed by erosion at the tip of the protective pipe and then molten copper came in from the hole and reached the wire. On the other hand, it was confirmed that when using the silicon nitride protective pipe with magnesium oxide film, the response was about eight seconds and 71
after 5OOtimes tests, the thickness was approximately same as before use and it showed a normal measured value. According to the SEM observation and the EDS analysis, it was found that the erosion of silicon nitride protective pipe by molten copper was prevented in the tens-of- p m magnesia layer which exists in the surface.
Results on molten aluminum The operation temperature of molten Aluminum is less than 973K and this temperature is much lower than that of molten copper and the cast iron. When doing temperature measurement repeatedly, it proved that it endured repeat use with equal to or more than 4OOO times in molten aluminum. And the response was almost l0seconds. As for thermocouple for molten aluminum , conventional stainless sheath type would become a competition with the developed one, but newly developed thermocouple is excellent about the response and durability. Moreover, the there is a merit that process of the powder spray before use becomes unnecessary because molten aluminum are hard to adhere to silicon nitride pipe.
CONCLUSION Tungsten-rhenium alloy wire which has high thermoelectro motive force, good linearity and high melting point, was used as sensor for temperature measuring and it was sealed by ceramics in silicon nitride pipe with film on the surface, and outer sleeve having layer structure was arranged around the silicon nitride pipe. On thus obtained thermocouple assembly, its performance was investigated.Ql For molten cast iron : it shows 6 seconds of response, 400 times of durability. @ For the molten copper, it shows 8 seconds of response, and more than 500 times of durability. 0 For the molten aluminum, it shows 10 seconds of response, and more than 4000 times of durability.
REFERENCES 1) K.Ogawa, S. Suzuki, Japanese Patent N0.05-034033 2)H.Kita, Dimensional change chracteristics of ceramics produced by simultaneous nitriding of Si and Ti powders and its application to adiabatic engine parts,lO0,4,PP.488-493(1992) Japan journal of ceramic society
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ADVANCES IN HOT GAS FILTRATION TECHNIQUE R. Westerheide*, J. Adler** (*)Fraunhofer-Institutfur Werkstoffmechanik, D-79108Freiburg im Breisgau, Germany (**)Fraunhofer-Institutfur Keramische Technologien und Sinterwerkstoffe, D-01277Dresden, Germany
ABSTRACT The efficiency of coal fired power plants is mainly a function of the firing process. However the process parameters have to meet the requirements of the raw gas filtration unit, when a gas turbine is used in combined cycle plants. Combined cycle plants are offering an increase of efficiency of about 4% by using new technologies compared to conventional plants. These technologies requiring at least higher operating temperatures. Therefore new materials for higher filtration temperatures in both oxidizing and reducing atmospheres are necessary. Further the lifetime of the filter elements regarding the increased temperature and regarding the element design has to be guaranteed up to 16OOO h operating time. Usually the filter elements are hollow candles made of siliceous bonded silicon carbide with an overall length of 1500 mm. However, the binder is often subjected to creep due to the softening of the binder at high temperatures and the candle may be damaged due to vibrations in the filter unit.
pressure drop in order to reach a maximum operation efficiency. These requirements are met by ceramic materials. Rigid ceramic filter elements are a leading technology in the removal of particles from gases, however their application has been seen to be highly dependent on the properties of the dust being separated from the raw gas [ 13. The structure of ceramic filter elements consists often of a highly porous support which ensures the mechanical strength and a layer which operates as the functional part for the particle removal (Fig. 1). The durability of the ceramic filter elements is limited in principal by chemical reactions between gas and ceramic material, the mechanical loading or thermal shock during pulse cleaning. Thereby the lifetime of the whole element is mainly determined by the lifetime of the support material and not by the functional layer.
The paper gives an overview about the work in developing new ceramic filter materials based on silicon carbide and alumina. The goal of the development was to increase the potential application temperature. It was found that the application temperature can be increased up to 900 "C due to an optimization of binder and sintering process. The analysis of stress distribution in the filter element has shown that the material damage can be reduced by slight optimization of the geometry.
INTRODUCTION The combined cycle technology offers the prospect of higher generation efficiency than the only steam cycle based technology. One of the central components in the combined cycle technology is the filtration unit in which the removal of particles from the feed of the gas turbine takes place. For a long term stable operation of the filter elements, their resistance against thermal and mechanical loading and aganist chemical attack has to be very good. Additionally, the filter media must ensure a minimum
Fig. 1: Structure of a ceramic surface filter element, Dia Schumalith" made by USF-Schumacher.
In this work different ceramic support materials with an increased resistance to high temperature creep deformation and increased strength were developed and investigated. The materials are oxide and non-oxide based ceramics. On the one hand they were produced as a conventional filter medium made of bonded silicon carbide or sintered oxides and on the other hand they were produced as self bonded silicon carbide by recrystallization or granulation of fine grained silicon carbide powder. As an alternative route for the production of rigid filter media, the formation of multilayered ceramic foam was investigated.
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MATERIALS DEVELOPMENT SILICON CARBIDE FILTER ELEMENTS WITH A SILICEOUS BINDER Most interest in the past was focused on silicon carbide filter elements bonded with siliceous material. The production of these filter elements is relatively cheap compared to self bonded silicon carbide due to the lower sintering temperatures and therefore the use of cheaper furnaces. An effective method is to add a claybased binder phase, e.g. alumindmullite which form glassy phases on firing and act as the binder. However, during long term exposure the so called clay bonded ceramics are subject to bending and creep, in which the candles have shown irreversible elongation. Due to the elongation the structure can be damaged and the element can fail during operation. During the last years new, high-temperature' SiCbased candles with improved resistance to creep were developed and extensively tested in demonstration plants [2] and in the laboratory scale [3] up to 870°C. The resistance against creep was found to be good; no element failed during operation. However, to meet the requirements of further increase of filtration temperature, the development of new materials was necessary. The work was focused on the optimization of the binder, the silicon carbide powder and the whole production process, i.e. the powder granulation, the powder compaction and the sintering. As the influence of the silicon carbide powder was found to be negligible in the range of the considered Sic-powders, the mechanical properties are strongly influenced by the kind and amount of binder. Fig. 2 shows the change of strength due to the variation of binder without changing the sintering process. The strength was measured at 4-point bending specimens at room temperature and at 900°C. All specimen have nearly the same porosity. An increase of the amount of binder leads to an increase of strength as it could be measured for an increase amount of flux. However a higher amount of binder has no positive effect on the deformation resistance at high temperatures, as it is shown for a constant load of 5MPa and temperatures from 850°C up to 1OOO"C (Fig. 3). Therefore compromises between strength and high temperature deformation behavior have to be found. After the optimization of the binder the variations of powder processing, powder compaction and sintering have shown that a further potential for the improvement of properties exists. Especially the creep resistance could be improved due to controlled crystallization of the binder during the sintering process (Fig. 4). Figure 4 shows the deformation behavior of the o p timized clay bonded silicon carbide filter material which is the basis for the newly developed Dia Schumalith"-N (USF-Schumacher) material. This material has a high
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strength at a similar porosity in comparison to the commercial available material Dia Schuma1ith"-T. The high strength is mainly a result of an optimal transition between silicon carbide grains and binder. -0- Strenglh at 20°C
p*
z:
- &Strength
at 900°C
- Binder: 30 - increaseof binder Binder: Alumina+clay constant flux increase of binder *
20-
Alumina4ay constant binder increase of flux
Fig. 2: Strength of different S i c materials with siliceous binder measured by 4-point bending at mom temperature and at 900°C. 150
lo00
p .-c
-f
900
125
800
'09s
E,
L
Q)
700
75
0
n
s 8600
504
500 25
400 0
0
5
10 t i e in h
15
20
Fig. 3: Deformation behavior of silicon carbide filter materials with two binder systems and different amounts of binder. L
0.15
Creep strain in &point-bending. All materialscontain the same SiCpowder and bioder.
- Firstoptimisationof sintering. Test at m.
Consequently the optimization of clay bonded silicon carbide filter elements resulted in a further improvement of mechanical strength, high temperature creep resistance maintaining the strong adherence of the ceramic membrane coating, the good cleanability and the high filtration efficiency as it is known for the earlier developed materials Dia Schumalith" F-40, Dia Schumalith" 10-20 and Dia Schumalith@-T(details are described in [4,5]).
SELF BONDED SILICON CARBIDE FILTER ELEMENTS The disadvantages of clay bonded silicon carbide filter elements due to the softening of the binder and resulting high temperature deformation can be avoided by self bonding of silicon carbide. An alternative route to produce a silicon carbide material without significant amount of siliceous secondary phases is the reaction bonding of silicon carbide (RSiC). This method is well known for less porous structures (15%) used for the setup in furnaces for sanitary and tableware. Also there were attempts in the US for using structures with higher porosity for the hot gas filtration technique. The results of intensive materials development and prototype filter tests has shown promising results [6]. The RSiC-route was applied in this project by the AnnaWerk GmbH. The bonding of silicon carbide results from the formation of &SIC during high temperature firing at 2000°C. Also for this material the drawback is the necessity of higher sintering temperatures and therefore a more expensive production route, however here the resistance against corrosive attacks seems to be very good and there is hardly any creep deformation measurable. The development of RSiC-filter elements resulted in the production of 1.5 m candles based on AnnaNox@CK225. The strength of the RSiC with a porosity between 25% and 30% is in the range of the strength measured for Dia
[email protected] this kind of silicon carbide filter elements a new membrane coating was developed on the basis of liquid phase sintered silicon carbide. At the moment the RSiC-candles undergo tests in demonstration plants to study the performance of this material. In the laboratory scale no material degradation due to oxidation, corrosion or any thermal or mechanical fatigue could be detected.
tion and the higher sintering temperatures. However it is expected that the creep resistance and the resistance against corrosive attack is improved, so that the much higher cost for production will in some cases be warranted.
OXIDE BASED FILTER ELEMENTS Oxide based filter elements should retain their stability in both oxidizing and reducing gases, providing that the free silica content is low, or not available for reaction. They are already in their most stable state of oxidation, and less likely undergo further phase transitions [7]. The use of filter materials based on spinel and corundum can led to a better adaption to specific filtration atmospheres in different power plant processes. The development of oxide based porous materials has led to strength values which are similar to those of the optimized clay bonded silicon carbide filter elements. In contrast to the clay bonded S i c the resistance against creep deformation was found to be excellent up to 1000°C. For the corundum material a pressure drop of 27mbar at 7cmls was measured. The successful development of oxide based filter materials resulted in tests at a hot gas filtration demonstration plant. After 1250h at 760°C under combustion conditions a corundum material has shown nearly no strength loss from a level of 24 MPa (measured by O-ring testing). However, the spinel materials have shown high strength loss. Details to the material production and to the differences between spinel and corundum are described in [6].
ALTERNATIVE CONCEPTS FOR CERAMIC FILTERS Most of the filter elements for the hot gas filtration have been formulated from rigid granular ceramic structures. Other concepts were focused on fibrous media or in whisker reinforcement [8]. The main advantage of such elements is their low density, however, their mechanical strength is generally lower as it can be observed for porous metallic structures. High temperature experience with fibrous candles is limited. Tests performed in the Tidd demonstration facility have shown an embrittlement and oxidation of the fibers [2].
Fig. 5: Necks between granulated sub-micron silicon carbide forming the open porosity for the filter support.
Another possibility of self bonding is the conventional sintering process of granulated sub-micron silicon carbide powder. This production method resulted in a very homogeneous structure (Fig. 5) with a similar porosity as it was measured for the clay bonded silicon carbide filter elements. The drawbacks of this production route are the cost of powder, the effort for granula-
Fig. 6: Silicon carbide foam structure with a 6Oppi support and a 6OOppi layer.
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Another possibility for the formation of low density filter materials is the production of ceramic net like structures made from polymeric foams. These foams are limited in the variability and controllability of the cell geometry as well as the strut shape, but it is possible to produce either isotropic or anisotropic structures with a wide variation of pore sizes. For the application as hot gas filter element layered structures of sintered silicon carbide as shown in Figure 6 were produced in the shape of plates. The structure is similar to those of coated granular silicon carbide candles consisting of a support material with 60 pores per inch (ppi) and a layer with 6OOppi. However, the production of the common long hollow cylindrical candles is not reasonable at the moment. Rather it is more suggestive to think about other geometrical structures for the application of foam structures as filter elements.
3:30
ANALYSIS OF GEOMETRY The geometry of a small hollow filter element with a wall thickness of about lOmm, an outer diameter of 60mm and an overall length of about 15oOmm with a porous support structure is not typical for ceramic materials. Therefore the risk of mechanical fatigue has to be investigated. A mechanical fatigue can be at first originated by the dead weight of the element and the additional thermal loading. This can be investigated by creep experiments to get the values for the high temperature deformation resistance. At second a mechanical fatigue can be originated by vibrations of the filter systems or vibrations of plant components which are connected to the filter systems. This has to be investigated by an analysis of the natural frequencies of the filter elements and the following local stresses. Also the results of the analysis have to be compared to results from experimental investigations at the filter element.
.
Fig. 7: Natural frequenciesof a ,standard"filter element.
Different possibilities to reduce the stress and to modify the stress distribution are existing. First of all one have to think of alternations of the material. Taking into account that a changing in materials properties are counterproductive for the resistance against high temperature deformation or the filtration efficiency one has to think only of changes in the geometry of the elements.
It can be assumed that the vibrations in the filter systems are in the region of 1 to 200Hz as the maximum value. The calculation for the above defined standard element has resulted in a first natural frequency of 19 Hz and a second natural frequency of 119 Hz in consideration of basic values for the density (1.9 g/cm3). the porosity (45 %), the Young$ modulus (34GPa) and the possions ratio (0.2). In regard to the order of natural frequencies shown in Figure 7 the local stresses are maximum at the top of the elements where the element is fixed at the mounting plate (see Figure 8). Thus it can be assumed that either during start-stop cycles of the plant or during operation the vibrations in the filter system are of the magnitude of the natural frequencies. Therefore high local stresses will be generated.
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Fig. 8: Stress distribution at the top of a filter element.
The simulation of other geometries has shown that a reduction of length leads to a significant shifting of the natural frequencies but the stress level is sometimes higher than for standard elements. A variation of the inner and outer diameter to acceptable values leads only to slight changes of natural frequencies but resulted in significant reductions of stresses. The results of further calculations for different geometries were that a change to conical elements or an adaptation of additional weights are promising favorable stress levels and stress distribution. This can lead to a reduction of the risk of mechanical fatigue. The calculations have to be verified by experiments. This has started in the middle of 2000 and will be published later.
CONCLUSION It has been shown that other structures than the common clay bonded hollow cylindrical candles are applicable for the particle removal from hot gas streams. For example oxide based candles, RSiC candles and materials with granulated structures made from sub-micron silicon carbide are offering new potentials due to the findings of high strength, high creep resistance and good oxidation resistance obtained in the laboratory scale. These materials are still under investigation. For clay bonded silicon carbide candles significant improvements in strength and creep resistance maintaining the good properties of Dia Schumalith@T, i.e. a strong adherence of the ceramic membrane coating, a good cleanability and a high filtration efficiency, were achieved. The new material Dia Schumalith@N will be extensively characterized in different demonstration plants soon.
REFERENCES S.K. Grannell et al.: Investigation into the Behaviour of Particle Compacts and Comparison with Industrial Experience. High Temperature Gas Cleaning, ed. by E. Schmidt et al., Karlsruhe 1996, 145156. M.A. Alvin: Performance and Stability of Porous Ceramic Candle Filter During PFBC Operation, Materials at High Temperature, Vol. 14, 3 1997, p. 355-364. P.J. Pastila et al.: Effect of Crystallization on Creep of Clay Bonded Sic-Filters. 22" Annual Cocoa Beach Conference on Composites, Advanced Ceramics, Materials and Structures, Jan. 1998, Fla, USA. To be published in Cer. Eng. and Sc. Proc. of the Am. Cer. Soc.. K. Schulz et al.: Improved Ceramic Filter Candles for Hot Gas Cleaning. High Temperature Gas Cleaning, ed. by E. Schmidt et al., Karlsruhe 1996, 843-845. R. Westerheide et al.: Filter aus Hochleistungskeramik fiir die Energietechnik.Final Report, Fraunhofer Institut fiir Werkstoffmechanik,W2199, Freiburg 1999. P. Eggerstedt, J. Zievers: High Temperature Particualte Removal Using Recrystallized Silicon Carbide Candle Filters. High Temperature Gas Cleaning Vol. II, ed. by A. Dittler et al., Karlsruhe 1999, 375-383. J.P.K. Seville et al.: Recent Advances in Particulate Removal from Hot Process Gases. High Temperature Gas Cleaning Vol. I, ed. by E. Schmidt et al., Karlsruhe 1996.3-25. Y.Sawada et al.: Evaluation on Fundamental Properties of Filter Materials at High Temperature. High Temperature Gas Cleaning Vol. 11, ed. by A. Dittler et al., Karlsruhe 1999,393-404.
The analysis of geometry has shown that different possibilities to reduce the risk of mechanical fatigue due to vibrations in the filter systems are existing. The variations are possible without changes of materials properties and filtration efficiency. The calculated results have to be verified in experimental work in the near future.
ACKNOWLEDGEMENT The authors gratefully acknowledges the Bundesministerium fiir Bildung und Forschung (BMBF) for financial support under contract No. 0326832 and under contract No. 0326866F. In addition we would like to thank the companies AnnaWerk I Rdental, USFSchumacher I Crailsheim, ESK-SIC GmbH I Grefrath and H.C. Starck I Laufenburg for their technical and financial support.
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NANO-SCALED CERAMIC MEMBRANES FOR THE FILTRATION OF FINE AND STICKY DUST B. Zimmerlin*, I. Miillner, H. Leibold, H. Seifert Forschungszentrum Karlsruhe GmbH, Institut f i r Technische Chemie, Bereich Thermische Abfallbehandlung, D-76021 Karlsruhe, Germany
ABSTRACT Hot gas filtration by using ceramic filters at temperatures up to 1000 "C is of crucial importance as key technology for numerous high-temperature processes. However, filter clogging by h e and sticky dusts as well as material degradation still remain problems, limiting the long term stable operation. In order to avoid filter clogging, multilayer grain ceramics were coated by a homogeneous and highly porous filter membrane consisting of nano-scaled A1203-particles.Filtration and recleaning tests with a fine and sticky model-dust revealed an excellent filtration and recleaning performance of the produced membranes, even at a temperature of 780 "C.
INTRODUCTION In advanced power generation systems fine dust separation at high temperatures is essential to protect down-stream components as gas turbines or heat exchangers. The integration of a hot gas filtration unit also offers an immensely optimization capacity for product recovery as well as for various thermal waste treatment processes. In the latter, thus enables among others the fiuther utilization of the product gases of pyrolysis or gasification of bio mass and waste. Furthermore, it improves heat recovery methods and reduces corrosion damages. Due to the high temperatures and chemical aggressive conditions, only ceramic filter media are suitable for this purpose. In their long-term stable performance, however, problems arise from material degradation - e.g. by corrosion, dynamic or static overload, etc. - and above all from filter clogging, based on fine and sticky dusts. Those dusts again are typical for thermal waste treatment processes, whereas softening of salts and tarry components and subsequently sticky or baking dust behaviour already occurs at temperatures below 500 "C. The performance of ceramic filter media is not only determined by the filtration conditions and dust characteristics. It is also determined by their microscopic structure. The structure of grain ceramic filters - mostly designed as filter candles - consists of a coarser grained supporting layer which ensures a high mechanical strength of the filter element and one or more coating layers with reduced grain size. Preferentially, filtration takes place on the top layer. For both, filtration and recleaning of fine and sticky dusts, a top layer is
required that minimises the penetration of the particles into the filter media and reduces the adhesion between filter and filter cake. Thus can be obtained by a homogeneous nano-scaled plus highly porous ceramic filter membrane. In general, the adhesion forces between particles are the lower the smaller one or both of the contacting particles are. As a consequence of the simultaneous high porosity not only the adhesion forces, but also the number of contact points are reduced by such a membrane. Investigations concerning the influence of the microstructure on filter cake formation revealed [l], that a highly porous filter structure or reduced contact points result in an enhanced porosity of the filter cake near the filter surhce. Corresponding to the porosity, the strength of the filter cake will be reduced in these areas and a complete detachment of the filter cake is promoted. In case of soft or sticky particulates it can also be expected, that wetting of the membrane will be hindered by the strong curvature of the spherical nano-particles and the high porosity.
EXPERIMENTAL SET-UP The experimental set-up of the filter device for the coating of the multilayer grain ceramics as well as for the filtration and recleaning studies is shown in figure 1. The main modules are the raw gas duct and the clean gas duct with the filter section, the aerosol generation, particle analysis and the recleaning installations. Aerosol generation is designed to produce polydisperse aerosols in two different ways. For the coating of the filters a spray type generator is used, with a maximum particle concentration of about lo7 particles/cm3. The filtration and recleaning studies are performed with a brush type dust feeder at a dosage of 2 g/m3. Coupling of other aerosol generators for monodisperse or narrow distributed aerosols is optional. Particle analysis, i.e. size distribution and number concentration, above 0.15 pm is performed at room temperature by a laser particle counter (LPC) with a dilution section during raw gas measurement at concentrations >1o5 particles/cm3. AI electrostatic classifier (EMS) combined with a condensation particle counter (CNC) enables the particle analysis fiom 10 nm to 1 pm. During a filter test run the aerosol or dust is fed into the filter section after the determination of the size and number concentration. The filtration velocity is controlled via mass flow controller (FC) and pressure drop
79
A Particle Analysis
-1
control unit EMS
air
1
control unit
EMS I
dilution
raw gas
L/dust feederv
Fig. 1: Experimental set-up for filtration and recleaning tests.
Recleaning
Aerosol Generation
as well as the clean gas particle concentration are monitored. The recleaning modus is started via maximum pressure drop and interval time, respectively. Therefore the filter section is set off-line via remote valves and the pressure in this section is increased by a loading pressure up to pL = 150 kPa. The recleaning itself is induced by the fast opening of a valve or a membrane against surrounding atmosphere, resulting in a fast pressure decrease on the raw gas side of the filter. After recleaning, the filter section is set on-line again and the next filtration cycle starts.
Table l: Structural data obtained by image analysis methods of mullite-coated Sic filter ceramic (3-2SPK).
1
I std.deviation IDore diamete? 1 meanvalue I
COATING As supporting structures for the nano-scaled A12@membranes served a commercial standard material, a mullite-coated SIC filter ceramic (3-2SPK). Structural data of this ceramic are shown in table 1, whereas the samples were used in form of discs (diameter = 45 mm). The nanometer-sized powder used for the production of the nano-scaled Alz03-membraneswas obtained by means of electrical explosion of wires [2,3]. A thin aluminium wire of about 0.7 mm is heated and evaporated in air by a short current pulse. Oxidation and condensation of the evaporated aluminium results in h e s t Alz03-particles.Due to an automatic wire feeding and a continuos air flow rate through the explosion chamber, this method enables production rates of some kg/h. The generated particles, however, consist mainly of y- and 6A1203and exhibit a wide particle size distribution between about 10 and 1000 nm. The relevant finer particles are captured in an electrostatic filter at a yield of about 10 %, corresponding to a production rate of approximately 100 g/h.For the coating only these particles were used.
80
*
I
2.4
I
I
%
1-
'
3*2SPK
std.deviation
ym ym
maximum
vm
I
I
I
I
I
1
54.9 50.2 653
maximum Feret's diameter [minimum 32 measurements per grain (angle V)] pore channel diameter by measuring the chord length (intercept lengths) of lines intersecting the pores perpendicular to the flowing direction. These lines are set every fourth pixel line of the micrographs (resolution: 1364 x 1024 pixels).
A suspension of 1 ml A12Q-powder / 1000 ml 2propanol was treated in a ultra-sonic bath for disaggregation of the powder, sprayed, dried, and then analyzed in the raw gas duct (fig. 1). Particle size distribution of the 2-propanol and the nano-aerosol is represented in figure 2. 10000 8000
ii n
5
6000
z
-0
*Q
4000
?i
Fig. 3: SEM-micrograph of the uncoated mullite toplayer of the supporting structure (3-2SPK).
[L
Particle Diameter d, , nrn
Fig. 2: Size distribution of particles generated fiom 2propanol and of the A1203nano-aerosol.
The particles resulting fiom the 2-propanol ranged between about 20 and 60 nm with a mean diameter of 47 nm. The particle size distribution of the A12Q*anoparticles was in a range of approximately 15 to 450 nm, at a mean diameter of 78 nm. Here, about 10 % of the particles were related to the 2-propanol. Coating of the multilayer Sic-discs (Fig. 3) was performed at a filtration velocity of uF = 2.5 a d s . Their filtration behaviour with respect to the A1203nanoparticles is shown in figure 4 as function of the additional pressure drop caused by the applied membranes. In case of a low additional pressure drop of Apd. = 91 Pa the surhce was only partly covered by a relatively thin nano-scaled membrane. Areas closed by glassy binder and larger mullite grains of the first coating remained on top (Fig. 4a). At higher additional pressure drop of Apadd.= 415 Pa the nano-scaled membrane was significantly thicker and covered the most part of the surface. But some larger areas closed by glassy binder and a few larger mullite grains of the first coating were also not completely bridged by the growing membrane (Fig. 4b). These uncovered areas are considered to reduce the beneficial effect of the nano-coated samples. To achieve uniform membranes the samples were coated up to an additional pressure drop of 700 Pa. Thus resulted in a multilayer filter ceramic (Nano-3-2SPK) with a homogeneous nano-scaled filter membrane on the top. The thickness of the membranes was in a range between approximately 2 and 7 pm with a porosity of about 80 to 90%. Maximum pore channel diameters were between 500 nm and 2 pm. Sintering studies on the so obtained Nano-3-2SPK filters were carried out in an electrical h a c e by varying the sintering temperature between Ts = 1000 "C and Ts = 1350 "C as well as the dwell time between 0.5 and 4 hours. To avoid sintering stresses, heating and cooling rates of 3 Wmin
Fig. 4 SEM-micrographs of the nano-scaled top layer after filtration of A1203-nanoparticlesas function of the additional pressure drop caused by the applied membrane, (a) Apad.= 9 1 Pa and (b) Apdd. = 4 15 Pa.
were chosen. Independent of the sintering parameters, the appearance of the sintered membranes (Fig. 5) showed no or only slightly differences to the unsintered membranes. A pronounced sintering neck formation was not visible. However, at sintering temperatures of Ts > 1300 "C diffusion of the clay-binder out of the supporting structure (Sic and mullite) led to a bubble formation on the top of the filters. Consequently, sintering temperature of the samples was limited to less than 1300 "C.
81
MODEL-DUST
Fig. 5: SEM-micrograph of the nano-scaled filter membrane (Apau.= 700 Pa, Ts = 1200 "C, ts = 4 h).
The sintering behaviour of the produced membranes reflects a low sintering activity in the investigated temperature range, in spite of their small particle size. Thus may be related to the y- and GA1203-powder used. In general, sintering of y- and 6A1203is characterized by a phase transformation of the unstable y- and 6-A1203into the stable a-A1203,followed by an accelerated sintering neck formation and grain coarsening [3]. The phase transformation is shifted to lower temperatures with increasing densification of the powder. Investigations concerning the sintering behaviour and the phase transformation of y- and 6A1203 [4]revealed, that phase transformation in powders required a temperature of 1300 "C, whilst in compacted bodies a complete transformation was already achieved at 1055 "C. Investigations of the author to the sintering behaviour of loose powder exhibited significant sintering neck formation and grain coarsening at sintering temperatures of 1200 "C for a dwell time of minimum 1 hour. However, the agglomerates of the loose powder were denser than the produced membrane. So, due to the structure of the membrane (high porosity, few contact points) phase transformation and subsequently sintering is hindered in the limited temperature range. On the one hand, this behaviour is favourable in that it results in a minimal additional pressure drop increase of the sintered membranes and in a minimal grain coarsening. But on the other hand, the obtained structure also leads to a low mechanical stability of the membranes. For all chosen sintering parameters the produced nano-scaled membranes could be damaged mechanically. Nevertheless, filtration and recleaning tests turned out, that the stability of the sintered membranes overcomes the appearing flow forces, more than sufficient for the relevant operating range. Therefore, it was decided to continue the filter performance tests with membranes sintered at Ts = 1200 "C for a dwell time of 4 hours. Further optimisation relating to the mechanical stability is part of actual work.
In thermal waste treatment processes, e.g. combustion, gasification or pyrolysis of bio mass and waste, the high number of input substances and possible interactions results in a wide variety of compounds in the fly ashes. In general, these particulates are characterized by a high amount of hazardous components. Moreover, they exhibit low softening temperatures, a high fraction of fine particles and - in combination with the flue gases - a high corrosion potential. The strongly fluctuating and heterogeneous waste composition additionally causes these properties to vary considerably. In order to exclude unknown toxic effects and to facilitate handling, a model dust was developed for the simulation of dusts from thermal waste treatment processes. Its composition is based on investigations of fly ashes and dusts from waste incineration plants [5-81. For real dusts, the airborne particle characteristics play a significant role in filter cake formation. The morphology of these particles can be classified in [5,6]:
- Fused Spheres (Spheroids of various colours) - Crystals (irregular shaped particles similar to calcite) - Polycrystullines (Dense agglomerates) - Opuques (Single, large irregular shaped particles) - Char (unburned, black fibrous particles) By mass, more than the half consists of large irregular shaped particles, which exists as single particles or as agglomerates. Critical to the hot gas filtration with ceramic filters are the micronhubmicron size particulates, mainly fUsed spheres and crystals. This fraction dominates the fiequency distribution with 64 to 80 %, whereas the mass is almost negligible. The better the operation of the incinerator the more the total particle size distribution will shift to smaller particle sizes and filter clogging due to the fine dust content becomes more stringent. The total dust loading in waste incineration ranges typically from 2 to 10 a m 3 [8,9] Table 2: Typical components of the fly ash from waste incinerators in comparison to the model-dust.
-
-
'CaO A1203 SO3
%
Fed-
%I
% %
1
GO
%I
CI
%
ZnO
%I
Po3
%I
Cr203
%
others
%
I
I 1
I 82
26.70 14.00 7.25 5.20 3.72 2.51 2.42 2.36 0.64 7.10
lCaO IA1203
IS03 IFeO* IkO lCl
ZnO
I
I
14.68 0.20 % 13.95 %I 0.10 % 10.05-11.56 % I 9.51 10.41 % 0-11.94 % %
I
1
-
I
INa20
%1
6.85
lMg0
%
I
1.65
As for the model-dust a composition of 50 wt.-% soda-lime glass spheres (A-glass), 30 wt.-% columnar gypsum (CaSO, 2H20), and 20 wt.-% KCI or ZnCI2 particles was chosen. In Table 2 this composition is compared to those of typical fly ash fiom waste incinerators. To promote the pollutant-forming and corrosive properties, the sulphate and chloride fractions in the model dust were increased. The chloride mixture also served to set various softening temperatures. Thermal properties of single and mixed dust components are shown in Table 3.
Table 3: Selected thermal properties of single salts, binary eutectic mixtures [101 and of A-glass.
The typical morphology of the model dust is reflected by figure 6. The mixture exhibited a bimodal particle size distribution with maxima at about 500 nm and 10 pm (Fig. 7). Typical mean particle size was about 6.3 pm at maximum values up to 50 pm. Hence, by using 50 wt.-% soda-lime glass spheres, 30 wt.-% columnar gypsum, and 20 wt.-% KCI or ZnC12 particles, not only a relevant particle size of < 30 pm could be set, but also the particle shape and dust composition of real dusts could be approximated. Furthermore, this model dust allows to simulate the softening behaviour of real dusts in a wide temperature range. For investigating the operating behaviour of filter media, the critical temperature ranges can be varied fiom about 230 "C up to about 800 "C - under exclusion of ZnCl?.
FILTRATION PERFORMANCE
1
14001
I
1
I
432 (mix.C\ Softening Point of A-Glass
I
I
I
= 600 "C
To assign the filtration performance of the nanoscaled filter membranes at conditions, relevant to hot gas filtration, filtration and recleaning tests were performed at a temperature of TF = 780°C. Based on the thermal properties of the model-dust components (table 3) at this temperature, the A-glass will be softened whilst the KCI is molten and the ZnC12 is evaporated. Therefore, a model-dust mixture of 50 wt.-% A-glass, 30 wt.-% gypsum and 20 wt.-% KCI (mixture 0) has to be used in these tests. The investigations were carried out at the experimental device represented in figure 1, at a filtration velocity of uF = 5 cm/s. Here, the dust was fed via the brush type dust feeder into the filter section. Recleaning was performed pressure drop controlled at a loading pressure up to pL = 130 kPa. 16000 150OOpL= 100 kPa
1400013000-
Fig. 6: SEM-micrograph of model-dust mixture A (50 wt.-% A-glass + 30 wt.-% Gypsum + 13.64 wt.-% ZnCl2 + 6.36 wt.-% KCI).
120001100010000-I
Fig. 8: Operating performance of the Nano-3-2SPK ceramic by filtration and recleaning of model-dust mixture 0 (50wt.-% A-glass + 30 wt.-% Gypsum + 20 wt.-% KCl) at TF = 780 "C and UF = 5 c ~ / s .
0.1
1
10
100
Diameter, pm
Fig. 7: Particle size distribution of model-dust mixture A, measured by laser scattering spectrometry (mean 6.3 pm, std. dev. 5.2 pm).
Figure 8 reflects the performance of the Nano-32SPK ceramic with respect to regeneration. The cycles are characterised by the formation of a dust layer until the maximum pressure drop for recleaning was reached. Subsequent, the differential pressure dropped down by the fast pressure decrease of the recleaning method applied and the formation of a new dust layer started.
83
Recleaning intensity of the formed filter cake could be enhanced by an increase of the loading pressure fiom 100 to 130 kPa. At pL = 130 kPa, the residual pressure drop was in the range of the initial pressure drop, corresponding to an almost complete detachment of the filter cake. Hence, with the produced nano-scaled ceramic filter membrane and the recleaning method applied, detachment of fine and sticky dust was enabled within the observed cycles, even at the chosen extreme conditions. These tests are continued with respect to long term stable operating behaviour.
CONCLUSIONS In high temperature dust separation using ceramic filter media, the filtration and recleaning of fine and sticky dusts is a major problem. For conventional filters, these dusts often lead to an enhanced filter clogging and long-term stable operation is limited. To overcome this problem, a standard filter ceramic was coated by the filtration of A1203-nanoparticleswith a mean grain size of 78 nm. Optimum membranes were achieved for an additional pressure drop of about 700 Pa. After sintering, the membranes exhibited a porosity of 80 to 90 % with pore channel diameters up to 2 pm. Membrane thickness was in a range between 2 and 7 pm. However, due to the limited sintering temperature of 1300"C, sintering of the membranes was hindered. Although the stability of the membranes overcomes the appearing flow forces, they could be damaged mechanically. Thus will be solved, among others, by optimization of the multilayer structure and by the use of supporting materials which allow higher sintering temperatures. Filtration and recleaning tests performed with a h e and sticky model-dust showed an excellent filtration and recleaning behaviour of the produced nano-scaled ceramic membranes, even at a temperature of 780°C. Therefore, a great potential for the separation of critical dusts is offered by those membranes, not only in thermal waste treatment processes but also in product recovery. The beneficial properties of the nano-particles will be further enhanced by using ceramic nanoparticles with smaller grain size distribution and smaller mean diameter. For investigating the operating behaviour of filter media, the critical temperature ranges of real dusts can be simulated with a four or three component model-dust fiom about 230 "C up to 800 "C. Here, especially the important softening, melting or boiling of local eutectics can be replicated.
84
ACKNOWLEDGEMENT The authors wish to acknowledge the European Commission DG XII for partially sponsoring this work under the Industrial and Materials Technology Programme (BRITE EURAM 111) in the BRITE EURAM Basic Research Project "HOTFIL" BRPR - CT97 0472.
REFERENCES T. Pilz, Particle Properties relevant for Hot Gas Cleaning with rigid barrier filters and their Characterization at High Temperatures, in Schmidt et al. (eds.): High Temperature Gas Cleaning, Institut f3r Mechanische Verhhrenstechnik und Mechanik der Universitiit Karlsruhe (TH), Karlsruhe, (1996), 132-143 I.V. Beketov et al., Synthesis of Nanometer-Sized Powders of Alumina and Titania Using the Electrical Explosion of Wires. Roc. Fourth Euro Ceramics Vol. 1, (1995), 77-82 A. Weisenburger, Anwendung der Hochleistungsimpuls- und Mikrowellentechnik zur Herstellung nanokristalliner pulver und Festkiirper. Dissertation Universim Karlsruhe, (1999) V.Ivanov, V. Khrustov, Proc. Fourth Euro Ceramics Vol. 2, (1995), 281-288 A.J. Chandler et al., Municipal solid waste incinerator residues. Elsevier Studies in Environmental Science 67, Elsevier, 1997, ISBN 0-444-82563-0 Waste Programme, Waste Analysis, Sampling, Testing and Evaluation Program, Phase 1 Final Draft Report Prepared for Environment Canada, US EPA and the International Lead Zinc Research Organisation, unpublished 1993 N.F. Glen and J.H. H o d , Modelling Refhe Incineration Fouling, Publication C 118/88 National Engineering Laboratory, Glasgow 1988 P.J. Hayes, R Schulz, H. Leibold, B. Zimmerlin, B. O'Reilly and P. Hahn, Hot gas cleaning using advanced ceramic filter technology for municipal waste incinerators(HOTFIL) - Review of ash and gas compositions, TRAWMAR Workshop Proceedings Sept. 1998, Rhodes P.J. Hayes, H. Leibold, B. Zimmerlin, R. Schulz, A. Zagorski and P. Hahn, Hot gas cleaning using advanced ceramic filter technology for municipal waste incinerators (HOTFIL) - Synthetic dust for modelling high temperature filter cake behaviour in MSW incinerators, TRAWMAR Workshop Proceedings Sept. 1999, San Sebastian (10) G.J. Janz, Molten Salts Handbook, Academic Press, N Y 1967
PRODUCTION AND CHARACTERIZATION OF TiCN-BASED MATERIALS FOR CUTTING TOOL APPLICATIONS F. Monteverde*, A. Bellosi*, C. Zancolij**, M. G. Faga** (*) CNR-IRTEC, Istituto di Ricerche Tecnologiche per la Cerarnica, 48018 Faenza (RA), Italy (**) CNR-ILM, Istituto Lavorazione Metalli, 10043 Orbassano (TO), Italy
ABSTRACT Nearly full dense Ti(C,N)-based materials were obtained by hot pressing and gas pressure sintering. The starting compositions were based on mixtures of TiCo.5No.5and WC(S%Co) powders with moderate content of metal additives (6.2wt% Ni, 2.lwt% Co) or additive fiee. Microstructure and mechanical properties were studied and related to starting composition and processing conditions. Cutting tools were obtained fiom the sintered materials and preliminary results on metal machining tests were reported.
metallic systems, the occurrence of liquid phase during sintering, the mutual solubilities between these complex carbideslnitrides and the binder are key factors that may bring effective benefits improving the expected performances [5-171. In this work TiCN-WC(Co)-based materials (with moderate addition of Ni-Co or additive fiee) were hot pressed (HP) and gas pressure sintered (GPS). Microstructure and some structural properties are described and discussed. Preliminary results of metal machining tests, in finishing operation using the additive fiee material as cutting tool, are reported.
EXPERIMENTAL PROCEDURES INTRODUCTION Titanium carbonitrides are very hard materials with increasing technical importance. They are mainly applied in composites with various metal carbides and/or binders (cermets) for metal cutting operations [I-41. Cermets can be described as composites of ceramic particles (carbides, borides, nitrides or a combination of them) bonded together with a metalbased system (i.e FeyNi Co, Mo or an alloy of them) [l-3, 5-71. The notable hardness and wear resistance are provided by the hard particles, while the binder phase improves the ductility and yields better toughness and thermal shock resistance. For many applications, included metal cutting tools, the ideal material should exhibit the highest hardness as well as the best toughness. However, as hardness and toughness are generally antagonist properties, a compromise has to be set or, keeping in mind the specific application of the component, to favour one of them. Titanium carbonitrides based cermets are characterized by high creep and difhsion wear resistance, due to the formation during sintering of the well known core-rim type structure [l]. The microstructure of TiCN-based cermets ranges fiom usual ceramic configuration with low content of metallic phases (i.e binder) at grain boundaries, to metal-ceramic composites where alloys constitute a frame with ceramic particles dispersed inside. Moreover, the composition of the binder is dependent on the overall stoichiometry and nitrogedcarbon content of the hard constituents [l, 8-10]. The physical and mechanical properties of these cermets can be adjusted within certain limits to meet the requirements of the application through the selection of the raw powders, the design of the compositions and the tuning of the processing parameters. Concerning these hard
The commercial raw powders used in this work are listed as follows: Ti(C,N) (HCST 50/50 grade E, Germany), WC-S%Co (Nanocarb, Nanodyne - USA), Ni (Carbonyle), Co (Sherrit). The Ti(C,N) powder has a stoichiometry TiCo.sNo.5and mean particle size 0.36
w*
The following two mixtures (amounts in wt%) mixture A: TiCo.5No.5 + 16.7 (WC-5Co) mixture B: TiCo.5No.5 + 15.3(WC-5Co)+6.2 Ni+2.lCo were prepared. Sintering of mixture A and B was accomplished through low vacuum hot pressing (HP) and gas pressure sintering (GPS). The thermal cycles were scheduled as follows: HP1: 20"C/min to 1700"C, 1700°C for 30 min 30 MPa; HP2: 20"C/min to 1650"C, 1650°C for 45 min, 5 MPa; GPS: 10"C/min up to 1650°C and 1650°C for 30 minutes in flowing gas (0.1 m a ) ; raising the gas pressure up to 2 MPa and holding it for 30 minutes. The microstructural characteristics were analyzed on the dense materials by XRD and SEM/EDX. Several mechanical properties were investigated using the following methods: -Young's modulus (E): resonance fie uency method on plate (28.0x8.0x0.8)mm ; -Vicker's microhardness (HV1.O) with a load of 1Kg; - 4-pt flexural strength ( 0 ) on bars (25.0~2.0~2.50)mm3 up to 1000°C; -fracture toughness (Klc) by the direct crack measurement method (load 10Kg) on polished surfaces using the models proposed by Anstis et al. [ 181 and by Evans et al. [19]. The cutting performance of inserts obtained fiom the material HPA were tested in longitudinal turning tests in comparison with a similar commercial tool. Continuous cutting tests were carried out according to the IS0 3685 standard for the part concerning tool wear criteria. During machining tests the flank wear
9
85
liquid phase. Carbon-poor Ti(C,N) residues survive and act as nuclei for the precipitation of (Ti,W)(C,N) upon solidification of the liquid phase (Figs. le-h). The melt resulting immediately after the onset of the liquid phase is notably rich of W. The inner layer surrounding the TiN-rich nuclei (Figs. le-h) is hardly enriched of W than would correspond to the composition of the powder mixtures [l]. Our SEMEDX analyses on these selected areas c o n fiied these features previously reported [ 11. The sintering conditions and particularly the nitrogen partial pressure of the atmosphere decisively influence the phase equilibria, and thus grain growth and the composition of the sintered phases [l]. Morphology of sample HPB (Fig. lb) and GPSBl (Fig. lc) as well as grain size (Table I) are similar. The binder is well distributed and forms thin layers around the hard ceramic grains. Concerning the sample GPSB2 an evident grain growth took place (Fig. Id) with the residual binder mainly located in pockets at the triple points of the hard grains (Fig. 1h). Ti(C,N) cores and (Ti,W)(C,N) outer rims, clearly revealed by SEM-BSED investigations (Fig le-h) are isomorphous (cubic) phases with fully overlapped XRD patterns (along the forward angular range). Semiquantitative calculations of the rim composition in the three cermets gave the following T i m atomic ratio: 6 for HPB, 6.5 for GPSB1 and 18 for GPSB2. Since the total volume of the rims in sample GPSB2 is larger than that in samples GPSBl and HPB (the amount of W introduced as WC is the same) the content of W taken up by the rim shows a reverse relationship with the rim volume in the materials. The binder, characterized by a cubic symmetry, was analyzed upon mirror-like flat surfaces in areas where binder is well localized. The results indicated that the binder of sample GPSB2 contains 3 times more more W than the cermet GPSB1, in a reverse relationship with the results aforementioned on the amount of W in the outer rim. In fact, fiom the XRD patterns the cell parameter of the binder is higher for sample GPSB2 (a= 0.357 nm) due to a higher W content, than for sample GPSB1 (a=0.355 nm) and HPB (a4.354 nm).
evolution was monitored using an optical microscope. SEM analyses were performed on the cutting tool after the tests.
RESULTS AND DISCUSSION Microstructure
The microstructure of the dense additive fiee sample HPA shows the typical core-rim structure; the Z-constrast imaging mode allows to distinguish the inner core Ti(C,N) fiom the outer rim (Ti,W)(C,N) (Fig. le). Ti(C,N) can be drawn as a solid solution between the boundary phases TIC and TiN, (Ti,W)(C,N) is a solid solution as well containing W ions. These two phases are isomorphous with fairly close cell parameter [7, 161. From the spectra obtained by windowless energy dispersive microanalysis, only part of the original TiCo.sNo.5grains constitute the inner core that in turn contains the main percentage of the nitrogen fiom the starting mixture. Any trace of W was found in the inner cores. On the other hand WC undewent complete dissolution during the heat treatment, and cubic (Ti,W)(C,N) phase precipitated as a rim around Ti(C,N) nuclei. Although the sample was hot pressed, some microporosity remained (Fig. la) and this could be ascribed to the nitrogen in the formulation partially released in the temperature range 1200-1300OC [l]. Grain boundaries appear clean with no evidence of intergranular phases. Dense cermets, HPB - GPSB1 - GPSB2, sintered through a Ni-Co based liquid phase, featured as faceted grains of hard phases embedded in a binder matrix. The microstructure of the as-sintered materials is compared in Figs lb-d (fkacture surfaces) and lf-h (polished surfaces). The grains of the hard phase show the typical core-rim microstructure already observed for material HPA. The solubility of TiN in liquid metals is much lower than that of T ic and WC. Therefore when at suitable temperature mixtures of TiCN and WC react with liquid metals, Tic is preferentially dissolved in comparison to TiN while WC is totally dissolved in the
Table I. Specifications of the sintering cycles and some microstructural features of the produced materials. Sample
Thermal cycle
Atmosphere
#
86
Density
Mean Grain size
gr/cm3
Y O
Cun
Crystalline phases
HPA
HP1
Vacuum
5.75
100
0.55
(Ti,W)(C,N), T K N )
HPB
HP2
Vacuum
5.82
98
0.75
GPSB1
GPS
N2
5.92
99.8
0.6
GPSB2
GPS
Ar
5.90
99.4
1.2
(Ti,WC,N), Ti(C,N) binder (a=0.354 nm) (Ti,W)(C,N), Ti(C,N) cubic binder (a=0.355 nm) (TiW)(C,N), Ti(C,N) cubic binder (a=0.357 nm)
Fig.
SEM micrographs from fractured (left column, seconury electron mode) and polished (right column, back scattered electron mode) of material HPA (de). HPB @If),GPSB1 (c/g) and GPSB2 (a) The .arrowed points correspond to: 1- inner core Ti(C,N), 2+ outer rim (Ti,W)(C,N), 3-+ W- rich inner rim, 4+ binder.
87
Table 11. Grain size and mechanical properties measured on the dense materials.
(*Anstis et al. model [18], "Evans et al. model [19]) These features appear strongly dependent on the processing atmosphere. It is well known that the grain growth rate in nitrogen-containing cermets is so much lower than in nitrogen fiee alloy [1,5], due to the fast achievement of saturation concentration of the liquid phase with respect to Tic and WC concentration, that a complete dissolution of the Ti(C,N) particles is never reached. At suitable temperature (namely >1600"C), the solubility of nitrogen is several orders of magnitude lower than that of carbon and titanium, therefore high solubility of carbon (into the liquid phase) compared with that of nitrogen enables a carbon flux and thus a close to uniform carbon activity throughout the sample [5]. On the contrary, a nitrogen-activity gradient will result depending on the sintering atmosphere. It was reported an enhanced grain growth, due to a high nitrogen-activity gradient favoured by the argon atmosphere [5]. The sample GPSB2 in fact has a grain size double than the sample GPSB1. The atmosphere condition during hot pressing should favour nitrogen activity as well, but the presence of an applied pressure had an effect to limit grain growth. Mechanical properties
The mechanical properties of the dense materials are reasumed in Table 11. They clearly depend on the starting compositions and processing conditions. Hardness: as previously reported [13], it increases with decreasing grain size of the hard ceramic phase. The additive fiee sample HPA in fact is the hardest while cermets fiom mixture B fit the expected behaviour on the basis of the grain size. Fracture toughness: the three cermets (fiom mixture B) resulted tougher than ample HPA. The alloy constituting the binder improves toughness and an effect of grain size is evident as well. The values obtained by the formula of Anstis [ 181 are higher than the values obtained by the formula of Evans and Charles [19], but the relative trends result similar. In any case the values estimated fiom the former model (measured Young's modulus as entry), besides the crack path length and hardness, can be assumed as more appropriate. Young's modulus: the measured values are in agreement with the ones reported in literature for similar compositions [l]. The slight variations of the experimental data could be due to residual porosity and
88
to specific properties of the binder alloy (only for cermets). Flexural strength room and high temperature strength (Table II) of the material HPA is lower than that of the three cermets from mixture B. (Ti,W)(C,N) phase has an inherent better wettability with respect to the binder not only in the liquid but also in the solid state, hence greatly improving strength of the cermets [13]. Due to the difference of the atomic radii, the binder alloy is strengthened by the W dissolved in it. Concerning the room temperature strength, the influence of fabrication defects and of machining often mask strengthening effects related to the properties of the hard phases andor the binders. In any case, the strongest sample resulted the GPSB2 (i.e gas pressure sintered in Argon), which contains more W in the binder phase than the other samples. High temperature flexural strength in air is affected not only by the intrinsic microstructural characteristics but also by new defects induced by extensive oxidation on surface after hot exposure. The latter represents the main degradating mechanism of the strength in sample HPA. As far as the three cermets concern, the tests up to 800°C fall in the low temperature domain where similar cermets exhibit brittle behaviour and non measurable plastic deformation. From 800°C up to about 1100°C limited plastic deformation was observed [15]. Our experiments c o n f i i a plastic deformation at 1000°C until fiacture occurs. This behavior may be ascribed to typical microstructural features (hard phase grain size, composition of binder). It was proposed in fact that the yield stress at constant plastic deformation is reversely proportional to the square root of the grain size and that dislocation motion occurs first in the metallic binder and is associated with an anelastic relaxation process controlled by solute atom diffusion in the metallic phase [151. Behaviour as metal cutting tools
The potential of the sample HPA as metal cutting tool in finishing operation was tested following the cutting parameters and the work materials shown in Table 111. The results of longitudinal turning tests (Table IV), in terms of tool life, observed until reaching a flank wear threshold VB~,,=0.30 mm, show the better performance of the HPA insert,
Table 111. Work materials, parameters for cutting tests Material Cutting speed Vc, (dmin) Feed (f, &rev ) Cutting depth (ap, mm 1 Lathe Type of insert Toolholder Lubricant Dimension of the workpieces
UNI c4 5 210 HB UNI 36NiCrMo4 300 HB UNI 42NiCrMo4V 3 15 HB
Table IV Results of turning tests (vB~,,,,=0.3 mm) Work Cutting conditions I Tool life (min) material f I a,, HPA [ C o r n .
I
250,300,400 0.15,0.20,0.25 0.5, 1.5 CNC UTITA (30 Kw ) SNUN 120408 T (Chamfered ) Sandvik Coromat CSRNL 3225P12-1C None 0 = lOOmm,L=450mm
under machining conditions of low feed rate and high cutting speed, due to the specific combination of mechanical properties of the produced insert. The wear morphology of the commercial and HPA inserts is compared in Figs. 2.
Tests at higher feed and cutting depth evidenced worse results than commercial materials. Further experiments under different metal machining conditions are in progress.
CONCLUSIONS Different Ti(C,N) based materials were produced with and without the addition of metallic binders (Ni,Co). The influence of starting compositions and of processing conditions on microstructure and properties was studied. A core-rim structure resulted: the core is constituted by partially dissolved raw TiCo,5No.s powder particles while the rim by (Ti,W)(C,N) solid solution. Grain size and composition of the rim depend
HPA insert. t=4 min. VBR=0.24mm
Commercial insert, t = 3.41 min, VBB= 0.47 mm Figure 2. Optical micrographs of the flank wear evolution from HF'A (a,b) and commercial (c,d) inserts. Workpiece material: 36NiCrMo4, HB 300; cutting conditions: Vc= 400 dmin, feed= 0.15 mmhev, depth of cut= 1.5 mm.
89
mainly on the sintering atmosphere. Gas pressure sintering under argon favours grain growth, less amount of W in the rim of the hard grains and more W in the binder phase, in comparison to cermet gas pressure sintered under nitrogen or hot pressed in vacuum. These features influence mechanical properties. Gas pressure sintered cermets exhibited the highest strength in comparison to the hot pressed one. The presence of a binder effectively strengthened and toughened these carbonitrides. Preliminary tests on material HPA as metal machining tool evidenced better performance than commercial inserts under finishing operations at low feed rate and cutting depth.
p
Ti
1
1
ca
0
-As-sintered
0
1
2
3
4
5
KeV
Fig. 3 SEM micrograph (upper) fiom the flank of HPA insert after cutting test. EDX spectrum (lower) was acquired upon the dark halo and compared to that of the as-sintered material.
Acknowledgements This work was supported by MURST and C N R under the National Project “Innovative Materials”, Law 5% ‘95.
REFERENCES (1) P. Ettmayer, H. Kolaska, W. Lengauer, K. Dreyer, Ti(C,N) Cermets- Metallurgy and Properties, Int. J. ofRefi. Met&Hard Mat, 13, (1995) 343-51. (2) S. Zhang, Material Development of Titanium Carbonitride-based Cermets for Machinig Applications, Key Engineering Materials, Vol. 138-140, TTP Switzerland, (1998) 521-43. 90
T. Watanabe, T. Dotsu, T. Nakanishi, Sintering properties and Cutting-Tool Performance of Ti(C,N)-based Ceramics, Key Engineering Materials Vol. 114, TTP Switzerland, (1996) 189-266. B. Clark, B. Roebuck, Extending the Application Areas for Titanium Carbonitride Cermets, Refi. Met&HardMat. 1 1 (1992) 23-33. P. Gustafson, A. Ostlund, Binder-phase Enrichment by Dissolution of Cubic Carbides, Int. J of Rep. Met&Hard Mat. 12 (1 993194) 129-36. H. Holleck, K. Kleykamp, Constitution of Cemented Carbide Systems, R&HM,( 1982)11216. T. Laoui, H. Zou, 0. Van der Biest, Analytical Electron Microscopy of the CoreRim Structure in Titanium Carbonitride Cermets, Int. J. of Refi. Met&Hard Mat. 1 1 (1992) 207-12. M. Ehira, A. Egami, Mechanical properties and Microstructure of Submicron cermets, Znt. J. of Ref?, Met&HardMaf. 13 (1995) 313-19. M. Rynemark, Investigation of Equilibria in the Ti-W-C-N System at 175OoC, Refi. MefhHard Mat. 10 (1991) 185-93. (10) T. Laoui, 0. Van der Biest, Effect of Tic addition on the Microstructure and Properties of Ti(C,N)WC-Co-Ni Cermet, J. Mof. Sci. Left. 13 (1994) 1530-32. (1 1) Z. Xingzhong, L. Jiajun, Z. Baoliang, M. Hezhuo, L. Zhenbi, Wear Mechanisms of Ti(C,N) Ceramic in Sliding Contact with Stainless Steel, J. Mat. Sci. 32 (1997) 2963-68. (12) S. Bolognini, G. Feusier, D. Mari, T. Viatte, W. Benoit, High Temperature Mechanical Behaviour of Ti(C,N)-Mo-Co Cermets, Int. J. of Re$-. Met&Hard Mat. 16 (1998) 257-68. (13) Z. Zackrisson, U. Rolander, G. Weine, H. 0. Andren, Microstructure of the Surface Zone in a Heat-Treated Cermet Material, Int. J. of Refi. Met&Hard Mat. 16 (1998) 3 15-22. (14) Liuying, L. Huayi, B. Hengzhen, L. Gouchun, Influence of CH4/N2ratio on microstructure and mechanical properties of PCVD-Ti(C,N,-,), Znt. J. of Refi. Met&Hard Mat. 13 (1 995) 369. (1 5) H. Pastor, Titanium Carbonitride-Based Hard Alloy for Cutting Tools, Materials Science and Engineering, A105/106 (1988) 401-9. (16) S. Zhang, C. D. Qin, L. C. LimySolid Solution of WC and TaC in Ti(C,N) as revealed by Lattice Parameter Increase, Int. J . of Refi. Met&Hard Mat. 12 (1993-94) 329-33. (1 7) Bellosi, V. Medri, F. Monteverde, Development and properties of Ti(C,N)-WC based materials, submitted to the J. Am. Cer. Society. (18) G. R. Anstis, P. Chantikul, B. R. Lawn, D. B. Marshall, A Critical Evaluation of Indentation Techniques for Measuring Fracture Toughness, 11, Strength Method, J. Am. Ceram. SOC.64 (1992) 539-43. (19) G. Evans, E. A. Charles, Fracture Toughness Determination by Indentation” J. Am. Ceram. SOC.59(1980) 571-2.
SENSITIVITY CHARACTERIZATION TO FLAMMABLE GAS Seong -Eun Sim, Sung-Churl Choi Department of Ceramic Engineering Hanyang University Seoul 133-791, South Korea
ABSTRACT There has been considerable interest in recent years in thin film gas sensing materials. In comparison with conventional sintered bulk gas sensor, thin film gas sensor materials have good sensitivity, optimum temperature and selectivity. The use of tin oxide based materials as gas sensing elements for a wide range of reducing gases has been extensively studied. Since the electrical conductivity of tin oxide (Sn02) thin film is changed by adsorbing gases, such as H:, C&, LPG, it has been applied for gas sensors. The SnOz is a n-hpe semiconductor in which excess electrons produced by o.xygen deficiency act as donors and therefore electrical conductivity depends on the changes of its lattice defects. The femc oxide (a-Fe203) has attracted interest of investigators as a material for semiconductor gas sensors. This interest is connected with the properties of a-Fe203which has a good gas sensitivity, a good long term stability and a lorn humid-effect without the application of noble metal doping.
EXPERIMENTAL In this study, the gas sensing device was fabricated by CVD. The fabrication process was as follows: first, washing the polished alumina substrate for the improved bonding effect. Second, sputtering the Pt electrode as 0,300 nm on the front. Third, coating aFe203 at 175 "C and Sn02 at 350 "C. Fourth, heat treatment at various conditions. Fifth,bonding the NiCr wire on the back. Si'uth, sensitivity measurement and analysis. Fifth, bonding the Ni-Cr vire on the back. Sixth, sensitibity measurement and analysis. Following the reaction to a-Fe2031Sn02 of deposing device.
a-Fe203 Fe(C0)s 175 " C
Precursor Deposition Temperature Deposition
Sn02 SnCL 350 "C
3,s minute
l,2,5,8,10 minute
Tim e
Pressure Flow Rate
Ar : 0
2
I I
I
1t o n 106: 26sccm
I I
I
1t o n 197 : 17 sccm
I
Table 1: Deposition conditions of variable factors Table 1 is shoning deposition conditions. The gas sensikity to flammable gases (C&, H2, LPG) was measured. The sensitivity "S" to gases is defined here as S = & - b) <7 k,,where R,and RgaJ are the sample resistance in clear air and in air containing testing gases respectively. This device was heat treated at 400 "C,450 "C, 500 "C, 550 "C, 600 "C for 2hr to enhance the gas sensitivity. The heat treated deh-ice at 500 "C for 2hr had the best properties and especially shows high sensitivity to H2 gas. The figures 1 and 2 are showing the SEM morphology. The sensitivity to gases was studied in the temperature range from 100 "C to 300 "C in order to find the optimum detection temperature. The optimum temperature was 175 "C.
Fig .1 . No heat treatment
4Fe(CO)s+1302+ 2Fe203+ 20C02 SnC14 + O2 + SnO-,+ 2 Ci A reaction ratio of the Ar and O2was controlled by MFC (Mass Flow Controller, MKS.Co.,)
F, : Flon7of Reactant F, : Flow of Carrier Po : Output Pressure P, : Vapor Pressure of Reactant
Fig. 2. Heat treatment at 500 "C
91
RESULTS The fabricated device which has been deposed with a - F e z q : 5 min and Sn@ : 8 min shows the best properties. The gas sensitivity for H2 ranges from 62% 2: 76%, for (2% from 16 2: 58 YOand for LPG from 8% 2: 37%, that is, the sensitivities are H2> CH, > LPG. The sensitivity espectally showed the better characteristics at low concentration range (500 ppm a 1,000 ppm) than high concentration range (3.000 ppm s 10,000 ppm). The heat treatment daice had the high sensitikity to 10% rather than not treated it in LPG. Although 21 sec for no heat treatment device, a sensing response time is 12 sec for heat treatment device. The multi-layer gas sensor device has a good long-term stability to converge the sensitivity 30% in LPG at 1,000 ppm.
Fig. 6 Sensitivity of the CH., Gas
Fig. 7 Sensitivity of the LPG Fig. 3 Sensor resistance on air as doping ratio
i
1
i
Tm~--aam (F)
Fig. 4 Sensor resistance in air at various temperature
Fig. 8 Comparison of the heat treated and the untreated device
r C
/--/ -
I
'
--cn,;
Fig. 5 Sensitiklty of the H2 Gas 92
Fig. 9 Long-term stability
CONCLUSIONS aFe~O3/ Sn02 thin film devices are prepared by a CVD process. The multi-layer device has the better sensitivity and selectivity than single-layer delices. 1.
2. 3.
4. 5.
The device which deposed at aFe203: 5 min and Sn02 : 8 min had the best properties. The selectively showed H2 > CH4 > LPG. The sensitivity especially showed the better characteristic at low concentration rang (500 ppm = 1,000 ppm) than high concentration range. The heat treatment device ha the h g h sensiti\ity to 10% rather than not treated it in LPG. The multilayer gas sensor device has a good longterm stability to converge the sensitivity 30% in LPG at 1,000 ppm.
REFERENCES R. Lalauze and C. Pijolat, A New Approach to Selective Detection of Gas by an Sn02 Solidstate as Sensor, Sensors and Actuators, B. 5,5563 (1985). A. A. Vasiliev and M. A. Polykarpom, Change of Ferric Oxide (Fe203) Semiconductor Conducti.r;ity Type in the Interaction "lth Reducing Gases, Sensors and Actuators, B 7, 626-629 (1992). C. C. Chai, J. Peng and B. P. Yan, Characterization of alpha-Fe203. Thin Films Deposited by Atmospheric Pressure CVD onto Alumina Substrates, Sensors and Actuators, B, 31,412-116 (1995). Sayago, J. Gutierrez, L. Ares, J. I. Robla, M. C. Honillo, J. Getino, J. Rino and J. A. Agapito, The Effect of Additives in Tin Oxide on Sensitivity and Selectivity to NO, and CO, Sensors and Actuators, B, 26-27, 19-23 (1995). C. F. Schaus, W. J. Schaff and J. R. Shealy, OMPVE Growth of GaXlnl-XP/GaAs(AlyGalyAs) Hetero-structures for Optical und Electronic Device Applications, J. Cvst. Growth, 77, 360-366 (1986). T. Suzuki, T. Yamazaki and M. Azumaya, Hydrogen Gas Sensing properties in Polycrystalline Tin Oxide Films of Submicron Thickness, J. Cera. SOC.Jpn., 97, 1263-1273 (1989). M. Kanamori. Y. Okamato. Y. Ohta and Y. Takahashi, Thickness Dependence of Senshity of Gas Sensor. J. J. Cera. Soc. Jpn.. 103, 113116 (1995). K. D. Schierbaum, U. Weimar and W. G pel, Comparison of ceramic thick-film and thin-film Chemical sensor based upon SnOz, Sensor and Actuators B. 7 (1992) 709-716. J. Watson, The tin oxide gas sensor and its appplication, Sensor and Actuators, 5 (1984) 2942. (10) K. H. Song, S. H. Shin, J. I. Park, K. J. Park and S. J. park, Effect of film oxygen content on gas SensiFity of tin dioxide thin film, Proceedings of The 31d East Asian Conference on Chemical
(11)
(12)
(13)
(14)
(15) (16)
Sensor, Seoul National University, 193-199 (1998). Zheng, Jiao, Zhang, Jinhuai, Liu, Stability of Sn02/Fe203 Multilayer thin film gas sensor, Proceedings of the 3rd East Asian Conference Seoul National Univ. ( 1998) 22 1-226. D. Kohl, Surface processes in the detection of reducing gases with SnOz based devices, Sensor and Actuators, 18 (1989) 71-114. M. H. Madhusudhana Reddy and A. N. Chandorka, Response study of electron evaporated thin-film hn Oxide gas sensor, Sensor and Actuators B, 9 (1992)l-8 R C. Evans, An introduction to crystal chemistry, Cambridge University Press. (1966) 148-149. P. A. Mulheran and J. H. Harding, The stability of SnO: surfaces, Modelling Simul. Mater. Sci. Eng. 1 (1993) 39-44 W. G. pel, K. D. Schierbaum and M.D. Wiemhofer, Solid State Ionics.
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11. Performance / Reliability
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EVALUATION OF MECHANICAL RELIABILITY OF SijNq NOZZLES AFTER EXPOSURE IN AN INDUSTRIAL GAS TURBINE H. T. Lin*", M. K. Ferber', M. van Roodeb (") Metals and Ceramics Division, Oak Ridge National Laboratory
Oak Ridge, TN 37831-6068, USA (b) Solar Turbines Incorporated San Diego, CA 92101, USA
ABSTRACT This paper provides a review of a recent study undertaken to evaluate the mechanical reliability of SN88 Si3N4ceramic nozzles exposed in an industrial gas turbine. Two field tests with exposures time of 10 and 68 h were carried out. The first 10 h field test revealed minor changes in both microstructure and strength of airfoil region. However, nozzles with crack generation, which initiated in low temperature airfoil region (< 1000°C), were observed after 68 h test, resulting in the termination of second engine test. Analyses of the cracked nozzles by scanning electron microscopy and x-ray diffraction revealed changes in microstructure and secondary phase due to environmental exposure. Dynamic fatigue tests on as received SN88 Si3N4at intermediate temperatures in air suggested that changes in secondary phase could result in the formation of an extensive damage zone, due to the generation of a large residual tensile stress, which substantially reduced the mechanical reliability and probably led to the failure of turbine nozzles.
INTRODUCTION Silicon nitride ceramics with reinforcing elongated grain microstructure are leading candidates for use as high-temperature, structural components in advanced gas turbines due to their superior thermomechanical properties [ 1-51. Recent ceramic gas turbine programs at both Solar Turbines and Rolls Royce Allison [6-81 have executed many field tests to increase the experience base concerning the behavior of ceramic components in industrial gas turbine environments. A key lesson learned in both programs is that environmental effects may severely limit the long-term reliability of Si3N4materials. In particular, the high velocities and presence of water vapor in the environment can lead to the volatilization of the normally protective silica layer. Researchers have shown that the presence of water vapor leads to the formation of a gaseous Si(OH)4 species via a reaction with the silica layer [9-111. The rate of Si(OH)., formation of this species and thus the rate at which the Si3N4 is consumed by its continued oxidation are a function of gas velocity, pressure of the water vapor, and total pressure of the environment. In addition to the volatility issues, which may lead to loss of function due
to excessive changes in component dimensions, the effect of the environment upon the long-term mechanical reliability must be understood as well. One approach addressing this issue is to evaluate the component properties directly using small test specimens. For example, the availability of the small dog-bone tensile specimen [ 121 has provided for the measurement of the tensile stress rupture properties of silicon nitride samples from first stage turbine blades. Results showed that there were significant differences in the creep response characteristics obtained for various parts of the ceramic blade due to variations in grain size and secondary phase content and chemistry. More importantly the data obtained from testing specimens prepared from the production billets greatly over-estimated the lifetime of components with complex shapes. This paper describes' the results of a recent component verification study involving SN88 Si3N4 nozzles exposed in an industrial gas turbine. Silicon nitride nozzles after exposure to 10 h and 68 h field tests were examined in this study. Scanning electron microscopy was used to elucidate the changes in the microstructure arising from the oxidation process. In addition, the stability of secondary phase(s) was evaluated using x-ray analysis. The exposed surface strength was measured as a function of field test time using a miniature biaxial specimen.
EXPERIMENTAL PROCEDURES The nozzles examined in this study werefabricated from NGK SN88 Si3N4 material (NGK Insulators Ltd., Nagoya, Japan). They were gas-pressure-sintered using rare-earth sintering additives and then post-heat-treated by NGK to form a protective silica surface layer. The predominant secondary phase present after densification is Yb4Si207N2,designated as the J-phase. The first stage ceramic nozzles (Figure 1) were designed for retrofit into a Centaur 50s turbine engine (Solar Turbines Incorporated, San Diego, CA) [ 131. The first engine test was carried out with one hour full load and 10 h total run time. The second engine test was planed for 100 h endurance test. However, the test was terminated after 68 h total run time including 15
97
Figure 1 Solar SN88 first stage Si3N, nozzles after engine test time of 10 h (a) and 68 h (b). start/stop cycles due to crack generation in the nozzles (Fig. lb). X-ray diffraction was used to identify the predominant secondary phases in the airfoil as a function of engine test time. Scanning electron microscopy (SEM) was first used to examine the surfaces of the airfoils and platforms after removal of the nozzles from the engine. Nozzles were subsequently sectioned by making a longitudinal cut parallel and adjacent to the trailing edge. These pieces were subsequently polished and examined with SEM. The biaxial flexure strength [ 14-16] was measured for samples from selected vanes using the ball-on-ring arrangement. Specimens were machined from both the airfoil and platform surfaces by first diamond core drilling small cylinders having nominal diameters of 5.5 mm. Each cylinder was then machined on one face only until the thickness was 0.4 to 0.5 mm. In this way the tensile face of each specimen always consisted of the exposed surface of either the airfoil or platform. The details of testing fixture and procedures can be found in Ref. 17
RESULTS AND DISCUSSION Macroscopically, the SN88 Si3N4nozzles after 10 h engine test revealed little change on airfoil surface features except some reddish deposits (Fig.la). On the other hand, the nozzles after 68 h engine test showed a white-colored, powder-like scale with crack generation (Fig. lb). X-ray analysis indicated that the whitecolored scale present on airfoil surfaces of 68 h nozzles was mainly the Yb2Si207 plus Yb2Si05, which formed due to the oxidation of initial secondary phase (J-Phase) present in the as-received nozzles. Note that the dominant crack always initiated in the low temperature (< 1OOO°C) airfoil region [181. SEM examinations of airfoil surfaces of both 10 and 68 h engine tested nozzles showed the presence of corrosion pits associated with Fe-Si-0 deposits (Fig. 2). The Fe element detected in the deposits was attributed the erosion of superalloy components in the engine.
98
Figure 2. SEM of airfoil surface features of SN88 silicon nitride nozzles after (a) 8 h and (b) 68 h engine test. Also, the 68 h tested nozzles revealed more surface porosity due to oxidation and volatilization of Si3N4as compared with 10 h tested nozzles. In addition, SEM analysis confirmed that the white-colored, powder-like scale observed in 68 h nozzles was indeed a secondary phase, which consisted of Yb, Si, 0 and trace amount of Y. The accumulation of the Yb2Si207 plus YbzSiOs on the surface was most likely due to the selective recession of Si3N4grains, as seen in Fig. 2b. The Si3N4 grains were removed by a combination of oxidation and volatilization processes, resulting in the accumulation of secondary phases on the airfoil surface.
The SEM micrographs of the polished crosssections are shown in Figure 3. The SN88 Si3N4 nozzles showed minor changes in microstructure, e.g., pores in secondary phase and very limited intergranular cracks, in the region adjacent to the airfoil surface after 10 h engine exposure (Fig. 3a). However, extensive
generation of intergranular cracks plus pores in the secondary phase was observed in nozzles after 68 h engine exposure (Fig. 3b). The changes in microstructure, i.e., formation of environment-induced damage zone, were attributed to the oxidation of Si3N4 materials in the presence of high-pressure water vapor. A previous study had showed that the presence of hightemperature, high- pressure water steam would lead to formation of an extensive subsurface damage zone in Si3N4 materials [19]. Also, results of Vickers indentation in oxidation zone in the vicinity of crack initiation site showed a preferential crack growth along the oxidation z o n e h l k material boundary, which was in sharp contrast to the symmetrical indentation feature observed in the bulk material region (Fig. 4). The feature of preferential crack growth in the oxidation zone suggested the presence of a large residual stress. SEM examinations also showed that the airfoiVplatform transition region, where the dominant crack normally Figure 4. SEM micrographs show indentations in (a) oxidation zone and (b) inner bulk region. The inserted photo shows the presence of an oxidation zone (light region). initiated and led to failure of nozzles, exhibited a more extensive damage zone formation (-30-40 pm) than the middle airfoil surface region (-10 pm). Note that the temperature and air velocity in the transition region were lower than the middle airfoil surface region. The finite-element-analysis of temperature distribution at steady state during engine operation indicated that the temperature in the airfoiVplatform transition region ranges from 800 to 95OOC [18]. The SEM observations just described are in contrast to the experimental results and oxidation models [9-111. It was possible that in the middle airfoil region the high gas velocity aggravated the removal of silica due to volatilization of the silica layer, resulting in less extensive damage zone evidence. On the other hand, in the transition region little, if any, protective silica would form after the initial silica layer present in the as-received nozzle receded and could lead to an extensively accumulated damage zone (Figure 3b). Strength test results at room temperature of samples extracted from nozzle airfoils after engine test are summarized in Figure 5 . Note that the disk samples were tested with exposed, pressured sides as tensile surfaces. Results showed that the 10 h tested samples exhibited -10% decrease in strength, while the 68 h tested samples exhibited -30% strength degradation. The extensive strength degradation observed in nozzles after 68 h engine exposure is due to the extensive formation of subsurface damage zone (Fig. 3b).
Figure 3. SEM micrographs of polished cross-section of nozzle airfoil surface region after (a) 10 h and (b) 68 h engine test.
The crack initiation in low temperature airfoil/platform transition region has led to a hypothesis that SN88 Si3N4might exhibit a mechanical instability at intermediate temperatures in air. Note that previous creep studies indicated that SN88 Si3N4 exhibited excellent creep performance at temperature L 1038°C in air [20]. Recently, studies of dynamic fatigue by
99
1000
results in the oxidation zone of cracked nozzles (Fig. 4). Note that there is a minor material volume increase (-5%) for J-phase changing to Yb2SiOs. The development of a high tensile stress would then lead to fracture of Si3N4grains and, also, generation of multiple intergranular cracks. Similar mechanical instability due to phase changes was previously reported for Si3N4 materials sintered with Yz03 additive [23].
800 n
2
8
600 400
b cn 200 0
0
0.5
1
1.5
2
2.5
3
3.5
(mm) Figure 5 . Stress versus displacement curves of biaxial samples extracted from airfoils of SN88 nozzles after 10 and 68 h engine test. Wereszczak et al. have shown that some commercial Si3N4 materials exhibited substantial strength degradation at temperature 5 850°C in air [21]. Therefore, a dynamic fatigue test in four-point bending was conducted to evaluate the mechanical reliability of SN88 Si3N4at 850°C in air [22]. Figure 6 summarizes the earlier dynamic fatigue response of SN88 Si3N4 tested at 85OOC in air at stressing rates of 30 MPds and 0.003 MPa/s. Note the NT154 Si3N4,which contained YzSi207as secondary phase, is used as a benchmark in this study. Results showed that the fracture strength of NT154 Si3N4 was similar those obtained at room-temperature and was independent of stressing rates accompanied by a high fatigue exponent of 87. On the other hand, strength of SN88 Si3N4was very sensitive to stressing rates. For instance, samples tested at 30 MPds exhibited similar strength as those obtained at room temperature, while samples tested at 0.003 MPds exhibited 43% strength decrease. In addition, SN88 Si3N4 exhibited a low fatigue exponent of 16, indicative of a high susceptibility to slow crack growth (SCG) process at high temperatures in air.
-
The fracture surface of SN88 samples tested at 85O0C/0.O03 MPds in air revealed a light-colored ring surrounding the bend bar (as shown in Fig. 7), indicating the presence of an environment-affected zone (EAZ). SEM examinations showed that a 30 pm damage zone, containing multiple cracks plus pores in secondary phase, developed in the light-colored EAZ. Note that the EAZ developed in all four-side surfaces, suggesting the development of EAZ was not a stresspromoted phenomenon. Results of x-ray analysis for 0.003 MPds tested samples also indicated a change in secondary phase from J-phase to Yb2Si207 plus YbzSiOs.
-
The phase transformation phase from J-phase to Yb2Si207 could introduce an extreme high residual tensile stress in the EAZ due to a -64% decrease in material volume [23], consistent with the indentation
100
The results of dynamic fatigue tests at 85OOC in air suggested that an EAZ developed in the airfoil/platform transition region, where the temperature was low and no protective silica layer would form during the exposure, after the initial silica layer present in the as-received nozzles recessed. The generation of extensive multiple cracking in EAZ would significantly reduce the mechanical reliability and, thus, greatly increase the susceptibility of nozzles to SCG processes. Consequently, a critical crack would readily initiate at the transition region and lead to the failure of nozzles. On the other hand, the reason why the EAZ was not observed in the higher temperature regions; i.e., middle airfoil region, appeared to be due the elimination of the EAZ by oxidationlvolatilization processes and thus continuous material recession. 1000 900 =
.
800
.
-=.
850°C in air
. . . 0
300
0 C
200 0.001
SN88 ' NT154
~~
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Stressing Rate (MPds) Figure 6. Fracture strength versus stressing rate results of SN88 silicon nitride tested at 85OOC in air.
SUMMARY Silicon nitride nozzles were designed and fieldtested in an industrial gas turbine. Dominant crack developed in the airfoil/platform transition region after 68 h test, resulting in the termination of engine test. Both SEM and x-ray analyses indicated changes in microstructure and secondary phase due to the exposure in engine environment. An extensive damage zone, containing multiple cracks and pores in secondary phase, developed in the subsurface region after 68 h engine test. A supporting dynamic fatigue test at 85OOC in air suggested that changes in secondary phase could result in the formation of an extensive damage zone, due to the generation of a large residual tensile stress, which substantially reduced the mechanical reliability and led to the failure of turbine nozzles.
ACKNOWLEDGEMENTS The authors thank Drs. M. Lance, T. N. Tiegs, and P. F. Becher for reviewing the manuscript. Research is sponsored by the U.S. Department of Energy, Office of Power Technologies, Microturbine Technologies Program, under Contract DE-AC05-000R22725 with UT-Battelle, LLC.
( 7 ) V. Parthasarathy, M.van Roode, J. Price, S. Gates, S. Waslo, and P. Hoffman, “Review of Solar’s Ceramic Stationary Gas Turbine Development Program, Pro. 6th Int. Symp. Ceramic Materials and Components for Engines, Arita, Japan, (1997) 259-264. (8) R. Wenglarz, S. Ali, and A. Layne, “Ceramics for ATS Industrial Turbines,” to be published in the Proc. of ASME Turbo Expo ‘96 Conference, Birmingham, UK, June 10-13, 1996. (9) E. J. Opila and R. E. Hann, Jr., “Paralinear Oxidation of S i c in Water Vapor, “J. Am. Ceram. SOC.,Vol. 80 [l] (1997) 197-205. (10) E. J. Opila, “Variation of the Oxidation Rate of Silicon Carbide with Vapor Pressure,” J. Am. Ceram. SOC.,Vol. 82, No. 3, pp. 625-36, (1999). (11) J. L. Smialek, R. C. Robinson, E. J. Opila, D. S. Fox, and N. Jacobson, “Sic and Si3N4 Recession due to SiOl Scale Volatility Under Combustor Conditions, Advanced Composite Materials, Vol. 8, No. 1 (1999) 33-45. (12) H. T. Lin, P. F. Becher, M. K. Ferber, and V. Partasarathy, “Verification of Creep Performance of a Ceramic Turbine Blade,” Proc. 2nd Int. Symp. Science of Engineering Ceramics, Osaka, Japan (1 998). (13) A. F. Kirstein and R. M. Woolley, “Symmetrical Bending of Thin Circular Elastic Plates on Equally Spaced Point Supports,” Journal of Research of the National Bureau of Standards - C. Engineering and Instrumentation, Vo17 1C, No 1, Jan-Mar 1967. (14) R. Thiruvengadaswamy and R. 0. Scattergood, “Biaxial Flexure Testing of Brittle Materials,” Scripta Metall. et Mater., Vol 25 (1991) 25292532. (16) D. K. Shetty, A. R. Rosenfield, P. McGuire, B. K. Bansal, and W. H. Duckworth, “Biaxial Flexure Tests for Ceramics,” Ceram. Bull., Vol 59 [12] (1980) 1193-97. (17) M. K. Ferber, H. T. Lin, V. Parthasarathy, and R. A. Wenglarz, “Degradation of Silicon Nitrides in High Pressure, Moisture Rich Environments,” to be published in the Transactions of ASME, 2000 (IGTI, ASME TURBO EXPO 2000, May 8-11, 2000, Munich, Germany). (1 8) 0. Jimenez, Solar Turbines, private communication. (19) M. K. Ferber, H. T. Lin, and J. Keiser, “Oxidation Behavior of Non-Oxide Ceramics in a HighPressure, High-Temperature Steam Environment,” Mechanical, Thermal and Environmental Testing and Performance of Ceramic Composites and Components, ASTM STP 1392, American Society for Testing and Materials, West Conshohocken, PA, 2000. (20) A. A. Weresczak and T. P. Kirkland “Creep Performance of Candidate SIC and Si3N4 materials for Land-based, Gas Turbine Engine Components, ASME Paper 96-GT-385, to be published in the Proc. of ASME Turbo Expo ‘96 Conference, Birmingham, UK, June 10-13.1996. I‘
Figure 7. Fracture surface of SN88 bend bar tested at 85O0C/0.O03 MPa in air exhibited an environmentaffected zone as indicated by arrows.
REFERENCES D. Anson, K. S. Ramesh, and M. DeCorso, “Application of Ceramics To Industrial Gas Turbines”; DOE/CE/40878- 1 &2; Battelle Columbus, March 1991. H. E. Helms, R. A. Johnson, and L. E. Groseclose, “AGT 100-Advanced Gas Turbine Technology Development Project,” Proc. 23rd Automotive Technology Development Contractors’ Coordination Meeting P- 165, Warrendale, PA, March (1986) 137-55. M. van Roode, “Ceramic Retrofit Program,” Proc. Joint Contractors Meeting: FEIEE Advanced Turbine Systems Conference FE Fuel Cells and Coal-Fired Heat Engines Conference, DOE/METC93/6132 ( 1993) 77-93. D. Carruthers and L. Lindberg, “Critical Issues for Ceramics for Gas Turbines,” Proc. 3rd Int. Symp. on Ceramic Components and Materials for Engines, Westerville, Ohio (1988) 1258-1272. P. J. Haley, R. L. Holtman, L. E. Groseclose, S. J. Hilpisch, and A. H. Bell 111, “Advanced Turbine Technology Applications Project (ATTAP)”, Proc. 29th Automotive Technology Development Contractors’ Coordination Meeting, Warrendale, Penn. (1992) 19-29. M. van Roode, W. D. Brentnall, K. 0. Smith, B. D. Edwards, L. J. Faulder, and P. F. Norton, “Ceramic Stationary Gas Turbine Development - Third Annual Summary”, ASME Paper 96-GT-460, to be published in the Proc. of ASME Turbo Expo ‘96 Conference, Birmingham, UK, June 10-13, 1996.
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(21) A. A. Weresczak, T. P. Kirkland, H. T. Lin, and S. K. Lee, "Intermediate Temperature Inert Strength and Dynamic Fatigue of Candidate Silicon Nitrides for Diesel Exhaust Valves," to be published in Proc. Ceram. Eng. Sci. (2000). (22) H. T. Lin, M. K. Ferber, and P. F.Becher, "Mechanical Reliability of Yb-containing Silicon Nitride at Intermediate Temperatures," presented at Am. Ceram. SOC.Annual Meeting, St. Louis, April 30-May 3,2000. (23) F. F. Lange, "High Temperature Deformation and Fracture Phenomena of Polyphase Si3N4 Materials," Proc. 2nd NATO Advanced Study Institute of Nitrogen Ceramics (1981) 467-490.
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FRICTION AND WEAR OF ADVANCED CERAMICS K.-H. Zum Gahr UniversitBt Karlsruhe, Institute of Materials Science I1 and Forschungszentrum Karlsruhe, Institute for Materials Research I P.O. Box 3640,76021 Karlsruhe, Germany
ABSTRACT Tribological performance of structural ceramics is of great interest because of the unique combination of properties promising high potential for components in unlubricated or marginal lubricated sliding contact. As a function of operating conditions a mild to severe wear transition can occur. Alumina ceramics offer high hardness, stiffness, chemical inertness and thermal stability combined with the economical advantages of relatively low costs and good availability in a lot of different qualities. Hence, the experimental results presented are focussed on this class of advanced ceramics. But many conclusions can be also transferred to other ceramics. The mechanisms for the wear transition are generally discussed. Results are presented from laboratory tests about the mfluence of microstructural parameters such as grain size, second phases etc. of monolithic and multiphase alumina ceramics on friction and wear in unlubricated sliding contact against A1203counterbodies. It is shown that microstructural design can be very effective in reducing friction and wear in a given tribosystem
INTRODUCTION Advanced structural ceramics such as A1203,Sic, Zr02, Si3N4or Sialon are increasingly used for components with high tribological, mechanical, chemical a n d or thermal requirements. Their favourable properties are high stiffness and hardness, compressive strength, temperature stability, corrosion and wear resistance in n m y cases, as well as low density. However, monol i h c ceramics can suffer severe damages under high loads, owing to their inherent brittleness and lack of defect tolerance. As a consequence during the last decade those ceramics were primarily developed for achieving hgh strength and fracture toughness. Most applications of engineering ceramics are, however, tribological, e.g. draw-cones, guides, cutting tool inserts, seal rings, cylinder liners, cam roller followers, bearing parts, medical prostheses, sand blast nozzles etc. The microstructural demands for ceramics used as mechanical load-bearing components of high strength and fracture toughness and for those used as tribological components resulting in low friction and wear can be substantially different. Using bend tests the long crack resis-
tance is measured and effects such as crack shielding or microcracking, crack branching and crack deflection may be favourable. In contrast, tribological interactions are concentrated on substantially smaller contact areas like the size of surface asperities or smaller than the grain size of the ceramic material. Only during the last years the microstructural design of ceramics has gained increasing considerations for tribological parts [ 1-41. Tribological behaviour of monolithic ceramics such as alumina is characterized by a transition from mild to severe wear, whereby mild wear is used to indicate a relatively low friction coefficient and the value of wear coefficient k (wear volume divided by normal load and sliding distance) smaller than mm3/(N.m) while the term severe wear is used for k > 1O4 mm3/(N.m). In the following, wear transition, wear m e c h s m s and microstructural features influencing the tribological behaviour in sliding contact of self-mated ceramics are discussed. The experimental data presented were measured in laboratory tribometers using ball-on-block or ring-on-block geometries during unlubricated oscillating sliding contact. Alumina ceramics were chosen as test materials because of their unique combination of hardness, corrosion resistance, thermal and hydrothermal stability as well as their economical advantages and broad acceptance in many applications. Properties of Contacting Bodies Chemistrymopcgraphy of Surfaces
Materials: Porosity Hardness, Strength Grain size Young’s Modulus Second Phases Fracture Toughness
Performance
Contact Geometry Contact Mode
Vibrations
Fig. 1 : Tribological performance depending on properties of the contacting solid bodies, loading parameters of the hibosystem and environmental conditions.
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WEAR MECHANISMS AND WEAR TRANSITION Tribological behaviour depends strongly on the type of tribosystem, environmental and interfacial conditions as well as microstructural features and resulting properties of the materials mated in sliding contact (Fig. 1). Hence, friction and wear are not inherent material properties but a h c t i o n of parameters of the complex system considered. It follows, that selection of a suitable ceramic and also the design of the microstructure have to consider the specific application. Fig. 2 shows schematically several types of solidsolid and solid-liquid-solid interactions on ceramic surfaces during sliding contact. Transgranular (a) or intergranular (b) microcracking due to mechanical interactions of surface asperities is promoted by pores or preexisting flaws in the surface or subsurface region. Plastic deformation (c) limited on a thin surface zone, mutual transfer (d) of material or wear debris and micro-abrasion (e) can occur depending on the ceramics involved and the operating conditions, e.g. load, temperature, humidity etc. Third-body layers (f) of aggregated wear debris or tribochemical reaction films (g) can be formed on the surfaces, e.g. aluminium hydroxide films on alumina or amorphous silicon oxide films on Sic or Si3N4,respectively. Increasing humidity or adsorbed water molecules can result in hydrated surface films (lubricious oxides) and hence low values of friction coefficient. Smooth surfaces may be formed by continuous removal of these thin reaction films (h) that favours locally elastohydrodynamic lubrication (i) in the presence of liquid media.
Fig. 2: Schematic representation of interactions between ce-
ramic solid surfaces during unidirectional sliding contact. Mild to severe wear transition occurs during unlubricated sliding of monolithic ceramics, e.g. alumina, depending on operating conditions such as contact load, speed and time of sliding, contact temperature and environment. Mechanisms of mild wear are mainly characterized by formation and detaching of tribochemical reaction films, plastic micro-deformation and micro-abrasion. With increasing friction power (Fig. 3) the wear intensity or the amount of wear at a given sliding distance increases owing to initiation and propagation of cracks resulting in localized spalling of
104
material by fatigue processes and pull-outs of single grains. Densification of grain fragments in the contact area lead to so-called third-body layers between the mated ceramics. Severe wear is connected with high friction coefficient and wear as a result of trans- andor intergranular cracking under the high tensile stresses at the rear end of the tribological contact. Wear damage occurs at the surface or below the surface owing to micro-fracture with grain pull-outs and detachment of surface layers by delamination fracture. The tribological contact between the mated solids strongly influences formation and time depending removal of thick thirdbody layers consisting of densified wear debris.
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Fig. 3: Amount of wear (or wear intensity) and mechanisms of interaction as a function of developed friction power per unit area. According to modelling [3] the load dependent wear transition is caused by the mechanical tensile stresses at the rear end of the contact area and the speed dependent wear transition by thermal stresses owing to frictional heating. The severity of tribological contact can be described by loading parameters and is influenced by microstructure and resulting properties of the mated materials. More general, the severity of tribological contact should increase with friction power per area, i.e. the product of contact pressure p, sliding speed v and friction coefficient p (Fig. 3). Without considering thermal effects, e.g. at low sliding speeds in unlubricated sliding contact, wear transition occurs when the mechanically applied contact pressure exceeds the critical value which increases with fracture toughness (short crack resistance, particularly), and hardness but decreases with increasing fixtion coefficient, Young’s modulus and initial defect sizes of the mated ceramics, e.g. grain sue, pores and microcracks [4]. For estimation of the relevant value of fracture toughness, i.e. the material resistance to propagation of transgranular cracks or cracks at weak grain boundaries or at interfaces of second phases, the small scale of contact, the complex stress situation and the short crack length of equal to or less than grain size has to be considered. Remembering that tribological behaviour depends on parameters of the tribosystem used, the type of counterbody material can be an important factor. Balls of AlzO3, steel 100Cr6 or cemented carbides WC6Co were mated with alumina (A123) blocks in unlubricated oscillating sliding contact in laboratory air at the normal load of 80 N [5]. Fig. 4 shows that mating A l z 0 3 balls with the alu-
mina block resulted in a sharp increase in wear intensity if a critical applied contact pressure was surpassed. This wear transition, however, was shifted to an about four times greater value of contact pressure if using the WC6Co ball and wear intensity increased gradually with further increasing contact pressure. A far different behaviour was measured if mating the 100Cr6 balls with the monolithic alumina Al23. No sharp wear transition occurred and wear intensity was relatively high even at low values of contact pressure. Increase in wear intensity was only moderate with increasing contact pressure. In the order of A1203, WC6Co and 100Cr6, hardness decreased and fracture toughness of the balls increased. Due to their low fracture toughness, the A1203balls suffered a lot of microcracking and pullouts of single grains during high contact pressure and the amount of linear wear occurred to about equal parts on the balls and the blocks after a sliding distance of 144 m. The lowest fnction coefficient was measured on pairs with the WC6Co balls. Worn surfaces of these pairs displayed evidence of thm tribochemical reaction layers which were rolled up during oscillating sliding contact. These layers (maybe W03, C0W04 or more complex [6]) offered lubricating properties and reduced amount of wear substantially compared with the alumina block mated with A1203balls. Two third of the total linear wear was measured on the A1203blocks and only one third on the WC6Co balls. Sliding pairs with steel 100Cr6 balls exhibited extensive wear (about 90 %) on the balls. The steel worn by a combination of micro-abrasion and tribo-oxidation. Material transfer from the steel balls to the ceramic block led to sliding contact between steelheel and/or iron oxideliron oxide, respectively, and resulted in high friction coefficient and high amount of linear wear after 144 m of sliding. However, the sensitivity to micro-fracture controlled wear processes at high contact pressure, i.e. above about 1000 MPa (Fig. 4) was substantially lower for the 100Cr6/A1203-pairsthan that for the other pairs owing to the lowest Young's modulus and hardness (or yield strength) as well as highest fracture toughness of the steel balls. E 100000
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MICROSTRUCTURAL DESIGN Important microstructural parameters of engineering ceramics are porosity, size, shape and size distribution of grains, grain boundary phases, dispersed particles of different morphology or fibers of a second phase, and last but not least surface or subsurface flaws such as isolated large scale pores, cracks or inclusions of a foreign substance. The effect of porosity was measured on self-mated 3 mol.% Y2O3-Zr02(TZP) ceramics [7] during sliding contact in laboratory air and at presence of distilled water. Friction coefficient was independent but wear intensity decreased clearly with increasing porosity up to about 6 vol.%. At greater porosity both friction and wear increased substantially. Studies [8,9] on self-mated dense alumina and zirconia ceramics showed that amount of wear in unlubricated sliding contact was strongly reduced by decreasing average size of globular grains while friction coefficient was not influenced. Figure 5 shows the effect of average size of globular grains on the critical contact pressure for onset of wear transition [S]. Wear intensity (slope of the curve of amount of linear wear versus length of the wear path) of the self-mated alumina increased drastically above a critical contact pressure. The critical value for this wear transition was shifted to greater contact pressure with decreasing average grain size fiom 12.2 pm to 0.85 pm. Values of wear intensity in the region of severe wear, i.e. at the high wear level, were substantially lower with the fine-grained material.
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Elongated grains with high aspect ratios offer effects such as crack bridging and crack deflection while small equiaxed grains reveal greater energy absorption ability. Multimodal grain structures can result in self-reinforcing effects as in Si3N4 ceramics. Heterogeneity of the ceramic microstructure promotes crack initiation but may reduce crack propagation processes depending on the loading conditions. The property of a grain boundary phase is one key factor for the onset of wear transition. Grain boundary phases influence grain boundary strength and toughness as well as residual stresses or stress and strain accommodation, respectively, at the grain inter-
105
faces under the applied tribological load. A ductile grain boundary phase resulting in high grain boundary strength and low elastic mismatch to the grains retards crack initiation and hence the transition to severe wear. Crack bridging and crack deflection are more important for propagation of long cracks and hence improve the bulk fracture toughness of the material.
the microstructure in a surface zone of a thickness of a few hundred micrometers only. For this reason, surface modification was carried out on commercial A1203 ceramic by using a C02 laser beam. Figure 6 shows the commercial dense A1203ceramic AD3 and two surfacemodified ceramics Al24Hf and Al24TiN, respectively. The commercial alumina ceramic A124 was surfacemodified (200 pm thick) by laser treatment [10,11] resulting in a multiphase structure of soft and hard phases. The ceramic Al24Hf (1750 HV5w) contained about 25vol.-% of a soft fine lamellar eutectic A1203-Hf02 phase along the boundaries of the A1203 grains. The ceramic Al24TiN (1900 HV5w) consisted of about 12 v01.-% of TiN particles (size of 1 to 5 pm), 16 v01.-% of a Ti-0-A1 grain boundary phase and balance alumina. Hardness of the microstructural modified ceramics was clearly lower than that of the commercial reference ceramic A123 (2030 HVSW).
Fig. 7: Wear path on the ceramics A123 and A124Hf, respectively, caused by a Rockwell diamond sliding unidirectionally under the normal load of 60 N and the sliding speed of 10 d m i n .
Fig. 6: Scanning electron micrographs (SEM)of commercial monolithic alumina ceramic A123 and the laser surface-
modified alumina ceramics A124Hf and A124TiN, respectively. From the foregoing it follows that a designed multiphase microstructure containing small grain size and a proper grain boundary phase should improve tribological performance. All tribologically induced interactions between two solids mated in unlubricated sliding contact are concentrated on a relatively thin surface zone. Hence, it may be sufficient and economical to modify
106
Figure 7 shows an example how the load-bearing capacity of alumina materials under unlubricated sliding action of a Rockwell diamond can be substantially improved by microstructural design. Creating the soft lamellar eutectic phase of 25 vol.% in the ceramic A124Hf (Fig. 6) reduced the hardness to 1750 HV, compared with 2030 HVSWof the commercial monolithic A1203 ceramic A123 used as reference. At the given normal load of 60 N and the sliding speed of 10 d m i n the ceramic A123 failed by intergranular cracking while the
laser surface-modified ceramic A124Hf showed a smooth groove (Fig. 7) owing to plastic deformation but no failure by microcracking. The beneficial effect of embedding hard (TiN) and soft (HfOz, Ti-0-A1) phases in alumina on the onset of wear transition during unlubricated sliding wear is shown in Fig. 8. It becomes obvious, that for both A124Hf and Al24TiN the transition fiom mild to severe wear or reverse occurs at substantially greater values of the contact pressure than for the ceramic A123. Different benefits of the second phases for the tribological performance of the surface-modified alumina ceramic in unlubricated sliding contact against A1203balls can be distinguished, e.g. grain refinement and partially alignment of columnar grains about normal to the loaded surface owing to laser treatment, reduction of Young's modulus by embedding soft phases in the thin modified surface zone, improvement of grain boundary strength and toughness as well as formation of lubricious oxides (Ti02J by embedding TiN particles.
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Instantaneous Contact Pressure p, MPa Fig. 8: Instantaneous wear intensity dWl/ds of the ceramics A123, A124Hf and A124TiN during oscillating sliding contact against unlubricated A1203balls (0 10 mm) at the normal load of 80 N versus calculated instantaneous contact pressure p. Figure 9 shows the influence of rel. humidity of the surrounding air at room temperature on friction coefficient and amount of linear wear of the commercial monolithic alumina A123 and the microstructurally designed ceramic A124TiN in oscillating sliding contact against A1203balls [lo]. Both friction coefficient and linear wear of the pair with A123 decreased continuously with increasing humidity down to the value at presence of distilled water in the contact area. In contrast, for the ceramic A124TiN these tribological characteristics were independent of humidity at and above about 30 % r.h. to a very good approximation and the values were also substantially lower in humid air than those of the pairs with the ceramic A123. The favourable performance of the sliding pair A1203/A124TiNwas attributed to the fine-grained microstructure and the grain boundary phase avoiding micro-fiacture controlled wear processes. Instead tribochemical films on the alumina matrix as well as on the TiN particles and Ti-Al-0 grain boundary phase were formed at sufficient humidity and resulted in low friction and wear.
Fig. 9: Stationary values of friction coefficient p and amount of linear wear W,* of (a,b) A123 and (c,d) A124TiN after 144 m of oscillating sliding contact against A1203balls at the normal load of 40 N and the sliding speed. of 0.02 m/s as a hnction of relative humidity or distilled water, respectively. Mechanisms of mild wear were explained by plastic deformation, micro-abrasion and formation as well as detaching of tribochemical films of low shear resistance (lubricious oxides or hydroxides). Studies [lo] on the A1203ceramics A123 and A124TiN as a function of temperature up to 500 'C, showed substantially lower values of friction coefficient and amount of linear wear of A124 TiN at ambient temperature and above of 400 "C than those values of the monolithic A123 (Fig. 10). In unlubricated sliding contact against A1203balls, friction and wear increased with increasing temperature from 28 "C to 110 "C for both ceramics. The beneficial effect of tribochemical films such as aluminium hydroxide on alumina and/or TiOz.x on TiN particles with A124TiN occurred in humid air (above about 30 % r.h., Fig. 9) only. Increasing test temperature reduced formation of AlOOH surface films on A1203 and led to enhanced friction owing to desorption of water molecules. At 400 "C and above Ti02.x films could be formed on the TiN particles and/or the Ti-0-A1 grain boundary phase and reduced friction coefficient. 5
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Fig.10: Stationary values of friction coefficient p and amount of linear wear WI* of (a,b) A123 and (c,d) A124TiN after 144 rn of unlubricated oscillating sliding contact against Alz03 balls at the normal load of 40 N and the sliding speed of 0.02 m/s as a function of test temperature.
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CONCLUSIONS Transition from mild to severe wear can be very detrimental for the reliability of tribological performance of ceramic materials in unlubricated sliding contact. This transition is caused by change of interacting mechanisms from plastic micro-deformation, microabrasion and tribofilm formation to micro-fracture with grain pull-outs, grain comminution, formation of thirdbody layers and delamination fracture of surface layers. The examples of experimental results on alumina ceramics showed that this transition in wear behaviour can be delayed or even avoided by microstructural design depending on the operating conditions of the ceramic components considered.
ACKNOWLEDGEMENTS These studies were supported by the Ministerium
flir Wissenschaft und Forschung des Landes BadenWiirttemberg in the scope of the KKS programme and are now supported by the Deutsche Forschungsgemeinschaft (DFG) in context with the Sonderforschungsbereich SFB 483 ‘‘High performance sliding and friction systems based on advanced ceramics”.
REFERENCES S. Jahanmir, X. Dong, Mechanism of mild to severe wear transition in a-alumina. J. Tribol., 114(1992)403-411.
S.M. HSU, Y.S. Nagarajau, Haiyan Liu and Chuan He, Microstructural design of ceramics for wear. Proc. Intern. Tribology Conf. Yokohama 1995,427-432. N. Chen, K. Adach and K. Kato, Transition mechanisms of wear modes in sliding of ceramics. Proc. Intern. Tribology Conf., Yokohama 1995,409-414. K.-H. Zum Gahr, Modelling and microstructural modification of alumina ceramic for improved tribological properties, Wear 200 (1996) 215224. K.-H. Zum Gahr and J. Schneider, Multiphase A1203 ceramic with high resistance to unlubricated sliding wear. Proc. Intern. Tribology C o d , Yokohama 1995,397-402. H. Engqvist et al., Tribofilm formation on cemented carbides in dry sliding conformal contact. Wear 239 (2000) 219-228. W. Bundschuh and K.-H. Zum Gahr, Influence of porosity on friction and sliding wear of tetragonal zirconia polycrystal. Wear 151 (1991) 175-191.
K.-H. Zum Gahr, W. Bundschuh and B. Zimmerlin, Effect of grain size on fiiction and sliding wear of oxide ceramics. Wear 162-164 (1993) 269-279. A. Krell, D. Klaftke, Effects of grain size and humidity on fretting wear in fine-grained alumina, A120JTiC, and zirconia. J. Am. C e m SOC.79 [5] (1996) 1139-1146. I.T. Lenke and K.-H. Zum Gahr, Mit TiNPartikeln lasermodifizierte A1203-Keramikim reversierenden Gleitkontakt unter Variation der Luftfeuchte und der Temperatur. Mat.-wiss. u. Werkstoffiech. 29 (1998) 57-65. K.-H. Zum Gahr, J. Schneider: Surface modification of ceramics for improved tribological properties. Ceramics International 26 (2000) 363-370.
Lifetime Prediction for Silicon Nitride Sheldon M. Wiederhorn National Institute of Standards and Technology 100 Bureau Drive. Gaithersburg, MD 20899-8500
ABSTRACT This paper reviews lifetime prediction methodologies for high-temperature structural ceramics. The methodologies consider failure from subcritical crack 0 to 1000 "C, and growth at low temperatures, ~ 8 0 "C creep, or creep-rupture at high temperatures. The paper discusses three methods of characterizing crack growth in silicon nitride: dynamic fatigue, static fatigue and a statistical analysis of lifetime. These techniques yield comparable results, suggesting that dynamic fatigue is the preferred procedure for characterizing crack growth. In the second part, the paper discusses a strain-based methodology for lifetime prediction. The effect of creep on acceptable engineering temperatures and stresses is considered for an allowable strain of 0.5% and a failure probability of 0.0001. For a given stress and failure time, a reduction in operating temperature of as much as 40 "C is needed when strain rather than creep rupture is considered as the failure criterion. An additional 25 "C reduction is needed to limit the failure probability to 0.0001.
INTRODUCTION Silicon nitride was developed originally as a material for gas turbines and is supposed to be competitive with air-cooled superalloys that can operate in turbines with inlet temperatures in excess of 1300 "C. To be economical, silicon nitride has to be used in an uncooled state and must be resistant to creep, creep-rupture and corrosion at high temperatures. In addition, because stresses can be very high during heatup and "trip" conditions, it also has to be resistant to failure at low temperatures, where brittle rupture and subcritical crack growth are problems. Many of these goals have been achieved. Today, several commercial grades of silicon nitride have the requisite fracture toughness and resistance to subcritical crack growth needed for structural reliability at low temperatures, and, therefore, are also suitable for high temperature operation. To be sure, there are still serious problems that must be overcome before silicon nitride can be used in turbines. Coatings must be developed to protect silicon nitride from water attack by high-pressure steam at elevated temperatures, and improved toughness under severe impact conditions is necessary to resist impact by metal pieces ingested or generated by turbines during operation'. These problems are currently being investigated in a number of industrial and government laboratories.
To ensure long term reliability, techniques to predict lifetime also had to be developed. Over the years, two general methods of lifetime prediction were developed; one is based on toughness, strength and crack growth susceptibility, the other on parameters related to creep and creep rupture behavior. Other modes of failure in ceramics (pit formation due to oxidation*, evaporation of S i c and Si3N4due to water attack', surface damage due to hard particle impact', etc.) are not considered here. No general method of handling these causes of failure has yet been developed. The first set of methods, is fracture mechanics based. Subcritical crack growth from microstructural defects normally present in the body is assumed to limit lifetime. Under the influence of external stresses, cracks grow from the defects until they reached a critical size, at which point crack growth accelerates and failure occurs. Using the framework of fracture mechanics, relationships were developed to describe lifetime as a function temperature and applied stress. Although the theory of fracture mechanics is used, no prior knowledge of the underlying causes of subcritical crack growth is required for the method. Simple power-law equations of applied stress and failure stress are used for lifetime prediction. The second method of lifetime prediction applies to failure at elevated temperature and is based on the creep and creep rupture behavior. Lifetime is predicted from the creep behavior, either by assuming a fixed strain to failure (the part becomes too large for its function), or from an empirical relationship between creep rupture lifetime and creep rate, e.g., the Monkman-Grant relationship3. The form of the creep equation is based on a theoretical treatment of creep behavior. Therefore, the better the prediction of creep rate, the more accurate the lifetime prediction. Finally, both techniques just mentioned give a median or average prediction of lifetime. In real design situations, knowledge of the scatter in lifetime prediction is needed in addition to the mean or median prediction. This requirement necessitates the use of statistics in the lifetime prediction scheme so that lifetime can be predicted at low failure probability. The techniques that are used for high and low temperature predictions differ because the modes of failure differ substantially. In this paper, we discuss the techniques that are used to predict component lifetime and the scatter in lifetime. Two modes of failure are assumed4. At low temperatures subcritical crack growth from preexisting defects is assumed to be the main cause of failure. At
109
high temperatures failure is assumed to occur as a consequence of creep, which can either cause the part to be too long for the engine, or can result in rupture from creep induced damage. A discussion of these two modes of failure starts the paper. The paper ends by dealing with scatter in lifetime, showing how low temperature scatter in strength determines the maximum stress that can be applied to a component, while high temperature scatter establishes the maximum temperature. All data discussed in this paper were obtained on current commercial grades of silicon nitride'.
FAILURE BY SUBCRITICAL CRACK GROWTH In ceramics, fracture at low temperatures is usually preceded by crack growth from flaws or cracks present at surfaces. Crack growth is normally a consequence of a corrosive environment (water for glass), but can also occur in a vacuum or another inert environment'. Because crack growth precedes catastrophic failure, a delay to failure is observed in components subjected to static loads. The delay is the time required for a crack to grow from the initial to the critical crack size. Subcritical crack growth also leads to a time dependence of the strength: the faster a material is loaded, the stronger it is. A logical framework for understanding subcritical crack growth and for designing structural materials is provided by the science of fracture mechanics, which is summarized in many excellent texts6*'. The principal assumptions behind the lifetime prediction methodology are: 0
0
0
Microstructural defects are present from which cracks grow. The crack growth velocity is determined by the crack-tip stress intensity factor, KI. Failure occurs when the crack reaches a critical size, such that KI=Klcand dKJdc>O, where c is the size of the crack and KIC is the critical stress intensity factor.
The crack tip stress intensity factor is related to the crack size by K,=OYC'~,while the crack velocity, v, is related to Ki by v=v,(Ki /KO)" exp(-Qf /RT). The constant, U,depends on the geometry of the crack. The constants v, and K, are normalization c o n s t a d . For any size crack, the strength or failure stress of the material, S, can be defined in terms of the size of the crack: S= KI&c'R. In case of a stress gradient, the failure stress is defined at the site of the fracture origin. The equation also assumes the material has a flat Rcurve. The failure stress can be measured in an inert environment so that crack growth is negligible prior to failure. The equation relating failure stress, S, and crack size, c, is important since it shows that as the crack gets longer the failure stress decreases.
These fracture mechanics relationships can be combined to give an equation relating the failure stress, S, at any time, t, to the applied stress, 0, in the asreceived part': t/
S"-' =Sin-' -[exp(Q f / R T ) ] ( l / B )j o " ( t ) d t , (1)
c-2
0
where B = 2K ,"/[(n - 2 ) v , Y K ] is constant for a given environment. The equation is integrated from the time at which load is first applied to the time for failure, q. The equation demonstrates that failure stress can be determined if the time dependence of the applied stress, o, is known. Equation 1 can be used to determine lifetime for constant load, constant loading rate and cyclic load (of various forms). It can also be used to determine constants, n and B, which are needed for the prediction of lifetime. The remainder of this section discusses the use of equation 1 for determining the constant n and for evaluating the statistical uncertainty in predicting lifetime. The time to failure, q, under a constant load is: t f = Ba-"S1-2exp(Qf/RT), (2) where Q is the applied stress and Si is the initial failure stress. The equation also suggests a method of obtaining these constants. Applying a range of stresses to a number of tensile or flexure specimens the times to failure are measured. Then, expressing equation 2 in logarithmic form, a plot of the failure time as a function of the applied stress yields a straight line with a slope of -n and an intercept of BS:-2exp(Qf/RT). An estimate of n for AS800 (AlliedSignal, Torrance, CA) is given in Fig. 1, where a constant stress was applied to tensile specimens at 900 "C and stresses of 450 MPa, 475 MPa and 500 MPa. The straight line was fitted to the median value of each set of stress data, yielding a value for n of 75. A problem with the use of static loading for the determination of n is the enormous scatter in lifetime often observed for silicon nitride', about 6 orders of magnitude in Fig. 1. To reduce the effect of scatter, median values of the data are usually fitted to obtain n. An advantage of the static fatigue approach is that real flaws and cracks grow at the rates likely to be experienced in service failure conditions. 1000
2s
1
+
100 c &
7
r ; (
m
10
;
1 ;
9
400
500
600
A p p bed Stress, MPa * The use of commercial designations is for purposes of identification only and does not imply endorsement by the National Institute of Standards and Technology.
110
Fig. 1 Static fatigue of AS800 tested at 900 T I 3 .
The most popular approach to obtaining the crack growth sensitivity, n, is the so-called dynamic fatigue technique, in which specimens are loaded to failure at several constant-stressing rates. This technique has the advantage over the constant load technique in that experiments are relatively rapid and scatter in the failure stress is relatively small. Furthermore, there are no run-outs or loading failures. Substituting CJ = bt into equation 1, yields an equation for the time to failure under constant loading rate: t f = (n -2)Ba-"S;-*exp(Qf/RT).
(3)
But for the coefficient, 12-2, equation 2 and 3 are the same. Thus, the failure time for the same breaking stress in a constant stressing rate experiment is substantially shorter than in a constant stress experiment. However, the use of short-term data to represent a long-term process can be questionable if long-term processes occur that are undetected by the short-term test. Substituting t = d 6 into equation 3, yields the equation that represents failure stress as a function of stressing rate:
Weibull analyses of the failure times were obtained for the data shown in Fig. 1 and straight lines were fitted to the data to obtain m', Fig. 2.13 Values for m' are given in Table 1 for AS800 (Honeywell, Torrance, CA) and SN88 (NGK Insulator, Nagoya, Japan). The Weibull modulus for AS800 at room temperature was 17, that of SN88 was 9.7.14 Using rn'=m/(n-2), values of n are found to range from 60 to 81, for the data shown in Fig. 1. These compare favorably with the value for n of 75 obtained by a linear regression of the median values of the times to failure. The parameter, n, was determined to be 55 by constant stressing rate experiments (dynamic fatigue). Values of n at the 0.05 and 0.95 confidence level were determined for the Weibull in time method (no5=37; n95=83. 900 "C, 475 MPa) and for the dynamic fatigue method (no5=40; n95=88). There is no statistical difference between values of n obtained by the techniques.
I
h
m'4.30
+ From a logarithmic plot of applied stress as a function of b , the slope of the line fitted to the stressing rate data is l/(n+l). Another approach to obtain n is to statistically analyze the distribution of failure times; we refer to this technique as the Weibull in Time technique. Because failure initiates from pre-existing flaws, the distribution of these flaws controls both the times to failure, q, and the distribution of failure stresses, Si,when the failure stress is measured in an inert environment. Failure stresses are often expressed in terms of a two-parameter Weibull distribution":
where rn is the Weibull modulus and So is the characteristic strength. Substituting equation 5 into equation 2, gives the failure time in terms of failure probability: tf = t o [ l n ( l - P i (6) The two parameters, m' and to, are given by rn'=rn/(n-2) and t0=BS,"'dnexp(Qf RT), respectively. Equation 6 will be used later to obtain constant times to failure at a fixed probability level. These are the curves that are needed for component design. Normally, a failure probability of one part in ten thousand is the minimum acceptable level for the design of industrial gas turbines. The crack growth parameter, n, can now be determined if both the failure stress and the time to failure are expressed in terms of Weibull distributions. Once rn and rn' have been determined, n can be calculated from the relationship rn'=ml(n-2)6v*, ' I v12.
I
I
I
I
I
I
I
I
In (Time to Failu re, h ) Fig. 2: Weibull diagram, failure times for AS800, 900 "C, 475 MPa13. Table 1: Static fatigue Data on AS800 and SN88I3 T "C 900
700
Stress, MPa 500 475 450 425-475 0.05 - 50 MPdS 500
900 900
525 500
AS800 m' n 0.24 75 0.30 60 0.22 81 75
0.19
55
Method Weibull in Time Weibull in Time Weibull in Time Equation2 Equation4
93
Weibull in Time
61
Weibull in Time Weibull in Time
SN88 0.22 0.18
55
FAILURE BY CREEP Two boundary conditions exist for the prediction of lifetime by creep. The first is that the component breaks during the creep process. This is an important mode of failure in ceramic materials, which are inherently much more brittle that metallic materials. The second mode of failure is that the dimensions of part exceed a critical strain during the creep process, so that it becomes too large for the particular application. A complete solution of failure by creep-rupture requires knowledge of the failure mechanisms and the relation between the failure mechanism and the creep rate. Such knowledge of failure modes is not available
111
for structural ceramics, so that a more empirical approach to lifetime prediction has to be adopted. One approach uses the Monkman-Grant equation3, which correlates secondary creep rate, &, with the time-tofailure. The time-to-failure is expressed as a power law function of the creep rate. For silicon nitride, the timeto-failure can also depend on the test temperat~re'~: 109
10-10
where Qm is the temperature dependence of the Monkman-Grant plot and tl and k are constants. Equation 7 fits creep rupture data for silicon nitride exceptionally well, as can be seen from the data for SN88 in Fig. 3.16 To obtain failure time as a function of applied stress and temperature it is necessary to characterize the creep rate as a function of these parameters. This has been done recently by Luecke and Wiederh~rn'~, who suggested that tensile creep in silicon nitride is a
100
10
5M)
Applied Stress, MPa
Fig. 4: Creep data for SN88 fitted to equation 8. of two over the entire range of measurements. The deviation of the data from the expected lifetime is systematic, positive at both ends of the plot and negative in the middle, which suggests that equation 9 does not capture the entire functionality of the lifetime.
9.P Bc a
t=f,% exdQJRT)
b1.14
loo
Qm=137kJ/mol 10'0
I
10
'
'
"'""1
100
" ' I '
loo0
loo00
Fig. 3: Creep rupture data obtained on SN88. The curves are fits of the data to equation 7. cavitation-controlled process. They demonstrated a strong connection between creep strain and cavity volume fraction. Based on this finding, they derived a relation between applied stress, 0, and the secondary creep rate, i :
.
.
= E , . u .exp(a. a).
exp( -Q,/RT),
(8)
where, Q, is the temperature dependence of the creep rate. An excellent fit of equation 8 to an extensive set of data was obtained for SN88, Fig 418. Similar fits have been obtained on other sets of silicon nitride16*17. By substituting the creep rate in equation 8 into equation 7, an equation is obtained that relates the creep-rupture lifetime to stress and temperature:
A comparison of the predictions of equation 9 with actual measurements, Fig. 5, shows that the predicted and measured values of the lifetimes lie within a factor
112
' """"
'
10
""""
100
' '"*""
' """"
1000
' 'a'LyJ
10000
100000
0bserv ed , h
Fig. 5: Lifetime prediction of equation 9, predicted versus observed times, SN88.
Failure Time, h
E
'
1
' """"
""""
Equations 2 and 9 both give relations between lifetime, temperature and stress: equation 2 for crack growth as the life-limiting mechanism, equation 9 for creep rupture as the life limiting mechanism. These equations can be used to create a fracture mechanism 20, For map, which identifies regions of safe de~ign''~ each equation, the allowable tensile stresses, the solid curves in Fig. 6, are plotted as a function of temperature for a fixed time to failure. Separate curves are plotted for 100 h, 1,000 h, and 10,OOO h. Each failure mechanism is marked on the diagram. Within the envelope for each time curve, allowable temperatures and stresses for safe design are indicated.
STRAIN LIMITED LIFETIME At elevated temperatures, lifetime can also be limited by the maximum allowable strain in a component. For normal design criteria, the strain limit can be as low as 0.5% strain, which means that even a fairly brittle material, such as silicon nitride, can be strain-limited. The problem of defining the strain to failure for silicon nitride was discussed recently by Wiederhorn et a1.22,who developed a general equation to describe the creep of silicon nitride in the primary
700
- Creep Rupture --_-_0.5% Strain
stress of 100 m a , the allowable temperature has to be decreased by 60°C which is substantial for high temperature operation.
PROBABILISTIC LIFETIME PREDICTIONS
0' 600
'
"
900
"
"
1200
'
1500
Temperature, "C
Fig. 6: Fracture mechanism map for SN88. and secondary stages of creep. Because the critical strain of 0.5% was reached towards the beginning of the secondary creep stage, tertiary creep was ignored in this treatment. The final equation had three temperature and stress dependent parameters, the values of which were obtained by a fit to all of the creep data. The form of the equation was: E = E .t + E M p [ I - exp(t/r )], (10) where E is the secondary creep rate given by equation 8; eMp, is the maximum primary creep strain; z is a relaxation time, which can be expressed as a power function of the minimum creep rate (in a manner similar to that of the Monkman-Grant plot). Each of these constants can be obtained from the creep data22. A comparison between the calculated and measured times to reach 0.5% strain is shown in Fig. 7. In all cases, the predicted lifetime lies within 50% of the measured lifetime. From the deviation, it is possible to determine the standard deviation of the individual data points, SD, which can then be used to give the probability for short term failures.
The curves relating temperature to applied stress shown in the fracture mechanism map in Fig. 6 were calculated for either the mean (creep rupture and strain) or median (crack growth) failure times, which are measures of the central tendency for failure. To establish a safe engineering stress or temperature, it is not the central failure time, but the failure time at a low probability that is important. One part in ten thousand failures is normally considered to be minimum acceptable number for commercial ~ e r v i c e ' ~ . The allowable stress and temperature for this minimum number is easily determined from the equations presented in this paper. For a constant strain to failure, the difference between the predicted and measured lifetime is used to determine the standard deviation for the scatter in the lifetime of components. The standard deviation, SD, of the difference between N measured and calculated lifetimes, At,, is given by:
r
2- 300.0
600 -
.& b
3
100.0
71/2
D=OS
500
A
61 2
400
rr 0
w
300 200
v)
.5%
100
-
" 600
Logarithm of Measured Timeto 0 . 5 % Stram,h
900
1200
1500
Tempemhue, C
Fig. 7: Lifetime prediction for 0.5% strain to failure, predicted versus measured time.
Fig 8: Fracture mechanism map for SN88, 10,000 h lifetime. The figure contains failure probabilities of both 0.5 and 0.0001.
Once the creep strain is expressed as a function of stress, temperature, and time, equations relating stress and temperature can be obtained for 0.5% strain and a fixed time. Fig. 6 shows the result of this calculation for 100 h, 1,000 h and 10,OOO h as dashed lines. For the 0.5% strain limited failure, the allowable stresses and temperatures are significantly less than those obtained for the stress rupture limited failures. For an applied
For subcritical crack growth as a failure mechanism, equation 6 can be used directly to determine the allowable stress for a given lifetime. Substituting 0.0001 into this equation, and, expanding t , in terms of stress and temperature, yields an equation between failure time, applied stress and temperature. The relation between stress and temperature can now be
113
obtained by fixing the failure time to a specific value, say 10,OOO h. The relation exhibits a strong dependence on the Weibull modulus. Curves of constant time to failure, 10,OOO h, were determined for both the crack growth failure regime and the creep regime, assuming 0.5% strain to failure for the creep regime and actual rupture time for the crack growth regime, Fig 8. In the creep regime, the allowable stress and temperature for the 10,OOO h lifetime are degraded by a modest amount. For 100 MPa applied stress, the decrease in temperature needed to assure a failure probability of less than 0.0001 is about 25 "C. For SN88 the maximum temperature allowed for a strain of less than 0.5% and at a probability of less than O.OOO1 is less that 1300 "C. The effect of the Weibull modulus, m, on the allowable stress is very strong, as can be seen from Fig. 8, where 10,000 h curves for a failure probability of 0.0001 have been drawn for Weibull modulus of 10 and 20. These values cover the range usually reported for silicon nitride. At a temperature of about 900 "C, a Weibull modulus of 20 reduces the allowable stress from the median value of about 450 MPa to about 350 MPa. For a Weibull Modulus of 10, the allowable stress is reduced to less than 200 MPa. These results demonstrate how important m is to component performance, since an allowable stress of less than 200 MPa lies very close to the design stresses in some ceramic components. The higher value of m yields a much greater flexibility for safe design.
REFERENCES (1) M. Van Roode, Design and Testing of Ceramic Components for Industrial Gas Turbines, in the International Symposium, Ceramic Materials and Components f o r Engines, June 19-23,2OOO,Goslar, Germany.. (2) N.J. Tighe and S.M. Wiederhorn, "Effects of Oxidation on the Reliability of Silicon Nitride," pp. 403-23 in Fracture Mechanics of Ceramics 5, Surface Flaws, Statistics, and Microcracking, R.C. Bradt, A.G. Evans, D.P.H. Hasselman, and F.F. Lange, eds., Plenum Press, New York (1983). (3) F.C. Monkman and N.J. Grant, An Empirical Relationship between Rupture Life and Minimum Creep Rate in Creep-Rupture Tests, Proc. ASTM 56 (1956) 593-620. (4) M.K. Ferber and M.G. Jenkins, Evaluation of the Elevated-Temperature Mechanical Reliability of a HIP-ed Silicon Nitride, J. Am. Ceram. SOC.,75[9] (1992) 2453-62. (5) S.M. Wiederhorn, "Subcritical Crack Growth in Ceramics," pp. 613-646, in Fracture Mechanics of Ceramics, Vol. 2, eds., R.C. Bradt, D.P.H. Hasselman and F. F. Lange, (Plenum Pub. Co., New York), 1974. (6) D. Munz and T. Fett, Ceramics, Mechanical Properties, Failure Behaviour, Materials Selection, Springer (1999). (7) J.B. Wachtman, Mechanical Properties of Ceramics, John Wiley and Sons, New York (1996).
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(8) S.M. Wiederhorn and E.R. Fuller, Jr., Structural Reliability of Ceramic Materials, Muter. Sci. Eng. 71 (1985) 169-186. (9) G.D. Quinn, "Review of Static Fatigue in Silicon Nitride and Silicon Carbide," Ceram. Eng. Sci. Proc. 3[1-21 (1982). (10 ) W.Weibull, A Statistical Distribution Function of Wide Applicability, J. Appl. Mech. 8 (1951) 293-297. (11) A. Paluzny and P.F. Nicholls, Predicting TimeDependent Reliability of Ceramic Rotors, pp. 95112 in Ceramics f o r High Performance Applications - 21, J.J. Burke, E.N. Lenoe and R.N. Katz, eds. Brook Hill Publishing Co. Chestnut Hill, MA (1978). (12) J.D. Helfinstine, Adding Static and Dynamic Fatigue Effects Directly to the Weibull Distribution, J. Am. Ceram. SOC. 63[1-21 (1980) 113. Correction 63[11-121 (1980) 716. (13) S.M. Wiederhorn and R.F. Krause, Jr., Static Fatigue of Silicon Nitride at Intermediate Temperature, to be published. (14) S.M. Wiederhorn, R.F. Krause, Jr., W.E. Luecke, J.D. French and B.J. Hockey, Ceramics for Gas Turbines Program: Final Report to the General Electric Company, December 1996. (15) W.E. Luecke, S.M. Wiederhorn, B.J. Hockey, R.F. Krause, Jr. and G.G. Long, "Cavitation Contributes Substantially to Tensile Creep in Silicon Nitride," J. Am Ceram. SOC.,78 [8] (1995) 2085-96. (16) R.F. Krause, Jr., W.E. Luecke Jr., J.D. French, B.J. Hockey and S.M. Wiederhorn, Tensile Creep and Rupture of Silicon Nitride, J. Am. Ceram. SOC., 82[5] (1999) 1233-41. (17) W.E. Luecke and S.M. Wiederhorn, A New Model for Tensile Creep of Silicon Nitride, J. Am. Ceram. SOC. 82[101 (1999) 2769-78. (18) K.J. Yoon, S.M. Wiederhorn and W.E. Luecke, A Comparison of Tensile and Compressive Creep Behavior in Silicon Nitride," J. Am. Ceram. SOC. 83 [8]2017-22 (2000). (19) M. Matsui, Y. Ishida, T. Soma and I. Oda, Ceramic Turbocharger Rotor Design Considering Long Term Durability, pp. 1043-50 in Ceramic Materials and Components for Engines, W. Bunk and H. Hausner eds., Verlag Deutsche Keramische Gesellschaft, D-5340 Bad Honnef, Germany (1986). (20) G.D. Quinn, "Fracture Mechanism Maps for Silicon Nitride," pp. 931-9 in Ceramic Materials and Components for Engines, W. Bunk and H. Hausner eds., Verlag Deutsche Keramische Gesellschaft, D5340 Bad Honnef, Germany (1986). (21) G.D. Quinn, Fracture Mechanism Maps for Advanced Structural Ceramics, Part 1, Methodology and Hot Pressed Silicon Nitride Results, J. Muter. Sci., 25 (1990) 4361-4376. (22) S.M. Wiederhorn, W.E. Luecke and R.F. Krause, Jr., A Strain-Based Methodology for High Temperature Lifetime Prediction, Ceram. Eng. Sci. Proc. 19[4] (1998) 65-78.
STANDARDISING MEASUREMENT AND TEST METHODS FOR ADVANCED TECHNICAL CERAMICS R. Morrell National Physical Laboratory, Teddington, Middlesex, UK, T W l l OLW
ABSTRACT The increasing diversity of engineering applications for advanced technical ceramics of all types means that there is increasing expectation placed on suppliers for assurance of quality and performance against specification. In addition, specifications are becoming more stringent as the applications become more critical, and this means that standardisation becomes essential for effective trade. The industry has supported standardisation of basic testing procedures which are purpose-designed for advanced ceramics, at first notably in the UK, Japan, USA, France and Germany, and since 1989, internationally in CEN and ISO. The emphasis has been on providing high-quality methods appropriate to the wide range of product types that are available, or which may appear commercially in the future. In addition, application-based standards incorporating specifications for materials and device performance are also appearing, notably in the electrical and biomedical fields. This paper briefly reviews the status of the standardisation programmes, and the effect they are having on the ability better to specify requirements, to control quality, and to provide reliable data. In this paper some specific examples of how the development process has occurred, and some of the outstanding problem issues relevant to ceramics for engines are discussed, including: 0 strength and toughness testing; 0 microstructural assessment; 0 ‘damage resistance’; 0 surface roughness. The value of international cooperation is emphasised, such as through the Versailles Agreement on Advanced Materials and Standards (VAMAS) [l], the European Structural Integrity Society (ESIS), and the International Energy Agency (IEA) programme, in identifjing and solving some of the problems, and in providing a platform for validation of test procedures and provision of statements of confidence in results from the test methods.
material development - consistent measures of what is achieved data acquisition - baseline property information for design purposes data exchange - making data sheets believable by others, also laboratory testing accreditation material selection - for property/performance comparison quality assurance - does the product meet a specification?
PROGRESS Particularly for thermomechanical applications, the properties of critical interest are wide-ranging and different from those typically used for traditional electrotechnical ceramics, with a strong emphasis on understanding mechanical properties. Thus a whole new suite of test procedures has been required, and huge steps have been taken over the last decade to lay the basis for a testing methodology that is appropriate to the wide diversity of material types to be addressed, including monolithic ceramics, composites, and coatings. As the first stage, the effort has concentrated on testing materials, rather than components, although some component standards are being produced in specific areas, such as orthopaedics. Independent national work commenced in a number of countries during the 1980s, including Japan, UK, USA, Germany and France. Unified European action took place from 1989, and an IS0 committee was established in 1994 to promote global activity. Table 1 provides an indication of the scale of the work achieved to date as well as future targets. Full details of published documents are available fiom standards bodies and are not covered in this paper
Table 1 - Advanced technical ceramic material test method standards production (mid 2000) (approximate data, excludes electroceramics) Published
In development
CEN TC 184
70
15
ASTM C28
30
8
JIS IS0 TC206
40 2
?
Organisation
THE FUNCTION OF STANDARDS Standardised ways of approaching testing and data acquisition are key to the effective use of and trade in any material. Once the research phase has matured, the target is the effective exploitation of the materials as real products. Standardised procedures are helpful for:
20
115
FORMULATING STANDARDS The standardisation process is not a trivial one when done properly. It takes considerable time and effort, as well as background knowledge and consensus-building abilities among the committee members, to produce a document which provides a usable test method. The document is often a balancing act between technical exactness and practicality. Make the method too complicated or too expensive, or too specific to a particular type of material, and it will have limited take-up by industry and the testing community. Modem standards making requires more than a simple test procedure. It needs clear definition of: the applicability of the method to a particular range of materials or material types. Not all test procedures work effectively on all materials, so it is often difficult to be certain of this point. the test-piece preparation procedure, especially when the results can be dependent on the procedure employed; the test method itself, particularly any variations in procedure that are required for particular materials features or responses; what to do if there is a problem; clear acceptance criteria for the result; any known ‘interferences’ which can lead to misleading or erroneous results; the required levels of precision of measurement of each parameter, which are practically and economically realisable; the anticipated accuracy of any individual result as a deviation ffom the ‘true’ result; the anticipated within-lab repeatability and between-lab reproducibility on the same material; this can often be in the form of a summary of an interlaboratory study of the method. Such information is particularly helphl when assessing whether differences in apparent properties are truly significant or not. In addition, there is an increasing expectation that test reports incorporate an error assessment to guide the user in providing an overall confidence statement on the results.
TECHNICAL BACKGROUND TO SOME SPECIFIC EXAMPLES FLEXURAL STRENGTH TESTS Prior to 1990 there were numerous test-piece crosssections and spans, as well as a variety of test-jig functionalities in three- and four-point bending. The situation has been reviewed in [2]. The current procedures in ASTM C116 1 and CEN EN 843- 1 for monolithic ceramics were developed over more than fifteen years of national and international collaboration with improving rationalisation of different national views. The details of the Japanese equivalent, JIS R1601 are different (30 mm rather than 40 mm span,
116
polished rather than machined surfaces). Further rationalisation is now taking place with the publication of IS0 14704, in which the ASTW CEN approach is preferred. Despite the fact that flexural strength is strictly not a deterministic property, owing to an inability to control the shapes and sizes of ffacture initiating features, it is often used as the key mechanical yardstick for a material. Consequently, in creating formal standardised test procedures, a large number of factors needed to be taken into account, including: the relevance of using material fabricated solely for testing relative to material fabricated for the enduse; this appears to be particularly the case for materials such as silicon nitride; the method of machining the surfaces and edges of the test-pieces; the test geometry, including test-piece dimensions and test spans; the rate of applying a force; the test environment employed; the method of applying the force to the test-piece. An optimisation exercise by the US Army during the early 1980s [3] was conducted to minimise the errors introduced by the method of loading the testpiece and its geometry, starting on the premise that the simple ‘thin-beam’ flexure equation should be retained for calculation of strength. A narrow window on the span-to-depth and span-to-width ratios of the test-piece was found in which errors could be maintained individually at less than 1%, with a total overall error of less than 2%, typical of the level needed for confidence in a test result. The accuracies of definition of loading geometry were also obtained, typically 0.1 mm being required. Risks of factors such as twist, either in the test-piece itself or in the manufacture of the test-jig could be reduced by articulation of the loading rollers. Free-rolling rollers, not constrained rods or knife-edges are also required to reduce or eliminate fiiction effects. These and further considerations were built into the prototype method described in MIL-STD-1942, which subsequently became ASTM C I 161. Collaborations through the International Energy Agency programmes involving interlaboratory comparisons encouraged the same format to be employed in Europe, resulting finally in EN843-1. The Japanese standard JIS R1601 defined in 1981 does not take all these factors into account, but it is to be hoped that the formalising of the IS0 standard will bring all countries into line. The importance of this work cannot be overemphasised. By adhering to the geometrical and h c t i o n a l criteria within the standards, it should be possible to see the true effects of changing material formulation, fabrication route, or machining condition, rather than being misled by extraneous effects within a poorly controlled test.
FRACTURE TOUGHNESS Although fiacture toughness is not a parameter that can be used for engineering design in brittle materials, it is an essential measure of material behaviour for a variety of applications since it controls thermal shock resistance, abrasion and erosion resistance, and in some measure local impact resistance. Compared with metallic materials, the principal problems in testing are:
focusing on methods for small test-pieces, and hence small cost in materials terms; difficulties in creating a clean, stress-fiee pre-crack for the applicability of well-defined fracture mechanical calculations; small crack opening displacements, so that crack face tractions can exist in long cracks, giving Rcurve behaviour; the simplest methods have the greatest errors due to poor calibrations.
Table 2 - Comparative attributes of bar test-piece methods for fracture toughness
Criterion Confidence in calibration
Method (see key)
’
~
Good
SCF
SEW2
IF
IS
Good
Good
Poor
Poor
Relative ranking of materials
Yes
Yes
I
Yes-
I
Determining long-crack toughness
Yes
No
I
No
I
No
Determining R-curve effects
Yes
No
Possiblv
I
No3
I
N O T N O
Determining short-crack toughness
No
No No
Producing a fast-fiacture toughness
Yes
Yes toughness value
I
No
I I
yes
No
I I
PossiblyTYes
No
No
I
I I
No
Yes
No
Possibly
Possibly
Use for fine-grained materials
Yes
Yes
Yes
Yes
Yes
Yes
Use for coarse-grained materials
Yes
Yes
Unlikely
Yes
Unlikely
Unlikely
I
Yes
Yes
Yes
Use for materials of KIc> 6 MPa m1’2
Yes
Yes
Sensitivity to crack growth before fast fiacture
Yes
Method uses a moving crack fiont
(Producing low scatter of results
lUse for materials of K,c < 6 MPa m”2
Sensitivity to notchlpre-crack geometry
Possibly
Suitability for use at elevated
ICost of facilities required for test Standards
I Medium ASTM, JIS ISO=CEN
I
Yes
I
Yes
I
Yes
method
Probably
Yes
Possibly for fine-grained materials
N/A
N/A
Yes
Yes6
Yes
No
Possibly6
Medium
High
Medium
Low
Medium
Medium
High
Low
I Medium
ASTM CEN
I
Medium
I
ASTM I m a f t i n ~ ~ ~JIS I ISO=CEN
I
-
Key: SEPB = single edge pre-cracked beam, CNB = chevron notched beam, SCF = surface crack in flexure, SEVNB = single edge vee-notch beam, IF = indentation kacture, IS = indentation strength, NIA = not appropriate. Footnotes to table: I On a kacture mechanics basis. Quality of experimental results influenced by the nature of the material. Attributes given are appropriate for notch root radii of typically 5 pm or less, i.e. sharpened by honing. Unless used initially to grow a crack fiom the notch. If the notch is considered to have a short flaw at its tip, R-curve effects are avoided. 5 There is an increasing likelihood of poor pre-crack generation with increasing toughness; the upper limit may be at K,< = 5 MPa m”*, but there are reports of successful and valid results obtained kom tougher materials. 6 The tests can usually be performed up to a temperature at which the material begins to soften or crack-heal, or where oxidation commences in the environment used.
117
In addition, not all materials can be tested by a given method; each method has a limited range of applicability, especially regarding the generation of a suitable length of pre-crack. Consequently a range of methods is needed, and these have needed to be developed to a reliable state before standardisation. Table 2 lists a variety of bar test-piece methods and summarises some of the advantages and disadvantages (being incorporated into [2]). Those which have been selected so far as formally standardised methods are also indicated. So far, plate and short rod methods have been ruled out. It is clear kom Table 2 that this field is complex and needs carefkl thought. A key concern is that different methods will not give equivalent results, primarily because they employ different lengths of pre-crack at the point the measurement is made. However, in finegrained materials which show no R-curve effects or significant subcritical crack growth, equivalence has been demonstrated. Several round robins have been conducted on the methods under the auspices of VAMAS Technical Working Area 3 and ESIS Task Group 6 in order to determine practical capabilities as well as reproducibilities of testing in order to validate procedures before standardisation.
Set of test-pieces
Razor blade with diamnondpaste
(b)
Figure 1 (a) SEVNB method schematic, and (b) a 3 pm tip radius notch produced in a silicon nitride.
118
Tests based on simple sawn notches have not been included because of substantial evidence that a notch is not equivalent to a sharp crack for many materials, particularly the harder, tougher varieties. However, producing pre-cracks is a skilled task; the SEPB method requires use of a special compression anvil, and is not always successfkl, and the SCF method requires fiactographic skills. To overcome this problem, recent research has been aimed at refining a method in which a sawn notch is honed with diamond paste to have a tip with an extremely small radius, typically 2 - 10 pm, which is a much better approximation to a sharp crack than a sawn notch. Figure 1 shows a schematic of the technique and a typical notch produced by it. This S E W method has recently been evaluated through an ESIS/VAMAS round robin [ 5 ] with excellent reproducibility, and it has been agreed to move the method towards a formal standard. MICROSTRUCTURALPARAMETERS Quantifj4ng microstructure is becoming a key element of specifications, and since microstructure is critically important in advanced technical ceramics in defining performance of the product, efforts have been made in CEN to write standardised procedures tailored for ceramics but based on the general principles laid down in ASTM E l 12 for metals. So far, a standard has been published on grain size determination by the manual line intercept method (ENV 623-3), and one on phase volume fi-action by a manual grid method has been approved for publication (prENV 623-5). In both cases the procedures have been tested out in CENNAMAS round robins in order to determine reproducibility [6,7]. For grain size, an update has recently been agreed which additionally incorporates grain size distribution measurement. Other parameters that might be considered for the future are grain shape and phase dispersion. The increasingly widespread use of automatic or semi-automatic image analysis (AIA) to speed up the analysis of micrographs has posed something of a dilemma for standardisation because there are no basic standards for systems or software. Digitisation of images permits a range of analysis techniques different fi-om those typical of manual methods, for example those based on area or nearest neighbour distance measurement. In order to take advantage of the power of AIA considerable additional care needs to be taken in the preparation of suitable images so that the software can correctly interpret them, e.g. that revealed grain boundaries are continuous and narrow, but well defined, and that each phase has uniform brightness with grey level discrimination fiom other phases. In the CENNAMAS round robin [7], images in which these conditions were not upheld resulted in much larger scatter for the AIA assessment of volume fiaction than for the manual method. The human eye still remains much better at compensating for poor micrograph preparation than a computer, despite considerable advances in pattern recognition programming. Manual
intervention to delineate for example, poorly defined grain boundaries or phase areas, is possible, but slow. Because of the proprietary nature of most AIA software, it may be necessary to ensure conformance with recognised analysis methods by use of digital reference micrographs.
PROBLEM AREAS As the experimental evaluation of materials widens with increasing attempts to substitute ceramics for metallic materials in extreme environments, issues of how to characterise for such uses arise. In many cases testing to simulate the application can be straightforward, but it is usually inappropriate to attempt to standardise such procedures, because specific test conditions tend to have little general value. For example in the field of corrosionlerosionloxidation, guidelines on the basic steps in conducting tests and the potential pitfalls to be guarded against are all that is currently practical. Two fiu-ther areas which continue to pose problems for standardising are ‘damage resitance’ and roughness determination, which are discussed in further detail below.
‘DAMAGE RESISTANCE’ Advanced technical ceramics, particularly in the engineering field, are widely employed for their hardness and dimensional stability. To perform as engineering materials they must withstand severe conditions such as thermal shock, impulsive loading, localised loading, edge damage, erosion, abrasion or sliding wear. It has proved difficult to standardise assessment methods for many of these performance attributes, primarily because the results are particularly sensitive to the details of the test method. In thermal shock, results are dependent on parameters such as emissivity, heat transfer coefficient and geometry, which are difficult to model exactly in a test. In fact it is often opined that a component trial is by far the safest way of assessing possibilities. In impulsive loading the impacting energy is dissipated in a manner dependent on geometry and support. In localised loading, some progress is being made using Hertzian indentation as a model [8]. Resistance to edge chipping is becoming recognised as being limited by material toughness (GI=) [9] and although quasistatic chipping is a simple test to perform, the mathematical basis for it remains uncertain. In abrasive and erosive wear, test results are system dependent, and rank different materials in different ways (e.g. Figure 3). However, recent work at CRIBC in Belgium [lo] has shown that experimentally based material rankings can vary with mode of material removal. The development of correlation equations for different modes of contact
may offer a tool for material selection which is independent of the actual test apparatus, as long as the mechanism of removal is appropriately modelled by the test. In sliding wear, critical evaluation of test methods reveals that machines are often not relevant to the application and tests are performed under conditions which do not give the same mechanism as would be seen in use. Most ceramics suffer fiom breakaway wear conditions above a critical speedload combination, so attention needs to be focused on determining these conditions in order to provide information on usability and a useful materials comparison. Overall, for these categories of test, it is unlikely that universally accepted standardised methods will be developed in the shodmedium term.
i
A
a
1
Hardness, HVO.l
Figure 3 ASTM G65 test results on various ceramics using a standard sand abrasive and different wheel surfaces, indicating different rankings
SURFACE ROUGHNESS IS0 standards lay a basis for the measurement of roughness using traditional stylus machines, but these are designed primarily for metallic materials. Ceramics have very short range roughness, and this has posed problems of reproducibility for many years. EN 623-4 restricts the flexibility offered by the IS0 standards in an attempt to improve reproducibility, but a recent EC sponsored round robin indicated that there is still considerable variability between different machines. It is believed that the origins lie in: stylus high-fiequency dynamics - affects the ability of the stylus to follow the surface at the typical traverse speed of 0.5 mm/s; inappropriate digitisation length - if too short this acts as an inadvertent short wavelength filter, possibly removing short-range information. Further investigation of this problem is underway at NPL at this time.
119
CONCLUDING REMARKS Good progress has been made in developing appropriate standards for ceramics for engineering, but this has only been achieved by the research effort dedicated to solving problems of consistency, accuracy and reproducibility. Much remains to be done, but the work so far has already had a considerable positive impact in the quality of testing and data to the benefit of the manufacturing and user communities.
ACKNOWLEDGEMENTS The author’s involvement in this work is supported by the UK Department of Trade and Industry through its materials metrology programme.
REFERENCES G. Q u i , VAMAS after twelve, h e r . Ceram. Soc. Bull., 1999, 78(7), 7tG83. G.D. Quinn, R Morrell, Design data for engineering ceramics: a review of the flexure test, J. Amer. Ceram. SOC.1991,74(9), 2037-66. F.I. Baratta, G.D. Quinn, W.T. Matthews, Errors associated with flexure testing of brittle materials, Technical report TR 87-35, July 1987, US Army Materials Technology Laboratory, Watertown, USA. prEN XXXX-I: Advanced technical ceramics monolithic ceramics - test methods for determination of apparent fiacture toughness Part 1: guide to test method selection (to be published). J. Kubler, Fracture toughness testing using the SEVNB method; round robin, VAMAS Report No. 37, EMPA, Switzerland, 1999. L.J.M.G. Dortmans, R Morrell, G. De With, Round robin on grain size measurement for advanced technical ceramics, J. Eur. Ceram. Soc., 1993,12(3), 205-13. M. Hendrix, E.G. Bennett, R Morrell, L.J.M.G. Dortmans, G. De With, CENNAMAS phase volume fraction round robin, Brit. Ceram. Trans., 1998,76(6), 293-6 S.G. Roberts, Hertzian testing of ceramics, Brit. Ceram. Proc.2000, Wl), 31-8. R Morrell, N.J. McCormick, Edge chipping as an indicator of toughness, PaoRim 2 Conference, Cairns, Australia, July 1996, Int. Ceram. Monographs, edited by P. Walls, C. Sorrell, A Ruys, Australasian Ceramic Society. (10) D.Beugnies, P Descamps, J T i r l q , et a/., Performance standards: a new concept for advanced ceamics wear performance, Sixth Euroceramics, Brit. 1999,60(1), 21 1-2. C a m . ROC.
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0 Crown Copyright 2000, reproduced by permission of Controller, Her Majesty’s Stationery Office
INVESTIGATIONS ON THE STABLE CRACK GROWTH OF INDENTATION CRACKS T. Lube Institut fur Struktur- und Funktionskeramik, Montanuniversitat Leoben, Peter-Tunner-StraQe5, A-8700 Leoben, Austria
ABSTRACT The stable growth of indentation cracks under an externally applied load offers an practicable possibility to study the fracture properties of ceramics. Under simplifying assumptions - constant geometry factor Y and residual stress parameter - a simple relation can be derived for the dependence of the strength oIs of an indented specimen on the indentation load P: oIsoc P". The exponent n equals 1/3 for materials with constant toughness and is less than 1/3 for materials exhibiting R-curve behavior. For a more realistic approach, the dependence of Y and x on the crack length c has to be considered. In most cases the value of Y for half-elliptical surface cracks in bending specimens decreases with crack extension. A method to calculate the expected decrease is proposed and the results are compared with experimental data. With the help of model calculations it is shown how these changes in Y influence o1s as denoted in common graphic representations. It becomes apparent that it is indispensable to account for the decrease in Y when oIsversus P data are evaluated: the decrease in Y may lead to exponents n < 1/3 even for materials with constant toughness.
wedged open by a residual stress field after removal of the indenter. From the assumption that the residual stress intensity factor K,, for a crack co produced by a load P is equal to the fracture toughness K,, it is possible to determine K, [7]. The residual stress can be written as:
x
The parameter x is a measure of the strength of the residual stresses and is usually obtained by calibration with a material with know fracture toughness. If a specimen containing such an indentation crack is additionally loaded by an external stress o,,(e.g. a bending stress) a stress intensity factor Kapplis caused. The total stress intensity K,, is given as
K , , , = K , + K ~ ~ I = PX ~ + o , Y J F '
where the geometry factor Y depends on details of the crack shape and the loading geometry. As K,,, decreases initially more strongly than Kappl increases with increasing c, an indentation crack can grow stably for a certain distance before failure. Applying the conditions for equilibrium,
Ktot = Kc and for stable crack extension, 3
INTRODUCTION
(2)
(3)
(4)
Indentation methods are widely used to determine the fracture properties of ceramics. The inherently stable growth of indentation cracks subjected to an external load and the presumably easy evaluation makes them especially suitable for the determination of R-curves. A variety of techniques has been established for this purpose [l-61. The differences in these techniques are mainly founded in the way the unknown geometry factor and residual stress parameter are determined. A review of these techniques shows that in some cases empirical corrections have to be introduced to obtain reasonable results. The question arises if the correlations that determine the behavior of indentations crack in such a test are maybe much more complex and interconnected than assumed.
THEORETICAL BACKGROUND Surface cracks produced by sharp indenters like Vickers pyramids are usually described as half penny cracks. Due to the irreversible plastic deformation and damage beneath the indenter point, these cracks are
to eq. (2) allows the calculation of the stress ols, the "indentation strength", and critical crack length c,, at failure:
Plotting the experimentally easily accessible quantity oIsversus the independent variable P (as a measure for the severity of the flaw) in logarithmic coordinates ac-
cording to eq. (5a) yields a straight line. Satisfaction of the P1'3-dependence is usually taken as validation of the assumption that the tested material has a single-valued toughness [8]. R-curve behaviour The inherent stable growth of indentation cracks under an externally applied load offers an interesting possibility to use them for the measurement of R-curves.
121
A combination of eqs. (2) and (3) for a gives:
In principle, with known values of x and Y, eq. (6) allows direct determination of the R-curve by monitoring the crack length c as 0,is increased. Several methods have been proposed to evaluate the parameter x form special graphic representations of eq. (6) [6,9]. In analogy to eq. (5a), eq. (6) can be solved for the maximum stress at failure, o,s(P), by applying eq. (4) and a suitable hnction for K,(c). The analytical analysis of KRAUSE[3] who used a power law function as a representation of the R-curve shows that the presence of an R-curve leads to oIsIXP” with n < 113. Comparable results. i.e. a deviation from the power law behavior with n = 1/3 are achieved by using more physically based relationships for the R-curve [ 10-1 11. The following simplifications are implied in the presented analysis concerning the geometry factor Y and the parameter x: a) The geometry factor Y is assumed to be constant during the stable growth of the crack. It has been recognized that this is not necessarily true for half-penny surface cracks, but it has also been argued that the changes in Yare small enough to be neglected [3]. Only a few experimental data exist on the evolution of the crack shape [9, 12, 131. They indicate that Y decreases during stable crack extension up to as much as 25%. b) The parameter x is also considered to be constant during crack growth. There is some empirical [13, 141 as well as some theoretical [ 151 evidence that increasing the applied load causes the residual stress to decrease. Furthermore, a dependence of x on the indentation load cannot be excluded [16]. It is the aim of this work to investigate the influence of possible changes of Y on the indentation strength and IS-plots with the help of simple model calculations. To back up the theoretical considerations, experiments to determine the evolution of the crack shape of indentation cracks during stable crack growth are performed using a multiple decoration technique.
EXPERIMENTAL Experiments were performed on a commercial gas pressure sintered silicon nitride doped with approx. 3 weight% MgO provided by ESK, Kempten, BRD. Young’s modulus E = 306,s f 0,3 GPa was determined by a resonant beam technique [17], Fracture toughness K, = 6,4 f 0,6 MPadm was measured using the S E W - B method [ 181, hardness is H = 14,6 GPa. Flexural specimens ( 3 ~ 4 x 4 5mm3) were machined from large plates and mirror polished to 1 pm diamond finish on the tensile surface. Vickers indentations were placed in the center of the tensile surface with the indentation diagonals aligned parallel and normal to the specimen’s long axis. Loads between 49 N and 294 N were used. During the indentation process the cracks to be used to study the evolution of the crack geometry were decorated with lead
122
acetate solution according to a procedure described elsewhere [19,20]. The growth of indentation cracks was monitored by interrupted indentation strength (IS-) tests on an electromechanical universal testing machine. The decorated specimens were loaded to different load levels below the indentation strength and then immersed into saturated lead acetate solution to decorate the newly formed portion of the crack which is in each case typical for the applied load level. After drying, these specimens were broken to make an investigation of the crack profile possible. By placing two indentations within the region of constant bending moment it was possible to produce stably grown cracks at relative applied stresses up to o,/o1s= 0,97. Due to the flat trend of the o,(c)-curve in this region, these cracks may still be considerably shorter than cnl.Note that the indentation strength of the decorated specimens was the same as for non-decorated specimens. All experiments were performed in air at room temperature.
MODEL CALCULATIONS Incorporating the possible influences mentioned above into eq. (2) leads to a general expression: (7)
Applying the conditions for equilibrium and stable crack growth, eqs. (3) and (4), allows the calculation of oIsand c,,, and of corresponding pairs of o,and c values for known trends of K,(c), ~ ( c P), , Y(c) and chosen values for P. The starting crack length co is calculated from the condition Kt,(O, CO,P) = K,(CO). Calculations following this scheme were performed for three conditions. Where it was necessary to define material parameters or specimen geometries, values that correspond to the material and specimens described in the experimental part were chosen. i) crack-length independent Y: K,, x and Y are independent of c. K, = 7 MRadm, x = 0,07 (with E and H of the investigated material after ANSTIS[7]) and Y = 1,26. The results for this case should be identical with analytical solutions presented earlier. ii) crack-length dependent Y: the geometry factor decreases with crack length during stable extension. The trend for this change has to be calculated separately, a possible procedure will be proposed below. K, = 7 MPadm is constant. iii) crack-length dependent Y & R-curve: both, the fracture toughness and the geometry factor are considered to depend on crack length. Any possible interactions between the decrease of Y and the increasing fracture toughness are neglected. A pronounced R-curve in the form of an exponential law was assumed (for c in mm and K, in MPadm):
Since there is not much quantitative knowledge about the influence of crack length and indentation load
x,
on the factor calculations.
x was kept constant during the model
The change of crack shape and geometry factor Y during stable crack growth NEWMAN & RAJU[21] proposed an empirical equation to calculate the geometry factor Y for semi-elliptical cracks in bars with rectangular cross section subjected to flexural loading. They defined Y as Y(alc,alh, clb, cp). where c is the crack-half-length at the surface, a the crack depth, h the bar's height, b the bar's half-width and cp the angular coordinate, see fig. 1. For many usual crack geometries (i.e. eo = a0 l co close to eo = 1, crack length small compared to the specimen dimensions) the value Ys= Y(cp= 0) at the specimen's surface is larger that Yo= Y(cp= 7d2) at the deepest point of the crack. Therefore such cracks tend to grow more along the surface of the specimen than depthwise. Their aspect ratio e changes during stable extension. An indirect evidence of such an decrease can for example be found in a paper by BRAm ET. AL. [4]. Typical as-indented cracks with eo P 1 have Y(co)= 1,27 & R A J U formula. The calculated with the NEWMAN value Y(c,) = 0,77 they determined for critical cracks via a calibration procedure can only be obtained for cracks with a lower aspect ratio. Direct experimental evidence of the decrease of the aspect ratio of indentation cracks has recently been given by SGLAVO [9].
1
El.
2c
For a given specimen cross section (i.e. fixed h and b) eq. (10) can be solved for the crack depth a numerically, whereby the crack half-length c is varied within & R A J U equation: the validity range of the NEWMAN cdb I clb 5 0,5, as already proposed [22]. Here the asindented shape of the crack is additionally taken into consideration by allowing changes in the crack depth a only if they lead to a crack extension, a 2 ao. The only experimental data that have to be known for this calculation are crack length co and crack depth a0 of the asindented crack. An example of the change of the crack aspect ratio e with crack length is plotted in fig. 2 for starting cracks with different sizes but the same aspect ratio eo = 1. co= 50 pm co = 200 pm h=2mm h=3mm
075
k l,o 1,5
2.0 c I C"
2,5
38
Fig. 2: Typical change of the crack aspect ratio for crack with co= 50 pm and co = 200 pm and eo = 1 in a flexural specimen with b = 2 mm and h=3mm.
J
D
RESULTS AND DISCUSSION
II
Evolution of crack shape and geometry factor Y An example of a decorated crack that has stably grown for some amount is shown in fig. 3. Fig. I: Schematic of a semi-elliptical indentation crack in a flexural specimen. The change of the crack shape can be determined experimentally but the experimental expense of such a procedure calls for a method to estimate the crack shape with the help of calculations Some possibilities to calculate the evolution of the crack geometry with the help of this formula have been proposed [12, 13, 21, 221 but they do not take into account the starting geometry of the cracks. It is reasonable to assume that the stress intensity is constant along the crack front during stable crack extension. For the investigated indentation cracks this can be expressed as: K,,(cp =O)=K,,(cp = x / 2 ) = K c .
(9)
For a first approximation, and to avoid too many assumptions in the calculation the residual stress intensity will be neglected here, so that eq. (9) reduces to: a a c x
Fig. 3: SEM micrograph of a multiple-decorated crack (P = 294 N) with a length co < c < cnl.The contour of the as-indented crack is indicated by a dashed line The hardness impression can be made out at the tensile surface of the specimen. The dark zone below the impression is not decorated because of the high compressive stresses within this zone. The geometry of the crack immediately after the indentation is indicated by a dashed line (the geometry of as-indented cracks is determined separately [23], the contour indicated here
123
12
" '
'
'
"
'
'
'
"
'
'
_
' .
_
1,1 .u
' 1g
.-0
Y
1,o
0
-
P
0,9
'
'
' .
4 9 ~ 0 98N A 196N V 294N
o
.
.
A mO
0,8
A
0,7 -
0 -
v.r..o
0
A
v
0
c I c,
Fig. 4: Evolution of the crack aspect ratio during stable crack extension for cracks introduced with different loads.
indented geometry of the cracks. According to eq. (10) an equilibrium crack aspect ratio exists for every crack length c. During the first segment of the bi-linear trend, this configuration is not yet reached, the crack depth u is too deep to fulfill the condition imposed by eq. (10). As soon as the equilibrium aspect ratio is reached, the crack grows strictly according to eq. (10). A calculation that does not take into account the as-indented geometry results in the right-hand segment of the bi-linear curve only and will therefore only predict smaller changes in e (or Y) as observed. Because of the excellent results achieved with this simplified approach (which neglects the influence of the residual stresses on crack growth) there seems to be no need to improve the calculation by proceeding according to eq. (9). Since K,, is rapidly decreasing with crack length and also with increasing Kappl[15] it will only influence the begin of the stable crack growth. The influence of K, tends to favor larger crack aspect ratios. The characteristic bend in the curve tends to disappear but the aspect ratio of critical cracks will not change much. Another change in the presented trends for the evolution of the crack shape during stable crack growth may arise if the possible R-curve effect is taken into consideration. The rising fracture toughness will have more influence on portions of the crack that have already grown to a further extent. R-curve effects tend to operate against the influence of the gradient in the bending stress and reduce the tendency towards a decreasing crack aspect ratio.
Calculation of indentation strength data for model cases The indentation strength q s and the critical crack length c, are affected by the crack-length dependences of Y and K,. Fig. 6 shows the relative critical crack length as a function of starting crack length (expressed as indentation load P)for the three investigated cases.
190
r
- - - - _- _
- _- - - -
- _- -..
-c49
Y= Y(c),K, = const. 98
196
294
indentation load P The lines correspond to the calculations, the symbols represent values for Yscalculated with the experimental data from fig. 4. The correspondence is reasonable, the deviations are less than 6% in all cases. The same calculation was repeated for alumina data from literature [9] with an similar good correlation. The bi-linear trend of the curve in figs. 2 and 5 is the main difference that discerns the present calculation of the crack geometry change from previous ones [12, 21, 221. It can be explained by the influence of the as-
124
Fig. 6: Influence of a decreasing geometry factor and a R-curve on the relative length of critical cracks. It is interesting to note that for all but the simplified case cracks reach more than 2,52c0 during the stable extension. Pronounced stable crack growth is favored by the R-curve but also by the decrease of Y. Since the cracks for the case iii) are longer than for the other case, the indentation strength in lower. Critical cracks with
c , > 2,52 co have been reported by several authors [4,
13, 161. Fig. 7 shows the indentation strength plot for all three calculated cases. As expected the data of the simplified case can be found on a straight line with the slope ns = 1/3. Data for the case ii) show a slope n y = 0,26. The data for case iii) can be fitted with n R y = 0,l 1. I
a;” 400
I
t
.-
q.. 0
A
~=const.,~~=const. Y = Y(c), K, = const. Y = Y(c) & Rcurve
I’
20
critical crack length and will not experience the predicted change in shape. For the practical evaluation of indentation strength tests it seems to be indispensable to account for the change of Y with crack length. Since the actual decrease of Y is material and geometry-dependent it should at least be checked if Y does indeed change. The measurement of the indentation strength as a hnction of indentation load alone will not suffice to decide if this is the case. A simple means to do so is offered by the representation of crack growth data proposed by ET AL. [6]. Rearrangement of eq. (7) under DRANSMANN the condition K,, = Kc(c)yields
J
40
60 80 100
200
400
indentationload P
Fig. 7: Indentation strength plot for the three investigated cases. According to the simplified theory a slope less than 1/3 in the indentation strength plot and critical cracks longer than 2,52 co are only expected if the investigated material shows R-curve behavior. The model calculations show that at constant fracture toughness a typical decrease of Y alone (case ii)) may lead to IS-data that are typical for an R-curve material according to the simplified theory - i.e. a slope < 1/3. If the evaluation of such data is carried out using the principles of the simplified theory an R-curve will be deduced even though none exists. It is also confirmed, that if the tested material shows R-curve behavior, neglecting the decrease of Y leads to an overestimation of a possible R-curve effect deduced from IS-experiments. Examples that confirm these conclusions can be found in literature. For a fine (4 pm) grained SSiC, the & LAWN[24], an control material used by PADTURE alumina [9], or on the sapphire presented by COOK ET AL. [lo] the IS-plots exhibit slopes less than n = 113 even though the materials are believed to have no Rcurve. On the other hand there are also examples that perfectly match the simplified approach with a Y independent of crack length, as for instance the glass data in [S]. This is not necessarily a contradiction to the results of the model calculations. How much the crack aspect ratio and the geometry factor change during the stable extension depends strongly on the as-indented geometry and the specimen cross section. Cracks that are close to the equilibrium shape in the as-indented state do not change their shape a lot. Together with the influence of the residual stresses this may lead to a geometry factor that is in fact constant during IS-experiments. If the residual stress intensity K,, is smaller than the indentation stress intensity at full load, i.e. if the lefthand equity in eq. (1) is not fulfilled, then cracks will not grow instantly once the external load is applied [16]. Consequently such cracks will not reach the expected
-Y _r---
‘
= const., K, = consf.
P / c3’*[M Padm]
Fig. 8: Model crack extension plot after [6] for case i), dashed line, and case ii), trough line. the error bars correspond to a measurement uncertainty off yo5 for crack lengths.
Included in the diagram are error bars which refer to a error in the measurement of crack lengths. The scatter between test runs with different specimens may even be exceed this value. It becomes clear, that the measurement error ‘and data scatter prevent one from spot an increasing slope from such diagrams. As the most efficient way to get information about a possible change of the crack shape and the geometry factor Y, direct investigations of as-indented and (nearly) critical cracks have to be recommended.
f 5%
CONCLUSIONS AND SUMMARY The stable growth of indentation cracks in silicon nitride has been monitored using a multiple decoration technique. A decrease of the crack aspect ratio occurs for the four investigated indentation loads. It was possible to calculated the decrease of the crack aspect ratio assuming that the stress intensity
125
factor along the crack contour is constant. The asindented crack shape was taken into consideration. The correspondence of the calculated decrease with the experimental one is good. Model calculations of the stable crack extension of indentation cracks were performed to investigate the influence of the observed change of the crack shape and the consequent decrease of the geometry factor Yon the indentation strength. The results show that such a decrease has an additionally stabilizing effect on the cracks during the stable extension. As a consequence the commonly accepted relation between indentation strength and indentation load, 01s a P" with n = 113 is not longer fulfilled for materials with a single-valued toughness and an exponent n < 113 is observed. An evaluation of such data according to the standard theory will lead to the erroneous conclusion that the material exhibits R-curve behavior. For materials that exhibit a true rising fracture resistance, the decrease of Y will lead to an overestimation of the R-curve effect. The complex mutual influences of a possible variable residual stress parameter, the tendency of indentation cracks to change their geometry during flexural loading and a possible R-curve on the evolution of the geometry factor Y during stable crack extension depend on the tested material as well as on the test-geometry. It is recommended to investigate the evolution of the crack shape if fracture properties are to be deduced from stable crack growth experiments on indentation cracks.
REFERENCES Chantikul, P., Anstis, G.R., Lawn, B.R., Marshall, D.B., A Critical Evaluation of Indentation Techniques for Measuring Fracture Toughness: 11, Strength Method. J. Am. Ceram. SOC.,64, (198 1) 539-543 Cook, R.F., Lawn, B.R., A Modified Indentation Toughness Technique. J. Am. Ceram. SOC., 66, (1983) C200-C201 Krause, R.F., Rising Fracture Toughness from Bending Strength of Indented Alumina Bars. J. Am. Ceram. SOC.,71, (1988) 338-343 Braun, L.M., Bennison, S.J., Lawn, B.R., Objective Evaluation of Short-Crack Toughness Curves Using Indentation Flaws: Case Study on Alumina-based Ceramics. J. Am. Ceram. SOC.,75, (1992) 3049-3057 Ramachandran, N., Shetty, D.K., Rising-CrackGrowth Resistance (R-Curve) Behaviour of Toughned Alumina and Silicon Nitride. J. Am. Ceram. SOC.,74, (1991) 2634-2641 Dransmann, G.W., Steinbrech, R.W., Pajares, A. et al., Indentation Studies on Y203-stabilized Z r 0 ~ :I, Toughness Determination from Stable Growth of Indentation-Induced Cracks. J. Am. Ceram. SOC.,77, (1994) 1194-1201 Anstis, G.R., Chantikul, P. Lawn, B.R., Marshall, D.B., A Critical Evaluation of Indentation Techniques for Measuring Fracture Toughness: I, Direct Crack Measurements. J. Am. Ceram. SOC..
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64, (1981) 533-538 Lawn, B.R., Fracture of Brittle Solide, Cambridge University Press, Cambridge, (1993) 269. Sglavo, V.M., Pancheri, P., Crack Decorating Technique for Fracture-Toughness Measurement in Alumina. J. Eur. Ceram. SOC.,17, (1997) 16971706 Cook, R.F., Lawn, B.R., Fairbanks, C.J., Microstructure-Strength Properties in Ceramics: I, Effect of Crack Size on Toughness. J. Am. Ceram. SOC.,68, (1985) 604-615 Lawn, B.R., Padture, N.P., Braun, L.M., Bennison, S.J., Model for Toughness Curves in TwoPhase Ceramics: I, Basic Fracture Mechanics. J. Am. Ceram. SOC.,76, (1993) 2235-2240 Ramachandran, N., Shetty, D.K., Prediction of Indentation-Load Dependence of Fracture Strengths from R-Curves of Toughened Ceramics. J. Mat. Sci., 28, (1993) 6120-6126 Smith, S.M., Scattergood, R.O., Determination of Short-Crack Toughness Curves. J. Am. Ceram. SOC.,79, (1996) 129-136 Li, C.-W-, Lee, D.-J., Lui, S.-C., R-Curve Behaviour and Strength for In-Situ Reinforced Silicon Nitrides with Different Microstructures. J. Am. Ceram. SOC.,75, (1992) 1777-1785 Fett, T., An Analysis of the Residual Stress Internsity Factor of Vickers Indentation Cracks. Eng. Fract. Mech., 52, (1995) 773-776 Bleise, D., Steinbrech, R.W., Flat R-Curve from Stable Propagation of Indentation Cracks in Coarse-Grained Alumina. J. Am. Ceram. SOC.,77, (1994) 3 15-322 Lins, W., Kaindl, G., Peterlik, H., Kromp, K., A Novel Resonant Beam Technique to Determine the Elastic Moduli in Dependence on Orientation and Temperature up to 2000 C. Rev. Sci. Instrum., 70, (1999) 3052-3058 Damani, R.J., Schuster, C., Danzer, R., Polished Notch Modification of SENB-S Fracture Tougness Testing. J. Eur. Ceram. SOC., 17, (1997) 1685-1689 Jones, S.L., Norman, C.J., Shahani, R., CrackProfile Shapes Formed Under a Vickers Indent Pyramid. J. Mat. Sci. Let., 6, (1987) 721-723 Lube, T., Riaerzeugung und Bruchzahigkeitsmessung in Siliziumnitrid mit Hiirteeindriicken, Dissertation, Montanuniversittit Leoben (1999) Newman, J.C., Raju, I.S., An Empirical StressIntensity Factor Equation for the Surface Crack. Eng. Fract. Mech., 15, (1981) 185-192 Krause, R.F., Flat and Rising R-Curves for Elliptical Surface Cracks from Indentation and Superimposed Flexure. J. Am. Ceram. SOC, 77, (1994) 172-178 Lube, T., Indentation Crack Profiles in Silicon Nitride, submitted to J. Eur. Ceram. SOC. Padture, N.P., Lawn, B.R., Toughness Properties of a Silicon Carbide with in Situ Induced Heterogeneous Grain Structure. J. Am. Ceram. SOC.,77, (1994) 2518-2522
PREDICTION OF THERMAL SHOCK RESISTANCE OF COMPONENTS USING THE INDENTATION-QUENCH TEST M. Collin*and D. Rowcliffe Materials Science and Engineering, Royal Institute of Technology, 100 44 Stockholm, Sweden
ABSTRACT Two applications of the indentation-quench test have been studied. The thermal shock resistance is evaluated for three ceramic materials (alumina, silicon carbide whisker reinforced alumina and silicon nitride) and the ranking is in agreement with results in the literature. The results show that the sensitivity of the indentationquench test is higher compared to other methods. The indentation-quench method is then applied to a simplified component (tapered plate). It is demonstrated how the test can be used for making a rough estimate of the pattern of thermal stresses generated on the surface at quenching.
INTRODUCTION Ceramic materials have generally good high temperature properties and are prime candidates for many applications involving thermal excursions. Examples are as varied as engine components, electronic devices and cutting tools. In many of these applications the ceramic material will meet rapid temperature changes, which will cause transient thermal stresses and risks for thermal shock damage. For simple geometries, the transient stresses at the surface will be tensile on rapid cooling and compressive on rapid heating. Normally rapid cooling is the most dangerous temperature change for ceramics, because the tensile strength is lower than the compressive strength. The components will fracture if the tensile stress exceeds the strength of the material and insufficient thermal shock resistance often limits the application area of ceramics. It is thus necessary to consider the thermal shock resistance and there are needs both for evaluation of the thermal shock resistance of materials themselves and for evaluation of the risks for thermal shock damage in components. A simple approach to estimating the thermal shock resistance of ceramic materials is to use the traditional thermal shock resistance parameters [ 1,2]. These parameters permit ranking of materials for general
working conditions and are useful for the first selection among materials. However the parameters treat heat transfer and material properties in a very simplified way and the parameters are of limited value for anisotropic materials and fiber composites. For more extensive investigations, practical measurements are needed. A common approach is the quenching-strength method [3,4]. Heated samples are quenched into a coolant, generally water, and the remaining strength after quenching is evaluated by bending. The remaining strength is plotted as function of the temperature difference, AT, over which the sample is quenched, and the critical temperature difference, AT,, is defined as the point at which the material shows a drastic drop in remaining strength. The drop in strength is caused by a limited growth of at least one of the most harmful defects. The harmfulness of a defect depends on both size and location and both of these are statistically distributed. A consequence of this is that a great number of specimens are needed for good statistics in the quenching-strength test. The specimens have to be manufactured with standardized shape and size. Alternatively, it has been suggested to use NDE techniques for evaluation of thermal shock damage after quenching [5]. Examples of such techniques are measuring Young’s modulus and internal friction by sonicresonance [6]. NDE evaluation is particularly advantageous for studying cumulative effects of cyclic thermal shock. However, most NDE techniques demand a rather large number of specimens in order to obtain good statistics. Statistical effects can be reduced by introducing precracks with known size and location and an indentation-quench test to evaluate the thermal shock resistance of brittle materials is currently being explored in detail [7-lo]. This indentation-quench test treats localized cracks with known crack geometry, which makes modeling possible. As the artificial cracks are made in the center of the sample surface, edge effects are avoided. Another advantage is that the method in most cases can be applied directly to components.
127
Table 1. Material properties Alumina" (130°C)
Reinforced Aluminab (200°C)
Silicon Nitride (400°C)
Elastic Modulus (GPa)
E
410
430
300d
Poisson's ratio
V
0.23
0.22
0.24d
Thermal Expansion ( K ' ) Thermal Conductivity (W/(m K))
a k
5.4 x 10" 24"
5.2 x 10" 24'
3.1 x -
'Ref. [ll]. bCalculatedfrom the components. 'Ref. [lo]. dRef. [12]. 'Ref. [13].
For the component engineer it is important to evaluate if there is any risk for thermal shock damage in the component and what part of the component that is most at risk. The first step is usually a finite element calculation of transient thermal stresses. Commonly the next step is practical measurements using thermal shock or thermal cycling of components under severe conditions and evaluation by inspection for visible cracks. In a similar way as for the quenching-strength type of tests, many components are needed for good statistics. By using the indentation-quench test it should be possible to obtain good statistics from one single component. Originally the indentation-quench test was designed and evaluated for cutting tool inserts but the results suggest that the test could be used for more types of components. In this paper the indentation-quench test is used to compare the thermal shock resistance of three ceramic materials. Further, the potential of usage of the method as a diagnostic tool for components is illustrated by measurements and calculations on a tapered plate.
test. The definitions of the crack length, c, and the crack depth, a, have been given elsewhere [lo]. The quenches were made by heating the specimens in a furnace with air atmosphere and then quenching them by free fall into water at 30°C. The furnace temperature was selected to give the desired temperature difference, AT. The water bath did not show any measurable change in temperature during quenching. In all cases the specimens were kept in the furnace for 20 minutes to assure temperature uniformity before quenching.
SIMPLIFIED COMPONENT An airfoil-like tapered plate was manufactured. The material was the same grade of silicon nitride as specified above. In all, 12 indents arranged in 4 rows were made in the plate. The peak load and the mean crack size after indentation were 200 N and 200 pm. Figure 1 shows the arrangement of the indents and the dimensions of the plate.
EXPERIMENTAL PROCEDURE Row
COMPARISON OF MATERIALS Three different materials were investigated: (1) high-purity densely sintered fine-grained alumina with an average grain size < 5 pm (Procera Sandvik), (2) alumina reinforced with 30 vol.% of silicon carbide whisker (Sandvik Coromant), and (3) silicon nitride, grade CC690, (Sandvik Coromant). The samples were in the form of plates (1) 13 mm diameter x 4 mm, (2) 13 mm square x 4 mm and (3) 25 mm square x 8 mm. Material properties are listed in Table 1. The measurements were performed using the procedure previously established for the indentationquench test [7, 101. Precracks were made through indentation with a Vickers diamond for 20 s. The peak load and the mean crack length after indentation for each material were (1) 35 N and 114 pm, (2) 60 N and 102 pm, and (3) 70 N and 105 pm respectively. Approximately 16 indentation cracks were measured in each
128
mm
Fig. 1. Diagram of the tapered plate showing the arrangement of the indents.
RESULTS AND DISCUSSION COMPARISON OF MATERIALS Figure 2 shows the crack growth as a function of the temperature difference, AT, for thermal quenches of precracked samples of the three materials. The mean percentage crack length increase, with respect to the as-indented crack length, has been calculated from the growth of the individual cracks along the surface. The error bars show 95% confidence level.
ATu = 160°C
m al
+I
300
60
100
80
120
140
160
Temperature Difference, "C
150
100
200
250
300
Temperature Difference, "C
300
0
100
200
300
400
500
600
Temperature Difference, "C Fig.2. Crack growth in indented alumina (a), reinforced alumina (b) and silicon nitride (c) quenched over a range of temperature differences. The bars show 95% confidence level of the mean value.
Regime A . At very low AT no significant crack growth can be detected. Regime B . In a medium AT interval the crack growth is stable. The variation in percentage crack growth between the individual cracks is rather small as shown by the mean values at 95% confidence level. One reason for the scatter between the individual crack sizes is the variation in the microstructure at the tip of each crack. It has been observed that the cracks grow in a stepwise manner at the microstructural scale [lo]. Regime C. At a certain AT some of the cracks grow unstably out to the sample edge or stop at another indent, while the other cracks still grow stably or do not grow at all. Within the tested interval of temperature difference, Regime C is reached for alumina and reinforced alumina but not for silicon nitride. The regimes have been modeled and explained elsewhere [lo]. The major reason for the absence of crack growth in Regime A is the very low surface heat transfer coefficient when the temperature on the surface of the sample is below the boiling temperature of water [14]. The regime with stable crack growth occurs due to the combination of residual stress from the indent and thermal stress from the quench. The stable crack growth is mainly advantageous, because it makes it possible to define specific values of AT for the evaluation of the test, such as the lowest temperature difference to cause a stable crack length increase of lo%, ATlo. Another suggested way to evaluate the test is the temperature difference when Regime C starts, ATu. Values for ATlo and ATu are listed in Table 2. A high value of A T l o or ATu indicates a high thermal shock resistance. The ranking of the materials according to thermal shock resistance is: Silicon nitride > Reinforced alumina > Alumina. This result is in accordance with literature data using other techniques [5,6,15] and shows the potential of the indentationquench test for ranking materials. It is also important to notice that thanks to the residual stress, the sensitivity of the indentation-quench test is higher than that of other methods. This can be illustrated by comparing our results for alumina (AT10 = 120°C and ATu = 160°C) with results reported in the literature from investigations using other techniques for evaluation of thermal shock damage after quenching alumina samples. In these investigations higher temperature differences are needed to get significant response. The critical temperature difference (ATc) was determined to be 200°C in a quenching-strength investigation (4mm thick samples, 22°C water) [15]. In another investigation changes in retained strength, elastic modulus and internal friction were reported to occur for AT = 250-300°C (2 mm thick samples, ice water) U61.
The pattern of crack growth is similar for all materials and can be divided into three regimes:
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Table 2. Fracture toughness and thermal shock resistance Alumina
Reinforced Alumina
Silicon Nitride
Toughnessc=lOOpm (MPadm)
KIc.100
2.9 a
6.4 a
6.9
Toughnessc=ZSOpm (MPadm)
KIC.ZS0
3.9 a
7.6a
-
Thermal Shock Resistance (“C) Thermal Shock Resistance (“C) a
ATIO
120
180
260
ATU
160
280
-
Ref. [lo].
An expression for the prediction of ATu has been derived elsewhere [lo]: (1-v) 1 AT, = A [ K / c ( ~ u ) ] - - E a f’p
1
&
Here K I , is the fracture toughness, cu and au are the crack length and the crack depth at the onset of unstable crack growth (250 pm and 175 pm respectively when the precracks are 100 pm [17]), v is Poisson’s ratio, E is the elastic modulus, a is the thermal expansion coefficient, p is the Biot number, h is the surface heat transfer coefficient, r is half the thickness of the sample, k is the thermal conductivity, F is a geometric function (F = 0.71 for semielliptical surface cracks with a/c = 0.6-1.0 [18]) and finally a,, b, , c, and d, are constants with the values 3.15, 1.33, 0.266 and 51.4 for an infinite plate [lo]. The parameter f 2 has a value between 0 and 1 and shows the part of the total temperature difference that determines the maximum stress at the surface. In this investigation we have used Equation 1 for estimations off , B and h and the results are summarized in Table 3.
>
Table 3. Predicted values off’B,
P and h
Alumina
Reinforced Alumina
f‘s
0.40
0.44
P
2.7
3.4
Wk (m-’)
1350
1700
h (W/(mZK)) 32 000
41 000
The estimated values for the surface heat transfer coefficients are 32 000 and 41 000 W/(m2K) for alumina and reinforced alumina respectively. There is an encouraging agreement between the estimated values and values found in the literature. By using a combination of fast response thin film thermocouples and calculations according to the lumped capacitance method
130
the surface heat transfer coefficient for alumina has been determined to be 10 0 0 0 4 0 000 W/(m2K) in the temperature range 10O-25O0C [ 191. This illustrates how Equation 1 can be a valuable tool to roughly estimate the surface heat transfer coefficient at AT”. If we instead knew the Biot number and all material properties we could use Equation 1 for prediction of ATu. Below, f; for silicon nitride will be estimated to be 0.27. By using this value together with Equation 1 , ATu for silicon nitride can be predicted to be 990OC. Calculated Value o f f
>
B
200
400 600 800 Temperature Difference,’C
Fig.3. Values of f 2 calculated from stable crack growth in non-tapered and tapered silicon nitride plates using Eq. 2.
>
Figure 3 shows f calculated for silicon nitride using the stable crack growth in Regime B. These calculations are based on an assumption that the sum of residual and thermal stress intensity is equal to the fracture toughness of the material after each quench:
Here KIc, F , E , a, v, f have the same meaning as is called the residual in Equation 1. The constant stress factor and is a material dependent constant. In this paper we have used a value of 0.106 for silicon
x
nitride in accordance with values used for silicon nitride in the literature [20,21]. P is the indentation load, c is the crack length at the surface, A T is the temperature difference and a is the crack depth. For measurement of the crack depth, fracturing of the specimen is required. This has been done for alumina and reinforced alumina and it has been suggested that the crack depth can be estimated from the crack length according to the following expression [17]:
Here co is the initial crack length and c is the actual crack length. It is assumed that the crack depth can be estimated in the same way for silicon nitride. Measurements with crack growth below 10% have been excluded because the figures of crack growth are uncertain. In the AT-interval 300450°C the mean value off is 0.27 for the non-tapered sample, which corresponds to a pvalue of 1.4 and a value of h/k of 350 m-I. This value is much lower than the corresponding values for alumina and reinforced alumina (1350 and 1700 m-' respectively). Thus the value of h is lower and/or the value of k is higher for silicon nitride. According to the heat transfer literature a lower value of h could be reasonable. The heat transfer situation when a plate is quenched is very complicated and there is no easy way to predict the surface heat transfer coefficient, h. The heat transfer mechanism changes while the excess temperature above the boiling point (the temperature difference between the surface and the boiling temperature of the liquid) increases. Initially the heat transfer rate will increase, but after the critical excess temperature is reached, the heat transfer rate will decrease [14]. The conditions when quenching alumina and reinforced alumina are close to the critical excess temperature and a high value of h can be presumed, while the quenching of silicon nitride is above the critical excess temperature and the value of h can be assumed to be lower.
>
SIMPLIFIED COMPONENT Figure 4 shows the results from quenching the tapered plate with A T = 100, 200, ...., 600°C. As predicted, the crack growth increases with increasing severity of the quenches and the crack growth is greatest in the thickest part of the plate. These results illustrate how the indentation-quench test can be used for components. When quenching with A T = 500°C an edge crack started to grow close to Row 4. As the crack intersected the indent cracks of Row 4, further measurements on this row would not have been significant. It is
interesting to note that the edge crack did not grow at the relatively mild quenches with AT= 100-400 "C. Thus, if there is a risk for edge cracks the investigation should be planned with mild quenches. % Crack Length Increase
1
2
3
4
3
4.5
6
7.5
(Rownumber) (Thickness, mm)
Fig. 4. Mean crack growth in indented tapered silicon nitride plate quenched over a range of temperature differences. Rows 1 - 4 are described in Fig. 1 .
>
Values o f f have been calculated for Row 1-3 (Equation 2) and the results are included in Figure 3. The mean values o f f in the temperature interval 300400°C are 0.15, 0.20 and 0.35 respectively. The value at Row 3 is surprisingly high compared to the value obtained for the non-tapered plate. We have not found the explanation for this yet. The crack growth pattern is effectively linked to the thermal stress pattern and could therefore be used for a rough estimation of thermal stresses. The maximum transient thermal stress at the surface of a plate, C T ~ ~ , , ,can ~ ~ , be calculated using the following expression [ 101:
>
(4)
>
Here E , a, v and f have the same meaning as in Equation 1. Equation 4 has been derived for an infinite plate, but it can also approximately be utilized for a tapered plate. Using Equation 4 and values off from Figure 3, the maximum thermal stresses at the surface after quenching with AT = 400°C have been calculated as 73, 98 and 169 MPa for Row 1, Row 2 and Row 3 respectively. This illustrates how the indentationquench test can be used to obtain information about the thermal stress pattern and it is also possible to compare these results with FEM calculations. The test was originally developed for cutting tools, which have a very suitable shape for polishing and indentation. We
131
think however that the test could be useful for more types of components on condition that practical solutions for polishing and indentation can be obtained. The condition that localized and well-defined cracks are measured, makes it possible to evaluate the crack growth after a specified number of thermal quench cycles. In this way the indentation-quench test could be used to study cumulative damage that might arise from thermal cycling. This could then form the basis of a method to estimate the lifetime of components based on a specified maximum crack extension in a critical region of a component.
CONCLUSIONS The indentation-quench technique has been studied and the results confirm that the method is appropriate for measurements of thermal shock resistance. Three different materials were investigated and the ranking according to thermal shock resistance was in good agreement with results reported in the literature. Measurements were made on a simplified component in the shape of a tapered plate. The results showed the greatest crack growth in the thickest part of the plate. The crack growth pattern is effectively linked to the stress pattern and a rough estimation of the thermal stress pattern on the plate was made. As a general conclusion the indentation-quench test has a variety of uses. It can be applied to all types of brittle materials in which it is possible to insert indentation precracks. It is possible to study both stable and unstable crack growth and to investigate the crack path afterwards. Further, the test can normally be applied directly to components and good statistics can be obtained by making many indents in one single component. The presence of residual stress makes the test sensitive compared to other methods and it is possible to get information from materials with very good thermal shock resistance. By combining measurements of crack growth and calculations it is possible to roughly estimate the thermal stress pattern in components. Another potential application for the test is to study the cumulative damage that might arise from thermal cycling. In these ways the test can be a valuable tool for the component engineer. Acknowledgement-This work has been performed within the Center Inorganic Interfacial Engineering, supported by the Swedish National Board for Industrial and Technical Development (NUTEK) and the following industrial partners: Erasteel Kloster AB, Ericsson Cables AB, Hoganas AB, Kanthal AB, OFCON AB, Sandvik AB, Seco Tools AB and Uniroc AB.
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D.P.H. Hasselman, Thermal Stress Resistance Parameters for Brittle Refractory Ceramics: A Compendium. Am. Ceram. SOC.Bull., 49,(1970)1033-1037. D.P.H. Hasselman, Figures-of-Merit for the Thermal Stress Resistance of High-Temperature Brittle Materials: A review, Ceramurgia Int., 4,(1978) 147-150. D.P.H. Hasselman, Strength Behavior of Polycrystalline Alumina Subjected to Thermal Shock, J. Am. Ceram. SOC.,53, (1970)490-495. D. Lewis, Fracture Mechanics of Ceramics 6, ed. Bradt and Evans. Plenum Press, New York, (1983)487-496. H.Wang and R.N. Singh, Thermal Shock Behaviour of Ceramics and Ceramic Composites, Inter. Mater. Rev., 39,(1994)228-244. W.J. Lee and E.D.Case, Thermal Fatigue in Polycrystalline Alumina, J. Mater. Sci., 25,(1990)5043-5054. T. Anderson and D. J. Rowcliffe, Indentation Thermal Shock Test for Ceramics, J. Am. Ceram. SOC.,79, (1996) 150S1514. T. Anderson and D. J. Rowcliffe, Thermal Cycling of Indented Ceramic Materials, J. Eur. Ceram. SOC.,18, (1998)2065-207 1. M. Collin and D. Rowcliffe, Analysis of the Indentation-Quench Test for Ceramics, Ceram. Eng. Sci. Proc., 20,(1999)301-308. M. Collin and D. Rowcliffe, Analysis and Prediction of Thermal Shock in Brittle Materials, Acta Mater., 48, (2000) 1655-1 665. R. Munro, Evaluated Material Properties for a Sintered a-Alumina, 80, (1997) 1919-1928. D. Richerson, Modern Ceramic Engineering, Marcel Dekker Inc., New York, (1992) 166-169. Y.S. Touloukian, Thermal expansion - nonmetallic solids. Thermophysical Properties of Matter, Plenum, New York, (1977)1140. F.Kreith and M.Bohn, Principles of Heat Transfer, Harper & Row, New York, (1986)512-517. P.Becher, D.Lewis, K.Carman and A.Gonzales, Thermal Shock Resistance of Ceramics: Size and Geometry Effects in Quench Tests, Am. Ceram. SOC.Bull., 59, (1980)542-545,548. K.Matsushita, S.Kuratani, T.Okamoto and M.Shimada, Young’s Modulus and Internal Friction in Alumina Subjected to Thermal Shock, J. Mater. Sci. Lett., 3, (1984)345-348. M. Collin and D. Rowcliffe, to be submitted. T.Fett, D.Munz and J.Neumann, Technical Note. Local Stress Intensity Factors for Surface Cracks in Plates under Power-shaped Stress Distributions, Engng. Fracture Mech., 36,(1990)647-651. Y. Kim, W.-J. Lee and E.D. Case, The measurement of the surface heat transfer coefficient for ceramics quenched into a water bath, Mat. Sci. Eng., A145, (1991)L7-Lll. N. Ramachandran and D. K. Shetty, Rising CrackGrowth-Resistance (R-Curve) Behavior of Toughened Alumina and Silicon Nitride, J. Am. Ceram. SOC.,74 (1991)26362641. T. Ohji, K.Hirao and S.Kanzaki, Fracture Resistance Behavior of Highly Anisotropic Silicon Nitride, J. Am. Ceram. SOC.,78 (1995)3125-3128.
CERAMIC COMPONENTS FOR METAL FORMING TOOLS Eckart Doege, Lutz Barnert, Torsten Hallfeldt, Steffen Kulp, Tobias Neumaier Institute for Metal Forming and Metal Forming Machine Tools, D-30167 Hannover, Germany
ABSTRACT In metal forming processes tribological conditions are the result of many factors. The tool material for example is of major importance. Improvement of existing ceramics has led to various applications in production engineering. However, until now the use of ceramics for hot massive or sheet metal forming tools has not been broadly investigated.
the surface of the die. Therefore, cooling of the tools by spraying water-based lubricants is mandatory in order to reduce high thermal loads. [ 11. In combination with high mechanical loads, the forging tool made from hot-work steel shows abrasive wear and deformation. The movement is initiated by thermally caused tempering effects, which soften the surface of the dies.
In hot forging conventional methods to increase the tool life of forging dies made from hot-work steel can basically not avoid wear. Ceramics, however, are expected to show better tribological characteristics like greater resistance against wear. Investigations in flat crush tests proved the basic suitability of ceramics as material for forging tools. Further investigations with a contoured tool indicated that silicon nitride ceramic does not show any visible wear after S O 0 0 forging cycles.
In order to achieve a notable progress in the tool life of hot forging dies, the use of ceramics is preferable. Ceramic materials are well-tried in various technical areas, especially concerning critical wear. An example is the use as valves in internal combustion motors [2]. They resist the high mechanical and thermal loads in the combustion chamber. Because of the comparability of the thermal shocks of valves and forging dies it seems reasonable to test ceramics also as material for dies in hot forging processes.
In sheet metal forming processes ceramics as tool material are expected to improve the tribological system. The process is influenced by various effects which determine the quality of the drawing part. Substantial influence on the forming process has the drawing tool with its geometry, the accuracy of the drawing tool, the tool material and their interaction in the tribological system “tool material - lubricant - sheet metal”. This interaction influences the beginning of tool surface defects due to adhesion or abrasion. These effects lead to aesthetic defects on surface of the drawing part and fracture. An example of the deep drawing process, using a ceramic die, is shown in contrary to the often used hard metal die.
In 1990 Ohuchi [3] tested the application of ceramic materials in hot massive forging. The wear of different ceramic materials was tested on segmented tools during isotherm forging. The material of the workpiece was the titanium alloy Ti-6 Al-4V. In conventional drop forging of steel materials the occurring loads are different from those of isotherm forging, so the results can not be used without further consideration.
CERAMICS IN FORGING TOOLS Tool wear is mostly the reason for the failure of hot forging dies. Therefore, extra costs for finishing processes of the tool arise in the following period. In order to increase the lifetime of the tools it is necessary to find new or improved materials for hot forging. Because of this, ceramics have to be considered because they are superior in terms of wear and thermal stability. At the Institute for Metal Forming and Metal Forming Machine Tools at the University of Hanover (IFUM) investigations are dealing with the application of ceramics in forging dies. The main loads on forging tools that cause wear on the tool are thermal and mechanical loads especially when combined. During each forging cycle the tool surface is exposed to extreme thermal stresses. While forging temperatures between 600°C - 800°C occur on
TEST EQUIPMENT The forging tests at IFUM are carried out on a 3.05 MN eccentric forging press which is equipped with an automated handling system for the parts. An integrated heating device assures a constant tool temperature of 200°C. The heating of the workpieces up to forging temperature of 1100°C is performed in an induction pusher type furnace. The frequency of the forging cycle is 13 seconds. Cold sheared carbon steel slugs are used for the tests. In the cylinder upsetting tests the diameter was 20 mm and the original height of the ingot was ho = 30 mm (Fig. 1). I
1
I
1
Fig. 1 Test geometry: left upsetting test, right cone die
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These are formed to the final height of hl = 6,8 mm, so that in the end the strain was cp = 1,48 lengthwise. In flash forging tests specimens with an initial height of h,,=40mm and a diameter of 30mm are used. The deformation ratio Q ,,is roughly 25 s-'.
CERAMICS IN UPSETTING TEST In order to get a good idea of the applicability of ceramic inserts as material for dies cylinder upsetting tests are preferable. In this process the mechanical and thermal loads of cylinder upsetting are similar to those loads of the forging process. A simple geometry for the test specimen can also be used. The ceramic body is fixed in the die by thermal shrinking. If a material fails in this trial it will fail even more when used in complicated dies.
mechanical load, the forging temperature was reduced from 1100°C to 800°C. For this lower temperature the ceramic broke after 400 cycles (Fig. 3).
FORGING DIE WITH n A S H In further tests hot pressed silicon nitride was investigated along with variations of silicon carbide. In these forging tests the three dies with silicon carbide inlays were destroyed after 14, 40 and 160 cycles (Fig. 4, left). Probably a growth of cracks appeared or a bending load acted on the ceramic. This bending load may result from the inaccurate positioning of the work piece into the die. This aspect will be further looked into in relation with Finite Element Analysis investigation.
The A1203-ceramic is already destroyed after 150 forging cycles probably due to thermal stress (Fig. 2, left). A further investigation of oxide-ceramics does not appear reasonable because of their physical properties and the stress within the forging process. Fig. 4 Fracture of silicon carbide after 14 forging cycles and no wear even after 5000 forging cycles for silicon nitride
after 150 forging cycles Fig. 2 Fracture of the Al~O~-ceramic and no wear after lo00 forging cycles for silicon nitride Contrary to the Al203-ceramic hot pressed silicon nitride shows no visible wear after lo00 strokes. In order to raise the stress on the ceramic specimen other tests were carried out with stress-loads pointing away from the centre of the inlay. The great flow of material causes extra strong mechanical abrasive loads on the ceramic. Here also, no visible wear was detected on the ceramic part, but the surrounding hot working steel showed notable scarring (Fig. 2, right). The silicon nitride ceramic shows an extraordinary resistance in the forging process. In order to find out more about the possible way of applicating this material a cylindric structure was produced in the ceramic. This structure stands for a deformation-specific additional element in forging processes. The mechanical load in forging processes leads to tensile stress in the slot caused by the pressure inside the structure.
Fig. 3 Si3N4-ceramicinlay with depression, fracture after 400 forging cycles (forging temperature 800 "C) Even after lo00 forging cycles no visible wear or fracture could be detected. In order to increase the
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The silicon ceramic which had already shown good results during the cylinder upsetting test also shows positive results when forging with flash. After 5000 forging cycles no wear was visible, figure 4 right. Such results can not be reached with conventional hot working steel.
DETERMINATION OF LOADS USING FINITE ELEMENT ANALYSIS Finite Element Analysis is a good instrument to discover loads and stress during the forging process at ambient and working temperature. The rotationally symmetrical ceramics of this investigation were bound together by thermal shrinking with the parent tool made from hot work steel. Considering the tensile stress sensitivity of ceramic materials a sufficient precompressive strain of the ceramic inlays is necessary. This pre-compressive strain must not be relieved otherwise the ceramic would fall out of the composite. The silicon carbide ceramic was not destroyed in the very first cycle. Therefore, the main cause for fracture is not only a non-critical breakage but also a transverse load acting on the inlay. When the work piece is not placed exactly in the centre of the die an unsymmetrical load on the upper die occurs. This results in transverse strains which are presented in Fig. 5 for silicon carbide. Starting with the hypothesis of the transverse strain for the determination of the reference stress by Mises (0, = 2 T-), a transverse strain of 2- = 170MPa occurs after a deformation of only 8.5 mm. Exceeding the flectional resistance this causes a breakage [4].
of the ceramic and steel ring. They were mounted after heating them to up 400°C using the lower coefficient of thermal expansion of ceramic compared to steel. Both were then pressed into the armouring. The pressure of about 1300 MPa in the die during one forging cycle led to a main load of about 950 MPa in the ceramic ring, which is much higher than the bending strength of SN. A closer look to the contact pressure of the system shows that a tool temperature of 180°C compensates the contact between armouring and ceramic ring (Fig. 8).
fracture of sillcon carbide
Fig. 5 Calculated transverse strain in forging cycle for silicon carbide with decentralised load With the help of Finite Element Analysis the reason for the failure of silicon carbide and for the nonbreakmg of silicon carbide during the 5000 forging cycles was found.
ANALYSIS OF TOOL FAILURE An industrial partner tested a tool for hot extrusion with a segment of silicon nitride (Fig. 6). That ceramic ring was applied in an area, where the work piece has to maintain high precision. armouring 2ndarmouring
(steps
"38 (SN) ceramic
b
I ~
pressure -160 MPa
)
die (2OOC)
Fig. 8 Contact pressure at room temperature and working temperature (1 80°C) Variations to improve the armouring were carried out. It was found that thermal shrinking is not suited to assemble both components. The maximum temperature of steel for thermal shrinking to avoid tempering effects is 550°C. The resulting oversize and armouring is not sufficient to give an initial stress.
CERAMICS IN SHEET METAL FORMING Fig. 6 Design of a hot extrusion tool with a silicon nitride ceramic ring The silicon nitride ring failed after 800 strokes (Fig. 7). The direction of the fractures indicates an overload.
Deep drawing is a technology often used in sheet metal forming. In deep drawing processes parts are produced for different industrial areas mostly with a number of parts amounting to one million, sometimes even more. During a deep drawing process the sheet material is normally formed with a tool system consisting of a punch, a blank holder and a drawing die. Especially the tribological system between sheet metal and tool system has an important influence on the deep drawing process, the quality of the formed component and tool life [ 5 ] . In figure 9 a conventional tool system with three friction zones in the drawing die, the blank holder and the punch is shown.
Fig. 7 Ceramic ring and fracture To analyse the relevant loads and the interaction of all components of the forging cycle the finite element method was used. Relevant data are the maximum tool temperature of 180°C and the forging temperature of 1250°C. Important for the analysis also was the oversize
Fig. 9 Friction zones in a deep drawing tool system
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In general at deep drawing processes the coefficient of friction p in the punch head zone should be high to allow a good transfer of the punch forces to the work piece. In the flange area and the die radius a material flow with small forces should be realised. Therefore, low coefficients of friction p between sheet metal and the drawing die are necessary [5]. In this context different kinds of forming lubricants and several coatings of the sheet metal material have been examined [6]. Furthermore, besides conventional steel dies, new tool concepts were developed and tested at IFUM [7,8]. For the processing of e.g. higher strength steels, aluminium and other light weight alloys new tool materials are of demand for the improvement of the tribological system in deep drawing processes. The main objectives are to reduce friction and forming forces as well as the required lubricants. Another important aspect is to avoid adhesion and pick-ups, that occur e.g. when uncoated aluminium sheets without lubricants are drawn with steel tools. Instead of these tool materials, ceramics have favourable qualities especially increased hardness, good wear resistance and excellent tribological behaviour [9. 10, 113, On the other hand, ceramics have a reduced ductility so that spontaneous cracks can appear. Furthermore, the treatment of ceramic materials is expensive, mostly it will be done with grinding methods [ 121 sometimes e.g. for silicon-nitride ceramics with laser cutting [ 131. At IFUM experimental examinations with different kinds of ceramics and steel materials have been carried out. To investigate the applicability of axisymmetric ceramic drawing dies, silicon-nitride, circonium-oxide and silicon-carbide ceramics have been applied. In comparison conventional hard metal drawing dies have been analysed under the same conditions. Figure 10 shows pick-ups in the area of the die radius, which affect the drawing results. Surface failures, geometric deviations and in the worst case cracks on the work piece can be found.
Deep drawing dies of silicon-carbide ceramics have been used with the same geometry. As shown in Figure 11 no adhesion or pick-ups can be detected.
Fig. 11 Ceramic drawing die The reason for this behaviour is the type of binding of the ceramics, which differs from the metallic bonds of steel materials. During the deep drawing process adhesion results in higher punch forces. A comparison of the maximum punch forces with the drawing die materials silicon-nitride ceramic and hart metal GE 50 without lubricants is shown in figure 12. The punch forces with the silicon-nitride ceramic are up to 23 % less compared to a hard metal tool system. I
5" Y
hard meta
g3 g 2
(GE50) ceramic
5 1
(Si3NJ
P
d o B0 = 1.06, SO = 0.30mm E
material: St4 LG Ni plat.
referem Ordder
lubrication:
dry @ IFUM
Fig. 12 Punch forces with different drawing die materials In examinations with lubricants the punch forces obtained with ceramics were 6 % less compared to hart metal drawing dies. Another reason for these results are the roughness of the surface of ceramic materials. They are lower than of conventional steel drawing dies (Fig. 13).
Fig. 10 Surface failures in the area of a die radius
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nitride showed no visible wear after even 5000 forging cycles. The use of ceramics as die inlays results in a considerably longer tool life. Only silicon nitride was convincing in these test as a suited material. In sheet metal forming ceramics as tool materials improve the tribological system. Compared to conventional dies, maximum punch forces are reduced by 23% for an non lubricated system. In consequence, higher drawing ratios are possible, which means process limits are extended. With ceramic dies reduced or no lubrication is necessary. Ceramics lead to a longer tool life. In the end maintenance costs are reduced. In order to gain more experience in the application of ceramics in forming tools further investigation are necessary. Fig. 13 Surface roughness of the examined materials
The examinations resulted in some advantages for the usage of ceramic materials. The amount of lubricants in deep drawing processes can be reduced. Consistent ceramics are more environmentally friendly. Due to the wear resistance the tool life is very high, so the expense for maintenance is smaller than for conventional steel tools. In deep drawing tools, ceramics are applied for inlays in high loaded areas, e.g. comers [14]. They are also well suited in progressive tools or components, e.g. drawing dies with small geometry. Ceramics also have advantages for processing aluminium or titanium sheet materials and other materials which show an adhesive reaction to steel tools. The inlays should be placed in areas with a high contact stress and a high relative velocity between drawing die and sheet material. Appropriate positions for ceramic inserts in large tool systems are displayed in figure 14.
References W. Stute-Schlamme: Konstruktion und thermomechanisches Verhalten rotationssymmetrischer Schmiedegesenke 198 1 . - Universitat Hannover, Fakultat f i r Maschinenwesen, Dissertation
G. Wotting, G. Leimer: Siliciumnitrid-Keramik, deren Eigenschaften und Anwendungen in der Umformtechnik.- In: Umformtechnik an der Schwelle zum nachsten Jahrtausend. 16. Umformtechnisches Kolloquium Hannover 25./26. Februar 1999 K. Ohuchi, S. Sasaki, K. Matsuno: Isothermal Forging with Ceramic Die on Industrial Basis.-In: Advanced Technology of Plasticity 1990. Proceedings of the 3rd Int. Conference on Technology of Plasticity, July 1/6 1990, Kyoto, Japan E. Doege, T. Neumaier, C. Romanowski: EinfluB der Rohteilpositionierung auf die Werkzeugbelastung beim Gesenkschmieden am Beispiel eines KeramikMetall Verbundes. In: Marc Benutzertreffen, 15.-16. Oktober 1997, Miinchen E. Doege, Wichtige EinfluBgroBen beim Tiefziehen. wt-Z. fiir industrielle Fertigung (66), Heft 11, S. 615-619,1976 W. G. Brazier, R. W. Thompson, The effect of zinc coatings, die materials and forming lubricants. SAE Technical Paper 860432, 1986
Fig. 14 Ceramic inserts located in high loaded areas of a die
SUMMARY Ceramic components are tested in a wide range of technical applications. But so far, their use in forming processes has not been broadly investigated. Obviously ceramics offer an enormous potential in several fields. In hot massive forming high thermal and mechanical loads lead to wear of the forging tools made of hot-work steel. In comparison a ceramic tool inlay of silicon
C. Frank, K. Droder, Reduced friction in deep drawing by means of organic sheet coatings. Birmingham, IDDRG Working Group Meeting, 07.-09.06.1999 E. Doege, C. Frank, Polymers as tool materials in sheet metal forming. 2nd Int. Assiut Conf. on mechanical engineering, Advanced technologies for production, Assiut, Egypt., 02.-04.03.1999 W. D. Kingery, Introduction to ceramics. John Wiley & Sons, New York, 1976 D. H. Buckley, K. Mikyoshi, Fundamental
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tribological properties of ceramics. Roc. 9" Int. Conf., Composites and advanced ceramic materials, Amerc. Ceramic Society, 20.24.01.1985 (1 1) E. A. Almond, L. A. Lay, M. G. Gee, Comparison of sliding and abrasive wear mechanisms in ceramics and cemented carbides. 2"6 Int. Conf. Science Hard Materials, Rhodes, 1986 (12) B. G. Koepke, R. J. Stokes, Effect of workpiece properties on grinding forces in polycrystalline ceramics. Proc. Symp. National bureau of standards, Gaithersburg, Maryland, 13.15.11.1987 (13) G. Warnecke, Schleifen von Hochleistungskeramik-Werkstoff, Anwendung, Bearbeitung, Qualittit. Verlag T W Rheinland, Koln, 1994 (14) J. Muller, R. Heinze, Einsatz von Keramik in
Werkzeugen fur die Blechumformung. Int. Konf. "Neuere Entwicklungen in der Blechumformung" Fellbach, 23.-24.05.2OOO
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EFFECT OF GRAIN BOUNDARY COMPOSITION ON HIGHTEMPERATURE MECHANICAL PROPERTIES OF HOT-PRESSED SILICON CARBIDE SINTERED WITH YTTRIA Donna Hermanub* and Hagen Klemm Fraunhofer-Institut Keramische Technologien und Sintenverkstoffe, Winterbergstralje 28, 0 1277 Dresden, Germany
ABSTRACT In materials with microscopic grain sizes, the grain boundary composition plays a decisive role in determining the resulting high-temperature oxidation, creep and slow-crack-growth behaviour. Therefore, it is crucial that the relationship between composition and properties is known in order to tailor the material to meet specific requirements. In the present study, the chemical composition in the system silicon carbide with Y2O3 as a sintering aid was varied. The resultant powders were hot pressed and tested in four-point bending at room temperature as well as at 1500 "C. The room-temperature strength and the fracture toughness (ICL method) were determined. The dynamic fatigue strength was measured at a stress rate of 0.05 MPds and an initial stress of 50 MPa. Hotpressed and tempered (100 h in air at 1500 "C) bending bars with either natural flaws or a single edge notch were tested. The fracture surfaces were examined and an attempt was made to correlate the observations with results of the mechanical tests. Differences were found between the Y203-rich and the Si02-rich materials, independent of the flaw type. At 1500 "C, the Si02-rich material exhibited slow crack growth, whereas the Y203-rich material exhibited typical fracture features such as a mirror, mist and hackle. The Y203-rich specimens failed at higher stresses than the Si02-rich specimens. Additionally, for notched specimens, the dynamic fatigue strengths were higher at 1500 "C than at room temperature and at both temperatures after tempering. It is assumed that processes involving blunting of the crack tip are responsible for the observed behaviour.
Silicon carbide, being a mainly covalently bonded material, is very dificult to sinter. Additives such as A1203 (or Al) and combinations of these with Y2O3 and such additives as SO2, CaC03 and MgO have been used as sintering aids [2-111, although the rare-earth oxide Dy2O3 has also been used [12]. An intergranular amorphous phase results from liquid-phase sintering or impurity segregation; this phase acts as a liquid phase and adopts an equilibrium thickness at high temperatures [ 131. Because oxygen, impurities such as Ca" and other cations such as rare-earth cations are essentially insoluble in Sic, they segregate into the grain boundaries to form a residual amorphous intergranular phase [14]. It has been proposed that the thickness of this film is related to the composition and not the amount of the glass phase present [ 151. Thus it is assumed that the chemical composition of the grain boundaries plays a decisive role in determining the mechanical behaviour at high temperatures. In this study, the effect of altering the grain boundary chemistry on the dynamic fatigue behaviour of notched and unnotched LPS-Sic bending specimens at room temperature and at 1500 "C was examined. Up to now, there have been relatively few studies published that relate to the high-temperature mechanical behaviour of LPS- Sic [16-191 or of LPS-SiC/Si3N4composites [20231. It is hoped that through this work, a better understanding of the way in which the grain boundary composition influences the crack-growth behaviour of LPS-Sic at room and elevated temperatures will be gained.
EXPERIMENTAL PROCEDURES INTRODUCTION Among silicon-based nonoxide ceramics, liquid-phasesintered silicon carbide (LPS-Sic) is a potential material from which combustion chamber linings and other stationary structural components in high-temperature gas turbines can be made. For these applications, the material must have low weight, long-term stability, high strength and good fracture toughness at room and high temperatures, no time-dependent strength degradation and minimal creep [I].
The materials were fabricated by mixing S i c powder with different amounts of Y203 and Si02 in an attritor in isopropanol, drylng in a rotary vacuum evaporator, and calcining at 450 "C. The additive amounts were chosen such that the grain boundary volume remained constant at 12 % and the molar ratio of Y2O3 to Si02 was 0.7 or 2.0. The resulting powders were hot pressed for one hour at a temperature of 2000 "C in a graphite-lined chamber in an argon atmosphere. The densities were measured by the Archimedes method and were found to match the theoretical values.
139
Bending specimens were produced from the materials thus obtained. The room-temperature strength and the fracture toughness (ICL method) were obtained for both materials. For each of the two materials, half of the remaining specimens were notched with a saw and 1pm-sized diamond paste to a depth of ca. 350 pm. Half of the unnotched and half of the notched specimens were tempered for 100 h in an oven in air at 1500 "C. The resulting specimens underwent dynamic fatigue tests in four-point (inner span: 20 mm; outer span: 40 mm) bending at room temperature and at 1500 "C with a loading ramp of 0.05 MPds and an initial stress of 50 MPa. X-ray diffraction (Cu Q was performed on hot-pressed and tempered specimens after dynamic fatigue testing at room temperature and at 1500 "C. XRD was subsequently performed on notched specimens tempered or tested at 1500 "C after removing an additional 100 pm.
Figure 1. Fracture surface of a Si02-rich specimen with natural flaws after dynamic fatigue testing at 1500 "C. The arrow points to the approximate location of the fracture origin.
The fracture surfaces were examined in a light microscope and in the SEM. EDS was used to determine the fracture origins. Polished and CF4 plasma-etched cross sections were also examined in the SEM.
RESULTS Room-Temperature Strength and Fracture Toughness The room-temperature strength and fracture toughness (ICL method) were 450 MPa and 3.8 MPa.m'" respectively for the Si02-rich material and 660 MPa and 3.O MPa.m"* respectively for the Y203-rich material. Both materials exhibited mainly intergranular fracture.
Figure 2. Fracture surface of a tempered Si02-rich specimen with natural flaws after dynamic fatigue testing at 1500 "C
Dynamic Fatigue Testing of Specimens with Natural Flaws The results of the dynamic fatigue tests for specimens with natural flaws are shown in Table 1. rable 1. Dynamic fatigue strengths (0.05 MPds starting at 50 MPa) for snecimens with natural flaws
Three or four results were averaged to obtain each of the above values. Figures 1 - 4 show the fracture surfaces of hot-pressed and tempered SiO2-rich and Y203-rich specimens after the dynamic fatigue tests.
140
Figure 3. Fracture surface of an Y203-richspecimen with natural flaws after dynamic fatigue testing at 1500 "C
Figure 4. Fracture surface of a tempered Y~Oj-rich specimen with natural flaws after dynamic fatigue testing at 1500 "C
Figure 6. Fracture surface of a SiO2-rich notched specimen after dynamic fatigue testing at 1500 O C
Dynamic Fatigue Testing of Specimens with an Edge Notch The results of the dynamic fatigue tests for notched specimens are shown in Table 2; the fracture surfaces for hot-pressed specimens tested at room temperature and at 1500 "C are shown in Figures 5 - 8. Table 2. Dynamic fatigue strengths (0.05 MPds starting at 50 MPa) for specimens with an edge notch of ca. 350 pm depth fracture stress at fracture stress at room temperature 1500 "C [MPa] [MPaI Hotmaterial Si02-rich 265
265
Figure 7. Fracture surface of an Y203-rich notched specimen after dynamic fatigue testing at room temperature
Three or four results were averaged to obtain each of the above values.
Figure 8. Fracture surface of an Y203-rich notched specimen after dynamic fatigue testing at 1500 "C Figure 5 . Fracture surface of a SiOz-rich notched specimen after dynamic fatigue testing at room temperature
The X-ray diffraction results are shown in Table 3 for notched specimens tested at room temperature and Table 4 for notched specimens tested at 1500 "C.
141
Table 3. Grain boundary phases found after performing dynamic fatigue tests on notched specimens at room temperature Surface ITemperd ITemperd 1
I
surface
I
100 pm below
I
Table 4. Grain boundary phases found after performing n notchec 100 pm Material Surface below below surface
Figure 10. Polished and etched cross section of the Y203-richmaterial after tempering for 100 h in air at 1500 "C
DISCUSSION
surface
SO2-rich Y ~ S ~ O S , y4.67 (Si04) 3 0 , Y2Si207, SiO2
Y2SiOs, y4.67 (Si04)
Y203-rich Y2SiOs, Y2Si207, Si02
Y2SiOs, Y2Si207
SO2
30,
Y2Si207
The polished and etched cross sections of the tempered specimens are shown in Figures 9 and 10. The thickness of the oxidation layer for the SiO2-rich material is ca. 5 pm and for the Y203-rich material, ca. 10 pm.
Figure 9. Polished and etched cross section of the Si02rich material after tempering for 100 h in air at 1500 "C
General Differences Between Y203-Rich and Si02-Rich Materials The two materials had different grain boundary compositions and liquidus temperatures (ca. 1820 "C for the Y203-richmaterial and ca. 1940 "C for the SiO2-rich material). This is believed to have led to a lower strength and higher fracture toughness at room temperature for the SiO2-rich material than for the Y203-richmaterial. With regard to the dynamic fatigue strength, differences were found between the Y203-rich and the Si02-rich materials, independent of the flaw type. The Y203-rich specimens generally failed at higher stresses than the SiO2-rich specimens, as shown in Tables 1 and 2. At 1500 "C, slow crack growth was observed in the Si02rich specimens, whereas typical fast fracture markings were observed in the Y203-rich specimens. Thus, slow crack growth during high-temperature testing appears to be associated with low dynamic fatigue stresses and with the Si02-rich materials. The steady-state creep rate for the Y203-rich material (3.5 x h-') was comparable to that of the SO2-rich material (2.9 x 10" h-' for a material with the same composition but a lower grain boundary volume fraction of 0.07; for a grain boundary volume fraction of 0.12, the steady-state creep rate should be somewhat higher, but still similar to that of the Y203-rich material); this indicates that creep probably did not play a significant role in strengthening the Y203-rich material during dynamic fatigue testing. X-ray difhction analysis of notched specimens (see Tables 3 and 4) showed that for both test temperatures, the two materials had different grain boundary phases on and 100 pm below the surface, but the same grain boundary phases on the surface after tempering. 100 pm below the surface of the tempered materials, however, different grain boundary phases were found for the two
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materials. Figures 9 and 10 c o n f m that the Y203-rich material was oxidized to a greater extent than the Si02rich material, but that neither material suffered significant oxidation damage after being tempered for 100 h in air at 1500 "C. It is assumed that differences in the condition and grade of crystallinity of the grain boundaries are mainly responsible for the differences observed in the two materials.
Dynamic Fatigue Testing of Specimens with Natural Flaws It can be seen that only at 1500 "C were there significant differences between the results (see Table 1). At 1500 "C, the dynamic fatigue strengths of the Y2O3rich specimens were much higher than those of the SO2-rich specimens. This is probably due to the differences in the condition and grade of crystallinity of the grain boundaries. The strengths at 1500 "C were lower than those at room temperature, especially for the SO2-rich material. This is probably related to the softening of the glassy phase at 1500 "C. Tempering caused the strength to increase for the Si02rich material and to decrease for the Y203-richmaterial, especially at 1500 "C. The SO2-rich material was not oxidized significantly during the tempering experiments (see Figure 9, showing an oxidation layer of ca. 5 pm). The fracture origins for the specimens tested at room temperature were surface defects; after tempering, the fracture origins were surface regions rich in Si02 and Y. At 1500 "C, the fracture origins were surface defects (see Figure 1); after tempering, the fracture origins were defects in the interior of the specimens (see Figure 2), the oxidation layer having effectively healed the surface defects. Thus, tempering had a positive effect on the dynamic fatigue strength, especially at 1500 "C. None of the typical markings associated with fast fracture (i.e., fracture mirror, mist, and hackle) could be observed on the fracture surfaces after testing at 1500 "C. The Y203-richmaterial was oxidized to a much greater extent (see Figure 10, showing an oxidation layer of ca. 10 pm). At room temperature, the fracture origins were surface defects such as large S i c grains; after tempering, the fracture origins were still surface defects, but were caused by build-up of the oxidation products. At 1500 "C, the fracture origins were either large particles or regions enriched in the grain boundary phase at or close to the surface (see Figure 3); after tempering, the fracture origins were generally regions consisting of Si02 and large Y2SizO7 crystals close to the surface (see Figure 4). Thus, tempering had a negative effect on the dynamic fatigue strength, especially during testing at 1500 "C. Oxidation of the grain boundary phase in the top 10 - 30 pm of the bulk occurred.
Dynamic Fatigue Testing of Specimens with an Edge Notch Significant differences between the results could be found, especially at 1500 "C (see Table 2). The dynamic fatigue strengths of the Y203-rich specimens were much higher than those of the SiO2-rich specimens at 1500 "C, whereas they did not differ considerably from those of the SiO2-rich specimens at room temperature. At 1500 "C, flow andor oxidation of the grain boundaries may occur to hinder crack propagation. Since oxidation is much more and flow is slightly more pronounced in the Y203-richmaterial than in the Si02-material, they may be responsible for the differences observed at 1500 "C. For materials with an edge notch, hgher strengths were obtained at 1500 "C than at room temperature and at both temperatures after tempering for 100 h at 1500 "C in air. It is assumed that processes involving blunting of the notch tip are responsible for the observed behaviour. X-ray diffraction analysis (see Tables 3 and 4) showed that after testing at 1500 "C, only additional crystobalite was present on the surface for the SO2-rich materials; for the Y203-richmaterials, crystobalite was present and Y2O3 was absent. After tempering, the crystalline grain boundaries found on and 100 pm below the surface after testing at 1500 "C were identical to those found after room-temperature testing. Thus, for the tempered specimens, no significant changes detectable by X-ray diffraction occurred during dynamic fatigue testing at 1500 "C. The 25% decrease in strength for the SO2-rich material and the 50% increase in strength for the YZO3rich material might be related to softening of the glassy phase during testing at 1500 "C. The edge notch changes the behaviour in the following ways:
1. The Si02-rich material underwent extensive slow crack growth during testing at 1500 "C; the dynamic fatigue strengths for hot-pressed and tempered specimens were correspondingly low.
2. Tempering causes the fracture stress to increase for Y20-rich and for Si02-rich materials, especially at room temperature. Since the amount of microstructural damage caused by tempering is the same for specimens with and without an edge notch, this increase must be due to blunting of the crack tip by tempering.
3. The fracture stress at 1500 "C is higher than at room temperature for SiO2-rich (see Figures 5 and 6) and YzO3-rich hot-pressed materials (see Figures 7 and 8) and for Y203-rich tempered materials. The increase may be due to blunting of the crack tip and lowering of the stress intensity factor by oxidation or flow of the grain boundaries during hightemperature testing. However, other mechanisms may account for the increase in dynamic fatigue strength at 1500 "C. Changes in the amounts of glassy phase, crystalline phases and microstructural
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damage in the grain boundaries due to oxidation may also have an effect on the dynamic fatigue behaviour.
Conclusions The dynamic fatigue behaviour of two materials, one rich in Si02 and the other rich in Y2O3, was examined at room and at elevated temperatures for unnotched and notched specimens. At room temperature, no differences were observed. At 1500 "C, however, significant differences were observed, regardless of defect type. The SiO2-rich specimens exhibited less oxidation, more slow crack growth and lower dynamic fatigue strengths than the Y203-rich specimens did. Additionally, for notched specimens, the dynamic fatigue strengths were higher at 1500 "C than at room temperature and at both temperatures after tempering. It is assumed that processes, such as oxidation, involving blunting of the crack tip are responsible for the observed behaviour. However, further studies must be performed in order to explain the high-temperature behaviour of these materials. Not only oxidation, but other mechanisms must be active at 1500 "C to cause the differences which were seen especially in the tempered materials between room temperature and 1500 "C. Dynamic fatigue testing could be performed in an inert atmosphere to remove the effects of oxidation. Additionally, the crystallization behaviour of the two materials could be tested using techniques such as differential thermal analysis, TEM and quantitative analysis of the X-ray diffraction results.
[6] Lee, J.K., Kang, H.H., Kim. Y.J., Lee, E.G., Kim, H., Effects of YAG-Phase Amount on the Microstructure and Phase Transformation during the Liquid-Phase Sintering of p-Sic, Key Engineering Materials Vols. 161-163, 1999,263-266. [7] Nader, M., Aldinger, F., Hoffmann, M.J., Influence of the alp-Sic phase transformation on microstructural development and mechanical properties of Iiquid phase sintered silicon carbide, J. Mater. Sci. 34, 1999, 11971204. [8] Kim, Y.-W., Kim, J.-Y., Rhee, S.-H., Kim, D.-Y., Effect of initial particle size on microstructure of liquidphase sintered a-silicon carbide, Journal of the European Ceramic Society 20,2000,945-949.
[9] Winn, E.J., Clegg, W.J., Role of the Powder Bed in the Densification of Silicon Carbide Sintered with Yttria and Alumina Additives, J. Am. Ceram. SOC.82 [12], 1999,3466-70. [lo] Kim, Y.-W., Mitomo, M., Fine-Grained Silicon Carbide Ceramics with Oxynitride Glass, J. Am. Ceram. SOC.82 [lo], 1999,2731-36. [ l l ] Zhan, G.-D., Mitomo, M., Kim, Y.-W., Microstructural Control for Strengthening of Silicon Carbide Ceramics, J. Am. Ceram. SOC.82 [lo], 1999, 2924-26.
REFERENCES
[12] Kim, S., Kriven, W.M., Preparation, Microstructure, and Mechanical Properties of Silicon Carbide-Dysprosia Composites, J. Am. Ceram. SOC.80 [ 121, 1997,2997-3008.
[I] Hecht, N.L., Goodrich, S.M., Chuck, L., McCullum, D.E., Tennery, V.J., Mechanical Properties Characterization of One S i c and Two Si3N4 Commercially Available Ceramics, Ceramic Bulletin 7 1 [4], 1992, 653.
[I31 Tanaka, I., Kleebe, H.-J., Cinibulk, M.K., Bmley, J., Clarke, D.R., and Riihle, M., Ca Concentration Dependence of the Equilibrium Thickness of the Intergranular Film in Si3N4, J. Am. Ceram. SOC.77 [4], 1994,911.
[2] Lee, J.-K., Tan&, H., Kim, H., Movement of liquid phase and the formation of surface reaction layer on the sintering of p-Sic with an additive of yttrium aluminium garnet, Journal of Materials Science Letters 15, 1996,409.
[14] Quinn, G.D., Fracture Mechanism Maps for Advanced Structural Ceramics Part 1 Methodology and Hot-Pressed Silicon Nitride Results, J. Mater. Sci. 25, 1990,4361.
[3] Falk, L.K.L., Microstructural Development during Liquid Phase Sintering of Silicon Carbide Ceramics, Journal of the European Ceramic Society 17,1997,983.
[15] Wang, C.-M., Pan, X., Hoffinann, M.J., Cannon, R.M., Riihle, M., Grain Boundary Films in Rare-EarthGlass-Based Silicon Nitride, J. Am. Ceram. Soc. 79 [3], 1996,788-92.
[4] Mulla, M.A., Krstic, V.D., Pressureless sintering of p-Sic with A1203 additions, Journal of Materials Science 29, 1994,934. [5] Mulla, M.A., Krstic, V.D., Low-Temperature Pressureless Sintering of p-Silicon Carbide with Aluminum Oxide and Yttrium Oxide Additions, Ceramic Bulletin 70 [3], 1991,439.
[I61 Wolf, C., Hubner, H., Adler, J., Mechanical Behaviour of Pressureless Sintered Sic at High Temperature, pp. 465-470 in Third Euro-Ceramics, Vol. 3. Edited by P. Duran and J.F. Femandez, 1993. [17] Jou, Z.C., Virkar, A.V., Cutler, R.A., High temperature creep of Sic densified using a transient liquid phase, J. Mater. Res. 6 [9], Sep 1991, 1945. [18] Keppeler, M., Reichert, H.-G., Broadley, J.M., Thurn, G., Wiedmann, I., Aldinger, F., High
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Temperature Mechanical Behaviour of Liquid Phase Sintered Silicon Carbide, Journal of the European Ceramic Society 18, 1998, 521-526. [19] Kinoshita, T., Munekawa, S., Tanaka, S.-I., Effect of Grain Boundary Segregation on High-Temperature Strength of Hot-Pressed Silicon Carbide, Acta mater. Vol. 45, NO.2, 1997,801-809. [20] Sajgalik, P., Hnatko, M., Lofaj, F., HvizdoS, P., Dusza, J., Warbichler, P., Hofer, F., Riedel, R., Lecomte, E., Hoffmann, M.J., SiC/Si3N4 nano/microcomposite - processing, RT and HT mechanical properties, Journal of the European Ceramic Society 20, 2000,453-462. [21] Dusza, J., Sajgalik, P. Steen, M., Fracture Toughness of a Silicon Nitride/Silicon Carbide Nanocomposite at 1350 "C, J. Am. Ceram. SOC.82 [12], 1999,3613-3615. [22] Cheong, D.-S., Hwang, K.-T., Kim, C.-S., HighTemperature Strength and Microstructural Analysis in Si3N4/20-vol%-SiC Nanocomposites, J. Am. Ceram. SOC.82 [4], 1999,981-86. [23] Rouxel, T., Wakai, F., Sakaguchi, S., R-Curve Behaviour and Stable Crack Growth at Elevated Temperature (1500 "C - 1650 "C) in a Si3N4/SiC Nanocomposite, J. Am. Ceram. SOC. 77 [12], 1994, 3237-43.
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THERMAL SHOCK PROPERTIES OF SIALON CERAMICS Pernilla Pettersson*,Zhijian Shen, Mats Johnsson, and Mats Nygren Department of Inorganic Chemistry, Stockholm University, S-106 91 Stockholm, Sweden ABSTRACT An indentation-quench method based on Vickers cracks for measuring thermal shock properties has been applied to different sialon materials. For sialons based on the P-phase, Si6-,AI,0,Ng., (0 I z I4.2), the best thermal shock resistance was found at low z values. The thermal shock properties could be further improved by adding an intergranular yttrium containing glass phase. For sialons consisting of a and the pmixture of the a- (YxSil~+,,+n~Al,,,+nOnN16n) sialon their thermal shock resistance improved with increasing amount of P-phase and adding an intergranular glass phase improved the thermal shock resistance.
INTRODUCTION
a-and p- sialons are two solid solution phases with compositions and crystal structures close to a-and pSi3N4 respectively. p-sialon is represented by the general formula Si6-zAl,0,Ns.Z, with 0 < z I4.2. a-sialon is represented by MxSi~Z.(m+n~lm+nOnN~~n where M represents metal ions (most often yttrium or rare earths) and x, m, and n are parameters that restricts the formation of a-sialon to small composition areas [ 11. An intergranular glass phase can be added in order to assist sintering. Sialon materials based on the a- and p-sialon phases or mixtures thereof are attractive for high temperature applications. However, at the moment there is little experimental data available in the literature concerning their thermal shock behaviour i. e. the effect of thermal cycling on mechanical properties of sialon based materials. In this article we present a qualitative investigation of the thermal shock properties of a- and p- sialon materials. An indentation-quench method based on Vickers cracks for measuring thermal shock properties [2] has been applied and evaluated. For P-sialon materials the thermal shock properties have been correlated with both different z-values within the solid solution range and the amount of residual intergranular glass phase present. For dpsialons the thermal shock properties are correlated with different ratios between the a- and the P-phase and with addition of increasing amount of intergranular glass phase. For those materials the a-phase (YxSiiz.(m+n,AIm+nOnNI,n) composition was fixed at x = 0.33, m = 1.0, and n = 1.2 while the P-phase was fixed at z = 0.6.
We have kept the sample size constant in order to make it possible to compare the thermal shock properties for the different sialon materials.
EXPEFUMENTAL Three different types of sialon ceramics were designed and prepared, see Table I. The first type consisted of pure p-sialons (Si6.,A1,O,Ns-,) with zvalues in the range 0.6 to 3.0. In the second ones the same 0-sialon compositions were prepared but with an addition of intergranular glass phase. The added glass phase had the nominal overall composition Y I . ~ ~ S ~ Z . ~ Z(28 ~ AY;~ 56 ~ .Si;O 16 ~~ Al;.80 ~ N0;I20 .N Z~ in equivalent % [I]). The third type of sialon material consists of mixed alp-sialons with various amounts of glass added. For those materials the a-sialon (x = composition was fixed at Y0.~3Si~.~A1~.20~.2N~4.g 0.33, m = 1.0 and n = 1.2); the P-sialon composition was Si5.4A10.600.6N7.4(z = 0.6) and the glass phase composition was designed to have the same composition as for the p-sialons.
Table I List of prepared compositions. The different materials are referred to in the text as samples A to Y, GP = vol % glass phase.
I
I
I
Specimens were prepared from commercial Si3N4 (UBE, SN-ElO), AlN (H.C. Starck-Berlin, grade A), YzO3 (99.9 %, Johnson Matthey Chemicals Ltd.), A1203(Alcoa, A16SG), and SiOz (99.9 YO,<325 mesh, Johnson Matthey Chemicals Ltd.), and corrections were made for the amounts of oxygen present in the Si3N4 and AlN raw materials. The starting-material
147
mixtures were milled in water-free propanol for 24 hours in a plastic jar, using sialon-milling media. The dried powder mixtures were hot pressed in order to ensure full density, at a temperature of 1750 "C and a pressure of 30 MPa, in nitrogen atmosphere, for 60 min. However, two monophasic samples (B and C) had to be hot-isostatically pressed at a temperature of 1820 "C and a pressure of 200 MPa for 60 min in order to become fully dense, due to the complete or nearly complete lack of intergranular glass phase. Samples A, 0 and P were densified by using spark plasma sintering technique (SPS) at 1700 "C, 50 MPa for 3 min. Characterisation The densities of the sintered specimens were measured according to Archimedes' principle. Before the mechanical and microstructural studies, the specimens were carefully polished by standard diamond polishing techniques. The hardness (HVlo) and indentation fracture toughness (KI,) at room temperature were determined by means of a Vickers diamond indenter with a 98 N (10 kg) load, and the fracture toughness was evaluated according to the method of Anstis et al. [3] assuming a value of 300 GPa for Young's modulus for all compositions. Five indentations were made on each sample for the hardness measurements. The microstructures of the samples were investigated in a scanning electron microscope (SEM, JEOL 880) equipped with an energy-dispersive spectrometer (EDS, LINK ISIS) that allows detection of boron and heavier elements. Micrographs of fractured samples were taken in secondary electron mode (SE). To obtain the best contrast between different phases on polished surfaces, the micrographs were recorded in back-scattered electron mode (BSE) at an acceleration voltage of 20 kV. The amounts of intergranular glass phase were evaluated with an image-analysing package supplied with the LINK ZSIS system. The contrast difference between the sialon grains ( a and p) and the yttriumcontaining glass phase was accounted for. The estimates of the minimum and maximum amount of glass phase gave an error of +2%. A focusing X-ray powder diffraction camera of Guinier-Hagg type with CuLI radiation (A = 1.5405981 A) was used to record the X-ray powder diffraction (XRD) patterns, and powdered silicon (a = 5.430879 A at 25°C) was added as internal standard. The computer programs SCANPI [4] and PIRUM [5] were used to evaluate the recorded films. The latter program was also used to determine and refine the unit cell parameters. Experimental z-values of the p-sialon phase were obtained from the unit-cell dimensions of the samples, using the equations given by Ekstrom et al. [6]. The amounts of a- and p-sialon were estimated by comparison of the intensities of the two strongest XRD peaks of each phase (1 0 2 and 2 1 0 for a-sialon and 1 0 1 and 2 1 0 for p-sialon).
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Thermal shock measurements Cylindrical samples with a diameter of 12 mm and a thickness of 3.76 0.40 mm with parallel surfaces, were carefully polished on one side. Well-defined cracks were initiated with a Vickers indenter (Instron, model 856 1, High Wycombe, England) allowing variation of the load and the loading-, holding- and unloading time. The loading and unloading times were 10 s each, and the holding time at maximum load was 20 s. Four indents were made on each sample, and each indent generated four cracks defined as the distance from the centre of the indent to the crack tip (see Figure l), so that a total of sixteen cracks were initiated on each sample. The crack lengths were measured in an optical microscope (Olympus PMG3, Japan). In order to compare different samples more easily, the original crack length was held constant, e.g. around 100 pm. As a consequence, different samples had to be indented with different. loads (see Table 11). A vertical tubular furnace was heated to the pre-set starting temperature, initially 190 "C, which was then normally increased in steps of 100 "C. The sample was hoisted into the furnace where it was thermally equilibrated for 20 min and was then quenched in a 90 "C water bath.
+
I
Figure 1 Definition of the crack length used in thermal shock measurements. The test pieces were heated in air, but in the experiments presented below the temperature were not raised above 1000 "C, in order to avoid oxidation. A 90 "C water bath was used as quenching media. This temperature was kept constant with a thermostat. The quenching procedure was repeated from stepwise increasing heating temperatures. The crack growth was measured at each temperature step and the total percentage growth was calculated. The critical thermalshock temperature (AT,) was defined according to Anderson and Rowcliffe [2], as the temperature where 25% of the cracks have grown more than 10% of their original length. For sample M, a comparison was made between the re-use of the same sample for a whole quenching series and the use of a fresh sample for each new quenching temperature.
Sample
Thermal shock measurements Load Initial AT, (N) crack (“C) length
Mechanical properties K1c
(MPa*m”)
HV (GPa)
Glass (~01%)
Density (g/cm3)
z value
d(a+p)*
* The amounts of a- and B-sialon were estimated fiom comparison between the intensities of the two strongest XRDlines for the two phases respectively (1 0 2 and 2 1 0 for a-sialon and 1 0 1 and 2 1 0 for p-sialon). ** These samples contain the B-phase, implying that no (or only a small amount of) a-sialon phase has formed.
RESULTS
Microstructural evaluation Measured densities are listed in Table 11. All samples were according to SEM studies hlly dense and the obtained densities for the p-sialons were also comparable to previously reported ones [ 6 ] . SEM micrographs are given in Figure 2a-d for some selected materials representing the different materials investigated: p-sialon, p-sialon + glass, alpsialon, and alp-sialon + glass. The morphology of the materials without additional glass phase added is equiaxed. The materials containing additional glass phase show elongated grains, in accordance with previous studies of a- and p-materials [7]. When no extra glass phase is added the p-sialons with z 5 1.5 fracture intergranularly, while the more Al-rich p-sialons (z > 1.5) fracture transgranularly. Thus the strength of the P-grains decreases with increasing Al-content. Besides the a- and P-phases also Y2SiA105N was identified in some samples of the alp-sialons with an
additional glass content (samples T, U, and Y). This phase is known as the B-phase, which most often crystallises from oxynitride glass in yttrium containing sialons [8]. The presence of this phase is due to the difficulties in designing the composition of the liquid that is in equilibrium with the a- and P-phases at the sintering temperature. For samples T, U, and Y it thus seems that some of the yttrium designed for the asialon phase has been used up in the glass formation during sintering. The unit cell dimensions of the a-phase were the same in all samples indicating that the composition of the a-sialon is similar in all samples. The z-values calculated from the unit cell dimensions of the P-phase are given in Table 11. Those values show a slight variation from the aimed values.
Mechanical properties The hardness values (HVIo)and fracture toughness (K,,)for the different samples are given in Table 11. The z-value does not influence on the hardness for the p-sialons. Addition of extra glass phase increases the
149
hardness and the fracture toughness somewhat for low z-values, but the hardness for high-z samples do not improve with addition of glass. A general tendency for the a l p sialons is that both hardness and fracture toughness increase slightly with increasing a-sialon content. The hardness decreases with increasing glass content, but fracture toughness increases. The measured hardness as well as the fracture toughness values are consistent with literature data [9].
Al-content in the P-phase makes it easy to fracture upon thermal cycling. For the alp-sialons the best thermal shock resistance was found at a low fiaction of a.Increasing the fraction of a-phase decreases the thermal-shock resistance. Presence of glass slightly improves the thermal shock resistance and the best resistance is found at the highest glass content investigated (20 vol%), see sample Y in Figure 4. High fracture toughness normally correlates with a good thermal shock resistance. The critical thermal-shock temperature (AT,) as defined by Anderson and Rowcliffe [2] is evaluated. AT, decreases with increasing z-value both for samples with and without added glass, see Figure 5a. For low zvalues the addition of glass has a pronounced effect on AT,, whereas for high z the glass has no influence. A high amount of glass seems to have a more pronounced effect for the p-sialon than for the alp-sialon, see Figure 5b. For the a l p sialons the AT, value decreases with increasing fraction of a-sialon, see Figure 5c.
0-0
5
v-v
h
4-A
Sample E Sample A Sample M
d 6 0 3
e
M
"
200
0
400
600
800
1000
A T ("C)
Figure 2a-d SEM micrographs of (a) fractured surface of sample A, z=O.6 (SE mode). (b) polished surface of sample M, z=O.6, 20 vol% glass. The p-sialon grains are black or dark grey and the glass phase is white (BSE mode). (c) Fractured surface of sample P, a/a+p=0.3 (SE mode). (d) Polished U, surface of sample ala+p = 0.3, 10 % glass. The p-sialon grains are black, the a-sialon grains are grey and the glass phase is white (BSE mode). Thermal shock properties For the p-sialon materials it is clear that samples with low z-values (low Al-content) are more resistant to thermal shock than those with higher z-values. For samples with additional glass added, the thermal shock resistance is improved for all z-values except for the sample with the highest z-value investigated, z = 3.0; see Table I1 and Figure 3. It is obviously so that a high
150
Figure 3 Crack growth in percent plotted versus AT (see text) for some selected p-sialon samples.
loo,, v-v 0-0
h
40
f
I 0-
-0
Sample R Sample X Sample P
I A T ("C)
Figure 4 Crack growth in percent plotted versus AT for some selected a/p-sialon samples.
For sample M (z = 0.6, 20 % glass) that show the best thermal shock resistance a comparison was made between the re-use of the same sample for a whole quenching series and the use of fresh samples for each new quenching temperature. The two series gave very similar results (see Figure 6). This concludes that the indentation method used is reliable for qualitative comparison of thermal shock properties of sialon materials.
A
' 0
1
800
I
15,
-
0 0
200
400
600
800
1000
AT (OC)
200
1
Figure 6 Crack growth in percent plotted versus AT for sample M (z = 0.6, 20 % glass); I The same compact has been used at each quenching temperature. I1 A new compact has been used at each quenching temperature.
0 I
1 .o
0.5
1.5
2.0
2.5
3.0
CONCLUSION
z value
1000
(b) 800
/
,z = 3.0 -----a
5
0
10
15
20
25
30
Glass (voIYo)
20
40
60
80
d(a+P)
Figure 5a-c The critical thermal-shock temperature difference (AT,) measured as a function of (a) the nominal z-value for p-sialons with and without additional glass phase added (b) the measured amount of intergranular glass phase for p- and alpsialons (c) the nominal amount of a-sialon for alpsialons.
100
An indentation-quench method based on Vickers cracks for measuring thermal-shock properties has been applied to sialon materials. The method allows the use of the same sample during a series of increasing temperatures. The percentage crack growth is measured at each temperature-quenching step, and the statistics are improved by making several Vickers indents on the same sample. For p-sialons (Si6.zAl,0,N8.z) the z- value has been varied in the range 0.6 I z I 3.0 and it is found that the thermal shock resistance is best at low z-values. The thermal shock resistance improves even further with addition of yttrium containing glass phase. However, for high z-values the thermal shock properties do not improve by addition of glass. For the dp-sialons the z-value of the P-phase were kept at 0.6 and the a-phase ( Y , S i ~ ~ ~ m + n ~ l m + n O n N ~ ~ n , ) had a composition of x = 0.33, m = 1.0, and n = 1.2. The thermal shock resistance improves with increasing fraction of P-phase. Also for these materials the thermal shock resistance improves with addition of glass phase. High fracture toughness normally correlates with a good thermal shock resistance. The best thermal-shock properties were found for a p-sialon with a z-value of 0.6, containing 20 vol% glass (ATc = 900 "C). The poorest resistance to thermal-shock was found for a p-sialon with z = 3.0 (ATc = 100 "C), and the thermal-shock properties of psialons with high z-values were not improved by addition of extra glass phase.
151
ACKNOWLEDGEMENT
This work has been performed within the Inorganic Interfacial Engineering Centre, supported by the Swedish National Board for Industrial and Technical Development (NUTEK) and the following industrial partners: Erasteel Kloster AB, Ericsson Cables AB, Hoganas AB, Kanthal AB, OFCON Optical Fiber Consultants AB, Sandvik AB, Seco Tools AB and Uniroc AB. REFERENCES
T. Ekstrom and M. Nygren: “Sialon ceramics”, J. Am Ceram. SOC.75 (1992) 259-276. T. Andersson and D.J. Rowcliffe, “Indentation thermal shock test for ceramics”, J. Am. Ceram. SOC.,79 (1996) 1509-1514. G.R. Anstis, P. Chantikul, B.R. Lawn, and D.B. Marshall: ”A critical evaluation of indentation techniques for measuring fracture toughness: I, Direct crack measurements”, J. Am. Ceram. SOC.,64 (1981) 533-538. K.E. Johansson, T. Palm and P.E. Werner: “An automatic microdensitometer for X-ray diffraction photographs”, J. Phys. E. Sci. Instrum., 13 (1980) 1289-1291. P.E. Werner: ”A Fortran program for leastsquares refinement of crystal structure cell dimensions”, Arkiv fdr kemi”, 3 1 (1969) 5 13516. T. Ekstrom, P.O. Kall, M. Nygren, and P.O. Olsson: ”Dense single phase p-sialon ceramics by glass encapsulated hot isostatic pressing”, J. Mat. Sci. 24 (1989) 1853-1861. C.M. Hwang and T.-Y. Tien: “Microstructural development in silicon nitride ceramics”, Mater. Sci. Forum 47 (1989) 84-109. H. Mandal, D.P. Thompson and T. Ekstrom: “Reversible a-P sialon transformation in heat treated sialon ceramics”, J. Eur. Ceram. SOC.12 (1993) 421-429. T. Ekstrom: “Effect of composition, phase content and microstructure on the performance of yttrium Si-Al-0-N Ceramics”, Mat. Sci. Eng., A109 (1989) 341-349.
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CORROSION OF NONOXIDE SILICON-BASED CERAMICS IN A GAS TURBINE ENVIRONMENT H. Klemm*, Chr. Schubert", Chr. Taut**, A. Schulz***, G, Wetting**** "Fraunhofer-Institute for Ceramic Technologies and Sintered Materials, IKTS Dresden, Winterbergstr. 28, D-01277 Dresden, FRG **Siemens AG, KWU, Wiesenstr. 35,45473 Mulheim, FRG ***Institut fut Thermische Stromungsmaschinen, Uni Karlsruhe, Kaiserstr. 12,76128 Karlsruhe, FRG ****Ceramics for Industry, CFI GmbH & Co KG, Oeslauerstr. 35,96472 Rodental, FRG
In the present study the corrosion behavior of nonoxide silicon-based ceramic materials was investigated. Hot gas tests were conducted on silicon nitride and silicon carbide materials in an atmosphere similar to that in a gas turbine. While some materials displayed a high degree of microstructural stability, all materials suffered surface degradation during the rig test. The oxidation surface layer of mainly Si02, which is essential for the oxidation protection of nonoxide materials because it induces a passive, diffusion-controlled oxidation mechanism, was found to be degraded by evaporation processes involving volatile silicon hydroxides. The partial pressure of steam was found to be the most important factor governing these processes.
found to have a great influence on the oxidation and corrosion behavior of the ceramic materials [7,8]. The main problem preventing silicon-based nonoxide ceramic materials from being applied in such environments was found to be the formation and evaporation of silicon hydroxides (Si(OH)4) described by Opila and coworkers, who investigated the corrosion behavior of CVD Sic in humid environments. These processes were observed to be enhanced in burner rig tests with severe environmental conditions, such as higher gas and vapor pressure and higher flow rate 19,101. In the present study nonoxide ceramic materials based on Si3N4 and S i c were tested in a real hot gas environment. The hot gas tests were carried out under various environmental conditions (temperature and water vapor pressure).
INTRODUCTION
EXPERIMENTAL, PROCEDURES
One of the main requirements necessary for ceramic materials to be applied confidently in gas turbines is the long-term stability of all relevant properties at elevated temperatures. Recently, nonoxide ceramic materials based on Si3N4 and Sic, which are featured by their superior mechanical, chemical and thermophysical properties, have been developed [ 1-41. Additionally, significant progress has been made in improving the long-term stability of these materials at temperatures up to 1500°C for application-relevant times of more than 10 000 h under laboratory conditions [5,6]. The main feature of these materials, and the key factor in stabilizing the mechanical properties, is their microstructural stability, achieved by minimization of oxidation-enhanced diffusion processes in the bulk of the material during long-term service at elevated temperatures. Nevertheless, employment of the results of these laboratory tests for gas turbine applications, for example, is of limited use due to the severe environmental conditions in a gas turbine combustor. The high pressure and flow rate water vapore and other corrosive components (Ca2+, V5+) of the hot gas were
The studies were conducted on a solid-state-sintered S i c (S Sic) with B and C as sintering additives, a liquid-phase-sintered (hot pressed) Si3N4 (SN), a Si3N4-MoSi2(SNMo) composite with Yb203 as the sintering additive and a Lu203-containing silicon nitride material (KY) produced by Kyocera (Japan) SN 281. Relevant properties of these materials are summarized in Table 1.
ABSTRACT
Table 1: Mechanical properties of materials investiga4-point bending strength at ted: oRTand oI4: room temperature and 1400°C; KIc: SENB F0.15mm; E: 200 MPa, 1400"C, 100 h.
S Sic
3.18
430
430
2.8
1.5~10-~
SN
3.53
840
720
8.1
1.7~10-5
SNMo
3.69
830
710
8.0
1.6~10-~
KY
3.39
660
530
6.2
~.OXIO-~
153
The corrosion resistance of the materials was investigated using bending bars in a high-temperature, high-pressure burner rig test apparatus developed by the Institute for Thermal Turbomachinery, ITS, University Karlsruhe. Details about the test facility are described in a former report [ 111. The ceramic materials underwent three burner rig tests for 100 h by changing temperature and water vapor pressure in the test facility. The test conditions are summarized in Table 2.
Table 3: Weight loss and evaporation rate of the materials tested in the high-pressure burner rig. Mat.
Table 2: Test conditions
I
I
I
I
1. Test 2. Test 3. Test K / m g / Am K / m g / Am K / mg/ cmZh cmZh Am cm2h
I
1
0.08 1.Test
I
2. Test
I
3. Test
Temperature / "C
1400
1300
1400
Pressure / bar
5
5
5
Flow speed / m / s
50
50
50
Part. steam pressure
0.9
0.9
0.4
I
After the tests, the weight loss of each bending bar was determined. Changes in the phase composition at the surface and the bulk region below the oxidation layer were investigated by XRD. To assess the damage resulting from the high-pressure burner rig tests, the room-temperature bending strengths of the specimens were obtained and compared with the strengths of the as-fabricated samples. Information about the microstructural alterations was obtained through observation of the microstructures (polished and CF, plasma-etched cross sections) and the surfaces after corrosion in the SEM.
RESULTS The results of the burner rig tests are summarized in Table 3. All materials exhibited a weight loss after the tests; this was found to be dependent on the test parameters used. From a comparison of the data, it can be seen that the Si3N4 materials behave alike. The surface layer, consisting of silica and the disilicates usually formed from the sintering additives during oxidation of siliconbased ceramic materials, was found to be degraded by the corrosive environment in the burner rig. The measured weight loss was due to the evaporation of silicon hydroxides, mainly Si(OH),, and the R2Si20, (R = Yb or Lu) formed during oxidation of the Si3N4 materials, independent of the sintering additive. As a consequence of the formation and evaporation of Si(OH),, the protective SiO, layer was destroyed during the test, leaving a very rough layer consisting of only disilicate grains which were partially spalled off due to the low stability of this surface layer, as shown in Fig. 1. From an application point of view, the materials' surface degradation is quite high. The weight loss of about 10 % observed for the S S i c material at 1400°C and a partial steam pressure of 0.9 bar led to a material loss of about 90 pm in 100 h under these conditions. After a time of about 10 OOO h or higher (relevant times for application in energy production), an unacceptably high material loss of about 9 mm would be observed.
154
I
KY
Fig. 1:
0.09
Surface layer Si N material with Yb203 as 3 4 the sintering additive after high-pressure burner rig test No. 1 (100 h, 1400°C, 0.9 bar partial steam pressure).
With respect to the microstructural stability of the bulk material, the results of the burner rig tests confirmed those of oxidation tests performed and reported previously [12,13]. During the rig test the microstructure changed due to oxidation processes that were dependent on the diffusion of oxygen into the upper region of the bulk and reaction of the oxygen with Si3N4 solved in the grain boundary phase. As in oxidation tests performed in previous studies, the microstructures of the Si3N4 materials and the Si3N4-MoSi, composites were found to be different after the test. In the Si3N4 materials, typical oxidation damage, i.e., surface and grain boundary phase inhomogeneities and pores up to the middle of the materials, was observed. This was the consequence of the reaction of the diffusing oxygen in the grain boundaries of the Si3N, materials, resulting in the formation of SiO, and leading finally to a grain boundary phase supersaturated with silica in the surface region of the Si3N4 material. The relaxation of this silica-rich grain boundary phase was considered to be the reason for the changed microstructure in the Si3N, material. The Si3N4-MoSi2behaved in a different manner. Similar to the S I ~ N ,materials, some of the oxygen diffused through the grain boundaries and triple junctions and penetrated into the upper region of the bulk material to react with the Si3N4 solved in the oxynitride glass. Instead of the SiO, found in the Si3N4 materials, crystalline Si20N2 was found to be the oxidation product in the composite material. The relaxation made necessary by the oxidation processes in the grain boundary
phase in the upper bulk region (formation of a silicarich grain boundary phase) was achieved by the immediate crystallization of Si20N,. As a consequence, the grain boundary composition remained nearly constant and the microstructure of the Si3N4-MoSi2 composite was stabilized after the burner rig test. The S S i c material was found to have a lower rate of oxygen diffusion. In this case, all of the oxygen reacted at the interface between the oxidation layer and the bulk, which consequently was not significantly changed or damaged. Due to the microstructural degradation in the bulk, the residual strength of the Si,N, materials was found to be much more degraded in comparison to the Si3N4-MoSi2 composite and the S SIC material. The results of the bending tests after the burner rig test are summarized in Table 4.
the conditions of the high-pressure burner rig tests is provided in Table 5.
Table 4: Residual strength of the ceramic materials after rig exposure; each value represents the average of at least 4 specimens. Mat.
1. Test
2. Test
3. Test
OTest
OT'Oo
OTest
OT/Oo
OTest
MPa
%
MPa
%
MPa
%
OT/Oo
S Sic
350
81
390
90
450
100
SN
400
48
490
58
520
62
SNMo
660
80
680
82
700
84
KY
320
48
430
65
420
64
As already mentioned, the weight loss observed for the S S i c material was the consequence of the formation and evaporation of silicon hydroxides from the surface of the ceramic material. As opposed to the surface layer found after oxidation in air with some crystobalite crystallites in a flat oxidation layer of amorphous silica (Fig. 2A), the surface layer of the S S i c after the rig test was found to be rough, indicative of the evaporation of the silicon hydroxides during the test (Fig. 2B). The weight loss obtained was dependent on the test conditions. Both temperature (comparison of Tests 1 and 2) and water vapor pressure (Tests 1 and 3) influence the degradation of the materials. As indicated in the literature, the water vapor pressure was found to have the highest influence on the evaporation processes at the surface of the ceramic materials. When the partial steam pressure was reduced from 0.9 to 0.4 bar, the weight loss was reduced by three or four times for all materials studied. Finally the data obtained in this study for the S S i c material are compared with the data of Robinson and Opila [9,10]. The materials investigated in the studies performed by Robinson and Opila, solid-state-sintered S i c with B and C and CVD Sic, are not the same as those investigated here; however, they can be used for comparison due to their similar oxidation behavior, i.e., the formation of an oxidation surface layer of pure SO2. Regarding the test conditions, however, several differences should be considered. For that reason, only a rough comparison of the data is possible. A summary of
Fig. 2:
Comparison of the oxidation surface of the S S i c material after oxidation in air (A) and corrosion in burner rig test No. 1 (B).
In Table 6 a comparison of the weight loss rates measured at 1400°C in both studies are given. Additional weight loss rates calculated by the model presented by Opila are used for comparison. Assuming nearly similar test conditions, calculated weight loss rates of Si(OH), for this study can be obtained by using the expression proposed by Opila:
with J as the weight loss rate in mg/cm2h, v the linear gas velocity, ptotalthe total pressure and PH20 the partial vapor pressure.
1. Test
2. Test
Ref.
Temp./"C
1400
1400
1400
pressurehar
5
5
6.3
fuel
nat. gas
nat. gas
jet fuel
airtfuel
2.0
2.0
1.1
flow speet / m l s
50
50
20
vapor press. / bar
0.9
0.4
0.6*
155
Table 6: Comparison of measured Ky and calculated & weight loss rates for S i c of this study with data published by Robinson [9] and Opila [lo] at 1400°C.
0.062
In principle, all data are of the same order of magnitude. The calculated values were found to be lower than the measured data in all cases, indicating the very complex character of the corrosion processes, e.g., the presence of other silicon hydroxides than %(OH),, as proposed by Opila. As a consequence of the addition of steam, which produced the highest partial vapor pressure, the highest weight loss rate was observed in Test 1 of this study. The weight loss rates from Test 3 and from tests in the literature in which a similar partial steam pressure was used seem to be more comparable. While the calculated data were found to be similar, the higher weight loss rate of Test 3 in this study may be the consequence of a higher influence of the flow speed in the burner rig.
CONCLUSIONS Silicon-based nonoxide ceramic materials were studied in a corrosive atmosphere similar to that in a gas turbine at high temperatures. Under these conditions, the microstructural stability of the bulk material can be stabilized by e.g. Si3N4-MoSi2composites by a mechanism similar to the mechanisms found during oxidation of these materials in air. Unlike oxidation in air, the tests in the burner rig caused all materials to suffer surface degradation. The weight loss measured was the consequence of the formation and evaporation of silicon hydroxides, mainly Si(OH)4.. The oxidation surface layer of mainly SiO,, which is essential for the oxidation protection of nonoxide materials because it induces the passive, diffusion-controlled oxidation mechanism, was found to be degraded. As a result of this, the material showed a higher rate of oxidation. The main factor influencing the evaporation processes seems to be the partial vapor pressure. In spite of the different burner rig set-ups and test conditions, the data obtained in this study were similar to the results for corrosion of CVD S i c reported in the literature. However, additional studies are required for a quantitative interpretation of the influence of test parameters such as pressure, partial vapor pressure and flow speed.
ACKNOWLEDGMENTS We acknowledge the interesting discussions with E. Opila, R. Robinson and J. Smialek, NASA Glenn Research Center, Cleveland. The studies were carried out on behalf of Siemens KWU Mulheim.
156
REFERENCES G. Pezzotti, "Si,N,/SiC-Platelet Compo- site without Sintering Aids: A Candidate for Gas Turbine Engines", J. Am. Ceram. SOC.76 [5] 1313-20 (1993). M.N. Menon et al, "Creep and Stress Rupture Behavior of an Advanced Silicon Nitride", J. Am. Ceram. SOC.77 [5] 1217-41 (1994). K. Watanabe et al, "Development of Silicon Nitride Radial Turbine Rotors", pp. 1009-1016 in Proc. 4th Int. Symp. Ceram. Mater. & Engines, Ed. by R. Carlson, T. Johanson and T. Kahlman, Elsevier London ( 1991). C. J. Gasdaska, "Tensile Creep in an in Situ Reinforced Silicon Nitride", J. Am. Ceram. SOC. 77 191 2408-18 (1994). H. Klemm, M. Herrmann, Chr. Schubert, "High Temperature Oxidation and Corrosion of Silicon-Based Nonoxide Ceramics", J. of Engineering for Gas Turbines and Power, 122 [ 11 13-18 (2000). H. Klemm, Chr. Schubert, "Long-term stability of non-oxide ceramics in an oxidative environment at temperatures above 14OO0C", 24 Annual Cocoa Beach Conference and Exposition, 23. - 28.01. 2000, Ceram. Eng. & Sci. Proc., 21 (3) (2000). E.J. Opila, R.E. Hann, "Paralinear Oxidation of CVD S i c in Water Vapor", J. Am. Ceram. SOC., 80 [l] 197-205 (1997). E.J. Opila, D.S. Fox, N.S. Jacobson, "Mass Spectrometric Identification of Si-0-H from the Reaction of Silica with Water apor at Atmospheric pressure", J. Am. Ceram. SOC.,18 [4] 1009-12 ( 1997). R.C. Robinson, J.L. Smialek, "Sic Recession Caused by SiO, Scale Volatility under Combustion Conditions: I, Experimental Results and Empirical Model," J. Am. Ceram. SOC.,82 [7] 1817-25 (1999). E.J. Opila, J.L. Smialek, R.C. Robinson, D.S. Fox, N.S. Jacobson, "Sic Recession Caused by SiO, Scale Volatility under Combustion 11, Thermodynamics and Conditions: Gaseous-Difhsion Model," J. Am. Ceram. SOC., 82 [7] 1826-34 (1999). D. Filsinger, A. Schulz, S. Wittig, C. Taut, H. Klemm, G. Wotting, "Model Combustor to Assess the Oxidation Behavior of Ceramic Materials under Real Engine Conditions," ASME Turbo Expo '99, Indianapolis, USA 1999, 99-GT-349. H. Klemm, K. Tangermann, Chr. Schubert, W. Hermel, "Influence of Molybdenum Silicide Additions on High-Temperature Oxidation Resistance of Silicon Nitride Materials", J. Am. Ceram. SOC.79 [9] 2429-35 (1996). H. Klemm, Chr. Schubert, "Si3N4-.MoSi2 Composite with Superior Long-Term Oxidation Resistance at 1500"C", submitted to J. Am. Ceram. SOC.
(ptSpecies
MECHANICAL PROPERTIES AND WEAR BEHAVIOUR OF DIFFERENTLY MACHINED SILICON NITRIDE AND SILICON CARBIDE CERAMIC SURFACES T. Hollstein*, W. Pfeiffer, R. Zeller Fraunhofer-Institut f i r Werkstoffmechanik (IWM) Wohlerstr. 11, D-79108 Freiburg, Germany
ABSTRACT Machining and wear of ceramics lead to specific surface topographies, and to damage, micro-plastic deformation, and residual stresses in the surface layers. Their effects on sliding wear under boundary friction conditions and rolling wear without lubrication are investigated. The analysis of these tribological systems show, that their characterisation and assessment - based on the main failure and wear mechanisms - allow in most cases for a straight-forward optimisation of the service behaviour. The life time of components operating with sliding or rolling contact under mixed friction conditions is increased at most by all those parameters, which generate compressive residual stresses and enhance micro-hydrodynamic conditions. A consequent design of the topography and an increased pressure on the lubricant may be most effective.
INTRODUCTION Machining and friction influence ceramics in bearings predominantly in surface-near regions. The load capacity and the wear behaviour of ceramic roller or sliding bearings depend on the machining parameters and the surface treatments applied during fabrication. The assessment of the tribological situation and the optimisation of component performance need the identification of the main wear effects and the improvement of the related near-surface characteristics of the components. These near-surface characteristics are mainly microstructure, topography, surface strength, and machining induced damage and residual stresses. The effects of the topography on the sliding and rolling wear behaviour of ceramics under mixed or fully lubricated conditions have been studied during the last years [l-31. Less attention has been paid for damage and residual stresses - which may be summarised as surface integrity - and their effect on the wear behaviour. It is well known that machining introduces both, damage and residual stresses, in the near surface layers of ceramics. Machining induced damage reduces the strength, while residual stresses, which are usually
compressive, leads to an improved strength. Both strength behaviour can be observed depending on which effect dominates [4-61. A very pronounced effect of near surface damage and residual stresses should be observable in contact loading conditions due to rolling or sliding, because the maximum load stresses occur near the surface. Of special interest are systems like water-lubricated face-seals or ceramic roller bearings designed to operate without lubrication. Theoretical and experimental investigations show that the failure behaviour of these bearings is controlled by peak stresses near the surface [3, 71. As a consequence, one should expect that the load bearing capacity of a ceramic component can be increased by introducing compressive residual stresses in the near surface region. It has been shown that - besides machining - this can be achieved by a shot-peening process [8]. Thus, the aim of the investigations was to understand better the wear behaviour of differently machined ceramics by applying surface sensitive characterisation methods like advanced X-ray diffraction methods and ball on plate tests.
EXPERIMENTAL PROCEDURE MATERIALS AND MACHINING The materials investigated where a commercial sintered silicon carbide (EKasicBD, Elektroschmelzwerke Kempten, ESK) and a commercial silicon nitride (GPSN, SN-N3208, Ceramics For Industry, CFI). The most important material data are given in Table 1. SN-N3208 Material parameter Char. tensile strength oo 900 MPa Weibull modulus 18 Young’s modulus 310GPa Fracture toughness KI, 4.7 MPadm (pre-cracked bend test) Vickers Hardness (HV10) 15.2 GPa
EKasicBD 350 MPa 10
400GPa 3.2 MPadm
> 24 GPa
Tab. 1: Material data of the investigated silicon nitride and silicon carbide.
157
Samples with cylindrical geometry were manufactured by grinding and lapping for the wear tests from the near-net-shaped sintered workpieces, see Figure 1. For face sliding wear tests with SIC specimens the relevant plane surfaces were finished by different lapping procedures (F500- and F220-B4C-abrasives), ultrasonic-assisted lapping (F70-Sic-abrasives) and polishing (diamond abrasives, 0.25 pm and 2-4 pm grain sizes). The surfaces for the Si3N4specimens were finished by a laser-beam shaping method. The circumference of the silicon nitride rings were finished using conventional grinding (D64-diamand wheel), grinding + partial polishing, and a laser assisted turning process (LAM) . The machining of the specimens was performed by the Fraunhofer-Institut fur Produktionstechnologie (IPT), Aachen.
E E 0
t
cant. The sample geometry and a sketch of the experimental arrangement are shown in Figure 1. More details of the experimental procedure are given in [9]. Rolling wear tests were performed in a so-called >>Amslerc< experimental set-up. Contact loads in the range of 1-3 GPa and slip ratios in the range of 1-3 % were used. No lubricant was used. The sample geometry and a sketch of the experimental arrangement is shown in Figure 1. More details of the experimental procedure are given in [9].
X-RAY DIFFRACTION ANALYSIS Residual stresses and micro-plastic deformations due to machining and wear processes were determined by X-ray diffraction techniques. The most important measurement parameters are given in Table 2. ceramic Si3N4 Sic
lattice plane (41 11-P { 213)-6ha
-45" I
ceramic Si3N4 Sic
Radiation CrKa FeKa
penetration depth zeff = 8 pm 0.6 pm c zeff c 18 pm
y-range
w I+45"
-69.6" 5 y SO"
Tab. 2: Measurement parameters for residual stress and micro-strain evaluations
J.
10 mm Fig. I : Specimen geometry and loading situation used for the sliding wear tests (left) and the rolling wear tests (right).
NEAR-SURFACE STRENGTH The near-surface strength of the machined specimens was determined using a ball-on-plate test. During this test the specimens are placed on a flat stiff area and loaded by an A1203 ball with a diameter of 10 mm. The load is increased stepwise until a typical cone-crack appears, which follows the maximum tensile stresses occurring at the surface just at the border of the circular contact zone. Because of the statistical behaviour of ceramics, the load at fracture varies from test to test within a certain scatterband. In this paper, the fracture load represents a probability of 50%. For more details of the ball-on-plate test see [7,8]. WEAR TESTS Face sliding wear tests were performed under boundary friction conditions using a contact load of 2 MPa, sliding speeds of 0.3 m/s and water as a lubri158
Due to the curved geometry of the rolling wear samples only the average values of near-surface residual stress components could be determined using the conventional sin$-method [ 10, 1 I]. The flat surfaces of the sliding wear samples allowed determination of the depth distributions of micro-plastic deformation and residual stresses within the first 3 0 p m using the inverse Laplace transformation method [ 1 11.
RESULTS AND DISCUSSION SLIDING WEAR OF SILICON CARBIDE In systems operating under boundary friction conditions (like water-lubricated face-seals during start/stop operations) the wear behaviour can be related to the surface integrity of the components in addition to micro-hydrodynamic effects. This may be illustrated by Figs. 2 - 4, in which the near surface strength, the depth distributions of residual stresses, and the material removal due to sliding wear of differently machined sintered silicon carbide specimens are shown. A steep increase in surface strength of nearly a factor of 2 is obtained by increasing the >>roughness<< of the machining procedure (Fig. 2): the lowest surface strength for the smoothest polishing with finest grains and the highest surface strength for the roughest ultrasonic lapping with coarse grains. Although rough machining conditions like lapping with coarse grained abrasives may introduce more damage than polishing processes, the near-surface strength is increased due to the machining induced high and relatively deepreaching compressive residual stresses (Fig. 3). Obvi-
ously, these compressive residual stresses are able to reduce the wear rate after the roughness peaks have been removed (Fig. 4). gOOiratureload in N
700
7-
600
,
500
400 300
ing to a higher wear resistance, lead to this lower wear rates. During the first stage of test up to 1 km wear of US oscillation lapped surfaces reduces the roughness peaks at about lpm. Thereafter only very few wear is observed during the stages of test up to 10 km. This can be attributed to the relatively far reaching compressive residual stresses for these specimens, which are still present after the removal of about 1.5 pm of the surface layer (Fig. 3) and to the existence of static lubrication reservoirs..
200 100
0
\
polished
lapped US-lapped F500 F280 polished lapped 2-4prn F220
0.25pm
Machining Figure 2: The near surface strength of differently machined S i c specimens measured by the ball on plate test
-100
-200 -300 -400 0.005 0.01 0.015 Depth in m
0.02
Figure 3: The depth distributions of the residual stresses of differently machined Sic specimens
Layer removal in yrn
10
1
-.
0 1.
polished polished lapped 0,25pm
2-4pm
F500
US swing lapped
0
2
4
6
8
10
12
14
16
Sliding distance in krn
I
0
100
4 2
Figure 5: Comparison of the wear of laser-beam structured surfaces with ultrasonic swing lapped surfaces of S i c specimens
loo,Residualstress in MPa
0
m 5
lapped US-lapped F220
F280
Machining Figure 4: The wear rates of differently machined S i c specimens. Each bar represents a mean value from three specimen pairs.
The polished specimens show very high wear rates at early stages of wear due to adhesive failure followed by a continuous material removal. The lapped specimens show significantly lower wear rates. Both effects, lubrication reservoirs due to higher surface roughness and especially high compressive residual stresses lead-
A comparison of the wear of laser-beam structured surfaces with ultrasonic swing lapped surfaces of Sic specimens is shown in Figure 5. The shaping by laser beams leads additionally to a reduction of wear due to the lubrication reservoirs in the first stages of test. But this beneficial lubrication effect seems to be present only for a limited period in contrast to the far reaching beneficial effect of the compressive residual stresses. The corresponding microstructures of an ultrasonic oscillation lapped surface and a laser-beam shaped surface of Sic specimens after a sliding distance of 10 km are shown in Figure 6.
MACHINING AND ROLLING WEAR OF SILICON NITRIDE Silicon nitride ceramics are appropriate materials for roller bearings designed for an operation without lubrication, e.g. for vacuum applications. Hence, silicon nitride balls and rings are being investigated under rolling contact and wear situations. The main effects of rolling wear on silicon nitride, generation of cracks and compressive residual stresses, are compiled in Figure 7. For different slip ratios the development of the residual stress components parallel and transverse to the rolling direction is shown. The test started with no significant machining induced residual stresses as the samples were in a fine-ground and partially polished condition. With increasing number of cycles, high compressive residual stresses are created by the cyclic contact loading of the surface. The tangential loading of the near-surface layers seems to be the most important driving force for the residual stress development. This can be concluded from the fact, that the amount of residual stresses is higher parallel to the rolling direction and from the higher residual stresses of the sample tested at a higher slip ratio.
159
being simultaneously introduced by the cyclic contact loading stabilise the surface layers by crack closure effects and maintain the performance up to 100,OOO cycles. The effect of machining induced residual stresses on the rolling wear behaviour of silicon nitride is shown in Figure8. Specimens in the ground condition (D64 wheel) with machining induced compressive residual stresses in the range of 400MPa and laser assisted turned (LAM) samples with negligible residual stresses were run at 1.5 GPa and 3 % slip under dry conditions.
100 200 300 400 Thousend of cycles of overrolling
500
Figure 6: Ultrasonic oscillation lapped surface (top) and laser-beam structured surface (bottom) of Sic specimens after a wear distance of 10 km n
-
I LUU
20000
40000 60000 80000 100000 Cycles of overrolling
Figure 7: Development of residual stress components in silicon nitride, parallel (RD) and transverse (TD)to the rolling direction (contact load = 3 GPa, slip ratios 1 % and 3 %, dry rolling) and micro-graph showing cracks in the 2 mm broad raceway after 300 cycles.
After about 30,000 cycles the residual stress level saturates and is nearly constant up to 100,000 cycles, when dramatic failure processes start (the sample tested at 1 % slip failed early due to a failure of the wear testing machine). The beneficial effect of the rolling-induced residual stresses on the life time can be imagined from the micro-graph of the sample tested at 3 GPa and 3 % slip. Already after 300 cycles large cracks are present. They are oriented transverse to the rolling direction and extend across the entire raceway. Apparently they follow the direction of the maximum tensile residual contact stresses. Obviously the compressive residual stresses
160
Fig. 8: Development of residual stress components parallel to the rolling direction in ground and in laser-assisted machined (LAM) silicon nitride due to rolling wear (contact load = 1.5 GPa, slip ratios 3 %, dry rolling) and micro-graphs showing cracks in the raceway after 5000 cycles.
Up to 20,000 cycles no significant change of the grinding induced residual stress state is observed whereas compressive residual stresses are developed in the LAM-sample. The micro-graphs in Figure 8 indicate that in the early stages of the cyclic contact loading the grinding induced residual stresses obstruct the development of cracks in the surface layer. After 20,000 cycles the residual stresses begin to diminish in both samples due to abrasive surface layer removal. At later stages of cyclic loading, the effect of the layer removal on the residual stress states is compensated by rolling induced residual stresses. This process is more evident for the LAM-sample as less subsurface damage is present than in the ground sample.
CONCLUSIONS The investigations have shown that finishing processes, abrasive wear and cyclic contact loading introduce substantial micro-plastic deformation and compressive residual stresses into the surface layers of high-strength ceramics. Compared to rough machining conditions using high material removal rates, relatively thin surface layers are affected. Thus, advanced X-ray diffraction methods for depth probing and surface sensitive strength tests like the ball-on-plate test are needed to determine and to assess the surface integrity. A significant beneficial effect of the residual stresses on the life time is obtained for sliding applications of silicon carbide under boundary friction conditions. After the removal of the topographic peaks during wear in the early stages surface layer with compressive residual stresses due to rough machining are able to withstand abrasive wear better than smoothly polished surfaces containing no significant machining induced stresses. The near-surface residual stress investigations performed on cyclic contact loaded silicon nitride have shown the development of surprisingly high compressive stresses up to the GPa-range. In addition, the cyclic contact loading introduces severe damage in the very early stages of loading. However, the specimens show acceptable life-times due to the crack closure effect of the compressive residual stresses. Compressive residual stresses due to moderate grinding conditions shift the generation of damage to higher cycle numbers.
ACKNOWLEDGEMENT The machining of the samples was performed by the Fraunhofer-Institut fur Produktionstechnologie (IPT), Aachen. The investigations were sponsored by the German Science Foundation (DFG).
Work, The Institution of Engineers Australia, (1998) 201 - 206. [41 W. Pfeiffer, E. Sommer, ‘Bewertung der Festigkeitseigenschaften oberflachen-bearbeiteter Keramikbauteile’, Werkstoffkunde, DGM Informationsgesellschaft mbH, Oberursel, (199 1) 575-584. [51 W. Pfeiffer, T. Hollstein, ‘Damage determination and strength prediction of machined ceramics by X-ray diffraction techniques’, Machining of Advanced Materials. NIST Special Publication 847, United States Department of Commerce, (1993) 235-245. [6] W. Pfeiffer, T. Hollstein, ‘Characterization and Assessment of Machined Ceramic Surfaces’, 2nd Intern. Confer. on Machining of Advanced Materials (MAM), VDI-Verlag, Dusseldorf, VDIBerichte Nr. 1276, (196) 587-602. [7] M. Rombach, ‘Experimental and Theoretical Investigations on Plastic Deformation and Brittle Cracking in a Ball-on-Plate Contact of Ceramic Materials’, Ceramics Charting the Future (Edit. P. Vincenzini, Techna Srl), (1995) 1055-1064. [8] W. Pfeiffer, M. Rombach, ‘Residual Stresses and Damage in Ceramics Due to Contact Loading’, Proc. of the Fifth Int. Conf. on Resid. Stresses, Linkoping, Univ., Sweden, (1998). 1,302-307. [9] R. Zeller, T. Hollstein, ‘EinfluO der Oberflkhenbearbeitung von Hochleistungskeramiken auf das GleitverschleiBverhalten im Grenzreibungsgebiet und auf das WiiIzverschleiBverhalten’, Technische Keramische Werkstoffe, (J. Kriegesmann, Edt.), 51. Erg.-Lfg., (1999) 5.2.2.3. [ 10lP. Muller, E. Macherauch, ‘Das sin$-Verfahren der rontgenographischen Spannungsmessung’, Z. angew. Physik 13, (1961) 305-312. [11]V. Hauk (Edt.) ‘Structural and residual stress analysis by non-destructive methods’, Elsevier, 1997, ISBN 0-444-82476-6.
REFERENCES G. Knoll, V. Lagemann, A. Radtke, ‘Stromungsmechanische Kennwertbildung zur Charakterisierung mikrohydrodynamischer Effekte sowie zum Festkorperkontakt keramischer Werkstoffe’, Werkstoffe fur die Fertigungstechnik (F. Klocke, B.-R. Hohn, Edts.), Wiley-VCH Verlag GmbH, Weinheim, ( 1999) 175-180. M. Geiger, W. Konig, H.K. Tonshoff, ‘Einsatz von Laserstrahlung zur Herstellung tribologisch optimierter Oberflachentopographien’, Werkstoffe fur die Fertigungstechnik (F. Klocke, B.-R. Hohn, Edts.), Wiley-VCH Verlag GmbH, Weinheim, ( 1999) 157-162. M. Zimmermann, M. Rombach, ‘Microhydrodynamic Effects in Sliding Contacts of Real Surfaces: A Numerical Approach’, Ausmb 98 - Tribology at
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DESIGN OF WEAR RESISTANT POLYCRYSTALLINE ALUMINA D. Galusek", F.L. Riley**, R. Brydson** (*)Institute of Inorganic Chemistry, Slovak Academy of Sciences, M. Razusa 10,911 05 Trench, Slovak Republic (**) Department of Materials, School of Process, Environmental, and Materials Engineering, University of Leeds, LS2 9JT Leeds, UK
ABSTRACT Alumina is often used in applications requiring high wear resistance. Ultra fine ultra pure aluminas (with mean grain sizes c 1 pm) exhibit the highest wear resistance, wearing primarily tribochemically. Coarse-grained alumina with the mean grain size > 1 pm wears predominantly by grain boundary microfracture, with rates an order of magnitude faster than submicrometre ones. Carefully designed microstructure and tailored grain boundaries, however, allow significant improvement of the wear resistance of coarse-grained aluminas. The paper attempts to identify the parameters responsible for improved resistance of liquid phase sintered alumina with grain size > 1 pm against the wet erosive wear.
INTRODUCTION Polycrystalline alumina ceramics are widely used for industrial applications where high resistance to wear is required. Ultra fine ultra pure aluminas (with mean grain sizes < 1 pm) exhibit the highest wear resistance. They wear very slowly, primarily by tribochemical processes. However, fabrication of aluminas with mean grain sizes below one micron requires expensive high purity starting powders, carefully designed processing routes, or special preparation procedures (e.g. sol-gel), which are usually not suitable for industrial production due to high production and raw materials' costs. The grain size of the most standard commercial aluminas is usually well over 1 pm. Such materials are prepared by liquid phase sintering, and normally contain several percent of a combination of calcium and magnesium silicates. The wear rates for wide range of wear modes (including erosive and abrasive wear, cutting, and grinding) of such materials are known to increase significantly with the grain size. - 6 The main wear mechanism is microfracture and crack interlinking, leading to grain detachment and the development of rough surfaces. In the case of pure alumina the fracture mode appears to be predominantly intergranular, suggesting fracture follows preferentially the grain boundaries. The use of silicate based sintering additives leads to significant improvements in the wear resistance of alumina, and this is well recognized industrially. Recent laboratory studies using model high purity aluminas of controlled composition have shown the generally beneficial effects of Group I1 metal silicate additions on the wear resistance of alumina
materials. The magnesium silicate appears to confer the greatest benefit. We report here results of work on magnesium silicate (MgSi03) densified alumina, with the focus on fully dense polycrystalline materials containing up to 10 % by weight of magnesium silicate of 1:l stoichiometry. The effects of additive quantity, mean alumina grain size, and residual stresses have been examined in detail.
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EXPERIMENTAL Alumina powders (CL3000SG, Alcoa International Ltd., Worcester, UK) with or without sintering additives were hot-pressed (HP)in a graphite die at 20 MPa and at various temperatures and times depending on the amount of additives to obtain fully dense discs (6 mm thick and of 25 mm diameter). The hot-pressing conditions are listed in Table 1. Magnesium silicate (MS) was added by mixing the alumina powder with isopropanol solutions of magnesium nitrate Mg(N03)2.4H20 (AnalaR grade, BDH Ltd., Poole, UK) and tetraethylorthosilicate (TEOS, AnalaR grade, BDH Ltd., Poole, UK). The amounts of these materials were calculated to achieve final compositions with 0.5, 1, 5, or 10 % by weight of MgSi03 with MgO : SiOz molar ratio of 1: 1. An aqueous solution of ammonium hydroxide (10 wt %) was then added to precipitate magnesium hydroxide, and to hydrolyse the TEOS and the suspension was dried under an infrared lamp. The dry powder was calcined 30 min at 900 "C.
Table 1 Hot-pressing conditions and mean grain sizes (MGS) of the materials studied.
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A1 A2 M0.5 M1 M5 M10
MS wt. % 0 0 0.5 1 5 10
HP conditions T,OC t,min 1620 45 1620 70 1500 70 1500 60 1450 30 1450 20
MGS pm 2.5 3.4 2.1 2.0 1.9 1.8
To induce grain growth, selected hot-pressed materials were subsequently annealed without pressure in the hot-pressing die at temperatures 50 "C higher than the temperature of hot-pressing for 3 or 6 h, or for 6 h at 1620 "C. Surface layers of the hot-pressed discs were removed by silicon carbide paper (240 mesh powder) and bulk densities (p) were determined by mercury
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immersion. Fragments of hot-pressed discs were milled in a tungsten carbide puck mill and used for a XRD measurement using a powder diffi-actometer (Philips APD1700) operating at 40 kV with C u b radiation within the range of angles 5" to 70". To avoid possible effects of porosity on wear rate, only materials with bulk density > 98 % of theoretical were used for further studies. Discs were weighed to k 0.1 mg and worn under carefully standardised conditions in an abrasive slurry consisting of 750 g of 0.5 to 1 mm fused alumina grit (Morganite 954 grog, Morganite Ceramics Ltd, Neston, UK) in 250 cm3 distilled water. The disc was clamped horizontally in a specially designed cylindrical polyurethane block so as to expose approximately one half of the disc surface. The block was attached to the shaft of an attritor mill, and rotated at 500 r.p.m., giving an effective alumina disc perimeter velocity of 1.9 ms-'. The final disc weight was measured to rf: 0.1 mg, and the weight losses occurring between the standard times of 2 h and 6 h were used to calculate a wear rate. Full details of the equipment and experimental procedures are given elsewhere. 8 After wear rate measurements, discs were cut by a diamond wheel and polished by diamond paste to 1 pm finish. Polished specimens were etched thermally (pure alumina at 1500 "C for 30 min), or chemically (MgSiO, sintered alumina, boiling phosphoric acid for 30 s), to reveal grain boundaries and examined by SEM. Micrographs were analysed by standard image analysis software (KS 400, Kontron Electronics GmbH, Germany) to give mean grain sizes and grain size distributions. For TEM examination specimens were core drilled and sliced to prepare 3 mm diameter discs of 500 pm thickness. The discs were further ground, polished, dimpled and ion-milled (Gatan PIPS low angle ion beam thinner) until transparent to the electron beam, and coated with carbon. Specimens were examined using a transmission electron microscope (Philips CM20 200 keV) fitted with a scanning (STEM) unit, ultra thin window EDX detector (Oxford Instruments) and PEELS (Gatan). Cr3+ luminescence spectra were determined on polished surfaces using a Raman optical microprobe (Renishaw Model 2000) and incident radiation from a HeNe laser at 633 nm (15797.79 cm-') in order to assess thermal expansion-anisotropy stresses in polycrystalline aluminas. The incident light was focused using a 5x objective giving a spot size much greater than the alumina grain size and a polished sapphire single crystal was used as the unstrained reference. Spectra were taken from at least 6 different positions on each specimen, including the sapphire single crystal, which was measured before and after each alumina specimen. The details of data analysis are given elsewhere.
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RESULTS Hot pressing followed by high temperature annealing allowed to prepare pure and magnesium silicate-sintered aluminas with a range of grain sizes. For all materials the wear rates increased with mean alumina grain size, and confirmed the strong influence of alumina grain size on wear rate. The measured values for the pure alumina agreed satisfactorily with those of the earlier study. The incorporation of small amounts of MgSi03 markedly reduced the wear rate. Figure 1 shows wear rate as a function of mean alumina grain size for different levels of MgSi03 addition.
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De1 Ctm Fig. 1 The wear rate of pure and magnesium silicate sintered aluminas vs. the mean grain size. The SEM examination revealed strikingly different morphologies of the wear surfaces of MS-doped specimens, compared to the pure aluminas. The facetting was an apparent feature of the wear surface of the pure alumina, and indicated clean grain detachment resulting from intergranular fracture, grain boundary cleavage, and crack interlinking. On the other hand, the roughness of the wear surfaces of magnesium silicate sintered materials decreases progressively with increasing magnesium silicate content. The smoother areas in the magnesium silicate sintered materials appear to be the combined result of smooth transgranular cleavage coupled with subsequent tribochemical polishing, in the absence of significant intergranular fracture leading to the detachment of entire grains. (Fig 2) TEM examination of materials containing MgSi03 showed a 2 nm thick continuous amorphous film at all alumina-alumina boundaries examined. Excess glass seggregated in triple grain junctions (triple pockets). A typical feature is shown in Figure 3. The composition of the silicate varied between the twograin boundary zone and the larger triple point regions, as had earlier been found for CaSi03 doped alumina materials. lo The two-grain boundaries (grain faces) had roughly equal proportions (expressed as atom %) of Mg and Si, whereas the triple point regions were
-
Fig. 2 SEM micrographs of wear surfaces of polycrystalline aluminas with various contents of magnesium silicate: O%-a,l%-b,5%-c,lO%-d rich in Si (Si : Mg ratio approximately 3 : 1). Alumina grains in materials containing MgSi03 showed evidence of considerable stress, in the form of dislocation networks. Supplementary XRD examinations showed small amounts of crystalline secondary phases formed by the reactions of MgSi03 and the alumina matrix. The main secondary phase in materials hot-pressed for shorter times were spinel (MgA1204) and sapphirine (Mg4AlloSizOZ3). High temperature annealing promoted the formation of mullite (A16Si2013),accompanied by the disappearance of sapphirine. These secondary crystalline phases detected by XRD, together with a low thermal expansion silicate film on the alumina grain boundaries, would be expected to lead to residual stresses on cooling from formation temperature, because of mismatches of thermal expansion coefficient. ” Analysis of the Cr3+ photoluminescence spectra showed obvious differences between the specimens. The measurements confirmed the presence of high local fluctuating compressive and tensile stresses on the level of approximately 450 MPa in the MgSi03 sintered materials, which were significantly higher than those in the pure alumina (- 200 MPa).
DISCUSSION This work reveals several features of the influence of the magnesium silicate sintering aid on the wear behaviour of polycrystalline alumina. An addition of -1 % MgSi03 reduced the wear rate more than 50 %. At the same time the predominant fracture path switches from inter- to transgranular. Associated with these aspects is a doubling in the maximum in the fluctuating residual stresses in the alumina microstructure. These features are examined individually. It appears that the erosion resistance of the material is primarily determined by the strength of grain boundary. The magnesium alumino-silicate grain boundary phase clearly has a strong influence on the microfracture behaviour of alumina. The greater difficulty in forming microcracks in the silicate containing material, together with suppression of intergranular fracture, suggests that the silicate has a marked grain boundary strengthening action. The calcium silicate densified alumina examined earlier lo showed amorphous silicate films at all the two-grain boundaries, of the order of 2 nm thickness. The TEM examinations of the magnesium silicate densified material used here also showed continuous films of amorphous silicate at two-grain boundaries. The excess over the amount needed to form the equilibrium thickness l3 grain face film will form pockets at threegrain edges (“triple points”) and grain corners, interconnected by the thin grain face films. Although
165
no calculations have been reported for MgSi03 sintered alumina, theoretical modelling of CaSi03-containing alumina material suggests that a silicate containing grain boundary of certain compositions should be 14 stronger than a pure alumina-alumina boundary. This is essentially because of the ability of the cations and their associated oxygen atoms to assist the relaxation of grain interface stresses caused by misalignment of Si-0 and Al-0 bonds. The analysis of the Cr3' piezoluminiscence spectra suggests that the high fluctuating local stresses in the magnesium silicate densified aluminas are not caused by thermal contraction anisotropy of the alumina grains but by thermal expansion mismatch between the alumina and localised pockets of intergranular silicate. A silica rich glass will have a much lower thermal expansion coefficient than the alumina. 15 It has earlier been assumed that hoop stresses across grain boundaries generated by thermal expansion mismatch are important factors influencing crack pro agation during the wear of two-phase materials. Large glass pockets, or particles of an appropriate intergranular phase, located at boundaries can be expected to generate such stresses. In the material, assuming a relatively rigid glass matrix, the alumina grains would be expected to develop internal tensile radial and hoop stresses during cooling from sintering temperatures (- 1600 "C). In practice, the glass matrix is likely to have a much lower modulus than the alumina and some stress relaxation will occur, particularly at temperatures above T, by viscous flow. However, as the Cr3+photoluminescence spectra show, the fluctuating residual stresses are still considerable, and much larger than those in pure polycrystalline alumina arising from thermal expansion anisotropy. The maximum stresses result from the isolated pockets of silicate at grain edges and corners. Moreover, because the silicate is the continuous phase, hoop compressive stresses (and radial tensile stresses) along the alumina grain boundaries will be developed. These provide the basis for a second grain boundary strengthening action of the magnesium silicate. Despite the fact that the addition of magnesium silicate alters the fracture mode from inter- to intragranular, the influence of alumina grain size on the wear rate is preserved. The explanation is based on an older model of Davidge and Riley developed for high purity aluminas where fracture occurs primarily intergranularly. For transgranular fracture it is assumed that the mean overall crack velocity is given by the mean grain size and by the time required for the crack to cross an average grain and the two-grain boundary. If we assume that the propagation of the crack through the grain is very fast, then the overall velocity of crack propagation is primarily controlled by the time the crack needs for crossing the boundary. In other words, the crack propagates slower if it encounters the grain boundary more often, i.e. if the alumina grains are finer.
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Fig. 3 TEM micrograph of 5 wt.% MgSi03-sintered alumina. 1 - triple pocket, 2 - array of dislocations with a strain fringe towards the triple pocket region, 3 - grain boundary with continuous layer of magnesium silicate. The last question is the influence of the amount of magnesium silicate on wear. It is obvious, from Fig. 1, that increasing amount of silicate in the material increases the wear resistance of polycrystalline alumina. If one assumes the equilibrium thickness of the grain boundary film, l 3 then the excess glass seggregates in triple pockets. The radial and hoop stresses within a pocket arising from thermal expansion mismatch between the alumina and a glass pocket, are independent of the size of the pocket. However, for larger pockets, the hoop stress will be maintained at larger distances in the two-grain boundaries, and the mean hoop stress over the whole boundary should be larger. Crack progression occurs only when the material in the vicinity of the crack tip acquires sufficient strain energy to overcome local constraints, supplied by a particularly severe local impact of the eroding particle. The assumption that the energy required to drive a crack across a boundary should be related to the magnesium silicate (and thus intergranular phase) content thus has justification in principle. The large tensile stresses in the alumina grains of the magnesium silicate densified material would also provide a basis for an enhanced rate of tribochemical wear. l9 Although no clear evidence for this was obtained in the present work, the increased smoothness of wear surface of samples with increasing magnesium silicate content suggests that this mechanism actually takes place. It seems, therefore, that the magnesium silicate intergranular phase influences the wear rate by two separate mechanisms, both involving the thermal expansion mismatch stresses. By applying hoop compressive stresses to the boundaries, and tensile
stresses to the alumina grains, it inhibits intergranular crack propagation and favours transgranular propagation. By raising the chemical potential of the alumina grains it can also potentially enhance the underlying rate of tribochemical wear.
CONCLUSIONS The mass loss caused by the wear of polycrystalline alumina with the mean grain size larger than approximately 1 pm occurs primarily by microcrack initiation and propagation, and crack interlinking followed by grain pull out. The magnesium silicate intergranular phase influences the wear rate by a mechanism involving the thermal expansion mismatch stresses. By applying hoop compressive stresses to the boundaries, and tensile stresses to the alumina grains, it alters the fracture mode from intergranular to transgranular, thus increasing the strain energy required for the crack propagation. Increasing magnesium silicate content improves the wear resistance via the increased average hoop compressive stressed that strengthen the grain boundaries. Despite the fact that the microcracks propagate primarily transgranularly, the grain size dependence of wear is retained. As a result, the magnesium silicate sintered alumina with the wear resistance two to three times higher than that of the pure alumina with a comparable mean grain size has been prepared. The work provides certain implications for design of wear resistant ceramics. Together with fine grain size, the secondary phases with suitable thermal expansion that results in formation of compressive hoop stresses across the two grain boundaries in the course of cooling from the processing temperature appear to improve largely the wear resistance. Increase of tribochemical component of wear on expense of microcracking is beneficial. The Group I1 element silicate grain boundary phases, continuously distributing at all two-grain boundaries and forming triple pockets with low-thermal expansion glass appear to meet the above requirements.
ACKNOWLEDGEMENT The work was supported by the NATORoyal Society Fellowship scheme, by EPSRC Grant GWJ86513, by the Slovak National Grant Agency under the grant NOVEGA 2/7055/20 and by the NATO Science for Peace Programme under the contract No SfP 974122. Alcoa International Ltd. donated the alumina powders.
REFERENCES 'R. Morrell, Handbook of Properties of Technical & Engineering Ceramics: Part 2 Data Reviews. Section I: High-alumina Ceramics, R.Morrel1, HMSO (London 1987). *R. W. Rice, B. K. Speronello, Effect of microstructure on rate of machining of ceramics, J.Am.Ceram.Soc., 59,330-333 (1976).
3S. M. Wiederhorn, B. J. Hockey, Effect of material parameters on the erosion resistance of brittle materials, J.Muter.Sci., 18, 766-780 (1983). 4D. B. Marshall, B. R. Lawn, R. F. Cook, Microstructural effects on grinding of alumina and glass-ceramics, J.Am.Ceram.Soc., 70, C139-C140 (1987). 5M. G. Gee, E. A. Almond, Effect of surface finish on the sliding wear of alumina, J.Mater.Sci., 25, 296310 (1990). %. Miranda-Martinez, R. W. Davidge, F. L. Riley, Grain size effects on the wet erosive wear of highpwity polycrystalline alumina, Wear, 172, 41-48 (1994). 7P. C. Twigg, R. W. Davidge, S. G. Roberts, F. L. Riley, Factors affecting the wet erosive wear of liquid phase sintered aluminas, in Euro ceramics V, Eds. J. Baxter et. al., Key Engineering Materials Volumes 132-136 Part 3, 1524-1527, Trans Tech Publications Ltd., Switzerland (1997). D. Galusek, P.C. Twigg, F.L. Riley, Wet erosion of liquid phase sintered alumina, Wear, 233-235, 588-95 (1999) D. Galusek, R. Brydson, P.C. Twigg, F.L. Riley, A. Atkinson, Y.-H. Zhang, Wet erosive wear of alumina densified with magnesium silicate additions, J.Am.Ceram.Soc.. submitted for Dublication lo R. Brydson, P. C: Twigg, S. C. khen, F. L. Riley, X. Pan, M. Riihle, Microstructure and chemistry of intergranular glassy films in liquid phase sintered alumina, J.Am.Ceram.Soc.,81,369-379 (1998). 11 C.A. Powell-Dogan, A.H. Heuer, Microstructure of a 96 % alumina ceramic: 11, crystallization of highmagnesia boundary glasses, J.Am.Ceram.Soc., 73, 3677-83 (1990). 12 C.A. Powell-Dogan, A.H. Heuer, M.J. Ready, K. Merriam Residual stress-induced-grain pullout in a 96 % alumina ceramic, J.Am.Ceram.Soc., 74, 64649 (1991). 13 D.R. Clarke, On the equilibrium thickness of intergranular glass phases in ceramic materials, J.Am.Ceram.Soc., 70, 15-22 (1987). I4S. Blonski, S.H. Garofalini, Atomistic structure of calcium silicate intergranular films in alumina studied by molecular dynamics simulation, J.Am.Cerum.Soc., 80, 1997-2004 (1997). "S. English, W.E.S. Turner, Relationship between chemical composition and thermal expansion of glasses, J.Am.Cerum.Soc., 10,551-570 (1929). 1 6p.C. Twigg, A. Castro, R.W. Davidge, F.L. Riley, S.G. Roberts, Indentation induced crack interaction in alumina ceramics, Phil. Mag. A , 74, 1245-1252 (1996). "R.W. Davidge, F.L. Riley, The grain size dependence of the wear of alumina, Wear, 186-187,45-49 (1995). "R.W. Davidge, T.J. Green, The strength of two-phase ceramic materials, J.Mater.Sci., 3, 629-634 (1968). '%. C. Twigg, R. W. Davidge and F. L. Riley, "The effects of silicon carbide nano-phase on the wet erosive wear of polycrystalline alumina", J.Eur.Ceram.Soc.,16,799-802 (1996).
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CRACK GROWTH OF CERAMIC MATERIALS IN SLIDING CONTACT Hatsuiko Usami*, Junji Sugishita and Hiroshi Murase Meijo University, 1-501 Shiogamaguchi, Tempaku, Nagoya, 468-8502, Japan Key Words: Brittle materials, Friction and wear, Crack growth
ABSTRACT In order to evaluate effects of crack growth on friction and wear of brittle materials such as ceramics, static contact for the crack initiation and flexural strength measurements were connected to the friction and wear experiment. Materials used in the present study were soda lime glass and alumina. A testing apparatus having a reciprocal movement was used for the friction experiment at a constant frequency with various normal loads in air. The results showed that the wear amount of the glass depended on the crack growth near the wear scar. Critical condition for the crack initiation during the sliding contact was estimated from the static contact test and was confirmed that the tangential traction resulted in the increase in the maximum stress at the trailing edge. The results of the strength measurement test revealed that the crack growth rate depended on the testing condition.
1. INTRODUCTION Ceramics are the candidate materials for anti wear systems because of their high hardness and chemical stability (1,2). Friction and wear behaviour of the ceramics has been studied past decade years (3). As results, basic wear mechanism of the ceramics are similar to those of metals and it is well recognized that the wear amount is considerably smaller. However, it has been confirmed that the wear amount becomes greater at the specific condition. For example, toughness of the material has often influenced on the wear amount when the abrasive wear is dominant on the wear mechanism (4). Although most of the results due to the sliding wear of ceramics are carried out at a low normal load to avoid the crack growth, it is important to evaluate the fiiction and wear in the condition where the applied load is enough for the crack initiation. Damage development due to friction test is usually characterized by a wear scar of the surface and the wear properties of the material is evaluated by the wear amount of the surface. As shown in the crack growth on the wear track of ball bearings, it is also well known that subsurface layer below the wear track has damaged (5). Therefore, it is important for evaluating the wear properties to investigate not only measuring the wear amount but also characterize
the damage development of the subsurface layer. The crack initiation is one of the most important factor for brittle materials such as the ceramics because of unstable failure. The present study is focused on the effect of the crack initiation on the wear properties of ceramics during sliding friction. Soda lime glass bars were used as model material. Static contact test to determine the critical condition for the crack initiation and flexural strength measurement to evaluate the damage development of subsurface layer near the wear scar were carried out. The critical condition for the crack initiation in sliding contact and the effect of the crack growth on the wear behaviour are discussed.
2. EXPERIMENTAL 2.1. Materials Soda lime glass bars and alumina spheres were mainly used for specimens. Table 1 shows mechanical properties of the specimens. The glass bar had rectangular shape (3x4~40mm) in conformity with JIS R 1601 (6). The radius of the alumina sphere was 9.5 mm.
Table 1: Mechanical properties of Specimens Young’s Glass Surfaces of the specimen were finished by polishing. The comer of the glass bar was bevelled to eliminate the stress concentration in strength measurement. Crack initiation under static contact and flexural strength measurements were carried out after the friction experiment using the same specimen.
2.2. Testing apparatus A testing apparatus having a reciprocal
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movement was used for the friction experiment. Figure 1 shows a schematic of the testing apparatus. The alumina sphere was fixed into the holder, then was attached to the octagonal shape dynamo-meter consisted from two curved beams and 8 strain gages. The glass bar was fixed to the guide way driven by a DC motor with a connecting link. The normal load was applied by mounting dead weights on the upper end of the dynamo-meter. The experiments were carried out in air. The glass bar was made in contact with the alumina sphere first, and the normal load was applied then the glass bar was driven. The testing condition was given in Table 2. The test was interrupted at the specific number of cycles to measure the weight loss of the specimens and to observe the surface.profile A commercial grade oil was supplied to reduce the tangential traction at the interface. The oil was supplied just before the test and did not add during the test. Since the oil was consumed with the friction, it was suggested that the contact condition at the interface had transformed with the increase in the number of cycles.
the load was removed. After the experiment, the testing surface of the glass bars was observed by an optical microscopy to measure the radius of the ring crack growth. 2.4. Inert strength measurement Measurement of fracture strength of the bar specimen was carried out after the friction experiment in conformity with 4 point bending test to evaluate crack growth below the wear scar. Schematic of testing configuration was shown in Fig. 2. The wear scar located to the lower surface and was subjected to the tensile stress. The cross head speed was 0.01 mm/min. CHS=O.Ol mm/m in
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Octagonal-shape dvamo-meter
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Normal load
//-
I Wear scar
Fig. 2 Schematic of flexural strength measurement
3. RESULTS 3.1. Coefficient of friction Friction force was measured continuously during the test and was evaluated as the coefficient of friction. The coefficient of friction as a function of number of cycles was shown in Fig. 3. In dry conditions, The coefficient of friction was greater at initial stage and indicated the maximum value at approximately 50 cycles, then decreased and reached to the steady state after 200 cycles. The values at maximum and at the steady state were 1 .O and 0.6 to 0.7, respectively.
Fig. 1 Schematic of testing apparatus
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Maximum speed
Number of cycles
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2.3. Contact failure Static contact tests were applied to evaluate the critical stress for the crack initiation of the glass surface using an electric-mechanical testing apparatus (AG-IOB, Shimazu. Co. Ltd.). The glass bar was mounted on the silicon nitride plate. The alumina sphere was put on the glass bar. The alumina sphere was pushed by the piston with a constant cross head speed of 0.01 mm/min. An acoustic emission (AE) sensor was set on the silicon nitride plate near the glass bar. The AE signal was measured continuously during the test. The critical load was determined from the AE signal, then
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In lubricated condition, the coefficient of friction was approximately 0.1 and was almost constant during test. The influence of the applied normal load on the coefficient of friction was small in dry and lubricated conditions.
3.2. Wear behavior The wear amount of the specimen was measured from the weight loss of the specimen and was indicated as the value per unit normal load. It was difficult to evaluate the wear amount of the alumina spheres and the glass bars tested in lubricated condition since the weight change was too small to measure.
In the lubricated condition (Fig. 6), cracks were not observed the initial stage at 40 N (Fig. 6a), but initiated after 1000 cycles. At 100 N, Ring cracks was seen after 1000 cycles (fig. 6b). The angle of the crack was different that of obtained in the dry condition. Therefore, it is estimated that not only applied normal loads but also tangential traction corresponding to the coefficient of friction have influenced on the critical value for the crack initiation.
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CI
* r
I
E"a L
25N
u I
3
+40N 0
200
40N(L) +l00N
400 600 800 Number of Cycles, n
(a), IOON, n=3 1000
Fig. 4 Wear amount as a function of number of Cycles (Filled points indicate the results in lubricated condition. ) Figure 4 shows the wear amount of the glass bar as a function of the number of cycles. Slopes in the Figure correspond to the specific wear rate. The specific wear rate was greater at initial stage and reduced with the increase in the number of cycles. After 200 cycles, it seemed that the specific wear rate reached to the steady state region. The average values at the initial stage less than 50 cycles and at the steady state region were approximately 1 O-* mm2/N and 1 O-'ommZ/N, respectively. The transition of the wear rate from the initial stage to the steady state region occurred at approximately 50 cycles at 15 N and 200 cycles at 100 N. There was small difference in the specific wear rate against the variations of the normal load. As results, it is estimated that the wear loss per unit normal load after 1000 cycles depends on the number of the cycles of the transition occurrence.
3.3. Wear scar and wear debris Optical micrographs of the surface after the test in dry condition were shown in Fig. 5. Many cracks grew near the wear scar were clearly seen. The cracks grew on the slant direction against the sliding motion in the initial stage (Fig. 4a). Since the increase of the crack growth region was smaller than that of the area of the wear scar, the ratio of the crack growth region and wear scar decreased with the increase in the number of cycles (Fig. 4b). The distance of adjacent cracks was about 100 pm. Such kinds of crack initiation was observed all cases in dry condition.
(b), IOON, n=
1000
Fig.5 Optical micrographs of the surface tested in dry condition Wear debris was collected from the surface near the wear scar of glass bars tested in dry condition (Fig. 7). It was difficult to collect the debris in lubricated condition because the oil was spread and consumed during the test. The shape of the wear debris was different depend on the number of cycles. At initial stage, the shape of the debris was flake whose size was approximately 100 pm. Comparing the interval of the cracks on the surface shown in Fig. 5, it seemed that the large wear debris resulted in the interaction of the crack growth. At steady state region, the wear debris was much smaller. Considering the wear behavior, the wear amount depends on the size of the wear debris.
171
(a), 40 N, n=20
(a), IOON, n=20
(b), IOON,n=1000
(b), 100 N, n=1000 Fig. 6 Optical micrographs of the surface tested in lubricated condition
Fig. 7 SEM images of wear debris
3.4.Crack initiation under sliding contact No tiastic behavior was observed on the loaddisplacement curve. The ring crack was clearly seen on the surface as shown in Fig. 8. The maximum value of the applied load was approximately 400 N, which was considerably greater than the applied load for the friction experiments. It was confirmed that the radius of the ring crack was about 15 % larger than that of the contact area predicting from the Hertzian contact theory (7). I t was also observed that several ring cracks formed outer region to the primary ring crack when the applied load exceeded the critical value predicted from the AE signal. Assuming that the contact condition between the alumina sphere and the glass bars was elastically, the critical stress (T, for the crack initiation was given as following equation (8),
Fig. 8 Optical images of ring crack
(3,
=
yp 2nx
Where, P, v and x are maximum load, Poison’s ratio of the glass (=0.22) and the radius of the ring crack, respectively. Critical stresses of the ring crack initiation was approximately 400 MPa and was greater than those obtained by conventional 4-point flexural test (=70 MPa).
172
3.5. Inert strength N o plastic behavior was observed on the loaddisplacement curve. The residual strength as a function of the number of friction cycle was shown in Fig. 9. The nominal strength before the friction experiment (n=O) was 70 MPa, In dry condition, the strength degradation was larger at lower friction cycles and at greater normal loads. In the lubricated condition, The degradation was relatively small
compared with the results tested in dry condition. Particularly, at 40 N The strength degradation after 20 friction cycles where the crack was not observed on the surface was negligible.
E a"
6 ' O0
t
that the area of the crack growth increased with the increase in the normal load and the number of cycles. Considering the strength degradation (Fig. 9), the crack mainly grows at initial stage then becomes complicated path.
0
strength
A
I
10
100
I000
Number ofCycles, n Fig. 9 Residual strength after the friction test
Coefficient of Friction, p Fig. 10 Maximum stress vs. coefficient of friction
4. DISCUSSION 4.1.Effect of shear stress on the crack initiation Relationship between the maximum tensile stress (T and coefficient of friction p was shown in Fig. 10. The maximum stress (T occurred at the trailing edge and defined as following equations (8).
o =o,+o,
Where, a is the radius of the contact area and Po is the maximum normal stress. (T n is the same as given in sect. 3.4. The value of the maximum stress is given at x=a, because the contribution of (T t on (T is greater than that of (T n and decreases with the increase in x at x>a. The value of the coefficients of friction p=l and p=O.l correspond to the dry and the lubricated condition at the initial stage of the friction test. Under the static contact (i.e. p=O), the stress for the crack initiation was 400 MPa. The maximum stress increases with the increase in the coefficient of friction, and exceeds the critical stress of the crack initiation ( 400 MPa )even in dry condition at a load of 15 N. The results of the friction experiments show good agreement with Fig. 10. Therefore, this suggests that the tangential traction results in the decrease in the normal loads for the crack initiation.
(a), 40 N, n=20, in lubricated condition
(b), 100 N, n=1000, in dry condition
4.2.Crack growth below the wear scar Figure 1 1 shows fracture surfaces. crack formed by the sliding contact below the wear scar was seen. In the lubricated condition at lower normal load, Crack growth was small (Fig. 1 la). At higher normal load, The crack path became complicated. It was confirmed
Fig. 11 Optical micrographs of fracture surfaces
5. CONCLUSION Effect of crack growth on the friction and wear of ceramics material was evaluated using the soda lime
173
glass bars as the model material. Static contact for the determination of the critical stress for the crack initiation and flexural test for the estimation of the crack growth below the wear scar were also were carried out. Followings are the summarize obtained results. 1) The addition of tangential traction at the interface results in the decrease in the normal load for crack initiation. 2) The coefficient of friction and the wear amount depend on the crack growth and are greater at initial stage at the test when the stress at the trailing edge exceeded the critical value of the crack initiation. 3) The crack growth rate is greater at initial stage in the friction test. REFERENCES I ) G. L. Boyd and D. M. Kreiner, “AGT10 1/Advanced Turbine Technology Project”, Proc. of the Twenty-Sixth Automotive Technology Development Contractors,2(1986) 895-899
174
2) A. Bennet, “Requirements for Engineering Ceramics in Gas Turbine Engines” Mater. Sci. Technology, Warrendale, PA (1 989) 275-280 3) S. M. Hsu, Y. S. Wang and R. G. Munro “Quantative wear maps as a visualization of wear mechanism transitions in cermica materials”, Wear134 (1989)l-11 4) K. Hokkirigawa and M. Mizumoto “Microscopic Fractural Aspects of Microtribology” Japanese Journal ofTribology 37,11(1992)895-901 5) A. Yoshida et al, “Fundermental study on applying fine ceramics to rolling bearings under water lubrication” International Tribology Conference Yokohama (1995) Vol. 1485-490 6) JIS R 1601 (1993) 7) B. R. lawn, “Indentation of Ceramics with Spheres: A Century after Hertz”, J. Am. Ceram. SOC.,81,8(1998) 1977-1994 8) R. D. Mindlin, “Compliance of Elastic Bodies in Contact”. J. Applied Mechanics Sept. (1944) 259-268
DAMAGE DETECTION IN TETRAGONAL ZIRCONIA POLYCRYSTALS (TZP) BY IMPEDANCE SPECTROSCOPY S. Wagner*, A. Tiefenbach**, R. Oberacker*, M. J. Hoffmann", B. Hoffmann** (*)Institut f i r Keramik im Maschinenbau, Universitat Karlsruhe, D-76131 Karlsruhe, Germany (**)Institut f i r Werkstoffe der Elektrotechnik, Universitat Karlsruhe, D-76131 Karlsruhe, Germany
ABSTRACT The influence of phase composition and artificial long cracks on the electrical behaviour of 9 mol% CeOz stabilised Zr02 (9Ce-TZP) has been investigated by impedance spectroscopy at room temperature as well as at elevated temperatures. Cracks were introduced and propagated by the bridge method. The crack propagation in 9Ce-TZP is combined with a tetragonal to monoclinic phase transformation in the process zone of the crack. These two processes correspond to a change in the electrical properties. It was possible to correlate the changes in the electrical properties to the phase transformation as well as to the crack propagation. Numerical field simulations were performed for this purpose. From experimental results and numerical field simulations a sensitivity analysis has been derived with respect to the detection of short cracks under load at room temperature. Closed cracks are detectable only at elevated temperatures.
possible on the basis of the present knowledge. However, ionic conductors such as zirconia seem to be suitable materials for this technique. In this study both, the influence of the tetragonal to monoclinic phase transformation as well the effect of single cracks on the electrical behaviour of 9 mol% CeOz stabilised ZrOz (9Ce-TZP) has been investigated from room temperature up to 55OOC. This 9Ce-TZP material develops large transformation zones at the crack tip. The electrical characterisation of the material result in restriction only for the room temperature characterisation, where only capacitive measurements are feasible. At elevated temperatures, the ionic conductivity becomes measurable and gives additional information with respect to defects.
EXPERIMENTAL PROCEDURE Sample preparation 3r the sample preparation 9 mol% CeOZstabilised (9Ce-TZP, Unitec, UK) has been used. Plates of 65x45~12mm3 were produced by pressing, cold isostatic pressing and sintering at 1400°C for 2 hours in air. From the plates four point bending specimens with a dimension of 4 8 x 9 ~ 4mm3 as well as specimens for the transformation and high temperature experiments of 1Ox8x 12 mm3 were manufactured by diamond cutting and grinding. The grinding induced monoclinic zone at the surface was removed by annealing the samples for one hour at 1000°C in air. By cooling the specimens to -85°C in cryogenic silicon oil, the transformation of the tetragonal to the monoclinic (m+t) phase was induced. A retransformation process from the monoclinic to the tetragonal state has been realised by heating to 600°C. -
INTRODUCTION In tetragonal zirconia ceramics (Tetragonal Zirconia Polycrystals TZP or Partially Stabilised Zirconia PSZ) subcritical crack growth under static and cyclic load is observed [1,2]. Phase transformation from the tetragonal to monoclinic phase, taking place in the process zone of propagating cracks, acts as toughening mechanism and plays an important role in these ceramics [3,4,5,6]. Lifetime predictions for TZP components under static and cyclic loading are subjected to statistical scatter, due to the unknown and often widely distributed defect population in the material. This limits the use of such components. Therefore the development of non destructive evaluation methods (NDE) is important for the detection of fatigue relevant defects prior to catastrophic failure. Electrical measurements offer some advantages over to the more common NDE methods, such as local measurement of the critical component sections and measurements from room temperature up to high temperatures. There is, however, only limited information available on the application of methods such as impedance spectroscopy for this purpose [7]. N o assessment of the potential of electrical methods is
,!rij2 powder
Microstructural characterisation SEM and TEM were utilised for investigating the microstructure. The mean grain size was determined from SEM micrographs taken from polished and thermally etched specimens by the line intercept method. Samples for TEM investigations were prepared by grinding, dimpling and ion milling. A Zeiss EM912 Omega microscope with an accelerating voltage of 120 kV was used for TEM investigations. The phase analysis was performed at an X-ray
175
difliactometer (SIEMENS model D500) using Cu-Lradiation and a secondary monocromator. The monoclinic fraction was derived from the (-lll),,,, (1 1l),,, and (1 1l)t - peaks according to [8].
Electrical measurements The electrodes for the electrical measurements at room temperature were produced by vacuum deposition of chromidnickel. For the experiments at elevated temperatures platinum has been used. At room temperature the sample capacitance was measured by a multi-frequency LCR Meter (HP 4275A, Hewlett Packard) over a frequency range from 10 kHz to 10 MHz. In the temperature range from 150°C to 550°C an impedance analyser (SI 1260, Solartron) was used over a frequency range from 1 Hz to 1 MHz. The samples were heated by high-intensity infrared line heaters. All experiments were performed in air.
RESULTS AND DISCUSSION Microstructure The material has a mean grain size of 1.1 pm. XRD measurements indicate a single phase tetragonal microstructure in the as sintered state. After the cooling induced tetragonal to monoclinic phase transformation a monoclinic phase content of 7 0 % has been determined. TEM investigations show a good qualitative correlation with the XRD results. TEM micrographs of the microstructure in the as sintered state and after the cooling induced tetragonal to monoclinic transformation are represented in Fig. 2.
In situ impedance spectroscopy experiments In situ impedance spectroscopy experiments were performed at an universal testing machine UTS during four point bending tests. The aim of the in situ experiments was to monitor the propagation of a crack and the process zone, which accompanies this crack, by electrical measurements. The ideal sharp starting cracks in the specimen was induced by the bridge method. Figure 1 shows schematically the experimental four point bending arrangement. The specimens were stepwise cyclic loaded and unloaded with a successive increase of the load in each step. The permittivity of the sample was measured before and after each loading step. Crack growth was measured by an optical microscope. The arrangement was designed for maximum sensitivity, although it does produce an inhomogeneous electrical field and thus numerical modelling is required for quantitative analysis. This modelling was performed using the software code MAFIA [9].
Fig. 2a: TEM micrograph of tetragonal zirconia a) tetragonal grain, b) grain boundary c) grain boundary phase in a triple point
8 Load
Fig. 2b: TEM micrograph of transformed zirconia with a) monoclinic grain b) tetragonal grain c) microcracks at the gain boundary
Electrodes
Fig. 1: Arrangement of the in situ impedance spectroscopy experiments
176
In the initial state the material has a predominantly single phase tetragonal microstructure with small amounts of a grain boundary phase situated in the triple points. In this state no microcracks were detected. After the cooling induced t+m transition, monoclinic as well as tetragonal grains are observed. Furthermore, microcracks at the grain boundaries have been developed due to the volume increase during the t+m phase transition. After the monoclinic to tetragonal
retransformation induced by heating the samples, the material shows a pure tetragonal microstructure with microcracks at the grain boundaries.
Electrical measurements at room temperature At room temperature, only the capacitive part of the impedance can be measured for the different materials. The effective permittivity of the as sintered material amounts about 38. There is a reproducible, drastic decrease in the effective permittivity after the cooling induced tetragonal to monoclinic transformation as shown in table 1. Differences of -30% compared to the specimens in the as sintered state have been measured. After the monoclinic to tetragonal retransformation, the permittivity is increased again to the initial permittivity. The cracks, which are still present in the retransformed 9Ce-TZP obviously cause no significant differences between the measured permittivities in the initial state and after the retransformation. This is easily explained by the small volume fraction of the essentially closed cracks, as the contribution of a second phase to the total permittivity is proportional to its volume fraction. Table 1: Permittivity of 9Ce-TZP before and after phase transformation 9Ce-TZP (microstructure)
Effective Permittivity E
tetragonal reference state without cracks after cooling induced t+m transition (70 % monoclinic phase) tetragonal state after m+t retransformation with a network of cracks
38,l
During unloading, the crack closes and this leads to an increase of permittivity approximately to the initial value. This confirms the results in table 1, where had been shown, that closed cracks are not measurable by impedance spectroscopy at room temperature due to their negligible volume fraction. Beyond 500 N can be recognised, that the permittivity of the unloaded specimen shows a residual loss, which increases with increasing the load and the crack length. This is caused by the tetragonal to monoclinic phase transformation in the process zone, which has a significant influence on the permittivity at room temperature (table 1). Contrary to the reversible contribution of crack opening and closing, there exists an irreversible contribution of phase transition. This irreversible proportion increases during crack growth and can be determined in the unloaded state. Comparing the loaded and unloaded state it can be recognised, that the main part of the changes in permittivity is caused by the contribution of crack opening. 0 -1
26,7
E- -2-3
38,O
a
$ 4
f=100kHz
-5
I
\
s !
Detection of long cracks under load Thus, it can not be expected, that closed cracks would be not detectable by impedance spectroscopy at room temperature. However, cracks open under mechanical load. This volume effect result in a decrease in the effective permittivity of a crack containing specimen under external load. To study this effect quantitatively, in situ impedance spectroscopy measurements were performed according to Fig. 1. Figure 3a demonstrates the changes in measured permittivity A& in dependence of the stepwise applied load. The results under load as well as after unloading are plotted in this graph. In Fig. 3b, the corresponding crack length and crack opening are shown. The crack opening was determined from the optically measured crack length by the relation [lo]: t i = a - %E. f ( a / W )
Considering only the specimen under load there is a monotonous decrease in permittivity with increasing load (Fig. 3a). This is caused by the air volume of the opened crack as well as by the mechanically induced tetragonal to monoclinic phase transformation in the process zone. Up to about 500N, crack opening increases at a stable crack length.
(1)
where a is the crack length, uBthe outer fibre stress, E the elastic modulus and f a geometric function of the crack length a /sample width W relation.
0
I
200
000
400
FI
800
1000
“I
Fig. 3a: Relative permittivity of a bending sample in dependence of the applied load during stepwise loading and unloading
0
200
400 FI
000
800
I000
“I
Fig. 3b: Corresponding measured crack length and crack opening in dependence of the applied load
177
For a quantitative discussion of the in situ experiment results, numerical simulations were performed [ 111 to correlate the changes in permittivity to the crack geometry and to the tetragonal to monoclinic transformation in the process zone. Figure 4 shows a comparison of experimentally measured and numerically calculated decrease in permittivity in dependence of crack opening for different crack lengths. The proportion of the monoclinic zone in the process zone (irreversible portion) had been subtracted from the permittivity and is not considered here. The diagram represents only the reversible portion. The curves, representing the numerical results, were determined for an a/W relation of 0.44 and 0.33 by varying the crack opening. Although the crack length measured in the experiment is smaller than in the simulations, the loss in relative permittivity is underestimated in the calculations. That means, the air volume of the crack is not sufficient to explain completely the permittivity loss. The reason for this difference could be a partial retransformation of the monoclinic phase during unloading.
With E=220 GPa for the investigated TZP ceramic one obtains crack opening-crack length relations shown in Fig. 5 for different stress intensities. For a sensitivity analysis these curves must be compared with the changes of permittivity for a corresponding crack geometry. The results of the electric field simulations can be fitted by the following analytical approximation: AC = const - 6 - 6
(4)
C
The corresponding results are represented as dotted lines in Fig. 5 for measurement sensitivities of 0.1 %, 0.05 % and 0.025 'YO.With the experimentally realised sensitivity of 0.05% and a fracture toughness of 15 MPam0.5, cracks with a minimum length of 30 pm and an opening <2 pm should be detectable. This is valid for a distance of the electrodes of 3 mm. The sensitivity can be fiuther increased by decreasing the distance of the electrodes. K, [MPam0'5)
0 -1
-Sirnulation for e 0 . 3 3
-2
g-3
0
4
8
12
16
Crack length a trCm1
6, I tlJm1
Fig. 4: Experimentally measured and numerically determined permittivity loss in dependence of the crack opening for different crack lengths
Sensitivity analysis with respect to short cracks The crack length used in the experiments is far away from practical relevance. Therefore a sensitivity analysis based on the simulation results has been derived to extrapolate the capability of the method for the detection of shorter cracks under load. The stress intensity factor K,at the crack tip is given by:
where Y is a geometric fhction. Using eq. (1) for the crack opening it follows:
s=x.,., Kl
178
f
(3)
Fig. 5 : Sensitivity of the room temperature permittivity measurements in dependence of the crack geometry for different stress intensities at the crack tip
Electrical measurements temperatures
at
elevated
Using impedance spectroscopy at room temperature, phase transformations and opened cracks can be detected. Damage detection by permittivity measurements at room temperature prior to catastrophic failure seems possible only under ideal conditions, e. g. positioning of the electrodes close to the damage zone of a component. For the detection of closed cracks in TZP ceramics, impedance spectroscopy at elevated temperatures has been studied. In the temperature range up to 55OoC, TZP is an ionic conductor [12]. Impedance spectra were recorded at temperatures up to 550°C on specimens in the initial tetragonal state as well as with a sharp single crack. In Fig. 6 the Bode plots of the impedance spectra of 9Ce-TZP in the initial tetragonal state are represented for temperatures from 20°C to 5OOOC.
11
10
500'
5
T
4 -I
Considering the spectra one can find two ohmic ranges in the damaged state too, similar to the undamaged state. The behaviour of the spectra can be described by 2 R-C elements too. In the high frequency range, there are no differences in the electrical behaviour between the damaged and undamaged tetragonal state. The crack only effects the low frequency behaviour. Here the resulting changes in the electrical properties can be observed as differences in the capacitive and resistive parts of the impedance. The impedance Z increases strongly with increasing the crack length.
Fig.6: Bode characteristics of 9Ce-TZP in the reference state for different temperatures At room temperature only a capacitive behaviour is observed, visible by a slope of (-1) of the log(Z)/log(f) plot. lncreasing the temperature ohmic ranges were detected too. These ohmic ranges can be recognised from the horizontal parts of the impedance spectra (slope 0). Considering temperatures above 300"C, the curves have a qualitatively similar behaviour. At low frequencies there is a first ohmic range, followed by a weakly pronounced capacitive area and a second ohmic range. At high frequencies, a pure capacitive behaviour had been measured. This behaviour corresponds to an electric circuit consisting of two R-C elements. In order to find a quantitative correlation between the microstructure and the electrical properties, modelling is required. Therefore the Brick Layer model, an idealised and often used model with a network of two R-C elements [ 131, has been applied. In this model the microstructure is described as an array of cubic-shaped grains, separated by plane grain boundaries. Such a microstructure of grains and grain boundaries is compatible to the findings of the TEM investigations (Fig. 2). In the model, the electrical behaviour of the grains determines the high-frequency capacitive and ohmic range of the spectra. Grain boundaries control the low frequency range of the impedance spectra. In the investigated 9Ce-TZP they are only detectable at temperatures beyond 300°C. Using this model specific conductances and capacitances for the grains and grain boundary phase were calculated for the undamaged 9Ce-TZP material.
5.8
E 2
5.4
3 m 0
-
5
m
Fig. 7: Influence of the crack length on the Bode characteristics of a 9Ce-TZP specimen In Fig. 8 the change of the capacitance and the change of the conductance in the low frequency range of the impedance spectra are shown as a finction of the a/W relation. Low frequency capacitance CLF and conductance oLF change in a linear manner with increasing crack length. It can be mentioned, that the low frequency capacitance is some more sensitive than the low frequency conductance. The difference is, however, rather small. With an experimental sensitivity of 4 for AcLF/cLF and AcJLF/oLFrespectively, cracks in the pm range should be detectable, even if they are completely closed. 0 n
g
-20
.g5
-40
:
-60
5
After investigating the undamaged tetragonal state, a sharp crack was introduced into the specimen and impedance spectra were recorded. The crack was propagated by the bridge method up to a/W=0.98 for successive measurement of impedance spectra. The corresponding Bode characteristics measured at 400°C are shown in Fig. 7. The curves represents only the influence of the crack. The monoclinic process zone at the crack tip was removed by an annealing process after each stable crack growth.
0
3
8
-80
-100
0
0.2
0.4
0.0
0.8
I
aMI
Fig. 8: Chances in low frequency capacitance CLF and conductivity oLF caused by sharp cracks of different length
179
For the interpretation of the electrical behaviour numerical field simulations has been performed. Therefore the Brick Layer model has been extended with a crack as an additional element of microstructure. This developed model is explained in detail in [14]. Using numerical simulations impedance spectra can be estimated. In Fig. 9 the measured as well as the simulated spectra are shown for the undamaged and the damaged state. Comparing the spectra one can see a very good agreement between the simulated and measured curves. That means, that the crack geometry can be derived from the measured spectra using the numerical simulation model.
Fig. 8: Comparison of the measured Bode characteristics with numerically determined spectra
CONCLUSIONS Using impedance spectroscopy measurements at room temperature tetragonal to monoclinic phase transformation of TZP materials can be detected. This is caused by a drastic difference in permittivity of 30 % between the tetragonal and monoclinic phase. Closed cracks, however, do not influence the permittivity. At room temperature only cracks opened under load are detectable. For the detection of closed cracks, impedance spectroscopy at elevated temperatures seems to have a high potential, as the low frequency capacitance and conductance of TZP is very sensitive to cracks. Impedance spectroscopy is, however, still far away from being a standard NDE testing method. Further work is required to investigate the influence of microstructural changes, which also contribute to the impedance spectra and which have to be separated from the effect of short cracks. A quantitative derivation of crack geometries from changes in impedance spectra requires further progress in modelling of the electrical behaviour.
ACKNOWLEDGEMENT This work was funded Forschungsgemeinschaft DFG
180
by the Deutsche (German National
Science Foundation) under contract No. Ho 69311 1-1 2 and Ob 104/4-2.
REFERENCES L. S. Li, R. F. Pabst: Subcritical crack growth in partially stabilized zirconia (PSZ), J. Mat. Sci. 15 (1986), pp. 2861-2866. T. Fett, D. Munz: Subcritical Crack Growth of Macro-Cracks in Zirconia. J. Mat. Sci., (1991), pp. 1103-1 106. P. Becher, M. V. Swain: Grain-Size-Dependent Transformation Behavior in Polycrystalline Tetragonal Zirconia. J. Am. Ceram. SOC.75 [3], pp. 495-502, (I 992). P. E. Reyes-Morel, J . 4 . Chering, I-W. Chen: Transformation Plasticity of CeOz-Stabilized Tetragonal Zirconia Polycrystals: 11, Pseudoelasticity and Shape Memory Effect. J. Am. Ceram. SOC.71 [8], pp. 648-657, (1988). D. J. Green, R. H. J. Hannik, M. V. Swain: Transformation Toughening of Ceramics. CRC Press, Inc., Boca Raton, Florida, 1989. R. Hannik, P. Kelly, B. C. Muddle: Transformation Toughening in ZirconiaContaining Ceramics. J. Am. Ceram. SOC.77 [5], pp. 1351-1356 (1994). M. Kleitz, C. Pescher and L. Dessemond, Impedance Spectroscopy of Microstructure Defects and Crack Characterisation, Science and Technology of Zirconia 5 (1 993), 593-608. H. Toraya, M. Yoshimura, S. Somiya: Calibration Curve for Quantitative Analysis of the Monoclinic-Tetragonal ZrOz System by X-Ray Diffraction; J. Am. Ceram. SOC.,67 (1984), 183184. T. Weiland, MAFIA Software-Dokumentation, 1992. (10) H. Tada. The stress analysis of cracks handbook Del Research Cop. Hellertown, Pa., 1973. (1 1) A. Tiefenbach: Elektrische Charakterisierung mechanischer Schldigungen in Zr02-Keramik. Dissertation Universitlt Karlsruhe, 1999, erschienen in: Fortschrittberichte, Reihe 5, Nr. 555, VD1-Verlag Diisseldorf 1999. (12) A.Barhmi, E.Schouler, A. Hammou and M.Kleitz, Electrical Properties of Tetragonal Partially Stabilized Zirconia, Adv. In Ceramics 24B (1 985) 885-894. (13) T. van Dijk and A. J. Burggraaf, Grain Boundary Effects on Ionic Conductivity in Ceramic Gd,Zr,x02-(x/z) Solid Solutions, Phys. Stat. Sol. (a) 63 (1981) 229-240. (14) A. Tiefenbach, B. Hoffmann: Influence of a Crack on the Electrical Impedance of Polycrystalline Ceramics. To bee published in the J. Europ. Ceram. Society.
RELIABILITY AND REPRODUCIBILITY OF SILICON NITRIDE VALVES: EXPERIENCES OF A FIELD TEST Gerhard Wetting*, Jiirgen Hennicke*, Helmut Feuer*, Karl-Heinz Thiemann**, Dieter Vollmer**, Erwin Fechter** Frank Sticher***, Andreas Geyer**** (*) CFI Ceramics for Industry GmbH&Co.KG, Rodental, Germany (**) DaimlerChrysler AG, Stuttgart, Germany (***) MAHLE Motorventile GmbH, Bad Homburg, Germany
(****) MAHLE Ventiltrieb GmbH, Stuttgart, Germany
ABSTRACT Within the last three years a field test with silicon nitride (Si3N4)engine valves was performed by DaimlerChrysler with 1628 Mercedes C200-compressor passenger cars. For 1000 cars Si3N4-valveblanks were produced by CFI on a pre-production scale and machined by EuroVal, now MAHLE. In order to realize production volume, the primarily developed gaspressure sintered Si3N4-gradehad to be down-scaled to allow low-pressure sintering while still meeting materials' requirements for engine valves. Machining had to be optimized to avoid damage of components, effecting life-time and reliability. Fitting in cylinder heads was adjusted for automatic assemblage. Effects of the Si3N4-valves on fuel consumption and emissions were analyzed on special test rigs of DaimlerChrysler whereas the majority of cars are in normal use by customers. Details of this development work were outlined and a preliminary summary on experiences will be given.
volume was 1,750 cars and CFI was chosen as supplier of SN-valves for about 1,000 cars. As all preliminarily developmental activities and tests were performed with a gas-pressure sintered SNgrade of CFI, this decision comprised the problem of a limited manufacturing capacity for this SN-grade in reasonable time. Thus, in accordance with DC a new SN-grade was developed, whch could be sintered under a low N2-pressure (10 bar) while still offering sufficient mechanical properties and reliability. With this sintering route, the manuf6cturing capacity was increased by a factor of 3-4 so that the test could be performed within the time schedule. Of course, this new material had to be qualified once more by DC to be a suitable valve material. Results of respectively experiences made along the whole manufacturing-, evaluation- and application-line within this field test are described in the following chapters.
EXPERIENCE AND RESULTS OF THE FIELD TEST INTRODUCTION During the last decade a lot of work was done to develop and realize automotive engine valves made of silicon nitride (SN). This is documented in various publications and also contributions to the last conference of this series [l-61. Daimler-Benz, now DaimerChrysler (DC), were pioneers in this area and got encouraging results with engines equipped with SNvalves on test rigs as well as in passenger cars, amounting in reductions in NO, up to 30 %, CO up to 20 %, fuel up to 6 %, noise up to 12 %, and additionally the observation of reduced coke layer formation on SNvalves in comparison with metallic ones. These positive results motivated DC to perform a field test with the Mercedes C200 passenger car (W202) equipped with the charged 4 cylinder116 valves-Ml l lML engine with SN-inlet and exhaust-valves. Target
GENERAL ASPECTS One main target of this manufacturing- and application-test was to prove the reliability and reproducibility of SN-valves. Details of processing and evaluation are already published [4, 71. In order to identify origins of possible damage, a tracking procedure starting from the raw materials along all processing steps and machining was created by an individual valve number. This number was also registered by mounting the valves into the cylinder heads of the individual engines, and with the engine number the relation to the car number was given. Thus, for each valve in any engine the whole route back to the raw materials and fabrication, sintering and machining lots is documented.
181
SN-MATERIAL GRADE As mentioned, a low-pressure sintered SN-grade (N7202) had to be developed due to manufacturing capacity reasons. Basic properties of this material necessary for structural design calculations are given in Tab. 1 in comparison to the former gas-pressure sintered SN (N3208). As already described in [4],bending bar samples prepared from the valve stems for product control during the manufacturing campaign even resulted in a mean strength of nearly 900 MPa and a Weibull modulus of 18 for 405 samples. This demonstrates that the low-pressure SN-grade N7202 is a highly reliable material also within components. This material had to be qualified once more by DC by rig tests as well as concerning its cyclic fatigue behaviour. Due to stress calculations, a fatigue-strength value of 2 450 MPa after 2-10' cycles was demanded. Fig. 1 shows that the SN-N7202 by means of fineground cylindrical samples fulfills h s demand completely and ends up with a strength of about 550MPa after the mentioned number of load cycles. Due to these results as well as the successful rig tests with valves, this material was qualified by DC to be used for the field test. COMPONENT PRODUCTION The formerly developed manufacturing steps had to be adapted for this new material and optimized. This comprised near-net shaping and minimal warpage during sintering to minimize the cost-effective machining effort. The sintering skin should be as thin as possible in order to remove it securely during machining. Blanks delivered to the machining were evaluated concerning density, dimension- and shapetolerances, cracks and big pores as well as chippings. This final procedure was scaled up to the fabrication of about 10,000 valve blanks per month, summing up to about 50,000 inlet and exhaust valve blanks within this campaign. T h s is probably the highest quantity of SNvalves ever produced. MACHINING Machining and finishmg of the SN-valve blanks was performed by EuroVal, now MAHLE, by the socalled quick-point grinding technique. This technique was a break-through in valve machining as it allowed to finish a SN-valve in about two minutes at reasonable costs [2]. The evaluation of such ground valves by the cyclic fatigue test, however, resulted in unacceptable low strength values after applying the specified load cycles. Reasons for thls were analyzed to be machining defects, hardly to be detected by non-destructive evaluations. Additionally, an unacceptable number of big edge chippings were formed. These results called in question the success of the whole field test. Therefore, enormous effort was spent to eliminate these defects; details are given in a separate paper of this conference [8]. This work comprised the optimization of the whole machining strategy and each individual machining step as well as of the grinding-wheel characteristic. By use of a special acoustic sensor it could be demonstrated
182
that the formation of machming defects and edge chippings is directly related with wear of the grindingwheel (Fig. 2). This example shows for the first machining step, the valve seat, the change of machining conditions with increasing time and number of parts, respectively. With this information, too h g h grinding forces and thus damaging of parts can be reduced significantly as represented in Fig. 3 [7]. Remaining small chippings are rated as non-critical. Accompanied with this optimization was a marked increase of the G-value of the grinding wheels, which indicates the ratio of the material volume removed to the wear volume of the wheel. Excellent Gvalues up to 8,000mm3/mm3 could be reached corresponding to a lifetime of the grinding wheel of 2 10,000 valve blanks. The effect of this grinding improvement on the cyclic fatigue behaviour is already shown in Fig. 1. Fatigue strength is now more than sufficient. Additionally, the machining time could be reduced by this optimization to about 25 % of the primary duration, being of enormous importance for the overall costs. SNvalves machined by this procedure were inspected concerning dimensions and shape-tolerances and sent to CFI for non-destructive evalution. NON-DESTRUCTIVE EVALUATION OF FINISHED SN-VALVES This was done mainly by a special ultrasonic analysis described already in detail in [4,7]. Fig. 4 shows the ultrasonic testing device established. A rather narrow scan net was applied especially in the regions of high stress in application, allowing a detection of surface and near-surface defects bigger than about 150 pm. The whole inspection time was about one minute per valve. T h s inspection procedure proved to be highly reliable as not a single SN-valve delivered to DC was rejected due to further defects found. MOUNTING OF SN-VALVES IN ENGINE CYLINDER HEADS This occurred on an automatic assembly line of DC (Fig. 5). In comparison with the standard procedure of metallic valves, lower stressed springs and softer valve fixtures were applied. As mentioned, the position of each SN-valve was documented in order to be able to track each part back from the car number to the raw material charges. After mounting, often a slightly higher leak rate was observed compared with the metallic system during the first check. This was obviously due to the much stiffer SN-ceramic. After adjustment and abrasion of machining tolerances of the valve seat ring, occurring after relatively short time, even lower leak rates than with the conventional valves were reached. Summing up, assembly did not result in any material-related problems, and could be changed without any alterations from metallic to ceramic valves.
APPLICATION EXPERIENCE In total, about 1,600 C200-ML cars were sold to customers without information on the valve material and additional 100 of such cars were used for DCinternal purposes. During the whole field test of about two years, only one engine failed due to a valve defect. By means of the tracking procedure this could be related to a very early machning campaign. Thus, there is a high probability that this valve failed due to an undetected machining defect. On the other hand, this result indicates that all other valves mounted in engines could be fabricated with h g h reliability and reproducibility. Of course, this is also due to the applied evaluations and can be rated as an important step towards a reliable mass application of highly stressed ceramic components. VALUATION OF THE FIELD TEST This field test with the probably highest quantity of SN-valves ever produced as well as number of cars ever equipped with SN-valves demonstrated that the SN-valve production and evaluation can be performed reliably and reproducibly. Assembling SN-valves into cylinder heads on a common production line proved to be possible with only minor changes. Unfortunately, however, the fuel reduction observed was diminishing low, probably due to the applied charged engine 200ML in comparison to about 4 % with conventional uncharged engines of this type. On the other hand, inspite of the one valve-related damage, DC rates this field test as successful. All processing steps could be improved and the evaluation techniques proved to be suitable to assure quality of the valves. Thus, t h s field test demonstrated the suitability to deliver reliable silicon nitride components in series.
SUMMARY AND OUTLOOK Though there is an overall positive result from the field test of SN-valves, at the moment there is no commitment of the automotive industry to apply this part in series in the near future. The reasons are mainly still higher costs compared with metallic valves. There are, however, various potentials for future applications, based on the facts that in hghly charged engines metallic valves reach their performance limits. This is also the case with engines running near h = 1 due to ecological reasons. Another potential for SN-valves is associated with the introduction of electro-magnetic valve control systems which favour very light weight valves. Finally, attempts to reduce the fuel consumption with the aim of the so-called ,,3-Liter-Car", force engineers to exhaust all weight reduction possibilities of the valve train components. On the other hand, the successful field test demonstrated that there were no higher stresses on the valve in application than calculated. This gives room for reducing material and component requirements and allows a further down-grading of the SN-material with associated lower costs. In this context it is important to
mention once more the successful reduction of the machining effort of SN-valves within this project which now approaches the machining time of metallic valves. It can be summarized that SN-valves prove their performance and reliability within this test. To improve their chances for application, however, further reduction of production costs is necessary to approach the costs of metallic high-performance valves, being the topic of various developmental activities.
We grateful& acknowledge financial support of this project by BMBF.
REFERENCES G. Wotting, H. A. Lindner, E. Gugel: Silicon Nitride Valves for Automotive Engines; in: Adv. Cer. Mats; Transtech Publications Vol. 122-124 (1996) 283-292 E. Gugel, G. Wotting, P. Claeys, P. Woditsch: Silicon Nitride Valves Available for Automotive Engines; Proc. 29* ISATA, Vol. I (1996) 677-684 U. Hoyer, P. Rahanavardi: Untersuchung mit Ventilen aus Leichtbau-Werkstoffen; MTZ (1999) Nr. 9,590-603 G. Wotting, H. A. Lindner, E. Gugel: Experiences with High-Volume Production of Silcion Nitride Automotive Engine Valves; Proc. 6~ Int. Symp. Ceramic Materials and Components for Engines, Arita, Japan (1997) 197-202 K. D. Morgenthaler: Ceramic Valves - A Challenge?; Proc. 6* Int. Symp. Ceramic Materials and Components for Engines, Arita, Japan (1 997) 46-5 1 T. Kinoshita, Y . Migairi, K. Kowasaki, S. Miwa, M. Masudai: Development of Silicon Nitride Valves for Automotive Engines; Proc. 6th Int. Symp. Ceramic Materials and Components for Engines, Arita, Japan (1997) 193-196
G. Warnecke (Ed.): Zuverlassige Hochleistungskeramik Proc. Project-Symposium ,,Prozel3sicherheit und Reproduzierbarkeit in der ProzeBkette keramischer Bauteile", Rengsdorf (2000); ISBN 3-00-005686-6 L. Schafer, K. Eichgriin, T. Magg: Process Design for High-Performance Grinding of Ceramics in Mass Production; Proc. this Conference
183
Tab. 1
Materials' characteristics of the low-pressure sintered SN-N7202 in comparison to the gas-pressure sintered SN-N3208
N7202
N3208
Density
g/cm3
3.23
3.24
RT-Bending Strength RT-Weibull
MPa
850 > 17
920 > 20
800°C-Bending Strength 800°C Weibull
MPa
730 > 17
750 > 20
62
> 100
> 80 > 50
Characteristics
SCG-Parameter n, RT SCG-Parameter n, 800°C Fracture Toughness, RT (SENB) Hardness RT
~~a~rn'" HVlO
7 14.8
7-8 15.0
Young's Modulus, RT Young's Modulus, 800°C Poisson-Ratio
GPa GPa
300 285 0.27
320 310 0.25
Thermal Expansion Coeff., RT-800°C Thermal Conductivity, RT Thermal Conductivity, 800°C
.lo6 1/K W1m.K W/m.K
3.4 22 13
3.2 25 17
A
Fatigue Sample, fineground
0
Valves, prior technology
Number of Cycles for Fracture
Fig. 1 Influence of machining conditions on the fatigue strength of SN-N7202 samples
184
2ndoptimization07/99
.
>*50~m
Fig. 2 Monitoring of grinding-wheel wear by an acoustic (AE) sensor
Fig. 3 Number and sizes of edge-chippings on SN-valves after different machining conditions
Fig. 4 Automatic ultrasonic analysis apparatus for SN-valves
Fig. 5 With SN-valves equipped cylinder head of the C200-engine
185
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SELF-MATED TRIBOLOGICAL PROPERTIES OF PLASMA SPRAYED CHROMIUM CARBIDE COATING Jianfeng Li, Chuanxian Ding Shanghai Institute of Ceramics, Chinese Academy of Sciences 1295 Dingxi Road, Shanghai 200050, China dimension were prepared by a Sulzer-Metco F4-MB
ABSTRACT
plasma spraying gun mounted on an ABB S3 robot
By means of stationary ring on moving ring
using optimized spraying parameters as listed in Table 1 .
arrangement, friction and wear coefficients of self-
The starting powder was commercially Sulzer-Metco
mated plasma-sprayed Cr,C,-NiCr coating were tested with respect to load and sliding speed. It was found that the friction and wear coefficients depended on load and
82VF-NS (93wt%Cr3C,-7wt%NiCr), and the Cr,C,NiCr coating with thickness about 0.5 mm was sprayed onto substrate of lCrlSNi9Ti stainless steel. Table 2
sliding speed. The wear coefficients of moving rings
lists some properties parameters of the coating.
were larger than those of corresponding stationary rings, and the higher the sliding speed, the larger the
Table 1 Optimized plasma spraying parameters Current ( A T 600 Voltage (V) 58 Argon (Lhin) 55 Hydrogen ( L N i n ) 12 Powder feed rate (g/min) 40 Spraying distance (mm) 130 ~
difference of wear coefficient between moving ring and stationary ring was. The results obtained were related to thermal shock and oxidation of the coating caused by frictional heat.
INTRODUCTION Thermal spraying technique is widely used to provide a wear-resistant coating on steel substrate in industry. Plasma spraying is one main method of thermal spraying in producing thick wear-resistant
~~~
Table 2 Some aroperties of the Cr,C,-NiCr coating 5.92 Density (Mg/m3) Porosity (YO) 6.3 Microhardness (Hv, 2 ) 890 Young’s modulus (GPa) 102 Bend strength (MPa) 148 Fracture toughness (MPam”’) 7.4
coatings [ 1,2]. Because of their thermal stability, Cr,C,NiCr coatings sprayed by plasma or other thermal spraying methods are often employed in high temperature atmospheres such as turbine engine and steel
industry, which
are
commonly used
for
temperatures ranging from 530 to 815 “C. Several studies have contributed to the processing, characterization and tribological properties of thermal sprayed Cr,C,-NiCr coatings [3-71. However, it is still necessary to further evaluate the tribological properties of this kind of coatings. This work investigated the selfmated tribological properties of plasma-sprayed Cr,C,NiCr coating with respect to load and sliding speed, and the results obtained indicated some new phenomena about the tribological properties of the Cr,C,-NiCr coating.
EXPERIMENTAL DETAILS Cr,C,-NiCr coating samples formed to the specific
Friction and wear tests were conducted on a ring-onring arrangement of an MM-200 wear tester in air at room atmosphere (see Fig. 1). Both the stationary and moving rings were 40 mm outer diameter, 16 mm inner diameter and 10 mm thick. The surface roughness Ra of the Cr,C,-NiCr coating before the friction and wear tests was 0.25 pm after polishing. The tests were performed under the following conditions: the loads of 100, 200, 400,600,800 and 1000 N, and two rotational speeds of 200 and 400 revlmin, which were equal to the two sliding velocities of 0.42 and 0.84 mls. The friction coefficients were obtained from the friction torques, which were directly read from the tester, being divided by the loads and ring radius. The wear coefficients were acquired from the wear mass loss, which were measured by weighing the samples before and after each of the wear tests with a TG328B analytical balance, being divided by the loads, sliding distance and density of the
187
,,Lx " I
200 revlmin
FI:
03-
IOON-L 600N-
4-
--t
-
0 2
.
.
.
*
200N-v-400N
800N
'
.
'
07
.
'
400 revlmin
Fig. 1 Schematic representation of ring-on-ring geometry coating. Prior to weighing, the samples were cleaned in
an ultrasonic bath with acetone for 30 min and then dried in an oven at 120 "C for 30 min. An EPMA-8705QHII type scanning electron microscope (SEM) was used to observe the worn surface morphologies, and an NIE-7 199C Fourier transform infrared spectroscopy (FTIR) to analyze the debris.
-.*
I 0I0
1 OON-A200N--&600N-0-8~0N--*-.
200 '
.
300
400N 1000N 400
,
500
0
Sliding dtstance (rn)
Fig. 3 Friction coefficient with respect to load and sliding distance at sliding speed of 400 rev/min inn Stationary nng 200 revlmtn -0-Moving nng 200 r e v h i n -8Stationary nng. 400 revlmin -.-Moving nng 400 revlmin -n-
RESULTS AND DISCUSSION Figures 2, 3 and 4 indicate the results of the friction and wear tests. From Figs. 2 and 3, it can be seen that the friction coefficient at steady state decreased with increase in load and sliding speed; however, the nature of wear-in varied with respect to load and sliding speed. A rapid initial rise in friction followed by a fall to the steady state for 100, 200, 400 and 600 N at 200 rev/min and 100 N at 400 rev/min, but a rapid initial fall in
Load (N)
Fig. 4 Wear coefficient with respect to load and sliding speed
friction followed a slight rise to the steady state for all the other test conditions. The above results somewhat disagreed with those that the fiiction coefficients of
speed, and the lower the sliding speed, the more the difference of wear coefficient between the stationary
thermal-sprayed Cr,C,-NiCr coatings were measured
and moving rings was. For the same load, the wear
under other counterpart materials and operating
coefficient of stationary ring at sliding speed of 200
conditions [3-71. This demonstrates that the tribological
rev/min was close to that of stationary ring at sliding
properties of thermal-sprayed Cr,C,-NiCr coatings are
speed of 400 revlmin. With respect to load, the wear coefficient of moving ring reduced from 100 to 600 N and then rapidly rose from 600 to 800 N at lower sliding speed. Although the previous studies [3,6,9J exhibited that thermal-sprayed Cr,C,-NiCr coatings were not very
rather complex and should be further examined. Figure 4 shows that the wear coefficient o f stationary ring was noticeably lower than that of the moving ring mated to it at the identical load and sliding
188
wear-resistant, the wear coefficient near lo4 mm3N-'m-'
was more remarkable than that on those of the
can be obtain in Fig. 4 at the sliding speed of 400 rev/min and the load among 200 and 400 N.
corresponding stationary rings. At low load and sliding speed, lamellar spallation apparently took place and the
Figures 5 and 6 are the SEM micrographs of the worn surfaces for the coating samples at different sliding speeds, respectively. The original polished surface and microstructure of the coating were similar to those of the Cr,C,-NiCr coating presented in a previous paper [ 6 ] . The coating consisted of flat plate-like lamella oriented parallel to substrate and possessed some pores and microcracks. From Figs. 5 and 6 , it can
interfaces of lamella appeared on the worn surfaces (Fig. 5 (a), (b), (d), (0 and Fig. 6 (b)). With an increase in loads and in sliding speeds, the lamellar spallation mainly transformed into particle fracturing within single
be sen that there were many holes, ffactures and cracking on all the worn surfaces, indicating that cracking occurred on the coating surfaces. However, the cracking on the worn surfaces of all the moving rings
lamellar, and the worn surfaces also expose abrasive score marking, plastic deformation and shear fracture (Fig. 5 (c), (e), Fig. 6 (a), (c)-(0). The greater the load and sliding speed, the more intense the plastic deformation took place. In order to explore the tribochemical mechanism of the Cr,C,-NiCr coating during the friction and wear tests, FTIR was used to analyze the debris carefully
Fig. 5 SEM micrographs of worn surfaces at sliding speed of 200 re vhin: (a) stationary ring of 100 N, (b) moving ring of 100 N; (c) stationary ring of 600 N, (d) moving ring of 600 N; (e) stationary ring of 800 N, (f)moving ring of 800 N
189
Fig. 6 SEM micrographs of worn surfaces at sliding speed of 400 revhin: (a) stationary ring of 100 N, (b) moving ring of 100 N; (c)
stationary ring of 600 N, (d) moving ring of 600 N; (e) stationary ring of 800 N, (0 moving ring of 800 N collected up. Fig. 7 gives the FTIR spectra of as-sprayed
and sliding speed. Adhesive wear readily occurred for
Cr,C,-NiCr coating and some debris. The as-sprayed
self-mated Cr,C,-NiCr coating during sliding [9]. At
coating and debris of 400 rev/min and 100 N had no
lower load and sliding speed, the rapid initial rise of
marked FTIR absorption peak, which suggested that they mainly consist of chromium carbides. The debris of
fiiction coefficient during wear-in may be explained in terms of overcoming initially high adhesive contact
200 rev/min, 800 N and 400 rev/min, 800 N showed the
between stationary and moving rings and also the
FTIR absorption peaks of Cr203 [S]. The results
absorbed contaminated layer on surface of the coating
revealed that Cr,C,-NiCr oxidated due to fiiction heat
[lo]. It was likely that work hardening of the coating
during sliding, and Cr20,formed at the sliding contacts. According to the above SEM and FTIR results, it can attribute the dependence of fiction and wear coefficients on load and sliding speed to the variation of flash temperature and adhesive intensity with the load
under compressive stress then led to a decrease in the
190
fiiction coefficient after the initial increase [ 5 ] . Work hardening caused the adhesive force between the rings to decrease, and thus the fiiction force and friction coefficient to decrease. At higher load and sliding speed,
I
I
d coating 200 rev/min and 800 N 400 rev/min and 800 N 400 rev/min and 100 N
induced thermal shocking deteriorated the wear performance of the coating. This is readily explained by the following fact: at the identical load, the higher the sliding speed, the more the difference of wear coefficient between the stationary and moving rings was because the higher the sliding speed, the weaker the thermal shocking was [ l I]. It is also easily explained that the wear coefficient of moving ring then rapidly rose from 600 to 800 N at low sliding speed. This results from that the increase in wear caused by more intense thermal shocking exceeded the decrease in wear due to decrease in adhesive force between the rings.
CONCLUSION Friction and wear coefficients of self-mated plasma2000
1000
1500
500
Wavenumbers Fig. 7 FTIR spectra of as-sprayed and some debris
on the one hand, more intense work hardening occurred as a result of higher compressive stress. On the other hand, Cr,C,-NiCr coating oxidized and harder Cr,O, formed at the sliding contacts due to higher flash temperature resulted from friction heat. The two reasons, especially for the later, resulted in that the initial rise of friction coefficient was too fast to expose in the friction curves, and the friction curves firstly showed a rapid fell during wear-in (Fig. 2 and 3). With the formation of Cr,O, and then embedding into the softer matrix, the tangential friction forces increased because of the plowing action, in turn a slight increase in the friction coefficient after the initial rapid decrease during wear-in.
sprayed Cr,C,-NiCr coating were tested with respect to load and sliding speed. It was found that the friction and wear coefficients depended on load and sliding speed. The wear coefficients of moving rings were larger than those of corresponding stationary rings, and the higher the sliding speed, the larger the difference of wear coefficient between moving ring and stationary ring was. The results obtained were related to friction-induced thermal shock and oxidation of the coating caused by frictional heat, which caused that the main wear mechanism of the coating transformed into particle fracturing within single lamellar from lamellar spallation with an increase in loads and in sliding speeds. At higher sliding speed and a certain range of loads, the coating exhibited rather wear-resistant and the wear coefficients close to 10-6mm3N-'m-'.
This can be proved from the abrasive score markings in
REFERENCES
Fig. 5 (el, (0,Fig. 6 (c), (d), (el and (0. Work hardening of the coating and the formation of
(1) R. W. Smith and R. Novak, Advances and
Cr,O, caused the adhesive force between the rings to decrease, which resulted in that the main wear mechanism of the coating varied with load and sliding speed, and hence the wear coefficient of the coating decreased with increase in a certain load and sliding speed (Fig. 4). As the surfaces of moving rings recurrently entered sliding contacts with
higher
temperatures and then were exposed to cooler atmosphere, they were necessarily subjected to thermal shocking in some degree. That the wear coefficients of stationary rings were noticeably lower than those of the moving rings mated to them disclosed that friction-
Application in U. S. Thermal Spray Technology, I. Technology and Materials. Powder Metallurgy International 3, (1991) 147-155. (2) G. Barbezat, A. R. Nicoll and A. Sickinger, Abrasion, Erosion and Scuffing Resistance of Carbide and Oxide Ceramic Thermal Sprayed Coating for Different Application. Wear, 162-164, (1993) 529-537. (3) Y. Wang, Y. S. Jin and S. Z. Wen, The Friction and Wear Performance of Plasma Sprayed Ceramic Coatings at High Temperature. Wear, 129, (1989) 223-234.
191
(4) G . Barbezat, A. R. Nicoll and Y. S. Yin et al, Abrasive Wear Performance of Cr3C,-25%NiCr Coatings by Plasma Spray and CDS Detonation Spray. Tribology Transactions, 38, (1995) 845-850.
1430. (8) R. A. Nyquist and R. 0. Kagel, Infrared Spectra of Inorganic Compounds, Academic Press, New York
( 5 ) M. Mohanty, R. W. Smith and M. De Bonte et al, Sliding Wear Behavior of Thermally Sprayed 75/25 Cr,C,/NiCr Wear Resistant Coatings. Wear, 198, (1996) 251-266. (6) J. F. Li, C. X. Ding and J. Q. Huang et al, Wear Mechanism of Plasma-Sprayed Cr,C,-NiCr Against TiO, Coating. Wear, 2 1 1, (1997) 177- 184.
and London, (1971) 94-95,216-217. (9) Y. S. Jin, W. Xia and H. Cheng, Experimental Research on Tribological Behaviors of Various Plasma-Sprayed Ceramic Coatings. Journal of Tsinghua University, 32, (1 992) 17-25. (10) S. Z. Wen, Tribological Principle, Tsinghua University, Beijing, (1990) 436-445. (1l)Z. B. Hou, S. J. He and S. X. Li et al, Heat
(7) J. Takeuchi, Y. Murata and Y. Harada et al, An
Conduction of Solid, Science and Technology Press,
Improvement of Cr,C,-NiCr
Sprayed Coatings
Followed by Chromium Diffusion Treatment. Proceedings of the 15th International Thermal Spray Conference, Nice, France, (1998) 1425-
192
Shanghai, (1984) 88-94.
ROLE OF GRAIN SIZE IN SCRATCH DAMAGE RESISTANCE IN ZIRCONIAS AND SILICON NITRIDES Seung Kun Lee, Robert P. Jensen, Michael J. Readey Advanced Materials Technology Caterpillar Inc., USA
ABSTRACT Scratch damage in zirconias (Mg-PSZ, Y-TZP, and CeTZP) and silicon nitrides (fine, medium, coarse grain) are investigated. Scratch testing is carried out using a conical diamond indenter. In all materials the damage mode changes from smooth plastic deformation to limited cracking with increasing scratch load: in MgPSZ, plastic deformation is predominant at lower loads, with microcracking at higher loads; in Y-TZP, plastic deformation is predominant over the range of the test loads-macrocracks initiate only at relatively high loads, but penetrate to a relatively large depth; again, Ce-TZP shows intermediate behavior, but with cracking patterns closer to that of Mg-PSZ. Bending tests on specimens subjected to scratch damage indicate a relatively high damage tolerance in the MgPSZ and Ce-TZP; Y-TZP shows the highest initial strength, but suffers relatively large strength loss above the critical load for cracking. In Y-TZP, fine Y-TZP showed smooth scratch groove without cracking. As grain size increased, the scratch groove became rougher with the development of microcracking. In Si3N4, the critical load for a transition from plastic deforamtion to microscracking decreased with increasing grain size. Coarse Si3N4 has a higher material removal rate than fine Si3N4. The material removal mechanism is attributed to grain pullout in this material. Implications concerning relative merits of each material for wear properties, contact fatigue, and machining damage are briefly discussed.
INTRODUCTION Advanced ceramics such as zirconia and silicon nitride have been identified as choices for sliding components in a variety of engineering applications, including engine components (bearings, rollers, dies, tappets, valves, fuel injectors), where contact, scratch, and wear damage are critical factors for lifetime performance [ 1 41. In zirconia ceramics, energy absorptive phase transformation from metastable tetragonal phase to monoclinic phase enhances fracture [5-91. Much progress has been made during the past 25 years in this kind of toughness enhancement in zirconias, by microstructural control through modification of selective additive phases (MgO, Y203, CeO) and heat treatments [ 10- 131. In silicon nitride ceramics, on the other hand, there has been a great success in improving strength and toughness by controlling composition and microstructure over the decade. Self-reinforced silicon
nitride with elongated grain structure provides an excellent mechanical-reliability associated with high toughness and strength [ 14,151. Properties like fatigue [16] and machining and wear [17,18] have been demonstrated to be strongly affected by such microstructural factors. Scratch testing using a translating sharp point provides more direct information on potential processes in wear and machining operations [17,18,19]. Data from this test procedure may be expected to provide complementary basic information on damage modes in engineering contact and sliding applications, as well as on material removal in wear and machining and flaw development in strength determination. In this study we investigate scratch damage in selected zirconia ceramics (Mg-PSZ, fine grained Y-TZP, medium Y-TZP, coarse Y-TZP, and Ce-TZP) and silicon nitride ceramics (fine grained Si3N4, medium Si3N4, coarse Si3N4). The results demonstrate a need for caution in materials selection for applications in severe contact conditions.
EXPERIMENTAL PROCEDURE Materials Commercially available zirconias and silicon nitrides were used in this study. In order to investigate effect of grain size, the three types of Y-TZP ceramics with different grain size were made by a routine ceramic processing. The zirconia powders (3Y-TZP, Tosoh Co. Tokyo, Japan) were compacted in a steel die of 50 mm diameter at a pressure of 50 MPa, followed by sintering at either 1450°C for 1 h, 150O0Cfor 2 h, or 1600 OC for 3 h. These heat treatments produced the “fine” (F-YTZP), “medium” (M-Y-TZP), and “coarse” (C-Y-TZP) microstructures. Microstructures of zirconias and silicon nitrides are presented in Fig.1 and 2, respectively. Mechanical properties of zirconias and silicon nitrides investigated in this study are summarized in Table 1 and 2, respectively. Mg-PSZ had a large grain structure (grain size = 58 pm) and higher toughness but lower strength (Table 1). Y-TZP had a fine microstructure (grain size = 0.4 pm) and higher strength and hardness but lower toughness. Ceshowed intermediate toughness and lowest strength. CFI3208 has a fine grain structure with a lower toughness. AS800 showed a coarse microstructure with large elongated grains. GS44 exhibited a bimodal grain structure with high strength and toughness.
193
Table 1. Characteristicsof zirconia ceramics
Al, Y
Grain size (P) Fine (0.5)
Strength (MPa) 730 f 34
Toughness (MPa-m'n) 5 .O
Hardness (GPa) 14.8
Al, Y La
Medium ( I . 1) Coarse (3.2)
959 k 44 707 +_ 39
7.5 7.2
14.8 14.2
Materials
Suppliers
Second Phase
CFI3208
Ceramic for Industries Honeywell Honeywell
GS44 AS800
Fig. 1. Microstructures of Mg-PSZ (a), Ce-TZP (b), and Y-TZP (c).
194
Fig.2. Microstructures of CFI3208 (a), GS44 (c), and AS800 silicon nitride (c).
Fig 1. Top and side views of scratch damage of Mg-PSZ (a), Y-TZP (b) and Ce-TZP (c), scratch load, P = I 0 0 N.
Scratch tests Scratches were made using a sliding conical diamond indenter with apex angle 120' and spherical tip radius 200 pm (Automatic Scratch Tester, CSEMREVETEST, Neuchatel, Switzerland). In automatic test mode, the load is increased continuously as the point translates across the surface. Both normal and tangential forces were recorded. Normal loads ranged from P = 5 N to 130 N over about 20 mm sliding distance, at sliding speed 20 mdmin, in air. Crosssectional area of the scratches were measured by a surface profilometer, to obtain an estimate of the volume removed per unit sliding distance. Some additional scratch tests were made on selected specimens, at prescribed constant loads. Bonded-interface specimens [ 16,18,19,20] were used to obtain section views through the indentation and scratch damage zones in each material. Specimens were cut into two half-blocks. The side surfaces of the half-blocks were first polished and then clamped faceto-face with an intervening thin layer of adhesive. The top surfaces were then repolished. Indentations were made along the surface traces of the bonded interfaces, at load P = 3000 N. Scratch tests were made perpendicular to these interfaces, at load P = 100 N. The adhesive joining the interfaces was subsequently dissolved in acetone. Separated half-blocks were gold coated for top- and side-surface examination in Nomarski illumination. Four-point bending tests were conducted on bars 3 x 4 x 45 mm (outer span 40 mm, inner span 20 mm) that had been subjected to scratch damage. The scratches were made on the prospective tensile surfaces, with the scratch axis parallel to the prospective bend axis, at loads P = 20N to 130 N. The damage areas were covered with a drop of dry silicon oil before flexure and broken in fast fracture (
specimens were examined fractographically to locate the source of failure, either scratch damage or extraneous flaws. Control strength tests were made on unscratched specimens, to determine baseline laboratory strengths.
RESULTS Zirconia ceramics Figure 1 shows surface and side views of scratch damage of the three materials from bonded-interface specimens at P = 100 N. In Mg-PSZ (Fig. 1 a) a diffuse microcrack zone extends well below the surface groove into the subsurface. In the near-surface regions, these microcracks have coalesced to produce extensive chipping. In Y-TZP (Fig. Ib) no microcracking is observed at all: instead, a deeply penetrating median crack is evident immediately below the scratch. In CeTZP (Fig. lc) limited microcracking is observed, on a smaller scale than in Mg-PSZ. Figure 2 is a plot of volume removed per unit sliding distance as a function of normal load for the three materials, evaluated from surface profilometer traces across the scratches. Mg-PSZ ultimately shows the most severe wear rate, above load 75 N where surface microcracking and chipping are manifest. Y-TZP shows the lowest removal rate, over the entire range of loads. Figure 3 compares strength degradation data for scratch damage on the three materials. In the Mg-PSZ and Ce-TZP, strength falls off beyond the load for extensive microcracking, but only slowly, indicating high damage tolerance. Y-TZP, on the other hand, show a typical brittle response; i.e. no perceptible degradation in initially high strength up to about 80 N, but with subsequent abrupt drop-off at higher loads,
195
highlighting the effectiveness of the median crack in Fig. 3c. I
I
I
I
scratch load P = 100 N. F-Y-TZP shows a smother surface track with no microcracking, indicative of plastic deformation. In C-Y-TPZ, on the other hand, extensive microcracking and larger debris (see arrows in Fig. 4c) are evident around the scratch groove, indicating brittle response in this material. M-Y-TZP shows an intermediate feature. Scratch damage mode depends strongly on grain size. Smaller grain size is responsible for the smoother scratch track associated mainly with plastic deformation in F-Y-TZP.
S a m h load (N) Fig 2. Volume removed for Mg-PSZ, Ce-TZP, and Y-TZP, as a function of load. Solid lines are empirical fits.
1200,
I
I
I
I
I
I
Fig. 3. Strength in scratching of Mg-PSZ, Y-TZP and CeTZP. Closed symbols represent failures from scratch origins, open symbols failures from other flaw origins; shaded boxes at left are strengths of as-polished specimens.
Fig. 4. Surface views of scratch damage in in F-Y-TZP (a), M-Y-TZP (b), and C-Y-TZP (c), produced at a scratch load P = l00N. Figure 4 presents surface views of scratch damage in
F-Y-TZP, M-Y-TZP, and C-Y-TZP, produced at a
196
Fig 5. SEM micrographs of surface views of scratch damage in CF13208 (a), GS44 (b), and AS800 silicon nitride (c).
Scratch load (N)
Fig. 6. Material removal rate as function of scratch load in silicon nitrides.
Silicon nitride ceramics Figure 5 shows SEM micrographs of the surface views of scratch damage in CF13208, GS44, and AS800 silicon nitrides. CFI3208 with a fine grain structure exhibited a smooth scratch groove with less cracking. AS800 with a coarse grain structure exhibited a rough scratch track with grain pullout and dislodgement. Microcracking and chipping are manifest in AS800, indicating intensive interaction between material and an indenter. GS44 showed intermediate cracking behavior. Figure 6 is a plot of volume removed per unit sliding distance as a function of normal load for CH3208, GS44, and AS800, evaluated from surface profilometer traces across the scratches. AS800 ultimately shows the most severe wear rate. CFI3208 and GS44 show lower removal rates, consistant with scratch groove in Fig. 5.
DISCUSSION We investigated the effect of grain size on scratch damage mode in zirconias and silicon nitrides. Scratch test techniques provide useful information to determine scratch damage and to simulate machining damage [ 17,18,2I]. It is also relevant for evaluating abrasive wear resistance. Mg-PSZ with large grain size (>50 pm) shows the least smooth scratch track, with extensive microcracking and associated chipping above a critical load, resulting in a high wear rate at higher loads. Y-TZP shows the smoothest track, and lowest wear rate. Ce-TZP is again intermediate, with modest microcracking. Y-TZP is the most brittle of the three materials, as reflected by the formation of a deeply penetrating median crack in the subsurface region above a critical load. Accordingly, Y-TZP is much more susceptible to abrupt strength loss in severe scratching conditions; its strength drops from almost twice to less than one half that of its counterparts over the data range. Mg-PSZ and Ce-TZP, on the other hand, while still subject to failure from scratch sites above a threshold load, are much more damage tolerant.
In Y-TZP, as grain size increased, the damage mode in Y-TZP changed from plastic deformation to microcracking. A similar trend has been reported on alumina [ 181. This result indicates the important implications on abrasive wear and machining. Grain size plays a major role in the properties determined at the microstructural level such as scratch, wear, and machining. Large grains lead to a rougher scratch track and a higher material removal rate. Again, hightougheness zirconia such as Mg-PSZ and C-Y-TZP would not necessarily enhance wear properties, and may even degrade them. F-Y-TZP with fine grains is likely to exhibit better performance where high wear resistance is an issue. In silicon nitride, as grain size increased, scratch damage mode changed from plastic deformation to brittle microcracking. Grain pullout is evident in AS800, indicating severe scratch damage. AS800 with coarse gains is expected to be less abrasive wear resistance than CFI3208 and GS44 that have smaller grain structures.
CONCLUSIONS In all materials the damage mode changes from smooth plastic deformation to limited cracking with increasing scratch load. In Mg-PSZ, plastic deformation and extensive distributed surface and subsurface microcracking; in Y-TZP, a smooth scratch track with limited plasticity and, above a critical load, deep cracks; in Ce-TZP, similar to Mg-PSZ, but less pronounced microcracking. Bending tests on specimens subjected to scratch damage indicate a relatively high damage tolerance in the Mg-PSZ and Ce-TZP; Y-TZP shows the highest initial strength, but suffers relatively large strength loss above the critical load for cracking. In Y-TZP, fine Y-TZP showed smooth scratch groove without cracking. As grain size increased, the scratch groove became rougher with the development of microcracking. In Si3N4, the critical load for a transition from plastic deforamtion to microscracking decreased with increasing grain size. Coarse Si3N4 has a higher material removal rate than fine Si3N4. The material removal mechanism is attributed to grain pullout in coarse grain materials.
Acknowledgments This research was sponsored by the Department of Energy in USA, under Contract DE-FC05-970R22579.
REFERENCES R.H. J. Hannink, M.J. Murray, H.G. Scott, Friction and Wear of Partially Stabilized Zirconia: Basic Science and Practical Applications, Wear, 100 (1984) 355-66. G.W. Stachowiak, G.B. Stachowiak, Unlubricated Friction and Wear Behavior of Toughened Zirconia Ceramics, Wear, 132 ( 1989) 15 1-7 1.
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(3) A. Gangopadhyay, H.S. Cheng, J.F. Braza, S. Harman, J.M. Corwin In Fiction and Wear of Ceramics; Said Jahanmir, Ed., Marcel Dekker, New York, 1994, p 329-56. (4) Y.M. Chen, B. Rigaut, F. Armanet, Wear Behavior of Partially Stabilized Zirconia at High Sliding Speed, J. Europ. Ceram. Soc., 6 (1990) 383-90. (5) R.M. McMeeking, A.G. Evans, Mechanics of Transformation Toughening in Brittle Materials, J. Am. Ceram. Soc., 65 (1 982) 242. (6) P.F. Becher; G. Begun; E.F. Funkenbusch In Science and Technology of Zirconia Ill, Advance in Ceramics, S . Somiya, N. Yamamoto, H. Yanagida, Eds., The American Ceramic Society, Inc., Westerville, Ohio, 1988, Vol. 24B, p 64551. (7) A.G. Evans, Perspective on the Development of High-Toughness Ceramics, J. Am. Ceram. Soc., 73 [2] (1990) 187-206. (8) D.B. Marshall, Reversible Stress-Induced Martensitic Transformation in ZrO2, J. Am. Ceram. SOC.69 [3] (1986) 215-17. (9) D.J. Green, R.H.J. Hannink, M.V. Swain Transformation Toughening of Ceramics, CRC Press, Boca Raton, FL, 1989. (10) R.C. Garvie, R.H.J. Hannink, R.T. Pascoe, Ceramic Steel?, Nature, 258 (1975) 703. ( I 1 ) Science and Technology of Zirconia, A.H. Heuer, L.W. Hobbs, Eds., American Ceramic Society, Cleveland, OH, 1981, Vol. 3. (12) Science and Technology of Zirconia 11, N. Claussen, M. Ruhle, A.H. Heuer, Eds., The American Ceramic Society, Columbus, OH, 1984,Vol. 12. (13) Science and Technology of Zirconia I l l , S. Somiya; N. Yamamoto, H. Yanagida, Eds., The American Ceramic Society, Westerville, OH, 1988, Vol. 24A,B. (14) C-W. Li, S-C. Lui, J. Goldacker, Relation Between Strength, Microstructure, and GrainBridging Characteristics in In-Situ Reinforced Silicon Nitride, J. Am. Ceram. Soc., 78 [2] (1995) 449-59. (15) P. Sajgalik, J. Dusza, M.J. Hoffmann, Relationship between Microstructure, Toughness Mechanisms, and Fracture Toughness of Reinforced Silicon Nitride Ceramics, J. Am. Ceram. SOC.,78 [ 101 (1995) 2619-24. (16) A. Pajares, F. Guiberteau, B.R. Lawn, S. Lathabai, Hertzian Contact Damage in MagnesiaPartially-Stabilized Zirconia, J. Am. Ceram. SOC., 78 [4] ( 1 995) 1083-86. (17) M.V. Swain, Microfracture About Scratches in Brittle Solids, froc. Roy. SOC. Lond., A366 (1 979) 575-97. ( 1 8) H.H.K. Xu, S. Jahanmir, Microfracture and Material Removal in Scratching of Alumina, J. Muter. Sci., 30 (1 995) 2235-47. (19) N. Axen, L. Kahlman, I.M. Hutchings, Correlations between Tangential Force and Damage Mechanisms in the Scratch Testing of
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Ceramics, Tribology International, 30 [7] (1997) 467-74. (20) F. Guiberteau, N.P.Padture, H. Cai, B.R. Lawn, Indentation Fatigue: A Simple Cyclic Hertzian Test for Measuring Damage Accumulation in Polycrystalline Ceramics, fhilos. Mag., A 68 [5] (1993) 1003-16. (21) H.H.K. Xu, S. Jahanmir, Scratching and Grinding of a Machinable Glass-Ceramic with Weak Interfaces and Rising T-Curve, J. Am. Ceram. SOC.,78 [2] (1995) 497-500.
CERAMIC COATINGS WITH SOLID LUBRICANT ABILITY FOR ENGINE APPLICATIONS M. Buchmann, R. Gadow, A. Killinger University of Stuttgart, Institute for Manufacturing Technologies of Ceramic Components and Composites (IMTCCC) Allmandring 5b, D-70569 Stuttgart, Germany
ABSTRACT Recent automotive engineering developments concerning fuel consumption regulations and decreasing material and manufacturing cost result in an increasing utilization of light metal components for automotive applications. Significant weight savings are obtained by a modification of the engine block material from cast iron to aluminum. Since all parts of a combustion engine interact as a system the single components must sustain the combustion pressure and temperature as well as wear and friction effects of moving surfaces in different environment and lubrication. Approaches to increase efficiency and lifetime of light metal engine blocks are protective thermally sprayed coatings on cylinder bores. The used thermal spray processes are high-energetic (Atmospheric Plasma Spraying) and high-energetic hypersonic processes (High Velocity Oxygen Fuel Spraying). The knowledge of the mechanical and thermophysical properties of composite materials is a key requirement for an optimized stable and repeatable manufacturing process as well as for reproducible high quality composites. This paper describes an actual overview about the material screening, the manufacturing technology and the measured coating characteristics. Residual stress measurements are performed and the effects on the coating properties like hardness, friction and wear is investigated.
INTRODUCTION Future engine design standards demand low environmental impacts, enhance operation lifetime as well as increased economy and energy efficiencies. Savings in fuel consumption require an improved motor design together with a reduction of the total vehicle weight. Nowadays the mass proportion of the engine block on the total automotive weight is in the range between 10 to 15 %. Therefor light weight engine design offers a great potential for a successful mass reduction [I]. The state of the art of light metal engines are cast aluminum crankcases. The worse corrosion, wear and friction behaviour of aluminum surfaces in a heavily loaded tribological system like cylinder liner - piston ring, make a wear and corrosion protection of the surface indispensable. Commonly used protection systems are integrated iron bushings, galvanic coatings based on chromium or nickel, as well as hypereutectic AlSi engine block alloys. Demands to reduce the production process complexity, the process steps and manu-
facturing cost on the one side and to improve the operation behaviour, the environmental safety and recycling process on the other side lead to the investigation of thermally sprayed coatings as protective surfaces on light metal alloys. Especially of tribological interest are coating systems with low friction and wear coefficients. In future tribological systems in engines will operate with a lubricant film thickness lower than 1 pm or with biologically degradable fluids or life-time lubrication or even under dry condition. If traditional liquid lubricants cannot be used, the tribological functions must be taken over by material surfaces with solid lubricant capabilities [2].
THERMALLY SPRAYED COATINGS FOR CYLINDER BORES The thermal spray process offers the possibility to apply a broad variety of metallurgical, cermet and ceramic coatings on the surface of machine components, even on systems with complex geometry. The coatings are commonly used to improve the wear and corrosion resistance, the operation temperature and thermal shock behaviour or to influence the electrical, magnetic and biological behaviour of the composite surface. The latest technological features are high energetic, high velocity coating systems. The HVOF process uses liquid fuels or fuel gases for high energetic combustion with oxygen 300 - 600 d s ; emax 2.500 - 3.200 "C). It leads to (vextremely dense coatings because of the high kinetic energy of the hot powder loaded gas jet. During the APS process temperatures up to 20.000 "C are obtained, therefor this system is mainly used for refractory materials. Precedent to the coating process a grit blasting and degreasing of the substrate surface is performed. The roughening of the surface with corundum of defined size improves the mechanical adhesion of the coating and induces compressive stresses into the substrate material. Following to the coating deposition a mechanical (grinding, polishing, honing) or thermal post-treatment of the coating surface takes place. In the case of cylinder bore coatings two principal different methods can be taken into consideration. State of the art for inside coatings are rotating plasma spray devices. The engine block is totally fixed and the bore is coated by means of a vertically moving rotating plasma torch. A second very promising method is the deposition of inside coatings with a fixed A P S or HVOF spraying
-
-
199
gun and a rotating engine block. The main advantages of the second technique are the improved coating microstructure by less porosity and coating adhesion, because of an enhanced jet propagation and thus particle velocity. Coating studies are performed on aluminum tubes (d = 95 mm, 1 = 150 mm) as well as on 4 cylinder aluminum crankcases. During the coating process the tubes or engine blocks are totally fixed on a rotating table with a rotation speed between 50 - 100 rpm. For the plasma coating process an F1 inside torch is used with a spraying distance of 45 mm perpendicular to the substrate surface. With this system a maximum operation power of 25 kW can be realized. For the plasma process a two axis feed drive system is utilised. The coating is deposited in several transitions with a constant vertical feed drive between 5 - 10 mm/s. Besides the plasma process also the HVOF process with the TOP GUN@ system is investigated for the deposition of inside coatings, compare figure 1. The TOP GUN@ system offers additionally the possibility to use acetylene as fuel gas, performing sufficient high temperatures to melt refractory materials. By means of a 7 axis robot system, including the regulation for the rotating table, variable spraying angles (30' - 90') as well as different spraying distances between 150 and 250 mm, dependent of the used spray powders, can be flexibly programmed. The HVOF coating process is adjusted in a way that the top of the cylinder bores, where the maximum wear mechanisms are assumed, are coated perpendicular, whereas the less loaded bottom regions are coated with an decreasing spray angle of 30 degrees. By adapting the feed rate of HVOF gun on the spraying angle a homogeneously coating thickness over the cylinder bore depth can be reached.
Fig. 1 HVOF equipment for the inside coating of light metal cylinder bores. To avoid an overheating of the substrate during HVOF spraying and to improve the coating microstructure a special liquid COz inside fan-shaped cooling system was developed. With a mass flow of 1.7 kg C 0 2 per minute the
200
maximal process temperature during HVOF spraying are lower than 250'C, measured pyrometrical on the aluminum tube outside. The use of a liquid COz system combines the advantages of a very efficient and low cost cooling system, compare figure 2.
Fig. 2 Left side - inside fan shaped CO2 cooling system, right side - pyrometrical measured temperature distribution on the outside of an AlSi tube after the coating process
MATERIAL SCREENING AND CHARACTERISTICS The material screening is focussed on the development of high efficient coating systems for tribological applications with regard to low friction coefficients and wear rates as well as to high corrosion resistance. Beside sliding speed, surface roughness and ambient conditions like temperature and humidity the tribological behaviour is mainly influenced by the used material systems and their characteristics. Of special interest are materials with solid lubricant properties or the capability to form lubricious oxides under tribological conditions, like oxides of titanium and vanadium as well as molybdenum and tungsten. Both types of materials can be characterized with low shear strengths caused by planar defects which are arranged in parallel crystallographic shear planes. For this kind of materials low friction coefficients and wear rates can be expected also under mixed and dry friction conditions [3]. To get a general material overview thirteen different ceramic, cermet and metallurgical coating systems were selected as coating material and investigated. In combustion engines with dynamic thermal load applications the operation behaviour of layer composites is mainly influenced by the thermophysical properties of the different composite materials. The thermal expansion coefficient a of the coatings represents a very important material property influencing the lifetime of coated components. Thermal expansion mismatches between substrate and coating material result in the development of residual stresses during thermomechanical load operations. Critical residual stresses caused by the manufacturing process in addition with operation loads can lead to component failure. Table 1 summarises the investigated coating materials and the measured average values of
a and the molar heat cp in the temperature range from 100 to 600 "C. Dilatometer measurements are performed on powder samples (ap)and specially prepared spraycones (@) as reference 'quasi-bulk' material and compared with the AlSi9Cu alloy data. The heat capacity cp is measured by differential thermoanalysis (DTA, DSC).
AISi9Cu Ti02
1 1
8.4
I I
28.5 9.4
1 1
0.73
A120fli02-60/40
5.97
5.4
0.92
Cr203
7.21
7.4
0.71
Cr203/Ti02-60/40
7.84
7.5
0.73
Cr20fli02-75/25
I
7.8
1
7.5
I
0.74
MOO^
5.0
5.8
0.73
AI203/ZrO2-60/40
6.9
7.2
0.83
Cr2C9/NiCr-80/20
9.1
10.7
0.55
(Ti,Mo)(C,N)/NiCr
6.0
9.4
0.42
CrB/NiCr-75/25
10.7
10.5
0.44
I Mo/NiCrBSi-30/70 I
Mo
MoICr-Steel-70130
4.7
-
I I
5.0 12.2 9.75
I I
1 1
0.91
0.24 0.42
I
I 1
0.30
Table 1 Material screening and thermophysical material properties
COATING CHARACTERIZATION The quality of thermally sprayed coatings with regard to the residual stress situation, microstructure, surface roughness, coating porosity, hardness and mechanical properties can widely be varied by tuning the equivalent spraying parameters like energy supply, substrate preheating and simultaneous process cooling. During thermal spraying the spray powder is partially or even fully molten within milliseconds, accelerated to high velocities and propelled onto the substrate surface. The more or less lamellar microstructure of many thermally sprayed coatings show different phases, due to the rapid solidification and quenching processes, a macro and microporosity as well as oxidation areas between the single splats if metals are applied in contact with the atmosphere. The correlation between microstructure and coating properties has to be known to optimize the spraying parameters.
RESIDUAL STRESSES AND BONDING STRENGTH Residual stresses arise during the thermal spray process and influence significantly the coating quality and composite performance. Critical residual stresses can
cause failure of coatings in form of delamination and buckling effects in the coating and the interface as well as they cause perpendicular cracking if the stress level reaches the ultimate strength of the coating or lead to plastic strain in the interface. Tensile residual stresses in the coating reduce the component lifetime since this favours crack formation and propagation. Furthermore condensed corrosive products can penetrate the coating through microcracks, destabilize the coating and attack the substrate material or the interface layer. Tensile stresses in the coating propagate stress corrosion cracks. Also the adhesion between coating and substrate is mainly influenced by the residual stresses in the interface. The final residual stress situation of thermally coated components is superimposed by several individual stress mechanisms. The reasons for residual stresses during manufacturing are temperature gradients in material combinations with originally incompatible thermophysical properties as well as mechanical loads which occur during substrate preprocessing, thermal spraying and finally composite postprocessing. With the hole drilling method, more exact the circular micromilling method, the residual stresses in components are determined partially destructive. In several drilling and milling processes a circular microhole is brought step by step (5 - 10 pm) into the component surface. The residual stresses are relieved due to this material removal, deform the surface around the hole and are measured as relaxed strains E at the surface by means of high precision strain gauges. Out of this strains the in plane stresses ox, oy are incrementally determined by Hooke's law [4]. The residual stresses of APS and HVOF deposited Ti02 cylinder bore coatings, sprayed with different cooling rates are measured up to a drilling depth of 0.4 mm, see figure 3. The coating thickness is constant 160 pm. Due to the low particle velocities (- 50 - 150 d s ) and the large degree of totally melted particles the occurring residual stresses of APS deposited coatings are mainly influenced by the temporary temperature distribution in the layer composite and the thermophysical material properties. Because of the high particle velocities and the low degree of particle fusion during HVOF spraying, here the thermal stresses are superimposed by additional compressive stresses induced by the high particle impact energy. It can be noted that the kinetic particle energy can be controlled via the total gas flow rate as a result of the oxygen / fuel ratio. For the APS process tensile stresses in the coating and the interface can be measured with increasing cooling rates by C 0 2 and air jets. For decreasing cooling rates, only air cooling resp. without cooling, compressive stresses can be detected. Decreasing cooling rates cause a higher uniform temperature level in the composite. Because of the mismatch in the thermal expansion coefficients between coating q and substrate as higher compressive thermal stresses arise in coating and interface. For the HVOF deposited coatings compressive stresses in coating and interface can be measured. Due to the larger
20 1
gas flow rates of the propane process (- factor 1.6) compared to the acetylene process and the less degree of particle fusion higher compressive stresses can be measured. In general it can be said that with increasing compressive stresses in coating and interface the hardness and the bond strength of the layer coatings increase.
-g
200,o
100.0 0,o
MECHANICAL COATING CHARACTERIZATION To assess coating quality, manufacturing reproducibility and the operation behaviour of the layer composites metallographic examinations were performed for all samples. By image analysis the coating porosity, expressed by the relative pore volume content Vp [%], is defined using cross section evaluation of metallografic samples. The hardness HV0.05 is measured with an automated universal hardness indenter equipment (F = 500 mN). Table 2 shows the measured porosity and hardness characteristics for both coating processes.
f -100.0
f
Coating
I
-200.0
I
I
HV0.05 HVOF
1
8.APS; w i l b u l mofing
a
g 300.0 dM n -w>"
I
Vp[Oh] V P ~ O ] HV0.05 APS APS HVOF
+HVOF.
pmpivm: CO2 moling
*
0,05
0
0.1
0.15
0.2
0.25
0.3
0.35
0.4
ddlllng depth [mm]
Fig. 3 Measured residual stresses in Ti02 coated cylinder bores, variation of the thermal spraying processes (APS, HVOF) The investigations of the bond strength are performed on a universal testing machine. A steel tension rod is glued on the coating surface. As soon as delamination of the coating occurs, the tension load is measured and the bond strength is calculated. Figure 4 shows the measured bonding strength results for the Ti02-A1Si9Culayer composite manufactured with different spray processes and cooling parameters.
CrB/NiCr-75/25
I
18
I
13
-
3-4
Mo/NiCrBSi-30/70
6-8
5
Mo/Cr-Steel-70/30
I 6 - 10 I
4-5
Mo
I
1500
I 1300- 1600 I 1000 - 1200
700 - 1100
1
400 - 500
1000 - 1200
I
700 - 1100
I
Table 2 Measured coating porosity and coating hardness
60
WEAR AND FRICTION
,---.I_.-,.
CQlidr ding
(pacum a i d '
mdlng
Fig. 4 Measured bonding strength of TiOz-AlSi9Cu layer composites In general it can be said that with increasing compressive stresses in the interface also the bond strength increases [ 5 ] . For the HVOF propane oxygen sprayed Ti02 system no bond strength can be measured because of the lower yield strength of the used glue. It just can be mentioned that the bond strength of the HVOF propane oxygen sprayed Ti02 coatings compared to the APS Ti02 coatings is by a factor of at least 1.8 higher.
202
The aim of the material screening is the detection of tribological systems with low friction coefficients as well as low wear rates on coating and counterpart (piston ring). Low friction coefficients are required to reduce the energy dissipation during operation. A decrease of the wear rate during operation increases lifetime and quality of the tribological system and reduces the pollution of the lubricating oil with particles. To get an overview about the wear and friction behaviour of the different coating systems, dry running oscillating pin on disc tests are performed. The coatings are finished to a surface roughness of R, 0.05 pm. As a counterpart A1203 balls with a diameter of 5 mm are used. The number of oscillating strokes is up to 50.000, sliding velocity 70 m d s , length of strokes 5 mm with a imposed measurement load of 10 N. For all thermally sprayed coatings friction coefficients p below 0.9 can be measured, compare figure 5. Very low friction coefficients between 0.1 an 0.2 can be measured for some Ti02 APS coatings. Additionally a very interesting correlation between hardness and friction coefficient can be detected for the Ti02 coating
-
systems. The coating systems with the low hardness HV0.05 values and friction coefficients are sprayed with an increased hydrogen percentage by optimized C o t and air jet cooling. Therefor it can be assumed that nonstoichiometric TinOzn., phases are responsible for this excellent friction behaviour. The wear investigations show similar results. .........................
:.........................
j A APSCr3C21NiCr
i AAPSCR03
REFERENCES [I] Hinz R., Schwaderlapp M.; "Potential zur Massenreduktion am Beispiel eines 4-Zylinder-Reihenmotors"; Leichtbau im Antriebsstrang (1 996), Expert Verlag, ISBN 3-8169-1336-9, pp. 162 - 173
[2] Woydt M., Skopp A., Dorfel I., Witke K.;"Wear Engineering Oxides / Antiwear Oxides"; Tribology Transactions, Volume 42 (1 999), 1, pp. 2 1 - 3 1
0 HVOF-Gr3CZNtCr
OHVOF-CQ03 0 HVOF-Ti02
[3] Woydt M.; "Materials-based concepts for an oil-free engine"; New Directions in Tribology (1997), pp. 459 -
468 ......................................................................
.....................
j ........................
:........................
4
0,1
0
1000
2000
3000
4000
coating hardness HV0.05
Fig. 5
II
Measured average friction coefficient p vs. coating hardness HV0.05
[4] Buchmann M., Gadow R., Tabellion J.; "Experimental and numerical residual stress analysis of coated CORIPOSites"; E-MRS Spring Meeting Strasbourg 1999, in print in Mat.Sci.Eng. [5] Buchmann M., Friedrich C., Gadow R. ; "Residual stress characterization of thermal barrier coatings - comparison of thermally sprayed, EB-PVD and CVD based coatings"; 24th Annual Cocoa Beach Conference & Exposition 2000, in print
CONCLUSION Intentions to reduce manufacturing cost, fuel consumption and waste emissions in the automobile industry result in increasing light weight design and applications. The poor tribological operation behaviour of light metal surfaces can be improved by protective coating systems. Beside galvanic coatings and reinforced materials, thermally sprayed coatings offer a broad material variety as well as flexible and cost effective manufacturing processes for refined crankcases. By means of APS and HVOF spraying wear resistant inside coatings are deposited on aluminum tubes and cylinder bores of combustion engines. During the material screening different ceramic, cermet and metallurgical coating systems were deposited and investigated. The thermophysical properties cx and cp were defined. Residual stress measurements have shown that the stresses in thermally coated layer composites strongly depend on the temperature level and history during the coating process as well as on the particle impact velocity. Wear investigations show friction coefficients lower than 0.9 for all coatings under dry conditions. In the case of plasma sprayed Ti02 coatings excellent friction coefficients between 0.1 and 0.2 can be measured under dry running conditions. The investigations during the material screening and first bench tests reflect positive results for most of the coating systems. A detailed coating selection can be done after further test runs of combustion bench tests in stationary reciprocating engines. It can be assumed that thermally sprayed coatings will play an important role as protective coatings for light metal engine applications and will continue to become more and more important in the near future.
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TRIBOLOGICAL BEHAVIOR OF SILICON NITRIDEBTEEL CONTACTS UNDER LUBRICATED CONDITIONS Hyung K. Yoon, Seung-Kun Lee,Frank A. Kelley, Michael J. Readey Advanced Materials Technology Caterpillar Inc., USA
ABSTRACT The tribological properties of some silicon nitrides for valve train components were evaluated under lubricated conditions. The silicon nitrides tested are SN235, SN237 (Kyocera, Japan) and AS800 (AlliedSignal Inc., USA). For comparison purposes, limited tests were also conducted with a hardened 52100 steel (60 Rc). These materials were tested against SAE 4620 steel in a block-on-ring tester with a polyalphaolefin (PAO) oil. The results show that AS800 gives lower friction and wear compared to other silicon nitrides at higher loading conditions. All silicon nitrides exhibited much higher wear resistance compared to a hardened 52100 steel. The effects of contact pressure (P) and sliding speed (V) on friction and wear characteristics were also examined. The friction coefficient and wear increase as the contact pressure increases. Wear also increased with sliding speed, however, the coefficient of friction remains about the same as the sliding speed increases. Silicon nitrides exhibited mild polishing wear in most conditions. However, a transition in the wear mode from polishing to abrasive wear seems to occur when P V is higher than 763 MPa*m/s. For all silicon nitrides tested, the friction transition was also observed when P V is higher than 763 MPa*m/s. Tribo-chemical films formed on the sliding surfaces seem to play an important role in the observed friction transition behavior.
INTRODUCTION Valve train components in heavy-duty diesel engines operate at high temperatures, high peak cylinder pressures and severe corrosive environments. These severe conditions have resulted in various tribo-related failures in valve train components in many diesel engines. Therefore, there has been increasing interest in developing better valve train materials that are more wear resistant to enhance the reliability and performance of these components. Advanced ceramics and emerging intermetallic materials are promising materials candidates for valve train components since they are highly corrosion and oxidation resistant, and possess high strength and hardness at elevated temperatures. In recent years, advanced ceramics have been widely used in sliding
components such as journal bearings, cylinder liners, piston rings and mechanical seals [ 1-31. Fine ceramics have also been successfully used as rolling bearings in machine tools and sliding bearings in water pumps 141. In the past two decades, silicon nitride-based ceramics have been targeted for valve train components, and some commercial successes have been reported in both automotive and diesel valve trains. They are currently being used in some diesel engine valves, valve guides, roller followers and tappet shims [5]. It was reported [6] that a significant wear reduction was achieved with a Si3N4 rocker-arm pad coupled with a cast iron cam. Silicon nitride ceramics are continually being developed to improve their physical and mechanical properties by controlling composition and microstructure [7,8]. In order to establish a more complete database for the performance characteristics of these materials in a sliding contact, the tribological properties for these ceramics need to be evaluated. The objectives of this work are to investigate the tribological properties of some silicon nitrideshteel sliding contacts under lubricated conditions, and to identify important properties to improve wear resistance of these contacts. In addition, the effects of contact pressure and sliding speed on friction and wear transitions are also examined.
EXPERIMENTAL SETUP Geometry of Contact This work mainly focuses on a cam roller/follower sliding contact in valve train components. To simulate the contact geometry of this critical contact, the friction and wear tests were conducted in a block-on-ring test rig. A flat block specimen (silicon nitride) is loaded against a ring specimen (4620 steel), which rotates at a given speed for a given number of revolutions. The ring specimen is partially submerged in the lubricant. A schematic of block-on-ring test geometry is given in Fig. 1. Materials Tested Three commercially available silicon nitrides were tested in this study. For comparison purpose, limited tests were also conducted with a hardened 52100 steel (60 Rc). These materials were tested against 4620 steel ring specimens, which have the hardness of 58-62 Rc and an average surface roughness of 0.23 pm Ra. Some
205
physical and mechanical properties for silicon nitrides
tested are given in Table 1.
Materials
Suppliers
Second Phase
Grain size (Pm)
Strength (MPa)
SN235 SN237 AS800
Kyocera Kyocera AlliedSignal
Al, Y Al, Y La
Fine (0.8) Finer (0.5) Coarse (3.2)
780 k 18 835 k 56 685 k 21
Toughness Hardness (MPa~rn”~) (GPa)
6.3 5.8 7.2
14.7 14.5 14.2
Thermal Cond. (W/m*K) 27.7
Young’s Modulus (GPa) 305 305 300
72.5
Load
RESULTS Block (ceramic)
Effect of Contact Pressure Friction results The friction data for silicon nitridedsteel contacts as a function of contact pressures are given in Fig. 2. For each condition, two tests were conducted. The results show that, for given conditions, the average coefficient of friction increases as the contact pressure increases for all silicon nitrides tested. AS800 that has a relatively coarse microstructure gives slightly lower friction coefficient at higher loading conditions (7 18 and 829 MPa).
Fig. 1 - Schematic of the contact geometry (block-on-ring)
Test Procedure All tests were conducted under submerged lubrication conditions for the test duration of 30 minutes. The coefficient of friction was monitored and recorded continuously throughout the test by a computer-based data acquisition system. The wear volume on the block specimens was calculated by measuring the wear scar width using an optical microscope. The worn surfaces of both ceramic block and steel ring specimens were also examined using a surface profilometer, optical interferometer and SEM. To examine the effects of contact pressure and sliding speed on the tribological characteristics of silicon nitrides, various loads and sliding speeds were used. The test conditions used are summarized in Table 2. Table 2 - Testing conditions 414,586,7 18,829 Contact Pressures (MPa) Sliding Speeds ( d s ) I 0.31,0.61,0.92, 1.23
I
Lubrication Condition
206
300
400
500
600
700
Initial Hatziaa C o n m Ressure.
800
P (MPa)
Fig. 2 - Friction data for silicon nitrides/steel contacts as a function of contact pressure The typical coefficient of friction plots obtained at various contact pressures are given in Fig. 3. In most conditions, the friction coefficient tends to decrease and then increase slightly with time. It should be noted that, when a contact pressure of 829 MPa was used, the friction transitions were obtained with all silicon nitrides tested. The initial increase in the friction data can be due to the local breakdown of the lubricant, and the following decrease is attributed to the formation of tribo-chemical films on the sliding interface. A tribofilm formed on the worn surfaces of silicon nitrides is given in Fig. 4. The SEM micrographs show that the area of the worn surface, which is covered by these films, increases as the contact pressure increases. Based on the EDS analysis, these films are a combination of silicon and iron oxides. It is known that a tribofilm formed on the worn surface remarkably alter the tribological characteristics of ceramic-steel
900
pairs. The study of Fischer and Mullins [9] has shown that a film of chemical resultant, such as amorphous SiOl or hydrated Si02 will be formed on the wearing surface of Si3N4.Consequently, the friction coefficient is reduced to a low value. It is believed that the tribofilms formed on the silicon nitride tested similarly affect the friction behavior of this material.
0.30
-P
= 414 M P n P = 586 MPa P = 718 MPa
0.25 C
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2
0.15
g
0.10
.-5
8
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20
30
P = 586 M P a
0.25
P = 718 M P s
.-B
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3 0.20
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.-tl 0
-
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E
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'
I
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'
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1
'
'
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5
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"
'
I
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15
1
'
"
I
'
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30
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T i m e . t (min) 0.30 c
0.2.5
._ .8 0 . 2 0
-
Fig. 4 - Tribo-chemical films formed on the worn surfaces of silicon nitrides (a) P = 414 MPa, (b) P = 829 MPa
Wear results The wear data for silicon nitrideskteel contacts as a function of contact pressures are given in Fig. 5. As previously indicated, the volume worn on a block specimen was calculated based on the wear scar width measured using an optical microscope. As expected, wear increases as the contact pressure increases for all silicon nitrides tested. For the exception of 414 MPa, the lowest wear was obtained with AS800. This is attributed to its higher thermal conductivity as shown in Table 1. Due to its higher thermal conductivity, the surface temperature of AS800 can be lower, resulting in better wear resistance. The thermal conductivity for SN237 is not available, however, it is believed that it is not significantly different from that of SN235.
i
;0.1.5
.-0 .-u
'=
8
0.10
0.0s 000
1
0
' '
'
' ' 5
'
' '
'
10
'
' ' ' 1 ' *
'
IS
'
I ' ' ' ' I '
20
'
'
25
30
Time, t ( m i n )
Fig. 3 - Typical coefficient of friction plots for silicon nitrideslsteel contacts obtained at various contact pressures
. .-2'. . ..
300
400
so0
600
700
.:.
.. .
8M)
900
Initial Henzian Contact Pressure, P (MPa)
Fig. 5 -Wear data for silicon nitrides as a function of contact pressure Lubricant: SHF-41 (PAO) The worn surfaces of both ceramic blocks and the counterface steel rings were examined using a surface profilometer, an optical interferometer and SEM. Surface roughness values of ceramic blocks before and after testing are given in Fig. 6 .
207
300
500
400
600
700
800
900
Surface roughness values of the counterface steel rings are also given in Fig. 8. Similar to silicon nitrides, surface roughness of the steel ring decreased when the contact pressure is lower than 718 MPa. The worn surface shows that the initial machining marks are removed during the rubbing process, and the tribofilms are formed on the steel ring surface. However, except for SN237, surface roughness values of the steel counterface sharply increase when the contact pressure of 829 MPa was used. Several local microgrooves were also seen on the worn surface of the counterface steel ring specimens at this condition. This indicates that a transition in the wear mode is also taking place on the counterface steel ring.
Initial Henzian Contact Pressure. P (MPa)
Fig. 6 - Change of surface roughness of silicon nitrides as a function of contact pressure The initial surface roughness of silicon nitrides was about the same. It was in the range of 0.26-0.35 pm Ra. However, except for SN235 and AS800 at 829 MPa, surface roughness after testing decreased one order of magnitude, indicating mild polishing wear. This is mainly caused by plastic deformation and microfractures of the asperities. However, surface roughness sharply increases for AS800 and SN235 when a contact pressure of 829 MPa was used. It was found that microgrooves were locally formed on the worn surface as shown in Fig. 7. This indicates that a wear transition from polishing to abrasive wear seems to occur at a given condition. ~ w . 3 ~m ’ a*
300
400
500
600
700
800
900
lnitial Henzian Contact Pressure. P (MPa)
Fig. 8 - Change of surface roughness of the counterface steel ring as a function of contact pressure Q* 24
22 2 1B 94 14
12 1
08
Pocehirng
0 0
0 2
.. I
0 4
0 6 Scm Lcngth. X
Ob 04 02
0
Abrasive Wear
0 8
1 0
1 2
(mm)
Fig. 7 - 3D surface map and surface profile of the worn surface of silicon nitride P = 829 MPa, V = 0.92 d s
208
Effect of Sliding Speed Friction results The friction data as a function of sliding speeds are given in Fig. 9. As previously indicated, the limited friction and wear data were also obtained with a hardened 52100 steel (60 Rc) for comparison purposes. For silicon nitrides, the coefficient of friction seems to remain the same as sliding speed increases. The data show that, for the exception of 0.31 d s , AS800 gives slightly lower. friction coefficient among silicon nitrides tested. For 52100 steel, the coefficient of friction at 0.31 d s is lower than those obtained with silicon nitrides. However, the friction and wear data for this material could not be obtained with higher sliding speeds since it scuffed immediately when the sliding speed of 0.62 m/s was used. The coefficient of friction plots as a function of time at various sliding speeds are given in Fig. 10. Similar to the friction behavior shown in Fig. 3, the friction transitions were obtained during the test when a sliding speed of 1.23 d s was used. The friction increase during the test at 1.23 d s is also attributed to the local breakdown of lubricant films at the sliding interface.
Wear results The wear data for silicon nitrides and a hardened 52100 steel as a function of sliding speed are given in Fig. 11. For the range of sliding speeds used, wear increases as the sliding speed increases. It is also seen that the lowest wear was obtained with AS800 for all conditions used. As previously noted, this can be due to its higher thermal conductivity compared to other silicon nitrides tested. However, the ranking of wear between SN235 and SN237 is inconclusive. At the sliding speeds of 0.61 and 1.23 d s , SN237 is clearly better than SN235. However, SN235 is slightly better than SN237 at the sliding speeds of 0.305 and 0.92 d s . Note that, at a sliding speed of 0.31 d s , 30-40% more wear was obtained with a hardened 52100 steel compared to silicon nitrides tested even though its friction coefficient is lower than those of silicon nitrides as shown in Fig. 9. As previously indicated, scuffing occurred immediately with a hardened 52100 steel when a sliding speed of 0.61 mls was used. Silicon nitrides, however, exhibited mild wear even with much higher sliding speeds, indicating superior scuffing resistance of these materials compared to a hardened 52100 steel.
.................................................................................
t
0.00
"
'
~
0.2
'
"
"
'
'
"
~
'
"
'
~
I .o
0.8
0.6
0.4
'
"
~
I .2
I .4
Sliding Speed. V ( d s )
Fig. 9 - Friction data for silicon nitrides and a hardened 52100 steel as a function of sliding speed Lubricant: SHF-41 (PAO)
0.30
,
0.25
, . . . . . . . .............,
,
I
,
I
,
,
V = 0.31 m/s V = 0.61 m/s V = 0 9 1 m/s
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......................
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Time, t (min)
0.25
.....
-I3
......
000
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00
02
~
~
04
06
08
'
10
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12
14
Sliding Speed, V ( d s )
Fig. 1 1 - Wear data for silicon nitrides and 52100 steel as a function of sliding speed Lubricant: SHF-41 (PAO)
T i m e , t (min)
-V -V
=0.31 mls =0.61 mls
-$
V = 0 . 9 2 m/s
0 0s
000 0
5
10
15
20
2s
T i m e , t (min)
Fig. 10 - Typical coefficient of friction plots for silicon nitrides obtained at various sliding speeds
30
Surface roughness data for silicon nitrides before and after testing, obtained at various sliding speeds, are given in Fig. 12. Data show that surface roughness of the worn surface decreases as the sliding speed increases up to 0.92 d s . Mild polishing wear was obtained for all silicon nitrides at those conditions. However, surface roughness of the worn surface sharply increases as the sliding speed of 1.23 d s was used. Again, a wear transition from polishing to abrasive wear was observed at this sliding speed. Surface roughness values of the counterface steel rings before and after testing are also given in Fig. 13. Note that surface roughness of the worn surface decreased if the sliding speed is less than 0.92 d s . However, surface roughness sharply increases when the sliding speed of 1.23 d s is used. Again, this is the indication of a wear transition of steel ring at this condition.
209
(4) For ail silicon nitrides tested, wear increases as the
contact pressure or sliding speed increases. (5) For given conditions, AS800 generally gives better wear resistance compared to other silicon nitrides tested. It is attributed to its higher thermal conductivity. (6) In most conditions, the dominant wear mode for silicon nitrides is mild polishing wear, which is mainly caused by plastic deformation and microfractures of the asperities. However, a transition in wear mode from polishing to abrasive wear seems to occur when PV is higher than 763 MPa*m/s. This transition might be related to the friction transition observed at this condition. (7) Silicon nitride ceramics show a superior wear resistance compared to a hardened 52100 steel (60 Rc) when they are tested against 4620 steel.
z v)
Fig. 12 - Change of surface roughness of silicon nitrides as a function of sliding speed
ACKNOWLEDGMENT This research was sponsored by the Department of Energy in USA, under Contract DE-FC05-970R22579.
REFERENCES
0.2
0.4
0.6
0.8
1.o
1.2
1.4
Sliding Speed,V ( d s )
Fig. 13 - Change of surface roughness of the counterface steel ring as a function of sliding speed
CONCLUSIONS Some silicon nitride ceramics for valve train components were tribologically evaluated. The silicon nitrides tested are SN235, SN237 (Kyocera Inc., Japan) and AS800 (AlliedSignal Inc., USA). For comparison purposes, limited tests were also conducted with a hardened 52100 steel. These materials were tested against SAE 4620 steel in a block-on-ring tester under lubricated conditions. The effects of contact pressure and sliding speed on friction and wear characteristics were examined. The results of this study are summarized as follows: ( 1 ) For a given sliding speed and the range of the
contact pressures used, the average coefficient of friction increases as the contact pressure increases. However, it remains about the same as the sliding speed increases with a given contact pressure. (2) AS800 ceramic gives slightly lower friction coefficient at higher load and speed conditions. (3) The friction transition was observed for all silicon nitrides tested when PV is higher than 763 MPa*m/s.
210
S. Ying, Advanced Ceramics and Applications of Ceramics", Beijing Press of Science and Technology, Beijing, (1990), 167170. S. Asanabe, Applications of Ceramics for Tribological Components, Tribological International, 20 (6), ( I 987). 355-364. M. Woydt and J. Schwenzien, Dry and WaterLubricated Sliprolling of Si3N4- and SiCBased Ceramics, Tribological International, 26 (3), (1993), 165-173. G. W. Stachowiak, G. B. Stachowiak, and A. W. Batchelor, Metallic Film Transfer During Metal-Ceramic Unlubricated Sliding, Wear, 132, (1989), 361-381. B. Dumont, P. J. Blau, and G. M. Crosbie, Reciprocating Friction and Wear of Two Silicon Nitride-Based Ceramics Against Type 316 Stainless Steel, Wear, 238, (2000), 93109. M. Kano and I. Tanimoto, Wear Resistance Properties of Ceramic Rocker Arm Pads, Wear, 145, (1991), 153-165. C-W. Li and J. Yamanis, Super-Tough Silicon Nitride with R-Curve Behavior, Ceramic Engineering and Science Proceedings, 10 (78), (1989), 632-645. C-W. Li, D-J. Lee, and S-C. Lui, R-Curve Behavior and Strength of In-Situ Reinforced with Different Silicon Nitride Microstructures, J. Am. Ceram. SOC., 75 (71, (1 992), 1777-1785. T. E. Fischer and W. M. Mullins, Chemical Aspects of Ceramic Tribology, J. Phys. Chem., 96, (1 992), 5690-5701.
OPTIMIZATION OF THE BRAZILIAN DISC TEST FOR CERAMIC MATERIALS A. Borger", P. Supancic and R. Danzer Department of Structural and Functional Ceramics, University of Leoben and Materials Center Leoben, A-8700 Leoben, Austria
ABSTRACT The Brazilian disc test is widely used for strength measurements of brittle low strength materials. It is also considered to be a simple and low cost method for the determination of the tensile strength of advanced ceramics. Because of the very high compressive stress components in the region of the loading point an undesired failure mode may occur. In order to prevent this mode a change of the test geometry is proposed and the stresses and the effective volume of the modified experimental set-up are evaluated.
specimen. Munz and Fett [6] presented analytical solutions for the stress distribution in BDT specimens considered as a 2D object. For a=O (point load) it holds:
F
INTRODUCTION The Brazilian Disc Test (BDT) is a method for testing the tensile strength of cylindrical specimen [ 1-41. A diametrical loading of a disc specimen via two parallel flat platens leads to a bi-axial stress state with tensile and compressive stress components in the specimen. Since the tensile stress maximum is in the centre of the specimen, it is generally assumed that failure occurs when the applied tensile stress components exceed the materials strength. In the region of the loading point very high compressive stress components exist [3]. Their magnitude is several times higher than that of the tensile stresses in the centre of the specimen. An undesired failure mode caused by these compressive stresses is sometimes observed. This occurs especially when testing advanced ceramics even if these materials have a much higher compressive then tensile strength. But since this test is applied to measure tensile strength this mode of failure has to be suppressed. In this work the geometry of the loading device has been modified by introducing a curved support in order to reduce the compressive stresses in the area of the loading contact and to provoke tensile failing from the centre of the sample. The stress distribution in the sample and in the platens is studied by a contactmechanical Finite Element (FE) analysis. Also the dependence of the maximum tensile stress in the sample on the curvature of the support is studied [ 5 ] .
STRESS DISTRIBUTION IN STANDARD TEST ASSEMBLY
THE
Fig. 1 Specimen geometry and loading status.
(2)
0,. =--
- (I + y * ) 2 x * .?
r2
]
(3)
with: =slR; y'=ylR , 5 r 2 =(q / R ) ~ ;r2.2 = ( r , / ~ ) ' .
X*
(4)
These equations result in an infinite compressive stress at the loading point. The maximum tensile stress occurs
in the centre of the disc. Therefore, using the maximum tensile stress failure criterion the disc would fail from their centre along the y-axis where the stress components are: (5)
Figure 1 shows the geometry and the loaded area of the specimen. As a result of loading an inhomogeneous stress distribution exists in the
211
Munz and Fett [6] also reported a modification of equs. 1-3 for a finite loading length a. The stress distribution in the disc along the y-axis (x = 0) is:
hardened steel support) at the contact area (Hertzian contact cracks). To be able to use this testing method for advanced ceramics a modification of the support geometry in order to increase the contact area and to decrease the compressive stresses has been carried out. FE calculations were done to determine an advantageous curvature of the support.
NUMERICAL SOLUTION In fig. 2 the ratio, R, , of the amplitudes of the maximum compressive stress components in the loading area to the maximum tensile stress components in the centre of the specimen is plotted against the ratio of the half contact length a to the disc diameter R. loo,
I
I
I
To study the effect of curvature of the support and of different support and test specimen materials a parametric FE study has been carried out using a 2D and a 3D contact mechanical model. The software used was ANSYS (Vers. 5.5). In the first approach a 2D model of a quarter of the test assembly has been modelled. To test the accuracy of the model in a first step a flat support was verified by comparison with the standard analytical solutions (fig. 4.a). In a second step the influence of a curved support on the stress distribution was evaluated using the curvature as a modelling parameter (fig. 4.b).
Fig. 2: Ratio of the amplitudes of the max. compressive stress at the support area to the max. tensile stress in the centre of the disc as a function of a/R.
Fig. 3: Valid failure mode (left) gained on the modified BDT and invalid failure (right) gained on flat support on electro-ceramic samples. A small contact length as it is the case for ceramic materials on a hardened steel jig the maximum compressive stress amplitude is much higher than the maximum tensile stress. Therefore, a different failure mode caused by these compressive stresses is often observed, (triple cleft failure) as shown in fig. 3 right. This happens also when testing advanced ceramics even if these materials have a much higher compressive then tensile strength. FE calculations have shown that compressive stresses can be up to 30 times higher than the tensile stresses [5]. Experiments on electro-ceramic samples show that nearly 100% of the samples fail in the triple cleft failure mode leading to invalid experiments. It was also observed that these large compressive stresses lead to crack initiation in the support (a
212
Fig. 4: Quarter 2D-model of the brazilian disc test with a) planar and b) curved support (the support radius is here 5% bigger than the disc radius). The results of the 2D model of the flat support show very good agreement with those of the analytical solution. The stress distribution along the x and y-axis fits nearly perfectly along the hole diameter of the disc. Only near the contact point between disc and support a small deviation of the compressive stresses achieved in the FE solution and in the analytical solution was observed. Also the stress distribution in the support was determined, showing large tensile stresses in the surface next to the contact area. This is exactly the area where
cracks were observed when the ceramic samples were tested on the hardened steel support.
EFFECTS OF A CURVED SUPPORT
r
l
0 10.0
102
+lo
curvature radius should be very similar to the disc radius. But if the curvature radius is to close to the disc radius small variations of the disc radius have a large effect on the introduced tensile stresses in the disc which is also undesirably (if for instance the radius is only 1% bigger than the disc radius the maximum tensile stress decreases to 66% of the stress value of the sample loaded with a flat support). As a compromise a support radius which is 5% bigger than the disc radius is used for further studies. This support geometry was machined and also subjected to an intense numerical analysis of stress distributions for different materials. All hrther reported work refers to this support radius. a)
10,4 10.6 108 11,O 11.2 11.4 support radius [mn]
0 11.6
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[mm] I
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010.2
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applied bad = 10 kN
b) 4001 10,O
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11.6
Fig. 5.a) Decrease of the maximum tensile stress o,in the centre of the disc and b) ratio R, of amplitudes of maximum compressive and maximum tensile stresses in the disc as a function of the support radius. The used material and geometric parameters can be found in the text. The curved support was introduced in order to reduce the compressive stresses on the disc and also to reduce the tensile stresses on the surface of the support that lead to crack initiation and failure of the support. This is achieved through the increased contact area between disc and support. It was observed that the stress ratio R, decreases if the curvature approaches the radius of the disc. Also a strong correlation between the stress ratio R, and the applied force was observed which again is caused by an increase of the contact area through the elastic deformation of the specimen and support. As an example this ratio R, is plotted in fig. 5.a over the support radius for several typical loads. The calculations have been made for a disc radius R = 10 mm, a disc thickness t = 3 mm, Young's moduli of specimen Especimr,,=100 GPa and support Esupporf= 200 GPa respectively and for both materials a Poisson ratio v = 0.3 is used. Compared to the testing assembly with flat supports a lower maximum tensile stress under the same load was observed (see fig. 5.b). Both above discussed effects have to be taken into account for the determination of an appropriate support curvature. For the reduction of the contact stresses the
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.
-1 0
Fig. 6: Stress distribution along a) horizontal (x) and b) vertical (y) axis of the disc for flat and curved supports. The support radius is 5% bigger then the disc radius. Same material parameters as in fig. 5 . The numerical solutions show that the stresses along the x-axis are not significantly affected by the curvature of the support except the decrease of the maximum tensile stress in the disc centre (figure 6.a). The calculations have been made for the same material parameters as before. On the other hand a comparison along the vertical axis shows that the analytical solutions totally fail to describe these stresses (fig. 6.b). Using the analytical solutions for a finite loading length 2a leads to a slightly better description of the stress distribution but still fails to predict stresses even nearly accurate for more than 70% of the radius of the disc. This is mainly caused by the fact that the analytical equations assume a constant load over the contact length 2a which clearly is not true in this case. As a result the evaluation of these experiments can not be done using the standard analytical solution for the maximum tensile stress in the centre of the disc.
213
has some influence on the effective loaded volume of the specimen.
EVALUATION OF THE MAXIMUM TENSILE STRESS IN THE DISC As already mentioned the maximum tensile stress in the centre of the disc is influenced by the experimental conditions. It can be expressed by modifying equ. 5 with a correction factor C:
which depends on the elastic properties of the testing and support materials, the radii of support and specimen and the applied load. It is plotted over the applied load in fig. 8.a for a disc radius of R = 10 mm. The support curvature is the parameter. The same material parameters as in the previous figures are used. It is obvious that the factor C reaches one if the support radius approaches infinite (flat support). With increasing applied load, F, the elastic deformation of specimen and support and also the contact area increase, causing a decrease of the maximum tensile stress in the disc centre compared to the flat support situation. In principle the factor C has to be evaluated for every geometry, for every material combination and every applied load. In Fig. 8.b the factor C is plotted over the maximum tensile stress in the specimen (defined by equ.10). Shown are results for several different specimen radii and ratios of the elastic moduli of testing and support material. For both materials the Poisson ratio is assumed to be 0.3. It can be seen that for a constant ratio of moduli the result does not depend on the disc radius. Analytical fit equations for the factor C for three different ratios of moduli are given in Tab. 1. The stress a, is given by equ. 10. These equations have to be solved iteratively.
I Fig. 7: Stress distribution in the disc (R = 10 mm) a) and b) on a flat and c) on a curved support (support radius 5% bigger than the disc radius). Plotted are q, ox,and q,respectively. Same parameters as before. Fig. 7 shows results of some of the performed FEcalculations. In figs. 7.a and b the flat support situation is modelled, showing the first principal stresses, oI, and the stress ox, respectively. Fig 7.c shows the first principle stresses, q,on a curved support. The curved support leads to a decrease of the maximum tensile stress in the centre of the disc and also to a decrease of the volume fraction that is under tensile stresses. This
214
I
Table. 1: Approximate equations for C for different combinations of Youngs moduli of specimen and support materials for a support radius 5% bigger than the disc radius. C EstIec1men 1 ESuLlDort I 112 1 -4.953.104-q +2.173.107.ax2 1 1 - 5.890-104.~,+2.699.10-’0,~ 312 I 1 -4.954.10~4.~+2.174~10-7~,2 I A 3D model of the test assembly shows no significant changes on the stress distribution except a negligible decrease of the tensile stresses near top and base surface of the disc.
EXPERIMENTAL INVESTIGATIONS For the fracture experiments disc shaped electroceramic samples (E = 100 GPa, v = 0.3) were used. The samples were tested in the as sintered condition with a diameter of 20 mm (+O,O 1 mm I -0,02 mm) and a height of 3,26 mm (k 0,059). Using the standard BDT support
geometry only ten samples were tested. All these samples showed triple cleft failure, fig. 3 (right). Also the support showed significant cracks on the surface near the contact area between the specimen and the support. Therefore all further tests were performed with the modified testing assembly. All 30 tested samples showed the tensile failure mode as seen in fig. 3 (left). As a reference testing method a miniaturised 4-point bending testing method described in [7] was used. For these tests mini-bending bars (30 specimens) with a geometry 1,5x2x15 mm were machined from the discs.
maximum tensile stress in the sample which is right below the loading point. Fractography proves that fracture starts - as in the case of bending specimens - at internal defects (pores). Table. 2: Results of the strength tests. The description can be found in the text.
confid. int. confid. int.
(103 - 108)
(107- 111)
(11 - 19)
(15-24)
1
1
crack origin
Fig. 9: Macroimage of a typical BDT fracture surface of a specimen tested using the modified support.
CONCLUSIONS
0
axin the centre of the disc [MPa]
Fig. 8:Correction factor C a) over the applied load for a disc with radius R = 10 mm and for several different support radii and b) over the max. tensile stress 0; for specimens with different radii and different ratios of Young's moduli of tested material to support material. The support radius is 5% bigger than the disc radius.
0
0
The test results are reported in table 2. For the BDT tests the maximum principal stress (in ceramics tensile failure is in general triggered by this stress component [6, 81) and for the bending tests the maximum outer fibre stress are given respectively. For each data set the characteristic strength, its 95 % confidence intervall, the Weibull modulus and its 95 % confidence interval1 are given (data evaluation with Weibull theory is common practice when testing ceramics [6, 8, 91). Within the scatter of the data no significant difference between the strength data can be observed. A comment on the need to apply Weibull theory in order to describe the Brazilian Disc testing of ceramic materials is given in the appendix. Fig. 9 shows a macroscopic image of the fracture surface of a BDT sample tested on a curved support. Fracture clearly starts from within the area of the
Brazilian disc testing with a flat support can result in undesired failure starting from the loading point, especially if testing high strength materials. A new design of the support reduces the compressive stresses in the support area and provokes specimen failure from the tensile loaded disc centre. In the newly designed jig the tensile stresses in the centre of the specimen are smaller than in the standard configuration if tests are performed at the same load. The reduction depends on geometrical conditions of the support, the applied load and the ratio of the Youngs moduli of specimen material and support material, respectively. For a support with a curvature 5% bigger than the specimen radius for a wide range of possible material parameters the tensile stresses have been calculated. Tests performed on an electro-ceramic (bariumtitanate) material give the same strength values when tested in 4-point bending and in the Brazilian disc test configuration.
APPENDIX: THE EFFECTIVE VOLUME The fracture probability of ceramics is commonly described by the Weibull distribution function, F(o,.), which for a homogenous tensile loading is [6, 91:
with oi, and m as parameters of the distribution. V, is a scaling volume. The Weibull hnction predicts a dependence of the probability of failure on the the 215
loaded Volume V. This is a consequence of brittle failure from sparsly distributed flaws [9, 101 (it is more likely to find large flaws in large than in small volumes). In order to take into account for a inhomogenous and multiaxiale stress state, a suitable multiaxial failure criterion (resulting in an equivalent stress, oe)and the effectively loaded volume have to be defined (for details see [6, 81 and the literature therein). In the following the maximum principle stress is used in place of the equivalent stress and the effective volume is given by:
yet not fully understood and investigations.
still matter
of
ACKNOWLEDGEMENTS The authors thank EPCOS AG/Deutschlandsberg (Austria) for supplying samples, and acknowledge the collaboration of I. Hahn (SIEMENSMunich), J. Riedler and G. Schoner (EPCOS) for providing some of the requested material data and the helpful discussions on related topics. This work was supported by the Austrian Kplus-program.
REFERENCES Using a method for numerical integration, the so called Gaussian quadratur [111, the effective volume can be calculated from the FE model. Fig. 10 shows the relationship between the effective volume and the Weibull modulus for the point loading case, an older analytical approximation found in the literature [ 121 (it describes also the point loading case) and the used modified BDT configuration. Of course the effective volume goes to zero if the Weibull modulus goes to infinity (indicating no data scatter; then the fracture origin must start in the maximale loaded volume element). The reduction of the effective volume in the tests made with a curved support compared with the point loading case reflects the fact that the tensile loaded volume is significantly reduced if a sample is loaded on curved supports. This can qualitatively also be seen by comparison of Figs.7.a and 7.c. poln load: -=rvtical(Neergaard) -flatSuppott(FE) Wruppoct R=10,5mm
-
s
I
.-
I
>
5%
modul m
Fig. 10: Ratio of the effective volume Vcfand of the disc volume as a function of the Weibull modulus. Shown are different approximations. In general (as shown e.g. for some structural ceramics [6, 7, 8, 13, 141) a proper comparison of BDT- and bending test data necessitates to take into account for the different effective volumes in both kinds of testing (the effective volume in BD testing is about ten times higher than that of the mini bending samples; a Weibull extrapolation of strength for m=17 and V,lV2=10would result in an 15% increase in the bending strength). But for the investigated material no influence of volume under load on strength has been observed within the experimental scatter [ 13, 141 and the comparison of data as made in table 2 seems to be fair. This behaviour is
216
(1) B. W. Darvell, Uniaxial compression tests and the validity of indirect tensile strength, J. Mat. Sci., 25, (1990) 757-780 (2) A. Briickner-Foit, T. Fett, D. Mum, Discrimination of Multiaxiality Criteria with the Brazilian Disc Test, J. Eur. Cer. SOC.,17, (1997) 689 - 696 (3) R.H. Marion, K.J. Johnstone, A Parametric Study of the Diametral Compression Test for Ceramics, Cer. Bull., 56, (1977) 998 - 1002 (4) U. Soltesz, G. Bernauer, R. Schafer, Spaltzugversuchs-Eignungzur ZugfestigkeitsErmittlung sprodbrechender Materialien, Fachberichte DKG, 72, (1995) 553 - 555 (5) A. Borger, Optimierung des Scheibendruckversuches fur keramische Materialien, Diplomarbeit an der Montanuniversitat Leoben, Austria, ( 1999) (6) D. Munz, T. Fett, Ceramics Mechanical Properties, Failure Behaviour, Materials Selection; Springer Verlag, Berlin, Germany, (1999) (7) T. Lube, M. Manner, R, Danzer, The Miniaturisation of the 4-point bend Test. J. Fatigue Fract. Engng. Mater. Struct., 20, (1997) 1605-1616 (8) D. Rubesa, R. Darner; The Pecularities of Designing with Brittle Materials-Weak Point and Deficiencies, Proc. of the 12'hBienniel Conference on Fracture - ECF 12, EMAS Publishing, West Midlands, Sheffield, U.K, (1 998) (9) R. Darner; Ceramics: Mechanical Performance and Lifetime Prediction; Encyclopedia of Advanced Materials; 1, Oxford, (1994) 358-398 (10)R. Darner, A general strength distribution function for brittle materials. J. Eur. Ceram. SOC.10, (1992) 46 1-472 (1 l)A. H. Stroud and D. Secrest, Gaussian Quadrature Formulas, Prentice-Hall, Inc.. Englewood Cliffs. N.J.,USA, (1 969) (12)L.J. Neergaard; Effective volume of specimens in diametral compression, J. Mat. Sci., 3, (1997) 2529 - 2533 (13)R. Danzer, T. Lube, Fracture statistics of brittle materials: It does not always have to be Weibull statistics. Proc. 6'h Int Symp. Ceramic materials and components, Arita, Japan, ( 1998) 658-662 (14) R. Danzer, Mechanical behaviour and reliability of Ceramics. in. P. Vincencini ed. gthCimtec-World Ceramics Congress, Faenza, Italy, (1999) 379-386
THE IMPULSE EXCITATION TECHNIQUE FOR RAPID ASSESSMENT OF THE TEMPERATURE DEPENDENCE OF STRUCTURAL PROPERTIES OF SILICON NITRIDE AND ZIRCONIUM OXIDE CERAMICS G. Roebben*, R G. Duan, B. Basu, J. Vleugels, 0.Van der Biest Department of Metallurgy and Materials Engineering, Katholieke Universiteit Leuven, de Croylaan 2, B-3001 Heverlee, Belgium
ABSTRACT This paper presents the Impulse Excitation Technique (IET) as a means of non-destructive high temperature mechanical testing. IET enables to determine both elastic and damping or internal friction properties of small laboratory samples as well as industrial components. As an example, the changes with temperature of stiffness and internal friction of silicon nitride engine valves will be shown, and their impact on the valve performance will be discussed. Further the effect of temperature on the internal friction of small test samples of different grades of silicon nitride as well as zirconia ceramics (Y-TZP type) is reported.
INTRODUCTION: THE IMPULSE EXCITATION TECHNIQUE (IET) Careful characterisation of the mechanical properties of ceramic materials and components is required before structural application can be attempted. However, often the amount of test-material - in laboratoryscale processing - or test-time - in the field of industrial production - is restricted. This can prohibit large-scale strength statistics and creep exercises. Also, the latter tests are destructive, and unsuitable for quality control. Alternatively, a considerable amount of structural information can be obtained from small laboratory samples as well as finished components, in a rapid and non-destructive manner, using the Impulse Excitation Technique (IET).
IET is a particular type of resonance frequency analysis. It essentially consists of inducing the vibration of a freely suspended test sample by a gentle, non-destructive mechanical impulse. Forster described the principal test procedure as early as 1937 [l]. Developments in digital signal analysis allowed substantial progress in the IETdomain in the last few years [2]. The results obtained by IET consist of the resonance frequencies of the tested sample, and the rate at which each of these frequencies loose amplitude with time after the impulse excitation. The resonance frequencies are related to the sample’s stiffness, whereas the (exponential) amplitude decrease reflects material damping or internal fiction [3]. IET, as well as other resonance frequency tests, have been widely acknowledged as an economically viable and sensitive means of product quality control [4,5]. Moreover, the non-destructive character of IET implies that a single sample can be tested periodically, e.g. during a temperature cycle. In this paper, it is shown how IET provides access to changes in stiffness and internal friction with temperature. These observations can be related to otherwise less easily accessible properties as fracture toughness and creep resistance. EXPERIMENTAL
Impulse excitation equipment Impulse excitation tests were performed with a Resonance Frequency and Damping Analyser apparatus (RFDA, Integrated Material Control Engineering (IMCE), Diepenbeek, Belgium), which has been
217
Figure I a) Set-upfor tests in air constructed on thefrontflange (5) of thefurnace chamber, before insertion in the chamber. The rectangular test sample (I) is suspended with NiCr-wires (2) that are guided along freely turning ceramic cylinders (3) between parallel A1203rods. Two of the A1203rods are driven by a stepmotor to prevent change of sample position by thermal expansion and creep of the wires; b) Graphite sample suspension platform (5) equiped with a shielded thermocouple (2), a tube (3) through which a ceramic projectile is fired, with a ceramic engine valve ( I ) suspended by graphite threads (4), readyfor insertion in bottom-loaded graphite-heating-elementfurnace.
previously described in [2]. Tests were performed both in a fi-ont-loaded a i r - h a c e (Fig. 1 a, J. W. Lemmens, Leuven, Belgium), and in a bottom-loaded graphite fiunace (Fig. 1 b, HTVP-l75O-C, IMCE, Diepenbeek, Belgium). An automated impulse excitation method is used, consisting of an either pneumatically (air furnace) or electromagnetically (graphite furnace) fired ceramic projectile. The sample vibration is detected by default - with a microphone. Laser vibrometry provides an alternative for challenging atmospheres (vacuum) or samples (lightweight). The RFDA software analyses the digitally stored vibration signal. After an initial Fast Fourier Transformation, the program uses an iterative algorithm to simulate the time-domain sound signal, thus determining the main frequencies present in the recorded vibration, as well as the exponential loss factor corresponding to each of the detected frequencies. This loss factor indicates how fast the amplitude of the vibration component of that frequency looses amplitude. If the sample is fieely vibrating, ie when the sample is suspended in the nodes of the investigated vibration mode, then this amplitude loss is due to internal friction or material damping. In the case of rectangular bar samples, the procedures described in ASTM C 1259 - 94 enabled us to calculate 218
the E-modulus of the material from the bending frequency of the bar.
Materials NGK Ceramics Europe supplied silicon nitride combustion engine valves (SN73K grade) for high temperature characterisation with impulse excitation. The impulse excitation response of these ceramic valves has been compared with that of standard, steel alloy engine valves, ‘explanted’ from a commercial car engine. Further, small rectangular bars of inhouse hot-pressed silicon nitrides with different sintering additives were compared. Details of the powder preparation and sintering parameters are given in [6]. The starting materials used to prepare the powder mixtures are commercial Si3N4, Y2O3 and A1203 powder. One grade (SNA14) was prepared with 4wt% A1203. A second grade ( S N Y S ) contained 8 wt% Y203. The third grade (SNY6A12) was sintered with a mixture of Y203 (6 wt%) and A1203 (2 wt%). The hotpressed disks were machined into suitable samples for IET (nominally 30 x 5 x 2 mm3). In addition, two grades of hot-pressed tetragonal ZrO2 polycrystalline (TZP) ceramics were tested. The preparation of these materials has been described in [7]. The first grade (3Y-TZP) was prepared by hot-pressing
. ...
1
0015 -
~
0.95
c
--E 50
001 --
r
steel alloy valve * silicon nitnde valve +
0.75
4
100
j 300
5w
700
-c
om-
900
Temperature ('C)
Figure 2 Comparison of the change of resonance frequency of ceramic and steel alloy engine valves with temperature.
Tosoh TZ-3Y powder. The second grade (2YM-TZP) was obtained by hot-pressing in a similar way a mixture of Tosoh TZ-3Y powder and Tosoh TZ-0 powder, to obtain an overall Y2O3 content of 2.5 mol%.
RESULTS Effect of temperature on the stiffness and damping of combustion engine valves . Both silicon nitride and steel alloy engine valves were tested in the graphite furnace in a N2-atmosphere. The valve, suspended by graphite threads as shown in Fig. lb, is excited approximately halfway along the length of its stem. The major bending resonance frequency of the ceramic outlet valve (valve head diameter 31 mm, length 104 mm) is near 10 kHz. The steel engine valve (valve head diameter 24 mm, length 92mm) has a major resonance frequency near 7 kHz. In Figure 2, the relative change of these frequencies with increasing temperature is shown, up to a typical maximum operating temperature of the outlet valves (900°C). Whereas the stifhess of the ceramic valve is hardly affected in this temperature range, the E-modulus of the steel alloy valve changes by a factor of about 0.8 x 0.8 = 0.64 (taking into account the quadratic relation between E-modulus and resonance frequencies). Figure 3 shows the change of the internal friction in the silicon nitride valve with increasing temperature. The internal friction is very low until 800°C. From this temperature on, damping increases rapidly, until it reaches a value which is higher than
Figure 3 Internal friction in silicon nitride engine valve as a function of temperature, indicating the existence of a damping peak near 1000°C with a maximum value larger than can be measured with IET (upper limit 1.5%).
the upper damping value measurable with IET. This upper limit value depends on sample geometry and material, and appears to be about 1.5% for the investigated engine valve. At a temperature of about llOO°C, damping has decreased and can be determined again. All together, the data mark the existence of a large damping peak centred near 975°C.
Effect of sintering additive composition on damping in Si3N4 Figure 4 compares the level of internal friction between 75OOC and 135OOC in three hot-pressed silicon nitrides. The SNA14 material-data do not reveal a damping peak. Internal friction does increase slowly from about 1100°C. The SNY8 silicon nitride is characterised by a small but distinctive internal friction peak near 1070°C. At approximately the same temperature a much larger damping peak occurs in the SNY6A12 silicon nitride, sintered with a mixture of Y203 (6wt%) and A1203 (2wt%). The internal friction peaks are superimposed on an exponentially increasing background. It can be observed that, in the investigated cases, a large background coincides with a high internal friction peak. Figure 4b displays the results obtained during a second heating. Clearly, the first temperature cycle (at 2"C/min to 1400°C) has affected the material's response: the internal friction peaks have decreased.
219
004
-
-
.:
SN-YB - A SN-Y6A12 003 -,OSN-AM +
g 3 r
2
I
I
F l n t Heaths
I .
E
E
5 e!
001
E
~~~~
-
I
-
Second Heating
I
I
-
002
- - - --
__--
-
SN-YB . A SN-YW2 c 0 03 0 SN-AM
1 1 P
A
b
+
0
u . 002
004
I
Il
i
-
001
Figure 5 Internal friction peak in three grades of silicon nitride, with diflerent sintering additive compositions; a)first heating at 2"C/min to 1400°C; b) second heating at 2OC/min to 14OOOC; resonance frequencies: SNA14: 26 kHz, SNY8: 25 kHz, SNY6A12: I0 kHz.
Internal friction in Y-TZP Figure 5 displays the IET-results obtained on the 3Y-TZP (Figure 5a) and the 2YM-TZP (Figure 5b) materials when heating at 2OC/min in the air furnace. In Figure 5b, a calculated damping peak is superimposed on the measured data. This simulated damping peak corresponds to a Debye-peak, representing an ideal point defect type of relaxation. The Debye-peak parameters (activation energy and relaxation time) have been taken from Weller [8], who identified this peak as due to the repeated switching of elastically anisotrope dipoles consisting of oxygen vacancies and yttrium substitutional atoms. The activation energy and relaxation time determine the temperature at which the maximum damping is reached, The maximum damping value was estimated by fitting the Debye-peak height to the measured data.
DISCUSSION Comparison steel and ceramic engine valves Hamminger and Heinrich [9] reported a significant reduction in noise emission (by 18 dB at 3000 revolutions/min) from the cylinder head of a combustion engine when replacing conventional valves by silicon nitride valves. It has been suggested [ 101 that the increased stiffness of silicon nitride (300GPa) over steel (200 GPa) should be reflected in improved noise characteristics of ceramic components. Indeed, for a given valve geometry, the ceramic valves are expected to produce significantly less noise due to the inaudibility (> 20kHz) of their elevated resonance fkequencies. Moreover, the IET-results of Figure 2 show that the stiffness of the ceramic valve is largely unaffected by
-
measured
*calculatedwiih literaturedata
50
150
250
350
450
Tempmtun ('C)
550
650
750
50
150
250
350
Figure 4 Internal friction in 3 Y - Z P and 2YM-ZP.
220
450
Temp.Rtum
rc)
550
650
-
750
temperature within the range of normal valve operation. Steel, on the other hand, does loose 1/3 of its stiffness between room temperature and 900OC. The IET-results therefore indicate that the advantages of high stiffness become even more pronounced at more elevated operation temperatures. The damping peak, which happens to occur near the operation temperature of the outlet valve head, also affects the noise produced by the valve. The origin of the characteristic damping peak near 1000°C is known to be the glass transition of intergranular pockets of amorphous silicate phases [ 111. The internal friction or vibration energy dissipation at this T,-peak can reduce the noise produced by the valves. At temperatures near or below the T,-peak, the behaviour of the silicon nitride is unaffected by creep deformation. Therefore, the intentional operation of the valve in the temperature region of the T,-peak can be considered safe practice, and could be exploited in other vibration-critical applications.
Comparison of SNA14, S N Y S and SNY6A12 The IET-results shown in Figure4 demonstrate the effect of the use of different additives on the high-temperature structural performance of silicon nitrides. The SNA14 does not show a T,-peak. According to the work of Donzel [ll] this implies that this grade is free of amorphous glass pockets. Indeed, it is known that the used composition leads to the formation of a sialon material, in which A13+ions are substituting for Si4+.To maintain electrical neutrality the o2--ions enter the Si3Ns lattice as well and substitute for the N atoms. The incorporation of the oxygen atoms decreases the amount of intergranular silicate glass, and renders the sialon material highly creep resistant. The comparison of the internal friction curves of the SNY8 and SNY6A12 materials reflects the different crystallinity of their intergranular phases. The height of the damping peak of the SNY6A12 material indicates that the amount of residual intergranular pocket phase is larger than in the SNY8 material. XRD-investigation of the
materials in the as hot-pressed state [6], confirms the presence of a crystalline intergranular phase (Y-N-apatite) in S N Y 8 , while the SNY6A12 material contains only aand p-Si3N4. Although creep tests have not been performed on the three tested silicon nitride grades, it can be inferred from literature that the creep resistance of the three grades ranks in the same way as the internal friction T,peak height. Monophase sialon indeed is more creep resistant than the materials with intergranular residues of the sintering additives. Also, the creep resistance is increased by the (partial) crystallisation of the intergranular phase. However, since the internal fiiction peaks are caused by relaxation phenomena with limited amplitude, there is no direct micromechanical relation between the T,-peak and creep resistance. Schaller has shown that the exponentially increasing background, which is due to microplastic deformation, correlates better with creep resistance [12]. In the silicon nitride grades we have studied, the ranking according to peak height and according to background coincide. However, this ranking can be affected due to grain morphology and intergranular phase distribution. Figure 4b shows internal fiction during the second heating up to 1400°C. The damping peak is significantly reduced for both SNY8 and SNY6A12 materials, but the background is largely unaffected. In-situ high temperature X-ray dimaction [61 confirms the crystallisation of the intergranular pockets.
Comparison of 3Y-TZP and 2YM-TZP The internal friction data show a peak near 2OOOC for both the 3Y-TZP and 2YMTZP samples. The location of this peak depends slightly on the resonance frequency which was 7 kHz for the 3Y-TZP and 14 kHz for the 2YM-TZP. Otherwise, the peaks are comparable as both are caused by elastic dipole relaxation [81. At temperatures above the dipole peak the two material grades behave differently. The 3Y-TZP is known to be fully stable and resistant to thermally induced transformation of the tetragonal to the monoclinic phase. The 2YM-grade, due to the
22 1
lower Y-content, transforms more easily. The thermally induced transformation of tetragonal to monoclinic phase in 2YM-TZP between 100 and 400°C has been confirmed by high temperature XRD [13]. This phenomenon is known as low-temperature degradation, and is seen to affect the internal friction. As such, IET can be a tool in the assessment of the transformability of the tetragonal ZrOz-phase, which is one of the factors determining the fracture toughness of the transformation toughened zirconia materials. Of the tested materials 2YM-TZP has the larger room temperature indentation fracture toughness (10.2 am'") in comparison with the 3Y-TZP (2.5 MPam’”). However, issues of thermal instability, as indicated by the IET-results, can inverse the ranking of suitability for applications under specific circumstances. CONCLUSIONS AND PROSPECTS
In this paper, tests performed with advanced impulse excitation devices have shown that the elastic and damping properties of ceramic materials and components can be accurately determined as a h c t i o n of temperature. It has been demonstrated that the advantageous noise emission properties of silicon nitride combustion engine valves can be explained by their stiffhess and damping characteristics. Further, IET was used to infer the presence or absence, and the amount of pockets of amorphous intergranular phase in Si3N4 from the internal friction T,-peak, while the creep resistance relates to the internal friction background. Finally, the martensitic transformation in Y-TZP, which determines the material’s fiacture toughness, has been monitored by JET as well. Future efforts will be concentrated on obtaining quantitative relations between the IET-accessible properties on the one hand, and creep and fracture resistance on the other hand. ACKNOWLEDGEMENTS The authors wish to thank Carine Dewitte (NGK Ceramics Europe) for supplying the Si3N4 engine valves. GR and 222
RGD are fellows of the Fund for Scientific Research - Vlaanderen. BB is supported by a fellowship of the K.U.Leuven Research Council. REFERENCES (1) F. Forster, Z. Metallkunde, Vol. 29 (1937) 109-115. (2) G. Roebben, B. Bollen, A. Brebels, J. Van Humbeeck, 0. Van der Biest, Rev. Sci. Inst. 68 (1997) 4511-4515. (3) A. S. Nowick, B. S. Berry, Anelastic relaxation in crystalline solids, Academic Press, New York (1972). (4) J. W. Lemmens, Dynamic Elastic Modulus Measurements in Materials, ASTM STP 1045, ed. by A. Wolfenden, American Society for Testing and Materials, Philadelphia (1990) 90-99. ( 5 ) H. Lindner, B. Capers, H. Feuer, J. Hennicke, P. Claeys, W. Kaesler, K. S p i d e r , K.-L. Weisskopf, Werkstoffivoche ‘96, Symposium 2, ed. U. Koch, DGM, Frankfbrt (1996) 1-6. (6) R.-G. Duan, G. Roebben, J. Vleugels, 0. Van der Biest, to be published. (7) B. Basu, L. Donzel, J. Van Humbeeck, J. Vleugels, R. Schaller, 0. Van der Biest, Scripta Mat. 40 (1999) 759-765. (8) M. Weller, H. Schubert, P. Kontouros, Science and Technology of Zirconia V, ed. by S. P. S. Badwal, M. J. Bannister, and R. H. J. Hannink, Technomic Publ. Co., Lancaster and Base1 (1993) 546-554. (9) R. Hamminger, J. Heinrich, Mater. Res. SOC.Symp. Proc., 287 (1993) 513-520. (10) S. Yang, R. F. Gibson, G. M. Crosbie, R. L. Allor, Ceramic Industry, October 1995, 117-121. (1 1) L. Donzel, A. Lakki, R. Schaller, Phil. Mag. A 76 (1997) 933-944. (12) R. Schaller, to appear in J. Alloys and Compounds. (13) B. Basu, J. Vleugels, 0. Van der Biest, to be published.
VAMAS Round Robin Testing of High Temperature Flexural Strength Akira Okada and Mineo Mizuno Japan Fine Ceramics Center, Nagoya, Japan ABSTRACT
JFCC. Most participants received twelve specimens
Thirteen laboratories in six countries participated in
while some participants received more. In some cases,
the VAMAS round robin to measure the flexural strength
specimens were supplied from the additional set of six
of silicon nitride at high temperatures. At each laboratory,
plates (lot numbers 11 to 16) of 51 x 51 x 11 mm3 for
the four-point flexural strength of silicon nitride was
additional tests.
measured in air at 1200°C. The results are summarized as
Each test specimen was placed on a four-point flexural
follows. (i) A comparison between semi- and fully
fixture having a semi- or fully articulated configuration.
articulating fixtures indicates that fully articulating
Specimens were then heated to a temperature of 1200°C
fixture has the greater strength; (ii) The strength of 30
in a furnace and soaked at this temperature for some time
mm x 10 mm spans is greater than that of 40 mm x 20
to ensure a uniform temperature distribution throughout
mm spans; (iii) The flexural strength in nitrogen at
the specimen. Most of the fracture tests were conducted
1200°C is found to be considerably lower than that in air.
using a crosshead speed of 0.5 m d m i n . The flexural
The mechanisms responsible for such behavior are
strength was calculated according to the equation:
discussed.
INTRODUCTION
0,=
3P@, - L, ) 2bh2
Flexural strength is one of the most important properties of advanced ceramics and the temperature
where P is the maximum load at fracture, L , is the outer
dependence is the principal indicator of mechanical
span, L2 is the inner span, and b and h are the width and
performance at elevated temperatures. Since domestic
height of the specimen, respectively.
standards for high temperature flexural strength have
Weibull
statistical
analysis
was
conducted
in
been established in several countries [ l , 21, the
accordance with the maximum-likelihood method [4-61
international standardization of flexural strength at high
and the 90% confidence intervals were calculated [5, 61.
temperatures is now being discussed [3]. Within the Versailles Project on Advanced Materials and Standards
Table 1. Participating laboratories
(VAMAS), the present round robin was proposed in
R Westemeide: FtaunhoferW t u t fiir Werkstohha& Germany
order to promote the international standardization
A. Okada and M. Mizuno: Japan Fine Ceramics Center, Japan
process. Thirteen laboratories from six countries took
T. Yonezawa:JMC New Materials Inc, Japan
part (see Table I). The main objectives of the round
S. -J. Cho:Korean Research Instituteof Standardsand Science, Korea
robin were to clarify the effects of using an articulated
I. Miyachi: Kyocera Corporation,Japan
configuration and the size of the spans in four-point
S. R Choi:NASA Glenn Research Center, USA
flexure.
S. Sakaguchi: National Industrial Research InstituteofNagoya, Japan
T Tanah: National Institutefor Research in Inoqganic Materials, Japan EXPERIMENTAL PROCEDURES Advanced silicon nitride (SN28 I , Kyocera Corp., Kyoto, Japan) was used for this round robin. Specimens were machined to dimensions of 3 x 4 x 45 mm3 from nine plates (lot numbers 1 to 9) of 50 x 70 x 12 mm3 at
R Morrell:National Physical Laboratory,UK
H.Sakai: NGK Insulators Lid, Japan K U m M NGK Spark Plug Co. Ltd, Japan K Bredm Oak Ridge National Laboratory,USA
V. M. Sglavo: Univmita d’egli Studi di Trento, Italy
223
The two-parameter Weibull distribution function is given
calculated using all the data shown in gray. The large
by
variation in Weibull modulus between each laboratory can be seen and some of the 90% confidence intervals lie far outside those obtained from all the data for the same category of test. The Weibull moduli determined from all
where F is the failure probability under the applied
the data for semi-articulating fixtures with 30 mm x 10
stress cr, m is the Weibull modulus and 00 is the Weibull
mm spans, semi-articulating fixtures with 40 mm x 20
characteristic strength. In the present study, the estimator
mm spans and fully articulating fixtures with 40 mm x
of the fracture probability F, at the stress cr was chosen in
20 mm spans are 11.7, 9.4 and 9.9, respectively. The
accordance with JIS R 1625 [4] to be:
90% confidence intervals for these categories are from 10.0 to 13.7, 8.2 to 11.0, and 8.1 to 12.2, respectively.
6. =- i - 0.3
n + 0.4
(3)
As a result, the values of m and their 90% confidence
01
9oo
LI
' ' ' '
I ' ' ' ' I
' ' ' I '
+
' ' I
'J
The flexural tests in air using semi-articulating fixtures of 30 mm x 10 mm spans were performed in Labs A, C,
D, E and F (The thirteen participating laboratories were labelled randomly as Labs A to M). The tests conducted in Lab B were performed in nitrogen atmosphere and similar tests were also carried out in Lab A.
Tests using
semi-articulating fixtures with 40 mm x 20 mm spans
:I1 , , * *
I , , , , , , ( , ,
l,,8,1j
0
Pamciponts (Labs A to M)
were performed in Labs A, G, H, I, K and M, and Labs H,
J, K and L performed tests using fully articulating fixtures with spans of 40 mm x 20 mm. The flexural strengths of as-machined specimens and as-heat-treated specimens were measured in Lab A at room temperature. Heat-treatment was conducted at 1200°C for 10 minutes in order to provide a similar
thermal history to the high temperature flexural tests. RESULTS AND DISCUSSION All the data are summarized in three categories: semi-articulating with spans of 30 mm x 10 mm, D
semi-articulating with spans of 40 mm x 20 mm and
~
F
!idly articulating with spans of 40 mm x 20 mm. According to this classification, the strength of specimens measured using semi-articulating fixtures with 30 mm x 10 mm spans is 661+75 MPa averaged over 70
specimens; the strength using semi-articulating fixtures with 40 mm x 20 mm spans is 622+89 MPa averaged over 83 specimens; and the strength using fully
Fig. 1. The results of statistical analyses: a) average
articulating fixtures with spans of 40 mm x 20 mm is
and standard deviation of flexural strength, b) Weibull
65 1+101 MPa averaged over 40 specimens (Fig. la).
moduli m with 90% confidence intervals,
Figure l b shows the Weibull modulus m and 90% confidence interval for each laboratory with the value
224
and c)
Weibull characteristic strengths, cro, with 90% confidence intervals.
intervals in the three categories are similar to each other,
fixture [9-111. In light of these analyses, a fully
indicating that the Weibull modulus is close to 10.
articulating fixture appears to be the most accurate and
Figure 1c shows the Weibull characteristic strengths
should be superior to the semi-articulating fixture
no with 90% confidence intervals. The Weibull
because of its ability to reduce errors due to twist. The
characteristic strengths, no, and their 90% confidence
maximum stress om,,due to beam twisting is given by [8,
intervals are significantly different between three
9, 111
categories. The Weibull characteristic strength from semi-articulating fixtures with 30 mm x 10 mm spans is 691 MPa and the 90% confidence interval is from 678 to
(4)
703 MPa. This is 5.3% higher than that for 40 mm x 20 mm spans since the characteristic strength is 656 MPa and the 90% interval is from 643 to 670 MF’a in this case.
where a
The 90% confidence interval, however, reveals that the
function of blh, and x is given by
=
(L I
-
L2)12, k2 is a numerical factor that is a
difference is significant because the lower bound for spans of 30 mm x 10 mm is higher than the upper bound for 40 mm x 20 mm spans. In the case of fully articulating fixtures with 40 mm x 20 mm spans, the characteristic strength is 689 MPa and the confidence interval extends from 670 to 708 MPa.
Although this is
where E is Young’s modulus, v is Poisson’s ratio and LT
only 5.0% greater than that for semi-articulating fixtures
is the total length of the beam. Since blh
with the same spans, the lack of cross-over between the
numerical values of kl and kz are 0.179 and 0.224,
90% confidence intervals indicates that the difference is
respectively.
significant.
length and QF is the angle of twist between a pair of load
The
strength
combinations of spans seems to be connected to the
opposite longitudinal faces of 0.015 mm, the maximum
statistical aspect of brittle failure. The differences in the
value in the twist angles is approximately estimated to be
mean
Qs
characteristic
the
Qsis the angle of twist along the specimen
and contact points. From the parallelism tolerance on the
Weibull
between
1.33, the
different
and
variation
=
strengths
in
=
0.015 m d 4 . 0 mm and QF = 0.015 m d 1 2 . 0 mm,
semi-articulating fixtures between 30 mm x 10 mm spans
assuming that the minimum length of the rollers is three
and 40 mm x 20 mm spans are 6.3% and 5.3%,
times the width of the specimens.
respectively. Assuming that the Weibull modulus is 10,
It is noteworthy that failure occurs after bottoming of
the difference in the effective volumes is calculated to be
the specimen when 0 < x < 1, and that failure prior to the
x
85%, which is equivalent to a 6.3% difference in Weibull
bottoming occurs when
characteristic strengths. In addition, an analysis of the
round robin, the maximum possible error in the present
effective surface area leads to exactly the same answer of
configuration of test specimens is calculated to be 0.9%
6.3% difference in strengths. Although the observed
because of
variation is slightly smaller than that predicted, the
strength error is expected to be much smaller because
results seem to be consistent since the details of each test
such high stresses are generated in a limited volume and
in the round robin differ slightly from each other.
the strength depends on the size of the stressed volume.
x
=
=
1.
Since, in the present
0.32 at the stress of 650 MPa. The
The difference in the alignment accuracies may also
The difference in strength between semi- and fully
affect the strength variation between articulating fixtures.
articulating fixtures therefore seems too large according
Marschall and Rudnick reported error analyses from
to twist error analysis.
four-point flexure tests [7] and a fixture design for minimizing such effects has been proposed [8]. Further
Some specimens obtained from lots 2, 5 and 6 exhibit considerably lower strength. Figure 2 shows strengths
analysis indicated that friction is the major cause of the
plotted against lot number. In all three categories,
error and it is reduced by using a semi-articulating
specimens with low strength are found for lots 2, 5 and 6.
225
-
700
t 8
n
5
600
5 ul
-
500
L
In
400
a u
1 -
0
1 PactMpants (labs A to M)
J
2 U
300
0
rn 2 00
-
-
-
100
0
I
.
1
.
I
.
I
.
I
,
I
.
I
,
I
.
l
Fig. 2. Strength plotted against lot number. Open circles: semi-articulating fixtures with 30 mm x 10 mm spans; closed circles: semi-articulating fixtures with 40 mm x
Pamclpants (labs A to N)
20 mm spans; closed square: fully articulating fixtures
with 40 mm x 20 mm spans. The possibility of severe processing flaws pre-existing in these specimens seems to be greater than for the other lot numbers. Figure 3 shows the results of eliminating data originating from lots 2, 5 , and 6. Mean strength with standard deviation is shown in Fig. 3a. In comparison
$ z m 4 Pamcipants (labs A to M)
with Fig. la, it is clear that the scatter in the data is considerably reduced and the mean strength slightly
Fig. 3. The results of eliminating the data in lots 2, 5 , and
increased as a result of the data selection. Weibull moduli
6: a) average and standard deviation of flexural strength,
obtained from the selected specimens are shown in Fig.
b) Weibull modulus m with 90% confidence interval, and
3b. The value of the Weibull modulus increases when
c) Weibull characteristic strength cso with 90%
only selected data, are used. Weibull characteristic
confidence interval.
strengths are plotted in Fig. 3c. The strength increases slightly as a result of data selection although there are no
x 20 mm spans and fully articulated configuration with
significant effects on the 90% confidence intervals. It is
40 mm x 20 mm spans. In all cases, the values of the low
clear that such behavior results from competition
strength data are inconsistent with the linear relationship
between two factors. One is the reduction in the scatter
of two-parameter Weibull statistics, suggesting the
of the data leading to a narrowing of the confidence
bimodality in strength distribution. However, selective
interval and the other is the decrease in the number of
removal of some of data changes the distribution,
data points leading to an expansion of the interval.
especially in the low strength region, leading to good
Eliminating data from lots 2, 5 and 6 leads to an
agreement with a two-parameter Weibull statistical
improved fitting to a Weibull statistical function. Figure
function. This indicates that the statistical distribution of
4 shows the Weibull plots obtained for the three
strengths is different between two groups: specimens of
categories: semi-articulated configuration with 30 mm x
lots 2, 5 and 6 have a wide distribution of strengths while
10 mm spans, semi-articulated configuration with 40 mm
the distribution of the other specimens is relatively
226
t L
Lab. B tested in nitrogen rn = 13.5 SO = 640 MPa n = 10
sO=691 MPa
selected data:
..
700 800 flexural Strensth [MPoJ
600
500
900
300
IMX)
400
500
600
700
800 900 1000
flexural strength (MPa)
Fig. 5. Flexural strength of silicon nitride measured in nitrogen atmosphere at 12OO0C (Lab B). The plots are
serni-ortiiulating
compared with tests conducted in air under the same a11 the data
conditions:
SO = 656 MPa
30
spans of
mm
x
10 mm
and
semi-articulating fixture. SO = 672 MPa 99
300
500
400
600
700
8M)
900 1000
f
spars: 40 mm. 20 mm fully articulating
70: 50
all the data rn = 9.9 SO = 689 MPa n = 40
30-
-
1
20
9
l o 5 -
f
SO = 656 MPa
50 3 3 0 20
as-heat-treated
f
10
SO = 774 MPa
P
5
I
500
n = 29
temperature
I
300
400
I
500
I
600
I
700
I
800
700
900
1000
Fig. 6. Weibull plots of flexural strengths at room
SO = 707 MPa I
600
flexural strength (MPa)
selected data: no lot NO.2. 5 and 6
a a a
I
as machined
70
g
-
90
bQ
90
2
flexural strength (MPol
99
b?
l
for
as-machined
specimens
and
for
l
600 900 1000
flexural strength (MPol
heat-treated specimens. Heat-treatment was carried out at 1200°C in air for 10 minutes.
Fig. 4. Weibull plots of flexural strength. Comparison is made between using all the data and using selected data
oxidation effects are particular to the measurements in air.
after eliminating lots 2, 5 , and 6. a) Spans of 30 mm x
Additional measurements of flexural strength were thus
10 mm, semi-articulated configuration, b) spans of 40
performed
mm x 20 mm, semi-articulated configuration, c) spans of
specimens and for heat-treated specimens. This indicates
40 mm x 20 mm, fully articulated configuration.
that a significant increase in the flexural strength of 17%
at
room temperature
for as-machined
occurs during exposure in air atmosphere at 1200°C (see Fig. 6 ) . It is therefore suggested that the higher strengths
narrow.
In the present round robin, the high temperature
in air result from oxidation during flexure testing.
The
in nitrogen
most plausible explanation is that the oxidative crack
atmosphere at 1200°C and it was considerably lower than
healing takes place, since in silicon nitride significant
that in air (see Fig. 5).
healing accompanying strength recovery has been
flexural strength was also measured
During high temperature
flexural testing, the specimens suffer both thermal and
reported [12, 131.
atmospheric attack. The thermal history is essentially the same for specimens in air and nitrogen atmospheres but
227
Ceramics, Advanced Technical
CONCLUSIONS
Ceramics) -Test
Round robin tests to measure the flexural strength of
Method for Flexural Strength of Monolithic Ceramics
silicon nitride were performed at 1200°C. The tests were
at Elevated Temperatures,” IS0 Technical Committee
carried out under four-point flexure with either spans of
206 (ISOiTC206iWG8).
30 mm x 10 mm or 40 mm x 20 mm, and using either
Data for Fine Ceramics,” Japanese Industrial Standard
semi- or fully articulating fixtures. The strength data obtained at each laboratory exhibit considerable scatter due to the limited number of test specimens.
However,
Weibull
4. JIS R 1625 (1996), “Weibull Statistics of Strength
statistical
analysis
considering the confidence intervals revealed that the
Committee, Japanese Standards Association, Tokyo.
5. D. R. Thoman, L. J. Bain and C. E. Antle, “Influence on the Parameters of the Weibull Distribution,” Technometrics, 11 [3], (1969) 445-460.
strength in the spans of 30 mm x 10 mm is greater than
6. ASTM Designation: C 1239-95, “Standard Practice for
that of 40 mm x 20 mm and the fully articulating fixture
Reporting Uniaxial Strength Data and Estimating
leads to greater strength than semi-articulating fixtures.
Weibull
This is consistent with the prediction from effective
Ceramics,” American Society for Testing and Materials,
volume/surface area analysis but inconsistent with twist
1995.
Distribution
Parameters
for
Advanced
error analysis because the close tolerance between
7. C. W. Marschall and A. Rudnick, “Conventional
parallel sides of the test specimens does not allow such a
Strength Testing of Ceramics”, pp. 69-92 in “Fracture
large difference in strength between the fully articulating
Mechanics of Ceramics, vol. 1,” edited by R. C. Bradt,
and semi-articulating fixtures.
D. P. H. Hasselman and F. F. Lange, Plenum Press,
The strength measured in nitrogen at 12OO0C was
New York, 1974.
result of oxidative crack healing occurring at 1200°C in
8. R. G. Hoagland, C. W. Marschall and W. H. Duckworth, “Reduction of Errors in Ceramic Bend
air, since the strength of specimens heat-treated in air at
Tests,” J. Am. Ceram. SOC.,59 [5/6], (1976) 189-192.
1200°C for 10 minutes was 17% greater than that of
9. F. I. Baratta, “Requirements for Flexure Testing of
considerably lower than that in air. This seems to be a
as-machined specimens.
Brittle Materials,” pp. 194-222 in “Methods for Assessing the Structural Reliability of Brittle Materials,
ACKNOWLEDGMENTS The authors are grateful to all the participants involved in this round robin. Helpful comments by George Quinn,
ASTM STP 844,” edited by S. W., Freiman and C. M. Hudson, American Society for Testing and Materials, Philadelphia, 1984.
Sung R. Choi, Stephen D u Q , Kristin Breder and Roger
10. L. R. Swank, J. C. Cacery and R. L. Allor,
Morrell are appreciated and special thanks are due to G.
“Experimental Errors in Modulus of Rupture Test
Q. and S. R. C. for their assistance in calculating errors
Fixtures,” Ceram. Eng. Sci. Proc., 11 [9/10], (1990)
of beam twisting. Support by the Science and
1329-1345.
Technology Agency in Japan is acknowledged.
11. G. D. Quinn, “Twisting and Friction Errors in Flexure Testing,” Ceram. Eng. Sci. Proc., 13 [7/8], (1992) 3 19-330.
REFERENCES 1. JIS R 1604 (1987), “Testing Method for Flexural
12. T. Ogasawara, T. Hori and A. Okada, “Threshold
Strength of Fine Ceramics at Elevated Temperature,”
Stress Intensity for Oxidative Crack Healing in
Japanese Industrial Standard Committee, Japanese
Sintered Silicon Nitride,” J. Mater. Sci. Lett., 13 [6],
Standards Association, Tokyo.
(1994) 404-406.
2. ASTM Designation: C 1211-92, “Standard Test
13. S. R. Choi, V. Tikare and R. Pawlik, “Crack Healing
Method for Flexural Strength of Advanced Ceramics at
in Silicon Nitride Due to Oxidation,” Ceram. Eng. Sci.
Elevated Temperatures,” American Society for Testing
Proc., 12 [9/10], (1991) 2190-2202.
and Materials, 1992. 3. ISO/CD 17565 (1999), “Fine Ceramics (Advanced
22 8
EFFECT OF HIGH VOLTAGE SCREENING METHOD ON TITANIA CERAMICS WITH DIFFERENT SURFACE FINISHING A. Kishimoto and T. Tanaka Institute of Industrial Science, University of Tokyo, 7-22-1 Roppongi, Minato-ku, Tokyo 106-8558, Japan
ABSTRACT Effects of high voltage screening were examined on perpendicular and parallel surface ground titania rectangular bars. A screening field at or below which 30 % of titania samples break electrically was applied to each group samples. After high voltage screening, the surviving samples were subjected to mechanical strength measurement and resultant strength distribution was compared to the original distribution. After screening, the Weibull plots of perpendicular ground sample bent to become a convex curve while plots in the high strength region remained almost the same, indicating that low strength samples were selectively eliminated by the high voltage screening. On the other hand, screening effect on parallel surface ground sample was very small.
INTRODUCTION We previously reported the effects of high voltage screening of dielectric ceramics, by which mechanically weak ceramic parts can be selectively removed by an electric method (dielectric breakdown) [ 1-41. The mechanism is thought based on the analogous roles of specific microstructures in mechanical fracture and in dielectric breakdown [5-81. For example, pores often behave as fracture origins in mechanical fracture, as they are usually stress concentrators. They also become the starting points of dielectric breakdown, because the electric field is inversely proportional to the ratio of electric permittivity (relative permittivity: the E, of an air gap is approximately unity, which is usually smaller than the ceramic's bulk) [8,9]. We previously examined the effect of high voltage screening of titania ceramics with a rough surface finish, and reported that a mechanically weak ceramic part has a relatively low dielectric strength [4]. We concluded that a surface crack beneath a ground groove is the decisive breakdown flaw, as with a mechanical fracture. These results suggest that such a flaw plays the role of an electric field concentrator, similar to the role of a crack in the concentration of mechanical stress. It is important to now examine high voltage screening of unidirectionally surface ground ceramics with crack-like surface flaws that are parallel to the direction of grinding, since in the previously reported study the surface grinding was in random directions [4]. In mechanical fracture, it is well known that thin surface
flaws perpendicular to the tension direction become serious stress concentrators [lo]. However, with a parallel electrode configuration in the electric breakdown experiment, surface flaws should play an identical role, We irrespective of their planer orientation [ 1 11. examined crack-like surface flaws, which may be a common weak spot in both mechanical and dielectric failure, and the influence of surface flaws on mechanical strength distribution after high voltage screening of unidirectionally surface ground titania ceramics.
EXPERIMENTAL PROCEDURE Sample Preparation Titanium dioxide ceramics were employed because they can be easily broken electrically at room temperature. Titanium dioxide powder (Kojundo Chemical Co. Ltd., Japan, rutile phase, purity 99.99%) was used as the starting material. Powder compacts were first formed by uniaxial pressing (30 MPa for 60 s) and subsequently fabricated by hydrostatic pressing (200 MPa for 90 s). The compact bodies were sintered at 1450°C for 4 h in air. The sintered bodies were cooled at 150 "C /h, to prevent reduction from occurring. The resultant bodies had relative densities around 98.5%. After removing the surface layer, they were cut into rectangular 13 X 4 X 0.3 mm bars with a precision cutting machine (Maruto Co. Ltd., Japan, MC-603). The cut ends of the samples were comparable to or smoother than a surface finished with 800-grit abrasive paper. To examine the effect of the direction of surface grinding, the 13 x 4 mm planes were ground using #400 abrasive paper. In one group, the grinding was perpendicular to the long axis and in the second group grinding was parallel to the long axis. The groups were called the perpendicular and parallel groups, respectively. In each group, 30 bars were used to measure strength distribution and approximately 45 bars were used for screening. The ground surfaces were observed by SEM
Strength measurements and high voltage screening Silver electrodes with a diameter of 2.5 mm were attached to both sides of each test piece. The electrodes were made with diffused edges to prevent concentration
229
.
2
.
.
,
.
.
. , . . . , . .
,
2
i *perpendicular
1
E=204(MPa) m=8.8
-0- perpendicular
d
L; -2 I= -
-4
.
- 5 t . ’ .
4.8
I
5
*
“
I
.
5.2 In0
.
.
I
5.4
”
’
3
m=8.1
-4 ’
5.6
Fig. 1 Weibull plots of mechanical strength on Ti02 ceramics ground with different direction.
of the electrical field at the edges of the electrode. The mechanical strength of 30 samples was measured using a three point bending test, with a span length of 10 mm and a crosshead speed of 0.5 mm / min. In order to compare the mechanical strength with the results of electrical screening, the maximum bending stress was applied to the centerline of the electrode. Fractured samples were subjected to dielectric breakdown measurement to determine the screening field. The breakdown test applied D.C. voltage, increasing at a rate of 50V/s. Test pieces were placed in silicon oil to prevent surface flashover. The electric field at which the current abruptly increased was regarded as the dielectric strength (Eb). The measured mechanical and dielectric strength distributions of the perpendicular and parallel groups were compared, and the screening field (Es) at or below which 30% of each group of samples broke electrically was determined from the dielectric strength distribution. The other 45 bars for screening were subjected to the breakdown (screening) test, in which the strength of the electric field was increased to that of the screening field (E,). Bars that broke during the screening test were discarded. The mechanical strength of the surviving samples was measured; during this test, maximum stress was applied at the centerline of the electrode. Mechanical strength with or without screening was compared using Weibull statistics. Evaluation of strength distribution The distributions of mechanical strength and dielectric strength were estimated using a two-parameter Weibull distribution, as follows: F = 1 - exp[-(s/so)mVl (1) where s is the mechanical or dielectric strength, and so, m, and V are the scale parameter, shape parameter (= Weibull modulus), and effective volume for each strength, respectively. The cumulative probability, F, was calculated using the mean rank method. For the mechanical stress screening, the probability of failure after screening, F(a) is expressed as, F(a) = (Ftotal- Fs)4 1 - Fs) (2). where Ftotalis the probability of failure without screening and F, is the probability of failure when the screening stress has been applied [13]. With this equation, the
230
-3
2.8
3
hE
3.2
3.4
Fig. 2 Weibull plots of dielectric strength on Ti02 ceramics ground with different direction.
strength distribution after screening is a convex curve approaching the screening stress, q..
RESULTS AND DISCUSSION Mechanical and dielectric strengths of parallel and perpendicularly ground samples Figure 1 shows the Weibull plots of the mechanical strength for the perpendicular and parallel surface ground samples. Both plots show good linearity (correlation coefficient r > 0.95), indicating that the scatter of each data set can be expressed by a single-mode Weibull distribution function. The sample with the surface ground perpendicularly has a relatively small Weibull modulus (8.8) and average strength, indicating that grinding introduces surface flaws, which play a role in the origin of mechanical fractures. Figure 2 shows the Weibull plots of dielectric strength for the two types of sample. There is no significant difference between the samples ground perpendicularly and those ground parallel. The Weibull moduli for the perpendicularly and parallel ground samples were 8.3 and 8.1, and the mean dielectric strengths were 242 and 240 kV/cm, respectively. For perpendicularly ground samples, the distributions of mechanical and dielectric strengths are very similar. One of the authors has already reported the similarity of the two strength distributions, indicating the analogous nature of the distribution of weak spots in the two types of failure [4-81. 3-2 High voltage screening of perpendicularly ground samples High voltage screening of the perpendicularly and parallel ground samples was conducted. From the dielectric strength distribution shown in Fig.2, the screening field (Es)was determined to be 222 kV/cm, for which the cumulative dielectric failure probability was 30% in both the perpendicular and parallel cases.
2
-co 1
4
y -1 u 4
c C -
-2
-3 -4
-5
48
5
52 ha
5.4
56
Fig 3 High voltage screening on perpendicular surface ground Ti02 ceramics
Figure 3 illustrates the Weibull plots of the mechanical strength of perpendicularly ground samples with or without high voltage screening. High voltage screening selectively rejects parts with a low mechanical strength. After high voltage screening, the Weibull plot bends to form a convex curve, although the plots in the high strength region are almost the same. This result shows that the electric method removed mechanically degraded samples. In other words, samples with a severe mechanical flaw tend to have weak dielectric strength. Incidentally, the convex curve drawn with a solid line in Fig. 3, denoted 0~30, is the theoretical mechanical strength at which 30% of the samples fail during stress screening. With high voltage screening, some samples have strengths lower than 0~30. This occurs because the correlation between mechanical and dielectric strengths is not perfect. Matsuo et al. reported that median cracks in a scratch groove are introduced parallel to the direction of grinding [lo]. A coarse abrasive particle would have increased compressive weight, increasing the depth and number of such flaws. For this reason, in perpendicularly ground samples the distribution is thought to spread to lower region and the average strength to decrease. Some authors have reported that when an electric field is concentrated on a concave surface, a critical electric field induces an electric avalanche and dielectric breakdown occurs [6,12], although there are few papers comparing breakdown strength and surface morphology. In our results, high voltage screening selectively eliminated mechanically weak samples, indicating that as with mechanical fracture surface cracks beneath the grooves caused by grinding are the decisive breakdown flaws. Analogous to the stress concentration in a mechanical fracture, electric field concentration occurs at the end of a crack due to the permittivity difference, leading to the analogy between the mechanical and dielectric strengths. In perpendicularly ground samples, crack-like flaws are formed perpendicular to the direction of tension. As a result, cracks that concentrate the electric field also concentrate mechanical stress. In other words, an
4.8
5
5.2 ha
5.4
5.6
Fig 4 High voltage screening on parallel surface ground Ti02 ceramics
electrically weak sample would also be mechanically weak, leading to the positive correlation.
High voltage screening of parallel ground samples Figure 4 illustrates the Weibull plots of the mechanical strength of samples with the surface ground parallel to the tension direction, with and without high voltage screening. Unlike the perpendicularly ground samples, no significant difference is observed with or without screening. The theoretical mechanical strength at which 30% of the samples should fail after stress screening is also very different from the experimental result. These results indicate that for parallel ground samples the correlation between mechanical and dielectric strength is very weak, unlike for perpendicularly ground samples. These results indicate that some of the microstructures introduced by grinding play different roles in the two types of failure. Role of cracks introduced by surface grinding The reason that high voltage screening appears effective only for perpendicularly ground samples can be explained in the following way. First, the geometric shape and dimensions of the surface cracks introduced, which can be derived from Fig. 2, are almost the same in parallel and perpendicularly ground samples. As shown in Fig. 1, the effects of such a crack on the three point bending strength differ with different crack directions. When cracks form that are perpendicular to the tension, they can become serious flaws, depending on their dimensions. However, cracks that form parallel to the direction of tension are less serious flaws, even if large. As a result, in the former case, if the fracture stress or the breakdown field is applied to the same area, both failures start at the same point, resulting in the correlation of the two failure strengths. In the parallel case, the starting points are not expected to agree, and therefore screening has no apparent effect. The connection between electric field distribution and stress distribution in our experiment explains why the correlation between mechanical and dielectric strength is not perfect in the perpendicularly ground
23 I
samples. In this study, parallel plates with the same dimensions were used as electrodes, forming an almost uniform electric field between the electrodes. In such a case, the effect of a microstructural weak point is equivalent, regardless of where it is positioned between the electrodes. On the other hand, in mechanical bending tests, it is widely known that there is a stress gradient in the thickness direction; tensile stress is generated on the bottom face and compressive stress is generated on the top face. A large crack on the top face might therefore not be recognized as a weak point in the mechanical bending test, whereas the same crack would act as a serious electric flaw in the parallel plates electrode system. In such a case, another weak spot, such as an internal void near the bottom face could act as a stress concentrator in a bending fracture. In conclusion, the results of high voltage screening are similar to those of stress screening when surface flaws are preferentially oriented perpendicular to the tensile direction, but surface flaws that parallel the direction of tension are much less apparent in measurements of bending strength. In other words, high voltage screening of ceramic parts can detect significant flaws that are parallel to the tensile direction, which cannot be detected by mechanical screening using flexure or tensile methods.
REFERENCES 1)A. Kishimoto, K. Endo, Y. Nakamura, N. Motohira, H. Yanagida, and K. Sugai, Effect of high voltage screening on strength distribution for Ti02 ceramics, J. Am. Ceram. SOC.,78 [8] (1 995) 2248-2250. 2) K.Endo, A.Kishimoto, N. Motohira, Y. Nakamura and H. Yanagida, High Voltage Screening on Ceramics Parts with Bimodal Strength Distribution, J. Ceram. SOC.Jpn., 103[lo] (1995) 1090-1092. 3) A.Kishimoto, K.Endo, N. Motohira, Y. Nakamura H. Yanagida, and M. Miyayama, Strength Distribution of Titania Ceramics after High Voltage Screening, J. Mater. S c i . , a (1996) 3419-3425.
232
T. Tanaka and A. Kishimoto, "Reliability improvement of titania ceramics with surface flaw through high voltage screening, Kor. J. Ceram. SOC., 5[4](1999)386-389. 5) A.Kishimoto, K. Koumoto and H. Yanagida, Mechanical and Dielectric Failure of BaTi03 Ceramics, J. Mater. Sci., 24 [2] (1989) 698-702. 6) A. Kishimoto, K. Koumoto and H. Yanagida, Comparison of Mechanical and Dielectric Strength Distributions for Variously Surface-Finished Titanium Dioxide Ceramics, J. Am. Ceram. S O C . , [8] ~ (1989) 1373-1376. 7) Y. Nakamura, M. Suzuki, N. Motohira, A. Kishimoto and H. Yanagida, Comparison Between the Mechanical and Dielectric Strength Distributions for Hardened Gypsum, J. Mater. Sci., 32 [I] (1997) 11 5-1 18. 8) A. Kishimoto, M. Nameki, K. Koumoto and H. Yanagida, Microstructure dependence of mechanical and dielectric strength distributions: I , porosity, Eng. Fracture Mechanics, 40 [4/5] (1 99 I ) 927-930. 9) N. Yoshimura, A. [to, J. Funaki, and T. Ogasawara, Effect of Pore on Electrical Conduction and Dielectric Breakdown for Dielectric Ceramics, Trans of. IEE. Japan-A 108 [4] (1988) 155-161. 10) Y.Matsuo, T. Ogasawara, S.Kimura and E. Yasuda, Statistical Analysis of the Effect of Surface Grinding on the Strength of Alumina Using Weibull's Multi-Modal Function, J. Mater. Sci., 22 (1987) 1482-88. 11) T. Asokan and T. S. Sudarshan, Dependence of the Surface Flashover Properties of Alumina on Polishing Abrasive Parameters, IEEE, Trans, Elect, Insl., 28 [4] (1993) 3419-3425. 12) E. K. Beauchamp, Effect on Microstructure on Pulse Electrical Strength of MgO, J. Am. Ceram. S O C . , ~ [lo] (1971) 484-487. 13)A. G . Evans and S. M. Wiederhorn, Proof Testing of Ceramic Materials. Analytical Basis for Failure Prediction, Int. J. Fract. Mech., 10[3] (1974) 379-392.
4)
IN SITU OBSERVATION OF TENSION AND CYCLIC FATIGUE DAMAGE IN Hi-Nicalon FIBEWSiC COMPOSITE Y. Kaneko", T. Mamiya**, M. Mizuno***, SJ. Zhu*, Y. Kagawa**, Y. Ochi*. (*) The University of Electro-Communications,Tokyo, Japan (**) Institute of Industrial Sciences, University of Tokyo, Tokyo, Japan (***) Japan Fine Ceramics Center, Nagoya, Japan
ABSTRACT In situ observation of cyclic fatigue crack initiation and propagation of Hi-NicalonTMfibers reinforced S i c composite at room temperature has been carried out by scanning electron microscopy. Fatigue crack initiates at the pores between fiber bundles. Delamination at the interfaces between longitudinal fiber bundles and transverse fiber bundles is the main damage process. This is because the inconsistent deformation causes shear stress. Due to delamination, the transverse crack deflects to the interfaces. As a result, crack propagation is retarded. There is no single crack propagation. Therefore, it is impossible to use fracture mechanics parameter to characterize crack propagation.
INTRODUCTION Ceramic matrix composites (CMC) have been designed for use at high temperature because they have good fracture toughness and other performances. In high performance turbine engines, high temperature components such as after-burner vanes, combustion chambers and even turbine blades may eventually be manufactured from CMCs. CMCs may also take an important role as fascinating devices for hot fractures [l-31. NicalonTM fiber/SiC composite has many good mechanical properties, but it has the low creep resistance owing to a viscous flow at temperatures as low as 1000-1200 "C. The presence of SiCxOy amorphous phase in NicalonTMfibers is responsible for Young's modulus degradation and the low creep resistance owing to a viscous flow at temperatures as low as 1000-1200 "C. The elimination of SiCxOy from the fibers by electron irradiation in vacuum instead of curing in air can improve creep resistance. The modified NicalonTMfibers are called Hi-NicalonTMfibers (Nippon Carbon, Tokyo, Japan). Hi-NicalonTMfibers with a very low oxygen content ( c 1 wt%.) exhibit improved thermal stability [4-71. Carbon coating layer in SiC/SiC leads to low oxidation resistance at high temperatures in air. A glassforming, boron-based particulate can be added to the matrix that reacts with oxygen to produce a sealant glass that inhibits oxidation. This technology is applied to SiC/SiC composites. The modified SiC/SiC is called enhanced SiC/SiC composite. To increase the
mechanical properties of SiC/SiC, Hi-NicalonTMfibers are used to reinforce the enhanced Sic matrix. Fatigue is responsible for the majority of failure of structural components. The fatigue mechanisms in composite materials are more complex and involved a multitude of spatially distributed and interacting mechanisms[8-11]. SiC/SiC composites exhibited a definite fatigue limit in classical S-N curves, hysteresis loops and the reduction of Young's modulus with number of loading cycles[12-131. It was reported that the fatigue limit was 65-80% of tensile strength. Although cyclic fatigue of CMCs was investigated and possible fatigue mechanisms were proposed, in situ observation of fatigue crack initiation and propagation in Sic-fiber/SiC composites is little [9]. The purpose of this paper is not only to present cyclic fatigue data, but also to describe damage evolution of cyclic fatigue of Hi-NicalonTM fibers reinforced S i c composite at room temperature. In situ observation of cyclic fatigue crack initiation and propagation has been carried out by scanning electron microscopy.
MATERIALS AND EXPERIMENTAL PROCEDURE The composites used in the investigation were processed by chemical vapor infiltration (CVI) of S i c into plane woven 0"/90" Hi-NicalonTMfibers preformed (Honeywell Advanced Composites Inc., Newark, DE, USA). Before the infiltration the S i c fibers were coated with carbon by chemical vapor deposition (CVD) to decrease interface bonding between fibers and the matrix. This coating acts to raise strength and toughness of the composites [14-151. The composite contained 40 vol% S i c fibers and about 10% porosity. The average diameter of fibers is 14pm and each bundle consists of 500 bundles. The testing specimen was machined from the panels using diamond-cutting tools. The shape and dimensions of the specimens for monotonic-tension and cyclicfatigue tests at room temperature (RT) are shown in Fig.1. The specimen has two types of orientations in fiber woven structure. One is called parallel orientation (Fig.1 (a)), and the other is perpendicular orientation (Fig.1 (b)). A difference between two kinds of specimens is the different number of layers. Perpendicular specimen has about 3 layers and parallel specimen has about 8 layers.
233
The optical micrograph of transverse cross section of Hi-NicalonTM / S i c composite that shows the fiber distribution, porosity, boron-based particulate and matrix are shown in Fig.2. In this micrograph the light regions correspond to the CVI S i c matrix, the fibers of from side to side are longitudinal bundle (0" direction), the oval shapes are fibers in transverse bundle (90" direction), and large black regions are pores between bundles. The pores are connected to each other with these matrix-rich regions. Such pores form a continuous network that facilitates matrix deposition during the CVI process [16]. The tensile and fatigue tests were performed using a floor-type Shimadzu hydraulic-servo fatigue testing system (Shimadzu Corp., Kyoto, Japan) combined with a scanning electron microscopy (SEM). Fig.3 shows configuration of the testing machine with a SEM. In site SEM observation of fatigue crack propagation on these specimens were carried out by this testing system at RT. The monotonic tensile tests were conducted in vacuum at RT under lading control. Strain was measured by strain gages. The tension-tension fatigue tests were carried out under load control with a sinusoidal loading frequency of lOHz in vacuum at RT. The stress ratio (R) was 0.1. After fracture, the specimens were examined by SEM.
Fig.2. Microstructure of Hi-NicalonTMfiber / S i c composite.
Fig.3. Configuration of the fatigue testing machine with a scanning electron microscopy.
RESULTS AND DISCUSSION
Fig.1. Woven fiber structure of 2D plain-woven HiNicalonTMfiber/SiC composite for monotonic tension and cyclic fatigue tests at RT. (a) parallel specimen, (b) perpendicular specimen. (c) shape and dimensions of perpendicular specimen.
234
Monotonic tensile behavior The stress versus strain curves of Hi-NicalonTM/Sic composites at room temperature are shown in Fig.4. The curves indicate a linear elastic behavior up to a proportional limit of 70MPa, and this stress is -30% of the average ultimate tensile strength (UTS) in parallel specimen and -35% of the average UTS in perpendicular specimen. Each exhibited nonlinearly at high stress because of matrix crack. Matrix crack propagates between layers because of delamination of layers. The number of layers in perpendicular specimens are more than those in parallel specimens, so that the tensile stress of perpendicular specimen is lower than that of parallel specimen. However the failure strain varied rather widely, from 0.2-0.7%. This variation is probably associated with effect of manufacturing stage. i.e. the fiber damage may occur either during the weaving or during another stage of the composites manufacture [ 171, hence the pores and 2D plain-woven architecture may certainly lead to a nonuniform stress and strain field under an applied load. The modulus calculated from the linear portion of the curves are 260GPa in two kinds of specimens. The average tensile stresses, fracture strain and Young's modulus are shown in Table I.
250 300
1
2 200 n
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100
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50
0.1
0
0.2
0.4
0.3
0.5
Fig.6. Crack deflection and debonding of the interfaces between longitudinal fiber and matrix.
0.6
n 0.0351
Fig.4. Monotonic tensile stress versus strain of HiNicalonTM fiber/SiC composite in vacuum at RT.
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arallel Hi-NicalonTM /Sic
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Table I. Tensile result of Hi-NicalonTM fiber/SiC comoosite Tensile strength
20pm
0
0.45 (+0.25/ 0.7.5)
205 (+30/-33)
50
100
150
200
250
Stress (MPa) Fig.7. Variation of the permanent strain versus stress of Hi-NicalonTMfiber/SiC composite.
260
Loading-unloading behavior The loading-unloading curves of Hi-NicalonTM/SiC composites at room temperature are shown in Fig.5 (a: parallel specimen, b: perpendicular specimen). From a reduction in modulus, as manifested in the slopes of hysteresis loops, the matrix cracking may initiate at about 70MPa. The loops appear wide in relation to those measured on other CMCs [181, implying that interfacial debonding and sliding occur during loading. Indeed, examination of fractured specimens reveals that the crack deflection and debonding of the interfaces between longitudinal fiber and matrix (Fig.6).
+parallel -0- perpendicular
0.8
0 0
50
100
150
200
250
Stress (MPa) Fig.8. Young's modulus reduction versus stress of HiNicalonTMfiber/SiC composite.
0
0.1
0.2
0.3
Strain (%)
0.4
0.5
0
0.1
0.2
0.3
Strain (%)
0.4
Fig.5. Stress versus strain curves and hysteresis loops of Hi-NicalonTMfiber/SiC composite in vacuum at RT. (a) parallel specimen, (b) perpendicular specimen
0.5
Cyclic fatigue behavior The cyclic fatigue life in vacuum at RT is shown Fig.9. The cycles to failure increase with decrease of maximum stress and the fatigue life of perpendicular specimen is lower than that of parallel specimen. The fatigue life for parallel specimen at 10' cycles is 200MPa which is 80 % of UTS, and for perpendicular specimen at lo6cycles is 140MPa which is 70 % of UTS. The fatigue limit of the composite at RT is much higher than the stress of the matrix cracking (about 70MPa). This means that the composites can avoid the unsteady propagation of the matrix cracks induced by the first loading during cyclic fatigue at the stress of the fatigue limit. 235
In site observation of fatigue damage shows that most of cracks initiate at the sharp corners of large pores at the crossover points of the fiber weaving (Fig.10) and matrix cracks do not appear at regular intervals likes other composites [19]. Figs.11 and 1 2 show in site observation of cyclic propagation processes in perpendicular specimens. Fig.11 (a) shows the original state. Because the maximum stress (145MPa) is higher than the proportional limit of stress, matrix cracks initiate at the pores (Fig.11 (b), N=1@). Scattering pores in rich matrix religions may make the S-N dates vary widely. Crack propagation from pores makes interface delaminate between layers (Fig.11 (c), N=105) and because of connection of cracks this composite fractures (Fig.11 (d), N=1.4x105).
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(b) N=103: Cracks initiate at the pores perpendicular 7 parallel tensile test for perpendicular . . A tensile test for parallel
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(a) N=O: Original state
lo2
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Cycles to failure Fig.9. Maximum stress versus cycles to failure of HiNicalonTMfiber/SiC composite in vacuum at RT.
(c) N=105: Delamination cracks between layers
Fig.10. Crack initiates at the pores at the crossover point of the fiber weaving of parallel specimen.
236
(d) N p 1 . 4 ~ 1 0 ~ Fig.11. Microscopic damage evaluation of perpendicular specimen under fatigue. (am,,=145MPa, R=O. 1, f=lOHz at RT in vacuum)
(a) N=O: Original state (Crack initiation by applying mean stress of 88MPa)
(b) N=10': Delaminaiton propagation
(c) N=104: Connection of transverse crack with delamination crack
Fig.12 (a) shows the original state, because of mean stress of 88MPa crack initiates. Delamination propagation between longitudinal and transverse fiber bundles under fatigue (Fig.12 (b), N=10"), Connection of transverse crack with delamination crack (Fig.12 (c), N=105) and fractures (Fig.12 (d), N=1.9x104). Under fatigue delaminations are observed (Fig.13) and they have three types. i.e. they initiate between (1) each fiber, (2) longitudinal and transverse fiber bundles, (3)fabric layers. From this observation, (3) type of delamination is main fracture mechanisms shown in Fig.14. This is reason that loading-unloading under fatigue severs fabric layers by the extend and contact of bending longitudinal fiber. This may be an original reason different from 3D-fiber/matrix composites. So perpendicular specimen has much fiber bundles and fabric layers, that fatigue life of perpendicular specimen is lower than that of parallel specimen. Delaminations in this composite easily occur, consequently the crack propagation is retarded, there is no single crack propagation and the unsteady of crack propagation are many observed at unbroken test piece at 10'. So the interface debonding strength in this composite is not strong between the fiber and matrix, that this fiber coating of Carbon is good property. As a result, the main fatigue mechanisms of the composites are interface debonding between layers and crack initiation at pores.
Fig.13. Delaminatrin at the interfaces between longitudinal (0") fiber and transverse (90") fiber bundles of perpendicular specimen.
90"bundle \
(d) N,=1.9x104 I Fig.12. Microscopic damage evaluation perpendicular specimen under fatigue. (am,,=160MPa, R=0.1, f=lOHz at RT in vacuum)
0" bundle
delanhated + -. layers loading-unloading Fig.14. Delamination model between fabric layers under fatigue.
237
CONCLUSION In situ observation of cyclic fatigue propagation of Hi-NicalonTMfiber/SiC composite has been carried out by scanning electronic microscopy attached to servo type fatigue testing machine. The following conclusions were obtained. 1) The tensile strength of Hi-NicalonTM/SiCcomposite at RT is 205-250MPa and Young’s modulus is 260GPa. 2) The tensile strength and fatigue life in parallel specimens are higher than those of perpendicular specimens. 3) The fatigue life is governed by delamination crack initiation and propagation between fabric layers. 4) The number of layers in perpendicular specimens is more than those in parallel specimens. This is the reason for the lower tensile strength and fatigue strength in perpendicular specimens.
REFERENCES ( 1)
(2)
(3)
(4)
(5)
(6 )
(7)
238
R. John, L.P. Zawada and J.L. Kroupa, Stress Due to Temperature Gradients in CeramicMatrix-Composite Aerospace Components, J. Am. Ceram. SOC.,82 [ 11 (1999) 161-168. T. Ishikawa, K. Bansaku, N. Watanebe, Y. Nomura, M. Shibuya and T. Hirokawa, Experimental Stress/Strain behavior of SiCMatrix Composites reinforced with Si-Ti-C-0 fibers and estimation of matrix elastic modulus, Comp. Sci. and Tech., 58 (1998) 51-63. W.R. Fohey, J.M. Battison, T.A. Nielsen and S. Hastings, Ceramic Composites Turbine Engine Component Evaluation, Ceram. Eng. Sci. Proc., 16 [4] (1995) 459-466. D.J. Pysher, K.C. Goretta, R.S. Hodder Jr. and R.E. Tressler, Strengths of Ceramic Fibers at Elevated Temperatures, J. Am. Ceram. SOC.,72 [2] (1989) 284-288. M. Takeda, Y. Imai, H. Ichikawa, T. Ishikawa, T. Seguchi and K. Okumura, Thermomechanical Analysis of the Low Oxygen Silicon Carbide Fibers Derived from Polycarbosilane, Ceram. Eng. Sci. Proc., 14 [7-81 (1993) 540-547. M. Takeda, J. Sakamoto, A. Saeki, Y. Imai and H. Ichikawa, High performance Silicon carbide Fiber Hi-Nicalon for ceramic Matrix Composites, Ceram. Eng. Sci. Proc., 16 [4] (1995) 37-44. C. Vahlas, Thermodynamic Approach to the Oxidation of Hi-NicalonTMFiber, J. Am. Ceram. SOC.,82[9] (1999) 2514-2516.
(8)
(9)
( 1 0)
( 1 1)
(12)
( 1 3)
(14)
(1 5 )
(1 6)
( 1 7)
( 1 8)
( 19)
L.P. Zawada and L.M. Butkus, Tensile and Fatigue Behavior of Silicon Carbide FiberReinforced Ahminosilicate Glass, J. Am. Ceram., 74 [ l l ] (1991) 2851-2585. S.J. Zhu, Y. Kagawa, M. Mizuno, S.Q. Guo, Y. Nagano and H. Kaya, In site observation of cyclic fatigue crack propagation of SiCfiber/SiC composite at room temperature, Materials Sci. and Eng., A 220 (1996) 100-108. A.G. Evans, F.W. Zok and R.M. McMeeking, Fatigue of ceramic matrix composites, Acta metall.mater. Vo1.43, No.3 (1995)859-875. S.F. Shuler, J.W. Holmes and X. Wu, Influence of Loading Frequency on the RoomTemperature Fatigue of a Carbon-Fiber/SiCMatrix Composite, J. Am. Ceram. SOC.,76 [9] (1993) 2327-2336. D. Rouby and P. Reynaud, Fatigue behavior related to interface modification during load cycling in ceramic-matrix fiber composites, Composites Sci. and Tech., 48 (1993) 109-118. M. Mizuno, S. Zhu, Y. Nagano, Y. Sakaida, Y,Kagawa and M.Watanabe, Cyclic-Fatigue Behavior of SiC/SiC Composites at Room and High temperatures, J. Am. Ceram. SOC.,79 [12] (1996) 3065-3077. F. Rebillat, J. Lamon and R. Naslain, Properties of Multilayered Interphases in SiC/SiC Chemical-Vapor-Infiltrated Composites with “Weak” and “Strong” Interfaces, J. Am. Ceram. SOC.,81 [9] (1998) 2315-2126. S. Bertrand, P. Forio, R. Pailler and J. Lamon, Hi-Nicalon/SiC Minicomposites with (Pyrocarbon/SiC), Nanoscale Multilayered Interphases, J. Am. Ceram. SOC.,82 [9] (1999) 2465-2473. M. Takeda, Y. Kagawa, S. Mitsuno, Y. Imai and H. Ichikawa, Strength of a Hi-NicalonTM /Silicon-Carbide-Matrix Composite Fabricated by the Multiple Polymer Infiltration-Pyrolysis Process, J. Am. Ceram. SOC., 82 [6] (1999) 1579-1581. A.J. Eckel and R.C. Bradt, Strength Distribution of Reinforced Fibers in a Nicalon Fiber/Chemically Vapor Infiltrated Silicon Carbide Matrix Composite, J. Am. Ceram. SOC., 72 [3] (1989) 455-458. J.M. Domergue, F.E. Heredia and A.G. Evans, Hysteresis Loops and the Inelastic Deformation of 0/90 Ceramic Matrix Composites, J. Am. Ceram. SOC.79 [l] (1996) 161-170. W. Kuo and T. Chou, Multiple Cracking of Unidirectional and Cross-Ply Ceramic Matrix Composites, J. Am. Ceram. SOC.,78 [3] (1995) 745-755.
MULTI-AXIAL STRENGTH DATA FOR Al2O3-AND MgO-Zr02-CERAMICS S. Kruger, T. Kentschke*, H.-J. Barth
Technical University Clausthal, Institute for Tribology and Machine Kinetics, D-38678 Clausthal-Zellerfeld,Germany
ABSTRACT
EXPERIMENTAL SET-UP
Multi-axial strength tests were carried out using tubular specimen made from A1203and MgO-Zr02. For biaxial tensile stresses the effect of the multi-axial stress state is negligible because of the large scattering of material strength. Therefore the computation of ceramic components can be simplified using the maximum principal stress hypothesis. A multi-axial stress hypothesis is not required. Under tension-compression the tensile strength decreases with increasing compression load, if the compression is higher than one third of the uniaxial compression strength. In this case a multiaxial strength effect exists and has to be taken into account when computing ceramic components.
The highly limited ductility of ceramics makes the determination of strength data difficult. Even small errors in load transfer can cause uncontrolled stress intensities and may lead to incorrect measurements. In the frame of a European Research Project funded by the European Commission a testing machine was developed, that allows multi-axial testing of tubular specimen (2). The gripping system of this machine, see figure 1, uses a conical bronze ring (3) to assure a safe friction contact with the specimen ( 5 ) . The vertical pistons (1) and (9) apply the axial loads (compression and tension). The internal pressure is applied by the inner piston (1 0). I
(10) piston internal
INTRODUCTION Multi-axial stress conditions are present almost in each component. For a computation a multi-axial stress hypothesis is required to compare the multi-axial stress state with uniaxial strength data. A large number of such hypotheses are known for ceramics in literature, but the differences are enormous, especially when compressive stresses have to be taken into account. Furthermore the validity is mostly unclear, such that the engineer has finally no possibility to choosing the “right” hypotheses for a certain application. It is the intention of the authors to supply multiaxial strength data of different ceramic materials and to make recommendations for practical computation of ceramic components. A theoretical analysis of the data leading to a new multi-axial strength hypothesis has been done in (1).
(1) piston axial
(2) sphencal plate
compression (8) protection shell
(6) strain gauges
figure 1: gripping system Three strain gauges ( 6 ) are located on the outer perimeter of the specimen with an angle of 120” between them. This allows measuring the bending stresses during the gripping and testing procedure. The geometry of the tubular specimen was designed with the help of Finite Element Method and enables tests in all quadrants of the principal stress field. Figure 2 provides an overview of typical ceramic testing methods using the principal stress diagram.
239
GI;CJ2 :principal stresses p. :extemalpresarre
pi
:intemlpnsSure
lube test
. I
344-pint bending
61
tesl
tube test
figure 2: ceramic testing procedures
To assure, that the specimen fails under all loading case in a desired and reproducible way, numerous parameters had to be taken into account. Parameters were the manufacturing process, the gripping loads, the application of strain gauges, the local risk of rupture and the total probability of failure. The result is a tubular specimen with a- reduced wall thickness in the middle (fracture zone), see figure 3.
TESTING PROCEDURES Table 1 illustrates the number of tests per load case and material (A1203,trade name F99.7, and MgO-Zr02, trade name FZM, manufacturer FRIATEC GmbH (4)). table 1: test plan
* 10 tests per manufaduringspdmen charge
figure 3: specimen geometry The risk of rupture reaches its maximum at the fracture zone at the inner surface for all load combinations. Therefore, the specimen is suitable for multi-axial tests.
240
The tests for instantaneous failure are performed with a maximum load rate of 50 MPds. This led to durations of 4 s (A1203)and 10 s (ZrOz) for tensile tests and 40 s (Zr02) for axial compression tests. The determined strength data were interpreted as the inert strength. For technical reasons, the axial compression strength of A1203 (3700 MPa) could not be reached for the specimen. The testing machine can only reach a maximum stress of 2300MPa in the axial compression test.
The step-load-procedure according to Nadler (3) was used to investigate the sub-critical crack growth. For uni- and multi-axial tests the stress was set to 65 % of the inert strength when starting the test. Accordingly the average duration of an experiment is 5 h (MgOZr02) and 4 h (A1203)for the tension tests and 7 h for the compression tests (MgO-ZrOZ). For multi-axial tests the combination of loads were chosen in a way that each data point in the principal stress field has almost the same distance from another point. The used multi-axial stress combinations are listed in table 2. The ratio of the combined axial loads and the internal pressure is an indicator for the most injurious stress case leading to a faster fracture of the specimen. table 2: stress combinations for the multi-axial subcritical crack growth tests
IcaJcaxOt I ctarjctanm I
lmaterial lstress
ratio
EXPERIMENTAL DATA INSTANTANEOUS FAILURE The stress-strain-behavior of A1203and MgO-ZrOz (figure 4) shows the different Young's moduli (A1203 410GPa f 13GPa and ZrOz 215GPa k 1GPa) and fracture characteristics of the ceramics. A1203 has an ideal-elastic behavior until rupture starts. MgO-ZrOz shows an increasing stress-strain-proportionin both the tension and the compression quadrant for high loads. The degressive curve results from non-reversible deformation of the specimen, which has been proved in a mechanical hysteresis test (1).
~
I
The maximum fiber stresses (tensile stress tsz for axial tension, tangential stress otiat the internal surface of the specimen for internal pressure) were used for the Weibull-statistical evaluation according to DIN EN 843-5. A comparison of internal pressure tests and tension tests is illustrated in the Weibull diagrams in figure 5a,b, wherein the probability of failure of the ith specimen Pfi = (i-O,5)/N is a function of the characteristic specimen strength t s c h a . The non-dashed lines represent the best approach of the two-parametrical Weibull distribution function to the experimental data using the Maximum-Likelihood-Method. The wide range of the 95 YOconfidence intervals (dashed lines) results from the small number of experiments. Furthermore the experimental data were analyzed assuming the volume (crack starts from defects in material) and the surface (crack starts from defects on surface) model. For calculations according to the volume model the effective stressed volume Vefi is needed. To compare the internal pressure and tension test the average specimen strength (oOV,specimen, Veff,,specimen = 659mm3) was transformed into the characteristic reference element strength (V, = lmm3) with the equation
IN is the Weibull integral.
When surface defects are assumed, the total risk of rupture is calculated as the sum of the local risks of rupture of the internal and external surface. The strength data were referred to a surface element of Oo= 1mm2.The calculated Weibull parameters are shown in table 3. table 3: Weibull parameters
char speclmen strength, 95 X confidence interval[MPa], max. lateral fiber stress Ichar. reference elem strength.
Itension
I
int. pressure
1000 Ah03 MgO-210~
tension
517<557<599
Ivolume model
Ishear
I
nodata
]char reference elem. strength.
Itension
I
458<495<538
lsurface model
Ishear
in1 Pressure
strain E
6 1
-1
tension adjusted Weibull modulus [-I
int pressure shear tension
max load reached
423<458<496
I
I
6M)c629<650
I
615480~755
I 1021~1058~1099I I 630~652<673 I
542483427
885<980<1086
nodata
I 1020~1056~1097I
5.1<8.5<13.8
11.5~20,0<27,9
6.oci3.9<29,6
4.5ci5.7e6.0
no data
12.544.0e73.0
214
449
-2500
figure 4: o-&-diagram of MgO-Zr02 and Al2O3
24 1
2
1
0
4- -1 --c C
-2
-3
-4 5.50
5,75
6,OO
6,25
6,50
IN4
+
......95 Oh confidence interval of tension tests
tension tests
---
internal pressure tests
95 % confidence interval of internal pressure tests
Weibull fit
figure 5: Weibull diagrams a) AI2O3, b) MgO-Zr02 The data were transformed with equation 1 into a theoretical 4-point bending strength (3khm4 and compared with manufacturers’ data. Of course this comparison should be carefully interpreted because of the different material charges, specimen geometry and specimen manufacturing processes. Nevertheless, the calculated data and the manufacturers’ data (4), shown in table 4, correspond very well.
table 4: comparison of measured and manufacturers data (4) this stud
manufacturers data 500MPa
F99.7 350MPa
The maximum tolerable fiber stresses in the internal pressure tests were much higher than in the tension tests. The same behavior could be observed for the characteristic reference element strength using the volume and the surface model. It can be concluded that
242
the reason is not the different stress state in the tests, but the anisotropisms of manufacturing (hot isostatic pressing) or mechanical processing. For shear stress the volume and the surface model lead to equal strength results. The Weibull modulus (m=44) is high. The Weibull distribution function was assumed for shear strength in order to compare compression with tension data. For shear strength, the Weibull parameters of MgO-ZQ are much higher than for normal stress failure. The scattering of the strength is accordifigly smaller. MgO-Zr02 (o& = -4) is more sensitive to shear load than A1203 (odoz>-10). The multi-axial behavior of both ceramic materials is characterized in the figures 6a,b, using the maximum fiber stresses om (axial tension and compression test) and ati(internal pressure test).
-_._---.---.--.-._-_-__-__-
-2
a)
-2
b)
0 6,
IMP4
- - - - - - - - - 95% confidence limits for avera e characteristic reference element strength (63,21% - value) - - - - equivalent stress hypothesis of ieierlein - - - equivalent stress hypothesis of Weibull-Stanley -
symmetry FE-calculatuions
figure 6: multi-axial diagrams a) MgO-Zr02 b) A1203 For both materials, no multi-axial effect can be recognized in the tension-tension quadrant. The specimens failed under the highest tensile stress. The tolerable tangential tensile stress of MgO-ZrOz decreases with increasing axial compression load, when the compression is higher than one third of the uniaxial compression strength. For this range a multi-axial strength effect exits and has to be taken into account in the calculations. This influence could not be obtained for Al2O3because of the limited test load.
SUB-CRITICAL CRACK GROWTH The time parameter n* was determined in the tension and the compression test. The specimen made from A1203 failed during 4 and MgO-Zr02 during 5 tensile load steps (figure 7 and 8). A time effect exits obviously resulting from sub-critical crack growth. In comparison to instantaneous failure, an identical stress leads to higher values of the probability of failure because of the longer load time. For compression of ZrOz a time effect was also detected (figure 9), but smaller than for tensile stress. The specimen failed on the 6" and 7* load step (after 6 and 7 hours). Thus n* is quite smaller.
figure 7: step load tests; A1203,m=l1,6 aov,=226MPa; n*=0,52
,
figure 8: step load tests MgO-Zr02, m=23,8, oOVt 407MPa; n*=0,58
243
and MgO-Zr02 the maximum principle stress hypothesis is sufficient in the tensile stress field. Compressive loads can be neglected as far as the compressive load is below 1/3 of the compressive strength.
REFERENCES 01
--
-----.
h&anhlroum (ihm t = 20 8
- - -,
-
0 5m
5y1
sm
am
?m
7 s
1 -
a0
sm
sm
ma
1mo
w.1
figure 9: step load tests MgO-Zr02 (compression), m=44,0, ro =913MPa; n*=0,45 n* was determined assuming a linear damage accumulation. The total risk of rupture BG is the sum of the risks of rupture per load step i (equation 2):
For each data point the time probability of failure was calculated and compared with the probability of failure for instantaneous failure. n* was determined by minimizing the sum of the square deviance between the both probabilities of failure. To check the congruence between calculation and experimental data n* was used to generate theoretical data points. A good congruence can be obtained (figure 8, 9). Transforming of n* into the crack growth parameter n with the equation
m
.-
nzn*
I
results in a good accordance to values given by the literature (5) (table 5). The conlusion is that Nadler's model is suitable for practical work.
L
n'
0.52 MgO-ZrOz (tension) 0.58 MgO-Zr02 (shear) 0.45
A1203
_____
_________~
lower value average value upper value llerature [5] n1 nm nu n
9
22
53
22 28
41 98
74 162
21 40 nodata
A multi-axial effect for time failure could be discovered for compression-tension stress combinations only. For combined tension-tension stresses the specimen always failed under the higher tensile stress.
SUMMARY For computation of ceramic components the multiaxial strength effect is from less importance. For AI203
244
Kriiger, S.: Ein Beitrag zur praxisgerechten Dimensionierung keramischer Bauteile bei mehrachsigen Beanspruchungen; Dissertation TU Clausthal, Papierflieger (1999), ISBN 3-89720345-6. Kriiger, S.; Barth, H.-J.: Entwicklung einer Priifmaschine fW die keramische Werkstoffpriifung; 5. Kongrefi fiir industrielle Mefi- und Automatisierungstechnik, Hiithing-Verlag (2000). Nadler, P.: Liisung von Problemen des ProofTestes in der Praxis; Mechanische Eigenschaften keramischer Konstruktionswerkstoffe; Hrsg.: G. Grathwohl; DGM Informationsgesellschaft mbH; Oberursel(l993); S.277-288. FRIATEC AG, Informationsbroschiire "Erzeugnisse aus Oxidkeramik", Mannheim (1 999). Wedingen, A; Grathwohl, G: Ermiidung keramischer Werkstoffe; Mechanische Eigenschaften keramischer Werkstoffe; Hrsg.: G. Grathwohl; DGM Informationsgesellschaft mbH; Oberursel (1993); S. 149-160.
111. Design, Modelling and Simulation
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FATIGUE DESIGN AND TESTING OF CERAMIC INTAKE AND EXHAUST VALVES C.M. Sonsino Fraunhofer-Institutefor Structural Durability (LBF) D-64289 Darmstadt/Germany
ABSTRACT Fatigue design of components subjected to cyclic loading require special testing methods, component related material data and an appropriate safety concept. The following contribution presents a newly developed methodology for the fatigue design and testing of highly loaded ceramic components. This methodology is verified by laboratory and in-service testing of Si,N,-intake and exhaust valves.
INTRODUCTION While design relevant data like S-N curve, slope and scatter, notch and mean stress sensitivity, influence of loading mode (axial, bending, torsion), environment (temperature, corrosion) and defects exist since many years for a large amount of metallic materials, fatigue properties and design principles for ceramic materials are almost unknown. The application of structural ceramics for functionally important components like turbine wheels, gudgeon pins or motor valves subjected to repeated loading require the knowledge of corresponding data. In the following, basic relationships about the fatigue behaviour of the structural A1,0,- and Si,N,-ceramics will be presented that were evaluated within former investigations (1, 2). After this, the fatigue design of Si,N,-intake and exhaust valves will be reported.
FATIGUE BEHAVIOUR OF CERAMICS Scatter and slope of S-N curves The normalized evaluation of 350 fatigue tests carried out with almost unnotched and notched specimens under filly reversed (R = omin/ompx = -1) and pulsating (R = 0) four-point-bending at room temperature, 800 and 1200" C resulted in a uniform S-N curve (1, 2) with following characteristics (3), Fig. 1: The slope of the S-N curve k = AlgN/Algo, = 85 (inverse Basquin exponent) is extremely flat and all results are covered within the probability of survival P, = 10 and 90 % by a scatter of stress amplitudes T , = 1 : [oa(lO%) : 0a(90%)] = 1 : 1.40. This scatter is also found for cast aluminium or nodular iron as well as aluminium and steel welds. The particular stress-amplitudes of the different test series were normalized with regard to the appertaining particular fatigue strength at 2.5 . lo6 cycles with P, = 50 %.
Fig. 1 Uniform S-N curve for different A1,0,- and Si,N,-ceramics The facts of the very flat inclination of the S-N curve and the determined scatter band are also confirmed by other until now statistically not evaluated results (4, 5, 6) which fit fairly well into the derived uniform S-N curve. Fig. 2 shows the evaluation of results obtained for a high-strength Si,N,-ceramic developed for turbo-charger turbine wheels and combustion engine valves (5). The fatigue strength of this material can be well covered by the uniform S-N curve. When 1000" C is exceeded the softening of the glas phase decreases the fatigue strength.
Fig. 2 Fatigue behaviour of a Si,N,-ceramic Up to 1000" C the static and fatigue strengths of structural ceramics like AI,O, and Si,N, are not affected by temperature. The very flat slope of the S-N curve and the very high life scatter (not stress scatter) are related to the ,,non
247
forgiving" brittleness of ceramics . This requires a well controlled manufacturing process with not exceeded allowable defect sizes. The determined uniform slope and scatter will remain for any ceramic, but the level of the fatigue strength will depend on the particular material quality. This knowledge facilitates the statistical evaluation even of few tests by covering them with the uniform slope and scatter band. The influence of the defect size on fatigue strength can be estimated by a threshold-stress intensity derived for a given endurable stress amplitude oal, defect depth a1 and width cl and loading mode dependent geometry factor Y(al, c1) :
AK~,,=2.oS1 - , / ~ - Y ( a , , c , ) .
I
I
0.5
.3;
0
4
Pm
4
10' l V I
U
aJ c
h
C
.2
a m m
10'0
2
a Y U
"h
Because of the very flat S-N curves the fatigue strength of ceramic materials cannot be separated into low-, finite- or high-cycle fatigue ranges. Ceramics are also not suitable for a design against variable amplitude loading which allows among others occasional exceedences of the fatigue limit. The fact that for ceramics even a fatigue limit in the classical sense does not exist is not a major problem because for a required fatigue life, e.g. 2 . lo9 cycles for valves, the decrease of fatigue strength is very small due to the shallow inclination of the S-N curve. All these characteristics suggest that for the fatigue design of ceramic components the allowable design stresses must lie below the scatter band in order to exclude failures. The distance of allowable stresses from endurable stresses with P, = 50 %, i.e. safety margin, is for a given scatter only a question of required probability of failure Pf, see section 3. Fracture mechanics The extreme low inclination of the SN-curves is also supported by fracture mechanical investigations on Si,N,-ceramics (7): Compared to conventional metallic materials the difference between threshold and critical stress intensities is very small. The threshold level AKh corresponds to the endurable high-cycle fatigue strength and the critical stress intensity AKc to the ultimate tensile strength. The ratio between A& and AK, is comparable to the ratio between ultimate tensile strength and endurable high-cycle fatigue strength. When AKh is exceeded, SipN, reveals due to its brittleness a very high crack propagation rate, Fig. 3. Only a small exceedance of the threshold level leads under cyclic loading to a total failure.
248
\AK~
i R=Q.J
*
- + 4
Q
For another defect geometry the fatigue strength is then
-
7
4 4 10" -
0.1
-
0
8 0
#
#
0
.
@
rn
0
I
@" 0
0
Material: Si, N, Rb=gOOMPa
I
&
LQdlrg: Compact tension, T=RT
f= 25s.'
m:Gilbert et al
Stress intensity AK, (MPa mfi)
Fig. 3 Crack propagation in Si,N,-ceramics
Mean-stress and notch sensitivity Tensile mean stresses, which may occur e.g. by a preloading, residual stresses, a centrifugal force or thermal stress transients diminish the fatigue strength significantly. For ceramics under pulsating loading (R = 0) only about half of the amount of the endurable stress amplitude for fully-reversed loading (R = -1) was found (1,2), Fig. 4. For ductile steels the decrease by pulsating load is only about 15 %. While ceramics ,,dislike" tensile mean stresses they ,,like" compressive mean stresses due to the defect (crack)-closure. Therefore, in presence of compressive mean stresses the endurable stresses will be much higher than the fatigue strength under fully-reversed loading as indicated in Fig. 4. The knowledge of the relationship between mean stresses and stress amplitudes is necessary for transforming of amplitudes superposed to different mean stresses to amplitudes with a constant R-value, e.g. R = 0. As it can be seen from Fig. 4 the fatigue strength of the almost unnotched state is reduced fully by the theoretical stress concentration.
\
\
\
0 K. = 108
*
K.
P.. N
T
Hmlarrl: A l p , SG Looding: Plan. bending ~
Fig. 4
190
.
50% Z S d
i
8W.C
amplitudes s = [lg(l/T,)] / 2.56 and z is the normalized safety factor which depends on the failure probability Pf, see Table 1. In fatigue design with ceramic materials Pf values smaller than are used. For a scatter T, = 1 : 1.40 and a required theoretical value Pf = 10' j, is 1.98. As long as the the ratio j* between endurable and occuring service stresses, i.e. safety, is higher than j, a failure can be excluded.
\I/
Mean stress-amplitude-diagram and mean-stress sensitivity
In Fig. 5 the comparison of the theoretical notch factors Ktb and fatigue notch factors Kfi reveal that they are within a scatter of + 10 % equal.
I
'
I
-Z(P,-)
5.20
5.61
6.00
6.36
6.71
7.03
Tab. 1 Failure probability Pf and normalized safety factor -z (Pf)
DESIGN EXAMPLE FOR S~~NJ-INTAKE AND EXHAUST VALVES R = -1. 0 R z -1, 0 R . -1 R . -I TV 1063. R -1 07 I PVA I. 8W.C. R . -1 R.4 wm sg1. RT. 9005 sg1. 800.C. R I - l RT. Ri-I HP. SG. SG: SG. TV 963.
2.0 ? .
r
RT
+
8OO.C
Rl.
8OO.C. RT. RT. 8OO.C.
RT R*-1
0
-
8OO.C 0
R.0
1.0 10
15 Th.Oi.11COI
20
nolch faclw K.
Fig. 5 Relation between theoretical and fatigue notch factors for Al2O3- and Si3N4-ceramics This indicates missing stress-gradient dependent volume effects. This maximum notch sensitivity can also be observed for brittle metallic materials. Although the rnaximum stressed volume of the notched specimens is smaller than the volume of the unnotched specimens, the local endurable stresses in the notch are not significantly higher than the endurable stresses of the unnotched state for a given fatigue life. Hence, volume effects known from static loading do not occur under cyclic loading or are at least of minor importance. This fact facilitates the evaluation of local stresses, in the sence that stress gradients and volume effects are not to be considered under cyclic loading.
The increased efforts of investigating the use of ceramic valves, Fig. 6, bases on following advantages compared to steel valves: Weight reduction by about 60 'YO. Improvement of valve dynamics and valve train efficiency by about 20 %. Smoothing of engine torque behaviour by changing the cam profile and increasing the valve acceleration as well as the volumetric effeciency. Saving of fuel consumption by 3 to 4 % and reduction of CO-emission by 20 %. Enaine W D ~ :2.0 t gasoline, 16 valves, p = 7 MPa, n = 6 800 rpm
SAFETY CONCEPT
Fig. 6 Valve train middle section with Si,N,-valves
In design, service dependent occuring stresses must lie below the allowable stresses which are derived from endurable stresses with P, = 50%: oc,dl = 0s (50%)/j,. The safety margin ja depends on the stress scatter T, and the required theoretical probability of failure (3). Under the assumption of a logarithmical Gaussian distribution jo = =, where the standard deviation of the stress
In the following, the design and evaluation methodology applied to ceramic valves within a collaborative work with Adam Ope1 AG and DairnlerChrysler AG will be reported (8, 9). The methodology was verified by fatigue testing of valves, laboratory test runs with six engines and field tests with five vehicles. All engines were gasoline heled and each equipped with 16 valves. The valves
249
manufacturered were gas pressure sintered (tensile bending strength 1200 MPa) and ground (R, < 2 pm). During the quality control all valves with defect sizes a > 40 pm detectable by x-rays and computer tomography were excluded from the investigations. Table 2 compares basic properties of the Si3Ncceramic used for the production of valves and of the substituted valve steel. Properties
Si,N+ (GPSN*)
-
I
X45CrSi93 (AISI W 3 )
-
Thermal elongation coefficient a
I
Thermo-shock parameter R .=R, 4 - P )
I
20 WlmK 1000 "C: I3 W/mK
Io6O
el0
Fig. 8 Analyzed points of a ceramic valve 20-700 "C:
20 "C:
I I
e9 e7
-
1.9.104K-' 20-1000 "C:
Specific thermal capacity h
~~
e4
I
13.24 3.25 g/ccm 7.70 g/ccm I290 310 GPa I I90 GPa 0.27 0.30 20 "C: 20 "C: Bending 1100-1300 MPa 880-1030 MPa 1000 O C : 800 OC: 850-950 MPa 70 MPa I5 Weibull modulus m . 22 Hardness HVlO 1200 1400 Fracture toughness K,c 8-10 MPa mln
Densitiy p Youne's ., modulus E Poisson constant p
The mean stress remains constant when the stationary temperature state is reached. For evaluating a valve, the stress states are calculated by FE for different locations, Fig. 8, i.e. valve seat, transition to the shaft, shaft and fillets for the attachment of the valve spring. Multiaxial stress are evaluated by the principle stress hypothesis valid for brittle materials.
20 "C: 21 WlmK
I
I -
While the highest stresses in the starting stage, at the maximum thermal stress and the stationary stage occur about 2 . lo9 cycles, the maximum stresses due to cold start-stop cycles may occur 3 . 1O4 times at most during a life cycle of an engine. All these service stresses must be assessed using mean-stress-amplitude diagrams which can be only determined experimentally from S-N curves with R = -1 and 0.
*) GPSN = Gas pressure sintered silicon nitride
Tab. 2 Material data of the ceramic Si3N4 and valve steel X45CrSi93 (AISI HNV3) Loading states and local stresses During one revolution of the camshaft a valve is subjected not only to the combustion pressure from fuel ignition but also to cyclic loadings from valve closing and opening and all of them are superposed to thermal stress transients which determine the mean stress level, Fig. 7.
Local endurable fatigue properties The local endurable stresses are determined using the manufactured ceramic valves. Two type of tests are necessary, valve seat bending by pulsating axial loading (R = 0, f = 170s-') of the valve head, and fully-reversed rotatory bending (R = -1, f = 17s') of the valve shaft, Fig. 9. a. Valve seat bending
b. Rotatow shaft bending
T = RT, 300°C,9OO'C
Fig. 7
250
Thermal stress and superposed cyclic stresses at a specific point in a valve
Fig. 9 Loading principles of valves
The failures are produced in the valve seat and in the transition from the head to the shaft, respectively. FEcalculations deliver the relationship between the applied loads and the local stresses in the failure locations. The data obtained for the mentioned critical locations under room temperature, 300 and 900" C, Fig. 10, contain already the influence of the material, maximum detectable defects (a < 40 pm), surface manufacture, residual stresses and stress gradients on the endurable stresses. All test results are covered by the proposed uniform slope and scatter of the S-N curve for ceramics, Fig. 10. The fatigue strength unter 900" C is slightly higher than for the lower temperatures, probably because of the beginning softening of the glass phase and increased ductility.
Diagrams for N = 2 . lo9 cycles
b.
-600
-200
-400
200
0
400 W a 600
Local mean stress q,,
Fig. 1 1 Comparison of Si3N4with valve steel in meanstress-amplitude diagrams Fatigue strength assessment and verification by laboratory engine testing and field tests Tables 3 and 4 compare the occuring local service stresses for the most critical loading states in different locations of intake and exhaust valves with the endurable stresses. a. Maximum stresses for start-stopp-events (N = 3 . 10')
T= 2 0 T T.300 C ' T=900Y
0
100
10'
A
A
0
0
'
'
2
4
.
.
I
68105
.
Vahesafterengmetestng + T=300'C(tntake) x T=900Y(uhaust) .
2
.
.
'
I
4 6 8 1 0 ~2
.
*
.
I
4 6810'
'
2
.
.
.
I
4 6810*
2
Nurrber of cyder to falure N, ...
I
Fig. 10 S-N curves of Si,N, (GPSN)-ceramic valves
e3
I
e4
I
The mean-stress-amplitude diagrams, Fig. 11, are derived from the S-N curves in Fig. 10 using the stress amplitudes for R = -1 and 0 at 3 . lo4 and 2 . lo9 cycles. The comparison in Fig. 11 shows the dependency of steel from temperature while ceramic is not effected.
a.
+67.7 -173.2
f
+31.7 +65.7 +58.2
f
f
f f
67.7 221.9
31.7 65.7 58.2
I I
67.7 24.3
31.7 65.7 58.2
I I
188 188
188
I I
2.8 7.7
188
5.9 2.9 3.2
I88
m
188
b. Maximum stresses during operation (N = 2 . lo9 )
Diagrams for N = 3 * lo4 cycles
e8 e9 el0 -600
-400
-200
0
200
Local mean stress a,
400 W a r n
+65.9 f 65.4 +58.2 f 58.2 -60.0 f 60.0
65.6 58.2 0
I70 170 170
2.6 2.9 m
Tab. 3 Evaluation of calculated occuring stresses on a Si3N4-intakevalve
25 1
p
a. Maximum stresses for start-stopp-events (N = 3 . lo4
pas,
el e2 e3 ._
e4 e5
__
eh
e7 e8 e9 el0
b.
Occuring local equivalent Stress [MPal a,,,* aa +70.3 f 70.3 +8.5 f 5 . 6 +71.6f 71.6 -1 15.6 f 274.8 +6.8 f 13.9 +44.1 .. f 44.1 +37.5 f 37.5 +68.1 f 70.0 +63.7 f 63.1 -60.0 f 60.0 ~~
Transformed Local endurable stress amplistress amplitude a, [MPa] tude asE[MPa] (R=O) (R=O, Ps=50%) 70.3 188 12.0 188 71.6 I88 79.6 188 10.3 188 44.1 188 37.5 188 69.1 188 63.1 188 0 I88
~
Safe& j'=a,/o, (R=o) 2.7 15.7 2.6 2.4 18.2 4.3 5.0 2.7 2.9
+47.3 +37.6 +69.1 +63.8 -60.0
f f f f f
40.9 34.5 69.1 63.6 60.0
I I
43.0 36.5 69. I 63.7 0
I I
170 I70 170 170 I70
5
I. 1 7 -
Fig. 12 Opel-load sequence for tests in a 2.0 1 engine In both cases, no failures or oxidation of the surfaces were observed (9). After the laboratory engine and field tests the valves were fatigue loaded. Most results fitted into the scatter band of the S-N curves in Fig. 10, some lie slightly below the P, = 90 %-line. However, it must be considered that these valves underwent before the valve seat tests already more than 10' cycles in the engines and may be predamaged.
I
3.9 4.6 2.5 2.7
I
Q)
Tab. 4 Evaluation of calculated occuring stresses on a Si,N,-exhaust valve Using the slopes of the mean-stress-amplitude diagrams all stress combinations are transferred to the stress ratio R = 0. All ratios j* between endurable and occuring stress amplitudes exceed the minimum safety margin j, = 1.98 for the failure probability Pf = 10'. This means for the most critical point Pf< lo-'. After this assessment each 16 valves (8 intake and 8 exhaust) were mounted in six 2.0 1 gasoline engines of Adam Ope1 AG and submitted to a standard laboratory loading sequence, Fig. 12. The required 587 repetitions of the sequence is reached after 55 000 engine kilometers in laboratory which corresponds to a damage equivalent driving distance of 300 000 km under usual service conditions. Parallel to this, five 16-valve 2.2 I gasoline vehicles of Daimler-Chrysler AG were equipped and driven all together for 730 000 km, Fig. 13.
252
m
9
Q)
Maximum stresses during operation (N = 2 . lo9
e6 e7 e8 e9 I el0 1
7
-
Vehicle number: 202F405
202F406
202F411
202F415
202F416
Fig. 13 Service performance of ceramic valves in DaimlerChrysler vehicles
CONCLUSIONS Structural AI,O,- and Si,N,-ceramics reveal a fatigue behaviour not comparable to conventional metallic materials, especially with regard to the very flat S-N curves, high mean-stress and notch sensitivity. The very flat S-N curves do not permit the exceedence of a technical fatigue limit, i.e. ceramics are suitable only for high-cycle applications. Despite these features, a fatigue design of functionally important parts like Si,N4-valves is possible when sufficient fatigue strength related to a maximum tolerable defect size can be provided. For this, the concepts of the uniform S-N curve and local stresses, supported by a probabilistic safety consideration are applied. Laboratory engine test runs and field tests with vehicles confirm the suggested methodology.
REFERENCES Sonsino, C.M.: Fatigue Strength of AI,O,- and Si,N,-Ceramics. Fraunhofer-Institut f i r Betriebsfestigkeit (LBF), Darmstadt. Report No. FB - 195 (1992). Buxbaum, 0.;Sonsino, C.M.; Esper, F.J.: Fatigue Design Criteria for Ceramic Components under Cyclic Loading. 1nt.Journ.of Fatigue l6( 1994)No. 4, pp.257-265. Sonsino, C.M.: Methods to Determine Relevant Material Properties for the Fatigue Design of Powder Metallurgy Parts. Powder Met. Int. 16 (1984) No.1, pp. 34 - 38 and 16 (1984)No. 2, pp. 73 - 77. Masuda, M.; Soma, T.; Matsui, M.; Oda, I.: Fatigue of Ceramics (Part I) - Fatigue Behaviour of Sintered Si3N4 under Tension-Compression Cyclic Stress. J. Ceram. SOC.Japan. Int. Ed., 96 (1988), pp. 275280.
NGK Insulators Ltd.: R & D Activities on Sintered Silicon Nitride Materials. Corporate R & D Group, October 1991,Japan. Socie, D. F.: Fatigue Behaviour of Ceramics under Static and Cyclic Loading. In: K.-T. Rie (Ed.), Low Cycle Fatigue and Elasto-Plastic. Behaviour of Materials - 3, Elsevier Appl. Science (1992), pp. 25-30. Gilbert, J.G.; Dauskardt, R.H.; Ritchie, R.O.: Behaviour of Cyclic Fatigue Cracks in Monolithic Silicon Nitride. J. Am. Ceram. SOC.78 (1999, pp. 2291-2300 Sonsino, C.M.; Brandt, U.; Storzel, K.: Theoretical and Experimental Investigations for Designing and Serial Development of Intake and Exhaust Valves from Silicon Nitride. Symposium 2 of Materials Week ‘98, Oct. 12-15, 1998, Munic, Germany
NOTATION depth of a defect or crack width of a defect or crack safety ratio safety factor mean slope of S-N-curve Weibull modulus standard derivation time normalized safety factor Young’modulus load hardness by Vickers theoretical notch factor by bending fatigue notch factor by bending bending moment number of cycles to failure probability of failure probability of survival stress ratio bending strenth thermo shock parameter room temperature temperature stress scatter geometry factor thermal elongation coefficient specific thermal capacity Poisson constant density stress amplitude endurable stress amplitude mean stress
Morgenthaler, K.: Ceramic Valves - a Challenge? Proc. 6th Intern. Symposium on Ceramic Materials & Components for Engines, Oct. 19 - 24, 1997, Arita, Japan.
Acknowledgement The author is indebted to the Federal Ministry for Research and Technology (BMBFT), Bonn, for the financial support, to the Robert Bosch GmbH, Stuttgart, for the research work with the A1203-ceramics and to the Adam Ope1 AG, Riisselsheim, as well as to the DaimlerChrysler AG, Stuttgart, for the research work with the Si3N4-valves.
253
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DESIGN AND TESTING OF A PROTOTYPE SiSiC HEAT EXCHANGER FOR COAL COMBUSTION POWER STATIONS H. R. Maier", K. Himmelstein Rheinisch-Westfalische Technische Hochschule (RWTH) Aachen, Germany
ABSTRACT The presented paper relates to a national founded project (BMWi) which is concerned with the design and testing of a ceramic high temperature heat exchanger for coal fired power stations (PPCCCC-process). The project is carried out in cooperation with the Institut fiir Warme- und Brennstofftechnik (IWBT) of the University of Braunschweig and four industrial partners. The Institut f i r Keramische Komponenten im Maschinenbau of the RWTH Aachen (IKKM) is engaged with the concept approach, computer aided design, with material characterisations and prototype testing. In parallel the IKKM is involved in a European project refemng to the design of a heat exchanger for an alternative coal fired cycle (EFCC-process). The main features are compared with emphasis on the PPCCCCprocess.
It is obvious that pressurised pulverised coal firing concepts will lead to highest efficiencies among all coal fired processes. Different concepts are pursued. Two of the most promising alternatives are the Pressurised Pulverised Coal Combustion Combined Cycle (PPCCCC) process and the Externally Fired Combined Cycle (EFCC) process. Figure 2 shows the circuit diagrams of these concepts with emphasis on the gas turbine and heat exchanger unit.
turbine
compressor
1. INTRODUCTION
Fig. 2a: Circuit diagram of Fig. 2b: Circuit diagram the PPCCCC-process /1/ of the EFCC-process /2/
All predictions referring to the energy consumption postulate that the primary energy source "coal" contributes to hture electricity generation and the existing coal combustion processes have to be improved concerning efficiency and COz-emissionreduction. The efficiencies of different power stations processes are shown in Fig. 1.
In both cases a ceramic high temperature heat exchanger is needed in order to achieve a gas turbine inlet temperature of 1300°C resp. 1500°C. The PPCCCC-process works without a pressure difference between the clean and rough gas side of the heat exchanger and with a maximum inlet temperature of 1400°C. Within the EFCC-process the heat exchanger is used to separate the steam and the gas cycle with a pressure difference of about 15 bar and a maximum inlet temperature of 1600°C provided that a suitable sealing technique is found that resists pressure, temperature and chemical conditions, the gas turbine is driven by clean gas. In contrast to this, within the PPCCCC-process a gas filtering unit has to be integrated to fulfil the high demands concerning the gas cleanness and low ash content at the gas turbine entrance. In this case the heat exchanger is used to serve the gas filtering unit with an inlet temperature of 850 to 900°C. Both heat exchanger concepts are based on tube bundles with the difference, that the rough gas side is in contact with the inside tube surface for the PPCCCC-process and with the outside one for the EFCC-process. Thus different slag removal techniques have to be considered. The paper concentrates on design, material characterisation, FEM analysis of structural reliability and testing operation of a ceramic heat exchanger for the PPCCCC-process and compares some design features of a heat exchanger to the EFCC-process. The thermodynamic layout and the operation of the testing is under the responsibility of IWBT, Braunschweig.
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St*M Cycle 1900 -1950
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200 400 600 800
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1000 1200 1400
Process temperature YC]
Fig. 1: Efficiencies of different power station processes /1/
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compressed air
Economical aspects such as heat exchanger surface, volume weight and costs per kW electrical power output considered seperately.
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2. DESIGN OF THE HEAT EXCHANGER UNIT The design process of the heat exchanger follows the IKKM guidelines referring to the design management, shown in Fig. 3. The overlapping design phases are impacted by a bundle of alternatives and each measure taken has to be carefully evaluated for its interactions. In addition it has to be taken into account, that each ceramic component, module or functional unit is linked to its ceramic or non ceramic environment by - force flow, - heat flow and - mass flow which is illustrated in Fig. 4.
Fig. 3: Design Phases and impacts 13/ On a macro scale the mass flow includes the rough and clean gas side, the leakage between them and the environment and on a micro scale chemical reactions and difision transport mechanisms. Special attention has to be paid to the fact that constrained thermal expansions due to heat flow can be turned into risky changes in force flow and stress distribution.
Fig. 4: Functional unit 131
2.1 Objectives and Requirements With reference to the PPCCCC-process in Fig. 1 the circuit diagram of the testrig, as developed by IWBT, is shown in Fig. 5.
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water
1 2 3
fire box ceramic heat exchanger quench
4 5 6
cyclone filter air preheater
Fig. 5: Circuit diagram of the test rig I41 The coal combustion is simulated with a gas burner of 20kW in combination with a feeding device for coal ash. In the first stage the unit is running at lbar, but the concept definition has to consider the feasibility of 15bar in principle. The sealing requirements within the heat exchanger unit remain comparable, but the metallic pressure vessel has to cope with a pressure level of 1Sbar in the second stage. The thermodynamical layout of the core of the heat exchanger is supplied by IWBT. The concept consists of a bundle of 38 tubes with a length of 1200mm. The length of the tubes is subdivided into six sections by five baffle plates and the tube bundle is surrounded by a liner diameter of 300mm. By these means a cross counter flow principle is achieved. The layout power level of about 3kW corresponds with the following temperatures: - Rough gas inletfoutlet: 1400"C1900"C. - Clean gas inletfoutlet: 80O"Cll300°C This care unit remains unchanged during design optimisation and is characterised in Tab. 1. Table 1:Heat exchanger core characteristics: power level 3kW 20x13x1200mm3 dimension of tubes 40,4mm distance of tubes to centre inner diameter of liner 300mm 1.54 outerlinner diameter of tubes 0.39 inner surface linedouter surface tubes free cross section outside1 11.65 inside of tubes inner surface of tubes I 0.022 [llm] inner volume of liner 1.58 [kWlm2] power levellinner surface of tubes 3.57 [kW/m3] power levelfinner volume of liner The main reauirements for the testrig layout are: compatibie with building situation at IWBT, easy access for measurements, inspection and repair, easy handling due to light weight design, fast thermal response due to low heat capacity, unconstrained thermal expansion between mechanically linked components and defined force, heat and mass flow.
2.2 Concept Definition
2.3 Structural Design Features
Based on the main requirements given above the first concept draft based on conventional brick insulation (chamotte and refractories) has been replaced by dense engineering ceramic liner tubes in combination with a so called superinsulation with ceramic fibre material, Fig. 6. Flexible wrapping and adjustable preforms allow a close and force free surrounding of the liner tubes as well as of the adapters for inspection, measurements and gas supply. The total weight of the A1203 fibre insulation amounts to 34kg only. In addition to weight the diameter of the metallic pressure vessel for the second stage can be reduced from about 1200mm to 460mm with positive impacts on safety and packing density (kW/m3)considerations. In order to reduce residual stresses during processing and thermal stresses during operation and taking into account the given facilities at the manufacturer as well, the ceramic liner has been subdivided into six modules. The length of the modules are identical with the distance between the baMe plates. The design concept is illustrated in Fig. 7.
As a 3-leg table does not wobble, is always in correct position and under defined loads the 3-point-support principle has been applied to - the main support to frame, - the pretension, relief and removal support, - the radial on the tangential positioning of the liner modules and - the axial and radial positioning of the baffle plates (Fig. 9). In order to achieve unconstrained thermal expansion, - the main 3-point support to frame is based on the roller bearing principle (Fig. 8) and takes care about controlled heat losses, - the connectors between liner adapters and pipe system are highly flexible in all directions (Fig. lo), - the connection between heat exchanger tubes and tophottom plate is force limited by a AI2O3-fiber preform, which acts as a predetermined breaking point even under slag infiltration (Fig. 9) and - the liner modules are axially buffered by A1203-fiber preform, which are under defined axial loads. (Fig. 7 and Fig. 9). The A1203-fiberpacks act as sealings as well and the permeability will decrease during operation by infiltration of ash particles and slag.
brick-insulation
fibre-insulation liner fibre insulation
bottom liner segment
steel jacket
fibre pack
frame fixpoint to frame
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refractories charnotte
Fig. 6: Impact of insulation principle on overall layout
frame
Fig. 8: Detail of 3-point main support to frame
Y liner fibre pack
pressure control
AP, fibre pac
Force. heat and mass flow
movable ash container
Fig. 7: Design concept of heat exchanger (fibre insulation not shown)
tube
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Fig. 9: Connection between liner modules, tubes and baMe plates 257
high temperature A@, sleeve allows axial and radial fibre pa&
I---Fig. 12: SiSiC wetted with lignite ash coal, 1400"C, 12h gap forseen for different thermal expansions
Fig. 10: Flexible connector (detail of 1 of Fig. 10)
2.4 Computer Aided Design will follow after material processing and joining alternatives.
3. MATERIAL ALTERNATIVES AND CHARACTERISATION Based on the manufacturing capacity of the industrial partner involved, the ceramic alternatives - Siliconnitridebonded Siliconcarbide (NSiC), - Recrystallised Siliconcarbide (RSiC) and - Silicon infiltrated Siliconcarbide (SiSiC, Si: 20%) have been taken into account. The most promising candidate has been preselected by slag wetting tests with three different hard coal ashes and one brown coal ash. The depth of infiltration and the amount of the corrosion products were determined by Scanning-Electron-Microscopy-analysis. This analysis showed that lignite coal ash is more corrosive to all materials candidates than hard coal ash and that SiSiC demonstrates the best corrosion behaviour against the different coal ashes. NSiC is positioned at the end of the ranking tail. SEM-pictures of SiSiC wetted with hard coal ash and with lignite coal ash for 12h at 1400°C are exemplary shown in Fig. 11 and Fig. 12.
As SiSiC offers additional advantages concerning zero shrinkage, minor risk of residual stresses and insitu joining of SiSiC components, SiSiC has been used for carry on characterisation. In accordance with the real composition at the clean gas side, enhanced gaseous conditions for quick motion exposure tests have been studied /5/. The impact on statistical strength parameters is summarised in Tab.2. Table 2: 4-point bending strength of SiSiC at RT after gaseous exposure tests
I Exposure condition
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SOz, 0 2 , 14OO0C, 12h HCI, Nz, 1400"C, 12h HCI. Nz,02,1400"C, 12h Pre-oxidized, 1200°C. 200h HCI. N2,02,1400"C, 12h
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143 244 252
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130 226 242
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397 458 325
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3,7 5-4 10,6
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The FEM parameters E(T), v(T), a(T), h(T) and cp(T) measured up to 1400°C on untreated samples are in first approximation not effected by the exposure tests. The positive effect of silica layers due to preoxidation treatment is demonstrated in Fig. 13. The strong corrosion effect due to HCl containing atmosphere (which is similar to SO2 containing atmosphere) is illustrated in Fig. 14. As evaluated a reducing atmosphere can become critical but it seems to be possible to counteract by alternating oxidation exposure, e.g. during removal of liquidised slag by means of short time increase of temperature up to 1400°C.
Fig. 11: SiSiC wetted with hard coal ash, 1400"C, 12h
Fig. 13: Preoxidised SiSiC, air, 12OO0C,200h 258
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cantilever - this value increases to 62.5MPa but at changed stress distribution. Taking into account the Weibullparameter m=lO.6 and o,= 325MPa according to Tab. 2 (optimistic approach) the 15MPa peak corresponds with a fracture probability of 0% (
I 3200N
1AOOOC
Fig.14: Corrosion effects of SiSiC, 5%HCI, 95% NZ, 1400"C, 12h
4. PROCESSING AND JOINING ALTERNATIVES The reaction bonding process of SiSiC offers the advantages of zero shrinkage, minor risk of residual stresses, in-situ joining SiSiC components and is therefore highly suitable for manufacturing large and complex shaped modules. It has to be taken into account that the content of free silicon is lost at temperatures between 135OOC and 1400°C (melting point ) and turns the dense structure into an open porous structure with reduced mechanical strength and open for corrosion attack. From this point of view a high amount of free silicon (presently about 20%) is of disadvantage. Although there are SiSiC qualities with less than 10% free silicon commercially available the chosen high contents in the first stage supports the feasibility of large and complex shaped components and modules. The basic shaping processes applied are extrusion (e.g. for tubes) and slip casting (e.g. for liners, adapters, plates and fixation elements). The elements are joint to modules with an identical SiSic-c-sluny in the green stage which turns during reaction bonding to SiSiC. The quality of this in-situ built material fit joints is highly homogeneous, limits residual stresses and reaches strength values almost comparable to those of the components itself. The force and form fit joints used are briefed in chapter 2.
5. COMPUTER AIDED DESIGN An example for FEM based strength and reliability analysis is given for the bottom liner module with the main 3-point support to frame. The worst thermal loading case is assumed during the liquidised slag removal procedure which is plant to reach 140OOC at steady state and can be used for the renewal of the protective silica layer. Using an axial mechanical load of 3200N, an axial temperature difference across the length of module of 30°C and no heat loss to frame via the 3-point supporters, the measured values of E(T), v(T), a(T), h(T) lead to a max first principle stress of 1SMPa, as illustrated in Fig. 15. Assuming an additional heat loss - presented as a temperature difference of 200°C along the 3-point
1370'C
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lmax
Fig. 15: FEM analysis of bottom liner module
6. TEST RIG OPERATION STATUS The test rig as installed at IWBT is shown in Fig. 16. In the first stage the test rig was only fired by gas without any ash additions. To proof the operativeness of the test rig in principle a medium temperature range of about 700"-800°C was chosen. No major leakage occurred, an unconstrained thermal movement of all components was verified and a heat transfer within the heat exchanger has attuned. A number of cycles were driven and different reproducible steady states have been achieved. The maximum heat exchanger entrance temperature was enhanced by degrees. One actual temperature distribution for a steady state is shown exemplary in Tab. 3. Table 3: Temperature distribution for a steady state
HE-entrance secondary HE-outlet secondary
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575°C 505°C 875°C
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In addition first tests with different ash types and varying ash contents have been carried out. To avoid slagging and corrosion the temperature was distinctly below the melting point of the ashes. This procedure was chosen to verify the fbnctionality of the purifier of the filter candles in dependence of the ash content and particle size. Within the next test step the temperature of the ash fired test cycle is to be approached to the melting point of the ashes. Correspondingly slagging and corrosion effects have to be expected and different slag removal strategies as well as material alternatives with improved corrosion resistance have to be taken into account.
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Fig. 17: Submerged entry nozzle, endpiece on MgPSZ+Spinel
Fig. 16: PPCCCC Testrig at IWBT
Fig. 18: Corrosion stability of MgO-PSZ+Spinel against hard coal ash, 1400”C, 12h
7. CONCLUSIONS AND OUTLOOK
8. ACKNOWLEDGEMENTS
Combining all factors of the design phases in Fig. 3 and taking into account the testing results so far it can be concluded, that the heat exchanger unit has a good chance to survive a large range of the targeted test profile. The defined force, heat and mass flow concept did not show any principle weaknesses at short time tests up to a rough gas inlet temperature of 1105°C. The selected SiSiC material with about 20% free silicon will exhibit its limits in reducing atmosphere and in contact with coal ashes at temperatures above about 1250°C. Especially the removal of slag from the inside of the heat exchanger tubes - by liquidising due to short term over heating - has to be proven. Precaution measurements have been taken into account in replacing critical SiSiC components by oxide ceramic ones with higher chemical stability. Based on micro crack tayloring a MgO stabilised ZQISpinel compound has been developed. It shows the highest corrosion resistance of all available ceramics and in addition the thermal shock resistance has been improved by factor 3 concerning the remaining strength after thermal shock treatment. These features have been transferred to large complex shaped components for thin steel slap casting 161., Fig. 17. The corrosion stability against hard coal ashes is illustrated in Fig. 18. The modular approach allows to replace or change critical components (e.g. tubes, upper liner module, removal ash container etc.) without changing the design concept. Comparing the heat exchanger design concept for the PPCCCC-cycle with that of the EFCC-cycle the first one shows clear advantages in structural feasibility. The thermodynamic evolution is left to the experts in this field.
Special thanks for the constructive and successfbl cooperation are adressed to our colleagues at the IWBT at Braunschweig, Prof. R. Leithner and C. Ehlers, as well as to our industrial partners.
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9. REFERENCES 111 Wang, J.; Leithner, R.: “Konzepte und
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Wirkungsgrade kohlegefeuerter Kombianlagen”, Brennstoff W h n e Kraft - BWK Bd. 47 (1999, Nr. 112, S. 11-17 Brite-Euram Project UHTHE/EFCC/B4 “Development and application of design and integration technologies for industrial sub-critical components based on CMC-materials” (BRPRCT97-0426) Maier, H.-R.: IKKM I IPAK SEMINAR 1996 I Institut fiir Prozess- und Anwendungstechnik Keramik an der RWTH Aachen (Hrsg.). Aachen 1996 : Trans Aix Press. -ISBN 3-931814-51-3 BMWi-Projekt “Entwicklung eines keramischen Whnetauschers fiir eine Kombianlage mit Kohlenstaubdruckfeuerung” (Az. 0326835 A/B) Himmelstein, K.; Maier, H.-R.: “High temperature corrosion of SiC-materials in heat exchangers for coal combustion”, Conference and Exhibition of the European Ceramic Society Vol. 1 (British Ceramic Proceedings No. 60) / IOM Communications (Hsrg.), Bd. 718, S. 437-438 ISBN 1 86125 093 2 s. a.: ECerS 1999 Brighton (England), 20.-24.06. I999 Maier, H.-R.: “Porous ceramics functional cavities suitable for system innovation”, 7” International Symposium “Ceramic materials and components for engines”, 19-21 Juni, Goslar, Germany
DESIGN AND TESTING OF CERAMIC COMPONENTS FOR INDUSTRIAL GAS TURBINES M. van Roode*, J.R. Price, 0.Jimenez, N. Miriyala, and S. Gates, Jr. Solar Turbines Incorporated, San Diego, California, USA ABSTRACT
Solar Turbines Incorporated (Solar), under the Ceramic Stationary Gas Turbine program sponsored by the U.S. Department of Energy, has designed and tested uncooled silicon nitride blades and nozzles and continuous fiber-reinforcedceramic composite (CFCC) combustor liners in the Solar Centaur-50s engine. The paper will review the iterative component design and testing in rigs and the engine.
TRIT is planned towards the end of 2000. Summaries of the CSGT program progress have been reported in the literature [l-51.
INTRODUCTION
The Ceramic Stationary Gas Turbine (CSGT) program was initiated by the United States Department of Energy (DOE) Office of Industrial Technologies (OIT) to develop ceramic technologies aimed at the improvement of fuel efficiency and output power, and reduction of NOx and CO emissions in stationary gas turbines by replacement of cooled metal components with uncooled ceramic parts. Solar is the prime contractor on a team, which involves suppliers of ceramic components, test laboratories, industrial end users, and consultants. The three-phase program started in September 1992, and is scheduled to be completed at the end of 2000. Phase I focused on concept/preliminary engine and component design, ceramic materials selection, technical and economic evaluation, and concept assessment. Phase I1 which began in April of 1993, addressed detailed engine and component design, ceramic part procurement, long term materials property testing, component rig testing, engine testing, and destructive and nondestructive component evaluation. Phase Ill which commenced in October of 1996, and which is still ongoing, involves field tests at the cogeneration site of industrial end users. Solar selected the Centaur 50S, a single shaft engine for electrical generation and cogeneration applications. The engine has a two-stage gas producer turbine and single stage power turbine. The engine is being retrofitted with uncooled first stage ceramic blades, uncooled first stage ceramic nozzles, and ceramic combustor liners. Figure I shows the Centaur 50s engine layout with the components targeted for ceramic insertion. A field test, an integral aspect of the program, was originally envisioned to be for 4,000 hrs at a design TRIT of 1 121OC. Because of the high risks involved with the monolithic ceramic components, the field test goals were subsequently relaxed to operation at the baseline TRIT of 10IO°C of the Centaur 50s engine. Short term in-house testing at the design target
Figure I . Solar Centaur 50s Gas Turbine with Components Targeted for Ceramic Insertion CSGT DESIGN-DEVELOPMENT STRATEGY
Based on the industrial gas turbine end user expectations, i.e. a service life between overhauls of 30,000 hrs or more for continuous duty operation, risk mitigation requirements were established which were consistent with ceramic design practice. These included limiting the number of ceramic components, using well established and characterized materials as well as promising new materials with potential costeffective scale-up to production applications, iterative testing with stepwise increases in firing temperatures to the final design TRIT, and minimizing steady state and transient stresses in the ceramic components and adjacent metal structures. Figure 2 is a schematic of the CSGT engine hot section showing the three key ceramic components: the CFCC combustor liner, first stage turbine nozzles and first stage turbine blade. Table 1 summarizes the total number of engine test hours for the three ceramic components.
Figure 2. Schematic of CSGT Engine Hot Section
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Table 1. CSGT Engine Test Summary (Aug. 1995- June 2000) Components Ceramics
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Total Test Time Time at Full Load Max. Time on Single Assembly
I Blades I Nozzles I Combustor Liners* 1 NT164,GN-IO,AS-800 Si3N4 I SN-88 Si3N4 I DLC/H-ACI SiC/SiC (CVI and MI)
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1512 hrs 1483 hrs 1052 hrs (AS-800)
91 hrs
74 hrs 68 hrs
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BFG SiC/SiC (CVI and MI) > 2 1,000 hrs > 16,000 hrs > 10,000 hrs (H-ACI SiC/SiC CVI Outer LinerhlI Inner LinerEBC)
BLADE DEVELOPMENT
Figure 3 shows the evolution of the blade design. The Centaur 50s has 62 cooled MAR-M247 metal blades with an oxidation resistant Pt-aluminide coating. The CSGT blade design has an airfoil shape that is almost identical to that of the metal blade, except for the absence of cooling passages. Two attachment designs were developed. The first design is that of a “dovetail” which requires the use of a compliant layer to act as a buffer for the ceramic/metallic interface. Maximum steady state stresses in a 55” dovetail design were predicted to be 214 MPa at the blade root neck under the platform at a temperature of 682°C. An alternative design concept was a blade with a double-lug support attachment to the disk with a ceramic pin. A compliant layer is not necessary for this “pinned root” design due to the “round-round” interface geometry. This design was abandoned because: ( I ) the maximum steady state stress was higher than for the dovetail blade (283 MPa vs. 214 MPa), (2) manufacturing was very challenging because of tight tolerances and intensive machining, and (3) some of the fabricated parts failed in cold spin proof testing.
While rig and short term engine tests of dovetail blades fabricated from GN- I0 (AlliedSignal Ceramic Components - CC) and NT-164 (Norton Advanced Ceramics) silicon nitrides were successful, these first generation design materials could not provide the required service life based on the CARES/Life life prediction analysis [6]. For the 30,000 hrs of service life required for commercial operation, a set of 62 GN-I0 or NT-I64 blades has a probability of survival (POS) of about 0.10 (based on fast fracture and slow crack growth failure modes), insufficient for acceptable component and engine durability. AS-800 (CC) and SN-281 (Kyocera Industrial Ceramics Corporation) silicon nitrides, which became available later in the program, were selected for second generation component fabrication. Both of these materials have POS values in excess of 0.995 for a set of 62 blades, sufficient for meeting the program life and reliability goals. Details of the blade design and development have been reported [7,8].
Figure 4. Centaur 50s first Stage Disk with Failed Ceramic Blade Residues
Figure 3. Blade Design Evolution in the CSGT Program Various compliant layer systems were evaluated in cyclic attachment testing. A nickel-base 0.13 mm thick compliant layer system withstood 5000 hrs of testing (> 200, 24 hr cycles, from room temperature to 682°C) indicating that it was likely to perform satisfactorily in an engine operating in continuous duty mode. Cold spin testing at 125% of design centrifugal force (CF) load was conducted as a standard proof test for all blades prior to engine testing.
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A milestone 100 hr cyclic endurance engine test with AS-800 first stage blades with two compliant layer systems (Ni-alloy + BN, and Ni-Alloy + Pt) terminated after 58 hrs due to failure of the first stage blades. There was no evidence of foreign object damage (FOD) as a source of component failure initiation. While most of the blades had ruptured above the platforms, only segments of the blade roots remained in alternating slots in one sector of the disk spanning twelve blades (see Figure 4). The more heavily damaged blades were associated with the Ptcoated compliant layer system, which caused the dovetail to pinch in the disk following cyclic engine testing. The attachment system was subsequently modified and the failed compliant layer system was eliminated [8]. These modifications to the blade attachment system were evaluated in a successfil
second 100 hr test, which included 12 cold starts and 15 hot restarts. Final inspection revealed the ceramic and metal components to be in excellent condition. The first CSGT field test was launched in May 1997 at the ARCO Western Energy (currently Texaco) Bakersfield enhanced oil recovery site when the Centaur 50s SoLoNOx engine retrofitted with AS800 (CC) first stage blades and enhanced SiC/SiC (DuPont Lanxide Composites) CFCC liners started operation. Figure 5 shows the CSGT engine being installed in the package at the ARCO site. The CSGT engine accumulated 948 hrs over the duration of the first field test in addition to 104 hrs qualification/ endurance testing at Solar, giving a total test time of 1052 hrs at full load. At 948 hrs of operation the engine shut down as a result of a blade failure, caused in all likelihood because of impact damage from a dislodged metallic locating pin from the inner combustor liner. Failure analysis indicated that all blade failures had occurred above the platform, which excluded failure of the dovetail blade attachment/ compliant layerldisk interface. A second engine field test with the same ceramic component configuration was started in March 1998. After only 352 hrs of field testing at full load, the field test was again terminated due to FOD of the ceramic blade airfoils. Various methods of improving the impact resistance of the silicon nitride airfoils are being examined, improving the materials fracture toughness, redesigning the airfoils to have more robust leading edges, and energy absorbing coatings.
MPa) at the hottest temperature location (1288OCL
FS-414 Metal Nozzle Figure 6. Program
CSGT Straight Vane Nozzle
CSGT Bowed Nozzle
Nozzle Design Evolution in the CSGT
Alternative design concepts considered included airfoils detached from either the inner or the outer shroud, segmentedholted airfoils (radial and axial), segmented shrouds, various pinned attachments and various clamped attachments, as well as a cooled nozzle. The reasonable alternatives required complete redesign of the nozzle airfoil. The airfoil chord was reduced by 50%, and the airfoil was bowed, axially and tangentially, in order to reduce the steady state stresses to the desired level of about 200 MPa. A cantilevered "hook" support system accommodates the ceramic-tometal interface at the first stage diaphragm. There are 42 single vane ceramic nozzle segments compared to 15 two-vane metal nozzles to avoid rotor vibratory issues. A peak transient stress of 242 MPa (25% span, 10% chord location) was predicted 3 s after a fast stop shutdown is initiated. CARES/Life analysis and thermal gradient proof testing indicated that these stress levels were acceptable for long term operation. SN-88 OHexoby SA 0 SN 253
Figure 5. CSGT CentaurSOS Engine Being Installed at ARCO Bakersfield Site
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NOZZLE DEVELOPMENT Figure 6 schematically represents the nozzle design evolution in the CSGT program. The Centaur 50s metal engine has 15 two-vane cooled first stage FS-4 14 metal nozzles, with a WRh-aluminide coating for oxidation protection. The ceramic nozzle development started with a single vane segment, representing one half of the current metal nozzle with tip shoe removed, for ease of ceramic fabrication. The initial nozzle design strategy was to maintain the airfoil profiles of the metal nozzle, thereby maintaining the basic aerodynamic performance of the turbine. However replacing. the metal airfoils with a ceramic vane of the same geometric configuration would result in an unacceptably high steady state stress level (480
Figure 7. Long Term Creep Test Data for Candidate Ceramic Nozzle Materials The POS was 0.9934 for a set of 42 SN-88 nozzles. Figure 7 summarizes creep data for candidate nozzle materials, which formed the basis for selection of SN88 (NGK Insulators, Ltd.) for first generation nozzle fabrication.
Figure 8. SN-88 Nozzle - Chipping at the Inner Attachment Shroud Occurred During Engine Test
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All nozzles were proof tested in a thermal gradient rig and in a mechanical attachment rig prior to engine testing. A full set of 42 SN-88 nozzles was tested for one hour at full load in the CSGT Centaur 50s developmental gas turbine. Some minor chipping, found at the inner and outer shroud areas contacting the metal, suggested a localized excessive contact stress condition (see Figure 8). The ceramic nozzle/metallic attachment was redesigned (see Figure 9) by faceting the diaphragm and removing material at the diaphragm contacting edge locations of each ceramic nozzle. Edge loading on the nozzle inner shroud was reduced and no chipping has been observed in subsequent nozzle engine tests. Details of the ceramic nozzle design have been documented in the literature [7,9].
Figure 9. Redesigned Diaphragm with Faceted Nozzle A planned 100 hour nozzle engine test was initiated which included cold and hot engine restarts and shutdown cycles that progressively increased the severity of stresses and thermal gradients on the nozzle. Borescoping after 68 hours of cyclic engine testing revealed cracking in several airfoils, and the test was terminated for detailed failure investigations. The cracking likely originated from environmentally induced degradation (temperature, water vapor in the hot section) involving removal of the protective silica layer from the nozzle airfoil followed by stresscorrosion type cracking of the silicon nitride. Application of an environmental barrier coating (EBC) will be required to retain nozzle materials integrity in service. An alternative nozzle material, SN-282 silicon nitride (Kyocera Corporation), was selected for the fabrication of the new nozzles, because of its reported improved oxidation and creep resistance compared to SN-88. A 100 hr engine test with the new nozzles is scheduled for the fall of 2000.
of their relatively low fracture toughness, monolithic liners had to be segmented, and designs based on tiles and rings were conceived. Because of the much higher fracture toughness values CFCC liners can be fabricated as integral parts. The materials were tested first in a subscale configuration using 20 cm diameter by 20 cm long cylindrical cans, sequentially for 1, 10, and 100 hrs. SiC/SiC CFCC materials performed best in the subscale combustor testing, completing successfully the 100 hr test which included 200 cycles (15 minutes between -48OOC and -1177°C). Following the test visual inspection, and NDE at Argonne National Laboratory showed no evidence of spallation or delamination. The AI2O3/Al2O3CFCC (McDermott, Inc., formerly Babcock and Wilcox B&W) suffered longitudinal cracks because of matrix shrinkage, and Hexoloy@SA (Carborundum) Sic tiles cracked because of contact stresses between the tiles and the adjacent metallic support structure. Also, these tiles were too thick to accommodate the thermal shocks during the test. NT-230 Sic rings (Norton Advanced Ceramics) were never tested because of fabrication defects. SiC/SiC CFCC materials were selected for the primary zone cylinders. A layer of compliant insulation between the CFCC parts and the metal housing minimizes radial contact stresses. Small axial gaps between the ends of the CFCC parts and the metal housing prevent end loads. All pressure loads are carried by the metal housing, giving the CFCC cylinders an easy mechanical ride at steady state conditions. A variety of SiC/SiC CFCC’s have been tested including 2-D and 3-D CFCCs fabricated by BFGoodrich Aerospace and 2-D CFCC’s fabricated by DuPont Lanxide Composites (DLC, currently Honeywell Advanced Composites, Inc., H-ACI). A set of enhanced SiC/SiC CVI CFCC liners (DLC) is shown in Figure 10. The inner and outer liners are 33 cm and 75 cm in diameter, respectively. They are 20 cm long, and are 0.2 - 0.3 cm thick. The surface of the liners is coated with a CVI Sic seal coat (- 0.25 mm thick).
COMBUSTOR LINER DEVELOPMENT
The ceramic combustor was designed to be interchangeable with the production Centaur 50s dry low-NOx lean-premix (SOLONOX)combustor, which incorporates louver-cooled Hastelloy@ X combustor liners. [lo]. To achieve the desired emissions reductions, only the cylindrical sections in the primary combustor zone were required to be replaced with ceramic parts. The ceramic combustor liner evolution is shown in Figure 10. Initially, both monolithic (silicon carbide) and CFCCs (SiC/SiC, AI2O3/Al2O3) were considered as combustor liner materials. Because
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Figure 10. Combustor Liner Evolution in the CSGT Program Having successfully completed the proof testing in atmospheric and high pressure rigs, full scale SiC/SiC CFCC liners were subsequently tested, either by themselves, or with silicon nitrides blades and
nozzles in the Solar Centaur 50s CSGT test engine at the 1010°C TRIT. Test conditions were representative for a wide range of operating modes in commercial operation including steady state, as well as the transient conditions associated with start-ups, shut-downs and trips. Inspections during and following the test showed the liners generally to be in good condition with no evidence of degradation. Favorable NOx (< 15 ppmv) and CO (< 10 ppmv) emissions were observed. The first 948 hr CSGT ARCO Western Energy field test, mentioned in the blade development section, also included a set of enhanced SiC/SiC (DLC) CFCC liners. Post-test inspection revealed that both liners, but in particular the inner liner, had oxidized. The mechanism of oxidation involves the formation of volatile SiO and Si(0H)x species under the conditions of gas velocity, temperature, total pressure, and water vapor partial pressure in the combustor. The phenomena of accelerated surface recession of silicon based ceramics in the gas turbine combustion environment has been reported [ 111. Several material changes were incorporated in later generations of CFCC liners to further mitigate the effects of environmental degradation, including replacing the ceramic grade Nicalon fibers with HiNicalon, increasing the density of the CFCC liners, and increasing the thickness of the protective seal coat used for the CFCC liners. Also, in later testing the inner liner has been fabricated using the melt infiltration (MI) process, which gives an inherently denser and higher conductivity CFCC than the CVI process. Figure 11 shows a composite photograph of an enhanced SiC/SiC outer liner after 5028 hrs of cumulative exposure in three subsequent field tests and pre-field test qualification tests. The surface oxidation which is clearly visible is highly correlated with the temperature of the liners monitored with thermocouples, which in turn is correlated with the fuel injector impingement locations. The CFCC material degradation from the field test closely matched the degradation encountered in simulated gas turbine environmental exposure testing at Oak Ridge National Laboratory (ORNL) [ 121, which enables the lower cost rig test at ORNL to be used to effectively screen various CFCC liner systems. Additional details of the rig and field testing of the CFCC liners have been presented elsewhere [ 131.
11111 11111
OeSFFloW
Figure 1 1. Composite Photograph of Enhanced SiC/SiC (DLC) CFCC Outer Liner after 5,028 hr Engine Testing In the most recent field test which started in April 1999, environmental barrier coatings (EBCs) to further increase the life of the SiC/SiC CFCC liners, have been incorporated in the liner system. Figure 10
shows a set of SiC/SiC CFCC liners with EBCs developed by Pratt & Whitney and United Technologies Research Center. As of the middle of June 2000, EBC coated liners at the Texaco (ARCO) site had accumulated over 10,000 hrs of service life. Apart from spallation correlated with pre-existing discontinuities in the inner liner and minor surface imperfections in the outer liner, the EBC appeared to have afforded adequate protection. Solar is planning continued field testing at the Texaco site in the 2000-2001 time frame. CFCC liners with EBCs have also been incorporated into a Centaur 50s engine at the Malden Mills Industries textile factory in Lawrence, Massachusetts. The test was initiated in August 1999 and is scheduled to continue for at least 8,000 hours. As of the end of May 2000, the CFCCs at the Malden Mills site had accumulated in excess of 5,000 hrs of service life. A second Centaur 50s engine at the Malden Mills site is scheduled for a CFCC retrofit in August 2000, bringing the number of Centaur 50s engines with CFCC retrofits operating in continuous duty service to three. While the Malden Mills engines operate with many more starts and stops (1 09 as of end of May 2000) than the Texaco engine (38 as of end of May 2000), the EBC apparently is equally protective under both sets of conditions. SECONDARY COMPONENT DEVELOPMENT
Secondary component redesign was made necessary by the following changes relative to the production Centaur 50s engine: (1) increase of TRIT from 1010°C to 1 121"C, (2) first stage nozzle design changed from integral to separate tip shoe to reduce fabrication complexity, and (3) compliant clamping and sealing for the uncooled ceramic nozzle. Materials for the stage 1 disk, the nozzle support housing and diaphragm were changed to accommodate the higher metal temperatures. A more compliant face seal was selected to minimize all mechanical loads imparted to the ceramic nozzle. The design changes have been documented in the literature [7,9]. SUMMARY
Functional ceramic designs were developed for the Solar Centaur 50s first stage blade, nozzle, and combustor liners. The ceramic components were fabricated and evaluated in rigs and engines. AS-800 silicon nitride blades were evaluated in a field test at the Texaco (formerly ARCO Western Energy) industrial cogeneration site in Bakersfield, California. A maximum of 948 hrs of full load field testing on one set of AS-800 blades was accumulated before the test terminated because of impact failure of the blades. Improving the impact resistance of monolithics will be required for rotating components of gas turbines. One set of SN-88 Silicon nitride nozzles was tested in-house for 68 hrs. Material degradation
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associated with the removal of the protective oxide layer was observed. The incorporation of an environmental barrier coating (EBC) will be required for extended engine life. Testing of SiC/SiC CFCC combustor liners has been most successful to-date. Over 2 1,000 hrs of engine testing had been accumulated by midJune, at the Texaco (Bakersfield, California) and Malden Mills (Lawrence, Massachusetts) test sites. The engines with the SiC/SiC CFCC liners have shown significant emissions reduction compared to the all-metal engine. NOx and CO levels below 15 ppmv and 10 ppmv, respectively, have been observed. The longest engine test with one set of SiC/SiC CFCC liners with a protective EBC system had accumulated 10,000 hrs as of mid-June 2000.
ACKNOWLEDGEMENTS The authors wish to acknowledge the support of personnel of the U.S. Department of Energy William P. Parks, Jr., Stephen Waslo, Patricia Hoffman, and Debbie Haught - for the work performed under the Ceramic Stationary Gas Turbine Development program (DOE Contract No. DE-AC02-CE40960).
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REFERENCES 1.
2.
3.
4.
5.
6.
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M. van Roode, W.D. Brentnall, P.F. Norton, and G.P. Pytanowski, “Ceramic Stationary Gas Turbine Development”, ASME paper 93-GT-309, ASME TURBO EXPO, Ohio, USA., May 24-27, 1993. V. Parthasarathy, M. van Roode, J.R. Price, S. Gates, S. Waslo, and P. Hoffman, “Review of Solar’s Ceramic Stationary Gas Turbine Development Program, 6* International Symposium on Ceramic Materials and Components for Engines, Arita, Japan, October 19-23, 1997. pp. 259-264, 1997. M. van Roode, J.R. Price, S. Gates, S. Waslo,and P. Hoffman, “Ceramic Stationary Gas Turbine Development - Opportunities and Challenges “Proceeding of the 7* International Gas Turbine Congress, Kobe, Japan, November 14-19, 1999. V0l.1, pp.65-72 J.R. Price, 0. Jimenez, V. Parthasarathy, N. Miriyala, and D. Leroux, “Ceramic Stationary Gas Turbine Development Program - Seventh Annual Summary”, ASME paper 2000-GT-0075, ASME TURBO EXPO, Munich, Germany, May 8-1 I , 2000. S. Gates, Jr., J. Price, V. Parthasarathy, N. Miriyala, and 0. Jimenez, “Recommended Direction of the Solar/DOE Ceramic Stationary Gas Turbine Program”, ASME paper 2000-GT0076, ASME TURBO EXPO, Munich, Germany, May 8-1 I , 2000. N.N.’Nemeth, L.P. Powers, L.A. Janosik, and J.P. Geykenyesi, “Durability Evaluation of Ceramic
Components Using CARES/LIFE, ASME paper 94-GT-352, ASME TURBO EXPO, The Hague, The Netherlands, June 13-15 , 1994. Transactions of ASME, Journal of Engineeringfor Gas Turbine andPower, 118 (I), 150-158, 1996. 7. P.F. Norton, G.A. Frey, H. Bagheri, A. Fierstein, C. Twardochleb, 0. Jimenez, and A. Saith, “Ceramic Stationary Gas Turbine Development Program - Design and Life Assessment of Ceramic Components”, ASME paper 95-GT-383, ASME TURBO EXPO, Houston, Texas, USA, June 5-8, 1995. 8. 0. Jimenez, J. McClain, B. Edwards, V. Parthasarathy, H. Bagheri, and G. Bolander, “Ceramic Stationary Gas Turbine Program Design and Test of a Ceramic Turbine Blade”, ASME paper 98-GT-529, ASME TURBO EXPO, Stockholm, Sweden June 2-5, 1998. 9. L. Faulder, J. McClain, B. Edwards, and V. Parthasarthy, 1998, “Ceramic Stationary Gas Turbine Development Program - Design and Test of a First Stage Ceramic Nozzle”, ASME paper 98-GT-528, ASME TURBO EXPO, Stockholm, Sweden. June 2-5.1998. 10. K.O. Smith, and’ A. Fahme, “Testing of a Full Scale Low Emissions, Ceramic Gas Turbine Combustor”, ASME paper 97-GT- 156, ASME TURBO EXPO, Orlando, Florida, USA, June 25,1997. 11. J.L. Smialek, R.C. Robinson, E.J. Opila, D.S. Fox, and N.S. Jacobson, “Sic and Si3N4recession due to volatility under combustor conditions”, Adv. Composite Mater, 8[ I], 33-45, 1999. 12. K.L. More, P.F. Tortorelli, M.K. Ferber, L.R. Walker, J.R. Keiser, W.D. Brentnall, N. Miriyala, J.R. Price, 1999, “Exposure of Ceramics and Ceramic Matrix Composites in Simulated and Actual Combustor Environments”, ASME paper 99-GT-292, ASME TURBO EXPO, Indianapolis, Indiana, USA, June 7-10, 1999. 13. N. Miriyala, and J.R. Price, “The Evaluation of CFCC Liners after Field-Engine Testing in a Gas Turbine - II”, ASME paper 2000-GT-0648, ASME TURBO EXPO, Munich, Germany, May 8-1 I , 2000. 14. H.E. Eaton, G.D. Linsey, K.L. More, J.B. Kimmel, J.R. Price, and N. Miriyala, “EBC Protection of SiC/SiC Composite in Gas Turbine Combustion Environment”, ASME paper 2000-GT-063 1, ASME TURBO EXPO, Munich, Germany, May 8- 1 1,2000.
AN INVESTIGATION ON PASTE FLOW IN A PRESS-MOULDED CERAMIC DOME X. Huang*, A. S. Burbidge*’** and S. Blackburn*9** (*). IRC in Materials for High Performance Applications, (**). School of Chemical Engineering, The University of Birmingham, Birmingham, B15 2TT, U.K.
ABSTRACT Paste flow in a press moulded ceramic dome behaves as a non-rheometric flow, for which a fully theoretical analysis is still difficult since ceramic pastes are highly viscous non-Newtonian fluids with a yield stress. Experiments and computer simulations, on the other hand, possess many advantages for the analysis of the behaviour Of such paste flows in the dome’ Flow patterns, which are a reflection of the rheological properties of ceramic paste, are now recognised as in determiningthe performance Of the playing a key final product. This paper presents some preliminary results of an investigation into the flow behaviour of a ceramic paste in a press-moulded ceramic dome using both experiment and computer simulation. In order to clearly display the flow patterns, a paste sample consisting of three different coloured layers was pressed in a dome mould and the pressed sample was analysed. The experimental visualisations show three regimes of paste flow in the pressed-mould dome, which should be characterised by different models. The flow pattern in each regime is highly dependent on the location of the paste front as it advances. The flow patterns obtained using Flow-3D computer simulations are consistent with the experimental results at large strains.
INTRODUCTION Modern manufacturing is being forced to reduce costs and material wastage through energy saving and better processing. Net shape manufacturing addresses these requirements by developing process routes, which give the desired shape with minimum waste and reworking. Ceramic domes are frequently made of piezoelectric materials, which generate a charge when strained and change shape when an electrical field is applied. Thus they can be used as a transducer material. Lead zirconate titanate (PZT) is one of such materials. Items made from PZT ceramics can be formed in a wide variety of shapes by pressing and moulding processes. Typical shapes include discs, plates, curved plates, cylinders and hemispheres, depending on the desired applications. Perfect hemispherical dome structures are difficult to produce in ceramics and have conventionally
been produced by machining. This method wastes energy and material and is consequently expensive. Precision press moulding offers a near net solution to the manufacture of this component. Ceramic domes are used in hemispheres and spheres for omnidirectional hydrophones and may potentially be applied in Hi-fi tweeters. In addition by the application of electromagnetic fields during the moulding process it is possible to influence the functionality of the material, a feature not possible by the machining route. Ceramic paste flow in a press mould is complicated. The deformation mode in the paste is triaxial, where the paste suffers compressive strain in the normal direction and elongational strains in both circumferential and radial directions. The stress acting on the paste in such a deformation consists of six independent components and is a second order tensor quantity. When the principal axes of deformation of the paste are aligned with the chosen coordinate frame, the analysis becomes possible using Benbow and Bridgwater equations [I] as illustrated in the work by Huang and Oliver [2, 3, 41. In the authors’ previous study [ 5 ] , a simple construction of ceramic dome, which was called a “top hat”, was built to approximately model the paste flow in the real dome. In this work, the paste movement in the “top hat” was divided into three steps as shown in figure 2. Each of these steps could then be considered to be a near rheometric flow, with which some approximate solutions could be obtained and could then be superposed to estimate the overall Solution. In the first step, the paste flowed as a squeezed film. Then, there existed both a squeeze film flow and an annular flow. When the paste front reached the rim of the “top hat”, three flows, squeeze film flows both in the top part and the rim, and annular flow in the space between two cylinders occur simultaneously. For each step, a semi-empirical equation was given by the analysis of the paste flow, which gave good agreement with experimental results For paste flow such as those being investigated here, it is more convenient to use visualisation and computer modelling for characterisation. This paper presents some preliminary results obtained using both experimental visualisation and computer simulation. Since the moving front of the paste in the mould has a sharp free surface,
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the code Flow-3D, which incorporates a free surface tracking algorithm, was employed in this investigation. Upper Part Paste
stopping the pressing process at various times and sectioning the paste through the central point it was possible to visualise the flow behaviour during the triaxial deformation. Paste
V
......
‘
Lower female part
Figure 2. A real press mould for making ceramic dome.
(b) step 2
04 0
.
.
.
1
. 2
.
. 3
4
shear rate(l/s)
Figure 3. Rheological property of the PZT paste. (c) step 3 Figure 1. A simplified press mould apparatus showing the paste flow development.
EXPERIMENTAL A commercial lead zirconate titanate (PZT) powder (PZT-SA, Morgan, UK) was used to make a paste by mixing with an aqueous binder system consisting of polyvinyl alcohol (PVA, Nipon Gohsei)/water. Premixed PZT powder and polymer bindedsolvent was milled under high shear stress on a twin-roll mill for 1015 min to form a tape with the thickness of 4 mm. Press forming was carried out by placing the press mould, as shown in figure 2, in a Denison mechanical testing machine. The pressing process was performed at cross head speed of 5 mm/min. The rheological properties of the pastes were investigated using a Rosand capillary rheometer. Three dies (3 mm in diameter) with the die length to die diameter ratio (UD)of 16, 10, and 0.08 were used. The range of the extrusion velocities was 0.25 to 2 mm/min. The experimental data are shown in figure 3. In the flow visualisation work, food colourings were used as dyes and no change in the paste flow properties was caused by the colour change. It was possible to make a “sandwich” of three coloured discs of paste. By
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SIMULATION AND NUMERICAL MODEL The Flow-3D simulation software package has been used for this study since previous studies have demonstrated its capabilities in producing reasonable predictions of macroscopic flow behaviour during mould filling in the casting process [6]. Compared to the conventional mould filling process, paste flow during dome press moulding has similar features, characterised by a sharp free surface. Coupled with the fluid flow and free surface tracking algorithms, the code provides insight into the details of flow patterns in the pressed mould. In this study, the effect of non-Newtonian behaviour of the ceramic paste on the deformation has been taken into account. Assuming the paste to be incompressible, the governing equations and boundary conditions can be written. Paste free surface evolution:
where F is the fractional volume-of-fluid contained in a given cell; ui is the velocity. Ai is the open area fraction associated with the flow in the ith direction and V stands for the open volume fraction of the cell to flow. The right hand side of the equation has been taken to be zero
since no phase change has been assumed to take place during the pressing process. Momentum conservation:
(2) where P is the averaged pressure, and g, and are gravitational acceleration and external body forces, respectively. In equation (2), p denotes apparent shear viscosity of the paste, which can be described by the Carreau equation:
fi
(3) where and are shear viscosity at the highest shear rate and zero shear rate, respectively. 4 and ill stand for parameters relative to time constants for the given material. denotes the shear rate. n represents the flow behaviour index. The above parameters for the ceramic paste used in this work have been experimentally determined, as lu, = 0.05 MPa.s, r(10 = 73.5 MPa.s, 4 = 0.405, ill = 700 s and n = 0.2. The Carreau equation is fitted to the flow curve for the PZT paste as shown in figure 3. The boundary conditions specified in the simulation are that the upper male part is being pressed down with a velocity of 5 d m i n and the lower female part is stationary. The paste flows with no shear stress on the contact surface. The simulation starts when the upper mould just touches the paste, which has been assumed to be placed in the lower mould. A mesh with the minimum grid size of 0.2 mm was employed in this work as shown in figure 4.
shown in figure 5. It can be seen from figure 5(a) that when the top mould touches the paste, the deformation rate of the upper part of the paste is obviously larger than the lower. This deformation process continues as the paste spreads to fill the space in the mould as the upper punch moves down (figures 5(b), (c) and (d)). Ultimately this deformation process results in the upper layer being thinned to such an extent that in places the central layer can be seen at the surface as shown in figure 5 (d). The upper layer is largely displaced to the rim by the operating mechanism. It should be noticed here that there are some cracks in the samples as seen in figure 5 (a), (b) and (c), which are caused by removing the sample from the mould. Also, cutting the samples after the pressing process can sometimes make two different colours of the paste mix on the surface of the cross section as shown in figure 5 (a).
t = 30 s
@) t = 60 s
(c) t = 102 s
(d)t= 138s
(a)
Figure 5. Paste flow visualisation during a pressing process with the downward speed 5 d m i n
Figure 4.The mesh used in the computer,simulation of the press mould.
RESULTS FLOW VISUALISATIONS The materials used were pre-pressed into a sheet and cut into a circular cake of constant thickness. Three discs of the paste were placed together with the central layer being a different colour. Flow visualization results are
FLOW3D SIMULATIONS Figure 6 shows numerical simulation results. It can be seen that the simulation results are consistent with the experimental visualisations at high strains. The deformation patterns observed in experiment are reasonably modelled except at early time intervals as in figures 6(b) and (c), in which the numerical results deviate significantly when compared with the flow visualisation. The reason for this case is unclear, but this result is highly suspect due to the high stress at the contact with the stationary female mould block. After the whole paste is deformed, the simulation result can truly reflect the real paste flow during pressing. The maximum deformation is taking place around the moving punch. Comparing the photograph shown in figure 5(d) with the simulation result in figure 6(g), both show that there exists a recirculation zone in the rim of the mould. However, the existence of such recirculation
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zone does not affect the final product (a sphere combining two domes) since the rim part has to be removed once the pressing process is complete.
(a) t = 0 s
(c) t = 60 s
(b) t = 3 0 s
(d) t = 90 s
(e) t = 108 s
(f)t=126s
(g) t = 144 s
(h) t = 156 s
Figure 6. Computer modelling of the PZT paste flow patterns during a pressing process with the downward speed 5 mndmin.
DISCUSSION The paste deformation first occurs near its top surface, which contacts with the moving upper male part, and the bottom edges, which are in contact with the concave surface of lower female part as shown by the visualisation method. The visualisation and computer simulation results show that the paste gradually deforms to match the surface shape of the upper mould and reduce the gap between the paste and the lower part of the mould. However in the early stages the simulation is unable to reflect this as all the deformation takes place at the top punch and only later is the bottom die filled. This cannot be correct as the point stress where the paste meets the bowl must be exceptionally high.
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When the bottom surface of the paste has not fully touched the whole concave surface of the lower mould, the paste flow is mainly taking place in the upper part of the paste where the paste is squeezed due to the downward movement of the upper mould. Once the bottom of the paste is in full contact with the concave surface of the lower female part, i.e. the gap between the paste and lower mould has been filled by the paste, paste starts to flow along the decreasing gap between the upper and lower mould. This flow pattern can be regarded as an annular flow where the radial dimension and thickness of such a flow are dependent on time. Based on this assumption, the paste movement in the mould can be divided into two flow regimes with regard to the lowest point of the convex surface of the upper mould as shown in figure 7. The flow pattern of the paste section located under this point may be regarded as the squeeze film flow while the motion of the other part above this point may be approximately treated as an annular flow. The latter changes rapidly with time. Also, the deformation of the paste is noticeably three dimensional. It is therefore manifest that it is unlikely to find a general analytical formulation to describe the flow of the paste in such situation because the paste is a nonNewtonian fluid with a yield stress. However, it is interesting to note that the movement of the paste in this area is highly organised as can be seen from the visualisation and simulation results. The upper layers of paste thin and move towards the rim while the lower layers tend to remain in the dome.
Figure 7. The flow regions of the paste in the dome mould. When the paste front enters the rim zone, the flow approximates a squeeze film flow between two parallel plates. Therefore the paste flow in the mould during pressing can be approximately described by these three modes, squeeze film flows in the rim and in the region below the lowest point, and an annular flow in the curved gap above the lowest point. In summary, it is evident that the deformation of the paste in the dome mould during pressing is a process in which the paste is extended in both radial and circumferential directions whereas it is deformed compressibly in the vertical direction. These deformations are dependent on time and the relative movement of the moulds or gap between the upper and lower moulds. Compared with the previous work [ 5 ] , simplifying the real press dome into a “top hat” shape is reasonable to allow approximate theoretical analysis.
Therefore, semi-empirical formulations developed there may be applied to the true dome with some necessary modifications; this will be addressed as part of further study in this project.
simulation although a full analysis remains to be developed.
ACKNOWLEDGEMENT CONCLUSIONS Paste flow in a press moulded ceramic dome has been investigated using both visualisation and computer modelling. It is found that the paste flow in such mould behaves like a non-rheometric flow, for which a fully analytical analysis is difficult since ceramic pastes are highly viscous non-Newtonian fluids with a yield stress. Preliminary experiment and computer simulation results have been presented and conclusions have been drawn as follows: (1) Experimental visualisations indicate the existence of three regimes of paste flow patterns in the pressedmould dome, which should be characterised using different models. (2) Flow patterns, which reflect the rheological properties of ceramic paste, are found to play a key role in determining the performance of final product. (3) The flow pattern in each regime is highly dependent on the location of the paste front and time. (4) The flow patterns obtained using Flow-3D computer simulations are consistent with the observed results at high strains. This demonstrates that it is possible to use computer modelling for the design of the paste forming processes but care in their application is needed at low stains. The rheological behaviour of such ceramic paste in such flows can be analysed using numerical
This work is financially supported by European Regional Development Funding and IRC in Materials for High Performance Applications at The University of Birmingham. The authors would like to thank Drs B. Su and D. Pearce for the help in providing the PVA paste and the pressed mould.
REFERENCES (1) J. Benbow and J. Bridgwater, Paste Flow and Extrusion, Clarendon Press, Oxford, (1993) 3 1-36. (2) X. Huang & D. R. Oliver, Processing properties of ceramic paste in radial flow. British Ceramic Proceedings (1997) 58, 103-112. (3) D. R. Oliver and X. Huang, Squeeze film testing of ceramic pastes. British Ceramic Transactions (accepted for publication in 2000). (4) X. Huang, Rheology of ceramic pastes. PhD Thesis, the University of Birmingham. (1998). ( 5 ) X. Huang, A. S. Burbidge, D. R. Oliver and S. Blackburn, Approximate flow analysis of the paste forming process for a simplified ceramic dome. British Ceramic Transactions (to be submitted). (6) X. Yang, T. Din and J. Campbell, Liquid metal flow in moulds with sprue off-set relative to runner. International Journal of Cast Metal Research, (1998) 11, 1-12.
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DESIGN STANDARD FOR ADVANCED CERAMIC MATERIALS AND COMPONENTS Hiroshi Kawamoto*, Minoru Matsui**, and Hideo Kobayashi***
*, ** Japan Fine Ceramics Center, Nagoya, 456-8587 Japan *** Tokyo Institute of Technology, Tokyo, 152-8552 Japan * Sending Researcher from Toyota Motor Corporation ** Sending Researcher from NGK Insulators, LTD. ABSTRACT Design technologies for advanced ceramics could be required to confirm sufficient strength and reliability of the components under the service conditions. Consequently, developing and applying these design methods, widespread uses of the materials and components should be accelerated in various structural systems. Whereas design methods suitable for advanced ceramics have been steadily making progress for the structural applications, it is extremely meaningful to incorporate the cumulative experiences in the individual field into design standards. However, there is no standard for these design methods in Japan and advanced countries until now. Therefore, constructing the design standards for ensuring reliability of ceramic materials and components has so far been conducted as an important subject for the Technology Infrastructure Project in the Ministry of International Trade and Industry, Japan. This paper introduces the present activities.
materials technologies for the development of high performance materials and the uses of these materials. In systems, it is essential that those components maintain high reliability with an excellent balance between cost and performance. Consequently, design methods are important for making sure of the reliability in components. Objectives ofthis work are as follows : 1) Accumulating and passing on experiences in structural design methods suitable for advanced ceramics systems. 2) Constructing Design Standards for the reliability of ceramics systems as an important subject ;Mass-produced component
Rockerarm seat Shim
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BACKGROUND AND OBJECTIVES Development efforts have been made in automobile technologies for heat engines to improve the power performance and the fuel economy, and so on. It is well recognized that ceramic applications are key to succeed in such advanced heat engines, owing to the mechanical and the thermal properties [ l , 21. Fig. 1 shows ceramic components developed or under research for reciprocating engines. Many ceramic components have been studied for a long time. However, there are not so many mass-produced components now. Few attempts have so far been made at practical applications of ceramic materials and components to present-day automobiles. An example is the amount of Si3N4 for mass-produced components. From FY 90 to FY 93, the amount was increased considerably because of the practical uses of turbocharger rotors. After that, the amount decreased according to a reduction in the numbers of turbo-charged cars. The use of tribological components such as rollers and roller-bushings in fuel injection pumps for diesel engines has recently increased. In Fig. 2, R&D directions for structural ceramics are recommended for the widespread uses. Those are
/ C a m roller
Swirl chamber
Fig. I Ceramic components developed or under research for automobile reciprocating engines
Materialtechnologies balanced between performance, reliability and cost Low cost manufacturing Processing, Machining, Inspection High damage tolerance materials Technologies in uses of materials for fracture resistance
Analyses a simulations
Fig. 2
R&D items expected from now on
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for the Technology Infrastructure Projects in the Japanese government. 3) Accelerating the wide-spread practical uses of ceramic materials and components in various structural systems. Applying the standards to these systems.
ACTIVITIES ON TECHNOLOGY INFRASTRUCTURE PROJECTS Technology infrastructure projects in the Japan Fine Ceramic Center are basically composed of development or construction of design standards, testing methods and data bases. These are supported as important subjects for the Technology Infrastructure Projects in the Japanese government. Members in the committee of design standards are as follows. Most of these members have experience with R&D on advanced ceramic components in private companies or enterprises. Chairperson H. Kobayashi (Tokyo Institute of Technology) Members N. Okabe (Ehime University) T. It0 (Tokai University) Y.Kimura (Kogakuin University) N. Nakazawa (Tsukuba College of Technology) T. Teramae (Tokyo Electric Power Company) A. Suzuki (Ishikawajima-HarimaHeavy Industry) T. Machida (Hitachi) T. Tatsumi (Kawasaki Heavy Industry) M. Hattori (NGK Insulators) S. Nakagawa (Kyocera Corporation) S. Suzuki (Isuzu Ceramics Research Institute) Secretariat M. Matsui (Japan Fine Ceramics Center) H. Kawamoto (Japan Fine Ceramics Center) One of the activities is to construct the design standard according to the program shown in Fig. 3. The design standard has been almost constructed for fundamental design methods and components under steady state service conditions.
I
I
I
1
FY1997 FYl998 FY1999 FY2000 FY2001 FY2002
Gener8l Design Standard Stationary components Design Standard under Particular Service Conditions Contact components Pileup structure components
The design standard consists of following technology elements. 1. scope 2. Design methods Development and design Design by rules Design by tests 3. Materials 4. Analyses of service loads and stresses 5. Design standards Temperature restriction Stress restriction for static fracture Stress restriction for time-dependent fracture Proof tests 6. Evaluation for environmental factors Corrosion and oxidation Acceleration effects depending on service conditions and the evaluation processes 7. Evaluation for machining factors of materials Residual stresses Machining and surface roughness 8. Component designs Design processes Evaluation items Static strength Time-dependent strength Resistance to contact damage Resistance to foreign objects Degradation and corrosion 9. Appendix In the appendix, there are the following explanations for these technology elements. A2. Design methods A3. Materials Comparison between metals and ceramics Kinds of structural ceramics A4. Analyses of service loads and stresses AS. Design standards A6. Evaluation for environmental factors A7. Evaluation for machining factors of materials A8. Component designs Stationary components Rotational components A9. Design examples Automobile components Blades for ceramic gas-turbines Ceramic components for a gas-turbine CGT302
DESIGN METHODS I
Fig. 3 Program for design standard development
274
CONTENTS OF DESIGN STANDARD
I
Design methods are applied to structural designs for advanced monolithic ceramics. Design processes for ceramic components are generally expressed in the flow chart shown in Fig. 4. This flow is divided into design by rules and design by tests. The design by rules has the processes in Fig. 5. The design by tests is classified into component evaluation and assembly evaluation. After these, evaluations of equipment and system follow.
.
to assure the design objective. Proof testing damage must be considered, when the unloading rate is slow. There is a critical unloading rate not to fracture in the unloading process [3]. Ceramics have the property that strength degrades according to applied stress time. Consequently, the strength must be estimated after the proof testing to establish whether that is lower than the proof stress. The critical rate depends on fracture toughness and crack propagation properties of the material, and the proof stress level [2, 31. Table 1
Kinds of tests
I: Equipment &? system Durability evaluation Design by rests
Total evaluation
Fig. 4
Design processes for ceramic components
Objective Portion b Service conditions
I
t
specifications Material, Structure, Component
t Load analysis Stress analysis Stress restriction Environment evaluation
I
Design by tests
Testing methods
Samples Simple components Sonfirmation Models Assemblies tests Number of samples Testing apparatuses Simplified systems Actual systems Loading methods Proof tests Stationary Non-stationary Analyses of results
I
M
I
Objectives
Design qualities
Production qualities
Initial static strength r o o f stress 0, S- t curve
v)
S
N curve
2t, vice
Design for manufacturing
I
I
Time t, Number of cycles N
Trial manufacturing
Fig. 5
a) Stress - strength model under proof test
Processes for design by rules
Design by tests is mainly composed of confirmation tests and proof tests, as shown in Table 1. Testing samples should be selected for simple components or models, and for those components fixed in assemblies. As the testing apparatuses, simplified systems and actual systems are employed for strength and durability tests. Loading methods are for stationary and nonstationary conditions. Non-stationary loading is important for the reliability of ceramic components due to the low fracture toughness and the large scattering of strength. Proof tests should be done for assuring the minimum strength or the life-tiqe of components, because of the limitations for detecting the defects. Fig. 6 shows the mechanical meaning of proof tests [ I , 21. The proof stress should be higher than the applied maximum stress
Time t b)
Proof test pattern and strength after the test
Fig. 6
Mechanical meaning of proof tests [ 1, 21
275
DESIGN STANDARD BY STRESS RESTRICTION Design methods for ceramic components are restricted by the service conditions of temperature and stress for fracture modes [4]. Considering the temperature, the restriction is for non-creep behavior or creep behavior. Stress restriction is defined as the following equation for static and fatigue fracture. The design stress has to be lower than the maximum equivalent stress am.
UmaX ; Maximum equivalent stress, S ; Minimum design strength, Koand KT; Basic and classifying safety factors, /3 and 7 ;Design factors for stress gradient and the region, u and u N ;Distribution and nominal stresses, SU;Specifled minlmum statlc stength SW;Fatigue iimR Sr ;Specified minimum tlme-dependent strength tr ;Deslgn life
The slopes of static and cyclic fatigue curves are obtained by fatigue tests using specimens with knoop cracks.
m
W
tl
m ;Weibull shape parameter, Vmt and VE ; Reference and effective volumes, V; Volume, 5 ; Equivalent stress, u T, u 2 and 43 ;Principal stresses, t IIIU ; Maximum shear stress, ; Shear parameter, u R ; Vertical stress As shown in the diagram between strength and time in Fig. 7, the minimum design strength s is defined in the
range from static strength to time-dependent strength under the non-creep behavior of static fatigue and cyclic fatigue [4]. For the example of Sf, this strength is a time-dependent one, corresponding to the time tr. s w expresses the fatigue limit for static and cyclic fatigue. The slopes of the static and cyclic fatigue curves are obtained by fatigue tests using specimens with h o o p cracks. These values are based on many experimental data for Si3N4, and so on. Su is the specified minimum design tensile strength. The non-fracture probability is more than 99% under the condition that the stress in components is lower than the stress value of Su. Su is calculated from the mean flexural strength and the Weibull probability technique. Table 2 shows a classification for important components [4]. From class 1 to class 4, safety factors should be defined according to the importance of fracture. For example, for class 1, the safety factor K1 has the highest value, because the fracture leads to a failure of the whole system. Table 2
Classification for fracture of components The fracture leads to a failure of the whole ‘lass system. The fracture leads to a local failure of the system. ‘lass 21The operation must be suspended for the1 lremoval of the damaged subsystem. lThe fracture leads to a local failure of the system. ‘lass lThe operation of system can be resumedl lafter the replacement of the component. lThe fracture leads to decreasing of the function. Class 4 The reoair of the comoonent can be made lwithoui the system stop.
’
1
tf
Time t Static fatigue strength
Classification for important components [4]
I
R ;Stress ratio
A DESIGN EXAMPLE OF AUTOMOBILE COMPONENTS
R
;sw = 0.5 su
tt
Time t Cyclic fatigue strength Fig. 7 Minimum design strength [4]
276
A strength design example is introduced in Fig. 8 for automobile components and the design methods. Various service loads are generated for every component in automobile systems. So, strength design methods are applied to confirm the reliability under the service conditions, such as design by rules, simulations and tests. On these occasions, safety factors, endurance limits for fatigue and empirical rules are employed. The design by tests includes the fracture behavior analysis under more severe conditions than those for vehicles that are usually used. In the tests, there are methods
using components only, those put in assemblies and in vehicles. Fig. 9 shows the representative mass-produced ceramic components, such as turbocharger rotor, swirl chamber and exhaust gas control valve [5, 6, 7, 8, 91. The Si3N4 turbocharger rotor is one of the most massproduced ceramic components. However, the amount is less than we would have expected it to be. This is due to a lack of balance between cost and performance, Fig. 10 shows a fuel injection pump of the distributor type for diesel engines. in this pump, Si3N4 roller bushings are installed for the purpose of increasing the durability performance [ 101. Fig. 11 shows a strength design example for ceramic components for automobile engines. First, the design in common with materials is done for strength and lifetime predictions. Secondly, the design suitable for ceramics should be conducted to over-loading and damage-tolerance. The criteria are mostly that the strength and the service life should exceed those of metal components. Thirdly, the assurance by proof tests must be performed by mechanical and thermal loadings. Fig. 12 shows a design example by tests for wear durability of Si3N4 roller-bushings [lo]. Taking the maximum load where no seizure takes place for the SUJZ bushing to be 1, it can be seen that the ceramic bushing does not seize even under a load which is three times as large as that for the metal one. The amount of wear for metal roller-pin, ceramic roller-bushing and metal roller is expressed as a ratio ofthe prescribed
value for the conventional SUJZ bushing, which is defined to be 1. Based on the design criteria of seizure load and wear rate for metal bushings, ceramic bushings have three times durability performances of metal ones.
/Roller bushing
Fuel Injection-pump assembly
Roller Bushing
Fig. 10 A fuel injection pump of distributor type for diesel engines and Si3N4 roller-bushings [ 101
j Components
1
Power plants
Components
Englne,TIM, TIC, W, Shafts
chassis
3ervlce loads Engine torque Engine vibration Thermal loads
-
Roller-bushing for fuel pump Swirl chamber
, Static strength
i II
i
Fatigue strength [Design Methods I Deformation Design by rules FEM simulations etc. (Loadsunder 1 Testing8 Components servlce conditionsl _I '
I
Assemblles
vehicles
Fig. 8 A strength design example ; Automobile components and the design methods
Criteria
Service loads and under fluctuations Thermal stress Thermal shock Contact load and Impact load Pressure f luctuatlons Foreign object damage
Strength design items a) Design in common with materials Strength and life predictions b) Design suitable for ceramics Over loading, Damage-tolerance c) Assurance by proof tests Fracture tests (Mechanical, Thermal) Non-destructive inspections Fig. 9 Mass-produced Si3N4 components for automobile engines, as turbocharger rotor, swirl chamber and exhaust gas control valve [5,6,7,8,9]
Fig. 1 1 A strength design example ; Design methods of ceramic components for automobile engines
277
1 Test Fuel :Kerosene (90%) I
-
Design criteria
0 Non-sebute
x seizure
expected to become a JIS (Japanese Industrial Standard). 5 ) A Design Handbook is also going to be put intc circulation after several years, expanding and fulfilling the examples in the design standard for structural ceramic materials and components. Well, this standard is written down in Japanese at the present time.
REFERENCES
SUJ2
s3N4
Matelials of roller-bushing
Seizure load for ceramic and metal roller bushings Ro Basic wear rate
0 Roller-pin
0 Roller-bushlng (Inner) 0 Roller-bushlng (outer) 0
K
1.0
A Roller
\
K
Test time, (h)
Wear rate vs. test time for ceramic and metal roller bushings Fig. 12 A design example by tests ; Wear durability of Si3N4 roller-bushings [ 101 In general, the design criteria for ceramic components are established, compared with the durability of metal components.
SUMMARIES 1) A Design Standard for advanced ceramic materials and components for the reliability of ceramic systems has been almost constructed for fundamental design methods and components under service conditions in steady state. 2) Widespread uses of the Design Standard are expected for the research and development of ceramics materials and components. 3) Consequently, many applications of practical components would be accelerated in various energy and environment systems. 4) A Design Standard for particular service conditions has been investigated to complete the standard. This Design Standard includes contact components, pileup structure components, and so on. Finally, the standard is
278
[I] H. Kawamoto, “Development of Structural Ceramic Components for Automobile Applications”, Ceramic Transactions Vol. 49, Manufacture of Ceramic Components, The American Ceramic Society, (l995), 173. [2] H. Kawamoto, “Strength and Fracture of Ceramic Components for Automobile Engines”, Science of Machine, V01.40, No.1, Yokendo, LTD, Tokyo, (l988), 131 (in Japanese). [3] M. Ichikawa, “Extremely Sensitive Dependence of Strength after Proof Testing of Ceramics on Initial Strength”, Material Science Research International, Vol.1, No.1, (1995), 59. [4] A. Suzuki, “Design Standard for Preventing Static Fracture of Fine Ceramic Components”, IHI Technical Review, 28-2, (1 988), 82. [5] H. Kawamoto, T. Shimizu, and H. Miyazaki, “Strength and Reliability of Silicon Nitride Ceramic Turbocharger Rotor for High Performance Automotive Engines”, Proceedings of International Symposium on Reliability and Maintainability 1990 - Tokyo, (1 990), 199. [6] K. Takama, S. Sasaki, T. Shimizu, and N. Kamiya, “Design and Evaluation of Silicon Nitride Turbocharger Rotom”, 91-GT-258,ASME, (1991). [7] H. Kawamoto, T. Shimizu, M. Suzuki, and H. Miyazaki, “Strength Analysis of Silicon Nitride Swirl Chamber for High-Power Turbocharged Diesel Engines”, Ceramic Materials and Components for Engines, Deutsche Keramische Gesellschafi, ( 1986), 1035. [8] S. Kamiya, M. Murachi, H. Kawamoto, S. Kato, S. Kawakami, and Y. Suzuki, “Silicon Nitride Swirl Lower -Chamber for High Power Turbocharged Diesel Engines”, I985 SAE Transactions,No. 850523,3.894. [9] T. Shimizu, Y. Hirata, and K. Tanaka, “Development of Ceramic Exhaust Gas Control Valve for the Two-way Twin Turbo Engines”, Proceedings of Society OfAutomotive Engineers ofJapan, No. 9433533, (1 994). [lo] K. Noda, S. Kamiya, T. Fujimura, and K. Taniguchi, “Development of Ceramic Roller-bushings for Diesel Distributor-type Fuel injection Pump”, JSAE Review, Society of Automotive Engineers of Japan, 20 (1 999), 197.
STATIC AND CYCLIC STRESS-LIFETIME CURVES OF CERAMICS Hideo Awaji, Jin Gang, Atsushi Honjoh Nagoya Institute of Technology, 466-8555, Gokiso-cho, Showa-ku, Nagoya, Japan
ABSTRACT A statistical technique for estimating static and cyclic fatigue limits of stress-lifetime curves was proposed. Fatigue tests were performed on alumina smoothsurface specimens to estimate stress-lifetime curves by means of a three-point flexure test in air at room temperature. These tests were carried out under several stress levels with constant maximum stresses, and the data of fatigue lifetimes under each constant stress level including imperfect data were analyzed statistically to estimate the Weibull’s original distributions of the lifetime. The statistically defined fatigue limits were estimated from the relationship between the residual crack length in the specimens surviving after fatigue tests and the stress level applied. The results revealed that the ratio of the fatigue limit to the flexure strength in air was 0.53 for the cyclic fatigue and 0.71 for the static fatigue.
INTRODUCTION Recent investigations on the cyclic fatigue behavior of structural ceramics clarified the following mechanisms of crack growth[1-51. (A) Time-dependent crack extension due to an environmentally activated process under tensile stress cycles, (B) crack extension caused by degradation of stress shielding in the process zone wake under cyclic loads, and (C) mechanical crack growth caused by debris or asperity of the crack surface under an unloading cycle. Mechanism (A) operates when a material is sensitive to the testing environment, while mechanism (B) operates when a material exhibits a rising crack-resistance curve (R-curve). Additional effects on the crack extension might be produced by temperature increase at the crack tip caused by a hysteresis loop of a cyclic stress-strain curve[6]. Due to the inherent brittleness of ceramics, the static and cyclic fatigue lifetime data of structural ceramic specimens with smooth surfaces exhibit quite a wide scatter, which sometimes obscures the fatigue tendency of the materials[7,8]. Thus statistical treatment is advantageous for the analysis of the fatigue data of ceramics.
Determination of a fatigue limit in stress-lifetime (St) curves is essential for evaluating the macroscopic fatigue phenomena occurring in the industrial structural ceramics. Several researchers have estimated the fatigue limits of ceramic specimens with smooth surfaces, but the techniques used are not convenient because they require a large number of specimens and are timeconsuming[6,9, lo]. In this study, static and cyclic fatigue tests were performed in air at room temperature, using a threepoint flexure specimen with smooth surface of alumina ceramics which is expected to show both timedependent and cycle-dependent crack extension behavior. The tests were conducted under several constant stress levels[7], and the fatigue lifetimes for each stress level including imperfect data were analyzed statistically to estimate the original statistical distribution function of the fatigue lifetimes. Residual crack lengths in the specimens surviving after the running time is also estimated from the retained strength distribution measured by a three-point flexure test in an inert environment. Then the statistically defined fatigue limits of static and cyclic fatigue lifetimes are estimated fiom the relationship between the residual crack lengths in specimens surviving after the fatigue tests and the stress levels applied.
EXPERIMENTAL PROCEDURES The material used in this study was a commercial polycrystalline alumina exhibiting 99.5% purity (Mitsui Mining Co., Ltd.) and mean grain size of about 2 1.1 m. The mean values and standard deviations of several mechanical properties are shown in Table 1. Specimens 3 X 4 X 40 mm in size were machined out, chamfered and finished by polishing with diamond paste. The intrinsic or initial fracture toughness of the alumina with rising R-curve behavior, K,c, measured by the SEVNB (single-edge V-notched-beam) method[ 11,121 in an inert environment (in dry N, gas) was 3.19 MPam”*. Static and cyclic fatigue tests were performed in three-point flexure (outer span 30 mm) using a piezoelectric bimorph-type instrument[131. The cyclic stress was a sinusoidal wave with a frequency of 127
279
Table 1 Mechanical properties of alumina Three-point flexure strength in air /MPa
51 0-1-58.0
Three-point inert flexure strength /MPa
Intrinsic fracture toughness in inert environment /MParn"'
51 4k49.0
3.1 9f0.10
Lifetimesh
Hz,a constant amplitude and a stress ratio R of 0.2. The fatigue tests were conducted at room temperature in air for a running time of lo5 s. The retained strength of the specimens surviving after the running time was also estimated to obtain the residual crack length in the surviving specimens using a three-point flexure test in a dry nitrogen atmosphere.
0.1 C 0
RESULTS S-t CURVES During fatigue tests on smooth-surface specimens of ceramics, some specimens will fail before the maximum stress amplitude is reached, and some specimens will survive even after the running time. Data for the specimens which failed before the maximum testing stress was reached and the specimens surviving after the running time are termed imperfect data in this study. All the data including the imperfect data should be analyzed together to obtain the original statistical distribution function of the lifetimes, fiom the viewpoint of the scattered distributions of flaw population in the specimens. Figure 1 shows the Weibull plot for the lifetimes of alumina in the cyclic fatigue test carried out under the constant maximum stress level of 412 MPa, assuming that the lifetimes can be described by a two-parameter Weibull distribution. The solid circle (0)with arrow pointing to the left (+) indicates the imperfect data for the specimen which failed before the maximum stress level was reached, the solid circles with arrows pointing to the right (+) denote the imperfect data for the specimens surviving after the running time, and the other circles with no arrow indicate the perfect data for the lifetimes of the cyclic fatigue tests. Only the perfect data are used to estimate the Weibull's distribution function using a least-square approximation, and the imperfect data are taken up only for counting the order. Cumulative failure probability was calculated using the following symmetric rank method,
i - 0.5
F, =-
P 5
10
lntl s Fig. 1 Weibull plot for the lifetimes of alumina in cyclic fatigue tests carried out under a o max of 4 12 MPa the Weibull plots in this study were constructed in the same way. Figure 2 shows the S-t curve of alumina obtained in the cyclic fatigue tests conducted under the four constant stress levels of 450,412, 354 and 290 MPa, in which the solid circles ( 0 )indicate the cyclic fatigue lifetimes of each specimen, and the empty circles (0) the median of the distribution of the cyclic fatigue lifetimes for each stress level, estimated by means of the technique described above. The empty circle with an error-bar' denotes the median and the standard deviations of the strength distribution of alumina measured using a three-point flexure test in air, where the median value was 514 MPa. It is noticeable that the fatigue lifetimes dispersed over a wide range, hence statistical treatments are a requisite for analyzing the
6001
I
I
I
I
I
500 b
m m
400 -
2 m
4-l
5
300-
E
.d
n
where i is the order of the lifetimes, n the total number of the specimens including the imperfect data, and F', the cumulative failure probability. The shape parameter, m, and the scale parameter, , for the two-parameter Weibull distribution of the cyclic fatigue lifetimes are shown in the figure, and the value of in the figure indicates the median of the distribution function. All of 280
lo2 lo3 104
loo 10'
j-9.99
Lifetime Is Fig. 2 Stress-lifetime curve for alumina in the cyclic fatigue test
I
I
fatigue properties. The cyclic fatigue parameter, n,, was calculated to be 14, using Munz and Fett's method[ 141.
FATIGUE LIMITS A new statistical technique for estimating fatigue limits defined upon a statistical aspect is proposed here. The retained strengths in the specimens surviving after the running time in the cyclic fatigue tests were measured using a three-point flexure test in an inert environment. Figure 3 shows the Weibull plot for the original distribution of the retained strength in the specimens surviving after the running time in the cyclic fatigue tests carried out under the maximum stress level of 412 MPa. Of the ten specimens in the figure, one specimen failed before the stress level intended was reached, five specimens fractured during the cyclic fatigue tests, and four specimens survived after the running time, as shown in Fig. 2. Then the Weibull's distribution function was analyzed using the four retained strengths. The shape and scale parameters, shown in Fig. 3, indicate the Weibull parameters of the original distribution of the retained strength in the surviving specimens, and So denotes the median of the distribution function.
S/MPa
I
0.1
- 1
6
6.1
6.2
6.3
6.4
1nS/MPa Fig. 4 Weibull plot for the original retained strength distributions in the surviving specimens after the running time in the cyclic fatigue tests
S(MPa) running time in the cyclic fatigue tests. By comparing the median of these distributions, it is clear that the medians of each strength distribution hnction increase with a decrease in the stress level loaded.
540 -19.99
s
DISCUSSION
\
Under the assumption of small scale yielding, subcritical crack growth rates of ceramics in region I of crack extension rate vs. stress intensity factor relation fblfill a power-law rule,
and the stress intensity factor is given as I
I
0.11' 6 6.1 '
1
I
6.2
I
' ' 6.44
6.3
1nS Fig. 3 Weibull plot for the inert retained strengths in the surviving specimens after cyclic fatigue tests carried out under a u IMx of 4 12 MPa Figure 4 shows the original distributions of the retained strength in the surviving specimens exposed to cyclic fatigue loads with the maximum stress levels of 412, 352 and 290 MPa, together with the inert strength distribution shown by the broken line and the strength distribution in air shown by the dot-dashed line. The three solid lines indicate the original distributions of the retained strength in the specimens surviving after the
K, = YOU' ' (3 1 where a is a crack length, c fatigue time, K,(.intrinsic or initial fracture toughness of the materials with Rcurve[8], K, stress intensity factor, A and n fatigue parameters, IJ applied stress and Y shape factor of the specimen used. Assuming that the fatigue lifetime and the strength of ceramics are caused by the same flaw, the initial pre-existing crack length, ai,can be conducted directly from the inert strength, S,, and the intrinsic fracture toughness, K,. u,
=[$)
(4)
If the inert strength can be described by the twoparameter Weibull distributions,
28 1
flexure test in an inert environment, and critical crack length in specimens test in air were calculated from Eq. (8) to be 37.1 and 38.7 p m, respectively, and the median values of the residual crack length distribution function in the surviving specimens after loading the stress levels of 290, 352 and 412 MPa to be 39.6, 49.1 and 61.6p m, respectively. Figure 6 shows the relationship between the median values of the residual crack length in the surviving specimens and the maximum stress levels loaded in the cyclic fatigue test. The three solid circles ( 0 )indicate the median of the residual crack distributions for each stress level, and the broken horizontal line denotes the median value (37.1 ,u m) of the initial crack length of alumina. The intersection of these lines, therefore, gives
-9.99
s \
0 . l L 4.8
I
4.9
‘
1
5
‘
I
5.1
‘-I
5.2
(-1/2)lna Fig. 5 Weibull plot for the residual crack length in the surviving specimens after the running time in the cyclic fatigue test under each stress level Inln-=m, 1 1- F
s1d
In(?)
the distribution of the initial crack length is derived from Eqs. (4) and (9,as
Equation (6) indicates that the Weibull plot of u;’” gives
Fig. 6 Determination of the cyclic fatigue limit, in which the solid circles indicate the residual crack length.
the two-parameter Weibull distribution with the shape parameter, m,, and the scale parameter, YS,,/K,,. In the same way, if the retained strength of the surviving specimens measured in an inert environment can be described by the following two-parameter Weibull distribution, (7) then the residual crack length distribution in the surviving specimens is expressed as
where, S, is the inert retained strzngth of the surviving specimen, m, and S,, Weibull parameters, and a, a crack length in the surviving specimens. Figure 5 shows the Weibull’s distribution functions of the residual crack length in the surviving specimens exposed to the cyclic fatigue loads with the several constant stress levels. The median values of the crack length pre-existing in the specimens measured by a 282
Fig. 7 Comparison of stress-lifetime curves obtained in the static and cyclic fatigue tests, in which the dotted line indicates the S-t curve of static fatigue, the solid line the apparent curve, and the dot-dashed line the modified curve of cyclic fatigue
the fatigue limit of alumina under cyclic fatigue loads which was calculated to be 270 MPa, because no crack extension is expected under the stress level. The same technique is applicable to the estimation of the static fatigue limit, which will be presented elsewhere[9]. The ratio of the fatigue limit to the flexure strength in air was calculated as 0.71 for the static fatigue and 0.53 for the cyclic fatigue. Figure 7 shows the complete S-t curves obtained in the static and cyclic fatigue tests, in which the empty squares (0)indicate the median of the distribution of the static fatigue lifetimes, the empty circles (0)the median of the distribution of the apparent cyclic fatigue lifetimes, and the dot-dashed line the modified S-t curve in the cyclic fatigue test calculated using the following g, modification[ 151. (9) where u max is the maximum stress of cyclic wave, u the cyclic stress and ;1 the cyclic period. By comparing the static and modified cyclic curves, it is apparent that the difference between them in the figure is relatively small at higher stress levels, but more significant at lower stress levels. This suggests that at higher stress levels, cracks propagate so quickly that the cyclic effects will not be sufficient to reduce the stress shielding by bridging and/or interlocking on the process zone wake. In addition, at higher stress levels, the time-dependent mechanism of cyclic fatigue due to an environmentally enhanced crack growth is predominant, and the difference between the lifetimes in the static and cyclic fatigue tests becomes small. On the other hand, at lower stress levels, the rate of the crack extension caused by cyclic fatigue loading is so low that the bridging at the process zone wake is extensively weakened due to the cyclic stress, which accelerates the crack extension. Hence, the difference between the lifetimes in the static and cyclic fatigue tests becomes larger. It is also evident that stress corrosion cracking plays no role in the determination of the cyclic fatigue limit because the cyclic fatigue limit is obviously lower than the static fatigue limit. Figure 7 indicates that there are mainly two crack extension mechanisms in alumina ceramics under a cyclic fatigue load, the one is timedependent crack extension caused by stress corrosion cracking, the other is cycle-dependent crack extension caused by degradation of stress shielding in process zone wake.
CONCLUSION Static and cyclic fatigue tests were performed on 99.5% pure alumina of moderate grain size (about 2 p m) using smooth-surface specimens under several stress levels. A statistical treatment, including imperfect data, was used to estimate the original distributions of fatigue lifetimes. A technique for estimating the static and cyclic fatigue limits defined upon a statistical aspect was also proposed using the relationship between the residual crack length in the specimens surviving after
the running time and the stress level loaded. The results revealed that the ratio of the fatigue limit to the flexure strength in air was 0.53 for the cyclic fatigue and 0.71 for the static fatigue. The statistical technique for estimating the fatigue limits proposed here was simple and quite reasonable compared with other techniques which require a large number of specimens and are time-consuming.
REFERENCES (1) S. Horibe and R. Hirata, Cyclic Fatigue of Ceramic Materials: Influence of Crack Path and Fatigue Mechanisms, Acta Metall. Mater., 39 (1991) 13091317. (2) D. S. Jacobs and Lwei Chen, Mechanical and Environmental Factors in the Cyclic and Static Fatigue of Silicon Nitride, J. Am. Ceram. SOC.,77 ( 1994) 1153- 1 16 1 . (3) H. Kishimoto, A. Ueno and H. Kawamoto, Crack Propagation Characteristics of Sintered Si,N, under Static and Cyclic Loads, 3. SOC.Mater. Sci. Japan, 36 (1987) 1122-1127 (in Japanese). (4) R. H. Dauskardt, W. Yu and R. 0. Richie, Fatigue Crack Propagation in Transformation-Toughened Zirconia Ceramic, J. Am. Ceram. SOC.,70 (1987) C248-252. (5) R. H. Dauskardt, A Frictional-Wear Mechanism for Fatigue-Crack Growth in Grain Brdging Ceramics, Acta Metall. Mater., 4 1 (1 993) 2765-278 1. (6) D. A. Krohn and D. P. H. Hasselman, Static and Cyclic Fatigue Behavior of a Polycrystalline Alumina, J. Am. Ceram. SOC.,55 (1972) 208-211. (7) H. Awaji, S. Yamamoto, S. Honda and T. Nishikawa, Static and Cyclic Stress-Lifetime Curves of Ceramics, J. Mat. Sci. Letters 19 (2000) 713-715. (8) G. Jin, A. Honjyo and H. Awaji, Static and Cyclic Fatigue Limits of Alumina Ceramics, J. Ceram. SOC. Japan 108 (2000) 614-617. (9) T. Sakai and K. Fujitani, A Statistical Aspect on Fatigue Behavior of Alumina Ceramics in Rotating Bending, Engng. Fract. Mech., 32 (1989) 653-664. (10) B. J. S. Wilkins and R. Dutton, Static Fatigue Limit with Particular Reference to Glass, J. Am. Ceram. SOC.,59 (1976) 108-112. (1 1 ) H. Awaji and Y. Sakaida, V-Notch Technique for Single-Edge Notched Beam and Chevron Notch Methods, J. Am. Ceram. SOC.,73 (1990) 3522-3523. (12) H. Awaji, T. Watanabe, Y. Sakaida and H. Nakagawa, Fracture Toughness Measurements of Ceramics by N Notch Technique, Ceramics Int., 18 (1992) 11-17. (13) K. Ohya and K. Ogura, Cyclic Fatigue Testing Device for Fine Ceramics by Using Piezo-electric Bimorph Actuator, J. SOC. Mater. Sci. Japan, 38 (1989) 44-48. (14) D. Munz and T. Fett, Ceramics, Springer-Verlag, Berlin (1 999) 89-9 1 . (15) A. G. Evans and E. R. Fuller, Metal. Trans., 5 (1974) 27-33.
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LIFETIME PREDICTION OF CERAMIC THERMAL BARRIER COATINGS BASED ON LIFETIME ANALYSES OF CLOSE TO REALITY TESTS M. Bartsch", G. Marci, K. Mull, Ch. Sick German Aerospace Center (DLR), Institute for Materials Research, D-51147 Koln, Germany ABSTRACT A lifetime prediction concept for thermal barrier coatings (TBC) based on lifetime analyses of close to reality tests is given. Basic idea is to ensure that failure mechanisms during testing are the same as under service conditions in an engine. Since, it is impractical to simulate service load cycles of about at least one hour in real-time, number of cycle dependent fatigue damages and time-at-high-temperature dependent damages by kinetic mechanisms are imposed separately. With a newly built testing facility the cyclic load of an aircraft engine service cycle is simulated but with drastically reduced holding times. First test results on specimens with an electron beam physical vapour deposited (EB-PVD) TBC exhibit that the as-coated properties of the ceramic coating are sufficient to survive the fatigue loading under service conditions. The influence of kinetic damage mechanisms on the TBC-lifetime will be determined in tests with pre-aged specimens.
hot gas to the metallic substrate under stationary conditions, depending on the heat flux in the turbine engine [2, 31. The temperature gradient across the wall of a coated and internally cooled turbine blade is schematically shown in Fig. 1. As additional important effect TBCs decrease the temperature transients, that means the thermal shock load, in the metallic substrate during start and shut down of the engine. Temperature
L 100°C
r Hot Gas
INTRODUCTION In advanced gas turbines for aircraft engines or electric power generation, internally cooled turbine blades are coated with a ceramic thermal barrier to reduce the thermal load of the metallic substrate. Exhausting the full potential of the ceramic thermal barrier coating to increase the turbine inlet temperature (TIT) requires reliable lifetime prediction methods, since after failure of the TBC the remaining lifetime of the blade will be extremely short.
The thermal barrier coating system The commonly used TBC material is yttria stabilised zirconia applied by electron beam physical vapour deposition (EB-PVD) or plasma spraying (PS). Zirconia TBCs are applied on airfoil surfaces coated with a metallic corrosion resistant layer which imparts good adhesion for the TBC on the substrate. This so called bond coat (BC) consists frequently of MCrAlY- or PtA1-alloys. The thermal conductivity of zirconia TBCs is relatively low and depends on the deposition process. Plasma sprayed coatings have initial values for the thermal conductivity between 0.8 and 1.2Wm-'K-' while EB-PVD TBCs have somewhat higher conductivity of 1.5-1.8Wm-'R1[1]. Due to the low thermal conductivity zirconia coatings of about 150-300pm thickness maintain a temperature drop of 60" to 17OOC from the
Ceramic .
Gas Film
. Metallic
Substrate
Bond Coat
Fig. 1. Thermal gradient across the wall of an internally cooled turbine blade with TBC Stresses by thermal mismatch between TBC and nickel- based super-alloy substrates are relatively low due to the similar thermal expansion coefficient which K1for the is at max. service temperature about 9-1 1. TBC and 14 - 16.10-6K-'for the substrate. The initially Young's-modulus of zirconia TBC is reasonably low with about only 35 - 60 GPa [4,51 and supplies high tolerance against straining and thermal shock.
Damage mechanisms Under service conditions the properties of the TBCsystem change. Progressive damages are caused by stresses in and between the different layers of the TBC system, generated by the complex thermal and mechanical cyclic loading and the mismatch of several physical properties of the substrate and the coating materials. During the lifetime of a TBC system, local loading conditions and resistance against damage and failure change by reason of changing physical properties of the individual layer materials. For example, densification of the ceramic topcoat by sintering leads to increased Young's-modulus and thus decreased strain tolerance. Thermal cycling above 1200°C cause phase transformations from partially stabilised tetragonal
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zirconia to monoclinic phase resulting in volume changes and hence increase of local stresses [6, 71. A significant portion of residual stresses perpendicular to the coated surface develops due to the growth of the oxide scale (TGO) on the bond coat. Diffusion of alloy elements and impurities from the substrate to the interface between bond coat and ceramic reduces the adhesion of the TBC [8]. The failure criterion of the TBC is defined as spallation of the TBC from the substrate since the function of the thermal barrier is impaired most seriously by local loss of the TBC. Spallation within the ceramic topcoat is called ‘white failure’. If the separation of the TBC propagates within the TGO or at the interface between TGO and BC it is called ‘black failure’. White failure is more often observed in PS coatings and highly dependent on interfacial roughness with evidence to slow (subcritical) micro-crack growth, crack linking and eventual buckling [9, lo]. Black failures are dependent on the extend of bond coatoxidation and frequently related to EB-PVD-coatings with their typical smooth interface between TGO and bond coat. The damage mechanisms and their interaction leading to spallation of the TBC, whether PS- or EBPVD processed, are not well understood. Generally we can distinguish between fatigue mechanisms, which are number-of-cycle dependent and kinetic mechanisms depending on time-at-high-temperature. Which mechanisms are dominating and which synergy effects are working depends on the materials parameter of the certain coating system and the service conditions. For example in aircraft engines the maximum temperature during one service cycle is higher than in stationary gas turbines while the time at high temperature is longer in the latter.
Lifetime prediction concept For TBC lifetime prediction in [4, 11,121 MansonCoffin type damage accumulation rules were developed which base on empirical damage parameters related to certain damage mechanisms. Kinetics and dependence of cyclic loading of the damage parameters are determined separately in simplified experiments. Subsequently the contribution of every damage mechanism to the consumption of initial lifetime under more realistic testing conditions will be accumulated. For example in [4] as time-at-high-temperature dependent damage parameter the thickness of the TGO is set in relation to the degree of damage due to the mechanism ‘oxidation of the bond coat’. Kinetics of TGO-growth is determined in oxidation experiments. The decrease of number of thermal cycles to failure due to TGO-growth was calculated with the cumulative damage rule for burner rig tests and experimentally verified. However, service conditions are quite more complex and consequently the number of parameters which have to be taken into account for estimating the TBC lifetime is reasonably high. Additional fitting parameters are needed to include interaction effects of the different mechanisms. The required high number of parameters needed to describe the damage situation makes this
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concept impractical for lifetime prediction under service conditions. Instead of accumulating the contributions of separate damage mechanisms determined under simplified testing conditions our idea is to make the testing conditions as realistic as possible. Thus, we ensure that the damage mechanisms including interaction effects are identical to that under service conditions [13, 141. However, it is impractical to simulate a service cycle in real time since, turbine blades in modern gas turbines are designed for about 20.000h in the case of aircraft engines and up to 50.000h for stationary engines. The time of one test cycle can be decreased drastically if we determine number-of-cycle dependent fatigue damages separate from damages by kinetic mechanisms because fatigue damages are induced by changes of the load and not by the holding time on a certain load level. Damages induced by time-at-high-temperature dependent kinetic mechanisms will be considered by testing heat-treated specimens.
EXPERIMENTAL Requirements for realistic test cycle Under service conditions coated turbine blades are loaded simultaneously by thermal and superposed mechanical cycling. The thermal and the mechanical load depend on the blade location. E.g. in rotor blades the mechanical load due to centrifugal forces increases from the blades tip to the blades fillet while the thermal load at the fillet is the lowest. With simple thermal cycling we are not able to simulate variable straining at a certain temperature since the strain amplitude due to thermal cycling is a function of the temperature difference. The strain of the compound of metallic substrate and TBC-system is compulsorily identical over the cross section. The strain compulsion consequently entails a complex stress distribution over the cross section due to the different physical properties of the single materials and the thermal gradient over the cross section. The stress distribution for the same strain is quite different under isothermal conditions and in a thermal gradient respectively. Since, the stress amplitudes at the failure location of the TBC-system are the driving force for the TBC-failure, it is necessary to impose a realistic temperature gradient during testing. On the other hand the temperature gradient has consequences for kinetic damage mechanisms because the time constants of kinetic mechanisms depend on the temperature.
Testing facility and load cycle A testing facility was built which simulates closely the conditions for a gas turbine blade, i.e. cycles with simultaneous mechanical and thermal loading, including thermal shock and temperature gradients over the cross section of the specimens. Hollow cylindrical specimens with a metal substrate wall of 2 mm, similar to the typical geometry of the convex side of a turbine blade near the leading edge, are tested. A sketch of the specimen is shown in Fig. 2.
M14x1
Fig. 2. Hollow specimen for internal cooling during simultaneous thermal and mechanical loading
Mechanical loading is imposed on the specimen by a servo-hydraulic testing machine. For simultaneous heating, the radiation of four quartz lamps is focused onto the specimen by elliptical mirrors. A temperature gradient is generated by internal air cooling of the specimen. Mass flow and temperature of the cooling air are controlled. For simulating the thermal shock during shut down of the gas turbine, the specimen is rapidly cooled with cold air blown onto the surface from 36 vents which are integrated in two sliders enclosing the specimen during cooling. A detailed description of this so-called Thermal Gradient Mechanical Fatigue (TGMF) testing facility is given in [ 151. With this testing facility a flight cycle of an aircraft engine is simulated. The whole flight cycle including take off, climbing, cruise, thrust reverse and shut down takes about 1-5 minutes, that is about 1% of the real flight time (Fig. 3).
-
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Fig. 3. Load and temperature spectrum of a test cycle (The relative stress is related to the yield stress at stationary cycle temperature.) Mechanical load and temperature are adjusted independently to simulate the loading conditions for several blade locations (e.g. blade-tip or -fillet).
Test program and results First experiments were carried out with specimens from directionally solidified nickel-based super-alloy IN100 DS to get an anisotropy effect like in single crystal blades. All specimens had a llOpm thick
NiCoCrAIY- bond coat. Some specimens had an additional 220pm thick ZrOz - TBC stabilised with 8wt%Y203. Both layers were processed by EB-PVD. The substrate material as well as the coatings were processed by the DLR Institute of Materials Research in Koln. The thermal cycle during testing was rapid heating until a given stationary temperature was reached and rapid cooling at the end of the cycle, simulating the shut down of the turbine engine. In contrast to the mechanical loading no additional temperature changes were imposed since, by reason of the thermal inertia of the system it would take too much time to reach certain intermediate temperature levels. The stationary cycle temperature was 920°C at the surface of the bond coat for specimens without as well as for specimens with ceramic topcoat. Mass flow and temperature of the internal cooling air were kept constant during all experiments. The resulting temperature gradient for the chosen conditions was measured with thermocouples at a specially prepared reference specimen. Between the bond coat surface and the surface of the ceramic topcoat the temperature difference was about 35°C and between bond coat surface and the cooled side of the substrate about 85°C. The mechanical load cycle included several certain levels with one maximum representing the load during start and climbing of the aircraft and a second maximum representing the thrust reverse after landing. Experiments with specimens only with bond coat were conducted to determine a realistic load level. For this purpose the level of the mechanical load spectrum was subsequently decreased until we got a number of test cycles to failure about 3000 cycles which is comparable to the number of flights a turbine blade was designed for, Fig. 4. The relation between the number of cycles to failure and the maximum of the mechanical load spectrum of the cycle can be described by a power law. Specimens with additional TBC were tested on the same load levels and identical bond coat temperature of 920°C like the specimens only with bond coat. Lifetime of the specimens substrate was significantly higher with TBC than without TBC. All experiments were terminated without failure of the specimens substrate. No spallation or cracks were found on the surface of the TBC.
287
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Measured number of test-cycles to failure (TBC-coated specimens did not fail until the test was terminated)
DISCUSSION The results of this work show that the investigated EB-PVD thermal barrier system survive a requested number of realistic load alternations. However, the test cycle includes the realistic load alternations but only extremely reduced holding times of about only 1% compared with a typical service cycle. The maximum testing time at high temperature was about 150h which is not enough to induce remarkable damages to the TBC system by kinetic mechanisms. Since, TBC’s in service show premature failure [3, 16, 171 damages by kinetic mechanisms are implied. In service fatigue damage mechanisms as well as kinetic damage mechanisms are working simultaneously. Significance of mechanically induced cyclic strain on the lifetime of TBC was demonstrated by [18] in Thermal Mechanical Fatigue (TMF) experiments. On the other hand it is known that kinetic damage mechanisms decrease the strain tolerance of TBC’s entailing premature failure [4, 12, 18, 191. For a lifetime prediction in service we need information about the interaction between fatigue and kinetic mechanisms. Driving force for spallation are the stresses at failure location. The externally imposed load cycle does not change during service. In contrast due to exposure time at high temperature and due to the number of cycles, or more exactly number of load alternations, we have changes of materials properties with the consequence of increased local stresses or/and decreased resistance against failure. However, there is a lack of data about the local situation at the failure location in service. Thus, for designing turbine blades with designed-in TBC-systems lifetime prediction methods are needed which do with the available information about externally imposed load and with data about lifetime behaviour of the special substrate/ TBC-system
288
compound as function of the variable external loads. Our strategy is to simulate with the TGMF-testing the externally imposed load alternations as realistic as possible - as there are T,, Tmi,, heating and cooling rate (dT/dt), thermal gradient across the TBC cross section (as a function of heat flux and geometry) and mechanically induced load. With information of lifetime behaviour of variable (time and temperature) pre-aged specimens in TGMF-testing we can calculate the lifetime of the TBC for TGMF-cycles or comparable service cycles with extended holding times. Ageing of the TBC-coated specimens also has to be conducted with a thermal gradient because the time constants of the kinetic mechanisms depend on the temperature. Since the temperature difference over the thickness of the TBC is about 100°C in service, isothermal testing would either impose too high temperatures at the bond coat in relation to the surface of the ceramic topcoat or vice versa. An acceleration of kinetic processes is possible by increasing the temperature. In this case we have to ensure that not damage mechanisms will be activated which are irrelevant for the TBC-life in service. Thermal load during ageing will be imposed cyclic like in service because evidence was found, e.g. in [20, 211, that cyclic and constant thermal load of same exposure time at high temperature cause different damages due to oxidation. For calculating lifetime curves of TBC’s, the thermal cycling under thermal gradient until TBCfailure supplies the data for determining the lifetime limit towards decreasing superposed mechanical load. Lifetime curves calculated for different holding times may look like sketched in Fig. 5 .
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According to our experimental results no TBC failure is expected for short-term loading, i.e. short holding times at high temperature even at high mechanical loading. In this load regime the lifetime of the TBC will be limited by the lifetime of the substrate, like it is sketched in Fig. 5 . Increased cycle time is supposed to reduce the damage tolerance of the ceramic top coat due to kinetic mechanisms like sintering. The premature failure of the TBC will limit the lifetime of the component in this case. During long term exposure acceleration of TBC failure may occur by change of dominating damage mechanisms by reason of different time constants, e.g. oxidation of the bond coat after sintering has been almost completed. Activating new or changing the dominating damage mechanisms will consequently change the slope of the TBC lifetime curve, as indicated in Fig. 5 . In order to determine an interaction parameter between fatigue and kinetic damage mechanisms and verification of lifetime calculations we inte.nd to test specimens with as coated TBC's in TGMF-cycles with extended holding times. Beside the cycle time and mechanical load level also the cycle temperature level has to be varied to achieve lifetime data for conditions at different blade locations. That combinations of thermal and mechanical load levels are relevant which entail premature failure of the TBC at blade locations which lifetime after TBC failure is less than the design life of the component or less than the inspection intervals of the engine. How the relevant load regime for TGMF-tests of TBC's can be delimited by lifetime curves of the substrate without TBC has been discussed in [22].
However, we found that the load capability of a specimen with TBC is increased compared to an specimen without TBC even if the substrate temperature during holding time is the same. This beneficial effect of the TBC is due to a reduction of thermal shock.
CONCLUSIONS Specimens with a ceramic thermal barrier coating were tested in thermal gradient mechanical fatigue with realistic alternating loads but extremely short holding times. The results show that the as-coated properties of EB-PVD TBC-systems are sufficient to survive the required number of load alternations during service life. Thus, the premature failure of TBC's observed in service is due to time-at-high-temperature dependent kinetic mechanisms. Fatigue damages due to cycling are supposed to reduce TBC lifetime simultaneously to damages by time-at-high-temperature dependent kinetic processes. To determine the interaction effects between fatigue and kinetic damage mechanisms TGMF-testing of pre-aged specimen as well as TGMF-testing with extended holding times are supposed to be suitable.
ACKNOWLEGEMENT Portion of this work is financially supported by the HGF strategy fund project 'Innovative Material Systems for Increased Efficiency of Stationary Gas Turbines and Aero Engines' of the 'Hermann von HelmholtzGemeinschaft Deutscher Forschungszentren' .
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REFERENCES R.B. Dinwiddie, S.C. Beecher, W.D. Porter, B.A. Nagaraj, ASME-Report 96-GT-282, (1996). G.W. Goward, Surface and Coatings Technology, 108-109, (1998) 73-79. A. Mariccochi, A. Bartz, and D. Wortmann, Proc. ‘Thermal Barrier Workshop’ March 27-29, NASA Conf. Publ. 3312, (1995) 79-89. S.M. Meier, D.M. Nissley, K.D. Sheffler, ,,Thermal Barrier Coating Life Prediction Model Development“, Phase 11, Final Report, NASA Contractor Report 189111, United Technologies Corp., Pratt & Whitney, (1991) G. Marci, M. Bartsch, K. Mull, Determination of the Young’s modulus of TBC-materials for thermally and mechanically highly loaded components in gas turbines, (in german) DVMBericht “Werkstoffpriifung 1999“, 2.13. Dez. ‘99, Bad Nauheim, (1999) 271-280. U. Schulz, Phase Transformation in EB-PVD Yttria Partially Stabilized Zirconia Thermal Barrier Coatings during Annealing, J. Am. Ceram. SOC.,83 [4] (2000) 904-910. L. Lelait, S. Alperine, C. Diot, R. Mevrel, ’Thermal Barrier Coatings: Microstructural Investigation after Annealing‘, Mater. Sci. Eng. A, A121, (1991) 475-82. U. Kaden, C. Leyens, M. Peters, W.A. Kaysser, Thermal Stability of an EB-PVD Thermal Barrier Coating System on a Single Crystal Nickel-Base Superalloy, in : Elevated Temperature Coatings: Science and Technology 111, J.M. Hampikian & N.B. Dahotre (Hrsg.), TMS, (1999) 27-38. A.G. Evans, J.W. Hutchinson, On the Mechanics of Delamination and Spalling in Compressed Films, Int. J. Solids Structures, Vol. 20, No. 5 , (1984) 455-466. [lo] M.Y. He, A.G. Evans, J.W. Hutchinson, Effects of morphology on the decohesion of compressed thin films, Mat. Sci. Engineering A245, (1998) 168-181. [ 113 R.A. Miller, Oxidation-Based Model for Thermal Barrier Coating Life, J. Am. Ceram. SOC.,67 [8] (1984) 517-521.
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[ 121 J.T. DeMasi, K.D. Scheffler, M.Ortiz, Thermal
Barrier Coating Life Prediction Model Development, Phase I - NASA Contractor Report 182230, United Technologies Corp., Pratt & Whitney, (1989). 131 W. A. Kaysser and M. Bartsch, Fatigue of Thermal Barrier Coatings, Proceedings of the 7* international conference Fatigue’99, P.R. China, Beijing June 8-12, 1999, ed. X.R. Wu & Z.G. W a g , Vol. 3, (1999) 1897-1904. 141 M. Bartsch, G. Marci, K. Mull, Ch. Sick, Fatigue Testing of Ceramic Thermal Barrier Coatings for Gasturbine Blades, Advanced Engineering Materials, 2 [ll] (1999) 127-129. [15] G. Marci, K. M. Mull, Ch. Sick, M. Bartsch, New Testing Facility and Concept for Life Prediction of TBC Turbine Engine Components, 3” Symposium on Thermo-Mechanical Fatigue Behavior of Materials, 4.-5. Nov 1998, Norfolk, Virginia, ASTM STP 1371, (2000) 296-303. [16] Z. Mutasim and W. Brentnall, Proc. ‘Thermal Barrier Workshop’ March 27-29, NASA Conf. Publ. 3312, (1995) 103-112. [17] M. Arana, J. Goedjen, Perspectives on TBC Life Prediction for Industrial Gas Turbines, Proceeding of the TBC Winterworkshop, Jan. 6-8, University of California, Santa Barbara, (1999). [ 181 P.K. Wright, Influence of cyclic strain on life of a PVD TBC,Mat. Sci. Engineering A245, (1998) 191-200. [19] R. Herzog, F. Schubert, L. Singheiser, The influence of substrate deformation on the damage of a TBC compound under thermomechanical loading, Proccedings of Euromat 99, 27-30 Sept. 1999, Miinchen, (2000), in press. [20] M. Schutze, Materials and Corrosion 47, (1996) 103-105. (in german) [21] M. Schiitze, Protective oxide scales and their break down, John Wiley & sons, Cichester, UK, (1997). [22] M. Bartsch, G. Marci, K. Mull, Ch. Sick, Lifetime Prediction for Ceramic Thermal Barrier Coatings, Proceedings of the Euromat’99 Conference , 27.30. Sept. 1999, Munchen, Volume 11, Surface Engineering (ed.: H. Dimigen), (2000) 25-30.
CREEP OF A SILICON NITRIDE UNDER VARIOUS SPECIMEN/LOADING CONFIGURATIONS Sung R. Choi, Lynn M. Powers, Frederic A. Holland, and John P. Gyekenyesi NASA Glenn Research Center, Cleveland, Ohio 44135, USA
ABSTRACT Extensive creep testing of a hot-pressed silicon nitride (NC132) was performed at 1300°C in air using five different specimenlloading configurations, including pure tension, pure compression, four-point uniaxial flexure, ball-on-ring biaxial flexure, and ringon-ring biaxial flexure. Nominal creep strain and its rate for a given nominal applied stress were greatest in tension, least in compression, and intermediate in uniaxial and biaxial flexure. Except for the case of compressive loading, nominal creep strain generally decreased with time, resulting in less-defmed steadystate condition. Of the four different creep formulations - power-law, hyperbolic sine, step, and redistribution models - the conventional power-law model provides the most convenient and reasonable means to estimate simple, quantitative creep parameters of the material. Predictions of creep deformation for the case of multiaxial stress state (biaxial flexure) were made based on pure tension and compression creep data by using the design code CARESICreep.
INTRODUCTION Advanced ceramics are candidate materials for hightemperature structural applications in gas turbine engines and heat recovery systems. The two major limitations of these materials, slow crack growth and creep, are generally encountered in high-temperature applications. At higher temperatures, particularly at lower applied stress, enhanced creep takes place in the form of permanent deformation and/or damage accumulation, leading to loss of structural integrity and possibly the eventual rupture of components. Therefore, for higher-temperature applications, the accurate determination of creep behavior including creep and rupture parameters and associated mechanisms is important for ensuring structural integrity and component life. Because of difference in creep between tension and compression [ 1-31, parameters that are derived from relatively easy flexural testing using the conventional simple beam theory, which assumes the neutral axis to be fixed, can be misleading. Many structural components are subjected to multiaxial stresses, typically combined with tensile and compressive stresses. Therefore, in order to accurately predict or estimate creep deformation and rupture of multiaxially stressed components, creep and rupture parameters of a material should be determined individually both in tension and in compression,
together with an appropriate predictionldesign methodology. The immediate objective of this work was to develop and conduct creep testing to determine creep behavior of a hot-pressed silicon nitride under various loading configurations at 1300°C in air. The loading configurations used in this study include pure tension, pure compression, four-point uniaxial flexure, ball-onring biaxial flexur e, and ring-on-ring biaxial flexure. Creep displacements for each specimenlloading configuration were determined as a function of time with several different levels of applied loads. NASA Glenn Research Center has developed an analytical methodology and an integrated design program named CARESICreep (Ceramics Analysis and Reliability Evaluation of Structureslcreep) to be used for predicting the creep deformation and rupture life of structural ceramic components [4,5]. The second objective of this work was to validate the CARESICreep design code on the basis of the creep database generated from the experimental work. This was done by predicting creep deformation subjected to multiaxial (biaxial) stress state based on the basis of pure tension and pure compression data and by comparing with actual biaxial flexure data.
EXPERIMENTAL PROCEDURES The material used in this work was a hot-pressed silicon nitride (designated as NC132, Vintage 1990, fabricated by Norton Co., Northboro, MA) containing MgO as primary sintering aid. This material was chosen since it has shown for decades a controlled uniformity in mechanical and physical properties such as hardness, fracture toughness, strength, and slow crack growth, etc. The material also has been extensively characterized previously to determine fatigue andlor creep life prediction parameters as a candidate gas turbine material [6-81. Creep testing was performed in dead-weight creep frames (Applied Testing Systems, Inc., Butler, PA) at 1300°C in air using five different specimenslloading configurations, including pure tension, pure compression, four-point uniaxial flexure, ball-on-ring biaxial flexure, and ring-on-ring biaxial flexure. The test specimenlloading configurations used in creep testing are depicted in Figure 1 . Detailed experimental procedures and techniques can be found elsewhere [9]. Briefly, in pure tension, the dog-bone-shaped, pinloaded tension specimens with round cross-section
29 1
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e) Ring-on-ring biaxial flex.
I
Figure 1. Five different specimedloading configurations used in this study: (a) pure tension; (b) pure compression; (c) four-point uniaxial flexure; (d) ball-on-ring biaxial flexure; (e) ring-on-ring biaxial flexure. were used that were nominally 2.0 mm in diameter, 70 mm in overall length, and 20 mm in gage length. The pin holes in the tension test specimens were also tapered to minimize misalignment. Creep displacements were monitored using a scanning laser extensometry system (LaserMike, Dayton, OH) with two Sic flags (typically apart from 12 to 15 mm) that were attached to the gage section of each test specimen by friction. Six different nominal tensile stresses ranging from 10 to 81 MPa were utilized. Typically one specimen was used at each applied load. Testing was conducted in accordance with ASTM Test Method C1291 [lo]. In pure compression testing, cylindrical test specimens 3.5 mm in diameter and 7.5 mm in height (= gage length) were used. The choice of such dimensions for the compression test specimens was intended to minimize undesirable phenomena such as buckling and barreling, with a suggested ratio (height to diameter, Wd) of 1.5 to 2.0 [ll]. Since creep deformation of ceramic materials, in general, is significantly smaller in compression than in tension, higher applied stresses, ranging from 160 MPa to 630 MPa, were used in compression testing. A total of five different nominal applied stresses were used, with one specimen tested at each stress. The creep displacement measurements in the gage section were made using a three-probe LVDT extensometer placed between the two ends of each test specimen. In four-point uniaxial flexure testing, flexure beam specimens were loaded in a Sic four-point flexure fixture with 20 mm- inner and 40 mm- outer spans. The nominal dimensions of test specimens were 3 mm by 4 mm by 50 mm, respectively, in height, width and length. A total of 7 different nominal stresses, ranging from 29 to 162 MPa, were used, with one test specimen tested at each nominal stress. An LVDT probe was placed underneath the center point (tension side) of each test specimen to measure creep deflection. In ball-on-ring biaxial flexure testing, 2-mm-thick, 45-mm-diameter disk specimens were loaded in the 292
biaxial test fixtures. The lower support fixture consisted of a total of nine 9-mm-diameter Sic balls mounted on the flat Sic support block, evenly spaced peripherally to form a pitch diameter of 40 mm. The upper 9-mmdiameter S i c loading ball was slightly ground to form a small, flat, round surface of 1.3 1 mm diameter that was to be in contact with the center of each test specimen. The center creep deflections of test specimens were determined with an LVDT probe. A total of four applied loads ranging from 45 to 182 N, which corresponded to nominal maximum stresses of 29 to 115 MPa, calculated based on idealized elasticity solution [ 121, were utilized, with one disk specimen tested at each applied load. Finally, in ring-on-ring biaxial testing, each disk specimen was loaded between the upper loading and lower support biaxial flexure fixtures. The upper loading fixture consisted of seven 9-mm-diameter SIC balls, evenly spaced peripherally forming a pitch diameter of 15 mm. The lower support fixture was the same as that used in the ball-on-ring biaxial flexure testing. The dimensions of test disk specimens used in ring-on-ring biaxial flexure were the same as those used in ball-on-ring biaxial flexure. The center deflection of each test specimen was monitored using an LVDT probe. A total of four applied loads ranging from 200 to 565 N, with corresponding nominal maximum stresses (calculated based on Ref. 12) ranging from 29 to 81 MPa, were used. One disk specimen was tested at each applied load.
RESULTS AND DISCUSSION EXPERIMENTAL RESULTS A summary of creep deflections (or deflections) as a function of time determined from the five different specimedloading configurations is presented in Figure 2. The end point of each curve corresponds to the end of the test. No test specimen failed up to its time of interruption except for the test specimen fractured in compression testing with a high nominal stress of 630 MPa. In most cases, surfaces of the tested specimens revealed some degree of oxidation, depending on the test time period. However, no visible signs of the microcracks associated with enhanced creep were observed on the specimens, even after the removal of their oxide layers. In tension creep testing, test times spanned fi-om 20 days at 81 MPa to 60 days for 41 MPa. The displacement curves for the specimens tested at 10 and 28 MPa seemed to show the steady-state creep regions where displacement rate remained relatively constant. By contrast, the curves for the specimens tested at 57 and 81 MPa showed no well-defmed steady-state region, but rather a change in displacement rates with time. In compression creep testing, most of the curves (except the one at 500 MPa) exhibited the steady-state creep region, which is in somewhat good contrast with the tensile creep curves. However, no distinct primary region was observed for most of the test specimens. At the highest compressive stress of 630 MPa, the specimen failed in a manner similar to fast fracture, leaving many tiny
1000
. . . . , * . . NC132 SILICONNITRIDE
,
,
,
,
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8oo
,
16001
1
.
.
.
.
,
,
,
,
.
,
"2132 SILICON NITRIDE 1400 COMPRESSION1 1300°C 1200 1000 800
600 400
200 1000
500
0
n
1500
0
500 TIME, t [h]
TIME, t [hl 4000,
.
.
.
.
.
.
.
.
.
,
~ ~ 1 SILICON 3 2 NITRIDE UNI-FLEXURE/ 1300°C 3000
1000
800
162 M P a
1000
c^^
1 I
(4
/
NC132 SILICON NITRIDE BALL-ON-RING BIAXIAL/ 1300'C
/,lMPa /
2000 4
1000
n
2 1 00
0
2000
1000
..
TIME. t Ihl ~
2000
I000
2000
TIME, t [h]
NC132 SILICON NITRIDE RING-ON-RING BIAXIAL/ 1300°C
(4
loool& 500
J
Or 0
I
100
200 300 TIME, t [hi
400
500
Figure 2. A summary of creep deflection (displacement) as a function of time for different levels of applied loads, determined at five different specimedloading configurations for NC132 silicon nitride at 13OOOC: (a) pure tension; (b) pure compression; (c) four-point uniaxial flexure; (d) ball-on-ring biaxial flexure; (e) ring-on-ring biaxial flexure.
fragments. The data for this specimen were not used in the analysis. In spite of appreciable creep deflection in uniaxial flexure, no specimen failed before the interruption of the testing period, which ranged from 17 days for the highest nominal stress of 162 MPa to about 100 days for the lowest nominal stress of 29 MPa. Regardless of the nominal applied stresses, creep rate changed with time, so that no definite steady-state creep region could be determined. Since no explicit secondary steady-state creep region existed, a clear-cut transition between the primary and secondary regions was hardly definable. As in tension, uniaxial flexure, and ring-on-ring biaxial flexure testing, no test (disk) specimen in ball-onring biaxial flexure testing failed before the end of the testing period. The testing period ranged from 20 days for the highest nominal stress of 81 MPa to 100 days for the nominal stress of 41 MPa. At the lower nominal stresses of 41 and 29 MPa, the primary creep region seemed to be accompanied by the secondary, steady-state creep region. However, it is not clear yet whether the primary or the
secondary creep region exists in the curves for higher nominal stresses of 57 and 81 MPa. It was observed from the crept disk specimens that the loading ball was in nearly full contact rather than in ring contact with the test specimen surfaces, which was attributed to the increased plasticity and creep of test specimens during the tests. Hence, the use of the stress formula [ 121 based on the fullcontact assumption seemed to be justified in this context. However, the exactness of the stresses actually distributed in the disk specimens during creep testing is in question and consequently is another separate issue. The testing period in ring-on-ring biaxial creep ranged from 10 to 12 days for the lower nominal stresses I 57 MPa (but for 45 days for the highest nominal stress of 81 m a ) . Considering the relatively short test times, it would be premature to draw any conclusion regarding the creep curves obtained from stresses I 57 MPa. For the disk specimen tested at the highest nominal stress of 81 MPa, its displacement rate changed with time, making it difficult to defme the primary and the secondary steadystate regions.
293
0.04
.
- -
0
I
-
*
-
I
*
-
’
-
NC132 SILICON NITRIDE 130OoC/57MPa
0.02 UNIAXIAL FLEXURE
BALL-ON-RING BlAXIAL
0.00 0
lo00
500
1500
TIME, t [h]
Figure 3. A comparison of nominal creep strain as a function of time determined at four different specimedloading configurations at a nominal applied stress of 57 MPa for NC132 silicon nitride at 1300°C.
CREEP STRAIN-CREEP STRAIN RATE The equations to determine creep strains for the five different specimedloading configurations were all based on the idealized elasticity solutions. Therefore, if significant creep, plasticity, relaxation and/or a shift of neutral axis toward the compression side are present, the application of those equations to flexure loading may not be appropriate. As a result, the flexure strain and stress, calculated based on the idealized solutions, were termed in this paper ‘nominal) creep strain and ‘nominal’ applied stress, respectively. A typical example of test results for a given nominal applied stress is shown in Figure 3, where nominal creep strain was plotted as a function of time for different specimen-loading configurations, determined at 57 MPa. The overall nominal creep strain as well as its rate was highest for tension loading, intermediate for uniaxial flexure loading, and lowest for biaxial flexure (ball-onring and ring-on-ring) loading. (Of course, the lowest nominal creep strain would be for compression loading if compression data were available at 57 MPa.) The degree of constraint in creep due to the presence of compressive stress in flexure was increased from uniaxial flexure to biaxial flexure. The results of Figure 3 clearly indicate a significant difference in creep between tension and compression loading. Since, in general, there was no explicit secondary, steady-state region (except for compression loading), the actual nominal strain rate as a function of time was determined using a best-fit nominal strain curve and then taking the derivative of the strain curve with respect to time. Although many advanced ceramics have exhibited well-defined secondary creep region, there still exist many silicon nitrides that have not revealed defmite steady- state but that have exhibited a monotonic decrease in strain rate with time [13]. Particularly, a rapid decrease in strain rate from the earliest observations (
294
during creep. The hardening effect, a phenomenon of strain rate decreasing with time, or enhancement of creep resistance, is also manipulated as a result of grain contact during deformation [131. When creep strain curves exhibit neither steady-state nor tertiary region, the minimum creep strain rate at the end of the ‘extensive’ primary creep is often used synonymously as the steady-state rate. The decrease in nominal strain rate observed from these experiments, however, was not significant (even after 1500 to 2000 h) compared with cases in the literature in which the decrease in creep strain rate for some silicon nitrides amounted to about 2 orders of magnitude [13]. A decrease in nominal strain rate of about 1 order of magnitude was observed for cases of tension and uniaxial flexure at 57 MPa. Otherwise, the decrease was generally less than 1 order of magnitude. Therefore, a steady-state nominal creep strain rate for a given curve was approximated by averaging the nominal creep strain rates at two different times: at the start of testing, from 50 to 400 h, and the end of testing, from 220 to 2000 h. The choice of the start time, which was assumed to represent the initiation of steady-state region, depended on the loading condition and applied load. CREEP STRAIN-RATE VS. APPLIED STRESS Several models to describe creep strain rate as a function of applied stress have been proposed. In this section, some representative models such as power-law [141, hyperbolic sine [15,161, step [ 171 and redistribution [18] models will be reviewed and applied to the experimental data. A summary of nominal creep strain rate as a function of nominal applied stress using the four models is presented in Figure 4. Note that the data on ring-on-ring biaxial flexure were not included in this figure; they were produced during a test period of less than 300 h cannot be compared with data obtained over a longer period. There was no significant difference in curve fitting between the power-law, hyperbolic sine, and redistribution models. The hyperbolic sine model in particular yields the almost same result as the redistribution model. The power-law model shows a
.
10-5 10"
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(a)NORTON'S LAW
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IO-lll
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10-'0
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NOMINAL APPLIED STRESS, r~[MPal 10-5
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(d) REDISTRIBUTION M O D E L
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lo-"
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lo-"
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NOMINAL APPLIED STRESS, u [MPnl
10
100
1000
NOMINAL APPLIED STRESS, u [MPa]
4. A summarv of nominal creeD strain rate as a function of nominal amlied stress with different specimedloading confi&rations for NC 132*siliconnitride at 13OO0C,fitted to four differ&t creep models: (a) powerlaw (Norton) model; (b) hyperbolic sine model; (c) step model; (d) redistribution model.
somewhat worse fit to the tension data at high stress, not revealing the actual curvature in the data. Similarly, the step model also gives a worse fit to the tension data, but gives a reasonably good fit to other loading configurations. It should be noted that the models employed were developed primarily for pure tension loading. Hence, strictly speaking, it is questionable whether the models would be applicable for other than tension loading, since the creep mechanism in one loading configuration may vary in different loading configurations. In view of its mathematical simplicity, the conventional power-law model has been widely utilized for decades, regardless of the type of loading (tension, compression or flexure) for quantitative description of creep behavior of most advanced ceramic materials. Many analytical models also have been proposed on the basis of the power-law, constitutive relations. The stress exponent IPS in the power-law model, determined with the experimental data (see Figure 4a), were N = 1.72k0.34 in tension, 2.51k0.36 in compression, 2.01M.45 in uniaxial flexure, and 2.49f 0.28 in ball-on-ring biaxial flexure. The somewhat negligible difference in stress exponent N between the four loading configurations implies that the mechanisms associated with creep of this material would not be significantly different with any specimen-loading configurations. It has been observed that dislocation and cavitation mechanisms yield values of N = 3-6, diffusion creep andor viscous flow of the glassy phase yields the value of N = 1, and grain boundary sliding results in N = 2 [ 19,201. Cannon and Langdon [191 showed that many silicon nitrides (and silicon carbides) exhibited a stress exponent of N = 2 regardless of tension, '
compression, or uniaxial flexure. The mechanism was elaborated as grain boundary sliding with intergranular cavitation. Although the stress exponent was not significantly different between tension and compression, the difference in creep strain (and strain rate) between the two for a given applied stress was significant. The strain in tension was about 1.5 to 2 orders of magnitude greater than that in compression, if the compression data are extrapolated toward lower stresses. This difference in creep strain between tension and compression eventually results in a neutral axis shift toward the compression side of a flexure beam or disk specimen. It has been reported that flexure beam specimens of many ceramic materials have exhibited a neutral axis shift for equilibrium because of difference in creep strain, often accompanied by appreciable void andor cavitation formation on the tension side [ 1-3,21-231. Because of this asymmetric creep behavior, the data determined from simple uniaxialflexure creep testing are considered inappropriate for use as design parameters of structural ceramic components. A recent study [24] using the data generated in this work demonstrated that with individual tension and compression creep data and taking into account the neutral-axis shift, a reasonable prediction of creep deformation in uniaxial flexure was obtained with the CARESICreep design code. The same approach will be made for the case of more complicated multiaxial (biaxial) flexure in the next section. CARESICreep DESIGN CODE The CARESICreep integrated design computer program [4,5] predicts the service life of a monolithic
295
ceramic component as a h c t i o n of component geometry and loading conditions. The CARESICreep couples commercially available finite element programs, with design methodologies to account for creep rupture. The code is divided into two separately executable modules, CARESICWEST and CARESICreep, which perform: (1) calculation of parameters from experimental data using laboratory specimens; and (2) damage evaluation of thermo-mechanically loaded components, respectively. Finite element heat transfer and nonlinear stress analyses are used to determine the temperature and stress distributions in the component. The creep life of a component is discretized into short time steps, during which the stress and strain distributions are assumed constant. The damage is calculated for each time step based on a modified Monkman-Grant creep rupture criterion. Failure is assumed to occur when the normalized accumulated damage at any point in the component is greater than or equal to unity. The corresponding time will be the creep rupture life for that component. A schematic representation of the integrated design process is shown in Figure 5 [5]. The CARESICreep algorithm makes use of the nonlinear stress analysis capabilities of the ANSYS finite element program. Before building a model in ANSYS, the creep response of the material must be known. An input file containing these parameters is generated by the parameter estimation module of CARESICRPEST. This module is written in FORTRAN 77 and has as its input data h m creep tests. After the parameter estimation and nonlinear analysis has been completed, the second half of the CARESICreep program may be run. This module is executed from within the ANSYS program and is written in APDL (ANSYS Parametric Design Language). APDL routines usually take the form of an ANSYS macro which is a sequence of ANSYS commands recorded on a file for repeated use. By recording these commands on a macro, they can be executed with one ANSYS command. When this execution is completed, a damage map of the component is displayed in the graphics window. This map consists of a contour plot of the component's damage at the time when failure has taken place, or at any design life.
Figure 5. Block diagram for the creep analysis of a monolithic ceramic component using the CARESICreep design code
Figure 6. Fine element meshes employed in CARESICreep predictions.
1500
The prediction of multiaxial creep deformation in ball-on-ring biaxial flexure was made using the CARESICreep code, on the basis of the tension and compression creep data obtained. The fmite element mesh used in this analysis is shown in Figure 6. Parameter estimations were determined with the NortonBailey laws. The result of predictions is presented in Figure 7, where both the predicted and experimental displacements were plotted as a function of time at four different levels of nominal applied stress. The prediction somewhat underestimates the data. Notwithstanding the discrepancy, the overall agreement in displacement between the two seems to be reasonable, in view of general observations that some degree of scatter in creep deformation even at the same applied load has been common in many ceramics. Figure 8 shows the corresponding state of stress (or stress relaxation) occurring in the biaxial disk specimen subjected to a nominal applied stress of 28 MPa. Two different times of t = 0 and t = 1000 h were considered. Note that the
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750 41 500
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Figure 7. CARESICreep predictions of creep displacements for multiaxial stress state in ball-on-ring flexure at five different nominal stresses. The experimental data were included for comparison.
the CARESICreep design code, were in reasonable agreement with the actual experimental data.
Acknowledgements The authors are grateful to R. Pawlik for the experimental work during the course of this study. This work was sponsored in part by Strategic Research Fund Program, Glenn Research Center, National Administration of Space and Aeronautics, Cleveland, Ohio. I
(a) t = 0
(b) t = 1000 hr Figure 8. Change in the state of stress for the ball-onring disk specimen from time t = 0 to t = 1000 hr. The nominal applied stress was 28 MPa. Note that the neutral axis was shift toward the compression side at 1000hr.
neutral axis was shifted toward the compression side at 1000 h (with the redistribution of stresses), which is in good agreement with frequent observations for uniaxial flexure specimens of various advanced ceramics [ 1-3,21231. Generation of more extensive creep database as well as subsequent verifications of the design code are needed for better establishment of creep deformatiodife assessment methodology.
CONCLUSIONS Creep testing of a hot-pressed silicon nitride (NC 132 silicon nitride) was carried out using five different specimen-loading configurations at 1300'C in air. The specimen-loading configurations used included pure tension, pure compression, four-point uniaxial flexure, ball-on-ring biaxial flexure, and ring-on-ring biaxial flexure. Except for compression loading, nominal creep strain rate decreased gradually with time, resulting in less-defined steady-state condition. The magnitude and the rate of nominal creep strain for a given nominal applied stress were greatest in tension, least in compression, and intermediate in uniaxial and biaxial flexure. In flexure loading, uniaxial flexure was greater in creep strain than biaxial flexure. Four different creep models - power-law, hyperbolic sine, step, and redistribution models - were applied to the experimental data. All of the models gave a reasonable data fit, unclear which would be most appropriate for the NC 132 material. The conventional power-law model showed an overall stress exponent of about 2 (ranging from 1.7 to 2.5) irrespective of specimen-loading configurations. Predictions of creep deformation in response to multiaxial (biaxial) stress state, made utilizing
REFERENCES 1. S. M. Wiederhorn, D. E. Roberts, T.J. Chuang, and L. Chuck, Damage-Enhanced Creep in a Siliconized Silicon Carbide: Phenomenology, J. Am. Ceram. SOC.,71[7] (1988) 602-608. 2. M. K. Ferber, M. G. Jenkins, and V. J. Tennery, Comparison of Tension, Compression, and Flexure Creep for Alumina and Silicon Nitride Ceramics, Ceram. Eng. Sci. Proc., 11[7-81(1990) 1028-1045. 3. T. J. Chuang, Estimation of Power-Law Creep Parameters from Bend Test Data, J. Muter. Sci., 21 (1986) 165-175. 4. 0. M. Jadaan, L. M. Powers, and J. P. Gyekenyesi, Creep Life Prediction of Ceramic Components Subjected to Transient Tensile and Compressive Stress States, ASME Paper # 97-GT-3 19 (1 977). 5. L. M. Powers, 0. M. Jadaan, and J. P. Gyekenyesi, Creep Life of Ceramic Components Using a FiniteElement-Based Integrated Design Program, Trans. of ASME J. Eng. for Gas Turbines & Power, 120 (1998) 162-171. 6. G. D. Quinn and J. B. Quinn, Slow Crack Growth in Hot-Pressed Silicon Nitride, Fracture Mechanics of Ceramics, vol. 6, R.C. Bradt, A. G. Evans, D.P.H. Hasselman, and F.F. Lange, Eds., Plenum Publishing Corp., NY (1983) 603-635. 7. D. G. Miller et al., Brittle Materials Design, High Temperature Gas Turbine Material Technology, AMMRC CTR 76-32, Army Materials & Mechanics Research Center, Watertown, MA (1976). 8. R. K. Govila, Ceramic Life Prediction Parameters, AMMRC TR 80-18, Army Materials & Mechanics Research Center, Watertown, MA (1 980). 9. S. R. Choi and F. A. Holland, Silicon Nitride Creep under Various Specimen-Loading Configurations, NASA TM 210026-0, National Aeronautics & Space Administration, Glenn Research Center, Cleveland, OH (2000). 10. ASTM C 1291, Standard Test Method for Elevated Temperature Tensile Creep Strain, Creep Strain Rate, and Creep Time-to-Failure for Advanced Monolithic Ceramics, Annual Book of ASTM Standards, vol. 15.01, American Society for Testing & Materials, Philadelphia, PA (2000). 11. J. M. Birch, B. Wilshire, D.J.R. Owen, and D. Shantaram, The Influence of Stress Distribution on the Deformation and Fracture Behavior of Ceramic Materials under Compression Creep Conditions, J. Muter. Sci.,11 (1976) 1817-1825. 12. D. K. Shetty, A. R. Rosenfield, P. McGuire, G. K. Bansal, and W. H. Duckworth, Biaxial Flexure Test
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for Ceramics, Am. Ceram. Soc. Bull,. 59[12] (1980) 1193-1197. 13. D. S. Wilkinson, Creep Mechanisms in Multiphase Ceramic Materials, J. Am. Ceram. SOC.,8 1[2] (1998) 275-299. 14. F. H. Norton, Creep of Steel at High Temperatures, McGraw Hill, New York, NY (1929). 15. Nadai, The Influence of Time upon Creep; The Hyperbolic Sine Creep Law, in S. Timoshenko Anniversary Volume, Macmillan, New York, NY (1938). 16. R. M. Hazine and C. S. White, Multiaxial Internal Variable Modeling of the Creep Deformation and Fracture of an Advanced Silicon Nitride, Ceram. Eng. Sci. Proc., 18[3] (1997) 455-465. 17. C. W. Li and F. Reidinger, Microstructure and Tensile Creep Mechanisms of an In situ Reinforced Silicon Nitride, Acta Muter., 45[ 11 (1997) 407-421. 18. R. F. Krause, W. E. Luecke, J. D. French, B. J. Hockey, and S. M. Wiederhorn, Tensile Creep and Rupture Silicon Nitride, J. Am. Ceram. SOC.,82[5] (1999) 1233-1241. 19. W. G. Cannon and T. G. Langdon, Review: Creep of Ceramics. Part 1 Mechanical Characteristics, J. Muter. Science, 18 (1981) 1-50. 20. W. G. Cannon and T. G. Langdon, Review: Creep of
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21.
22.
23. 24.
Ceramics. Part 2 An Examination of Flow Mechanisms, J. Muter. Sci., 23 (1988) 1-20. J. A. Salem and S. R. Choi, Creep Behavior of Silicon Nitride Determined from Curvature and Neutral Axis Shift Measurements in Flexure Tests" Lge Prediction Methodologies and data 35.0 ceramic Materials, ASTM STP 1201, C. R. Brinkman and S. F. Duffy, Eds., American Society for Testing and materials, Philadelphia, (1994) 84-97. S. R. Choi and J. A. Salem, Creep Behavior of Silicon Nitride Evaluated by Deformation Curvature and Neutral Axis Shift Determination, Silicon-Based Structural Ceramics, B. W. Sheldon and S. C. Danforth eds., Ceramic Transactions,42 (1994) 285293. K. Jakus and S. M. Wiederhorn, Creep Deformation of Ceramics in Four-Point Bending, J. Am. Ceram. SOC.,71[10] (1988) 196-199. S. R. Choi, L. M. Powers, and J. P. Gyekenyesi, Creep Behavior of Silicon Nitride with Various SpecimenLoading Configurations, and Creep Prediction Using the CARESICreep Program, Proc. of the I" China International Conference on HighPerformance Ceramics (CICC-I), Tsinghua University Press, Beijing, China (1999) 612-615.
MATHEMATICAL MODEL OF MICROHARDNESS OF PLASMA SPRAYED CHROMIUM OXIDE COATING Chuanxian Ding, Jianfeng Li, Heng Ji Shanghai Institute of Ceramics, Chinese Academy of Sciences 1295 Dingxi Road, Shanghai 200050, China process parameters on the quality of the coatings [6-81.
ABSTRACT Thermal spraying involves a great number of
Because thermal spraying involves many process parameters, it is very difficult to investigate the
process parameters. The dependence of microhardness
dependence of properties of thermal-sprayed coatings
of plasma-sprayed Cr,O, coating on process parameters
on the process parameters by employing conventional
including plasma gas flow, current, spraying distance
experimental methods, such as the optimum seeking
and powder feed rate was investigated using the
method and the orthogonal design method. For the
experimental method of uniform design with ten grades
purpose of mathematical modeling, the thermal spraying
for each of the process parameters. The mathematical
process is often divided into three distinct regions: the
model of microhardness with respect to the investigated
torch, the spraying and the substrate. A number of
process parameters was formulated employing stepwise
assumptions are then made to simplify the mathematical
regression to a third order polynomial of the parameters.
representations
Finally, it was experimentally confirmed that the
phenomena occurring in these regions [6-91.
mathematical model demonstrated the relationships between
microhardness
of the
coating and
its
parameters.
INTRODUCTION Ceramic materials have received much attention in
of
complex
physical-chemical
K. T. Fang and Y. Wang developed the uniform design method, which is especially suitable for experimental research involving a great number of parameters. Compared with other experimental design methods, the uniform design method needs much less experimental time and can achieve almost the same
industry, such as in ceramic engines, due to their high
efficiency [10,11].
hardness, chemical stability, oxidation-resistance and
important property of thermal sprayed coatings [12]. In
thermal barrier properties at high temperatures. Their
this paper, the uniform design method with ten grades
high cost of production and brittle character, however,
for each of the process parameters was introduced into
will restrict the application of bulk ceramics in industry
plasma spraying and used to investigate the dependence
to a certain extent. For this reason, ceramic coatings
of microhardness of plasma-sprayed Cr,03 coating on process parameters including plasma gas flow, current, spraying distance and powder feed rate. The
Microhardness is a basic and
onto metal substrates, which are cheap and reliable in shock, are more widely employed [ 1-31. Thermal spraying, such as plasma spraying and
mathematical model. of the microhardness with respect
detonation spraying, is one of the most important
to the investigated process parameters was then
processes for ceramic coatings since coatings with
formulated employing stepwise regression to a third
thickness of 0.3 mm or more can be obtained. The quality of thermal-sprayed coatings primarily depends on powder characteristics and spraying conditions. The spraying conditions themselves are dependent on a large number of process parameters [4,5]. Experimental and
order polynomial of the parameters. Finally, it was experimentally confirmed whether the mathematical model demonstrated the relationships between the microhardness of the coating and the process parameters.
analytical research were conducted over the past three decades to control the quality of the coatings. In this respect, mathematical models offer a cost-effective and highly flexible means for analyzing the effects of
EXPERIMENTAL DETAILS Plasma sprayed chromium oxide coating was prepared using a Sulzer-Metco F4-MB gun mounted on
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an ABB S3 robot. A schematic illustration of plasma spraying is shown in Fig. 1. The starting powder had the particle size distribution indicated in Fig. 2, and the substrate was 1Cr 18Ni9Ti stainless steel dimensions of 6 0 x 3 0 ~ 2mm.
with
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Particle size (pm)
Fig. 2 Cumulative particle size distribution of the starting chromium oxide powder
Table 1 Scopes of experimental design for process parameters The smallest valueThe largest value Parameter Current (A) 480 660 Ar (L/min) 22 49 7 16 H, (L/min) Powder feed rate (g/min) 8.0 32.0 Spraying distance (mm) 80 170
2 -Cooling water 3 -Plasma gas Coating 6- Spraying distance 8-Plasma jet 9- Powder 10 -Anode I 1- Plasma an:
1-Cathode
4-Nozzle 57-Substrate
The microhardness of the coating samples were measured on their polished cross sections employing the same polishing process for all the samples. The measurements were conducted by indenting at a load of
Fig. 1 Schematic illustration of plasma spraying
In order to reveal the effects of important process
200 g for 15 s using an HX-1000 Vickers hardness tester
parameters on the microhardness of the coating, the process parameters including argon gas flow, hydrogen gas flow, current, spraying distance and powder feed
with an optical microscope which has a magnification of
rate were investigated. Scopes of experimental design for each of the parameters are listed in Table 1. The Table U*,,(108) of uniform design was employed to arrange the investigated process parameters as listed in Table 2. Each parameter was graded into ten grades, and the other process parameters were fixed as constants. Detail on the uniform design method are presented in
600x. For each sample, the measurement series was
comprised of 40 readings, which were randomly located on the cross sections and didn’t interfere with each other.
MATHEMATICAL MODEL The microhardness of the coating with respect to the investigated process parameters can be expressed by the following formula [ 131:
reference [ 111. Table 2 Sample codes and process parameters arranged using uniform design for plasma-sprayed chromium oxide coating Sample code Al, B1, C1 A2. B2. C2 A3. B3, C3 A4, B4, C4 AS, B5, C5 A6. B6, C6 A7, B7, C7 A8, B8, C8 A9, B9, C9 A10, B10, C10
300
Process code 1 2 3 4 5 6 7 8 9 10
Current (A) 520 580 640 480 540 600 660 500 560 620
Ar (L/min) 22 25 28 31 34 37 40 43 46 49
H* (L/min) 15 13 11 9 7 16 14 12 10
8
Powder feed rate (g/min) 18. 7 32. 0 16. 0 29. 3 13. 3 26. 7 10. 7 24.0 8. 0 21.3
Spraying distance (mm) 110 150 80 120 160 90 130 170 100
140
5
5
5
where F(x,, x2,..., XJ is the microhardness, x,, x2, ..., x5 stand for current, argon gas flow, hydrogen gas flow, powder feed rate and spraying distance respectively; and m is the order of polynomial for the investigated parameters. Once the microhardness of the coating samples are obtained with respect to the investigated process parameters, one can formulate the mathematical model by regressing the experimental results to the above formula. However, two aspects must be considered. On the one hand, thermal-sprayed coatings may present
inhomgeneities,
pores
andor
anistropic
properties [12]. As a result, measured microhardness data values typically appear rather scattered, and must be evaluated employing statistical estimation procedures. In the present study, assuming that the microhardness data values obtained for the coating samples follow the normal distribution, estimated means and confidence intervals were calculated with the standard deviation and sample size [13]. For each process code, an experimental microhardness value was obtained by taking the mean of the overlap of the two 95% confidence intervals which were the closest to each other among the three samples. On the other hand, the number of process codes, 10,
in Table 2 is too small and it is impossible to regress the experimental results directly to the formula (1) polynomial with an order more than 1. So the elements in formula (1) must be sifted to reject the elements which affect the microhardness negligibly and only retain those which significantly affect the microhardness for formulating the mathematical model. The methods available to sift the elements include the best subset regression, the backward elimination procedure and stepwise regression [13]. In this study, stepwise regression was used to model the experimental results of the uniform design method.
RESULTS AND DISCUSSION Fig. 3 gives the 95% confidence intervals of microhardness obtained with the standard deviation and
I 4OOJr 1
A5
I
I
I
2
3
4
, , , 5 6 7 Process code
,
8
, 9
,
10
Fig. 3 95% confidence intervals of microhardness for all the samples sprayed using process parameters of uniform design
sample size for all the coating samples sprayed using the process parameters listed in Table 1. This figure really verified that the microhardness of coating samples sprayed using the identical process parameters appeared rather scattered. The widths of the confidence intervals were in the range of 40 to 140, and most of them were between 50 and 90. From Fig. 3, the experimental values of the microhardness used to formulate the mathematical model can be obtained by taking the mean of overlap of the two 95% confidence intervals which were the closest to each other among the three samples. These results are listed in Table 3. Table 3 Comparison of experimental results of the uniform design method for microhardness with values calculated from formula (2) Process code 1
Experimental results Values from formula (2) 738 816 832 816 1006 1026 944 926 677 617 1198 1216 750 780 783 778 900 855
Table 4 exhibits the results of stepwise regressing th e e x p e r i m e n t a l v a lu e s o f m ic r o h a r d n e ss t o polynomials of the investigated process parameter with different orders. According to regression principles, the overall F can test whether the regressed equation is significant or nonsignificant [ 131. The values of overall F in Table 4 are more than F( 1, 8, 0.95)=5.32 and F(2,7,
301
Table 4 Results of stepwise regressing the microhardness to polynomials of the investigated process parameter with different orders Order Equation Source Sum of squares Degrees of freedom Mean square Overall F
microhardness of the coating and its process parameters,
0.99)=9.55, respectively. It can be concluded that the regressed first, second and third order equations are all significant [ 131. The residual mean square in Table 4 can provide a measure of the error with which microhardness could be predicted fiom given values of
the chromium oxide coating samples were sprayed using six sets of process parameters randomly selected among the scopes of experimental design in Table 1. The experimental microhardness results for the coating samples are compared with the values calculated fiom
investigated process parameters using the determined equation, i.e. if the process parameters are x,, x,, ..., x5, there would be probability of 0.95 that the predicted microhardness are among F(x,, x,, ..., x5)k.2x(residual mean square)”’ [13]. From Table 4, it can be hence concluded that the higher the order of the regressed equation, the smaller the error of predicted microhardness. For the third order equation, the predicted microhardness has an error of &2x20001n= k90, which nears the widths of the confidence intervals of microhardness in Fig.3. Therefore, the third order equation was adopted as the mathematical model of microhardness with respect to the investigated process parameters in this study and is listed as formula (2).
formula (2), as listed in Table 5. Table 5 indicates that the errors between experimental results and the calculated values were less than 90. Thus it is experimentally confmed that the mathematical model of microhardness for the chromium oxide coating expressed as formula (2) really demonstrated the relationships between the microhardness of the coating and the investigated process parameters. From formula (2), it can be seen that the microhardness of the chromium oxide coating was influenced by all current, argon flow, powder feed rate and spraying distance parameters. Moreover, the influences of current, argon flow and powder feed rate are interdependent. The interdependent influence can not be revealed by the conventional mathematical models which often divide the thermal spraying process into three distinct regions: the torch, the spraying and the substrate [6-lo]. Therefore, the mathematical model of plasma-sprayed chromium oxide coating formulated using uniform design experiments and stepwise
F ( x , ,x2)...)X, ) = 1140 - 4 . 4 5 ~ ~
+ 7.73
1 0 - X,XzX4 ~
(2)
Table 3 compares the experimental results of the uniform design method with the values calculated fiom formula (2) also lists in Table 3. It can be seen that the experimental results and the calculated values agreed with each other very well. In order to c o n f m whether formula (2) demonstrated
the
relationships
between
regression in this study can further the understanding of the relationships between the properties of the coating and its process parameters.
the
Table 5 Experimental test of the mathematical model expressed as formula (2) with process parameters randomly selected among the scopes of experimental design for the plasma-sprayed chromium oxide coating Ar H, Powder feed rate SDraving distance Values from ExDerimental Process code Current (A) (L/min) (mm) formula (2) results (L/iin) (ghin) .
I
I1 I11 IV V VI
3 02
660 660 660 660 650 660
48 48 48 48 40 40
16 16 16 16 13 12
32 32 32 32 32 27
*
v
80 100 120 140 110 110
1543 1441 1339 1238 1309 1193
1568 1479 1390 1301 1294 1202
CONCLUSION The mathematical model of microhardness for plasma-sprayed Cr,O, coating with respect to the process parameters including plasma gas flow, current, spraying distance and powder feed rate was formulated using experimental method of uniform design and stepwise regression. The mathematical
Abrasive Wear Performance of Cr3C,-25%NiCr Coatings by Plasma Spray and CDS Detonation Spray. Tribology Transactions, 38, (1995) 845-850. (5) E. Lugscheider, I. Rass and H. L. Heijnen et al, Comparison of the Coatings Properties of Different Types of Powder Morphologies. Thermal Spray: International Advances in Coatings Technology, C. C. Berndt, Rd., ASM International, (1992) 967-973.
model was expressed as a third order polynomial of the parameters. It was experimentally confirmed that the mathematical model demonstrated the
(6) D. R. Mash, N. E. Weare and D. L. Walker, Process
relationship between the microhardness of the coating and its process parameters. The mathematical
(7) M. Vardelle, A. Vardelle and P. Fauchais, Spray
Variables in Plasma-Jet Spraying. Journal of Metal, 13(7), (1961) 473-478.
Application in U. S. Thermal Spray Technology, I. Technology and Materials. Powder Metallurgy
Parameter and Particle Behaviour Relationships During Plasma Spraying. Journal of Thermal Spray Technology, 2( I), (1 993) 79-9 1 . (8) H. Fukanuma, A Porosity Formation and Flattening Model of an Imping Molten Particle in Thermal Spray Coatings. Journal of Thermal Spray Technology, 3( l), (1 994) 33-44. (9) S. Das, V. K. Suri and U. Chandra et al, OneDimensional Mathematical Model for Selecting Plasma Spray Process Parameters. Journal of Thermal Spray Technology, 4(2), (1995) 153- 162.
International 3, (1991) 147-155.
(1 0)K. T. Fang, Uniform Design and Tables of Uniform
model revealed that interdependent influences on the microhardness existed between the process parameters, and it can further the understanding of the relationship between the properties of the coating and its process parameters.
REFERENCES (1) R. W. Smith and R. Novak, Advances and
(2) G. Barbezat, A. R. Nicoll and A. Sickinger,
Design, Science Press, Beijing, (1 994).
Abrasion, Erosion and Scuffing Resistance of
(1 1)K. T. Fang, Y. Wang, Number-Theoretic Methods
Carbide and Oxide Ceramic Thermal Sprayed Coating for Different Application. Wear, 162-164, (1 993) 529-537. (3) J. E. Fernandz, Yinglong Wang and R. Tucho et al, Friction and Wear Behaviour of Plasma-Sprayed Cr,O, Coatings Against Steel in a Wide Range of Sliding Velocities and Normal Loads. Tribology International, 29(4), (1996) 333-343. (4) G. Barbezat, A. R. Nicoll and Y. S. Yin et al,
in Statistics, Chapman and Hall, London, (1993). (12)C. K. Lin, C. C. Berndt, Statistical Analysis of Microhardness Variations in Thermal Spray Coatings. Journal of Materials Science, 30, (1995) 1 1 1-1 17. (13)N. R. Draper, H. Smith, Applied Regression Analysis, John Wiley & Sons, New York Chichester (1981)
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LIFE TIME PREDICTION MODEL FOR PLASMA-SPRAYED THERMAL BARRIER COATINGS BASED ON A MICROMECHANICAL APPROACH R VaSen, G. Kerkhoff, M. Ahrens, D.Stover ForschungszentrumJulich GmbH, Institut fur Werkstoffe und Verfahren der Energietechnik 1, D-52425Jiilich, Germany
ABSTRACT A typical thermal barrier coating (TBC) system consists of a dense bondcoat layer (e.g. an oxidation resistant NiCoCrAIY-alloy) and a porous zirconia stabilized with 7 - 8 wt. % m a . A frequently used method to produce these coatings is the plasma spraying process. In order to obtain sufficient adherence of the ceramic top coat to the bondcoat a certain bondcoat roughness (Ra 5-8 pm) is necessary. It is now found that during thermal cycling operation these systems often fail within the top coat by crack initiation and propagation close to the bondcoat - top coat interface. This type of failure mechanism can be explained by stresses arising from the formation of a thermally grown oxide (TGO) layer on the rough bondcoat surface. TGO formation, creep effects in both bondcoat and top coat and roughness of the bondcoat lead to a rather complex stress state. All these factors have been taken into account by using a finite element method (FEM) to calculate stress development during thermal loading. These data were then introduced into a crack propagation model to estimate crack development during thermal cycling operation.
INTRODUCTION Thermal barrier coating (Tl3Cs) systems are often applied to metallic components in gas turbines to improve the performance of the engines. Without a change of the turbine inlet temperature (TIT) the use of TBCs leads to a reduction of the metal temperature and hence to an increase of the components life time. Alternatively, a TBC system enables an increase of the TIT and, as a result, a higher efficiency of the engine
PI.
TBC systems consist in most cases of a bondcoat and a 7-8 wt. % yttria stabilized zirconia topcoat. The bondcoat protects the substrate from oxidation and enables the bonding of the topcoat. Two different techniques to produce the top coat are widely used. One process is the electron beam physical vapor deposition (EB-PVD). For these EB-PVD topcoats Ptaluminide or also peened MCrAlY-bondcoats (M=Co, Ni) with a rather smooth surface are often used. The second process is the atmospheric plasma spraying (APS). In most cases the APS-TBC is deposited on MCrAlY bondcoats with a rather rough surface (5-7 pm) produced by vacuum plasma spraying (VPS). More details are found in several review articles [2, 3, 4, 5 , 6 7 1 .
The use of TBC systems in industrial applications is connected with the availability of a reliable life prediction model which is essential for the design of turbine engines. At present, life time models are mainly based on empirical approaches [8, 9, 10 ,11 ,121 which need the knowledge of the rather complex and partially difficult to determine thermo-mechanical data of the specific TBC system. In the present investigation a life time model for plasma-sprayed TBC systems which is based on their microstructure is presented. We assume that this model can lead to more reliable predictions and to a better understanding of the system. The improved understanding should lead to the identification of the key factors controlling the life of the TBC system. Detailed strategies for an optimization of the system are expected. In order to develop a microstructural based life time model the stress state in the plasma-sprayed TBC system is evaluated. As it is generally accepted that the formation of the thermally grown oxide (TGO) during high temperature exposure is a significant factor which determines life time of the system it is included in the model. Furthermore, the roughness of the bondcoat as well as the curvature of the substrate are considered. The results of the stress evaluation are then used to develop a model which describes the growth of the microcracks in the TBC. The assumptions made in the model and the predictions are compared with experimental results from thermal cycling tests.
EXPERIMENTAL Thermal cycling tests have been performed with disk-shaped specimens with 30 mm diameter and 3 mm hckness. The outer edge was beveled to reduce stresses at the outer free edge of the sample [13]. One side of the samples was coated with a 150 pm thick NiCoCrAlY bondcoat by VPS using an F4 gun in a Sulzer Metco plasma spray unit. A typical roughness value of 6.4 pm was found for most of the produced coatings. Some of the samples have been pre-oxidized before cycling at 1000 "C for 100 h in air. In this case the bondcoat roughness varied between about 4.9 and 7.2 pm. The 8 wt. % YSZ topcoat was applied by an APS process with a Triplex gun in a Sulzer Metco APS facility. The thickness of the coatings were about 300 pm, the porosity was between 12 and 13 %. Thermal cycling was performed in a gas burner test facility which operated with natural gas and oxygen. The specimens were heated from the TBC side while
305
cooled with compressed air from the back. After the heating period of 5 min the burner was moved away from the surface and the sample was also cooled from the front with compressed air for 2 min. In addition to this standard cycle a shorter cycle with about half the heating and cooling time was used for specimens preoxidized prior to cycling. The surface temperature was measured with a pyrometer operating at wavelengths between 8 and 13 pm. In the middle of the substrate a 1 mm hole was drilled to the center of the substrate. In this hole a thermocouple was fixed to measure the substrate temperature. In the experiments it took about 20 s to reach the maximum surface temperature. The steady state surface temperature during the heating phase was about 1250 "C, the temperature measured at the thermocouple was about 970°C which corresponded to about 1000 "C at the interface. The lengths of 50 cracks in the TBC were measured at three different locations. The exact position of the areas, which had a width of 50 pm, are shown in Fig. 1. The position of the investigated areas were in the middle of the TBC 70 pm away from the interface (1 in Fig. 1) and close to the interface (2). Additionally, crack sizes at the beveled edge were measured (3). From the area which had to be evaluated to measure 50 cracks the crack density was calculated. I
i 3
As we did not observe segmentation cracks in the coatings after thermal cycling, relaxation by sliding along the interfaces of the spray lamellae, where a lot of microcracks are located, might occur. As no reliable data of this high temperature relaxation are available we assume that the relatively high stress levels are build up in the coating and are relaxed by creep mechanisms. These calculations performed with literature data for the creep behavior of TBCs give a fast relaxation of the high temperature stress state. These relaxation leads to a compressive in-plane stress in the coating after cooling down to room temperature. In fact, we measured with X-Ray diffraction techniques, that the surface stress in the TBC of a cylindrical specimen was reduced from about 5 MPa to about -20MPa after one cycle at 1000 "C for 2 h. These results made it reasonable to assume that the high-temperature stress level is very low after a limited number of thermal cycles. For a simplification of the discussion we will assume a stress-free TBC at high temperatures. At a later state of the model development, when precise relaxation data of TBC and bondcoat are available, these effects can be included. They will lead to lower stress levels in the TBC at room temperature. The first parameter which we will discuss is the curvature of the substrate. A cylindrical geometry was used as shown in Fig. 2.
i
i
i
bondcoat ( N i C o C r A l w substrate (IN734
I
Fig. 1 Schematic drawing of a part of the thermal cycling test specimens. The three locations, where the crack lengths were measured, are indicated ( 1-3)
MODEL DESCRIPTION For the development of our model it is essential to know details about the stress states within a plasmasprayed TBC. The calculations have been performed with the finite element code ABAQUS using materials data which are published elsewhere [14]. Of special importance for the understanding of the stress development in a TBC system are the thermal expansion coefficients. These are high for the substrate and the bondcoat (> 15.10-6 K),intermediate for the YSZ topcoat (10 - 11.10-6 K ) and rather low for the thermally grown oxide made of alumina (8-10'6 K). From literature data and also from our own XRD stress measurements it is known that the stress state in as-sprayed TBCs is rather low, typically of the order of 10 MPa or less [15]. If such a system is heated to elevated temperatures, a tensile stress state will be found in the TBC. For A T of 1000 K an estimation will give an approximate value of 200 MPa in-plane stresses. Probably the ceramic can not withstand such high tensile stress without some kind of relaxation.
306
bondcoat thickness 150 pm Fig. 2 Cylindrical geometry used in the finite element calculations. Cooling down such a cylindrical system from the stress-free high temperature state will result in in-plane compressive stresses and, more important, radial tensile stresses in the TBC. The origin of these radial stresses is the larger thermal expansion coefficient of the substrate compared to the TBC, as it is also the case for the in-plane stress. The inner metallic cylinder will contract more during cooling than the outer cylinder and hence, a radial saess will develop at the interface. As shown in Fig. 3 the stress level depends on the radius of the substrate, which is inversely proportional to the curvature. The higher the curvature the higher the stress level. It is also interesting to note that a thicker TBC will increase the stress level nearly
linearly. This increase can be explained by the higher stiffness of the thicker TBC.
a
Fig. 3 Radial stress levels in a cylindncal TBC after cooling down from 1000 "C to room temperature.
k
I
-100 1 -.J
Fig. 4 Stress levels within the TBC at the height of the peaks of the bondcoat roughness after cooling down from 1000 "C to room temperature without TGO and with 3 pm or 9 pm TGO. Additionally, the profile of the bondcoat roughness is shown. The next important parameter is the roughness of the bondcoat. As a result of this roughness regions of high positive and negative curvature exist at the bondcoat/TBC interface. In these regions stresses develop for the same reasons discussed for a substrate with cylindncal geometry. We choose for our calculations a sinusoidal profile as a description of the roughness profile. In Fig. 4 the stress levels (after cooling down) at the height of the peak location within the TBC are shown. At the peak locations high tensile stresses are found whereas at the valley locations compressive stresses exist. After the growth of the TGO to a certain thickness, which depends on the profile of the bondcoat roughness, the stresses in the TBC change their sign. They become compressive at the peak and tensile at the valley locations.
f'
BC+substrate
II-
BC+su bstrate Fig. 5 Four different phases describing crack growth and failure of thermally cycled APS -TBCs. The presented results on the stress development are used to develop a model for crack propagation and failure of plasma-sprayed TBCs. The basic ideas of the model are outlined in Fig.5. At an initial phase high stress levels induced by the roughness of the bondcoat exist at the peak locations (Fig. 5a).
307
They lead to a crack extension of the existing microcracks between the spray lamellae. As the interfaces between the spray lamellae close to the BC/TBC interface are similar shaped as this interface, the cracks will run into the valley locations. Due to the compressive stress level they will be arrested there. We propose that crack propagation is mainly controlled by the cyclic nature of the loading. In this case the crack growth per cycle dc/dN can often be described by the following empirical law:
edge. A variation in the amount of spallation existed between different specimens, however no systematic difference between samples cycled with short and standard cycles could be found. These results correspond with the results of crack size measurements. Samples with short and long cycles have similar crack lengths after 400 cycles (Fig. 6). These results c o n f i i the assumption that crack growth is determined by the cyclic nature of the thermal loading as proposed by equation (1). shortcycle long cycle
(1) dN E
E
AKI is the difference in stress intensity factors during the cycle, K I , ~is the critical stress intensity factor, and A, n are constants. This situation will change in the second phase when considerable TGO growth takes place (Fig. 5b). As stated above, the stress at the valley locations will be converted to tensile and cracks can grow further. Long cracks can be formed during this state by coalescence of cracks running from different peak locations into the same valley. Once these long cracks have been formed, the curvature of the substrate has to be considered (Fig. 5c). Stress levels due to typical curvatures existing in turbine blades (< 0.5 /mm) are below 30 MPa (s. Fig. 3). For the largest initial crack lengths of about 100 pm (s. below) the stress intensity factor can be estimated to by using the simple equation be below 0.4 derived for a crack of length c in an infinite plane [16]:
AK, =
451 center in TBC
T
center at interface curvature at interface
Fig. 6 Crack length of as-sprayed, pre-oxidized and cycled (400 cycles) specimens at the different locations. 14
1 12
--
---I+as-sprayed\
$10
$ *
w-
0
E
z4
As this value is considerably lower than the critical stress intensity factor in TBCs (about 1 - 2 MPamO.’ [17, 18]), we expect only a minor contribution of the substrate curvature induced stresses to crack growth during the initial phase. However, due to crack growth during phase 1 and 2 large cracks with crack size of 200 pm and longer are formed. The stress intensity factor will increase considerable (> 40 %), fast crack growth might be observed for the longest cracks. In the final phase (phase 3) the crack will reach a length at which spallation of the coating will occur (Fig. 5d). This crack length c can be estimated using elasticity theory [19].For a coating with thickness D, Young’s modulus E, Poisson ratio v under tangential compressive stress (T the critical crack length c is:
)
I/Z
ETBc
(1-2)ff
(3)
RESULTS AND DISCUSSION The cycling of the pre-oxidized specimens was stopped after 400 cycles. All investigated specimens showed spallation of the TBC in the area of the beveled
308
IOOh, 1000°C
$ 6
c+g
c = 2.21 D(
v
75
€
2
0 0
100 crack size [pm]
50
150
Fig. 6 Crack size distribution of as-sprayed and thermally cycled specimens at the beveled edge location (standard cycle). Samples in the as-sprayed condition and after 500 and 1080 cycles without pre-oxidation have additionally been evaluated. Only the sample cycled for 1080 cycles showed spallations in the region of the beveled edge. The crack size distributions of samples after 0, 500, and 1080 cycles measured at the beveled edge are shown in Fig. 6. The crack size distributions are rather broad with a number of large cracks up to 100 pm also in the as-sprayed coating. During cycling the distribution shifts to larger crack lengths. An increasing amount of cracks above 100 pm is found. This tendency is also obvious in Fig. 7, in which the mean crack lengths at the three different locations are shown. Only the cracks in the center within the TBC show a decrease of crack size with increasing cycle numbers. This might be a result of the sintering of the cracks. Close to the bondcoat interface the cracks grow
as predicted in our model as a result of the roughness induced stresses. The beveled edge, which corresponds to a substrate curvature of about 1.2 - 2 mm, further promotes crack growth (s. phase 3 of our model). Correspondingly, the largest cracks and the first occurrence of spallation is found in this area.
1
loo
+
B
s E
6oj 50
!
centre at interface curved interface
structural consideration of the stress state and is able to explain the influence of important micro-structural parameters ldce bondcoat roughness, TGO growth, substrate curvature and others on the life time of a TBC coating. In thermal cycling experiments a major assumption of the model, which is that crack growth is dominated by the cyclic nature of the thermal loading, found a first justification. The predictions of the model regarding the importance of bondcoat roughness and substrate curvature for the crack growth in plasmasprayed TBCs were also in accordance with the experimental results.
ACKNOWLEDGEMENT
i
I
0
500
1000
number of cycles
Fig. 7 Mean crack length at three different locations of the TBC (s. Fig. 1) of as-sprayed and thermally cycled specimens (standard cycle). In Fig. 8 a micrograph of the sample after 1080 cycles at the beveled edge is shown. The long cracks close to the interface within the TBC are obvious. Additionally, the dark TGO layer and a P-phase depletion in the bondcoat close to the TBC can be seen.
-1
*
'
*-
Fig. 8 Optical micrograph of a sample at the beveled edge after 1080 standard cycles. Additionally to the crack size the crack densities were evaluated. The number of cracks per area was between 650 and 850 cracks I mm2 in the as-sprayed coatings for all locations and reduced to 600 to 700 cracks / mm2 after 1080 cycles. Also this result can be explained by our model. The stresses close to the bondcoat will lead to crack growth of existing cracks instead of originating new ones. The coalescence of these cracks will result in a slight reduction of the crack density.
CONCLUSIONS A model was developed which describes the failure of plasma-sprayed thermal barrier coatings during thermal cycling. The model is based on micro-
The authors would l k e to thank K.H. Rauwald and R. Laufs for the preparation of the coatings and the performing of the thermal cycling experiments. The metallurgical specimen preparation of M. Kappertz and the crack sue measurements of Ms. S. Schwartz are also gratefully acknowledged.
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Developement, Phase I Final Report, Contract NAS3-23944, NASA CR 182230 (1989). 12 S.M. Meier, D.M. Nissley, K.D. Sheffler: Thermal Barrier Coating Life Prediction Model Development, Phase 11 Final Report, Contract NAS3-23944, NASA CR 189111, (1991). 13 C. Funke, J.C. Mailand, B. Siebert, R. VaBen, D. Stover: "Characterisation of ZrO2 - 7 wt?? Y2O3 Thermal Barrier Coatings with Different Porosities and FEM Analysis of Stress Redistribution During Thermal Cycling of TBC's" Suflace Coatings and Technologies, 94-95 (1997) 106-111. 14 G. Kerkoff, R. VaSen, D. Stover, Numerically calculated oxidation induced stresses in thermal barrier coatings on cylindrical substrates, European Federation of Corrosion Publications 27 (1999) 373-382. 15 G. Kerkhoff, Thesis, Ruhr Universitiit Bochum, Germany, (2000). 16 H.L. Ewalds, R.J.H. Wanhill, in Fracture Mechanics, Edward Arnold Publishers, London (1986) p. 32. 17 G.K. Beshish, C.W. Florey,F.J. Worzala, W.J. Lenling , Fracture Toughness of Thermal Spray Ceramic Coatings Determined by the Indentation Technique, Journal of Thermal Spray Technology 2, 1 (1993) 35-38. 18 R. Vassen, X. Cao, F. Tietz, D. Basu, D. Stover, Zirconates as New Materials for Thermal Barrier Coatings, to be published in J. Am. Ceram. s0c.(2000). 19 S. Timoshenko, J.M. Gere, Theory of Elastic Stability, Engineering Society Monographs, New York (1936).
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ANALYTICAL DESIGN AND EXPERIMENTAL VERIFICATION METHODS OF CERAMIC RADIAL, TURBINE ROTORS FOR A GAS TURBINE Norio Nakazawa Tsukuba College of Technology, Tsukuba-Shi,Japan ABSTRACT Radial turbine rotors were designed and evaluated for a prototype automotive gas turbine using several test rigs with a turbine inlet gas temperature of 1350°C. One of the most important goals of the gas turbine engine development was to demonstrate 40% thermal efficiency without dependence on any cooling systems for the turbine components. This program was conducted by the Petroleum Energy Center in japan.'')^'^) The materials of the rotors were Si,N4, and test pieces including those cut from actual rotors were evaluated through flexural and tensile tests for fast fracture and long-term life. Several types of rotors were designed and manufactured with the goal of improving strength reliability and aerodynamic performance. FEM models were produced to calculate temperature and elastic stress distributions, and the design procedures with probability calculations were then applied to analyze reliability and life span using the material properties. Many rotors were evaluated through cold and hot spin tests which included burst tests and these results were correlated with the theoretical analyses in order to verify the effectiveness of the manufacturing processes and the design, and also to provide feed back for the next step in design evolution. Some types of rotors were proven by hot spin and endurance tests and then assembled into the engines. Failures due to blade resonance vibrations were observed and analyzed in combination with the experiments to modify the blade design. Foreign object impact damage was also encountered in the engine tests and turbine test rigs. Trap systems were designed to be placed upstream of the turbine. These were installed in the engines and the test rigs and demonstrated their effectiveness. This work is partially supported by the Agency of Industrial Science and Technology (AIST) in the Ministry of Trade and Industry (MITI).
which were assumed to be included in the materials during the manufacturing processes, as well as limits due to creep phenomena, were considered in the design. In designing for static fatigue life, reliability analyses using statistical calculations in regard to material properties, such as the Weibull modulus and static fatigue exponents, were combined with elastic stress analysis. Defect sizes and their distributions in the actual components depend on their dimensions and shapes, and then on the manufacturing process, Therefore the strength and life of the components vary as a result of them. These variations were measured and reflected in the design. Creep phenomena of the silicon nitrides adopted for the components seem much different from those of super alloys. What occurs in the materials' microstructure is that the sintering additives in the grain boundaries move to the material surfaces and are oxidized. Pores in the boundary due to these additives' movements are newly generated defects that influence the strength of the material. These phenomena are not yet fully understood and it is difficult to determine an appropriate design method in the light of them. Therefore we avoided designing components which include these creep conditions. Figures 1 and 2 show the rupture lives of the silicon nitrides we tried to adopt for the turbine rotors.('),(*) SN88M indicates failure by static fatigue at up to 1200"C, and creep failure is indicated beyond that temperature. On the other hand, SN91 shows creep from 1100°C and its static fatigue exponent at 1200°C is very low. Rupture stress of SN 91 drops at a value of the Larson-Millar Parameter which is less than that of SN88M by about 5. Thus we chose not to adopt this material for the rotors. SN88M (2) 81038 01150 m1350'C
DESIGN AND EVALUATION METHOD 0.001
Long-term life A design goal of continuous running hours was established for the turbine under steady state operation at the rated engine conditions. Limits in rotor life due to cracks originating from small defects (static fatigue)
001
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iooo
ionno
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Fig.1 Long-Term Strength of Silicon Nitrides
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-Creep Failure ---- Static Fatigue Failure
(Material Temperature) OSN88M 12007: VlO387: ASN88M 1300 0 1 1 5 0 0 SNSlM 1400 0 1350 OSNg1
ASN91 SN91
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TTemperature (K) Tr:Time (hour) 30035
40 45 50 55 Larson Millar Parameter M=[T(3O+log T r ) ~ l 0 - ~ ]
Foreign Object Impact Damage There is a relatively higher possibility of generating foreign objects in ceramic gas turbine engines than in conventional metallic engines, because their construction usually requires many insulating layers, and these layers and stationary ceramic components have a much greater tendency to chip and delaminate. The prevention of objects from flowing into the passage leading to the rotor, and improving the impact strength of the blade were both considered in parallel.
-
Fig.2 Silicon Nitride SN88M, SN91 Creep Rapture Strength
Strength Reliability under Transient Stresses In gas turbine engines, stress changes due to starts and stops generally impact the rotors' reliability significantly. Thus we analyzed their transient thermal stresses combined with centrifugal stresses in terms of their possibility to induce failure.
Connecting Interface between the Rotor and the Shaft The connecting interface of the ceramic rotor to the metallic shaft was designed by predicting its contact surface pressure, stress, and temperature distributions to meet their allowable limits, and also to be able to transmit the required power (torque). This analysis covered all the steady state and transient running conditions taking into account the centrifugal stress combined with thermal stress caused by the thermal expansion difference between both materials.
Blade Strength under Resonant Vibrations Rotor blades experience high-cycle vibrations caused by aerodynamic exciting forces included in periodical velocity distributions generated at the upstream vanes of turbine stator nozzles (nozzle wakes). In the initial development stage, we observed rotor breakage due to these vibrations. In designing blades made of super alloys, engine operations at speeds that induce resonant blade vibrations with frequencies of aerodynamic exciting forces are usually analyzed. The safety of these operations is analyzed using data describing the material's fatigue strength and internal damping properties. Silicon nitrides have higher stiffness and lower densities than do super alloys, so a blade with the same profiles and dimensions as a super alloy blade has nearly 1.9 times higher natural frequencies than does the super alloy blade. Therefore its resonant operating points shift to a higher speed range with increased aerodynamic exciting forces and blade centrifugal stress. In addition, some measured data regarding the internal damping properties of silicon nitrides indicate far smaller values compared with those of metallic materials. Metallic materials also seem to demonstrate favorable behavior by exhibiting increasing damping properties with increasing stress exceeding their elastic limits, while restricting the rapid stress increase. Silicon nitrides demonstrate no such characteristics. At resonant speeds, the vibration stresses of silicon nitride blades increase in value and cause breakages such as fast fractures much like those induced by a static load. Because sufficient knowledge of material properties under high-cycle vibrations had not been obtained, we shifted the design policy and chose to avoid these resonant running points within the engine's operational range.
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TURBINE SPECIFICATIONSAND DEVELOPMENT GOALS The final turbine specifications and engine conditions in view of those obtained at the middle of the program are the following: Turbine Inlet Temperature (TIT); 1350°C Rotational Speed; 100000rpm * Rotor Blade Tip Speed; 665mlsec. Turbine life; 100 hours Figure 3 shows the turbine design, which had progressed, from the 1st to 6th models in order to improve both aerodynamic performance and reliability. The 6th rotor design is shown in Figure 4.
-
-.%6
Fig.3 Radial Turbine Structure
ScallopDiameter
Fig.4 6th Design Rotor
I
DESIGN AND DEVELOPMENT TASKS FOR ROTATIONAL STRENGTH AND LIFE Tasks and Development Steps Temperature and elastic stress were first analyzed after generating the FEM model based on the aerodynamic design data. By applying these temperature and stress distributions to reliability analyses using strength properties of the material test pieces, the average life and rupture speed, and failure probability at a rotational speed can be calculated. These analyses also show the rotor's failure probability distributions and the most probable point of failure. Design modifications after reviewing the results of these analyses were repeated until the design goals were achieved. In parallel with the design process, the design's required manufacturing technology was also evaluated and incorporated into the design. A manufactured component's strength and reliability depends on the efficiency and accuracy of the manufacturing processes. In order to evaluate the manufacturing processes, the following characteristics and materials were tested. 0 Strength of test pieces cut out from blocks manufactured together with the rotors. @ Strength of test pieces cut out from manufactured rotors. @ Rotating burst strength of the rotors The burst tests were conducted by spinning at room temperature (cold spin tests) and at high temperatures (hot spin tests) corresponding to the rated running conditions. The cold spin tests were carried out in order to study optimum factors in manufacturing processes for the initial phase of manufacturing together with test piece evaluations by After establishment of the processes, the cold spin tests were performed for quality assurance of the manufactured rotors. The hot spin tests provided short-term burst strength data under the same temperature conditions as those performed in the engine tests. These data were compared with the analyzed values and the burst strength data from the cold spin tests, and the design and manufacturing processes were then reviewed. The hot spin tests were also conducted to confirm the rotor's reliability (proof tests) before delivering them to the next development step which included endurance tests and engine performance tests.
analysis to obtain failure probabilities, and linking these calculation codes was considerably effective. Centrifugal - and thermal stress analyses, and vibration analysis (natural frequencies, displacement and stress distribution of blade) had to be repeated together in order to determine the optimum blade thickness distribution and hub profile.
Fig.5 FEM Model (Single Blade Segment)
Failure'ProbabilityAnalysis The criteria dealing with rupture due to static fatigue were applied to each finite element and the failure probability of the element was calculated. In this kind of analysis code, using the Weibull two-parameter distribution function, various rupture criteria can be included (using a calculation code such as CARES [Ceramic Analysis and Reliability Evaluation of Structures])(3). The basic input data of the materials under analysis are the four-point flexure test results such as the Weibull plots shown in Figure 6 and their dependence on duration and temperatures.
a,@.
Design and Evaluation Method Temperature, elastic stress, and vibration analysis The FEM analysis model was made for a single blade segment of a rotor with 14 blades as shown in Figure 5. This FEM model was used in temperature/stress/vibration analyses and reliability
500
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900 lo00
Flexural Strength (MPa)
Fig.6 Four-point Flexural Strength (SN88M)
Failure probability distributions per unit volume of the rotor at various operating conditions can be calculated by using temperature-stress data and material properties as shown in Figure 7. This kind of analysis and evaluation is useful to check for excessively high values of failure probability and to improve the design. The failure probability of the rotor as a whole obtained as a function of Weibull modulus by multiplying the failure probability values of all elements is shown in Figure 8 under various engine running and test conditions. By plotting the average burst strength (speed) test data on this figure, the Weibull modulus of a population of the tested rotors can be obtained, and the reliability or quality level of the population can be
313
estimated. These results were useful for verifying the manufacturing processes and also for modifying the design. v
1
1st Design Rotor
Rated Conditions (TIT=1350"C ,100000 m),
300hour Running, Weibull Modulus
m=8
Fig.7 Failure Probability Distribution per Unit Volume (l/mm')
Cold Spm Average
15 Hot Spin test Results
12'9x10
-15
Ti"
t,4Ro&kalysii
*3
!TIT=I200 Average 11.7xlO'rpm
2x v
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lo Hot S p i n t e s t i k
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%T=1350
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TITTurbe litlet Tempcrature F :FailUte Pmbability
I
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The correlation between the rotor burst strength and duration was also predicted in the same manner as that mentioned above. The predicted results are shown in Figure 9 together with test conditions and the design goal. The relationship between the burst data of the hot spin tests as short-term (around 0.Olhours) strength and the long-term life could be clarified by comparing the analytical results shown in this figure. By studying the test results and predictions shown in Figure 9, proof test conditions, in which the over-speed test was performed for a duration of 0.lhours before the rotors would be moved on to the engine performance and endurance tests, could be determined. Hot Spin Burst Results (TIT=1350 "c) 0 Endurance Test (TIT= 1350"c.without Failure) 0
.-
Prediction (TIT=1350"c,m=15) m:Weibull Modulus FFailure Probability
(TIT=1200 "c,F=1/100,m=l5)
0.01
0.1
1
Prediction flIT=1350"c.F=lR)
10
Life (hour)
100
Fig.9 Life Prediction and Test Results of 1st Design Rotors
3 14
Test Results of the Rotor's Initial Design Numerous burst test results were obtained for the rotor's initial design, with the Weibull modulus becoming 15-17 as shown in Figure 8. One rotor of the last manufacturing lot attained a maximum speed of 135000 rpm (900 m/sec. tip peripheral speed). The result of an endurance test of 100 hours of another rotor of the initial design completed in the turbine component test rig is shown in Figure 9. 6th Rotor Design The purpose of the 6th Rotor Design was to improve aerodynamic performance and to increase blade stiffness in order to avoid resonant vibrations caused by stator nozzle wakes. We incorporated more complex three-dimensional curvatures into the blade design in order to improve the internal flow, and also made the outer diameter of the hub disk (blade root scallop) larger in order to increase the blade stiffness. This design brought about an increased blade and hub stress level, but the maximum stress in the rotor at the hub center could be limited below the allowable value. The maximum stress at the hub center governs the rotor's strength reliability, and we set 5% as an allowable stress increase from that of the rotor's initial design. Based on the test results (Figures 8 and 9) that we had accumulated up to that point, we thought that a durability level of 100 hours could be attainable even with this stress increase by applying proof tests. This design was finalized by transient thermal stress analyses to attain the optimum configuration, because the larger the disk's outer diameter, the higher the hub transient stress level becomes. The analysis conditions involved the engine starting with a step increase of turbine inlet temperature (TIT) from room temperature to 1350"C, and engine stop with a step decrease of TIT from 1350°C to the regenerator outlet temperature of 900°C due to fuel cutoff at the rated running conditions. The changes in temperature distribution in the rotor during the engine start and stop are shown in Figures 10 and 11. The maximum combined thermal and centrifugal stress during start is below that at the rated running conditions, as shown in Figure 12. On the other hand, a thermal stress peak occurs at 10 -20 seconds after fuel cutoff during the engine stop. This peak stress depends on the dimension of the disk diameter. The stress distribution at the peak stress point indicates that maximum stress occurs on the surface of the disk's outer periphery (Figure 13). Reliability analysis performed to check the effect of this thermal stress shows that the failure probability during engine stop does not exceed the value of the rated conditions, as shown in Figure 14. The failure probability at this peak stress point is nearly equal to that at the hub center because the material strength at this point becomes higher than that at the hub center due to the sudden temperature drop around this area because of cooling after fuel cutoff (Figure 15). This rotor volume showing these high failure probabilities is limited to a region close to the material's surface and negligibly affects the
overall failure probability of the rotor. - - - -.. Scallop Dia. : 88mm
- Scallop Dia. : 105mm Blade Tio
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lime (scc.)
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Fig.10 Temperature Distribution Change during Engine Start
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.......... Scallop Dia : 88mm -Scallop Dia. : IOSmm
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Fig.15 Rotor Failure Probability Distribution per Unit Volume (20sec. after Fuel Cutoff, l/mm3)
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Time (sec.) Fig.11 Temperature Distribution Change during Engine Stop
- - Scallop Dia. : 8 8 m m - Scallop Dia. : 105mm
BLADE STRENGTH UNDER RESONANT VIBRATIONS
_.. -.
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Time (SCC.) Fig.12 Rotor Maximum Transient Stress (Centrifugal + Thermal)
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The Campbell Diagram of the rotor indicates the resonant speeds in the engine operating range (Figure 16). The rotor of the initial design broke at a resonant speed with the third vibration mode where the leading and trailing edges have large displacement responses (Figure 17). The conditions that caused the failure were experimentally evaluated using the turbine component test rig because there had been no actual data to predict this. The forces which excite the vibrations are nearly in proportion to the turbine output torque. Therefore the minimum turbine output torque that broke the turbine was measured and it became evident that the measured values were not sufficiently higher than the required torque performance curve for engine operation.
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We tried to make the third resonant point exceed the rated speed by increasing blade stiffness, but at the same time this would also raise the second resonant speed and approach the former third resonant speed. The minimum turbine output torque that broke the turbine was also measured for the second vibration mode and was slightly lower than that of the third mode. Accordingly, a major modification of the blade in the 6th Design was needed to make both the second and third resonant speeds exceed the rated speed. Mode
21Nozzle Va
mmzz 6th Design ............. ....... 1st Desib Nozzle Vanes
Resonance Speeds
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Y
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hour endurance test was completed. A major cause of FOD during the engine tests was delamination of the bonded insulator on the metallic housing. The insulator material was changed to a new one with suitable elasticity and a hard surface Iayer, and a new mechanical fastening device was adopted instead of the original adhesive. Like the component test rig, we also decided to install a trap system in its turbine scroll. In order to design the engine trap system, we conducted particle impact tests on the spinning rotors using the cold spin test rig and measured the minimum particle weight that would produce damage to the rotors. Then we analyzed the flow distribution inside the scroll and the particle movements to determine the most effective trap configuration. Figure 19 shows the scroll flow path design with the trap and the analyzed particle traces. The scroll and trap structure made of silicon nitride is shown in Figure 20. After installing it in the engine no FOD problem occurred.
z" 0
8 1 0 1 2 1 4 A Rated Speed Rotational Speed (xlo'rpm)
2
4
6
c~mbustor
~ r p system p
T e m p 9 6 ~ c I- X
Fig.16 Rotor Campbell Diagram
4 I 9 Rotor Fig.18 Turbine Component Test Rig
Stress Increase Fig.17 Blade Vibration Pattern (3rd mode)
FOREIGN OBJECTS IMPACT DAMAGE POD) Fig.19 Analyzed Particle Traces in Scroll with Trap
Foreign Objects Impact Damage was encountered both in the turbine endurance tests first and then in the engine tests. It was generally difficult to determine cause of failure to be FOD, but occasionally we discovered foreign objects in the engine and rig, and determined these as the objects that impacted the blades. The objects that caused FOD in the turbine component test rig were broken sensors, carbon deposits in the combustor, and oxides of the test rig's metallic materials. Consequently, design improvements were made in the rig's structure, materials, and cooling system. In addition to that, a trap system to prevent foreign objects from flowing into the turbine area was installed between the combustor and the turbine inlet as shown in Figure 18. After these measures had been applied, there was no failure due to FOD and the 100
316
Fig.20 Scroll Structure with Trap
/
ACKNOWLEDGMENTS The author is grateful to PEC for permitting of publication of this paper and the support of MITI is very much appreciated in the works constructing technoinfrastructure for ceramics.
REFERENCES (1) N.Nakazawa, H.Ogita, M.Takahashi, T.Yoshizawa and Y.Mori, Radial Turbine Development for the lOOkW Automotive Ceramic Gas Turbine. Transaction of the ASME, Journal of Engineering for Gas Turbines and Powers, V01.120, pp172-178(1998). (2) A.A.Wereszczak and T.P.Kirkland, Creep Performance of Candidate S i c and Si,N, Materials for Land Based Gas Turbine’ Engine Components. Transaction of the ASME, Journal of Engineering for Gas Turbines and Powers, Vo1.119, pp799-806(1997). (3) N.N.Nemeth, J.M.Mandersheid and J.P.Gyekenyesi, Ceramic Analysis and Reliability Evaluation of Structures (CARES). NASA Technical Paper 2916, (1990). (4) N.Nakazawa, K.Niwa and T.Sugimoto, The Automotive Ceramic Gas Turbine Development Results in Japan. 9th Cimtec-World Ceramic Congress, Ceramics: Getting into the 2000’s-PartD, pp229-240.
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THE EVOLUTION OF DAMAGE IN CERAMIC MATERIALS FOR GAS TURBINE APPLICATIONS UNDER COMPLEX LOAD CONDITIONS E. Nestle, R. Westerheide Fraunhofer Institute for Mechanics of Materials, D-79108 Freiburg, Germany
ABSTRACT By means of oxidation, four point bending and tensile tests performed on different silicon carbide materials the relevant damage mechanism was determined. For that purpose a test rig was developed in which the material characterisation under defined atmospheres and homogeneous normal stresses up to high temperatures can be performed. A model for calculation of the material parameters necessary for lifetime prediction was developed combining Weibull statistics and crack growth theory.
INTRODUCTION To increase the efficiency factors of gas turbines higher gas inlet temperatures are required. As state of the art coated Ni-based materials are in use. A further improvement is expected by the change from metallic to ceramic materials: ceramics keep their strength and corrosion resistance up to high temperatures. For moving components the lower density is an additional advantage. Especially non-oxide ceramics provide a high creep resistance due to their covalent bonding. But on the other hand ceramic materials possess obvious disadvantages that arise from their brittle characteristic: other than in metals stress peaks at crack tips cannot be reduced by local plastic deformations. This leads to undesirable properties as low fracture toughness, low fracture strain and - in comparison with metals - to a high scatter in the component strength. To ensure a safe application material data have to be available during the development process and the component use for being able to dimension ceramic parts corresponding to their load and to identify their crucial damage condition. The determination of material data requires defined boundary conditions: in prior investigations concerning the tensile creep behaviour of ceramic materials it was found that at temperatures above 1 500 "C frequently corrosive attack at the specimen's surface lead to a premature failure [I]. One aim of the presented work was to modify the existing test rig appropriate to the raised demands. To further increase the testing temperatures a redesigned clamping device was put in service and optimised. In addition the test rig was modified by components for testing under defined atmospheres. In this modified test
rig experiments with different Sic-materials were performed to characterise the interaction of the different damage mechanisms. At this the focus lay on the atmospheric influences. In addition to the tensile tests four-point-bending-tests and oxidation-tests were carried out. Based on these experiments a model for lifetime prediction was developed.
EXPERIMENTAL PROCEDURES The highly stressed components are subjected to a combination of thermal, mechanical and corrosive influences: thermal loads result from the direct impact by hot gases. Gas temperatures of about 1 700 "C lead to material surface temperatures of up to 1 300 "C (as state of the art). Static and dynamic mechanical loads are induced by steady state flow- and centrifugal forces and transient states by load changes and pulsations in the gas stream. Corrosive attack arises mainly from combustion products. The attack mode furthermore depends on the fuel-to-air ratio, because thus the character of the atmosphere (reducing or oxidising) is determined. From the combustion a water content of approx. 10 mol-% results. This proportion can be significantly higher: to lower the NO,-emissions by homogenisation of the combustion temperature in some plants additional water is injected. So a typical water content in the gas is about 20 mol-%.
Tensile Creep Tests To reproduce some of the relevant load conditions in the laboratory a tensile testing machine was modified: round specimens with an overall length of 270 mm and a diameter of 12 mm at the ends and 6 mm in the 30 mm long tested area were gripped outside a resistance-heated two-zone-furnace by a tempered gripping device. To separate the testing atmosphere from the atmosphere of the furnace the specimen were situated inside a tube of RSiC. At the ends of the tube the corrosive media (dry or moist air with a water content of 20 mol-%) were led in. The outlets for the gas were the holes for the contacting extensometry and the non-contacting temperature measurement in the middle of the tube. Because of the relatively high flow rates an inward diffusion of gases from the furnace could be excluded. The moisture content of the testing gas was
319
adjusted by bubbling air through a tempered bath of deionized water. All the specimens were tested in the temperature range from 1 450 "C to 1 550 "C in atmospheres containing up to 2Omol-% H20. A constant mechanical load (tension or compression) of 100- 290 MPa was applied. Some of the tests were stopped as a dwell time of 100 hours was reached. Other tests were performed until the specimens failed under the constant load at high temperature. The non-broken specimens from the interrupted 100 h tests were subjected to a residual strength measurement at room temperature.
Oxidation Tests For getting a larger data base concerning the atmospheric influence additional oxidation tests without simultaneous mechanical load were performed. Therefore a furnace was equipped with a RSiC-chamber through which the testing gas was led. Discs made from the end pieces of the tensile specimen were oxidised for 100 h. These tests were interrupted for weighting the specimens after 12,5, 25, 50 and 100 hours. Testing conditions were 1 450,l 500 and 1 550 "C under dry and moist atmospheres. The other type of oxidised specimens were fourpoint-bending-bars. Here tests at 1 500 - 1 550 "C also under dry and moist atmospheres for 100 h respectively 1000 h were carried out. These tests have not been interrupted.
RESULTS Oxidation Behaviour While for both SSiC qualities passive oxidation with parabolic behaviour under all conditions could be found, the hot-pressed Cercom PAD showed active oxidation starting even at 1 450 "C. The thickness of the silica scales lay in the order of 10 pm after 100 h. X-ray diffraction analysis of the oxidised discs showed for all conditions, i.e. dry and moist air in the temperature range from 1450 to 1 550 "C, always Cristobalite. To determine the parabolic rate constants two ways were chosen: on the one hand the scale growth was measured using ScanningElectron-Microscopy, on the other hand the mass change per surface unit was used. The gained values can be converted into one another according to: Am *12,9 = A S A
eq. 1
for Cristobalite where AmlA stands for specific mass gain in mg/cm2 and As for scale growth in pm. . Comparison of both measuring methods delivered a good agreement. The numerical values of k, lie in the range of 1,6 * 10" to 2,4 * 10" mgz/(cm4h)and are in good agreement with [4]. The oxidation behaviour of the tensile specimens was similar to that of the discs and the bending bars.
Four-Point-Bending-Tests Four-point-bending tests were performed with the oxidised and some as-machined bending bars at room temperature in a fully-articulating standard fixture with spans of 40 and 20 mm. At least 10 specimens per condition were tested.
Materials Based on previous investigations [l] in this work comparable materials were examined. In addition to the ESK' SSiC material tested in [l] Hexoloy SA from Carborundum2and PAD S i c from Cercom3were characterised. While the SSiC qualities from Carborundum and ESK have a high purity of Si + C > 993 % (residue mainly A1 and B) the hot pressed S i c from Cercom contains approximately 2 % of additives (residue mainly Al). All materials are of the a-Sic polytype. Further details are documented in [ll, [21 and C31.
Fig. 1: Etched Hexoloy SA after 100 h at 1 550 "Cin dry air
' ElektroschmelzwerkKempten (ESK), Germany Carborundum Corporation, Niagara Falls, USA Cercom Inc., Watson Way, CA, USA
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Fig. 2:Etched Hexoloy SA after 1 OOO h at 1 550 "C in dry air
By etching in hydrofluoric acid the Si02-scale was dissolved for analysing the surface damage of the oxidised specimens. While the surfaces after the interrupted oxidation tests up to 100 h did not show a significant difference concerning the surface modification, after 1000 h oxidation at 1 550 "C a plane, continuous attack could be seen. The pitting visible in Fig. 1 does not result from selective corrosion but from material inherent defects. Fig. 2 shows the same material after 1 000 h oxidation time.
pores found in the uninfluenced material must be the origin of fracture. Subcritical crack growth starting from these material inherent defects was found to be the active damage mechanism.
Flexural Strength Four-point-bending tests have only been performed on the Hexoloy SA material. The comparison of the strength values with and without preceding oxidation shows that for all oxidation conditions and times the strength is increased by 8 - 20 %. Similar results for Hexoloy SA are reported in [5]. As an explanation oxidation induced crack healing of surface damage resulting from machining is mentioned. Long term oxidation does not seem to weaken the material by pitting. Fracture analysis of the oxidised bending bars revealed that no pitting occurs by oxidation under these conditions.
Fig. 3: Etched Hexoloy SA tensile specimen after 830 h at 1 550 "C in moist air under tensile stress of 1 0 0 MPa
Creep Deformation
Modelling and Lifetime Prediction
Creep behaviour of the Hexoloy SA material investigated in this work was found to be similar to the ESK material tested in [l]. While the creep rates of the two materials differ by a factor of approximately two (for example at 1 550 "C, 100 MPa; Hexoloy SA: dddt = 2,Ol * lo-', ESK SSiC: dddt = 4,00 * lo-') the creep exponents and the activation energies are quite similar (about n = 2, Qact. = 500 kJ/(mol K)). This also indicates a similar creep mechanism: creep is supposed to be controlled by grain boundary diffusion processes without changing to self-diffusion at the temperatures investigated. Concerning the Cercom PAD HP-Sic no comparable values could be measured, because even at temperatures of 1450 "C the strong corrosive attack resulted in premature failure by surface damage.
A combined model for lifetime prediction was developed that assumes that crack-like defects - depending on the level of the mechanical load - cause fracture either immediately or after a preceding phase of subcritical crack growth. In this work no other damage mechanism could be verified. Therefore in the model all experiments were fitted at the same time: i.e. flexural and tensile tests, at room temperature and high temperature, dry and moist atmospheres, with and without dwell time. In the presented model the parameters for the Weibull statistics and the parameters for subcritical crack growth are fitted by a single method of approximation at the same time. The advantage oft this procedure is the possibility to get information about the material behaviour by a relatively small amount of specimens. The increase of mechanically loaded cracks is described by the crack growth law:
Damage Mechanism Unlike than reported for testing in undefined furnace atmospheres [l] in this research fracture of Hexoloy SA and ESK SSiC did not occur due to surface damage. Exclusion of the furnace atmosphere provides the opportunity for testing creep and crack growth mechanisms without premature failure by a changed flaw population at the surface due to atmospheric influences. Fracture surfaces revealed that defects inside the bulk were the origins of fracture. Even on micrographs of the etched surfaces it is recognisable that cracks are starting from pores (Fig. 3). Investigations by transmission electron microscopy showed that no additional pore formation by creep deformation occured. So the
da dt
-- - AK,"
eq. 2
The stress intensity factor is given by eq. 3
K, =Y&o
After elimination of Kr in eq. 2 eq. 3 can be integrated. According to the experiments the crack grows in this case during the dwell time t, under the applied testing stress o,,.Afterwards the residual strength DRF is determined. The specimen also may fail during the dwell time. Integration over the dwell time results in 1 (n-2) I 2 a0
1
-
(n-2)I 2
aP
-
n-2 2 AYno;tp
-
eq.4
32 1
.-1 Here a~ denotes the initial crack length, a, the crack length at the end of the dwell time or at the time of failure during the dwell time, respectively. If the crack lengths in eq. 7 are replaced by the corresponding strengths, i.e. the initial crack length by the initial strength
eq. 12
Aeff,B
0.z = g B L \ Aeff.z
)
where ozdenotes the strength of the tensile specimens and oBthat of the bending bars. The effective surface of the tensile specimens AeffZcorresponds to the cylindrical surface of across the tested length (here: 471 mm2). The stressed surface of the bending bars Aeff,Bdepends on the Weibull modulus rn itself:
and the final crack length by the residual strength
eq. 13 eq. 6
the initial strength, that the specimen would have had in a short time tensile test, results from eq. 4.
S2 is the outer span of the fixture, B the width of the specimen and k the ratio of the spans (here 20 mm : 40 mm). If only flaws in the volume are considered, for the equivalent strengths eq.'14
If fracture occurs during the dwell time om = o, has to be set; tp is the time to fracture. Some specimens have been tested under compressive loads. In this case $, = 0 has to be set, as under compressive stresses no crack growth is expected. For the initial strength a Weibull-distribution is assumed:
follows, where Veffz results from the tested length and the diameter of the tensile specimens to 707 mm3. The effective volume of the flexural specimens depends like the effective surface - on m.: Veff,B
eq. 8
=
km+1 2(m + 112
eq. 15
Here k denotes the ratio of the spans and V the volume between the outer rolls (3 x 4 x 40 mm3).
Elimination of (3i by eq. 7 in eq. 8 yields
Fitting Sequence and Calculation
whereas B contains parameters of fracture mechanics: 2 B = K[c2-n eq. 10 A Y ( n - 2) The experimental fracture probability usually is determined by allocation in the order of failure whereby the criterion is the time to fracture or the fracture stress, respectively: pe,,(oi)
=
i-0.5 Z
eq. 11
As tensile tests and flexural tests are to be analysed at the same time, the difference in the effective tested surface, respectively effective volume, has to be regarded. Here the geometry of the tensile specimens is considered as the reference: thus 00 in eq. 8 is the characteristic strength of the tensile specimens. If only damage at the surfaces of the specimens is regarded, the equivalent strength results in
322
Fig. 4 shows the applied fitting procedure. The parameters of the two-parameter Weibull-statistic (m and oo)and the parameters for subcritical crack growth (B and n) are calculated by eq. 9 at the same time. An experimental data set consists of the loading time t,, the testing stress o, and the residual strength ow. Since in regarding the residual strength ow and the time to fracture tp at the same time a definite ranking is impossible, for each data set a value for the initial strength (3i has to be calculated, assuming plausible initial values for oo,m, B and n. The tests are ranked according to their initial strength. Then the experimental fracture probability after Weibull (eq. 11) is associated. By means of the Levenberg-Marquardt-Algorithmeq. 9 is fitted to the experimental probabilities Pexp,whereas oo, m, B and n are the variable parameters. The obtained values are used to recalculate the initial strengths (3i. If flexural test data are included, the dependence of the effective surface, respectively effective volume, on the Weibull modulus m according to eq. 12 and eq. 13, respectively eq. 14 and eq. 15 has to be considered. Normally the new parameters lead to a different order ranked after o i - than assumed before the first run. The tests now are ranked regarding their new values of ( ~ and - after association of the respective probabilities
i
according to eq. 11 - are fitted again to eq. 9 by the Levenberg-Marquardt-algorithm. This step of iteration has to be repeated until the sequence of the tests regarding the initial strengths does not change anymore. The thus stabilised values for the parameters oo,m, B, and n are the wanted values.
'1 :
m ~ l spurnen. e fracture dunng uwptest msle spearnen. fracture at restdual strength te u, bending t 8 t rpearnm u,~,residualswngm % rnedlanica Imd dunng creep test Weitull rarneters Van Ongin Fit
* o
, 92
I
111
130
148
199
250
301
352
Fig. 5: Results under assumption of volume defects
*1
1
1
changed
c
-
Y
-
-1
-2
c -
I
associate P-
*
I
0 Strengtl kaf
A
1 LevenbergMarquardt-
d
+ 92
111
130
148
speamen. fradure a t rerldual rtmgth bsnding tee spearnen 0,. revdual strength ov rneChantca bad during a ~ e test p Welhll pararnetsr fmm Qlgin Fit 0,tmde
0,
199
250
301
352
NPaI Fig. 6: Results under assumption of surface defects 0,
w--.-....U..-
Fig. 4: Fitting algorithm of the combined lifetime model
The advantage of this procedure is the possibility to get four characteristic values for a material at the same time, requiring only a comparatively small amount of experiments. In an exemplary calculation for Hexoloy SA 16 four-point-bendingtests and 20 tensile tests were used. For three of the tensile specimen tested under compression the aforementioned assumption (no crack growth, o, = 0, t, = 0) was applied. The complete data sets are documented in [6]. Fig. 5 shows the stabilised ranking at the end of the described fitting process under the assumption of pure volume defects. The fitted parameters become 00 = 214,s f 1,2 MPa, m = 6,s f 0,4, B = 92 689 f 39 509 MPa2h and n = 12,4 k 1 3 . In the case of surface defects (Fig. 6) the results are oo= 258,s k 0,6 MPa, m = 6,7 f 0,2, B = 4 576 f 666 MPa2h and n = 13,7 k 0,4.
DISCUSSION Limits of the Model In these calculations it is assumed that the fracture mechanism for all the tested specimens is the same, i.e. brittle fracture starting from flaws after preceding slow crack growth dependent on the testing conditions. The assumption would be disproved, if the initial strengths of the high temperature and the room temperature tests arrange in discrete groups. This assumption is supported, if the initial strengths of the different types of testing are mixed up. That the initial strengths of the tensile tests are divided up into two groups was to expect and is no argument against the evaluation procedure: the specimens fracturing during the dwell time have necessarily lower initial strengths than the specimens reaching the given dwell time. These both cases are the limits within which the actual flaw population must lie. The lower divergence from the Weibull-straight in the case of surface defects is in good accordance with the fractographic investigations, which revealed that a big portion of the fracture origins lies near the surface, often with a noticeable oxidation on the crack surface. Although this evaluation method leads to a mixing up of the different types of testing one must be aware of the postulations. The following restrictions have to be taken into account:
323
-
-
-
REFERENCES Besides the surface defects fractures in the volume occur. The parameters of the crack growth A and n may strongly depend on the environmental conditions. However in the present investigations no identifiable influence of moisture was found. A and n are surely dependent on temperature. But the difference in temperature of 50 “C in the high temperature tests doesn’t influence the results in such a way that groups in the initial strength values are formed. If the value of the fracture toughness varies with temperature, a correction in eq. 7 has to be made: with knowledge of the temperaturedependent behaviour of KIcresults: 1
Klc,RT denotes the fracture toughness at room temperature, on which the initial strength q refers, KI~,BT the fracture toughness at that temperature, at which fracture of the specimen occurred. Parameter B (eq. 10) then contains K l c , ~For ~ . the tested sintered a-SSiC no correction is necessary: Investigations on comparable material show only very little shifting of KIc up to 1 600 “C [7,8].
ACKNOWLEDGEMENTS Parts of this work were funded by the “Deutsche Forschungsgemeinschaft” (DFG) in the framework of the “Graduiertenkolleg Technische Keramik”, Karlsruhe.
324
R. Westerheide, T. Hollstein, K. A. Schwetz, Tension-Compression Testing of HIP-Treated Sintered S i c for Gas Turbine Applications at Temperatures Between 1400 “C and 1 600 “C. Proc. 6” Int. Symp. Ceramic Materials and Components for Engines, Arita, Japan (1997) 253 - 258. Carborundum Corporation, S i c Ceramic Materials for Design of High-Performance Applications, product information on Hexoloy, Form A-12,047 ( 1997) Cercom Inc., product information on Cercom PAD Sic, Internet: http://www .cercom.thomasregister.com, 2OOO-06- 15 K. G. Nickel, P. Quirmbach, Gaskorrosion nichtoxidkeramischer Werkstoffe, Technische Keramische Werkstoffe, Ed. J. Kriegesmann, Kijln: Deutscher Wirtschaftsdienst, Chapter 5.4.1.1., 1 - 76 (1991) D. J. Baxter, Oxidation Round Robin (Monolithic Technical Ceramics), Final Report European Commission, EUR 19039 EN (2000) E. Nestle, Charakterisierung des Schadigungsverhaltens von Siliciumcarbiden fur den Einsatz in Gasturbinen unter komplexen Beanspruchungsbedingungen, doctoral thesis at the University of Technology Karlsruhe, Germany (2000) G. Himsolt, T. Fett, K. Keller, D. Munz, Fracture Toughness Measurements on Silicon Carbide, Materialwissenschaft und Werkstofftechnik, 20 148-153 (1989) G. H. Campbell, B. J. Dalgleish, A. G. Evans, Brittle-to-Ductile Transition in Silicon Carbide, Journal oft the American Ceramic Society, 72 [8] 1402-408 (1989)
MECHANICAL FAILURE OF ELECTROCERAMIC COMPONENTS P. Supancic, M. Lengauer, R. Danzer Materials Center Leoben and Institut fur Struktur- und Funktionskeramik Montanuniversitat Leoben, Peter-Tunner-Str.5, A-8700 Leoben, Austria
ABSTRACT Electroceramic current driven devices, such as varistors and power switching PTC- and NTC-components, are primarily loaded electrically. As a consequence of the intrinsic Joule selfheating process, heat expansion of the material occurs. The corresponding thermomechanical stresses can be so high that mechanical failure occurs. The stress amplitudes are governed by a large number of geometrical, electrical and thermoelastic properties. Two examples of electroceramic components are discussed in detail: a power varistor (where temperature changes arise so fast that inertial effects play a dominant role for the mechanical loading) and a high-power PTC-switching device (where a highly nonlinear electrical resistance characteristic causes strong thermal gradients).
INTRODUCTION Although electroceramic devices are primarily electrically loaded, high power devices can suffer mechanical failure [l-31. To a large extent, this is a consequence of Joule heating, which causes temperature differences and locally different thermal strains and stresses in the components. Due to the general demand of industry for increasing power densities, the mechanical loads of these components increase and problems with insufficient mechanical strength get more and more severe. Therefore, and despite the large costs, proof testing of highly loaded electroceramic components is a commonly used practice [22]. In the light of these facts it is surprising that relatively little literature exists about the mechanical failure of electroceramics. In principle the necessary procedure for a proper mechanical design of electroceramic components is quite similar to that for structural ceramic components: In a first step, the loading analysis, the temperature and mechanical stress fields in the component has to be determined. Input parameters are the external loading conditions of the component, its geometry and some material properties. In a second step, the reliability analysis or life time prediction, the loading of the component is compared with the relevant strength of the material.
Loading analysis In structural ceramic components the external mechanical loading is in general determined by the action of external forces or by the action of temperature differences or changes [4-61. The determination of the resulting strains and stresses is a standard problem. Numerical solutions can be obtained by using commercial Finite Element codes (e.g.: Abaqus, Ansys, Nastran, etc.). In electroceramics additional strains and stresses may occur due to a coupling of electrical with thermal andor mechanical properties. This fact is well known for piezoelectric ceramics, were an applied electrical field induces mechanical strains [7]. Other examples are thermistors (temperature dependent resistors), where the electrical current causes a Joule selfheating process, which may cause thermomechanical stresses [8]. In thermistors the resistivity (and therefore also the strength of the heat source) strongly depends on the temperature. In these materials thermal gradients (which occur in any component) cause a strongly inhomogeneous distribution of heat sources which again influence the temperature and the electric field. This highly nonlinear problem of mutual interactions has to be solved in order to determine the temperature field and the thermomechanical stress and strain field. Over a long time scale, the load spectrum might change due to degrading electrical properties [9,10]. These not fully understood kind of aging processes lead to a change of the heat source distribution and consequently also influence the thermo-mechanical loads. These additional changes in strains and stresses must also be accounted for, but up to now, only relatively little theoretical work on these problems were performed. In the present work a short summary on state of the art of the mechanical loading and failure of high power varistors and PTC-devices will be given. Reliability analysis In ceramics failure generally originates at flaws, which are sparsely distributed inhomogeneities in the material and which can, in general, be described like cracks [I 1, 121. Fracture starts if - at the location of the "crack" - the stress intensity factor exceeds the fracture toughness (this criterion is called the Griffith criterion). The failure stress (the strength) depends on the size and orientation of the flaws and on the fracture toughness of the material [13]. For a given material the strength is high if the most critical flaw is short and vice Vera. In general the size of flaws is not exactly known and differs from component to component. Therefore the strength
325
of ceramic materials scatters and has to be described by statistical methods [14, 15, 111. A widely used strength distribution function is the well-known so-called Weibull function [16, 171. On the basis of these ideas a probability of failure can be determined for any given stress field (the result of the loading analysis). Mechanical design in the field of structural ceramics is based on the ideas described above. The developed methodology should be applicable to electroceramics if the failure (the fracture toughness, the strength) is not influenced by an applied electric field (the development of mechanical stresses can be influenced). Up to now such an influence has not been reported for ceramic resistors but recently such an influence has been observed on piezoelectric materials [18-21]. Therefore the loading analysis can be applied for resistors but has to be modified for piezoelectric ceramics. In many ceramics cracks slowly grow under the influence of a subcritical load [14, 221. This results in a decrease of strength with time under load. If the strength approaches the loading stresses an increasing fraction of the components will fail. The life time is exhausted if this fraction reaches a predetermined limit. Although these ideas are well known for structural ceramics an application of a life time prediction on electroceramics has as yet not been reported in the literature.
FAILURE OF VARISTOR COMPONENTS
absorbs energy and heats up rapidly. After establishing the normal voltage, also the varistor returns to the highly resistive state. The reason for their unusual electrical behaviour can be located at the grain boundaries: they act as electrostatic potential barriers. A certain limit of electrical field strength has to be exceeded so that electrons can overcome these barriers. An excellent review on many different aspects of varistor ceramics can be found in [23,24]. The ceramic bodies of power varistor devices are typically shaped as cylinders. Both flat surface areas of the disc are coated with a metallic electrode layer to provide a conducting surface to the circuit. The width of an uncovered gap at the rim of the disc is in the range of about 1 mm. The peripheral surface is covered with an insulating glassy coating to protect the ceramic against chemical attacks and to avoid flash-overs.
Fig. 2a: Puncture failure of a power varistor (indicated by an arrow). Diameter of the disc: 70 mm.
Varistor ceramics, based on doped ZnO, show highly non-linear electrical fieldresistivity characteristics (voltage dependent resistor). A typical relation is plotted in Fig. 1 on a double logarithmic scale.
Fig. 2b: Separation in two parts, the fracture origins are located at the peripheral surface (indicated by arrows). Diameter of the disc: 34 mm.
Current density [A/cm2]
Fig. 1: Typical electric field/current characteristic curve of a high power varistor.
Varistor components are nearly electrical insulators at low electric fields and good conductors above a threshold field strength (i.e. a certain switching voltage). Typically varistors are used in parallel with circuits to protect them from voltage surges. In normal use, voltages below the switching voltage are applied to the varistors and they act as insulators. When the voltage exceeds the threshold (for instance during a voltage transient), the varistor becomes highly conducting and draws the current through it. By this process, the varistor
326
Fig. 2c: Separation in two parts, the fracture origin is located near the centre of the disc (indicated by an arrow). Diameter of the disc: 34 mm. Three modes of failure have been observed, see Fig. 2a2c. The first mode, the puncture failure, is commonly associated with thermal run away, the two other modes are consequences of thermal strains, which occur due to selfheating in service.
Puncture failure The varistor component is designed to absorb the Joule heat homogeneously during a voltage surge. In some cases (predominantly at longer voltage pulses) some current paths draw relatively higher currents and result in melting of the ceramic. A typical fusion channel of molten varistor ceramic is shown in Fig. 2a, which connects both electrodes. This phenomenon of current concentration is accounted for by inhomogeneities in the microstructure, essentially by large grains [23]. Since the material is destroyed after a breakdown of this kind, the failure causing defects cannot be found afterwards. The dimensions of high power varistors are in the range of several cm and there is no simple, non-destructive way of detecting potential microstructural inhomogeneities. For an experimental investigation, special thin plates of about 0.5 mm thickness, were prepared. By measuring the resistance of the plates at low field strength and successive splitting of the plates, the low resistive parts could be isolated. Finally the large grains, causing the resistance deficiency, were detected by serial sectioning. In Fig. 3 two large grains bridging the whole thickness of the sample could be made visible by etching techniques. It is a reasonable assumption, that large grains also exist in real high power components and cause current localisation at high voltages.
Fig. 3: Two large grains in the varistor microstructure, indicated by white arrows (approx. length: 300 - 400 pm).
Brittle failure from the surface As mentioned earlier, there is a thin glassy coating (width: 50 - 100 pm) at the peripheral surface of the disc. Since this layer is electrically insulating, the glass is not actively heated by a current pulse. The time span of these pulses are typically in the range of one or a few tens of ps. Since there is no significant heat conduction from the heated ceramic body to the glass layer within this short time interval, strong thermal mismatches are induced between the glass and the ceramic. A second geometrical feature of the varistor component can increase this thermoelastic mismatch: the uncovered gap between the circular electrode layer and the rim of the disc leads to a reduced current intake at the ceramicregions near to the whole peripheral surface. As a consequence less electrical energy is absorbed in these regions. The worst case concerning thermal stresses is
found immediately after a current surge: The glass layer has still nearly the ambient temperature, the varistor ceramics next to the glassy surface is less heated compared to the massive bulk of the varistor component. The cold glass and the cooler ceramic material at the peripheral surface are put into tension by the expanding core material of the varistor. For a typical mean temperature increase of 120 "C the corresponding thermal stress maxima are found in the range of the materials strength of about 100 MPa in the varistor ceramics and about 50 MPa in the glass layer. After reducing the temperature differences by heat conduction the amplitude of tensile stresses decrease. In Fig. 2b the fracture surface of a varistor is shown, which is split into two parts. The fracture origins are found in the highest loaded area, namely close to the outer surface of the cylinder. Brittle failure from inside The loading case described above leads to a tensile stress free field in the whole core of the varistor. From this point of view there is no reason for starting failure from the inner regions. But in experiments, especially in short pulse tests of duration spans of some ps, also these cases are found. An example is shown in Fig. 2c, where the cracking starts obviously in the centre of the cylinder. A pore was found to be the fracture origin on the one hand, and there was no evidence for any kind of puncture failure on the other hand. The mechanism of this fracture mode can be understood as the result of inertia forces in response to the rapid Joule heating. The conceptual basis for this effect can be found in [2], where a one-dimensional model is presented. By assuming a homogeneous heating process during a power pulse in the order of l o p s , compressive stresses are built up in the centre, since the top planes cannot follow the heat expansion that fast. This inertia effect can be made evident by evaluating the time span, which is used by an elastic wave to move from one top plane to the other: since the velocity of sound c, of varistor material is 4300m/s, a components height h,, of 44 mm is propagated within a time interval of z,= 10 ps. After this initialisation phase the compressive stress state changes to tensile stresses. Elastic waves are running
80 60
a 40 fn
3
20
fn a
0
c
.-n 0
c -20
'r
n s7 -40
-60 0
20
40
60
80
100
Time [ps]
Fig. 4: Stress oscillation in the centre of the disc, calculated by an one-dimensional [2] (dotted line) and three-dimensional model [25] (solid line).
327
through the whole varistor and lead to oscillations. The largest stress amplitudes are found in the centre of the cylinder. This stress oscillation for a typical test pulse, which leads to a temperature increase of 120"C, is shown in Fig. 4 as the dashed line. The solid line in Fig. 4 corresponds to the results of a three-dimensional model for the same heating process, published recently [25]. In this case the reflections of the initial compressive wave at different surfaces of the cylinder lead to a complicated interference pattern in the centre. The tensile stress maximum at about 80 MPa was found to be higher in this case compared to the one-dimensional model due to lateral contraction effects, which are triggered by Poisson's ratio. Analysing the relation between pulse duration and tensile stress maximum, by using the simpler one-dimensional approach, leads to the following essential results: for test pulses shorter than the time interval needed for an elastic wave to propagate through the length in the order of the mean varistor dimension (in most cases: height h,,), the stress is proportional to the product of the heat expansion coefficient a: (i.e.: 7.10-6OC-'), the Young's modulus Y (i.e.: 100 GPa) and the temperature change AT during the pulse. For longer pulses the influence of inertia forces decreases dramatically. In Fig. 5 the stress amplitude d (nonnalised to the maximum value @Y.AT) vs. the pulse duration is plotted for a typical varistor height of 44 mm. This relation explains the experimentally observed fact that short pulses in the order of z, = 10 ys and shorter can lead to failure initiated in the centre of the cylinder. For longer pulses in the order of multiples of z, this mechanism becomes unimportant concerning failure. Until now inertia forces are the only explanation for a central initiation of fracture of varistor components.
FAILURE OF PTC-DEVICES Positive temperature coefficient (PTC) resistors are characterised by an increase in the electrical resistance with temperature. Apart from special polymer composites, commercial PTC-components are based on a doped, barium titanate ceramic. A ferroelectric-paraelectric phase transition at a critical temperature (the Curie-temperature) in combination with screening effects of potential barriers at grain boundaries causes the exceptionally strong PTC-effect - encompassing a resistance increase of several orders of magnitude within a temperature interval of a few tens of degrees. A typical curve of the resistance/temperature characteristics is shown in Fig. 6. The actual PTC-range spans a temperature interval from 60 to 230 "C, but the strong onset
Temperature ["C]
Fig. 6: Resistance/temperature characteristics (normalised to p-) at various electrical field strength.
1
10
100
Pulse duration [ps] Fig. 5 : Calculated relation between the maximum normalised tensile stress and the pulse duration (one-dim. model).
328
of the resistance rise is found at about 120 "C, which is called the reference temperature. While the solid line corresponds to the resistance behaviour at nominally zero electrical load, the dashed line represents the resistance under an electrical field strength of 1 kV/cm. Generally the resistance drops by increasing the electrical field. This additional resistance lowering effect is called the varistor-effect of PTCs. Power PTC-resistors are used as self-regulating switching devices for degaussing, heating and overload-protection applications. However, mechanical failure can occur during high-power switching processes, which limits the magnitude of the applied electrical voltage. One main mode of failure, concerning the ceramic, is delamination fracture, which splits the whole component (which is usually formed as a cylindrical disc) into two parts. A typical example of a delaminated component is shown in Fig. 7.
. stat.
,;/. a .,j
Fig. 7: Delaminated PTC-component (split into two parts). In the following, the results of a theoretical analysis will be presented [3], which explain the electrically induced thermal stresses during the PTC switching process, that can lead to the catastrophical delamination failure. A mathematical model was developed, including all relevant electrical, thermal and mechanical aspects of the PTC. To demonstrate the main mechanisms, a simple trial PTC-component was chosen, which has a cylindrical geometry of 3 mm diameter and a height of 2.5 mm. A copper wire is attached to the electrodes by solder. The switching behaviour of the PTC-component in a circuit, consisting of a serial connection of the PTC and a fixed shunt-resistor, was measured by monitoring current and the voltage drop at the PTC. Taking all experimental boundary conditions into account, the mathematical model reproduces well the observed current and resistance behaviour in time. The comparison between experiment and theory is shown in Fig. 8. After a certain time span an initially strong electrical current is dramatically reduced, since the PTC has switched to the high resistive state due to selfheating. But the temperature rise within the PTC-component is not homogeneous. Although the heat production rate is nearly constant in the disc at the beginning of the switching process, heat conducting leads to a solder-induced cooling of the disc at both top planes. The central regions of the PTCdisc reach the reference temperature faster, so that the resistance increase starts locally in this range. A higher resistance leads to a concentration of the heat generation and consequently a strong temperature increase in the centre of the PTC-disc. This positive feedback causes in this case a transient temperature difference of more than 100 "C. The build-up of the axially varying temperature
0,30 0,25
- 0,zo s
10'
E
105
4j
E
L
C
E
0.15
c
L
r;
0,lO
n
0,051
Time [s] Fig. 8: Current and resistance vs. time in a F'TC switching process in air at 20 "C, the shunt was fixed to 1 kR. Applied total voltage: 300 V DC.
___.-.-._. . , ,.,-,-
all
t=2 s
'
-
t=0.25 S
-
t=o s
20
0 -1,0
-0,5
0.0
0,5
1,0
Axial position [mm]
Fig. 9: Transient axial temperature profiles at several points in time. field is shown in Fig. 9. Finally the electric power is reduced so much, that it is balanced by the heat loss through the surface and a stationary temperature distribution is achieved. Since the mechanical stresses have to be known for the loading analysis, the corresponding thermal stress field was calculated for that point in time when the largest temperature differences occur. In Fig.
at k0.75 s
0,lO
--
/!!SqEs+q compression
1 I
-0,lO
I
-0,15 -0.10 -0,05 0.00
.=loo
I
0,05
0,lO
0,15
Diameter [cm]
Fig. 10: Isostress lines of the 1'' principal stresses on a cross-sectional area of the PTC.
10 isostress lines on a cross-sectional area of the first principal stress field are shown. While the centre of the ceramic disc is under a compressive stress, the outer regions are in tension. In the considered case a significant stress maximum of about 100 MPa occurs near the peripheral surface in the plane of symmetry between the electrode layers. This significant stress concentration is caused by the constraining effect of the relatively cold disc-faces, which clamp the thermal expansion of the warmer central planar regions. The amplitude of tensile stress is in the range of the material's strength and it should be expected that delamination fracture is originated by flaws in this ring-shaped region. By investigating fracture surfaces using scanning electron micros-
329
copy, it was possible to iden@ these fracture causing defects. In Fig. 11 a typical fracture origin is shown (pore with a diameter of about 80 pm).
Fig. 11: A pore as fracture origin.
CONCLUSIONS Mechanical failure of electroceramics due to electrical loading is of great economic significance and has to be avoided by a proper mechanical andlor microstructural design or proof testing. In general mechanical loading of active electroceramics results from thermal strains, which are caused by Joule heating. In most cases the loading is “quasi-static”, but sometimes it is highly transient so that inertia forces have to be considered. There is a similarity to the situation under thermal shock or thermal cycling. The concepts used to optimise geometric and microstructural design in thermal shock-situations can be transferred to designing electroceramic components. In all investigated cases failure was caused by “large” flaws. Reducing their size will directly lead to a higher performance. In order to improve the components reliability, there is a demand for tools to analyse mechanical stresses in electrical loaded components. Details of understanding of the interaction between electrical load, heat production and transport and thermomechanical properties have to be taken into account.
ACKNOWLEDGEMENTS The authors thank EPCOS AGDeutschlandsberg (Austria) for supplying samples, and acknowledge the collaboration of I. Hahn (SIEMENShlunich), J. Riedler and G. Schoner (EPCOS) for providing some of the requested material data and the helpful discussions on related topics. This work was supported by the Austrian Kplus-program.
330
REFERENCES C. Dewitte, R. Elst and F. Delannay, On the Mechanism of Delamination Fracture of BaTi03based PTC Thermistors. J. Eur. Ceram. Sw., 14, (1994) 48 1- 492. A. Voita and D.R. Clarke, Electrical-ImpulseInduced Fracture of Zinc Oxide Varistor Ceramics. J. Am. Ceram. SOC.,80, (1997) 2086 - 2092. P. Supancic, Mechanical Stability of BaTi03-based PTC Thermistor Components - Experimental Investigation an Theoretical Modelling. J. Eur. Ceram. SOC.,(2000) in print. H. S. Carslaw and J. C. Jaeger, Conduction of Heat in Solids, Second Edition. Clarendon Press, Oxford, (1959). B. A. Boley and J. H. Weiner, Theory of Thermal Stresses. Krieger Publishing Company, Malabar, Florida, (1960). Mechanics and Mathematical Methods, Second Series: Thermal Stresses. ed. by R. B. Hetnarski, North-Holland, (1986). Lmdolt Bornstein, Numerical Data and Functional Relationships in Science and Technology, Group 111,Vol. 3. ed. by K.-H. Hellwege and A. M. Hellwege, Springer Verlag, (1969). D. S. Smith, N. Ghayoub, I. Charissou, 0. Bellon, P. AbClard and A. H.Edwards, Transient Thermal Gradients in Barium Titanate Positive Temperature Coefficient (PTC) Thermistors. J. Am. Ceram. SOC.,81, (1998) 1789 - 1796. R. C. Bradt and G. S. Ansell, Aging in Tetragonal Ferroelectric Barium Titanate. J. Am. Ceram. SOC., 52, (1969) 192 - 199. K. Wu and W. A. Schulze, Effect of the AC Field Level on the Aging of the Dielectric Response in Polycrystalline BaTi03. J. Am. Ceram. SOC.,75, (1992) 3385 - 3389. R. Danzer, A General Strength Distribution Function for Brittle Materials. J. Eur. Cer. SOC.,10, (1992) 461 - 472. D. Munz and T. Fett, Ceramics - Mechanical Properties, Failure Behaviour, Materials Selection. Springer Verlag, (1999). A. de S. Javatilaka and K. Trustrum, Statistical Approach to brittle Fracture. J. Mat. Sci., 12 (7), (1977) 1426 - 1430. (14) J. B. Wachtmann, Mechanical Properties of Ceramics. John Wiley & Sons, New York, (1996). (15) R.Danzer, Ceramics: Mechanical Performance and Lifetime Prediction. Concise Encyclopaedia of Advanced Ceramic Materials, ed. by R. J. Brook, Pergamon Press, Oxford (1994) 285 - 298. (16) W. Weibull, A Statistical Theory of the Strength of Materials. Proc. Ing. Vatenskaps Akad. Handlingar, Vol. 151, Royal Swedish Institute for Engineering Research, Stockholm, (1939) 1- 45. (17) W. Weibull, A Statistical Distribution Function of Wide Applicability. ASME J. Appl. Mech., 18, (1951) 293 - 297.
(18) T. Fett, D. Munz and G. Thun, Tensile and bending strength of piezoelecric ceramics. J. Mat. Sci. Letters, 18, (1999) 1899-1902. (19) G. A. Schneider and A. Kolleck, Spannungsinduziertes ferroelastisches Domhenklappen als Verst%rkungsmechanismus in Bariumtitanat. Lmgzeitverhalten von Funktionskeramik, WerkstoffInformationsgesellschaft mbH, (1997) 55 - 63. (20) T. C. Wang and X. L. Han, Fracture mechanics of piezoelectric materials. Internat. J. of Fracture, 98, (1999) 15 - 35. (21) F. Meschke, A. Kolleck and G. A. Schneider, RCurve Behaviour of BaTi03 due to Stress-Induced Ferroelastic Domain Switching. J. Eur. Ceram. SOC.,17, (1997) 1143 - 1149. (22) R. Danzer, Subcritical Crack Growth in Ceramics. Encyclopaedia of Advanced Materials, Vol. 4, ed. by R. J. Brook, Pergamon Press, Oxford, (1994) 2693 - 2698. (23) D. R. Clarke, Varistor Ceramics. J. Am. Ceram. SOC.,82, (1999) 485 - 501. (24) Yet-Ming Chiang, D. Birnie I11 and W. D. Kingery, Physical Ceramics. John Wiley & Sons, New York, (1999). (25) M. Lengauer, D. Rubesa and R. Danzer, Finite element modelling of the electrical impulse induced fracture of a high voltage varistor. J. Eur. Ceram. SOC.,20, (2000) 1017 - 1021.
33 1
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THE ENERGY AND THE POWER TIME DEPENDENCE ON THE ULTRASONIC WELDING PROCESS A WEIBULL STATISTICAL BASED MODEL I. A. Bassani", H. Kockelmann"", K. Kussmaul** (*) Laboratbrio Instrumentagio, Materiais e Ensaios (LIME), Regional Integrated University P.O. Box 44,98800-000 Santo h g e l o Rs,Brazil (**) Staatliche Materialpruefungsanstalt (MPA), University of Stuttgart Pfaffenwaldring 32,70569 Stuttgart, Germany
ABSTRACT Considering the energy balance of the fiiction conditions of the ultrasonic welding process and based on measurements of the growth rate of the bonded area. a Weibull statistical four parameter based model is developed in order to describe the energy and the power time dependence. The comparison between the developed model and the experimental results is made by welding compounds of partially stabilized Zirconia zIo2 (3% M,@) and the heat resistant steel X5 CrNi 18 10. using an interlayer of aluminium alloy AlMgSi 0.5.
INTRODUCTION In ultrasonic welding the parts to be welded are clamped together between a sonotrode and an anvil. By welding metals, glasses and ceramics the sonotrode vibrates parallel to the interface of the two parts and after a certain time a joint is formed. The variables involved are the fiequency of vibration. the vibration amplitude. the welding time. the clamping force. as well as the material properties and the roughness and shape of the parts and of the sonotrode 111. By welding the deformation. the welding time or the energy can be controlled. The formation of a bonded joint between materials surfaces has been attributed to (a) the localised melting or heating arising fiom fiiction. elastic hysteresis and plastic deformation. (b) mechanical interlocking. (c) interfacial nascent bonding and (d) a chemical bond involving diffusion [2]. The welding process is always accompanied by a temperature time variation. When joining ceramics and metals internal stress caused by thermal expansion mismatch must be compensated. otherwise fatal damage due to thermal shock or thermal cycling can occur in the ceramic. At least a residual stress pattern is induced which can cause plastic deformation and cracking. affecting the mechanical bond quality [3-51. Together with phase transformation the result is a final residual stress state that can cause the damage of the joint. Transmission electron microscopy analysis of the bonded areas of welded parts revealed the formation of a fine ,pined structure. resulting fiom both recrystallization and/or transformation of heavily deformed material and short time melting [a, 7.
For ceramic-metal joints this bonded area has been assessed to be of the order of a few nanometers [S]. According to the Hall-Perch relation. the mechanical strength can be expected to be about 2.5 times higher than the strength of the base material. without an embrittling effect [9]. The fact that the bonded area depends on welding time and all other parameters [l, 8,10,11]. leads to this work. If it is possible to correlate the growth of the joined surface with the energy involved for the generation of the welded joint. then there is a way to simulate with precision the resulting temperature time variation and hence the residual stress state ofthe surrounding bonded volume. i he main purpose of this work is (1) to present a model that describes the energy- and the power-time dependence based on the ultrasonic welding process energy due to the fiction conditions. (2) to establish a correlation between the growth rate of the welded area and the energy- and power-time dependence. and (3) to compare and validate the developed model with experimental results.
THEORETICAL PART Energy and Power Consumption The enera E involved in the bonded surface formation of hvo materials is dependent on an outer fiiction mechanism (sliding between pamers) and on an inner fiiction mechanism (elasto-plastic and viscoelastic deformation of the material) [S]. So E can be thought of as the s u m of an outer E.F and an inner energy part EI. Considering two solid surfaces A,, mo&g with respect to another one, both chemically compatible. for a welding time ,t and constant welding parameters. we assume that the outer energy E.F is proportional to the joined surface growth A(t), and the inner energy EiF is proportional to the integral of A(t) with 0 I A(t) I A,,. Relating E to A,, with the constants kland k2 leads to:
The factor k, represents the input energy per unit surface due to outer fiiction [J/m'] and k2 the input power per unit surface due to inner fiction[W/ m']. For t << ,t (A,, will be joined) the value of E(t)/A,, approaches zero. For t >> t,, (A,, is joined) the function E(t)/& goes toward a,
333
with a slope equal to kr. and QF/&assumes the constant value kl. The outer power time dependence P.F (t) is hence proportional to the joined surface growth rate dA/dt. The inner power time dependence Pp is proportional to the joined surface growth A(t). The total power P(t) related to & is:
For t << ,t just the outer power PeFinfluences the total power P(t) and for t >> t,, it does not influence it. In this case Pp/& and P(t)/& assume a constant value equal to k2.For 0 It I t, both E(t)/& and P(t)/& are dependent on the joined surface growth A(t). The concept of time constant [12]. leads to the weld system time constant T~ given by: -kl - &OF L t W (3)
dP*
k2
Then the bonded surface growth A(t) related to & can be represented by the distribution function: -=l-e A(t) -(at)" t>O (4) A0 where a is the inverse of a time constant and n the WeibullParameter. This function outlines the probability that the surface A(t) related to A,, is joined at the time t and is represented on Fig. l(a) as a function of the non-dimensionalparameter at. For a given Weibull-Parameter n the function given by Eq. (4) approachs 0.632 as t tends to l/a. i.e. l/a is the time to join 63.2 % of the surface &. The Weibull-Parameter n can be delineated fiom the slope of the linear function: (1- A(t)/ A0
=n
h (at).
(5)
The derivative of the distribution function Eq. (6) is the probability density function. i.e. the joined surface growth rate W d t :
Bonded Surface Growth and Bonded Surface Growth Rate It is presupposed that the probability of surface aspirities touching one another is an absolute continuous random variable. that obeys a Weibull distribution [13]. and that the joined surface A is directly proportional to the weld time &..
1.5
. -
a-
1
-1
t 1.0-
7
~
The Weibull-Parameter n outlines how swiftly the joined surface A(t) grows. Fig. l(b). Regard must be taken of this point while optimising welding processes: a high value of n means a low welding time. The value n = 1 gives the particular situation for an exponential distribution.
Description of the Energy- and Power-Time Dependence The energy time dependence E(t) /& is obtained fiom Eqs. (1) and (4). Eq. (7) is represented in a nondimensional form as a function of the non-dimensional and relative time at on Fig. 2(a).
0.63
-
0.0
1.o
2.0
3.0
at
When t >> l/a the outer energy E o reaches ~ a constant value kl& i.e. the outer energy appears just in the first weld phase. Fig. 2(b). The higher the Weibull-Parameter n. the sooner a constant value of kl is reached. For all values of n the outer energy reaches 6 3 2 % at the time t=l/a. During the beL@nhg of the welding process, the inner energy value E i is ~ small for high values of n. For t >> l/a it is linear uith time, and for infinite welding time it tends to infinity. Fig. 2(e). The pouer time dependence P(t)/& is obtained fiom Eqs. (6). (4) and (2). Fig. 3(a): p(t)
--40
at
(b) Figure 1 - Bonded surface growh (a). (Eq. (4)) and (b) bonded surface growth rate (Eq. (6))
334
[
- k, a n (at)"-'
,-(at'"
] [
+k2 1-e
. (8)
From the physical inteqretation of the welding process phenomenon involved. the power time dependence P(t) is just possible when n >1. This yields for the energy time dependence E(t) also. In this case for t = 0 the slope of the curve must be smaller than 45O. A constant slope of 4 5 O for n=l means a constant power value of 1. because the power is the derivative of energy and tan 4 5 O = 1. With the limiting value n > 1 the surface & is always welded after at = 3 for all values of the Weibull-Parameter n. i.e. after 3 times
the time constant. Fig. 1 (a).
1.5
-. w’ 4
1.0
0.5
0.0
1.o
2.0
3.0
1.o
2.0
3.0
at
1.5A
-
.
$
=-
1.0-
.
0.0
-. -
1.0-
-
.
0.0
0.0
1.o
2.0
3.0
at
(c) Figure 2 - Energy-time dependence: (a) total E(t). (b) outer EeFand (c) inner energy E ~ F The outer power PeF is represented in Fig. 3(b). High values of n correspond to high peaks of power. Independent of n, 63.2 % of the power consumption occurs for t Il/a. neinner power ps has a consmt value ofk, A. for t >> 0. Fig. 3(c). For high values of n the value ~i, is sooner than for a lower one. F~~ I 1iathe topower input p(t) amounts to 63.2 %. The value k, is calculated with the parameters a. n and kzkom Eq. (8). Quod erat demonstrandum.
1.o
2.0
3.0
at
(c) Figure 3 - Power-time dependence: (a, total P(t), (b) outer peF.and (c) inner power pW
EXPERIMENTS AND RESULTS Experimental h ~ c e d u r e Welded joints were produced on a ceramic cylinder of partially stabilized Zirconia zrO2/3%MgO (410x50 mm) and a metal sheet of austenitic steel X5 CrNi 18 10 (1%12~0.5 mm?. using an interlayer of aluminium alloy A1Mn0-5Mg0*5 (16s1 6xo.1 =’). The ceramic surface was prepared with coarse silicon
335
carbide grinding paper numbered 500 and 120. The ceramic sticking out 5 mm was firmly fixed on an anvil and its grooves were rotated 90" relative to the sonotrode movement. Aluminium and metal steel were placed over the ceramic and all parts were clamped with 800 N by the sonotrode. The sonorrode had square pyramidal grooves with 0.5 mm base length. The fiequency used was 20 kHz with an amplitude of 20 pm. These parameters were controlled and the power-time acquisition was done with a computational system using a commercial ultrasonic welding machine. All welding parameters cited above were optimised previously to the experiments [8].
rate is higher than that for a high one, i.e. a small roughness implies a better energy coupling. The graphical determination of the constants a and n in Fig.4(a) can be seen in Figs. Ya) and (b).
Bonded Surface Growth and Bonded Surface Growth Rate With the variation of the welding energy the welding time t, was varied too, and the welded surface A(t,,) was measured after the tensile test. The description of the bonded surface growth and of the bonded surface growth rate were compared with measured results. Fig. 4(a) and (b).
l/a = 0.077 0.0
0.2
0.1
0.3
Time [ s ] (a) 3 2-
1 0 0 f i I 1-
t
0
-
-1 -
-2
I
n =tan a = 2,6
~
-3.
-2
-
0.2
0.3
25
Time [ s] (a)
-l 0 O 0 t
2
(b)
I
0.1
1
In ( at )
---PaperNr. 120 a = 10 n = 1.3 -Paper Nr. 500 a = 13 n = 2.6
A
0
-1
-
20-
..
total
A Time [ s i ...-.. ..
0
0.1
0.2 Time [ s] (b)
___---_
...
0.3
Figure 4 - (a) Measured bonded surface growth and (b) bonded surface growth rate The bonded surface growth follows an exponential function and is dependent on the ceramic surface preparation. For a small roughness the bonded surface growth
336
Figure 5 - Determination of (a) the time constant l/a. (b) the parameter n and (c) the constant kz and the time constant of the weld system TW
Description of the Energy- and Power-Time Dependence The factor k2 is delineated fiom the measured power P(t). Fig 4(c) (ginding paper 500). For this purpose the case with the smallest possible upper part and the higher relative motion between ceramic and metal was used, i.e. a
condition with small energy loss and high fiction between the welding parts.
5
20 2
constant T, is then determined. Eq (3).The true kl value is calculated from k, = k2T~ or fiom Eq.(11). The total. h e r and outer power were calculated according to Eq. (8). and are summarised in Fig. 6(a) (another case with paper 500). The same procedure was carried out for the determination of the energy-time dependence (Eq. (7)). see Fig. 6(b).
CONCLUSION AND REMARKS
-
1 j
power input in theweld2edzone
k. =0.9 Jlm = 5.0 Wlm2
-
a = 13.0 .. n = 2.6
-
Based on an energy balance a Weibull statistical four parameter based model was developed in order to describe the energy-/power-time dependence for the ultrasonic welding process. A correlation between the growth rate of the welded area and the energy-/power-time dependence was established. Eqmimental measures of the bonded surface growth confirm the validity of the model. The model can be used for the description of other welding processes based on fiction. like fiiction welding for example, but it has to be experimentally proofed. The model is the starting point for a finite element simulation of the welding process. Residual stress analysis. as well as temperature measurements should be carried out to confirm these results.
REFERENCES 0.3
0.4
0.5
Time [ s ]
(a)
I
7
Time [ s ]
(b) Figure 6 - (a) total power input and calculated power input in the bonded zone.(b) the same for the energy With the known constants a. n and k2the factor k, is approximated and the curve in Fig 4(c) is obtained. The
(1)Bassani. I. A.. Optimierung der Ultraschallfuegetechnologie fuer Aluminium-Keramik-Verbind-ungen. MPA-Studienarbeit. Univ. Stuttgart.Germany. 1993. (2) Drews. P.. Beitrag zum Ultraschallpunhschweissen von Metallen. Dissertation. T. H. Aachen. Germany, 1966. (3) Suganuma. K.. Okamoto. T.. Koizumi, M.. Shimada, M.. Effect of interlayers in ceramic-metal joints with thermal expansion mismatches. American Ceramic Society. C-356-7. 1984. (4) Cao, H. C.. Thouless. M. D.. Evans. A. G.. Acta Metall. 36(1988)8.2037-46. (5)Munz D.. Fett. T.. Mechaniszhes Verhalten keramischer Werkstoffe. Springer Verlag. Berlin. 1989. (6) Furrer. P.. Structure of rapidly solidified alumi.lium alloys. Dissertation. University of Stuttgart. 1972. (7) Kreye. H.. Melting phenomena in solid state welding processes. Welding Research. (1977)154-157-s. (8) Bdssani. I. A.. Beitrag zur Ultraschallfuegetechnologie am Beispiel von Keramik-Metall-Verbind-ungen.Dissertation. University of Sruttgart. Germany. 1999. (9) Perch. IG. J.. Armstrong. R. Codd. I.. Douthwaite. R. M.. The plastic deformation of polycristalline aggregates. Philosophical Magazine. 7( 1962) 8. Ser. (lOY Reuter. M.. Das Ultraschallschweissen von Glas und Glaskerainik mit Metall am Beispiel des Aluminiums, Dissertation. Univ. Kaiserslautern. Germany, 1993. (1 1) Wagner. J.. Ermittlung mechanischer Festigkeitseigenschaften und thennischer Eigenspannungen an ultraschallgeschweissten Keramik Metall - Verbunden, Dissertation. Univ. of Kaiserslautem. Germany, 1997. ( 12) Bassani. I. A.. Kieckow. F.. Pazos. R. P.. Modelagem matematica de sistemas mecanicos. Anais do XXVI Cobenge. Brazil. 4( 1998) 1877-91. (13) Weibu1l.W.. The phenomenon of rupture in solids. Ingenioers Vetenskaps Akademien. Handlingar Nr. 1 53( 1939).
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AN ENGINEER’S APPROACH TO DESIGN CERAMIC COMPONENTS S. Kruger, T. Kentschke*, H.-J. Barth Technical University Clausthal, Institute For Tribology and Machine Kinetics, D-38678 Clausthal-Zellerfeld, Germany
ABSTRACT
PROCEDURE
By use of the Weibull function it is shown that the computation of ceramic components can be simplified, if the Weibull Modulus is high, as is for silicon nitride or zircon dioxide. An approximate solution is derived for notch effects that uses the notch factor and leads to a Weibull statistical estimation. Based on these results an engineer’s strategy for the design of ceramic components is proposed.
The aim of a component computation is to assure reliable function of the component during a limited life-time. For metallic materials the component is considered safe, if the highest equivalent stress of the component approximates an admissible stress value. However this can not be applied for ceramics because of the wide scattering of material strength. The admissible strength is then replaced by the admissible probability of component failure Pfli,. In practice values of pflimare lo4 to
INTRODUCTION In practice a calculation of ceramic components does not take place in most cases. Reasons are the wide scattering of material strength and life-time and the missing of a lower strength. These material properties make a statistical procedure necessary and lead to the result “the component resists with a certain probability”. Even for small loads the probability of failure is never equal zero, which is unusual for the engineer. For statistical computation the Weibull function (1) is available since the 30s but is rarely used in practice without computer software. Various research laboratories world-wide have developed software (e.g. 2,3), which allows a statistical analysis of a known stress state. The stress state is determined by a common Finite Element System (FE). Although these programs are of great advantage the results can not be better than the model and material data used. Furthermore FEcalculations require special knowledge and time. The result of this situation is that in practice neither a simple estimation of the probability of failure nor a FE-calculation is done when developing ceramic components (few exceptions exist). It is the intention of the authors to demonstrate that under certain conditions a calculation of ceramic components without using a FE-system is feasible and meaningful.
In collaboration with the ceramic industry a list of requests for a practical computation method was prepared: a) No FE-calculations (-+ “global viewpoint”, Weibull statistics), b) High reliability of computation results (“on the safe side”); the material efficiency is of small importance, c) Use of less and measurable material data, d) Consideration of multi-axial stress conditions when necessary, e) Defined scope, f) Simple mathematical procedures. Below the risk of rupture Bv is used instead of the probability of failure Pw. The relationship between both can be seen from the Weibull function:
Pw (0)= 1 - exp (- B Y ) where Bv is
for the two-parametrical distribution function. Herein is: component’s volume characteristic volume omax = maximum stress value oov= characteristic strength mv = Weibull modulus IN = normalised integral V
=
Vo
=
339
The difficulty of a computation without a Finite Element System is to determine the integral IN, which depends on the components geometry and the load case. Solutions exist for only a few cases of IN.
The limiting value for A 2:
lim
A+-
For a simplified computation procedure it has to be clarified, if in any case an integration over the total volume as well as the total surface has to be done. The alternative is to consider only the highest stressed component areas, which could reduce the computation expenditure enormously.
"J," (
--.
-+ 00 is according to equation
___
1"
= (4)
"'V
l i m - - . 1L ( L ]
Am,-'
A+-
vo
oov
HOMOGENEOUS STRESS STATE
because m > 1. This means:
For homogenous stress states the calculation of the probability of failure is simple, because the stress does not depend on the position. A typical example is a tension bar. The tensile stress (J
P,+O
(T=-
F
(3)
A
can be used directly in the Weibull function (equation 1,2) instead of (J,,,~ and IN is set to one. Because the Weibull function is only valid for tension a hypothesis has to be used for compression. A number of hypotheses are available in the literature. The tension bar is useful to demonstrate the consequences of the volume-dependent strength. It is important to know, that the characteristic strength o0v and the characteristic volume Vo are pairs of values. When setting the volume Vo in equation 2 the appropriate strength oovhas to be put in for which oovis valid. In contrast to metallic materials, the engineer has always to think in three dimensions while designing ceramic components. To calculate the probability of failure of the tension bar both the cross-section A and the length L has to be taken into account: (figure 1). For metallic materials only the cross-section is of importance.
Figure 2 demonstrates the consequences for designing. As shown before the risk of rupture decreases as the diameter increases (figure 2a). Therefore the risk of rupture of a shouldered bar (figure 2b) will decrease, if the smaller diameter is adapted to the greater one.
-
higher probability of failure lower probability of failure
figure 2: Influence of the design on the probability of failure
0=
FIA
d
IF figure 1: Tension bar Increasing the length L leads to a higher probability of component failure, but increasing the cross-section A as well as the diameter d results in a lower probability of component failure:
340
if A + m a n d
-----
4F
=o
INHOMOGENEOUS STRESS STATE Mostly the stress state of components is inhomogeneous, e.g. in case of bending, torsion or combined load cases. Furthermore inhomogeneous stress states arise from the geometry of the component, especially notches lead to local stress intensities. It is important to clarify the contribution of stress intensities to the risk of rupture. If the computation could be reduced to the higher stressed areas alone, expenditure for computation would sink rapidly.
The total risk of rupture of two volume elements, e.g. of a shouldered bar (figure 3), is according to equation 2:
Herein for simplification the notch effect is neglected, which would not be admissible for a real components computation. Additionally homogeneous load transmission is assumed. Division of equation 6 by BGresults in:
of 02/01and m allow to reduce the computation expenditure to the higher stressed component area. However a criteria has to be found, which helps the engineer to distinguish between neglectable and not negtectable areas. This is the relative calculation error r: (9)
r exists, when the risk of rupture BI is neglected in the calculations. A given maximum value for r ( e g 1 %) leads to the equation
(7) This sum was chosen in figure 3. With the assumption of equal volumes (V, = V2 = V), equation 7 gives the following:
BG
if V, = V2. Up to now the special case V, = V2 = V was discussed, which rarely arises in practice. Therefore the next step is to investigate the more general case Vl # V2 yet still with the restriction of different and homogeneous stress states in the volume elements. Once again the shouldered bar serves as an example (figure 3). Analogous to equations 8.1 and 8.2 is:
I+[$
B G
if A, > A2.
-B, -
1
B
1
2=
(1 1.1)
(11.2)
The limiting values for lim , lim are: m+O
lim :
m+m
BI/BG= 0,5 ;BJBG = 0,5
m+O
For small values of the Weibull modulus both the partial risk of rupture BI and B2contribute substantially to the total risk of rupture. The theoretical limiting value of m = 0 results in equal partial risks of rupture B, and B2 of 50 % independent from the stress state. For large values of the Weibull modulus the total risk of rupture BG is determined predominantly by the risk of rupture B2 of the higher stressed area. The limiting case m = co corresponds to the deterministic computation of components (B2 = 100 YO.)
Using r = BI/BG (see equation 9) it follows
if r < < 1 and
does not exceed a given value of r. Same examples have been calculated in (4).
The contributions of B1, B2 to the total risk of rupture depend substantially on the proportion of A2/AIas well as 02/olbecause of the exponential law (see equations l, 2). If the proportion of stress 02/o,is high the total risk of rupture BG is dominated by the risk of rupture B2, even for small values of the Weibull modulus. The conclusion is that certain combinations
34 1
DETERMINATION OF WEIBULL’S NORMALIZED INTEGRAL
I, =
In the last chapter it is assumed that the stress condition is “partially homogeneous” to show the influence of different stresses on the total risk of rupture. Actually this assumption is too broad for a components computation because local stress intensities are mostly connected with high stress gradients, e.g. if notches are present. According to equation 2 for inhomogeneous stress conditions the nonnalised integral IN has to be used to calculate the risk of rupture. Because of integration difficulties only few analytical solutions are known for relatively simple geometries and load cases (e.g. uniaxial tension / compression, bending of a beam). The evaluation of local stress intensities within a component is not possible with those solutions but would be necessary for a computation.
(15)
Equation 15 gives a general approximate solution for notches and enables simple evaluation per Weibull statistics of notch effects without the help of a FEsystem. This is a substantial request to allow an engineer’s computation of ceramic components in practice. Now the volume V corresponding to IN has to be determined. This is the area, where the maximum stress omax decreases to the nominal stress one,,,,. It is also shown in (4), that this area can be described by a semicircle with the radius rak: (figure 4)
’( b,j
r,, = - 1--
To find a simple approximate solution for a normalised integral IN, which can be used to calculate the risk of rupture of local stress intensities, two assumptions are made: a) Negligible influence of multi-axial stress conditions, b) linear stress gradient. x
Using the notch factor c(k
ak =-CJ max 0nenn
the approximate solution of IN is (4)
figure 4: radius rak
i \&
risk of rupture
4B
\L,
I
a..
..
0
5
10
15
20
25
35
30
Welbull modulus m [-I
figure 3: partial risks of rupture B1, Bz related to the total risk of rupture BG=B,+ Bz ,V1= V2 = V
342
For a practice computation a notched bar can be separated in three areas of damaging: Around the notch root is a circular zone with a very high risk of rupture BI. Because of the reduced cross-section in the middle of the notched area the risk of rupture B2 is higher than in the remaining area. The nominal stress oN in the notched area is assumed to be homogeneous. The risk of rupture B3 of the remaining area is caused by the operating stress oB. The total risk of rupture BG is the sum of the particular risks of rupture and is determined substantially by the risk of rupture BI.
For m = 5 the risk of rupture B3 is largest. Thus fracture of the plate is most probable in the non notched area. The reason is the very high scattering of material strength (m=5). table 1: results of FE-calculations and calculation “bv hand” FEM approximate solution M M=22 M = 10 M=5
Bl 0,012
1
I
0,054 0,088
B2
5,9” 0,002 0,022
I I
B3 4,0-’
0,002 0,682
I
I
Pf
Pf
0,012
0,012 0,055 0,657
0,057 0,631
EXAMPLE To evaluate the accuracy of the approximate solution a FE-model of a ceramic plate was analysed with the statistical processor “KerB” (2,3) (figure 5). The notch factor for this case was taken from the literature. As can be seen from table 1 the probabilities of failure according to the approximate solution and the FE-calculation are nearly identical. Furthermore the results show the relationship between the risk of rupture and the Weibull modulus: For m = 22 the probability of failure is determined substantially by the risk of rupture BI of the notch roof. For m = 10 the risks of rupture B2, B3 are equivalent and can not be neglected, when calculating the total risk of rupture.
DISCUSSION The investigations have lead to the following important conclusions: a) For m 2 20 the probability of failure is dominated by the highest stressed areas. Integration over the total volume is not necessary: computation is reduced enormously b) The elimination of local stress intensities (notches !) reduces the total probability of failure, even when the volume increases. The homogeneous stress condition is of most advantages, what offers probabilities for an automatic optimisation procedure (5). c) The use of average Weibull parameter can cause insufficient estimation of the risk of rupture (shown in (4)), even when using FEsystems. The calculations would be more reliable if the lower Weibull parameter or the 95 % confidential interval were used. Unfortunately in most cases those material data are not available, what makes an engineer’s computation difficult. For designing ceramic components the following procedure can be recommended to the engineer: a) Clarification of load state (even small deviations in load assumptions can lead to false estimations of the risk of rupture), b) Consideration of a safety factor regarding the tolerable stress, c) Use of reliable material data, d) Identification of the mostly stressed component area, e) Linearization of the stress state in this area, f) Computation of Weibull’s Integral respectively the effective volume, f) Summation of the risks of rupture and derivation of the total probability of failure. The proposed procedure can not replace a reliability analyses with FE, when the stress state is complex. But for a number of cases sufficient results can be obtained without FE and can support the engineer’s designing process in practice.
REFERENCES (1) Weibull, W.: A Statistical Theory of the Strength
figure 5: FE-model of a notched plate
of Materials, Ingeniors-vetensk. Handlingar Nr. 151; Stockh. (1939); s. 5 -45.
343
(2) Jakel, R.: Ein Beitrag zur Berechnung und konstruktiven Gestaltung keramischer Bauteile, angew. am Beisp. eines keramischen Ventilatorrades, Diss. TU Clausthal(l996). (3) Kriiger, S.; Jakel, R.; Rubio, D.; Barth, H.-J.: Berechnung keramischer Bauteile mit dem neuen statistischen Software-Prozessor ,,KerB"; Konstruktion 5 1 (1999); S. 33-36. (4) Kriiger, S.: Ein Beitrag zur praxisgerechten Dimensionierung keramischer Bauteile bei mehrachsigen Beanspruchungen; Dissertation TU Clausthal, Papierflieger (1999), ISBN 3-89720-345-6. (5) Kentschke, T.; Kriiger, S.; Barth, H.-J.: Shape optimisation of ceramic components by CAOmethod, International Conference on engineering Design, Proceeding Volume 2, Munchen (1 999), S . 1061 - 1064.
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IV. Cost Effective Manufacturing
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ADVANCES IN BRAZING OF CERAMIC MATERIALS FOR ENGINES P. Sire, A. Gasse, F. Saint-Antonin* CENCEREM-DEWService of Materials Engineering, 17 rue des Martyrs, 38054 Grenoble Cedex 9, France ABSTRACT This paper reports some recent advances, performed at CEAKEREM, in the brazing of S i c based materials, alumina, aluminium nitride and also aluminium metal matrix composites with Sic particles. Some applications for engines are presented.
Brazing technology has several advantages that have been extensively described in [3]. The more important ones for ceramic joining are the ability: - to achieve extensive joint area, - to perform complex or multi-component assemblies, - to join dissimilar materials.
BRAZING PARAMETERS INTRODUCTION Properties of ceramics make them very attractive for a large range of applications that are already significant in volume for engines. Their use requires the implementation of high-performance joining procedures to take hlly advantage of their remarkable characteristics. Brazing is a suitable solution, but performances of commercially available filler alloys are not always adapted to the high demanding requirements, and thus, development of specific brazes is necessary. Brazing alloys must be able to create a 'link' between the materials to join but also must have properties close to the joined part or must not be the weak point of the assembly during operation. In some case, the brazing alloy must exhibit a special property or ensure a specific function: for instance, in AINkopper joining, the brazing alloy must achieve a good heat transfer between the two materials for heat management. CEAKEREM has developed several brazing alloy grades for various ceramics taking into account the full requirements of the final applications.
JOINING OF CERAMICS IN ENGINE The main applications of ceramics in engines are, for instance: turbine, seals, bearings, blades, filters, electrical heaters, pump parts, nozzles flaps, combustion chamber, heat exchangers. The last two are generally made of several ceramic parts that must be joined together. For the other applications, the ceramic parts needs to be joined onto a metallic one which is generally a structural part. Several joining technologies are existing, but they are not always adapted to ceramics due to their inherent brittleness: for instance, mechanical joining such as riveting or bolting. Some others are in development and/or are devoted to specific applications for geometrical reasons: for instance, laser welding [ l ] or joule heating [2].
Brazing of ceramics and/or ceramic onto metal need to take into considerations several parameters in order to achieve a good joint quality. A brazing operation is described in figure 1 with a temperature versus time curve.
T
Time Figure 1 : the different steps of the brazing process. Three different phenomena are necessary to produce a sound joint: - after melting of the brazing alloy, all the surface to be joined must be covered bv the liauid, - then, a must be created during the liquid state of the brazing alloy andor, during the solidification at the beginning of the cooling step, - as the sample is finishing to cool down, the parts have to remain ioined. These three steps have been described in various books or proceedings: see for instance [4-61. COVERING THE SURFACE
The two phenomena controlling the surface covering are the wetting and spreading behaviours of the brazing alloy. Wetting is characterised by the contact angle 8 obtained during a sessile drop experiment (see figure 2). A good wetting behaviour is obtained for a contact angle lower than 90" elsewhere it is a non-wetting behaviour. For a contact angle lower than 30°, the
347
behaviour is called capillary: in this case, it is not necessary to pre-place the brazing alloy on the surfaces to join, the alloy can flow toward it by a good control of the gaps. liquid , solid
solid
Figure 2: sessile drop experiment. The spreading behaviour is a function of the reaction, occurring at the drop boundaries, between the solid, the liquid and the atmosphere. This parameter will control the duration of the brazing process.
strength of the base materials or the brazing alloy, cracks andor decohesion may occur. Thus, the joint design, taking into account the brazing alloy and the brazing cycle, is a key point in the development of such technology. Some design rules are described in [3, 71 for simple geometries. Finite Element modelling gives a more precise description of the cooling effect on the distribution of residual stresses at, and near, the joint. But it needs experimental validation before this modelling can be adopted with confidence for the design of joined components. The use of interlayer material with intermediate thermal expansion property or ductile behaviour for stress relaxation, can also limit the presence of residual stresses. The use of interlayer materials induces some geometrical constraint: for instance, it needs the design of larger joint area.
CREATION OF A LINK
There are classically two ways to induce a 'link' at the interface: creation of a bond at atomic level and mechanical anchorage. The last one is reached by a control of the surface roughening (figure 3). In order to take full advantage of the mechanical anchorage, the filling of the roughening defects needs a good wetting behaviour.
and the brazing alloy ones. This mechanism induces the formation of a reactive layer at the interface between the materials to join and the brazing alloy. The brazing alloy is generally made of a matrix element + a reactive one: for instance, in the well-known Ag-Cu-Ti alloy, Ti is the reactive element inducing the formation of titanium oxides when used for the joining of oxides ceramics. It is generally necessary to use this principle when the brazing alloy is of metallic type and the materials to join are of ionic or covalent type. When the materials to join and the brazing alloy are both metallic type, non-reactive brazing alloy can be used. KEEPING THE PARTS JOINED
Generally, there is a difference in the thermal expansion (or contraction) between the brazing alloy and the base material and, for ceramic-to-metal joints, between the two base materials. During the cooling, this thermal expansion mismatch induces residual stresses: these residual stresses can induce bending and/or deformation of the joined parts. When they exceed the mechanical
348
BRAZING ALLOYS DEVELOPED SILICON CARBIDE MATERIALS
The main difficulty for high temperature brazing of Sic is the strong reactivity of Sic with all pure metals or alloys: this reaction induces generally cracks, pores and possible decohesion (see figure 5). A brazing alloys family, BraSiC", has been developed for various applications [8]: high temperature applications up to 1600"C, resistance to corrosive media up to about 300°C, low activation under neutron irradiation. The typical joint structure is reported in figure 6.
Figure 5 : typical SiC-metal Figure 6: joint structure reaction at the interface. with BraSiC". The main properties of BraSiC@alloys are: - very good wetting with contact angles from 20 to 40" in vacuum and also, in neutral gas with oxygen partial pressure lower than lO"mbars, - no-reaction with the substrates (figure 7), - strong adhesion at the BraSiC@/SiCinterface. The mechanical strength of brazed joint measured on 4 points bending samples brazed with BraSiC@at various temperatures is reported in figure 8. The bending strength of SiC-BraSiC"-Sic is equivalent to those of Sic. The fracture occurs in bulk Sic and not in the joint. Corrosion resistance has been evaluated under hot air up to lOOO"C, in combustion gas up to 900' and in various acid solutions [8].
of interface between the SIC (bottom) and BraSiC" alloy (top). Atomic planes of the two compounds are visible: no interfacial compounds can be detected at this scale (distance between two white dots in Sic is about 219).
Figure 10: SiC/SiC composite brazed with BraSiC@. The brazing alloy thickness is about 20pm. ALUMINA MATERIALS
Capacitive sensors will be used in for aero-gas turbines [ 101 for the measurement of the clearance between the blade and the casing: the fuel consumption should then decrease. A brazing alloy based on the Pd-Mg system has been developed for alumindmetal joining. This brazing alloy can withstand high temperature up to 1000°C and corrosive or oxidising atmosphere. The typical structure of the BrasOx@/aluminajoint is reported in figure 11. The mechanism of the reactive layer formation is based on diffusion of Mg within the alumina [9].
brazed with Pd-Mg alloy (the diameter is about 6mm).
349
ALUMINIUM NITRIDE MATERIALS
Aluminium nitride (AIN) materials have very high thermal conductivity of about 200W.m-'.K-' that is half the copper one. The mechanical properties are about the same than Sic with a Young Modulus that is about 213 of the Sic one. CEAKEREM has developed a brazing alloy for AIN-to-metal joining able to transfer very efficiently the heat from AIN to the metal [ 1 I]: this heat transfer performance is about 6 to 8 times higher compared to the one obtained with the Ag-Cu-Ti type brazing alloy. This has been achieved by a very good control of 1) the reactivity between the brazing alloy and AIN and 2) the interface structure. Figure 13 gives a typical interface structure between copper and AIN.
performed with the brazing alloy BmAl@:the working conditions should not be higher than 250°C as the solidus is about 420°C.
Figure 15: Al globules within AI-Ge eutectic.
CONCLUSIONS AND PERSPECTIVES
Figure 13: CdAIN interface. The interface layer cannot be seen at this scale. CERAMIC-METAL MATRIX COMPOSITES
Composite materials made of Al as matrix and Sic fibbers or particles used for reinforcement are candidate for engine parts due to their high thermal conductivity, high strength and rigidity, and low coefficient of thermal expansion. An alloy family, BmAl@,based on the AI-Ge system, with a brazing temperature ranging from 480 to 550"C, has been developed [12-131. Figure 14 gives a typical view of AI-SiC/AI-Sic joint with the brazing alloy BrasAl@.
During the last decade, lot of progress has been achieved in the field of joining ceramics and ceramic-tometal by brazing for Sic, alumina, aluminium nitride, metal matrix composite with ceramic particles: it is then possible to integrate this joining technology for the assembly of engine parts and components. Medium and long term future researches should be focused on: - the definition of rational joint design rules as most of the existing ones are strongly based on empirical results, - the development of predictive tool for joint life estimation: this development is strongly related to the precedent point, - the definition of methodology for the joint strength evaluation as so many test methods have been proposed and none seems to be fully satisfactory ~41. The development of ceramic-to-metal joining should receive attention due to the difficult achievement of reliablejoints when there are high mechanical loading.
REFERENCES
a
(1) A.M. Nagel, H. Exner, Laser beam weldin
(2)
Figure 14: the brazing alloy thickness is about 200pm.
(3)
Due to the fabrication process, the brazing alloy is very ductile: the alloy can then be formed in various shape and size. The structure after brazing is globular (figure
(4)
15).
It should also be noticed that aluminidalumina and aluminidmetal (stainless steel, aluminium) joints can be
350
(5)
Of alumina: a new successful technology. Proc. 7 Int Symp. Ceramic Materials and Components for Engines, Goslar, Germany (2000) in press. J.P. Kay, J.P. Hurley, Joining of silicon carbide by joule heating. Proc. Materials Conference '98 on Joining of Advanced and Specialty Materials, Rosemont Illinois, USA (1 998) 7- 10. M. Schwartz, Brazing for the Engineering Technologist. Chapman & Hall (1 995) 1-7. Joining of Ceramics. Edited by M.G. Nicholas, Chapman & Hall (1 990). Ceramic Joining. Edited by 1. Reimanis, C. Henager, A. Tomsia, Ceramic Transaction,
published by the American Ceramic Society, vol. 77 (1 998). (6) M. Schwartz, Joining of Composite Matrix Materials. ASM International, (1 994). (7) Brazing Handbook. Published by the American Welding Society (1994) 9-4 1. (8) F. Moret, P. Sire, A. Gasse, Brazing of Sic using the BraSiC process for chemical and thermal applications. Proc. Materials Conference '98 on Joining of Advanced and Specialty Materials, Rosemont Illinois, USA (1998) 67-72. (9) 1. Guesdon, F. Saint-Antonin, L. Coudurier. N. Eustathopoulos, Morphological and chemical characterisation of reactive a-Al2O3 brazed joints. Proc. European Ceramics Symposium, Brigthon, UK, (1999) 333-334. (1O)P. Sire, F. Saint-Antonin, 1. Guesdon, Reactive brazing of oxide ceramics for high temperature applications with the BrasOx process. Proc. Conf. Materials Solutions, ASM International, Cincinnati Ohio, USA (1999) in press. (1 1)T. Baffie, F. Saint-Antonin,Joining of AIN to copper by Brazing. Proc. of the European Ceramics Symposium, Brigthon, UK (1999) 329330. (12)J. Valer, F. Saint-Antonin, P. MCnCses, M. SuCry, Influence of processing on microstructure and semisolid behaviour of AI-Ge alloys. Proc. of the SemiSolid Processing of Alloys and Composites, Denver Colorado, USA, (1998) 3-10. (13)J. Valer, F. Saint-Antonin, P. MCneses, M. SuCry, Microstructural and mechanical characterisation of an Al-28wt'?/&e brazing alloy with a globular morphology of the primary ALRich phase. Material Science and Engineering,A272 (1999) 342-350. (14)F. Saint-Antonin, How to characterisejoints ?, to be published.
35 I
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STEREOLITHOGRAPHY FOR CERAMIC PART MANUFACTURING C. Chaput*,T. Chartier**,F. Doreau*, M. Loiseau* (*) CTTC,
Ester Technopole, BP 6915,87069 Limoges, France (**) SPCTS, UMR CNRS 6638, ENSCI, 47/73 avenue Albert Thomas, 87065 Limoges, France ,ABSTRACT The fabrication of a prototype with complex shapes in ceramics, that are very hard and brittle materials, is very difficult and costly. In most cases, it is necessary to develop a mould for casting or injection moulding and subsequently to finish ceramic pieces by machining. Rapid prototyping presents a great opportunity to realise ceramic prototypes. The first step in fabricating ceramic pieces by stereolithography is to prepare adapted ceramic suspensions with the appropriate rheological behaviour. In our case, we have developed ceramic pastes with high viscosity and adapted yield value to work with a new French Rp machine. The main advantages are the stability of the suspensions, no relaxation time for the menisci and rapidity to realise each layer. After fabrication by stereolithography, ceramic parts are debinded and sintered. The obtained ceramic parts exhibit good mechanical properties, similar to those obtained by classical ceramic processing. This technology is very interesting to manuhcture complex ceramic shapes in a short time and by reducing costs.
INTRODUCTION Shaping of complex ceramic parts is performed by various methods using moulds such as slip. casting, centrifigal casting, gel casting, direct coagulation casting or injection moulding. High pressure injection moulding is largely used for industrial production of ceramic parts (generally small ones) of complex shape and high dimensional tolerances. One factor limiting a wider application of high pressure injection moulding is the high initial tooling cost that makes this technique not adapted for small runs and of course for the fabrication of prototypes. Over about fifteen years, many heform processing methods of 3D ceramic parts without moulds or tooling have been developed, for instance selective laser sintering, three dimensional printing, fised deposition modelling, laminated object manufacturing and tape casting techniques. An attractive route is to transpose the method of stereolithography, which is largely used for the fabrication of three dimensional polymer parts, toward the process of 3D ceramic pieces with final properties (mechanical, thermal, electrical...) close to those obtained by classical processing techniques [ 1-41.
Stereolithography involves polymerisation of a reactive system, generally based on acrylate or epoxy monomers, by a space resolved laser. The main applications of the fabrication, by stereolithography, of ceramic parts should be medical implants, the direct fabrication of refractory moulds and cores or of prototypes prior to defining an expensive mould of injection moulding. This paper describes the fabrication of complex alumina parts by stereolithography, removal of the polymer and sintering. It also gives some characteristics of the sintered parts, as flexural strength and dimensional resolution that can be reached today.
EXPERIMENTAL PROCEDURE Starting Materials The alumina powder used (CT1200SG, ALCOA, USA) has a mean particle diameter of 1.5 pm and a specific area of 3.4 m2.g-'. In order to increase the ceramic fraction, then to reduce the critical shrinkage during polymerisation and during sintering, an efficient dispersant which acts both by electrostatic and steric repulsion is used. A thickener confers a high yield value to the paste to prevent settling of particles and to support the piece during fabrication. The monomer is a diacrylate (HDDA, UCB, Belgium). The photoinitiator absorbs in the range of the UV laser emission (Irgacure 65 1, Ciba, Switzerland).
Preparation of the Suspensions The photoinitiator (0.5 % by weight of monomer), the dispersant (2 % by weight of powder) and the thickener (0.5 % by weight of monomer) are first dissolved in the diacrylate, then the alumina powder is added. The suspension containing 62 vol Yo alumina, is milled during 30 min to break down the agglomerates and to achieve a good homogeneity.
Fabrication of Green Ceramic Pieces by Stereolithography Homogeneous layers, with different thickness (from 25 to 100 pm), with smooth surfaces are spread by means of a specific device before polymerisation. Contrary to classical stereolithography equipment, where a fluid monomer is required, our machine uses a
353
paste with a high viscosity, thus avoiding the use of a container. Using CAD information, laser beam radiation (Arion laser, Coherent, h = 351-364 nm), focused on the top surface of the deposited paste, is deflected by galvanometric mirrors. When a layer of the part is performed, the deposition of a subsequent layer of paste on the already polymerised part allows to continue the manufacturing process. This procedure is repeated until the polymer part is built.
.t
9 [Pa4
F 1
3500
14ooo
P
2000 1500
Binder Removal and Sintering Based on thermogravimetric analysis of green samples, the debinding is performed with a heating rate of 1"C.min-' up to 120°C, then of 0.2"C.min-' up to 550°C with a stage of 3 h. The polymeric phase is then completely removed. The parts are finally sintered with a heating rate of 5"C.min-I up to 1700°C with a 1.5 h plateau. The density of sintered pieces is 97% of the theoretical density.
Characterization Rheological measurements are performed with a controlled stress rheometer (RS150, HAAKE, Germany) using a cone-plane configuration. The reactive system has to verig two main requirements, i) the cured depth must be high enough to avoid an excessive time of fabrication and, ii) the cured width must be low enough to ensure a good resolution. In this respect, cured depth and width are measured on small polymerised lines cured in one layer by one scanning of the laser beam. The brittle polymerised lines are included in an epoxy resin in order to cut sections for observation. The values of depth and width correspond to the average of four measurements. The flexural strength of as-sintered and polished alumina bars (3.5~5.5~40mm3), fabricated under similar UV exposure conditions but with six different scanning patterns is measured by threepoint bend tests (average of five values). Two loading directions, with respect to layer planes, are also tested. These values are compared to flexural strength of the same alumina pressed samples sintered in the same conditions.
RESULTS AND DISCUSSION The Paste A rheological behaviour adapted to the principle of the specific stereolithography technique used is required. The flow curve exhibits a typical shear thinning behaviour (Fig. 1). The high yield value (1200 Pa) allows to prevent non-insulated surfaces fiom flowing during the building of the piece and then to support it. The shear thinning behaviour allows spreading of homogeneous layers with a thickness ranging fiom 25 to 200 pm. At a shear rate of 100 s-', corresponding to the minimum value generated by the machine during layer deposition, the measured viscosity is 110 Pas, compared to 12000 P a s at rest.
354
500 0
20
40
60
r
80
100
120
140
I Fig. 1 : Suitable rheological behaviour of the ceramic paste 1s-l
Depth and Width of Photopolymerisation A first objective is to reduce the time of fabrication of pieces, then to use a high scanning speed while maintaining a sufficient cured depth. The cured depth (Ep) depends on the density of energy (DE) transmitted to the paste, on the depth of penetration of the beam (Dp), characteristic of the paste (particle diameter, volume fiaction of powder, difference between the refiactive index of the powder and of the resin ...) and on the critical density of energy (DEc) that represents the smallest DE for which polymerisation OCCUTS [5-71: DE Eq. 1 Ep = Dp In(-) DEc The density of energy DE is function of the power of irradiation P, of the scanning speed v and of the diameter of the laser spot wo (80 pm) on the working SUrface: 2.P Eq. 2 5c.w .v It has been shown that the depth of polymerisation does not vary with P/v ratio at constant density of energy, in agreement with equations 1 and 2. We can also notice that a cured depth of 200 pm can be achieved using a high scanning speed (1 m.d, DE FZ 0.3 J.cm-') for a concentrated system (62 vol % alumina). A second objective is to obtain a good dimensional resolution, then to control the width of polymerisation. The width of polymerisation could be greatly affected by scattering phenomena due to the presence of ceramic particles and could be significantly larger than the beam diameter. The classical theories of scattering (Rayleigh, Gans, Mie) are verified for diluted systems (
high density of energy whereas a high resolution requires a low density of energy, then a compromise has Osilica d50=10pm alumina d50=0.5pm
-5
Palumina d50=5pm szirconia d50=1Opm
1.8 1
R2 = 0.9989
E 1.5 E
p
g
2
,9855
1.2 -
,9938
0.9 -
0 Silica d50=1Opm
Q alumina d50=5pm alumina d50=0.5pm 0 zirconia d50=10pm
-
0.8
-5 0.6
i
R2 = 0.9793
-
R2 = 0.9868
$ 0.4 -
R2 = 0.9947
U
0.60.3
to be found, in particular for the scanning conditions of external surfaces of the pieces.
.9813
-
'0
f
0.2-
0
01 0
gfs
R2 = 0.9983
04 2
4 DE (Jlcm2)
6
0
2
4 DE (Jlcm2)
6
Fig. 2 : Cured depth and width according to density of energy DE for various powders ( 5 0 %~ )
In comparison with these results, the limit dimensional resolution of our system has been studied, by varying the scanning speed and the layer depth. Two scanning patterns have been tested : one with simple lines parallel to the spread axis (X), and a second representing a mesh
(parallel and perpendicular (Y) to X ) of 1 and 2 mm. Fig. 3 shows the results obtained with the simple lines for 10 layers of 25pm depth. A limit resolution of 170pm can be reached for simple lines.
1 mm
DE=80 mJ/cm2
I
I
I
D E 4 0 mJ/cm2
I
DE=30 mJ/cm2
230 pm 420 pm Honey comb (1 mm mesh) Honey comb (2mm mesh) Layer depth 2 5 ~20, layers. Layer depth 25pm, 120 layers. DE=50 mJ.cm-* DE=160 mJ.cm-* Fig. 4 : Honey comb parts for different densities of energy and mesh dimensions
I
I 355
Mechanical Characterization of Flexural Strength No significant differences in flexural strength are observed between the different layer orientations or scanning patterns used, attesting of a good homogeneity. An average value of 275 MPa is obtained for as-sintered samples, which is similar to that of pressed specimens (265 m a ) . When bars are polished before testing, an average value of 394 MPa is obtained for rapid prototyped specimens and 396 MPa for pressed specimens. It is then possible to fabricate
Cylinder head (350 layers) Fabrication time :5h for-2~ieces
complex shape parts by stereolithographywith similar mechanical properties than those obtained by classical shaping process like uniaxial pressing.
Sintered Ceramic Parts Obtained by Stereolithography Fig.5 presents some of the sintered pieces obtained by rapid prototyping process. The number of layers and the necessary time of fabrication are also specified for each piece, showing that complex and big parts with small details can be build in a short time.
SNECMA Core (250layers)
Break engine part (420 layers) Fabrication time :5h30 for 2 pieces
Turbine (100 layers) Fabrication time :2h30 for 10 pieces
Drawplate with stars pattern (120 layers) Fabrication time :2h30 for 8 pieces
Fig. 5 : Examples of alumina parts fabricated by rapid pr,ototypingusing a ceramic paste (layer depth : 100pm)
CONCLUSION The major drawback of the fabrication of complex shape ceramic parts is the cost of the mould, the difficulties to obtain different complicated cross sections and the processing time required to develop the first prototype. Rapid prototyping offers a unique opportunity to fabricate, without using costly moulds, complex ceramic parts with mechanical properties similar to those obtained by classical processing routes, in a very reduced time.
356
A key point is the definition of a suitable photocurable paste which has to fulfil many requirements, in terms of rheological behaviour and in terms of W reactivity.
ACKNOWLEDGEMENTS We want to thank the French organisation for the research valorisation, M A R , for its financial support.
REFERENCES [ 13 C. Hinczewski, S. Corbel and T. Chartier, "Ceramic suspensions suitable for stereolithography," Journal of the European Ceramic Society, 18 (1998), pp. 583-90. [2] C. Hinczewski, S. Corbel and T. Chartier, "Stereolithography for the fabrication of ceramic threedimensional parts," Journal of Rapid Prototyping., 4 [ 3 ] (1998), pp. 104-1 1. [3] M.L. Griffith and J.W. Halloran, "Freeform fabrication of ceramics via stereolithography," Journal of the American Ceramic Society, 79 [lo] (1996), pp. 2601-08. [4] H.Liao and T.W. Coyle, "Photoreactive suspensions for stereolithography of ceramics," Journal of the Canadian Ceramic Society, 65 [4] (1 996), pp. 254-262. [5] Brady G.A., Chu T.M., Halloran J.W, "Curing behavior of ceramic resin for stereolithography", Proceedings of the Solid Freeform Fabrication Symposium, University of Texas, Austin, p. 403-4 10, 1996. [6] Grifith M.L., Chu T.M., Wagner W.C., Halloran J.W., "Ceramic stereolithography for investment casting and biomedical applications", Proceedings of the Solid Freeform Fabrication Symposium, University of Texas, Austin, p. 31-38, 1995. [7] Hinczewski, "Stereolithographie pour la fabrication de ckramiques", PhD Thesis, Institut National Polytechnique de Lorraine, October 1998.
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APPLICATION OF THE MOLD SDM PROCESS TO THE FABRICATION OF CERAMIC PARTS FOR A MICRO GAS TURBINE ENGINE Sangkyun Kang*, Jurgen Stampfl, Alexander G. Cooper and Fritz B. Prinz Stanford University, Stanford, CA 94305-3030, USA
Abstract A micro gas turbine engine with silicon nitride parts is being developed. In this project, the Mold Shape Deposition Manufacturing (Mold SDM) process is used to fabricate high quality ceramic parts with complex shapes such as the rotor group. The merits of micro gas turbine engines in general are described before focusing on processing and fabrication issues. The obtained silicon nitride parts are characterized concerning their mechanical and microstructural properties. The surface roughness, shrinkage during sintering, final density and achievable feature sizes have been determined. Using Mold SDM a functional rotor group has been successfully fabricated. 456,000 rpm rotational speed has been acheved during the spin tests at room temperature with nitrogen as driving gas.
INTRODUCTION Micro gas turbine engine A micro gas turbine engine with silicon nitride parts is being developed by the Rapid Prototyping Laboratory (RPL) of Stanford University and its industrial partners. The engine is designed by M-DOT Aerospace (Arizona, USA) and the RPL is responsible for the manufacturing and materials processing of the silicon nitride parts. Figure 1 shows the original design of the micro gas turbine engine developed by M-DOT. This one is similar to current engine but it does not use ceramic parts. This type of engine can be used as a portable energy source or in small flying vehicles.
Silicon nitride rotor group One of the key parts for the success of the micro gas turbine engine project is the rotor group. The CAD model of the rotor group is depicted in Figure 2. It is a monolithic part composed of the rotor shaft, compressor and turbine (from left to right). The diameter of the turbine is 12 rnm and the minimum blade thickness at the tip is 220 pm. The part is designed to operate at 800,000 rpm.
Figure 2 CAD model of the silicon nitride rotor group of the micro gas turbine engine.
Silicon nitride was selected as the material for the rotor group due to its superior high temperature properties and lower density compared to superalloys. The turbine blades of the micro gas turbine engine cannot have the sophisticated cooling channel systems which can be found in larger engines due to their small size. So, a ceramic material is the better choice for the turbine application because of the better temperature resistance. The turbine of the small engine must rotate at a higher speed than that of a bigger engine to achieve the higher power density. This results in a substantial centrifugal force. In this case, a higher strength to density (of/p) ratio is important. Silicon nitride has higher strength than superalloys at elevated temperature and it has 1/3 of the density. So, the use of silicon nitride will make the whole system lighter and, thus, increase the thrust/weight ratio of the engine.
Advantage of miniaturization
Figure 1 M-DOT micro gas turbine engine.
One of the possible merits of the small gas turbine engine is the increased power density which is defined as power/volume. Theoretically, the power density of a gas turbine engine can be increased as the size of the engine decreases. In that case, multiple small engines can be used instead of one big engine and the small engine system will occupy less volume due to the increased power density. Since the same output power can be achieved in a smaller volume using multiple small engines, a redun-
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dant system can be implemented by adding more engines into the saved volume. A redundant system can have higher reliability. Figure 3 illustrates a redundant system. It was assumed that 16 small engines in the right system will generate the same power as the one big engine on the left. In this case, four engines are redundant.
Single engine system
Redundant system (20 small engines)
Figure 3 Single engine system and redundant system. In this system, 16 small engines are necessary to produce the same amount of power as the large engine. Simple calculation of the probability of failure based on binomial distribution shows the higher reliability of the redundant system. (The formulation is in Appendix) Figure 4 compares the failure probability of one big engine and the redundant system illustrated in Figure 3. According to the computational result, assuming the failure probability of each engine is 2x10-*, the failure probability is 2 ~ 1 0 for - ~ one engine but 3 . 8 ~ 1 0for - ~ the redundant system. Pfs 1.E-2 1.E-5
Pf1
1 F-I1 ..
0.01
0.1
Figure 4 The failure probability comparison of one big engine (dotted black line) and 20 small engines (solid line) system. Pn: failure probability of each engine, Pfs:failure probability of 20 engine system. The probability of failure of a large ceramic specimen is greater than a small ceramic specimen under the same stress and the Weibull treatment of failure incorporates this relation [l]. So, in terms of reliability, the application of ceramic parts in a small gas turbine engine is more beneficial than a large engine. The volumetric relation of the reliability of the ceramic parts can be combined with the reliability of the redundant system. The failure probability of each small engine that has ceramic parts in it will decrease due to the reduced failure probability of the ceramic parts. This will reduce the Pfl value of each small engine in
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Manufacturing process for the silicon nitride rotor group The silicon nitride rotor group (Figure 2) has a complex shape. Design iterations are anticipated during the progress of the project. So, the ability to fabricate high quality ceramic parts quickly is required. The surface quality of the rotor is important because the strength of ceramic parts is hghly dependent on the surface condition. It must also be noted that the rotor geometry does not allow postprocessing, such as grinding, due to its shape complexity. A manufacturing process that can produce a ceramic part with complex shape and good surface finish is required for this project. In terms of shape complexity, commercial rapid prototyping techniques are good candidate processes for the manufacturing of the rotor group. However, the surface quality of rapid prototyped parts is limited due to the stair step effect. The rough surfaces become good crack initiation site and thls reduces the strength of the parts significantly. The surface quality can be improved by reducing the layer thickness but the improvement is linear, which means the stair step is halved as the thickness of the layer is halved [2]. The traditional ceramic fabrication processes, such as machining of green ceramic blanks, have limitations in achievable shape complexity due to the tool access constraint. However, machining can produce parts with good surface quality and the surface quality improves rapidly. For example, in ball endmill machining, if the space between two machining paths become 112, the scallop height reduces to 114 due to the quadratic convergence property of machining [2].
MOLD SHAPE DEPOSITION MANUFACTURING (MOLD SDM)
1 .E-8
0.001
Figure 4. Therefore, fiuther improvement of the reliability can be achieved form the miniaturization.
Mold Shape Deposition Manufacturing (Mold SDM) is a manufacturing process that can be used to build ceramic, metal and polymer parts. Mold SDM is a two step process. Fugitive wax molds are first built using an additive-subtractive layered manufacturing process. Then a variety of castable materials, including ceramic and metal gelcasting slurries as well as castable thermoset polymers, can be cast into these molds to produce parts [3].
I . .
5
6
7
0
Figure 5 Mold SDM procedures.
Sample Polished unpolished parallel scallops unpolished parallel scallops unpolished perpend. scallops 0.001mm unpolished perpend. scallops 0.005mm unpolished perpend. scallops 0.0lmm unpolished perpend. scallops 0.lmm
Slurry ACR ACR SRI ACR ACR ACR ACR
Figure 5 illustrates an example sequence for the fabrication of a simple part. The mold is built up layer by layer in steps 1 through 4. Each step represents one material deposition and machining cycle. The mold material forms the mold itself while the temporary support material defines the mold cavity and provides support for undercut features in the mold. The support material is removed in step 5 and the part material is then cast into the mold cavity in step 6 . After removal of the mold, in step 7, finishing operations, such as casting feature removal, are performed leaving the finished part in step 8. Sintering can be performed after either step 7 or 8 [4]. Mold SDM has several advantages over other manufacturing processes. As a layered manufacturing process Mold SDM is able to build geometrically complex shapes. As with other layered manufacturing processes, no part specific tooling is required. This reduces lead times and costs. Short production runs or prototype parts can be built rapidly and economically. There is minimal cost for design changes since no new or modified tooling is required. Unlike most layered manufacturing processes, which are purely additive, Mold SDM is additive-subtractive. The subtraction step, performed using CNC milling, enables the creation of smooth accurate geometries and also accomplishes this without the need to use extremely thin layers which would increase build time. Mold SDM is also capable of building parts from a wide range of castable materials. To date parts have been made from structural materials including silicon nitride, alumina, stainless steel, epoxy, polyurethane and silicone. For the manufacture of complex ceramic parts, Mold SDM has two key advantages over other layered manufacturing processes. First, all surfaces are either machined or replicated from machined surfaces. This results in smooth accurate surfaces without any stair step effect. Second, the part material is cast monolithically. There are therefore no layer boundaries in the finished part. Layer boundaries are potential sources of defects due to incomplete bonding between layers, or foreign particles or voids trapped at the layer boundaries. Both of these advantages are particularly important for flaw sensitive materials where surface roughness and internal defects can significantly reduce the mechanical strength. Mold SDM currently uses a variety of waxes as the mold material, the preferred material being KC3254A wax (Kindt-Collins Company, Cleveland, OH). The wax is deposited by casting, usually at a
Average Strength (MPa) 930 414 428 406 416 288 23 1
Max. Strength (MPa) 1012 600 550 500 495 360 257
temperature between 80 and 120OC. It is removed from the final part by a combination of melting and solvent removal using BioAct 280 Precision Cleaner (Petroferm Inc., Fernandina Beach, FL). The support material is a W curable polymer (Advanced Ceramics Research, Tucson, AZ) that is deposited as a liquid, cured under UV light, and then later removed by dissolution in water. Ceramic parts are made by gelcasting using silicon nitride or alumina gelcasting formulations developed by Advanced Ceramics Research.
PROPERTIES OF MOLD SDM PARTS In Table 1 the mechanical properties for various gelcast 4-point-bend specimens are summarized [S]. To obtain these results, wax molds were machined to produce different scallop heights on the machined surfaces. In some instances the scallops where parallel to the length of the beam, in some instances perpendicular. Perpendicular scallops are expected to degrade the mechanical properties since they serve as notches for crack initiation. This expectation is reflected in the results shown in Table 1: The highest strength was achieved for polished beams. The best “as sintered” beams were the ones with parallel scallops. The larger the scallop height for perpendicular scallops, the more the strength is decreased. By generating machining paths according to the requirements for the surface quality of the mold, Mold SDM can achieve surface qualities superior to other Solid Freeform Fabrication processes. Especially for ceramic parts, this high surface quality significantly improves the part performance. Since many features on Mold SDM parts are too delicate to be processed by grinding or polishing, it is necessary to achieve a smooth surface finish directly during manufacturing. One way to improve surface smoothness is by adjusting the furnace atmosphere so that the a-p phase transformation takes places only in the bulk material, but not on the part surface. In Figure 6 these two types of microstructure are shown: The bulk microstructure consists of the typical P-needles which give the material the necessary strength and toughness. On the part surface the grain growth is inhibited, and therefore a smoother surface is obtained. By using a profilometer, the surface roughness on the bottom and top side of a sintered but not polished
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test part were measured. The bottom surfaces, where the geometry was replicated from machmed surfaces, had root mean square (RMS) roughness of 0.5-0.7pm. The upper surfaces, where the casting features were cut off by manual machining of the green part, had RMS roughness of 1.3-1.8 pm. These values compare favorably with values of 4 pm reported for ceramic parts produced by stereolithography [6].
RESULTS: SILICON NITRIDE ROTOR GROUP The sintered turbine rotor group is shown in Figure 6. Small features such as the 220 pm thick blades were successfully fabricated. This one is slightly different from the CAD model illustrated in Figure 2. This turbine rotor group has two identical turbines back to back instead of one compressor and one turbine. However, this part demonstrates the capability of the Mold SDM process to build the actual rotor group. The molds for the rotor group were prepared using 5-axis machining. Wax was used as the mold material and soldermask as the support material.
Figure 7 Sintered silicon nitride rotor group. Since the rotor group will spin at high speed, the straightness of the sintered part is crucial for the performance. A set of special sintering fixtures (Figure 8) was designed to keep the part vertical during sintering. It helped prevent distortion due to sagging. The sintering fixtures were made of green material so that the part and fixture would shrink at the same rate.
Figure 6 Microstructure of sintered silicon nitride in the bulk material (a) and on the surface (b). The linear shrinkage of Mold SDM parts depends on the solids loading of the gelcasting slurry and the final density of the sintered parts. The shrinkage was calculated by comparing the major dimensions of the sintered parts with those of the CAD model of the mold geometry. An average linear shrinkage of 18f0.5% was measured [2]. Since the radius of the rotor group is 6 111111,f0.5% results in f 3 0 pm difference in radius. The sintered Mold SDM silicon nitride parts achieved 97% of full density.
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Figure 8 Sintering configuration for the silicon nitride rotor group to avoid distortion during sintering. Table 2 shows the mold build time for the turbine rotor group. The deposition time per mold can be significantly reduced since the deposition can be done on multiple molds at the same time. It usually takes 4.5
days for the gelcasting, curing, drying, debinding and sintering.
Table 2 Build time of the turbine rotor group. Process Machining Soldermask deposition
I
Time(hr:min) 8:5 1 2:40 17:31
PostProcess
I
4.5 davs
500,OOo
450,000 400,000 350,000
300,000
250,000 200,000 150,000 100,000 50,000 0
kPa
0
20
40
60
80
100
120
Rotor group spin test The sintered turbine shaft geometry (Figure 9) has been spin tested in a test rig. It has the same turbine as the rotor group but does not have the compressor because it is not necessary for the spin test. The only post-sintering process for this part was the grinding of the shaft to make it fit into the bearings. Pressurized nitrogen (N2) was used as the driving gas. The mlet nozzle was made of polyurethane using SDM. A black line was marked using black ink on the side of the shaft to measure the speed optically. The reflected light was received by an optical fiber and the fluctuation of the amplitude of the received light was processed using a LabView spectrum analyzer to calculate the rotation speed.
Figure 9 The sintered turbine shaft used for the spin test. The achieved speed was 456,000 rpm (Figure 10). The design speed of the rotor group is 800,000 rpm. However, the turbine is supposed to be driven using hot gas which has higher energy and speed of sound. This result shows that the Mold SDM process can successfully build functional ceramic parts.
Figure 10 Spin test result of the turbine shaft geometry.
CURRENT PROCESS ISSUES Though it was demonstrated that the Mold SDM process can produce ceramic parts with complex geometry and good surface quality, there remain some issues. The first issue is the deposition of wax. The imperfect deposition of wax on the machined soldermask surfaces may generate micro bubbles on blade surfaces (Figure 9). If wax is deposited at too high temperature, soldermask will soften and sag. It causes the distorted final geometry such as sagged blades. Optimal wax deposition parameters must be found to minimize the micro bubbles and the distortion. The second issue is to preserve the straightness during the sintering. The relative angles among the shaft, turbine and compressor are crucial for parts operating at such high speeds. Currently, visible sagging is observed in the sintered parts even with the special sintering fixture set up and the repeatability of sintering process is poor. For h s issue, better alignment scheme for the sintering should be investigated.
CONCLUSIONS Mold SDM can achieve complex geometries and smooth surfaces at the same time by combining the merit of Rapid Prototyping techniques and conventional machining. Mold SDM was used as the manufacturing process for the silicon nitride parts of the micro gas turbine since it satisfies the requirements of the parts. Unpolished, sintered Mold SDM silicon nitride has shown 400 MPa mean strength, 0.5 pm RMS surface roughness, 18+_0.5%linear shrinkage and 97% of full density. Mold SDM has demonstrated its capability to build silicon nitride parts for a micro gas turbine engine. A turbine shaft geometry has been fabricated using Mold SDM and it was successfully spin tested up to 456,000 rpm.
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ACKNOWLEDGEMENT The micro gas turbine project has been b d e d by the Defense Advanced Research Projects Agency and the Office of Naval Research under contract number NO00 14-98-1-0734 titled “Solid FreeFrom Fabrication of Ceramic Components for Microturbine Engines”. The authors would like to acknowledge Tom Hasler and Tibor Fabian of the FWL for their effort on the spin test setup and measurements.. The authors also would like to thank Hansjorg Schilp of the Technical University of Munich for this effort in the fabrication of the molds for the rotor group.
APPENDIX Binomial formulation of the failure probability of redundant system The definition of variables. 0 Pfs: The failure probability of redundant system 0 Pfl: The failure probability of each engine in the redundant system. n: The number of engines in the redundant system. 0 r: The number of engines required to run the system successfully. The failure of the system occurs when more than (n-r) engines fail. Assuming that the failure of each engine happens independently, the failure probability can be calculated as the following.
References [ l ] John B. Watchtman. Mechanical Properties of Ceramics, John Wiley & Sons, 1” edition, 1996. [2] S. Kang, A. G. Cooper, J. Stampfl, F. Prinz, J. Lombardi, L. Weiss, and J. Sherbeck. Fabrication of high quality ceramic parts using Mold SDM. In D. L. Bourell et al., editor, Solid Freeform Fabrication Symposium 1999, pages 427-434, University of Texas, Austin, August 9-1 1 1999. [3] A. G. Cooper, S. Kang, J. W. Kietzman, F. B. Prim, J. L. Lombardi, and L. Weiss. Automated fabrication of complex molded parts using Mold Shape Deposition Manufacturing. Materials and Design, 20(2/3):83-89, 1999. [4] A. G. Cooper, S. Kang, J. Stampfl, F. B. Prinz, and L. Weiss. Fabrication of structural ceramic parts using Mold SDM. In Proceedings of the American Ceramics Society meeting, Cocoa Beach, FL, January 2000. [5] S. W. Nam, H.-C. Liu, J. Stampfl, S. Kang, and F. Prim. pMold Shape Deposition Manufacturing.
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M R S Spring Meeting 2000, San Francisco, April 2000. [6] G. A. Brady and J. W. Halloran. Solid freeform fabrication of ceramics via stereolithography. Naval Research Reviews, L(3):39-43, 1998.
POTENTIAL OF THE HYDROLYSIS ASSISTED SOLIDIFICATION PROCESS FOR WET FORMING OF Si3N4 CERAMICS Kristoffer Krnel* and Tomai KosmaE Joief Stefan Institute, SI-1000 Ljubljana, Slovenia
ABSTRACT The Hydrolysis Assisted Solidification (HAS) process was used in the slurry forming of Si3N4 green bodies. AIN powder was added to a highly loaded aqueous Si3N4 slurry, also containing Y2O3, prior to shaping by casting into preheated impermeable moulds. Once moulded, the suspension starts to solidify due to the reaction of AIN with water. In the course of this reaction the AIN consumes water from the suspension and an aluminium-hydroxide gel is formed, which is accompanied by an increase in viscosity until ultimately a saturated body is formed. Two Si3N4 powders were used in this study, which differ in quality, properties and price. The results indicate that the HAS process can be used for the slurry forming of high-performance as well as for lower performance Si3N4 ceramics since the properties of the sintered ceramic parts do not differ from conventionally prepared Si3N4ceramics.
INTRODUCTION Hydrolysis Assisted Solidification (HAS) is one of the Aqueous Injection moulding (AIM) techniques in which a highly loaded aqueous suspension containing a small amount (1-5 wt.%) of well dispersed AIN powder is poured, cast or injected into an impermeable mould, where it solidifies'. In the presence of water in the suspension, AlN will hydrolyse following the reaction scheme2:
AIN powder will not hydrolyse immediately, since a thin amorphous layer of hydrated alumina protects the surface of the powder3. An incubation time is needed for the water to dissolve this layer. This incubation time depends on the thickness of the protective layer and its solubility in aqueous environments. Once the incubation time has expired the hydrolysis will proceed resulting in an increase in the viscosity of the host slurry which finally leads to solidification. In the HAS process there are several mechanisms contributing to the solidification of the host slurry: water consumption, electrostatic destabilization due to ammonia formation and particles dissolution/ precipitation. Water consumption increases the solids content in the slurry; ammonia formation causes a pH change and a subsequent zeta-potential
change which leads to electrostatic destabilization, particles dissolution will change the ionic strength in the slurry and it can also lead to surfactant poisoning; particles precipitation causes an increase in the specific surface area, whereas precipitated gel is expected to form a stiff network by binding the host particles. The HAS process is applicable for any ceramics containing alumina, at least as a minor phase (alumina, ZTA, Si3N4, SiAIONs...); large and complex-shaped parts with a homogeneous micro- and macrostructure can be produced. The process works in acidic as well as in alkaline regions; the removal of the liquid medium (water) is easy - by drying and no organic binders are needed. In the present paper we report on the use of HAS in slurry-forming Si3N4bodies. Alumina, which is formed during thermal decomposition of aluminum hydroxide, in combination with added Y r 0 3 ,later on serves as a sintering additive to promote Si3N4densification. Two different Si3N4 powders were used in this study. The powders are produced using different production routes and therefore differ in surface composition and oxygen content as well as in price and quality. These two powders were chosen to determine whether the HAS process is dependent on the powder characteristics as well as to see if it can be used to produce highperformance ceramic parts.
EXPERIMENTAL PROCEDURE The Si3N4 powders used in the experimental work were: E-10 powder (UBE, Japan) with a nominal particle size of 0.5 pm and an oxygen content of 1.34 wt.% and SILZOT H Q (SKW, Germany) with a nominal particle size of 1,3 pm and an oxygen content of 0.7 wt.%. The setting agent was A1N Grade B powder (H.C. Starck, Germany) with a nominal particle size of 1.2 ym, oxygen content of 2.2 wt.% and a specific surface area of 3.2 m2/g. For the sintering additive Y203(99.99%, Ventron, Alfa products) powder was used. The hydrolysis tests were performed to determine the reactivity of AIN in Si3N4 slurries. Supernatants were obtained by centrifuging 10 wt.% solids containing aqueous slurries of both (UBE and SKW) powders. The reactivity of AIN in these supernatants was determined by monitoring the pH versus time at room temperature. The viscosity of the slurries prepared with the as received and leached Si3N4powders was measured after AIN addition to confirm its reactivity in the ,Si3N4 slurries. 365
A nominal starting composition of 85 wt.% Si3N4, 8.6 wt.%Y,O, and 6.4 wt.% A1203 was chosen for the preparation of sintered Si3N4 ceramics. Since it was intended to obtain the A1203 from the thermally decomposed Al(OH)3 reaction product formed during the HAS process, the starting powder mixture in the slurries consisted of 86 wt.% Si3N4, 8.7 wt.% of Y2O3 and 5.3 wt.% AlN. To avoid premature hydrolysis of the admixed AIN powder, this powder was added to the slurries after milling of the Si3N4 powder and its homogenisation . powders were leached to remove silica with Y 2 0 3Si3N4 from the Si3N4powder surface. Afterwards, the slurries were filtered and the cakes were washed, dried and redispersed in deionised water. A commercial polyelectrolyte, PC 33 (Zschimmer & Schwartz, Germany) was added as a dispersant with a concentration of 0.5 wt.% of the total solids weight. The pH of the suspensions was adjusted to a value of 10.5 by the addition of ammonia. Y 2 0 3powder was added and the suspensions were well homogenised for 1 hour by milling in a planetary mill. AIN powder was then incrementally added and the slurries were milled for another 15 minutes, followed by de-aring. The total solids loading in the slurries (after AIN addition) was 50 vol.% for the UBE E l 0 powder and 53 vol.% for the SKW Silzot HQ powder. Once homogenized, the slurry was poured into the moulds (small round containers with = 30 mm and h = 15 mm) for the solidification tests at four different temperatures (room temperature, 50, 60 and 70 "C). The time needed for solidification was estimated. After being solidified, the specimens were demoulded and dried in air at 80°C overnight and their fractional density was measured. Samples with the highest green densities were sintered in a graphite crucible at 1780°C for 2 hours in flowing nitrogen gas. The green densites of the HAS-formed parts were measured geometrically. Distilled water was used as an immersion liquid for the sintered-density measurements. Microstructural observations by SEM were performed on the fracture surfaces of sintered samples.
water
+SKW Silzot HQ +silicic acid
5 ! 0
I
2o
I
time(h)
40
60
Figure 1. Variation in pH with time of 2wt.% AIN suspension in different media (water, UBE El0 supernatant, SKW Silzot HQ supernatant and solution of silicic acid)
+
RESULTS AND DISCUSSION Our preliminary experiments using the HAS method in the slurry forming of Si3N4 bodies failed, because there was no reaction between the AIN and water and consequently the suspension did not solidify in the mould. Therefore, the hydrolysis tests were conducted in supernatants obtained from the two Si3N4 powders to determine the reactivity of AIN powder. In Figure 1 . the results of the hydrolysis tests are presented. Since ammonia is formed during the reaction of AIN powder with water, measuring the pH can give information about the powder's reactivity. The dashed line in Figure 1. indicates the change of pH during the reaction of AIN with water. The pH of the suspension of AIN powder in the supernatants of both Si3N4 powders did not change over a period of 56 hours after which the tests were stopped. The reason for such behaviour was ascribed to
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leaching method time, h 1 2 3 4 5
mixing in boiling water
milling in water at R.T.
mixing in hot alkaline solution
concentration of Si, mg/l 27 106 141 62 197 161 103 200 158 105 155 1 I5 1 I4
Table 1. Concentration of Si (mg/l) versus time in SKW Si3N4powder-suspension supernatants after leaching using different methods
90,oo
0,70
+T
("C) UBE E l 0 A.R. UBE E l 0 leached -E+-SKW Silzot HQ A.R.
+ 80,OO
-+-
0,60
70,OO
0,50
60,OO
0,40
h
9
W
5>
W
0,30
50,OO
0 u)
40,OO
0,20
30,OO
0,lO
20,oo 0,oo
20,oo
40,OO
80,OO
60,OO
100,OO
0,oo 140,OO
120,OO
t (min) Figure 2. Variation in viscosity with time at constant heating rate for concentrated (40 ~01%)aqueous Si3N4 suspensions prepared from different powders containing 5 wt.% of AIN To confirm the reactivity of AIN powder in the Si3N4 slurry prepared with leached powders, the viscosity measurements during constant heating of the highly loaded (40 ~ 0 1 % )aqueous Si3N4 suspensions, also containing 5 wt.% of AIN powder, were performed. The results are presented in Figure 2. together with viscosity measurements of the suspensions prepared from the as received powders. In the latter case, the viscosity did not change with rising temperature indicating that AIN did not react with the water. In contrast, when the leached powders were used, at approximately 50°C (i.e. after 40 minutes) the viscosity started to rise. With these experiments, the efficiency of leaching was confirmed and the information regarding the temperature at which the HAS process should be conducted was obtained. Furthermore, the data also show that leaching of the powder did not have any side effect on the viscosity of both Si3N4 (UBE and SKW) slurries This is in agreement with a study performed by Laarz et al., who reported on even better rheological properties of the leached Si3N4powder'. The results of the experiments, which were conducted in order to determine the most suitable solidification conditions, are presented in Table 2. The times needed for solidification at four different
RT UBE E l 0 SKW Silzot HQ
temperatures are listed together with the corresponding green densities of the solidified samples after drying. The solids loading of the slurries were 50 vol.% for the UBE E l 0 powder and 53 vol.% for the SKW Silzot HQ powder. The higher green densities were in both cases achieved by setting at room temperature, where the setting time is very long. With increasing setting temperature the solidification time will be shorter, but at the expense of a lower green density. At 70°C there are already large spherical pores (bubbles) visible due to ammonia evaporation. It is known that the solubility of ammonia in water rapidly decreases with temperature and the amount of ammonia formed is already above its solubility at 70°C. Based on these results 50°C was chosen as the setting temperature at which the samples need 2.5 hours to solidify i.e. these conditions represent a good compromise between green density and the time needed for solidification. The relative densities of sintered samples, which were solidified under these chosen conditions were 97.0% for the UBE E l 0 powder and 96.9% for SKW Silzot HQ powder. Since no attention was paid to the composition and sintering conditions, two samples prepared using other wet-forming techniques were also sintered, for comparison. An injection-moulded sample
time
Par
1 day 1 day
51 % 57%
time 2.3 h 2.5 h
70
60
50 Par
50 Yo 56 Yo
time 2h 2.3 h
Par
49 Yo 54 Yo
time lh 1.3h
PC!
47 Yo 53 Yo
Table 2. Setting times and corresponding green densities (% of theoretical density) for highly loaded Si3N4slurries containing 6.78 wt.% of Y 2 0 3 and 1.81 wt.% AIN for disk with 9=35mm and h=15mm.
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prepared from UBE El0 reached 98.0% T.D. and the sintered density of a sample prepared from SKW Silzot HQ powder using conventional slip casting was 98.1 YO T.D. The slightly lower sintered densities of HASformed ceramics is most probably due to the somewhat lower green density of the HAS formed bodies. It is worth mentioning that these values are considerably higher than our previous results obtained by solidification at a higher ( 8 O O C ) temperature6.
The SEM fractographs of sintered Si3N4ceramics at low and high magnification are presented in Figures 3 10. They show that the microstructures of HAS-formed ceramics do not differ from those obtained by other wetforming processes. The elongated P-Si3N4 grains are homogeneously distributed in all cases indicating that the presence of aluminium hydroxide in the green bodies has no influence on the densification and phase distribution. The results are similar for UBE E-10 and SKW Silzot H Q powders.
Figure 3. Fracture surface of UBE El0 formed by the low-pressure injection moulding process (low magnification)
Figure 4. Fracture surface of UBE E 10 formed by the low-pressure injection moulding process (high magnification)
Figure 5. Fracture surface of UBE El0 formed by the HAS process (low magnification)
Figure 6. Fracture surface of UBE E I0 formed by the HAS process (high magnification)
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Figure 7. Fracture surface of SKW Silzot HQ UBE E l 0 formed by slip casting (low magnification)
Figure 8. Fracture surface of SKW Silzot HQ UBE E l 0 formed by slip casting (high magnification)
Figure 9. Fracture surface of SKW Silzot HQ UBE E l 0 formed by HAS process (low magnification)
Figure 10. Fracture surface of SKW Silzot HQ UBE E l 0 formed by HAS process (high magnification)
CONCLUSIONS
REFERENCES
The hydrolysis assisted solidification (HAS) process can be successfully used in the slurry forming of Si3N4 ceramics, provided that the concentration of silicic acid in the slurry is sufficiently low to allow AIN hydrolysis. Silicic acid adsorbs onto the AIN-powder surface protecting it from hydrolysis. Removal of the silicic acid by leaching of Si3N4 powders in a hot alkaline solution proved to be a very effective method. The most suitable setting conditions were determined to be 2.5 hours at 50°C. The process can be used for the production of high-performance ceramics from high quality powder as well as for the production of lower performance Si3N4 ceramics from cheaper powder. In both cases the densities and microstructures did not differ much from those of conventionally prepared silicon-nitride ceramics using other wet-forming techniques. The presence of aluminum hydroxide does not seem to exert any negative effect on the properties and microstructure of liquid-phase-sintered Si3N4 ceramics and can thus be exploited as a sintering additive.
I T. KosmaE, S. Novak, M. Sajko, HydrolysisAssisted Solidification (HAS): A New Setting Concept for Ceramic Net-Shaping, J. Europ. Ceram. SOC.,17 (1 997) 427-432. P. Bowen, J.G. Highfield, A. Mocellin, T.A. Ring, Degradation of Aluminum Nitride Powder in an Aqueous Environment, J. Am. Ceram. SOC.,73 3 (1 990) 724-728. 3 W.M. Mobley, Colloidal properties, processing and characterization of aluminum nitride suspensions, PhD thesis, Alfred University, Alfred, New York, 1996. K. Kmel, T. KosmaE, Reactivity of AIN Powder in Dilute Inorganic Acids. J. Am. Ceram. Soc., 83 6 (2000), in print. 5 E. Laarz, G. Lenninger, L. Bergstrom, Aqueous Silicon Nitride Suspensions: Effect of Surface Treatment on the Rheological and Electrokinetic Properties, Key Engineering Materials, Vols. 132-136 (1 997) 285-288. K. Krnel, T. KosmaE, Use of hydrolysis assisted solidification (HAS) in slurry forming Si3Nj bodies., Ceramic Transactions, 83 ( 1 998) 257-264.
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Green Machining of Aluminum Oxide Ceramics Fritz Klocke; Dietmar Pahler; Christoph Schippers; Christian Schmidt Fraunhofer-Instituteof Production Technology IPT, D-52074 Aachen, Germany
Abstract The primary objective of the machining sequence prior to ceramic sintering is to produce a contour similar to the finished shape and a surface rim zone free from damage so that it is only functional faces which need to be refinished. Although this strategy has been in industrial use in many applications, where green machining constitutes an important step in the production of the majority of ceramic parts, there is, however, considerable development potential with regard to the economic process. The reason for this is mainly that the available results are predominantly based on empirical research within companies. Furthermore, the machining behaviour of ceramic green bodies is greatly dependent on the respective company specific and extremely complex material characteristics.
process and of the characteristic material properties on the machining process, the tool wear and the surface quality. The aim is to develop an optimum technological process design for economic machining in practical operation based on these fundamental interrelationships.
Choice of Material The task in hand made it necessary first of all to demonstrate significant interrelationships between the parameters of commercially available, thus representative granulates and the characteristics of the resulting structural arrangements. The main characteristics identified in this context were the chemical composition of the powder, its secondary grain size, the kind of added organic additives and the mould pressure during densification /10/,/11/,/12/.
Introduction In industrial practice, machining of green oxide ceramics with a defined cutting edge is an efficient shaping method as a part of the ceramics production process, which shows its economic advantages at small to medium piece numbers /1/,/2/,/3/. However, as green machining was mostly considered a subordinate step on the way towards the ceramic component, there was a considerable deficit in fundamental process knowledge in the past. Machining of green compacts was based on experience and empirical knowledge from the areas of commoditiy and refractory ceramics, which were transferred to the technical highperformance ceramics i7/,/13/. This fact and the emerging trend to regard green machining not only as a part of roughmachining and/or premaching in the production of functional surfaces, but also to produce fitting components under calculation of a reproducible amount of contraction /5/,/6/,/9/,indicated that an acute research need existed. The lack of process technological knowledge resulted in differing interpretations of the occurring process phenomena and in divergent and partly contradictory machining recommendations /8/,/14/,/15/,/16/.
Objectives and Methods It was the objective of the investigations to acquire from the analysis of the machining and wear mechnanisms a profound understanding of the process technological correlations of effects in green machining of oxide ceramics. The analysis of the machining and wear mechanisms is to lead to a profound understanding of the correlations of effect between influencing and outcome variables during green machining of oxide ceramics. As a result of these investigations it must be possible to interpret in particular the influences of the regulated conditions of the
Residual Moislure f, [%I 0.260.28 open porosity WMPa 32.7 150MPa 30.7 p. 151 210MPa 29.6
2.47 2.54 2.59
2.35 2.38 2.40
0.254.28
0.25-0.27
31.7 29.8 28.5
34.2 33.3 32.8
56/56/55
59/58/58 9W88 11O’Wunl d, 28Cb‘262/234405/364/3G+ I135/126/1 I =3,6l = 3.56 TrueDensity P, [ @ c M ] -3.60 = 3,70 = 3.90 Theo. Densuy P. IglcM1 3.70 = 2.5 -9.6 Closed porosity P, [a] = 2.8 ~aal Pwosityp, [a] 35.5-32.4 34.2-31.0 43,842.4 (fa90-210 MPa) 15/25/40 mw95 Ratio of FracNre ICI 45/65/80
Dimanofks (901150/210MF%)
d,,
84/82/79
4. 196/188/171253/232/214
-
Fig. I :
2.29 2.34 2.38 0.13-0, I4
53/51/50 96/93/91
-
3.80
o,260,2n
0.26-0.29
37.1 35.7 34.6 43/42/40 72/69/68 9519 1/89
31.6 29.9
3.63
-
= 3.90
= 3.90
= 2.6 41.541.0
44.542.0
7.4
72/258/238 743331288
-
3.60
-
= 3.88 7.8 39,636.6
98/Y8/100
Parameters of investigated green ceramics
It was possible to subdivide the examined structural arrangements into two divergent property profiles on the basis of two parameters, i.e. single granulate strength and porosity. From each of these a representative (2 and 5, see fig. 1) was chosen for the basic machining and wear tests in order to cover the spectrum of green ceramics. The mechanical behaviour of the green compacts can be described as predominantely linear-elastic with plastic components, the latter of which can be more or less neglected 141. However, under the increasing load speed, which is a factor relevant for the machining process, material embrittlement progresses, where the porosity of the structure proves to be a criterion of failure.
Analysis of the Machining Process From the phenomenological observation of chip formation and surface formation, two machining mechanisms could be
371
identified for green oxide materials, which were defined as fracture and shearing /lW. In the fracture-type separation process the elastic energy imparted to the material by the cutting edge is reduced in the shape of macrofissures when critical states of stresses are exceeded. The instable fissures run ahead of the cutting edge, always going deeper than the set cutting depths so that the created component surface is characterised by macroscopic break-outs. It was possible to describe the fracture progression in detail by describing the cutting edge engagement in the form of a mechanical fracturing substitution model as a superimposed mixed-mode-stress, after it had been possible to prove the creep brittle behaviour of green ceramics, by assigning it to a known brittle fracture criterion, and thus an adequate presentation of the regularitites of linear-elastic fractu mechanics.
The influcences of the relevant process and material parameters was discussed on the basis of a fracture mechanical analysis of both machining mechanisms. The fracture mechanical substitution model proved useful in the assessment, particularly with regard to the break-out depth, as did the produced surfaces and chips, subdivided according to break-out size and distribution. Thus, the proportion of material undergoing fracture-type machining, which is accompanied by a higher degree of surface damage, increases with rising machining depth and cutting speed (see fig.3) and if positive effective cutting angles are employed. The proportion of sheared-off material sinks proportionally. On the other hand, negative effective cutting angles and large cutting edge radii of curvature promote the shear-type cutting mechanism. The comparison of the different A1203-green compacts indicated that with increasing dispersivity of the material there was a tendency towards the fracture-type machining mechanism, for which reduced specific energies at break are primarily responsible. A reverse behaviour can be achieved by increasing the mould pressure during primary moulding. Furthermore it was proven that the knowledge gained can be applied to the group of zircon oxides.
Cuningspeed v. Imhniol
Fig. 3: Curring process in dependence on cutting speed Fig. 2: Chip formation during quasi-static penetrant testing and chip particles
Figure 2 presents some results of quasi-static penetration test, revealing the occurrence of the two mentioned mechanisms. Some resulting chips are also presented. The shear-type machining mechanism is characterised by slide faces caused by plastic material displacements and/or localised structural damage. The reasons for this were identified to be the volume effect, which contributes to an enlargement of the plastic deformation area in front of the cutting edge, and the fissure impediment as a consequence of compressive strains. In the machining operation this effect appears primarily at low effective engagement depths as a fissure causing stress concentration does not develop so well in small volumes. Apart from manifesting itself in a good surface quality the shear mechanism shows itself in small chip and minute dust particles.
372
The residual moisture of the green compacts has an important influence on the material behaviour. With increasing residual moisture the plastic deformation fractions rise so that the principles of linear-elastic fracture mechanics cease to be applicable at higher levels of residual moisture and the measurements become insensitive to the mechanisms of fracturing and shearing in the machining process. With regard to a reproducible design of the machining process this fact forces the user to store the compacts under defined conditions at low humidity.
Analysis of tool wear Besides ultra hard cutting materials, the relevant literature about machining of green compacts repeatedly recommends ceramics as well as hard metals and their coatings as suitable tool materials. Owing to the predominantly abrasive
/
wear contact it is the material hardness which must be closely scrutinised. The parameters of the composite materials and layers must always be classified in comparison to the non-directional a-AI2O3-ceramics crystallites, which have a hardness of 2500-3000DH and act upon the tool cutting edge in a loose and bonded state. Thus, a subdivision into different classes of material (according to the hardness of the cutting materials in comparison to the values of the ceramics crystallites) seems appropriate with regard to the analysis of the wear mechanisms. With increasing influence of surface ruin on the wear process, the resistivity of the cutting materials to alternating mechanical stresses, i.e. the material tenacity or flexural strength, generally gains importance. The cutting material comparison, which is more informative with respect to green machining is carried out using cylindrical surface traverse machining of the homogeneous material 5 as an example. Figure 4 gives the wear of cutting edges of wear land at two different cutting times (GI = 300 s and &2 = 600 s), the initial cost K and the economics factor W for each tool.
lml 0.4
0.3
0.2
0.1
Wcnr of Cutting Edge VB
0
0 0
25 10 COSS
SO 20
7S
30
of cutting edge of VB,, = 33 pm which increases to VBk2= 45 pm when the cutting length is doubled. The nonlinear increase of wear progression reveals the course which is, as the findings show, digressively shaped without exception. The comparison of the tools with regard to economic criteria is carried out on the basis of the costs arising per length of cut. The calculation is based on a tool life end of VB = 0.25 mm or the time of layer failure. Furthermore the fact is taken into account that all the indexable inserts have two usable cutting edges with the exception of the PCBN and PCD tools. In summary, there are two alternatives to the high hardness cutting materials: the K-group hard metals and, with limitations, mixed ceramics, which have an economics factor which is inferior by about the factor two (hard metals: K-group) and two and a half (ceramics) respectively in comparison to polycrystalline diamond. However, the extremely disadvantageous wear development of the cermets and the conventional coatings does not make economic green machining possible using these cutting materials. The moderate parameter of the crystalline diamond compound, too, can be put down to a failure of the CVD-layer even before the desired wear of cutting edge is reached.
IWlDMl [DMhml
K. cbsI-uTeaivencss w
Fig. 4: Comparison of cutting tool materials for turning operations of unsintered oxid ceramics
If the results are considered first of all from a technological point of view, it is not possible to determine at first sight a general dependence characterising the abrasion wear between cutting material hardness and wear development. Cermet materials and conventional coatings show extremely high wear progression, which is not comprehensible on the basis of the corresponding hardness. The coatings are worn or chipped off after only a short engagement time, so that there is direct contact between the green ceramics and the substrate material. Afterwards the wear increase is mainly characterised by the properties of the basic body. It remains to be clarified which are the mechanisms responsible for this premature tool failure. The connection between increasing hardness and decreasing wear, however, applies adequately to all the other cutting materials. This is why the diamond tools offer the highest possible resistance to the ceramics crystallites which strike and slide off. After a cutting length L,,= 500 m the PCD-cutting edge has a wear
Gliding on the Tool Flank
Free Floating on the Tool Surface
Fig. 5: Different impact zones of the cutting edge /18/
For the examination of tool wear it proved to be expedient to subdivide the cutting edge according to the mechanical stresses into the load zones abrasive blasting process (cutting edge and rear face region) and sliding process (flank and front face region) /18/ (see fig. 5 ) . Through the analysis of the tribosystems, adhesive wear proportions and tribochemical reactions could largely be excluded as a consequence of the low temperature level and the existing dry friction /18/. The wear behaviour of all the cutting materials is dominated by abrasion and surface ruin. Loss of cutting material on the tool flank is of priority importance in comparison to crater wear on the face due to the change of the contour producing dimensions, the surface impairment of the work piece and the determination of the tool life criterion.
373
Abrasion occurring in the form of micro-machining of crater wear on the face, which increases with rising cutting binder phase and hard material grains is the primary cause speed and cutting thickness value, is of subordinate interest of wear in the flank region of hard metal tools. It is for practical operation, as it neither influences tool life nor supported at the cutting edge by surface ruin as a secondary machining process. effect of binder phase wear. Thus it is broken out hard material particles which, besides the ceramics cristallites, M*M No. 5 (ISOMPa), Dried (2h. 50°C) cause so called natural wear of the slide face (see fig. 6). Sample Dimension b = 3 mm Cutting Conditions Cutting Tool rn HM €20,0
f = h Lc= MO m a. = 7*, 7, =o*
HM K10. A PCBN.
PKD D3-8, o CVD-coating
Carbii DiamCIec Approl2-3 pm
- 78% wc .14% (Ti-
- 8% c4%B
\
Fig. 7: Influence of process parameters on wear development Yg. 6: Abrasive wear on a tungston carbide insert
The wear of ultra-hard cutting materials can be put down to increased breakdown and breaking out of the hard material particles and the subsequent natural wear. If there is a binder phase there is also abrasive damage caused by particle grooves. The diamond crystals wear out due to thermochemical reactions at flash temperatures /18/. The wear of the faces is caused by breakdown under steep and through abrasive grooving under a shallow impact angle of the ceramics crystallites. The circumstances of impact change with the speed and the effective cutting angle. In the comparison of cutting materials the high strength materials, boron nitride and diamond, as well as, with limitations, WC/Co hard metals were demonstrated to be suitable for green machining. This means that at present the polycrystalline diamond represents the optimum cutting material from both the technological and economic point of view. In the case of complex tool geometries K-group hard metals and crystalline diamond layers should be reconsidered, as soon as it is technically possible to adhere these sufficiently and reproducibly to the substrate. Generally all the measures which promote the fracture-type machining mechanism cause a wear reduction on the tool flank, as the active slide face, slip plane and the rebound frequency of the crystallites on the cutting edge are reduced. This includes e.g. great cutting thickness values and high cutting speeds as well as positive effective cutting angles (see fig. 7). In addition to this, the use of coarse grained cutting materials is advantageous (see fig. 8), while in green oxide ceramics this structural property increases the active contact area and thus wear. With the formation of the wear the fractured surface typography of the workpiece is covered by roughnesses, which finally define the tool life end of the cutting edge in dependence on the required surface quality of the workpiece (see fig.9). The slight markedness of
374
The acquisition of an elementary understanding of the machining and wear mechanisms was followed by investigations into the effects of green machining on component quality. No damage penetrating into the structural depth could be detected in the samples infiltrated for this purpose. It is rather macroscopic fissures which develop at the base of a surface break-out and which predominantely progress parallel to the surface. Strength tests on unsintered samples do not allow conclusions to be drawn about the machining parameters, while it is proven that sintered components react to the adjustment of the process parameters in the green ceramics machining operation I1 81. Consequently, finish-machining in unsintered state is necessary also to increase strength. The surface is levelled according to the material specific sinter shrinkage.
@$
I
I
Maferinl SampleDimension Cuning Conditions Cutting Tool
Nr.2 + 5 (ISOMPa), Dried (2h, 50°C) b=3m f = h = 0.1 nun v. = 2 M mlmin PKD. a,= 7-, r, = oo
\
Fig. 8: Influence of PKD-grain size on tool wear
Grain sizc [pml
Material Sampk Dimension Cuning Tool Cutting Condition S p a Size
-
As the effect of tool wear on the workpiece surface is
No.2 (IS0MPa). Dried (2h.SOT) b = 3 nun HM KIO. a,= 7". r, = 0" f = h = 0.2 nun A = 25 mm*(for AJ. A = 0.27 mml (for R,)
of subordinate interest in the roughing process as long as it is possible to machine the damaged rim zone in the subseauent Dhase. it is not necessarv to set a maximum wear of cutting edge. If the ceramics can be freely chosen, materials with a low specific energy at break are preferable with regard to minimum tool wear. In case of the examined aluminium oxides this means a high A1203 content with a fine-disperse microstructure and a mould pressure which is as low as possible.
P
Wear of CuningEdge VB [mml
Wear of Cuning Edge VB lmml
9: Sugace formation of unsintered oxid ceramic depending on tool
wear
Recommendations regarding the machining of unsintered Qxide Ceramics
With the exception of the functional surfaces which require refinishing, the final contour of the component is achieved by finish machining. The primary aims of the fine finishing process design are to shorten the production time taking into consideration the given quality requirements and to sever the damaged rim zone. The basic process control strategies for a finishing process which is oriented towards quality and wear minimisation are in detail as follows /18/: 0
The recommendations for the process design of green ceramics manufacturing processes are based on the acquired knowledge. It is necessary to carry out a subdivision into the processing steps roughing and finishing for the divergent target dimensions of surface quality, production time and economic efficiency, in order to realise the optimum solution for each specific component through suitable process control strategies. Starting from the green ceramics compact, the roughing process is used to achieve an approximation of the component contour down to a uniform overmeasure for the following finishing procedure. The emphasis of process optimisation lies on achieving a maximum possible timecutting volume with minimum tool wear. The basic process control strategies for a roughing process which is oriented towards material removal and wear minimisation are in detail as follows /18/: High cutting thickness values and high cutting speeds cause an increasing time-cutting volume and decreasing tool wear. Flank wear, which is relevant to green ceramic machining, can be minimised by using positive effective cutting angles and increasing clearance angles. The cutting edge radius of curvature does not exert any measurable influence. The use of coarse grained PCD-cutting tools achieves maximum tool life together with the best economic efficiency. If the cinematic attack conditions necessitate geometries it is possible to back upon and*in future*upon diamond layers K-grOuPhard with improved substrate adherence.
Decreasing cutting thickness values and cutting speeds improve the surface quality of the workpiece, but simultaneously they increase production time and tool wear. Both parameters should thus be adjusted to values which are as small as necessary and as big as possible to satisfy the component requirements. Decreasing negative effective cutting angles improve the surface quality at the expense of tool wear. For optimum process design the rule is again: as small as necessary and as high as possible. Although an increasing clearance angle makes it possible to reduce wear it must on principle be kept small because of the arising cutting edge misalignment. The cutting edge radius of curvature must be set as large as possible, even if its influence is only of short duration due to wear progression. The same statements as to roughening are true for the choice of cutting material: the best choice are coarse grained diamond cutting edges which, apart from a tool life advantage, imply large cutting edge radii. Complex tool geometries on drilling tools or undercuts necessitate hard metal materials or diamond coatings.
0
In the finishing process the effect of flank wear on the surface quality must not exceed the surface roughness and damage caused by the cutting procedures. The presented investigations consequently indicate that flank wear should be a maximum of VB = 0.25 mm. Green ceramics geared to finishing have a high specific energy at break, which, however, negatively influences wear.
__ 'lhe considerations based on the preceding chapters indicate that there are potentials which can be developed to increase economic in green machining of ceramic components (see fig. Generally, cutting must be proportioned in such a way that the formation of component edges does not coincide with the tool emerging from the workpiece. Edge break-outs can also be minimised by fracture suppressing cutting edge
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I71 Lurson, D.: Green Machining. In: Engineered Materials Handbook, Vol. 4: Ceramics and Glasses, S. 181-185, ASM International, Materials Park, Ohio, 199 1. With the knowledge gained through this work, the user is 181 Mehlhose, J.; Schneeweg, M.: Wirtschaftliche und offered the necessary tools to configure the ceramic schadigungsarme Bohrbearbeitung technischer Keramik production process more efficiently through green im Griinzustand. In: IDR 30 (1996) 4, S. 225-232. machining. The use of technology oriented machining 191 Sentoku, E.; Tanaku, H.; Kawabata, H.: On the strategies minimises tool wear and improves the certainty of Machinability for Alumina-Green Compact Pieces the process and the quality of the ceramic components. Tool Shape and Roughness of Finished Surface. In: Journal of the Japan Society of Powder and Powder Metallurgy 44 (1997) 2, S. 185-189. I101Hornbogen, E.: Werkstoffe: Aufbau und Eigenschaften von Keramik-, Metall-, Polymerund Verbundwerkstoffen. 6. Auflage, Springer-Verlag, Berlin Heidelberg New York, 1994. I 111Dietz M. et al.: Charakterisierung und Bearbeitung keramischer Griinkorper. In: Effizienzsteigerung durch innovative Werkstofftechnik, VDI-Berichte, Band 115 1 (1995), S. 239-251, Werkstofftagung 1995 der VDIGesellschaft Werkstofftechnik, Aachen, 15.-16.3.1995. /12/ Claaflen, T.: Herstellung und Charakterisierung von A1203-Verbundkeramiken. Dissertation, TU HamburgHarburg, VDI-Verlag, Dusseldorf, Reihe 5, Nr. 309, 1993. I131Schneeweg, M.: Technologische Grundlagen fur die Gestaltung der spanenden Vor- und Endbearbeitung von Maschinenbaukeramik. Dissertation, TH Zwickau, Eigendruck, 1990. F,=3.ON F ,=2 7N F,- 21 N 1141Schulz, H.; Zhu, L : Teilgesinterte technische Keramik F,=3.l N F K-LSN F,=27N F,=LSN Fff.Po.79 FJF 1.17 FJF = 1.47 F f f = 1.62 hochgeschwindigkeitsfrhen. In: Werkstatt und Betrieb 129 (1996) 5, S. 370-372. I151Steiner, M.; Stingl, P.: Herstellung von Bauteilen fiir Fig. 10: Edge appearence at cutting edge exit point in dependence on Axialpumpen und Motoren aus reaktionsgebundenem tool angle Aluminiumoxid. Forschungsbericht, BMBF Nr. 03M2106C, CeramTec AG, Lauf, 1997. References 1161Thomas, K.: Weichbearbeitung von SiliziumnitridStrukturkeramik. In: Jahresbericht des Fraunhofer I11 Butler, N. D.; Dawson, D. J.; Wordsworth, R. A.: Institutes fiir keramische Technologien und Shaping Complex Ceramic Components by Green Sinterwerkstoffe, Dresden, 1995. Machining. In: British Ceramic Proceedings 45 (1990), 1171Klocke, F.; Gerent, 0.; Schippers, C.: Green S. 53-58. machining of advanced ceramics. In: Jahanmir, Said I21 Jaschinski, W.; Nagel, A.: Moglichkeiten und Grenzen e.a.: Machining of ceramics and composites. New York: der Formgebung keramischer Pulver. In: Keramische Marcel Decker Inc., 1999, ISBN 0-8247-0178-X, p. 1Zeitschrift 44 (1992) 10, S. 667-670. 10. I31 Klocke, F.; Gerent, 0.; Schippers, C.: Griinbearbeitung von Hochleistungskeramik als Alternative zur I 181Schippers, C.: Griinbearbeitung von Oxidkeramik mit definierter Schneide. Dissertation, RWTH Aachen, Anwendung von formgebenden Werkzeugen. Eigendruck, 1999. AbschluBbericht, AIF-Forschungsvorhaben 9613, Fraunhofer IPT, Aachen, 1995. 141 Klocke, F.; Gerent, 0.; Schippers, C.: Machining of Information Advanced Ceramics in the Green State. In: ceramic forum international 74 (1997) 6, S. 288-290. I51 Konig, W.; Verlemann, E.; Wagemann, A.: Bearbeitung For further information please contact und Charakterisierung von Bauteilen aus Fraunhofer IPT, Steinbachstr. 17, D-52074 Aachen Hochleistungskeramik. AbschluBbericht, BMFT Nr. 4 9 241 89040, Fax +49 241 8904198 Tel. 03M20422, Fraunhofer IPT, Aachen, 1993. Email
[email protected],Internet www.ipt.fhg.de I61 Konig, W.; Gerent, 0.;Schippers, C.: Griin- und WeiBbearbeitung technischer Keramik. AbschluBbericht, DFG-ForschungsvorhabenKO 3611981, Fraunhofer IPT, Aachen, 1996.
attack conditions, e.g. during turning the use of small rake angles which generate pressure.
P
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LASER-ASSISTED TURNING OF' SILICON-NITRIDE CERAMICS F. Klocke, S. Bausch*, T. Bergs Fraunhofer-Institutfur ProduktionstechnologieIPT, D-52074Aachen, Germany
ABSTRACT The production of high-precision parts made of silicon nitride ceramic such as roller bearing rings or valves, has required time and cost intensive finish-grinding operations up to now. This situation has resulted in a demand for more efficient machining techniques to be developed and made available to manufacturing industry. A prototype precision lathe with an integrated high power diode for laser-assisted turning has been developed at the Fraunhofer-IPT, which co-operated closely with industrial partners. Continuous heating of the workpiece via laser and the resultant localized loss of material strength, make it possible to machine ceramics using a geometrically defined cutting edge. Complex silicon nitride parts can be manufactured without cooling lubricant, considerably more flexibly and with surface qualities of up to R,=0.3 pm.
Intensive efforts focussing on both process and on machine technology, have resulted in the development of an innovative technique, laser-assisted turning, referred to in the following as >>LATa,which is on the threshold of industrial implementation. The higher flexibility and the considerable reduction in setting up time in comparison with conventional grinding operations, are only some of the points in favor of this technique. It also eliminates completely the need for cooling lubricants. The development of laser-assisted turning would not have been possible without intensive co-operation between the Fraunhofer-IPT, systems manufacturers and industrial users or without the support of the German Federal Ministry for Education and Research (BMBF). An insight into the principle underlying this technique will be given in the following as will the specific areas of application and technological limits in the manufacture of parts made of silicon-nitride ceramic.
INTRODUCTION Advanced ceramics are gaining steadily in significance due to increasingly exacting demands which technical products have to meet. The driving force in this development comes not only from the need to make existing components more efficient but also from the desire expressed by companies to access whole new areas of application. Hybrid or even ceramic roller bearings are now being manufactured, for example. These enhance the sliding and emergency running properties in high-frequency spindles and can withstand exposure to extremely aggressive media. Pilot applications are already in place in the engine building sector. Valves made of silicon nitride ceramic are significantly lighter in weight and boast considerably higher levels of wear resistance than steel valves. The comparatively high manufacturing cost of these materials is limiting the more extensive use of these materials. The costs are incurred firstly by the manufacture of the sintered parts on one hand and on the other, by the need for time-consuming and costly end-machining operations. Unfavorable overdimensions on sometimes very complex contours, usually require protracted grinding operations which are not expected to achieve any notable increase in the volume of material removal. From an ecological point of view, the use of large amounts of cooling lubricant is a further drawback of grinding. Since there have been hardly any alternative processes so far, manufacturers have simply had to put up with this situation [1,2,3,6].
LASER-ASSISTED TURNING OF SILICON-NITRIDE CERAMIC In principle, laser-assisted turning belongs to the group of hot machining techniques. The silicon-nitride ceramic is softened in the area directly ahead of the machining zone by heating the material via laser radiation (fig. 1).
fig. I: The temperature-dependent reduction in the strength of siliconnitride ceramic, is the basic requirement for effective application of hot machining techniques.
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This is attributable mainly to the high-temperature strength characteristics of the amorphous glass phase surrounding the rod-shaped silicon-nitride crystals. It looses its initial strength at any temperature in excess of approx. lo00 "C. The reduction in deformation resistance as the temperature of the glass phase increases, ultimately results in enhanced machining characteristics. From the point of view of wear and economic efficiency, it makes sense to machine siliconnitride ceramic using cutting tools with defined cutting edge geometry, only after this point has been reached [1,2,31.
industrial applications, had a glass phase proportion of approx. 10 96 and generally demonstrated very good machining characteristics. The specification of the machining operations and associated process parameters depends on the shape and over-dimension situation of the part [2].
Work carried out to identify suitable tools, has shown that the combination of high levels of hardness in conjunction with sufficient high-temperature strength is the principle requirement for tool life oriented machining. To date, these requirements have been met only by high performance cutting materials made of cubic boron nitride (CBN) and of polycrystalline diamond (PCD). The types concerned, are all commercially available. Laser-assisted turning is closely related to conventional turning in terms of the classical cutting parameters of cutting speed, feed and cutting depth. However, only additional variables, the so-called laser parameters, must be taken into account when the process is designed to reach a minimum level of heat. The level of laser power as well as the shape and position of the focal spot must be regulated so as to ensure that heat is generated selectively and in a controlled manner since this is the only means of ensuring a sufficiently high level of process reliability. The complex three-dimensional temperature profile which develops as a result of the relative motion between the workpiece surface and the focal spot, poses particular problems. The ultimate objective is to ensure that the thermal conditions which ensure good material machining characteristics are in place within the machining zone which is usually a few millimeters away from the focal spot. The cutting and laser parameters interact directly with one another. When, for example, the cutting speed is reduced at a constant focal spot temperature, the temperature level in the machining section drops, because the material is exposed to a longer period of cooling as a result of the lower speed at which it moves between the laser focal point and the cutting edge. Effects of this nature make it vital to ensure that the parameters are balanced precisely to suit one another (fig. 2). The material itself is an additional factor which influences process design. There are a number of material variants within the group of silicon nitride ceramics, whose thermo-mechanical characteristics can differ widely from one another. When the proportion of the glass phase in the material is too low, for example, the machining properties of the material may remain insufficiently favorable, despite the fact that considerable heat is generated. The material variants, which are widely used both in research work and in
378
fig. 2: The final outcome of the operation is determined by the precise co-ordination of all of the factors which influence the process.
As in other manufacturing processes, the ability to create the right balance and the knowledge of the ways in which all of the process variables behave together, form the basis of quality and cost-oriented process design. The use of tools on one hand and on the other hand, the level of part quality achieved, are the principal evaluation criteria. Tool wear is determined largely by the collective load to which the cutting edge is exposed. This load is produced firstly by the high machining temperature caused by both the heat applied and by the self-induced heat and secondly by the extremely abrasive effect of the siliconnitride crystals. Only very hard, heat-resistant materials such as cubic boron nitride (CBN) or polycrystalline diamond (PCD) can resist these load levels. Due to their better heat removal characteristics, the more roughly grained types of PCD have proved to be better than the finely-grained types in terms of wear. The percentage increase in tool life path when PCD is used rather than CBN, is measured in double figures. The reason for the better wear behavior of PCD is that the diamond is harder. As previously indicated by the investigations into cutting material, the capacity to remove heat from the cutting edge, is an important factor. An internal cooling system for the tool holder should, therefore be provided. This acts as a so-called heat sink, which promotes the removal of heat from the cutting edge. In extreme cases, failure to provide a tool cooling system, can result in the PCD cutting edge becoming unsoldered, causing process failure.
MACHINING DEMO-PARTS The results of extensive machining tests undertaken on rings made of normally sintered silicon-nitride ceramic, illustrate the influence exerted by various process parameters on tool wear. The linear focus produced by an innovative high power diode laser (HDL), was positioned at a distance of approx. 5 mm in front of the machining zone. The temperature of the focus can be kept constant using a pyrometric temperature measuring device and a special PID control. The cutting edge wear VB, of a roughly-grained PCD tool tip, was investigated at various focal temperatures, cutting rates, feeds and cutting depths. Whilst a significant decrease in wear with increasing temperature was observed within the parameter limits (fig. 3), the influence exerted by cutting speed was found to be of considerably less importance.
or, if it does occur, then only in isolated cases. This is attributable to the splinter-shaped chip formation associated with insufficient heat penetration into the material. Process stability decreases steadily. In contrast to this, an increase in cutting depth to 2 mm, has no adverse effects on chip formation. The high levels of material removal which can be achieved at this cutting depth, thus guarantee sufficient levels of process reliability [2]. This presents one of the major advantages of laserassisted turning operations conducted on silicon-nitride ceramic in comparison with the traditional grinding operations: The high cutting depths which can be achieved, permit even complex contours to be manufactured flexibly and rapidly in one cut. The part example shown in fig. 3 illustrates this on the basis of a cylindrical surface turning operation in which a complete contour was machined in one step. Cutting depths of 2 mm posed no process-related problems in the manufacture of either grooves or chamfers. Indeed, on conclusion of the machining operation, the maximum flank wear was only 8 0 p m with surface quality of R, = 0.3 pm, which equals typical grinding quality (fig. 4). Machining result
Process
fig. 3: On favorable process design, the laser-assisted turning enables the machining of ceramics with geometncally defined cutting edge using commercial available cutting tools
cuttingrpawl v.=Mmhnin feedf P 15pm
depth of cut
iz
0.3
...1.5 mn
temperatureT r 1200 Y
The severely limited heat penetration into the surface of the part despite high focal temperatures becomes particularly apparent when the feed rate is increased. At values in excess of f = 3 0 p m , cutting edge wear becomes serious and can result in cutting edge spalling in individual cases. The reason for this is the increasing loss of material strength in the direction of the feed movement and the associated high levels of mechanical load.
SURFACE QUALITY AND SUBSURFACE CHARACTERISTICS
Despite a similar pattern of wear progression in the case of the cutting speeds investigated, there are significant differences in chip formation. A low cutting speed of between 30 and 40 d m i n and a feed rate below 30 pm results in a chip shape which is familiar from metal machining operations, the so-called shearing chip, or continuous chip. The occurrence of this type of chip in laser-assisted turning operations is an indicator that there is a sufficiently high level of material melting and thus, that the machining process is stable. When the cutting parameters cutting speed and feed are increased, or when the temperature of the focal point is reduced, this characteristic chip form either does not occur at all
On closer examination of the levels of surface quality which can be achieved, it was found that, as in conventional machining operations, there is a strong correlation between the condition of the cutting edge and surface roughness. Given otherwise optimum cutting conditions, surface roughness values between Ra=0.25pm and R a = 0 . 3 p m can be achieved at a cutting depth of 0.5 mm, even after a cutting time in excess of 2.5 h (that corresponds to a turning path of over 5 km). The tool wear measured on conclusion of the machining operation, was less than 100 pm. When, in contrast, the cutting depth is increased considerably, the more serious cutting edge wear results in a
two cutting took
fig. 4: Given optimum process design, the laser-assisted turning technique offers considerable freedom in part design.
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deterioration in surface quality to values of approx. 0.6 pm. Since the process is always stable, premachining or rough-machining operations, in which surface quality is of much less importance, can be conducted under these conditions [2,3]. The extremely high localized heating of the part, presents the question as to the possibility of subsurface part damage. To date, none of the investigations conducted have confirmed any reduction in strength in sample parts manufactured in laser-assisted operations. In fact, the characteristic fracture strength was found to be higher than that of the samples which had been ground whilst the probability of fracture was virtually identical (Weibull Modulus). The tribological characteristics too, of rolling contact for example, proved more favorable than in the case of the demoparts machined in a grinding operations. This applied particularly to rolling stress in dry machining operations. However, the part heating factor assumes additional importance with regard to thermal expansion and the resultant form errors. Although these have not yet be measured in concrete terms, it can be assumed from the low thermal expansion coefficients of siliconnitride ceramic, that these form errors lie within acceptable boundaries [3]. On the basis of the technological information acquired so far, the following conclusions can be drawn in relation to the efficiency of the LAT technique: Despite laser penetration, parts made of silicon-nitride ceramic can be manufactured without exerting any adverse influence on the mechanical characteristics. The surface quality which can be achieved is comparable to that achieved in grinding operations, and in some cases, actually exceeds grinding quality. Greater cutting depths permit the flexible manufacture of a wide range of contours at high material removal rates. The use of
standard PCD tool tips is conducive to long tool life times and, due to the defined cutting edge geometry, CNC-controlled cutting operations. This operation does not require the use of cooling lubricants. However, it is important to ensure that the chips produced, can be blown out of the beam path of the laser [ 6 ] . At the same time as the technological development and optimization, the development of a machine with all process-specific components which would be suitable for industrial application, was promoted within the framework of the BMBF project in order to ensure that all of these advantages could be exploited to the full by industry (fig. 4).
DEVELOPMENT OF MACHINE PROTOTYPE The resultant prototype consists of a precision lathe with an integrated 1.2 kW high-power diode laser whose low weight, small construction space required and the considerably higher efficiency proved to be particularly suitable for this application. The limited working area in the machine makes it essential for the laser, including the movement kinematics, to have a compact structure. The tool and the laser can be moved flexibly in relation to one another with a total of four linear axes and two rotational axes. This is absolutely vital to any attempt to implement many of the turning operations such as transverse turning, chamfering, etc. NC programming is still conducted manually. It is anticipated that it will be backed up in the near future by a suitable CAM system. The most important additional components which will require to be integrated in the machine when this technique is applied, are described in the following:
emlsJrorrpse~
- safetycsngineering - compactco~ction ChateCte-ofmL focaf distance c a 80 mm
-
-
high
max. linear focal spot 1.5 x 4 md integrated pyrometer '
fig. 4: New compact-construction lathe permits flexible machining operations to be conducted on a wide range of parts
;80
-
-
-
The use of specialized materials and special cooling systems for the tool and the clamping chuck, The integration of a pyrometer, a PID controller to regulate the temperature of the focal point, Direct, automated detection of the laser beam focus position (comparable with tool length measurement), Implementation of force and structure-borne noise for additional process monitoring, Coupled emergency stop circuits, Protection against laser radiation, Various extraction facilities and Implementation of in-process monitoring modules with special sensor-assisted termination criteria and safety routines to guarantee process reliability [5,6].
It has already been shown, within the preliminary test phase, that with the exception of the need for a few, relatively minor improvements, the machine is suitable for deployment in a manufacturing environment. The outcomes of the technology and machine tool development are, a highly efficient machine prototype and the availability of extensive process knowledge about machining of silicon-nitride ceramic in laserassisted turning operations.
ACKNOWLEDGEMENTS
[41 Konig, W.; Zaboklicki, A.: Laserunterstutzte Drehbearbeitung von Siliziumnitridkeramik, VDI-Z, 1993
[51 Weck, M.; Kasperowski, S.: ,,Wie Butter in der Sonne", Einsatz von Hochleistungsdiodenlasern in Werkzeugmaschinen, fertigung, 1998 [71 Bergs, T.; Kasperowski, S: HeiBes Licht fur harte Keramik, Tools, Informationen der Aachener Produktionstechniker, Nr. 4, 1998
INFORMATION For further information please contact: Fraunhofer-Institut fur Produktionstechnologie IPT, SteinbachstraBe 17 52074 Aachen Phone: ++49 / ( O ) 2 411 89 04 - 0 Fax: ++49 / ( O ) 2 411 89 04 - 1 98 http://www.ipt.fhg.de
Financial support of this work by the German Federal Ministry for Education and Research (BMBF) under 02PV72077 is gratefully acknowledged.
REFERENCES [I1 Zaboklicki, A.: Laserunterstutztes Drehen dichtgesinterter Siliziumnitrid-Keramik, Berichte aus der Produktionstechnik. Bd. 16, 98 Shaker Verlag GmbH, 1998 ISBN 1-8265-3934-6 [21 Klocke, F.; Bergs, T.: Laserunterstutztes Drehen von Bauteilen aus Siliziumnitrid-Keramik; Werkstoffwoche'98, Munchen, 1998 [31 Klocke, F.; Konig, W.; Zaboklicki, A.: Einfluss der laserunterstutzten Drehbearbeitung auf die Eigenschaften der Bauteile aus Siliziumnitridkeramik, Werkstoffwoche'96, Stuttgart, 1996
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CERAMIC ENGINEERING WITH PREFORMS FOR LOCALLY REINFORCED LIGHT METAL COMPONENTS Ilka Lenke", Gert Richter and Dirk Rogowski CeramTec AG, D-73207 Plochingen, Germany Abstract
Metal-Matrix Composites
Light metal alloys (e.g. aluminium, magnesium and titanium) show a fast increasing significance in mass production especially for automotive applications (I, 2). To use the advantages of weight reduction and to eliminate mechanical, chemical or thermal disadvantages reinforced light metal components and their manufacturing are of rising interest (3,4).
The main motivations for using low-cost light MMC materials in the automotive industry are 0 the reduction in mass,especially in engine parts 0 improved fiiction and wear properties 0 improved mechanical and thermal properties (e.g. strength, stifkess, coefficient of thermal expansion, ...).
A very promising way to reinforce light metal components is the local modification realised by infiltration of highly porous ceramic or metallic particulate preforms. Those kinds of preforms can be easily designed to meet the specific requirements of the application's functional zone and offer the possibility for cost effective manufacturing.
Therefore, the major goal in developing MMCs is.to achieve improved material properties while retaining low weight using a low-cost material. (6).
Light metal alloys for automotive applications The amount of light metal components for automotive applications, especially made fiom aluminium, has significantly increased during the last years and this trend is expected to continue (Tab.1) (A). The motivation for using light metals is the wish to reduce overall weight and thus to lower petrol consumption and toxic/hydrocarbon emission. Moreover, it is desired to reach a high recycling quote. ~
1993 1998 2008 I2008
I 50 kg A1 in average per car
I 70 kg A1 in average Der per car I 120 kg A1 ? in average per car
I
Tab. 1: Increase of aluminium applications (1). However, new applications or cost effective manufacturing may be limited by critical properties of light metal alloys, such as low wear resistance, low strength, Young's modulus, insufficient fiction behaviour or high thermal expansion. In many cases these properties are only critical in partial areas of the component. To specifically improve these areas, a highly porous ceramic or metallic preform can be infiltrated and surrounded with light metal alloys ( 5 ) to form a so called metal-matrix composite (MMC). This way, only the critical area is modified.
Examples for applications are cylinder liners in the Porsche Boxster engine block ( 5 , 7, 8) or brake rotors (7, 9). Many fiuther ideas for MMC-components do exist (3, 5 , 6, 8, 10). However, if conventional materials are replaced by MMCs, the construction of the components often has to be changed and adapted to the new material's properties.
Types of MMC There are different types of Metal-Matrix Composites. They can be categorised as follows (7): 0 Particulate reinforced MMC (PRM) Dispersoid reinforced MMC Cermets 0 Short fibre reinforced MMC (SFRM) 0 Whisker reinforced MMC (WRM) 0 Continuous fibre reinforced MMC (CFRM) 0 Monofilament reinforced MMC (MFRM)
In the case of particulate reinforcement, the MMC components can be produced by casting an alloy which contains the particles or by infiltrated porous preforms as described later on. Advantages of the infiltration process are lower manufacturing costs and a more homogenous arrangement of particles in the metal matrix (5). Further on, the fabrication costs of particulate reinforced MMCs are lower than the cost for fibre reinforced MMCs because the fibres are usually more expensive than the powders ( 5 ) . Besides, the fibre reinforced MMCs often show anisotropic and worse recycling properties (5).
383
An example for a MMC application in the automotive industry One of the first practical application was realised with the LOKASIL-composite technique for cylinder liners in the aluminium engine block of the PORSCHE Boxster (8) as illustrated in Figure 1. The challenge was to improve the tribological properties.
Fig. 1: Pouros preforms for cylinder liners in the aluminium engine block of the PORSCHE Boxster. This was realised by producing a ceramic cylinder which can be infiltrated with metal melt by pressure diecasting. For this CeramTec Germany developed particulate preforms with porosity > 70 % and designed the process for the fabrication of the product. During pressure diecasting the engine block, the highly porous preforms (Fig. 2) are infiltrated with the metallic melt and the partial modification is constituted in a one step cost effective technology. The metal-matrix composite is formed locally with a perfect conjunction to the rest of the metallic component (Fig. 3).
Fig. 2: Highly porous preforms for cylinder liners.
3 84
Fig. 3.: Metal-Matrix Composite with a perfect conjunction to the rest of the metallic component (AlSi9Cu3, (226)).
Many design possibilities with porous ceramic or metallic preforms Not only tribological properties of light metal alloys can be improved by the ceramic preform reinforcement. Other critical properties like hardness, wear resistance, mechanical strength, creep resistance, thermal expansion and Young’s-Modulus can be designed to meet the customer’s requirements by choosing the right preform material, (Fig. 4 and 5 ) using different combinations of raw materials (Fig. 6 and 7) varying the particle size distribution (Fig. 5, 6 and 7) changing the preform’s pore size producing different porosity rates form 20 to 75 %. Figures 4 to 7 present examples for different reinforced A1Si9Cu3-MMCs. The materials used to form the highly porous preforms are metallic silicium (Fig. 4, 6, 7) and alumina (Fig. 5,6,7). The figures also show that different particle sizes and combinations of different raw materials can be utilised to design the prefoms and thus the metal-matrix composite.
Fig. 4: AISi9Cu3-MMC with (metallic) silicium (average particle size 50 pm, volume content 25%).
Fig. 7.: AISi9Cu3-MMC with silicium (average particle size 50 pm) and alumina (average particle size 10 pm), volume content of particle: 25%.
In Figure 8, the influence of particle size and material combination on the tensile strength is illustrated. The used metal matrix was A1Si9Cu3. By reducing the particle size the strength is improved. A higher strength is also achieved when silicium (which are needed when a good fiction and wear behaviour is required) is substituted by alumina. The tensile strength then rises fi-om below 220 MPa to values over 270 MPa and is higher than the value for the metal matrix material.
Fig. 5 : AISi9Cu3-MMC with alumina (average particle size 15 pm, volume content 25%).
These examples give just an idea of how wide the range is for special MMC-designs by varying the reinforcement parameters. The best and most cost effective particulate reinforcement can be chosen for the essential material properties.
Matrix
Serle
35pm
10vm
3Svm
ISpm
Al,~-Korngr&Ee
Fig. 6.: AlSi9Cu3-MMC with silicium (average particle size 50 pm) and alumina (average particle size 35 pm), volume content of particles = 25%.
Fig. 8: Influence of alumina particle size (KomgrOBe) on the tensile strength (Zugfestigkeit). Matrix = A1Si9Cu3, Referenz: AISi9Cu3with 25% silicium.
385
effective manufacturing. Close co-operation with metal casting specialists and application engineers ensures the best MMC development results.
Technical production To realise technical components in mass production, the following topics have to be considered additionally: cost effective processes for manuhcturing porous ceramic or metallic preforms, system requirements for application, casting requirements, final machining of the components, recycling process.
0 0 0
Thus, ceramic engineering has to be done in close cooperation with the metal casting specialists and the automotive engineers. For the preform shaping there are different standard manufacturing methods such as gel casting, isostatic pressing, axial pressing and extruding. In general, the production costs decrease fiom the first to the last mentioned method. Depending on the preform requirements and the needed amount of pieces per annum, the most cost effective approach has to be taken. In addition, the dimensional and shape tolerances needed for the preforms clearly influence the cost. Figure 9 shows schematically how time and cost input develops with the different production stages. Mass processing, Bisuit Sintering Product b Shaping, -b Drying hardening T”
REFERENCES (1) F. Venier, Leichtbau stimuliert den Absatz.
ATUMTZ-Sonderausgabe: Werkstoffe im Automobilbau (1 998/1999) 54-56.
(2) D. Brungs, H. Fuchs, Leichtmetall im Automobilbau - Trends und zukUnftige Anwendungen. ATZ/MTZ-Sonderausgabe: Werkstoffe im Automobilbau (1 998/1999) 50-53. (3) M. K. Aghajanian, et al., High Reinforcement
Content Metal Matrix Composites for Automotive Applications. SAE Technical Paper Series 950263 (1 995). (4) V. M. Kevorkijan, Commercial Viability of
MMCs in the Automotive Industry, The American Ceramic Society Bulletin, September (1 999) 6769.
(5) A. Nagel, Keramische Innovationen im Fahrzeugbau. Keramische Zeitschrift, 52 (2000) 406-4 11.
1
-
(6) M. V. Kevorkijan, MMCs for Automotive
Applications. The American Ceramic Society Bulletin, December (1 998) 53-59. (7) http://mmc-assess.tuwein.ac.at/,June (2000) what
are mmc’s.
processing Time and Cost Input
A
(8) E. KiShler et al., Einsatz von Aluminium-MatixVerbundwerkstoffen im Verbrennungsmotor. In K. Friedrich, Verbundwerkstoffe und Werkstoffverbund,DGM, Frankfurt ( 1997) 15 1163. (9) K.-U. Blumenstock, Gluhendes Verlangen. mot, 7 (2000) 74-76.
Fig. 9: Ceramic Production and time and cost input.
Summary Ceramic Engineering with porous forms for locally reinforced light metal components offer a high potential for product specific material design and cost
386
(10)H. P. Degischer, Schmelzmetallurgische Herstellung von Metall-MatrixVerbundwerkstoffen. In K. U. Kainer, “Metallische Verbundwerkstoffe”DGM, Oberursel (1994) 139- 168.
LASER BEAM WELDING OF ALUh4INA- A NEW SUCCESSFUL TECHNOLOGY A.-M. Nagel, H. Exner Laserinstitut Mittelsachsen e.V. an der Hochschule Mittweida, University of Applied Sciences D- 09648 Mittweida
Abstract This paper describes a very successful new method of laser welding of ceramics. The two beam laser welding technology, an additive free procedure, allows to create joints, e.g. of alumina parts, which have a strength of 85% of the original material. It enables to join parts of various shapes in only a few minutes without hrnaces and in a natural atmosphere. The achieved results as well as the advantages of laser material processing like small welding seams, high flexibility, high productivity and a high degree of automation make this technology ideally suited for industrial application. Applications based on this technology are expected in several branches, for instance for welding tubes or sensor elements, for the protection of electronic components against high temperatures, abrasion a n d or chemical attacks.
Introduction Ceramics are materials produced by a special sintering process. Depending on the composition it leads to properties such as high temperature resistance, extremely high hardness, low electrical conductivity and high thermal insulation, high chemical resistance and a lower density, compared with metals, can be achieved. These excellent properties are the reason for applying technical ceramics in wide fields of electronics, automotive and chemical industries. Currently there is no technology which produces joints of satisfactory quality between ceramical parts, preserving the excellent properties of the material. Brazing and adhesive bonding reduce the thermal and chemical stability of the system. These disadvantages are based on an additional material (glue or solder) with completely different mechanical, chemical and thermal properties than those of ceramics. That means a critical weak point is generated at the joint. Furthermore, brazing is usually only possible after metallisation of the ceramics to improve their wettability. This process needs time and is very expensive.
A very good quality of the joint, for example, can be
achieved by diffusion welding. The joining mechanism is based on diffusion processes at high temperatures. That means diffusion welding needs a long processing time of about one hour. The preparation of the material is very expensive (high quality of the surface) and a high bearing pressure is necessary (Therefore it is not suitable for joining small parts.). In addition, diffusion welding and brazing both make a vacuum atmosphere necessary. Our presentation will demonstrate that the limits mentioned above may be overcome by a technique using two laser beams. For the fust time it is now possible to produce geometries previously not practicable. Further development of this technology will lead to an enormous expansion of the application of ceramics.
Experimental Procedures Alumina (a-A&) substrates of 96.0% purity and a medium grain size of 3pm were used in our experiments. The substrates were 30 mm long and 10 mm wide, their thickness varied between 0.7mm and 1.2 mm. Because of the very low thermal shock resistance of alumina, a short local energy input by the laser beam will lead to cracks in the material. Thus the material has to be heated to minimise the thermal shock effect of the welding laser beam. To o v e r m e the disadvantages of preheating in a furnace, we are employing a second laser beam (fig. 1). This 600 W C02- laser beam scans across the surface of the material at a speed of about lm/s. Because of the very high absorption coefficient of the wavelength of 10,6pm, the material gets heated in seconds. The surface temperature of the parts is measured continuously by a pyrometer. An emission value of 0.75 was used for measuring. The power of the preheating laser is automatically adjusted to maintain the desired temperature. When the necessary preheating temperature is achieved the 1.2 kW Nd: YAG laser beam welds the parts together. It penetrates about 0.8mm deep into the material. That means that for such a low thickness' the generation of a welding bath is nearly independent of the thermal conductivity.
387
fibre
*
I
C02- laser beam
w-
Nd: YAG- laser beam
--
ceramics
I
J
Fig. 1: Experimental setup used for laser welding of ceramics In order to optimise the quality of the welding seams and to enhance the strength of the joints, investigations were concentrated mainly on process, preheating and welding parameters. The surface as well as cross- sections of the welding seams were investigated by optical and scanning electron microscopy (SEM). The strength of the joint was determined by the 4-pointbending test.
ResuIts Preheating Three major problems had to be solved: 1. What minimum preheating temperature is necessary to achieve crackfieejoints? 2. What maximum heating and cooling rate are possible? 3. What maximum temperature gradients across the surface are possible? A locally homogenous preheating across the whole substrate surface was carried out. The welding took place after the final stationary temperature had been achieved, which we varied in steps of 100K. The joined materials were investigated for cracks. It was found that 100% crackfiee joined materials were generated at a preheating temperature of 1500°C. However, investigations of cross- sections of the welding seams showed a minimum of porosity at a preheating temperature of 1600°C. Lower preheating temperatures as well as higher ones increased porosity within the solidified welding seam. For a cost effective technology the processing time is of special importance. Therefore the time for preheating and cooling was minimised to a degree, which still guaranteed a crackfiee result. It was not possible to define a specific heating rate for achieving 1600°C. A heating rate varying fiom 2OWs to 30Ws can be used up
388
to a temperature of 1400°C. The temperature range
around 1500°C is critical as the highest stresses occur there. A hold time permits stress reduction by reorientation of the grains. After that the Nd: YAG- welding laser beam can generate crackfiee welded parts. A thermally influenced cooling of the welded specimens is recommended in the upper temperature range above 1500°C. Below that temperature, no cracks were produced by normal cooling down in the surrounding air. For larger pieces lateral homogenous preheating is unsatisfactory because of the long processing time and high energy demand. Therefore maximum temperature gradients in relation to the distance to the welding seam were determined in a one- dimensional direction across the surface. The variation of the energy input was realised by changing the scanning lines of the C02- laser beam per area with a maximum concentration at the welding seam. As a result a maximum temperature gradient of 70 x 10Wm is possible across the surface. At a distance of more than 20mm fiom the weld the temperature is less than 500°C. Therefore, such pieces can be clamped and moved by conventional methods. In a furnace special and expensive high-temperatureresistant materials would be necessary.
Welding The achievable quality of the joints depends, apart from the wavelength, on four factors in general: 1. the mode of the welding laser beam 2. the focus position of the welding laser beam 3. the power of the laser welding beam 4. the velocity of welding. A laser beam can be generated as a continuous beam or
as a pulsed beam. Both variants are used in laser material processing. The mode is very important for laser welding of ceramics. In pulsed welding, more power is needed for welding material of the same thickness as it is necessary to compensate for the breaks between the pulses. This affects the temperature distribution in the welding seam as well as the solidification of the welding bath directly. The results are shown in figures 2 and 3.
a) Welding seam generated by a pulsed laser beam
The solidified welding bath shows a hemispheric form of a homogenous solidification and distribution of porosity. 3. The focus is positioned within the material. This leads to a material densification in the middle of the seam and to very large pores at the borders. In our opinion the second variant should be favoured.
Fig. 2 SEM- view of the laser welded surface of the welding seam
Laser beam power and welding velocity are related to the path energy. The path energy describes the powertevelocity ratio. However, a constant path energy at varying power and velocity leads to different solidification structures. For instance, a twofold increase in laser power, and velocity (same path energy) mean that only half time is available for heat transfer processes, mainly for thermal conductivity. The importance of this circumstance is shown in figures 4 a), b) and c).
a) pm welding seam b) cw welding seam Fig. 3 Cross section of laser welded A12Q (butt welding) Comparable parameters for welding 0.8mm thick tiles lead to a different solidification structure. Pulse mode (pm) welded specimens show an inhomogenous solidification of crystals. In the middle of every pulse (here the highest temperature existed) the structure is coarsegrained and fiiable. The border area of every pulse is characterised by columnar crystals oriented to the middle of every pulse, which are caused by radial temperature gradients within the pulses. The continuous (cw) laser beam welded specimens show a more homogenous structure. A grain growth of about five times that of the original could be obtained. The structure is dense and approaches that of the original at the border of the seams. In addition, figure 3 shows the distribution of pores in a cross- section. Pores are mainly located at the sides of the welding bath. They arise 60m vaporisation of impurities, and/or 60m agglomeration of pores, existing in the material. Pulsed-welded specimens showed higher porosity than continuous welded seams, probably because of the very high temperatures in the middle of the pulses. The position of the focus (the point of the highest intensity of a laser beam) influences the solidification of the melt, too. Three cases have been investigated: 1. The focus is positioned above the material surface. This leads to a flat welding seam of a homogenous solidification. 2. The focus is positioned on the material surface.
a)
Power: Velocity: Path energy:
22 w 0,22 mm/s 100 J/mm
b)
Power: Velocity: Path energy:
50 W 0,5 mm/s 100 J/mm
389
The investigation of 20 specimens makes a statistical evaluation possible. For industrial application the relation between breaking probability and breaking is of special importance (fig. 5).
c)
Power: Velocity: Path energy:
100 w 1,O mm/s 100 J/mm
Fig.4 Surface morphology at constant path energy and varying ratio of power and velocity Figure 4 a) shows a typical surface for flat welding seams. There is no root at the underside. Low laser
power and a low welding velocity lead to a rapid heat transfer into the base material. A solidification starting from the bath borders leads to the formation of columnar crystals oriented to the centre of the bath. There impurities will be concentrated melting at lower temperatures. At the same time a contraction occurs at both solidification fionts and leads to hot cracks known fiom welding of metals. In figure 4 b) a well-balanced ratio between energy input and energy losses due to thermal conductivity allows solidification in a homogenous and nearly isotropic manner. The cross-section of these joints is comparable with those shown in fig. 3 b). The crystal growth is limited to the threefold value of the original. These joints are also gas-tight. In figure 4 c) the centre of the welding bath is characterised by big grains of up to l o o p , enclosed by high porosity. This results fiom high temperatures induced by the high welding velocity and the thereby minimised thermal transfer. Impurities have been mostly vaporised. The borders of the seam show columnar crystals up to a length of 200pm.
0
I
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Fig. 5
20
0 Referenzl
30
40
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Io,/Wa cw- welded mecimen I 183 pm- welded specimen I 72 unwelded specimen I 191
390
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I100
70
Breaking probability and breaking load
The steepness of the curves characterises the spread of the measured data. As a result the continuously welded specimens show a smaller spread of breaking load compared with pulsed-welded specimens as well as with unwelded specimens. Up to now this result is not really clear but it has been confirmed several times. It is assumed, that the welding seam acts as a favoured point of fracture.
Welded assemblies The following laser welded assemblies can be presented:
Fig. 6
Laser beam welded wave structure
Fig. 7
Laser beam welded hollow cubes
Strength The strength of the welded specimens was determined by the 4-point- bending method. Two pieces measuring 30x7x0.8mm3 were welded together at their fiont in a pulsed and a continuous welding mode. The resulting strength ( 0 ) compared with that of the original material is shown in the following table:
60
BmMng load’
Reference Nagel, A.-M.: LaserstrahlschweiRen von Aluminiumoxidkeramik, TU Ilmenau, Fak. f Maschinenbau, Diss. (Theses) 1999
Fig. 8
Laser beam welded tubes
Summary Extensive investigations of laser welding of ceramics, mostly alumina of a purity of 96%, by Nd: YAG laser beam were carried out for the first time. The welding was done employing a two beam laser method. The preheating of the material, necessary to achieve crackfiee joints, was done by a second, scanning C 0 2 laser beam.This technology is very well suited for the laser welding technology. Particularly when compared with the alternative method of preheating in a h a c e , the advantages are obvious: - high processing speed - high flexibility - temperature fields can be generated and varied very quickly - creation of temperature gradients saves energy - material can be clamped by conventional methods - direct observation of the process is possible. On the other hand, material thickness is limited to about 2 mm because of the low thermal conductivity of the material used. Furthermore it could be determined, that the wavelength of the Nd: YAG laser beam as well as the continuous mode are very well suited for a homogenous solidification structure within the welding seam, especially for thin materials. The reason is the absorption behaviour of alumina, the energy distribution typical for a laser beam transmitted by a fibre and the avoidance of a vapourphase. The high quality is confirmed by a bending strength of 85% of that of to the base material. The mentioned results promise properties of the joints at high temperatures and/or corrosive atmosphere near them of the base material. First experiments to join alumina with transparent alumina of a purity of more than 99,9% and with some metals were very promising, and should be fiuther investigated. Due to the technology-related limitation of the material thickness and the very small welding seams of about lmm, multisectoral applications for small parts are expected in a number of fields like the chemical industry, analytic, measuring, mechanical engineering, and micro systems, leading to products and processes with improved performance for the user.
39 1
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SRBSN MATERIAL DEVELOPMENT FOR AUTOMOTIVE APPLICATIONS Biljana Mikijelj*and John Mangels Ceradyne, Inc. Costa Mesa, California USA
ABSTRACT
successful in high volume manufacturing of finished high performance components.
The development of a gas pressure sintered reaction bonded silicon nitride (SRBSN) manufacturing process capable of producing near net shape blanks for automotive applications is discussed.
SINTERED REACTION BONDED SILICON NITRIDE (SRBSN)
The effects of grinding parameters on material strength, rolling contact fatigue life and friction were determined. Based on this, cost-effective grinding and finishing techniques required to finish components to automotive application tolerances were developed.
SRBSN was developed on an R&D scale at the Ford Motor CompanJ.’ in the late 1970’s, and was scaled up for production by Ceradyne, Inc. in the late 1980’s. Figure 1. compares the SRBSN to the conventional sintered silicon nitride (SSN) process.
Examples of high volume SRBSN automotive components are presented. Components operating at high Hertzian contact stresses for extended periods, where metal components fail, have been proven as successful automotive applications for SRBSN ceramics.
INTRODUCTION Silicon nitride has been identified since the 1970’s as a material that would find wide application for various engine components due to its high temperature mechanical properties, thermal shock resistance, tribological and wear properties. Katz’ summarizes a number of engine components that did go into high volume production including glo plugs, turbocharger rotors and cam roller followers. However, the number of production components is small when compared to the number of prototypes that have been successfully evaluated. Cost and reliability are the principal barriers to introduction of silicon nitride components in high volume automotive applications relative to components made from metal or other advanced materials. The principal factors affecting the cost of silicon nitride components are the raw powder and component finishing costs. Efforts to reduce cost often lead to compromises that adversely affect the component reliability. This paper will describe the development of sintered reaction bonded silicon nitride (SRBSN) that has successfully addressed the cost barriers and has been
*
SRBSN
I
Raw Materials (Silicon + Additives)
I SSN
I Raw Materials
I
(Silicon Nitride + Additives)
Spray Drying
Spray Drying
Fabrication
Fabrication
i I
Nitriding (3% + 2Nz 9 Si3N4)
Gas Pressure Sintering 10.3 MPa (1500 psi) Nz pressure
Sintering
The SRBSN process begins with the use of an inexpensive raw material - silicon powder - providing a significant cost advantage: silicon costs are far less than $lOnCg, compared to over $Sokg for Si3N4 The SRBSN process also exhibits a significantly lower shrinkage ( 10-12% versus 17-21%) and better dimension control than the SSN process. This results in a reduced grind stock, and consequently lower machining costs. Lower component shrinkage also allows more efficient
presenting author
393
use of high temperature furnace space relative to the SSN processing route, providing an additional cost advantage. The gas pressure sintering process results in a material having a density of >99.3 % of theoretical, along with a microstructure composed of interlocking needles (Figure 2). This structure is responsible for the excellent fracture toughness (5.5 - 7.0 MPa m In) exhibited by SREJSN.
Grade
Ceralloy@ 147-31N
Density
3.21 dcm3
Characteristic Strength
750-830 MPa
Weibull Modulus
15-25
Hardness W 5 )
1500 kglmmz
Fracture Toughness
5.5 - 6.5 MPa m
Thermal Conductivity
25 W/m K
a.
b. Figure 2. Optical microstructure of 147-31N a. polished, b. etched sample. Use of neural networks were employed to optimize numerous sintering parameters'; the result being a material with a robust sintering window (Figure 3). This wide process window results in high process yields. The result of this SRBSN development is a material that exhibits the properties summarized in Table 1.
Machining of SRBSN Machining studies, using design of experiment (DOE) techniques, have been conducted on Ceradyne SREJSN materials. Effects of machining direction, diamond wheel grit size, table speed, and material down-feed and in-feed per pass on the strength of rectangular ASTM C1161 size B MOR (Modulus Of Rupture) bars were initially investigated4. In 1997, evaluation of effects of grinding conditions on cylindrical MOR bar specimens were
394
Strength Figure 3. Neural Network Sintering Response Surfaces for Density and Strength6. material removal rates (decrease machining costs) and optimize the material performance.
It was found that strengths of bars machined longitudinally were not sigmihntly affected by grinding conditions.
Based on this data, and the component application, grinding conditions and sequences can be generated to optimize the component performance and minimize the machining costs.
Machining damage was significant for ASTM C1161 Size B MOR bars with a transverse machining direction. Diamond wheel grit size was the dominant factor for transversely ground bars, reducing the material strength by an average 56 Mpa when an 80 grit diamond wheel instead of a 120 grit was used (Figure 4). Figure 4 also shows a strong interaction between table speed and down feed, indicating that higher removal rates (table speed x down feed) actually improve the strength.
In cases where the component strength is critical, postmachining treatments have been developed which restore the inherent material strength (Figure 6).
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Figure 6. Post machining treatment effects on transversely ground 147-31NASTM C1161 (B) bars. A 320 grit wheel was used on all bars.
Figure 4. Design of Experiments Major Effects and Interactionson the SRBSN Strength of Transversely Ground ASTM C1161 MOR Bars. Wheel grit also affected the transverse ground strength of cylindrical bend specimens (Figure 5). In this figure, longitudinally ground specimens were compared to transversely ground ones (600 and 320 grit diamond)'. Fracture origins of the transverse ground bars confirmed that subsurface damage (10-80 pm sharp cracks) was the predominant failure mode'. I
Cylindrical YOR S p . c i u n i
a
147-311
Stress
Testtype
I
Rig test of production component
Cydes(test time)
Result
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2.45~10~
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Figure 5. Effects of diamond grit and orientation on 147-31N strength - cylindrical bend specimed.
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(lo00 hrs)
(ma)
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(loo hrs)
395
FRICTION COEFFICIENTS AGAINST STEEL Friction torque of Ceradyne Ceralloy@147-31N tappet inserts and two other sintered silicon nitride materials were measured against a steel cam-lobe'. Engine oil without a friction modifier was used as a lubricant. Tappets were initially diamond super-finished. Some parts were then finished by two techniques (chemomechanical-CMP and "Ford" proprietary finish). Friction results showed that, of the materials tested, only 147-31N Si3N4 offered significantly reduced friction a&st the cam lobe at all RPM values tested and for all finishes Figure 7a. Sintered Si3N4(SSN) grades showed a friction coefficient reduction for the "Ford" proprietary finishing technique only, which was applied as an additional manufacturing step after diamond super-finishing (Figure 7b). The differences in friction between SRBSN and SSN materials and steel cannot be explained by their surface finishes, and are not completely understood. It is believed that the material composition and microstructural differences contribute to the different tribological behavior. The use of Si3N4 tappet inserts instead of steel could offer 0.5% fuel economy benefits'.', in addition to a wear reduction in the valve train.
2
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Silicon nitride tappet inserts for piston applications. Test results have shown that the use of silicon nitride inserts results in reduced friction relative to the standard metal insert^"^ (Figure 7 and 8).
396
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I250
I000
"
1500
0' P ' z
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Ceradyne has participated in a number of ceramic automotive component qualification programs in the U.S. since the early ~ O ' S , with the various prototype components illustrated in Figure 8. These programs include: Silicon nitride exhaust valves and clevis pins for heavy-duty diesel applications. The culminated in a successful NATO durability qualification test for the engine".
"
a
a
Ceradyne has been producing silicon nitride ceramics for industrial applications since 1990 and has produced over half a million parts annually for wear applications from 1996 through 1999.
'
750
Camshaft RPM
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COMPONENTS
'
'
0.0'' 500
750
"
"
1000
1250
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b. Figure 7. Friction between steel cam lobe and steel or SiJN4tappet insert. a for 147-31N Si&; b. for a European grade of sinkred Si& redrawn from
-
'.
a.
Figure 8. Examples of prototype components made from Ceradyne Ceralloy@147-31N. Ceradyne has been successful in qualifying its silicon nitride for high volume automotive production applications. These include: Silicon nitride flat and radiused valve lifters for racing applications, Figure 9a. The customer is selling these parts commercially for high performance after-market applications in both drag and stock car racing. Silicon nitride cam rollers for heavy-duty diesel engines, Figure 9b. These components have been tested at contact stresses of 2,400 Mpa for time periods exceeding loo0 hours. Silicon nitride rolling elements for fuel pumps in light duty diesel engines, Figure 9b. These components have been tested at contact stress levels of 1100-1600 m a . These examples demonstrate that Ceradyne's Si3N4 material has the capability of routinely withstanding high Hertzian stresses. This is a result of a virtually pore free microstructure of its Si3N4, achieved by gas-pressure sintering (Figure 2a). This combined with Ceradyne silicon nitride's high fracture toughness, results in a reliable material for high performance applications.
b. Figure 9. Examples of Ceradyne production components manufactured from Ceralloy@147-31N a. Valve Lifters b. Cam Roller Followers
CONCLUSIONS New emissions requirements for heavy-duty trucks will require engines to operate at higher fuel injection and cylinder pressures". This, combined with the trend for longer warranty periods, often extending to 1 million miles, will open up new opportunities for ceramic components in the wear and tribology areas. Ceradyne SRBSN is a proven, cost-effective material for the above demanding engine operating conditions. Integration of cost-effective SRBSN manufacturing with high production volume machining capabilities, and experience in proven engine component applications, make Ceradyne a leader in the field of high performance automotive ceramics.
397
References R.N. Katz, “Applications of Silicon Nitride Based Ceramics”, Ind. Ceram. (Franza, Italy), 17[3] 158-64 (1997). J.A. Mangels, G.J. Tennenhouse, “Method of Densifying a Reaction Bonded Si3N4 Article”, U.S. Patent No. 4,285,895, 1981. J.A. Mangels, G.J. Tennenhouse, “Densification of reaction Bonded Silicon Nitride”, Am. Cer. SOC.Bull.. Vol. 59, N0.12 (1980). p.1216. MST Ceramic Machining Consortium, Appendix to the 10” Program Review Meeting, Gaithersburg, MA, April 10-11, 1997 G. Quinn; L. Ives; P. Koshy; “Cylindrical Rod Flexure Test Method and Results,” MST Ceramic Machining Consortium, 14* Program Review Meeting, Worthington, OH, April 15-16,1999 Neural Network design and training by NA Technologies, 1997. George Quinn, Lew Ives, “Cylindrical Bar Test Results on Ceradyne SRBSN,” 16’ Program Review Meeting NIST Ceramic Machining Consortium, , Costa Mesa, CA, April 18-20,2000. G.M. Crosbie, R.L. Allor, A. Gangopadhyay, D. McWatt, P.Willermet, “Surface Finish and Composition Dependence of Valvetrain Friction with Silicon Nitride Tappet Inserts”. Cer. Eng. & Sci. Proceedings; Vol. 20, Issue 4. p25; ACS 1999. A. Gandopadyay et al., “ Effects of Composition and Surface Finish of Silicon Nitride Tappet Inserts on valve train Friction,” Proc. 25” Leeds-Lyon Symp. On Tribology, Sept 1998, Lyon France. lo E. Schwartz et al. “NATO Qualification of Detroit Diesel #8V71-TA Engine at 530 BHP with Advanced Ceramic Components” S A E 2000-01-0524 report. I1 “Heavy Vehicle Propulsion Materials Workshop”, DOWORNL,Knoxville, TN, August 1999
398
PROCESS DESIGN FOR HIGH PERFORMANCE GRINDING OF ADVANCED CERAMICS IN MASS PRODUCTION L. Schafer*, K. Eichgriin*, T. Magg** *University of Kaiserslautern, Institute for Manufacturing Engineering and Production Management (FBK), D-67653 Kaiserslautern, Germany **Diamant-GesellschaftTesch GmbH, D-71610 Ludwigsburg, Germany The aim of the process design for grinding in mass production is to lead the process to its technological borders in order to maximize the output at sufficient workpiece quality and low wear of the tools. Fig. 1 shows the different fields of measures for the design of high performance grinding processes of advanced ceramics in serial production. Starting with the process chain of the blank, the process elements of the grinding process - machine system, tools and setting parameters - have to be designed. Besides intensive measures of quality assurance, in particular during the startup-phase, an in-process supervision for the analysis of instationary and dynamic process phenomenons are a further precondition for high efficieny and reproduceability. The example of high performance grinding of sililcon-nitride engine valves demonstrates how the process-chain-widecooperation of a ceramic supplier (CFI, Ceramics For Industry, Roedental), a tool manufacturer (Diamant-Gesellschaft Tesch GmbH, Ludwigsburg), and a component manufacturer (Mahle Motorventile GmbH, Bad Homburg) accompanied by a research institute (Institute for Manufacturing Engineering and Production Management, Kaiserslautern) lead to lastingly improvements of reproduceability, cycle times, tool life times and standards of quality.
ABSTRACT This article discusses the fields of measures for the design of high performance grinding processes for a competitive mass production of ceramic components using the example of Si,N, engine valves. The process design for these grinding processes has to take into account the ceramic-specific characteristics of the blank, the technological boundary conditions and the influencing factors on the workpiece quality over the entire process chain. By measures of optimisation in the fields of blank accuracy, statistical quality assurance, process supervision, tool development, adaption of the machine system and process design and by employing methods such as multidimensional statistical pattern recognition, process simulation and in-process data analysis, the realisation of reliable and cost-effective production of advanced ceramics is possible.
INTRODUCTION Designing a grinding process for mass production of components of advanced ceramics with high reproduceability, low subsurface damage and high cost-efficiency is a great technological challenge. Only the intensive, holistic consideration of the process chain and all elements of the grinding process leads to reproduceable results with optimum quality and costs.
INFLUENCE OF THE BLANK In serial manufacture of ceramics, the geometric deviations of the unmachined sintered compacts constitute a
1
Blank and process Grinding process
---
Qualityassurance
I
I
I
Machine system
"-I__
1
I -
Number
1
I
Process analysis
p'
1
Process design
I
syndronisation
/II1
Digitized NCSets
:
Output voltage of sensor vs. time
I
Fig. 1: Fields of measures for the design of grinding processes for mass production of ceramic components
399
major influence on the efficiency of the grinding process. These deviations from the ideal geometry are typical of ceramics, but lead to variations of the clamping situation of the machined parts in the clamping tools. Radial runouts of the workpiece surfaces in relation to the rotational axis are one of the consequences in cylindrical grinding. At the same time, large deviations from the ideal geometry require large overmeasures on the workpiece to ensure complete and accurate machining of all surfaces. Both together lead to high wear load of the tool due to dynamically varying depths of cut and consequently to increased operating and nonproductive time. In order to facilitate economic grinding, manufacturing of the blanks has to be carried out with high accuracy and reproduceability. Increased efforts in early steps of the process chain can, therefore, lead to significant cost reductions of the grinding process, thus reducing the total workpiece costs. A major problem of reproduceable ceramics manufacturing lies in the distortion of the workpiece after sintering. For the reduction of workpiece distortion all steps of the process chain, particularly powder prepation, moulding process and sintering, have to be analysed. A statistical analysis can contribute effectively to finding and removing the causes for workpiece distortion [l]. If the reproduceability of the blank geometry can be sufficiently assured, a near-net-shape forming is economically feasible. Apart from the influences on the economic efficiency, the blank has important effects on the work results of the
A
Model
grinding process. Examinations of the process chain of Si3N4 engine valves show that the geometry of the ground finished part is decisively influenced already by the moulding process. This is illustrated in Fig. 2 for the valve plate. The basic contour of the finished valve is determined by moulding and reproduces itself through all subsequent steps of the process chain. The grinding process reduces the amount of the geometric deviations to approximately one tenth, but does not impose significant changes on the original contour. The reason for this phenomenon lies in the high process forces in high performance grinding. These forces lead to dynamic deflections between the workpiece and the machine system due to the roundness deviations of the blank. Accordingly, the workpiece is not exactly circular after grinding, but maintains its original shape.
MACHINE SYSTEM The machine system for economic manufacturing of ceramics in serial production has to allow high material removal rates. This requires high cutting speeds (> 120 d s ) and high depths of cut. Due to the high process forces, high static and dynamic machine stiffness and high performance of the drive are required. A coolant system filtering the coolant down to remaining particle sizes of a few microns and incapsulated bearings and guideways represent another precondition. A special problem in serial production constitutes the handling of the blanks, in particular the automated feed and clamping into the clamping chucks. On the one hand, the blanks show variations of the accuracy and the run of the workpiece surfaces, hence there is danger of mistakes and tilting of the parts. As a consequence, the gripping and clamping devices require particularly large opening widths and bearing widths, as well as clearly defined fixture points. On the other hand, the high hardness of the workpieces leads to increased wear load on the grippers and clamping devices. For this reason, friction and relative motion between the workpieces and the clamping devices have to be minimized by adequate constructive solutions. Furthermore, the wear of the grippers and clamping devices must be supervised regularly. Fig. 3 shows an example of a clamping chuck specifically designed for the machining of ceramic engine valves with flexible, but defined contact areas, large opening widths and maximum possible bearing width.
Blank
Fig. 2: Reproduction of the workpiece geometry over the process chain 400
Fig. 3: Ceramic-specific clamping chuck for the machining of ceramic engine valves
TOOLS The grinding tools are a central element of an economic serial production of ceramic components. The tools influence the functional characteristics of the workpieces with respect to rim zone, surface and geometry functions. Furthermore, the tools essentially affect the reproduceability and efficiency of the grinding process. Due to this, the grinding tools must be developed continously in cooperation with an experienced tool supplier. Frequent changes of the tool supplier should be avoided, since supplier-specific grit and bond characteristics lead to discontinuities in the behaviour of the grinding process, thus increasing costs.
Fig. 4 shows the simulation results of the resonant frequencies of a tool for manufacturing of ceramic engine valves. The calculations carried out corresponded well to the measurements of the wheel, in the in the case of a fixed wheel as well as in the shown case of a free hanging. If harmonic oscillations of the rotational frequencies of workpiece or tool lie close to the resonant frequencies of the wheel itself, regenerative oscillations can be the consequence. This leads to increased wear of the tool and to the creation of waviness on the grinding wheel, resulting in increasing subsurface damage of the workpiece. Hence resonant frequencies of the tool have to be considered for the selection of the setting parameters.
Tools for large scale production need high wear resistance, consequently the bond requires high grit holding loads. Metallic bonding systems fulfil this requirement very well, but are only suitable, if the requirements on profile accuracy of the workpiece are not too high. In this case, a vitrified bond has to be selected. Increased grit concentrations are also favourable with respect to wear resistance, but lead to reduction of the chip space volume and hence require good cooling conditions and cleaning of the grinding wheel. Small grit sizes offer the advantage of the absolute flattening of the grits remaining small, thus the sharpness of the grinding wheel remains more constant and process stability improves. Small grit sizes and higher concentrations tend towards reducing the uncut chip thickness of the single grits. This is essential for the machining of the material with low subsurface damage. But these advantages can only be used, if the tool preparation is performed with high accuracy, since run, profile accuracy and topographie depend decisively on tool preparation. Small grit sizes require high accuracy of the run of the grinding wheel in axial and radial direction, and moreover the right amount of grit extend has to be guaranteed. Therefore, detailed guidelines for mounting, dressing and measuring of the tools need to be developed and their fulfilment needs to be supervised.
SETTING PARAMETERS The main criterion for the process layout and the selection of setting parameters is the machining with low damage of the rim zone. Since break-outs on the edges of the workpiece result from the unification of radial and lateral subsurface cracks, the size and number of breakouts give a good hint on the degree of subsurface damage of the material due to machining. For the machining of ceramics with low subsurface damage, a ductile regime cutting process must be achieved. This can be realised basically by low uncut chip thicknesses of the engaging grits in the contact zone. For this reason, high cutting speeds and low feed rates are favourable for machining, so deep grinding has to be preferred against pendulum grinding. A machine system of high performance and stiffness is a precondition for this. Due to the combination of ultrahard tools, hard workpiece materials and high machine stiffness, process dynamics in grinding of advanced ceramics are of particular importance to the work result and its reproduceability. Apart from tool wear and efficiency aspects, the dynamic behaviour of the grinding wheel has to be considered for the selection of process parameters. This is influenced fundamentally by the grinding wheel hub and by the type of bond. Simulations of the dynamic behaviour of the grinding wheel can contribute to avoiding critical revolution frequencies of workpiece and tool.
Fig. 4: FEM-Simulation of natural frequencies of ultrahard grinding wheels for manufacturing of advanced ceramics Furthermore, the dynamic behaviour of the workpiece has to be taken into account. Simulations of the natural frequencies of the engine valve showed that the oscillation of the valve strongly depends on the clamping situation. Fig. 5 shows an example. The upper part of the figure displays a simulation result. Resonant frequencies between 1207 and 1348 Hz were obtained for various fixture points. Due to geometric deviations of the blank, the clamping situation of the valve in the chuck varies widely. Measurements of the same valve that was inserted into the chuck three times showed three very different resonant frequencies (lower part of Fig. 5). Due to
Fig. 5: Simulation and measurement of the resonant frequencies of a blank under various clamping conditions 40 1
those varying and hard to model clamping situations, the results of the simulation can only show partial good correspondence to the measurements of a very lightweight acceleration sensor. The unfavourable interaction of natural frequencies of the valve and the tool can lead to the genaration of regenerative oscillations with the consequence of a self increasing waviness on the tool over the tool lifetime. As Fig. 6 demonstrates by the example of a grinding tool used in the early phases of the engine valve production, a wave pattern of 15 waves was created due to the regenerative oscillation of tool and blank. This problem could be solved easily by a decrease of the workpiece frequency.
Frequency of workpiece 1170 Ulmin Calculated number of waves: 3,l
Frequency of workpiece 4130 Ulmin Calculated number of waves: 14,l 14
1
12
11
Tool 1, new
m
a
7
Frequency of workpiece 3730 Ulmin Calculated number of waves: 30.9
Tool 1, end of tool life time
Fig. 6: Generation of wave pattern on a grinding wheel for the grinding of the valve plate due to interaction of blank and tool
A third criterion for the determination of admissable workpiece and tool frequencies in cylindrical grinding processes, besides the resonant frequencies of the interactive partners, results from the rolling kinematics of the process. Due to the interaction of workpiece and tool, a kinematic waviness is generated on the workpiece surfaces that can be determined from the frequency relation and the radiuses of the two interacting partners. This kinematic waviness can be calculated according to the approach of MERZ [2]. Fig. 7 shows an example of the calculation by using the valve seat. As the figure shows, the comparison of calculation and work result shows a very good correspondence. Considerable differences can occur, if the rotational frequencies of the grinding wheel or the workpiece differ from the selected wheel frequency of the machine control and thus the basis of the calculation is wrong. Hence the results of the calculation are only reliable, if there is an inprocess supervision of the exact rotational frequencies.
PROCESS SUPERVISION The sensor based inprocess supervision serves the detection and control of dynamic and instationary processes
402
Fig. 7: Calculated number of waves on the valve seat and obtained work result during machining. For the supervision of grinding processes, acoustic emission (AE) sensors are particularly suitable, since the acquired sensor signals contain significantly more information which is more detailed than alternative measuring methods (deflections, forces, engine current), even at low sampling rates (between 1 and 50 kHz). Fig. 8 gives an example by comparing the acoustic emission signal to the engine current signal of the machining process of the valve plate. The example shows the significantly higher dynamics of the AE-signal. The different machining steps can be clearly differentiated using the AE-signal, whereas the engine current shows only an unprecise, inert run of the signal, since the high inertia of the drive leads to a strong low-band filtering of the signal. Using the acoustic emission signal for the process supervision of ceramic engine valves, a large number of process characteristics can be extracted. This includes the initial contact detection, the exact duration of the machining and the nonproductive time, the revolution frequency of the workpiece and the tool, the relative overmeasure of the blank, the relative run out of the blank, dynamic effects during machining and the exact time overlaps of the machining processes of multiple acting spindles. Furthermore the development of the tool
Process supervision on the basis of a multi sensorical supervision concept with various measuring principles supplements the acoustic emission signals and acquires extensive information about the process.
QUALITY ASSURANCE I I emission I
Acoustic emission (hydrophonic sensor) ::External did:Seat : Cham$ :Plane :
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Fig. 8: Configuration of multi sensor process supervision of valve plate machining, comparision of acoustic emission signal to engine current profile wear can be detected under certain circumstances. Fig. 9 shows an example of the machining of the valve plate over the tool life time. Due to the increasing width of the wear mark over the tool life time, the maximum depth of tool engagement is not achieved discontinuousliy, but increases continuously over the wear mark width until full contact. This can be clearly seen by the longer rising time of the AE-signal. Thus the current width of the wear mark can be deducted from the AEsignal and the feed rate.
In comparision to machining of metallic materials, the serial production of advanced ceramics requires extended measures of quality assurance. Due to the higher wear rates, more frequent tool corrections are necessary. Therefore the density of quality checks needs to be increased. At the same time, clear guidelines for the workers concerning handling, operations scheduling, tool preparation, process setup and work-piece measurements must be developed and their fulfilment has to be ensured. Experiences show that a 100 % check of all workpieces is unnecessary if the process is satisfactory stable. For a lasting improvement of process reliability, the characteristics of the intermediate- and end-products must be acquired correspondingly and correlated to the blank characteristics, respectively. This requires the identification of the individual workpieces over all steps in the process chain e. g. by using a workpiece number. If the values of the measured characteristics are consequently entered into an observation matrix, the collected data can be evaluated by employing multivariate statistical pattern recognition methods, such as analysis of variance, factor analysis, analysis of discriminance or analysis of correlation [3]. These methods enable the quick identification of causes for quality variations over the process chain that can later be removed in optimisation cycles.
RESULTS Employing the measures of optimisation introduced, significant progress with respect to product quality and efficiency of the grinding procoess of the engine valves was achieved. At the preliminary end of the optimisation, a reduction of the cycle time of up to 74 % for specific valve types could be realised. Thus, the cycle times for ceramic engine valves exceed the cycle times of similar machined steel valves only in a range of few seconds (Fig. 10). Cycle time station I
Starting situation-
1
2nd optimisation Unworn tool
100 %
e time)
3rd optimisation
End of tool life
(Optimisetion of feedretes)
._-__ - I1856 Remaining potential:
(Futthe; feed optimisation, tool adeption)
-
0,O
0.5
1.0
1.5
Cutting time [s] Fig. 9: Wear of tool profile and course of the AE-signal over the tool life time
Fig. 10: Development of cycle times for machining of Si3N4engine valves In spite of the quicker machining, the combination of the different measures of optimisation lead to a significant reduction of tool wear. The G-ratio (ratio of cut chip volume and wear volume of the tool) of the tools was raised with a factor between 2 and 10, depending on the various working areas of the tools. The results are Gratios of 3000 up to > 8000 (Fig. I I).
403
5
Edge'2
7
s P
D Plane
a
U
Diarneterichamfer
- 6 . . 0
starting
I0
z 11
12 13 14 15
0
16
Fig. 12: Reduction of break-outs on the workpiece over the course of the projekt
17
0
2ooo4oO06o0O8oO01oooo
Gratio I mms/mma
Fig. 11: Development of grinding ratio with progressing optimisation Besides improvements of the economic efficiency of the grinding process, the subsurface damage of the workpieces due to machining was decisively reduced. The breakouts on the edges of the workpieces decreased in a way such that only few, very small break-outs remain on the valve's edges. The size of these break-outs lie significantly below the quality-critical size. Hence measures of quality assurance for break-outs can be dropped largely. Fig. 12 shows this by using the example of the edge 2 between the chamfer and the outside diameter. Due to a less harming machining, the flexural strength of the valves increased about 40 % compared to the orignal state, so now approximately 75 % of the theoretical material strength are achieved [4]. Aim of the valve production was to demonstrate the feasibility of manufacturing ceramic parts in high volumes with high reproduceability.Regiwding the manufacturing targets, this could be achieved. A fleet test in 1660 cars approved the reliability of the valves in the everyday application [4].
CONCLUSIONS To achieve high quality standards, high economic efficiency and high process reproduceability in mass series production of advanced ceramics, a holistic optimisation of the entire process chain is required. For this reason, the close cooperation of the ceramic supplier, the ceramics machiner and the tool supplier is essential for the holistic understanding of the process chain. From the point of view of the grinding process, measures to optimise the blank, the machine system and the setting parameters are necessary. These measures become feasible on the basis of an extensive statistical quality management, simulation results and inprocess sensor signals.
404
5 I0 15 20 25 30 35 Number of break-outs on edge no. 2
The combination of all measures have shown that it is possible to achieve a degree of reproduceability suitable for serial production of ceramic engine valves. At the same time efficiency and workpiece quality improved, ensuring functional reliability. The valves were tested succesfully in a fleet test on the road.
REFERENCES G Warnecke (Edt.), Zuverlassige Hochleistungskeramik. Report on the Projekt "Reliability and reproduceability within the process chain of ceramic components", University of Kaiserslautern 2000, ISBN 3-00-005686-6. R. Merz, Konzept zur Auswahl der Abrichtbedingungen bei der Einsatzvorbereitung konventioneller Schleifscheiben mit Diamantprofilrollen. FBK Produktionstechnische Berichte (Edt.: G. Warnecke), Vol. 15, Dissertation, University of Kaiserslautern, 1994. M. Kendal, Multivariate Analysis. Charles Griffin & Company Ltd, 2nd Edition, London 1980. K.-H. Thiemann, Erfahrungen mit keramischen Ventilen im Ottomotor aus Sicht der Werkstofftechnik. Presentation held at the Kolloquium "Zuverlassige Hochleistungskeramik, Rengsdorf, Germany, 5.-6. April 2000. G. Wotting, H. Lindner, E. Gugel, Experiences with High-Volume Production of Silicon Nitride Automotive Engine Valves. In: Proceedings of the 6th Interational Symposium on Ceramic Materials & Components For Engines, Arita, Japan, OCT 1922, 1997. G. Warnecke, L. Schafer, K. Eichgriin, C. Barth, W. Pfeiffer, Manufacturing of Ceramic Components. In: I. Marinescu, H.-K. Tonshoff, I. Inasaki (Edt.): Handbook of Ceramics Grinding and Polishing. William Andrew Publishing, Norwich, NY, 1999.
NEW CERAMIC EXCELLENCE FOR COMPLEX MACHINING OF ENGINE MATERIALS A. Krell, P. Blank, L.-M. Berger, V. Richter
Fraunhofer Institut f i r Keramische Technologien und Sinterwerkstoffe, Dresden, Germany
ABSTRACT New tool ceramics on the basis of A1203. exhibit an outstanding performance on coarse, intermittent, and final precision machining of engine materials like vermicular or chilled (hard) globular cast iron or hardened steel. These tools do not only provide a chance to substitute grinding by time-saving cutting operations, they can also be integrated into new machining concepts which minimize the frequency of tool exchanges reducing costs.
1. INTRODUCTION Machining of hardened steel and hard (chilled) cast iron by grinding suffers from the costs associated with the productivity of this process which is time consuming up to a factor of three compared with cutting. Cutring of heterogeneous materials like vermicular or globular cast iron (GGV, GGG), on the other hand, is either impossible with today's ceramics or requires expensive fools like polycrystalline diamond or cubic boron nitride (CBN). But even on machining mild steel or grey cast iron problems arise when commercial ceramic tools have to be exchanged for hard-metal inserts on each first cut at the rough edge of a turned shaft. It was, therefore, the objective of the present investigations to develop new tools on the basis of sintered corundum which are reliably applicable to the different engine materials.
2. EXPERIMENTAL PROCEDURE With the large body of evidence for an improved hardness and wear resistance by decreasing grain sizes, many manufacturers started to develop submicrometer composite tools on the basis of A1203 with Tic or Ti(C,N). Unfortunately, the covalent nature of the carbide bonds prevents pressureless sintering at temperatures of 1600 "C or less. Then, the most fine-grained of these new composites with 25-35 % Tic or Ti(C,N) exhibit average sizes of A1203 and carbide subregions of about 0.8-1 pm associated with a hardness HVlO up to 23 GPa (measured at a testing load of 10 kgf). On the other hand, much more fine-grained composites with a higher hardness can be produced when oxygen is introduced into the covalent phases of T i c or Ti(C,N) [associated with a different behavior of the raw powders on milling1.l Hence, the performance of these new composites (Tab. 1) was investigated in the present cutting experiments; their manufacture and characterization has been described previously.'
Advanced commercial ceramics were used as references. SH1 and SH20F are well known A1203/Ti(C,N) composites manufactured by CeramTec (Plochingen, Germany). Additionally, a "submicronstructured" composite introduced into the market with special emphasis to the machining of hardened steel was included into the study. Its major microstructural merit is the elimination of slightly agglomerated carbide structures which are not perfectly avoided in other grades. Two commercial CBN grades manufactured by the DeBeers company (South Africa) with nitride and carbide bonding (DBN45 and DBC50, respectively) were used as references for machining vermicular cast iron. It is not clear if carbide-reinforced tools will ever meet the thermodynamic and chemical demands for turning hardened steel. On the contrary, pure sintered corundum ceramics exhibit a hardness that after pressureless sintering is in no way inferior to hot-pressed and carbidereinforced composites. Their strength is about 800 MPa2 and equals or even exceeds the strength of the composites, and they offer the additional advantage of highest chemical and oxidation resistance. Therefore, pure alumina inserts with a submicrometer microstructure were prepared from a 99.99 % A120, powder by cold isostatic pressing and pressureless sintering as described repeatedly.2 Note that the material is macroscopically sensitive to creep: at T 2 11OO...1200 "C, bending experiments proceed up to a bending angle > 45 " without any indication of crack initiation (even at a high crosshead velocity of 1 m d m i n [with a span of 30 mm and a specimen width of 4 mml). Tab. 1 characterizes the investigated inserts, all data for laboratory and commercial grades were measured in our laboratory. It is common to measure the hardness on polished surfaces, and these data are given to enable a wide comparison. However, the surfaces of technical ceramic inserts are ground, and the hardness data of polished and ground alumina or A1203/Ti(C,0,N) surfaces differ. Therefore, the data for ground surfaces are more representative to illustrate the cutting behavior. The cutting performance was tested on a 35 kW CNC turning lathe (NILES, Chemnitz, Germany, 1990). In most of the tests with cutting inserts SNGN120412 (12.7-12.7.4.76 mm3, 1.2 mm radius, 20" chamfer with width 0.2 mm) the feed rate f was 0.1 mm/min, the depth of cut a was 0.2 mm. The plates were positioned with a rake angle y = 6", an inclination angle h -4", and an entering angle K = 45 ". The following parameters were measured (1)Flank wear width VB, the usually evaluated wear parameter at the primary cutting edge.
-
405
Table 1. Tool ceramics for machining vermicular cast iron, hard (chilled) globular cast iron, and hardened steel. Composition
Density
Grainsize*
absolute relative
~ c m 3 1[ % I
Vickers hardness (testing load 10 kgf)
[ GPa 1
[pmi (*composites: average of all phases)
[ GPa 1
surface preparation: polished ground
Pure alumina AC41, AC56
A1203
3.960 (99.3%)
0.56
20.2 f
+ 33 vol-% Tico 73Oo,14 A1203 + 33 vol-% TiCo,4200,,No,35
4.325 (100 %)
0.70
20.8 i 0.2
22.8 f 0.8
4.355 (100 %)
0.70
20.2* 0.3
21.1 f
4.352 4.356
1.52 1.63
19.6i 0.3 19.6 f 0.2
21.2i 0.2 21.2i 0.5
0.2
22.1 f
0.9
Composites AT60A AT62
A@,
Commercial references: SH1, SH20F A1203 + 33 voL% Ti(C,N) "submicron" A1203 + 33 vol-% Ti(C,N)
(2) Cutting edge displacement CED (which determines the degree of precision in machining hard pieces). (3) The quality of the cut metal surface is important for precise machining, it was described by 2 parameters: (i) The roughness was measured as a function of the cutting time. Ra gives the statistical average depth of the profile. (ii) On turning hardened steel at high cutting velocities > 250 d m i n there is a significant input of heat not only into the formed chips but also into the cut surface of the shaft. Depending on the state of wear of the tool, cutting forces and process temperatures may increase by an extent that causes a softening of the hardened steel, and in the present investigations the hardness was observed to drop from HRC = 5760 to values of about 55. Therefore, changes in the hardness were recorded as an additional parameter in such tests, and an interim machining operation with a low velocity v = 180 d m i n was required to restore a surface with the original hardness before every new experiment with v > 250 d m i n . The tools were run up to an upper wear flank width VB in the range between 0.15 and 0.25 mm. A minimum of two tests series at least were performed with each grade. Three different iron-basis workpieces were machined (German standard notations): - The Brine11 hardness of the investigated vermicular cast iron GGV40 was HB 185 f 14. This GGV, a candidate for engine blocks in the automotive industry, is hard to machine with commercial tools because its heterogeneous microstructure with residual graphite is highly abrasive. - The alloyed hard (chilled) globular cast iron GX3OOCrMo153 exhibited an average Rockwell hard-
-
406
-
0.6
*
ness HRC 42.9 1.3 (which was controlled after each cut). The material contained about 0.3 wt-% C, 15 wt-% Cr, and 3 wt-% Mo. - The hardness of the hardened steel 9OMnCrV8 was HRC = 58.4 1.5. It contained 0.86 wt-% C and 0.2 wt-% Si. Important additives are 1.98 wt-% Mn, 0.43 wt-% Cr, 0.14 wt-% Cu, 0.10 wt-% A], and 0.08 Wt-% V. The surfaces of the as-delivered shafts of hard cast iron and hardened steel were rough, macroscopically uneven and had to be pre-machined to get an equally prepared state for all tested inserts. This preparation is difficult because it associates the high hardness of the counterpart with discontinuous cutting conditions of changing frequency and power of impacts. Cubic boron nitride inserts @reborid@, Lach company, HanadGermany) failed on machining the hard cast iron, and hardmetal inserts (WC / 6 % Co, Vickers hardness = 16 GPa, K, = 9 MPadm) were not able to cut the rough outer shell of the hardened steel shaft at velocities between 50 and 100 d m i n (or more). The same observation applied for both commercial ceramic composites. Finally, laboratory grade ceramics originally designed for precision turning tests had to be used for this severe operation.
*
3. RESULTS 3.1 Machining with moderate processing temperatures: cast iron On machining hard globular cast iron, a smaller grain size of Ti(C,N) or Ti(C.0) reinforced composites increases the hardness (Tab. 1) and may reduce the wear on turning hard cast iron, but the effect is small compared with the qualitative leap to greatly reduced wear
of the pure (sub-pm) alumina tools (Fig. la), associated also with a reduced roughness R, of the cut s u r f a ~ e . ~ Wear of sub-pm Al,O,/Ti(C,O,N) composites increases slightly when nitrogen is introduced into the lattice of the covalent phase.4 Similarly, on high-speed machining vermicular cast iron GGV40 it is again the white corundum tool which surpasses both commercial Ti(C,N) reinforced composites and CBN not by percents but by factors of whole units (Fig. lb). In this operation, however, an unusual influence of the depth of cut reduces the preference of the A1,0, tool with increasing depth, and at a 1 mm the performance of all of the investigated tools was similar.
-
0.3 cast iron HRC
0.251
v f
--
a
-
250 Wmin
-
0.1
-
1
43 - 45
r - 1.2 mm
*-&-
mdrev
0.2mm
= SH1
-:
0.2-
m
> g0.15
-
5
9
0.1
-
3
I 0.0
Fig. 2a displays some results. The high hardness of the steel and the choice of tools with low thermal conductivity caused a large input of heat into the formed metal chips which appeared rather flaming than red-glowing. For the reference composites (SHl and commercial subpm composite), increasing temperatures were associated with a wear which still was limited at 220 d m i n , but already at 235 d m i n a greater crater-wear was observed, and more than 50 % of the commercial inserts exhibited a sudden, strong increase in the flank wear during the first 10-20 minutes of cutting accompanied by local or global fracture at the cutting edge. On the contrary, no one event of fracture was observed with the new sub-pm grades from our laboratory, neither with the tools of pure alumina nor with those reinforced with Ti(C,O). This behavior correlates with the surprising performance on discontinuous cutting (cp. 0 3.2.3).
The influence of the cutting speed on the time of use was described in greater detail previ~usly.~ At v = 220 d m i n (which is the upper limit of a reliable use of the latest generation of commercial composites), best re-
Y
c
3.2 Machining hardened steel 3.2.1 Precision machining at 200-250 d i n
I
0.05
'"
hardened steel HRC = 57 - 60
I
I
I
I
0
10
20
I
I
30 40 Cutting Hme (min)
vermicular cast iron HB = 185 (HRC v = 500rnImin = mm f = 0.25 mdrev
'.' c - 75"
0
I
I
50
60
70
14)
..
0
g5
/,?
0.4-
ATW .
6fold cutting path wth A1203 (criterion: VB = 0.3 mm)
// /
Y
I
-
(AIZOJ
-
0.0
0
L
I
0.3 I
I
1
I
I
I
1000
2000
3000
4000
5000
6000
, ' * .* .4
*sub-pm Alfi/Ti(C,Ol
a
ACS6
0 0.1
**0 '
71
Travelled cutting path (rn)
5
c .
r 10
0
I
I
20 25 Cutting time (min) 15
I
30
I
( V8>0.5 mm)
I !
Fig. 1. Wear flank width on turning hard cast iron (1a - above) and vermicular cast iron (1b - below). Note different cutting conditions in Figs. la/b.
Whereas the sub-pm grain size is imperative for successful machining hard globular cast iron, additional investigations have shown that on high-speed turning of Vermicular cast iron (GGV) the special advantage of white corundum ceramics at small cutting depths is a general property of the oxide character and is obtained with other grades as well.
a
-
02mm
1-45.
407
sults are obtained with the new sub-pm Al2O3/Ti(C,0) composite (which exhibits the highest hardness of all of the tested ceramics, Tab. 1). Here, the cutting edge displacement of pure sub-pm alumina is intermediate between our sub-pm A1203/Ti(C,0) composites and advanced commercial tools (Fig. 2a). A similar ranking was observed for the flank wear.3
3.2.2 High-velocity turning (hardened steel) The relationship between the new sub-pm composites and the pure alumina tool again changes at still higher cutting speed, and at 300-400 m/min A1203 once more gives the greater profit (Fig. 2b) as already observed on machining hard cast iron (Fig. 1a). On the contrary, most commercial composites fracture at 300 m/min within the first 5-10 min, and the one continuous curve displayed in Fig. 2b up to 15 min was obtained at a softer part of the steel shaft (HRC = 52-57). At a velocity of v = 300 m/min, increasing with time wear deteriorates the cutting edge in a way that obviously causes an increasing input of heat into the surface of the machined steel. For some of the tests, the resulting decrease in hardness of the workpiece was measured and is given in Fig. 2b as an additional parameter.
3.2.3 Discontinuous coarse cutting of hardened steel It is generally assumed that sintered pure alumina ceramics with their typical toughness of K,, < 4 m a d m cannot be used as cutting tools for intermittent machining or with high feed rates. However, the rough, uneven hard shafts with their imperfect circular cross-sections used in the present investigations could not be premachined with commonly recommended tools (cp. 0 2): CBN inserts failed on machining the shafts of hard cast iron, hard metal inserts and the commercial A1203/Ti(C,N)composites were unable to cut the rough outer shell of the hardened steel shafts. Therefore, the new laboratory grades had to be applied and were tested for this purpose at v = 120 m/min, with a large feed rate f = 0.3 &rev, and a depth of cut a = 1.5 mm. In regions with substantial deviations from a circular crosssection, a = 1.5 mm is an average value when in a first cut some parts of the circumference where not cut at all whereas the cut was rather deep at other positions. Besides of hard impact loading, such intermittent conditions are associated with severe thermal shock (indicated by the fluctuating appearance of the red color of the cutting tip and documented by video recording). Both commercial grades (SH1, submicrometer grade) failed by fracture within 1-2 min after rapidly increasing tool wear (deteriorating the roughness of the cut metal surface: Ra = 4-5 pm after c 1 min). On the contrary, both sub-pm A1203 (AC41) and the new composites (AT60NAT62) were successfully applied, Fig. 3 gives the flank wear. These new cutting ceramics do not only not fracture macroscopically during one hour of severe cutting with one tip (!) but retain microscopically nearly perfect cutting edges (providing a surface quality of R, 1-3 pm of the hardened shaft for
-
408
/il 3 9 0.2
0.1
0.0
0
I
I
I
5
10
15
I
I
I
20 25 30 Cutting time (min)
I
I
I
I
35
40
45
50
Fig. 3. Flank wear on severe discontinuous cutting of hardened steel with sub-pm tool ceramics. t 5 1 h already in this first cutting operation with the coarse conditions notified in Fig. 3). Surprisingly, with these severe conditions the general performance of pure alumina ceramics (AC41) is even more prospective than the behavior of the new sub-pm composites - which again exhibit a similar behavior with Ti(C,O,N) and with Ti(C,O) reinforcements as it was observed on turning hard cast iron (Fig. 2a).
4. DISCUSSION When at moderate processing temperatures at the cutting edge pure sub-pm corundum tools exhibit less wear than both the advanced commercial and the new sub-pn composites (e.g. on turning cast iron, Fig. l), this difference is most probably caused by similar tribo-chemical wear mechanisms of the covalent Tic, Ti(C,O) and Ti(C,O,N) phases as previously observed on sliding wear at room temperature.5.6 Circumventing this problem by using pure, thermodynamicallyhighly stable alumina inserts, the wear resistance is improved as long as the process temperature at the cutting tip is low enough to prevent a detrimental influence of the low creep resistance of the submicrometer alumina microstructure.
On turning hardened steel, again both pure alumina and the laboratory grade sub-pm composite (AT60A) exhibit less wear than all of the commercial cutting ceramics, but the mutual ranking of A1203 and A1203/Ti(C,0)ceramics changes when different shaft materials are cut at different velocities. It is obvious that the technical ranking of these new tool ceramics is governed by some important basic properties: (a) With moderate conditions, chemical interactions cause preferential (local) wear of the covalent carbide constituents of the microstructure in A1203/TiCcomposites even at room temperature with the consequence that at equal crystallite sizes pure sintered corundum ceramics are more wear resistant than the composite.~*6The same ranking was observed here for the flank wear of ceramic tools on turning hard cast iron (Fig. la). (b) On machining hardened steel at 200-250 m/min,
flaming chips indicate a higher process temperature than on turning hard cast iron at the same velocity. Whereas pure sub-pm A1203 exhibits intense creep at temperatures > 1100 OC, a continuous network of covalent crystals reduces the creep rates in composites with more than 25 vol-% of Tic, Ti(C,N) or Ti(C,O). Indeed, comparing at v 5 250 d m i n the machining of cast iron and steel, the increasing temperature increases the absolute wear only of the alumina tool and changes the ranking in a way that on machining hardened steel it is now the sub-pm A1203/Ti(C,0) composite which (at similar grain sizes as in the A1203 insert) exhibits least wear and cutting edge displacement (Fig. 2a). (c) This advantage of composites is lost when at still higher temperatures oxidation of the Ti(C,O) phase or more intense chemical reactions start. With the results in Fig. 2b, the commercial composites with Ti(C,N) are more susceptible to such processes than the new sub-pm laboratory grades with Ti(C,O), but even the latter deteriorate more under the influence of high temperatures at 300 d m i n than pure alumina (AC41) - in spite of its susceptibility to (macroscopic) creep. The same behavior of a stronger influence of high cutting velocities on the flank wear of composites was also observed on machining hard cast iron.4 Beyond the advantage of the known thermodynamic stability of pure corundum inserts, the comparison with two different grades of CBN and with a commercial Al20,/Ti(C,N) cutting ceramic on high-speed machining of vermicular cast iron (GGV40) reveals a unique abrasive resistance of the sub-pm corundum microstructure. When on severe, discontinuous cutting the central goal is to remove as much of the hardened steel as possible, the resulting roughness of the metal workpiece reported above is technically not relevant. It is, however, another evidence of the wear resistance of ceramic tips with a sub-pm microstructure. The strange discrepancy between the known promotion of macroscopic high-temperature creep in oxide sub-pm microstructures and the exceptional performance of such sintered corundum tools just under those cutting conditions which should induce creep is perhaps an indication that at the cutting tip the spatial extension of the zone of really high temperatures > lo00 OC may be much smaller than often assumed. Under severe conditions like the machining of vermicular or hard globular cast iron or on intermittent or highspeed machining of hardened steel, the fracture of cutting edges is influenced by time dependent wear processes (cp. the observation of fast fracture of the commercial composites on high-speed turning of hardened steel afer preceding rapidly increasing flunk wear). Hence, the high frequency of fracture of commercial grades and the surprising lack of such events with new laboratory grades must not be compared readily with the bending strength at room temperature. For the commercial grades, manufacturer's data give a strength of 600 MPa for SH1 (no data for the submicrometer grade available). The bending strength of the new sub-pm
composite is about 800 MPa,' a lower strength of about 650-700 MPa was observed for the submicrometer A120, ceramic.2 The fracture toughness of all of the laboratory grades is 3.3-3.8MPadm, whereas unusually high values of 5.5 MPadm (SH1)and 6.6 MPadm ("submicron" commercial composite) are given by manufacturer's information for these reference tools. None of these data explain the much smaller fracture risk of the new laboratory grades on machining hardened steel. Obviously, complex wear-inducedprocesses of flaw-generation have to be considered to understand the high global (and microscopic) stability of the new submicrometer ceramics. For the further development of new cutting ceramics it is the unique message from all of the investigated parameters (flank wear, cutting edge displacement, and roughness) that the leap to cutting ceramics on a basis of sintered corundum which stand both - the high temperature on high-speed machining and - the mechanical impacts and the associated thermal shock on discontinuous cutting of hardened steel comes with the submicrometer grain size of the tools: a grain size D < 0.7 pm seems to be the first, most important requirement (possibly associated with a required strength > 600 MPa). Other parameters like fracture toughness (KI,)and creep resistance may give additional optimizations but, surprisingly and contrary to the common opinion, their influence is of secondary importance. The influence of high-velocity turning on the hardness ofthe workpieces (Fig. 2b) is important for different aspects in the development of new machining technologies. On the one hand, softening of the hardened steel promotes high material removal rates, and turning of hardened steel with cutting ceramics that are less expensive than cubic boron nitride (CBN) becomes possible for other operations than only f i e machining. If, on the other hand, for the final step of precisionmachining a constant high hardness of the cut surface is required, critical limits exist W for the cutting velocity (about 250 d m i n in the present investigations), or, respectively, W for the tolerable tool wear (Fig. 2b: VB 5 0.15 mm at v = 300 dmin). These technical demands are, of course, affected by the choice of the ceramic tool depending on its thermal and wear properties.
5. CONCLUSIONS Advanced sub-pm ceramics have been presented that machine vermicular cast iron, globular chilled (hard) cast iron, and hardened steel. Both covalent-phase reinforced composites und pure sintered alumina provide a leap in the cutting performance compared with CBN or with commercial cutting ceramic composites. For this end, control of tool microstructure with a low frequency of flaws and with an average grain size < 0.7 pm is fun-
409
damental to a successful technical performance. Associated with different advantages and limitations in important basic properties as chemical stability, creep resistance and hardness, the ranking of the groups of new sub-pm single-phase corundum and sub-pm composite tool ceramics is different depending on - the microstructural properties of the machined materials, andon - the cutting conditions (resp. the process temperatures that develop at the cutting tip). The most remarkable surprise is the excellent performance of pure sub-pm alumina, comparable for instance with recent advances in the grinding efficiency enabled by similar corundum microstructures7(where, until present, surprisingly few attention was devoted by the scientific ceramic community to the similarly unexpected stability of sintered sub-pm corundum grits against hard mechanical impacts!).Conventional A1203 cutting tools disappeared from the market 25 years before and were totally replaced by composites with ZrO, and Ti(C,N) reinforcements. The present investigations, however, indicate that new pure (and sub-pm) corundum inserts can be applied successfully for both - the precision-machiningof hard materials on continuous cutting (with small feed rates), and likewise - under severe discontinuous conditions. After the failure of hardmetal and CBN inserts on premachining the rough (as delivered) shafts of hard globular cast iron and hardened steel and with the large wear of CBN on high-speed machining vermicular cast iron, pure submicrometer alumina was the one reliable solutions for these discontinuous operations. The observation of a reliable use of sub-pm sintered corundum tools on discontinuous cutting operations is of outstanding importance for manufacturing processes in the automotive industry. For example, the surprising intermittent use of sub-pm corundum ceramics designed originally for high-speed precision turning of hard workpieces should enable the integration of these new tools into complex machining processes which omit the repeating exchange of ceramic tools when cutting the edge of a shaft. In this way, the benefit for productivity will be much higher than given only by the increased lifetime (or cutting path) of the improved insert.
REFERENCES lA. Krell, L.-M. Berger, and P. Blank. Submicrometer AI,O,/Ti(C,O,N) Composites for Tool Applications. Advanced Ceramics, Materials, and Structures. Proc. 22nd Annual Conference on Composites, Vol. A. The American Ceramic Society, Westerville/OH (1998) 139-146. *A. Krell and P. Blank. The Influence of Shaping Mettpd on the Grain Size Dependence of Strength in Dense Submicrometre Alumina. J. Europ. Ceram. Soc., 16 [ l l ] (1996) 1189-1200. 3A. Krell, P. Blank, L.-M. Berger, V. Richter. Submicrometer Alumina Cutting Ceramics for Continuous and Discontinuous Machining of Hard Materials. cfi I Ber. Dz. Keram. Ges., 76 [4] (1999) 23-29.
410
4A. Krell, P. Blank, L.-M. Berger, V. Richter. Alumina Tools for Machining Chilled Cast Iron, Hardened Steel. Bull. Am. Cerum. Soc., 78 [12] (1999) 65-73. 5V.N. Koinkar, B. Bwhan. Microtribological Studies of AGO,, Al,O,-Tic, Polycrystalline and Single-Crystal Mn-Zn Femte, and S i c Head Slider Materials. Wear, 202 [l] (1996) 110-122. 6A. Krell.D. M a e . Effects of Grain Size and Humidity on Fretting Wear in Fine-Grained Alumina, A1203/TiC, and Zirc0nia.J. Am. Ceram. Soc.,79 [4] (1996) 1034-1040. 'A. Krell, P. Blank, E. Wagner, and G. Bartels. Advances in the Grinding Efficiency of Sintered Alumina Abrasives. J. Am. COUIPZ. SOC.,79 [3] (1996) 763-769.
PROCESS STRATEGIES FOR GRINDING OF ADVANCED CERAMIC CUTTING TOOLS H. K. Tonshoff, T. Friemuth, D. Hessel Institute of Production Engineering and Machine Tools, D-30159 Hannover, Germany ABSTRACT In cutting tool manufacturing the tool grinding process is responsible for the properties of the cutting tool surface, the subsurface and the cutting edge geometry. The cutting performance and wear behavior of uncoated ceramic inserts is carried out in hard turning. This paper shows the influence of the tool grinding process on the surface and subsurface properties and the cutting edge quality of ceramic inserts as well as the consequences on the tool life time in a hard turning process. In-process sharpening techniques which minimize the influence of the grinding wheel wear on the grinding process and improve the grinding process behavior are presented.
to these results small negative rake angles should be used to improve the tool life in hard turning with ceramic cutting tools. Therefore it is necessary to reduce the size of the chamfers.
",.145lllWl
+. 02nm I
-
l a 0.1 mn
up,* W.N) 0 IW
Fig. 1 : Influence of chamfer size on tool life
INTRODUCTION The development of cutting tools is concentrating on new tool materials with enhanced mechanical properties, on optimized geometries and on coatings to improve the cutting performance. For applications in turning of hardened steel in most cases uncoated cutting tools like polycrystalline boron nitride (PCBN) or ceramics are used. In manufacturing of uncoated cutting tools the grinding processes are responsible for the surfaces and the cutting edge micro geometry as well as the surface integrity. Additionally in hard turning high mechanical and thermal loads are induced in the tool-workpiece-contact. Therefore properties like superior hardness especially at high temperatures and adequate toughness of the cutting tool materials are necessary to achieve a high tool life time.
Different grinding processes are necessary to manufacture cutting tools. In case of inserts the main steps are shown in Fig. 2. Starting from a sintered blank the surface grinding process is used to generate the exact thickness and the tool face. Afterwards a periphery grinding process creates the cutting tool flanks and radii. These grinding processes are responsible for the micro geometry of the cutting edge (Fig. 2 left bottom).
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THE MICRO GEOMETRY OF THE CUTTING EDGE 3zwzes15 0 VW
The main wear mechanisms in hard turning are the flank and the crater wear. The flank wear is usually used to determine the tool life time T,. In hard turning only small depths of cut a,, are used. As the contact area in hard turning is very small, the micro geometry of the cutting edge decisively influences the wear mechanisms. Fig. 1 shows the influence of the chamfer size, which is responsible for the effective rake angle yeff and the relative tool life in dry turning of hardened steel. The tool life criterion is defined as a width of flank wear land of VBc = 200 pm. A high negative rake angle leads to increasing thrust forces. Hence the load on the flank of the tool rises. This interaction induces a higher wear of the tool flank and finally a shorter tool life. Due
Fig. 2: Effects of tool grinding on the micro geometry of the cutting edge As a result of the brittle material behavior of ceramics excavations of the cutting edge can be observed. This rough and irregular formation can not be used as a cutting edge in the later application of the tool. In order to ensure the necessary cutting edge quality with a homogeneous and sharp micro geometry an additional chamfer preparation is carried out. These chamfers can be characterized by the chamfer width and angle. The minimum size of the chamfer is defined by the largest excavation after the grinding processes. Fig. 2 illustrates the close correlation between tool grinding, the created size of excavations,
411
the necessary size of the chamfer and the generated micro geometry.
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As the size, number and distribution of the excavations at the cutting edge is determined by tool grinding, this process has to be optimized to improve the cutting edge quality. The main aim in grinding of cutting tools is to increase the output and quality and simultaneously to decrease surface and subsurface damages of the cutting edges. Especially ceramic cutting tool materials lead to break-outs and micro cracks at the cutting edges which cause a short life time in the following machining operation. In grinding Ti(C,N) reinforced alumina ceramics the cutting edge quality can be positively influenced by using smaller grained grinding wheels (Fig. 3). The cutting edge quality is mainly determined by mechanical loads during grinding. By decreasing the grain size of the grinding layer, the number of active grains increases by the square of the grain size reduction. Typically resin bonded diamond grinding wheels are used for grinding of ceramics. The grinding process is heavily depending on the size of the grains in the grinding layer. For example due to the reduction of the average grain size fiom 6 = 41 pm (D46) to & = 10 pm (D10)about sixteen times more grains are available in the grinding layer. Because of the higher number of active grains there is a reduction of mechanical loads in subsurface layers induced by a single grain [ 1,2].
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Fig. 4: Influence of the diamond grain size on grinding forces However, as shown in Fig. 3, to improve the cutting edge quality of ceramic inserts fine grained grinding wheels are required. The economical use of these grinding wheels is only possible, if several restrictions are taken into account. It is necessary to reduce the influence of the microgeometrical wear mechanisms (blunting of grains and chip loading) on the grinding wheel performance by using short times between two dressing processes and by using lower material removal rates in grinding. These results show that in conventional grinding the two contrary targets - high efficient grinding and minimum cutting edge damage - have to be fidfilled (Fig. 5 ) .
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Fig. 5 : Conflicts in process design for tool grinding
Fig. 3: Influence of the diamond grain size on cutting edge quality In grinding with small grained grinding wheels typically a different process behavior is observed. While grinding with coarse grains the normal forces only increase slowly with a higher material removal, but in grinding with small grains the forces are higher and even raise significantly (Fig. 4). The grinding performance of fine grained grinding wheels is reduced because only a smaller grain protrusion and chip volume is available. These grinding wheels behave very sensitive against a variation of the grain protrusion. As a consequence of the microgeometrical wear mechanisms taking place in the grinding layer the grain protrusion and chip volume decrease.
412
One solution to this problems in grinding of ceramic cutting tools is to use small grained grinding wheels in combination with continuous in-process-sharpening techniques. The main aims of this parallel process are to preserve the grinding ability of the grinding layer and to reduce the influence of several microgeometrical wear mechanisms during grinding. Most important is a constant grain protrusion in the grinding layer especially in grinding with very small grain sizes. Therefore it is necessary to remove all blunt grains and carry out the workpiece material in order to avoid loading with chips. By using a well-balanced in-process sharpening the grinding process is characterized by constant grinding forces and workpiece qualities. This paper shows results in using two different sharpening principles (Fig. 6). On the one hand the typical abrasive sharpening method with a corundum stick is tested. This working principle is characterized by a high additional grain treatment during the sharpening process and therefore leads to an increased macrogeometrical grinding wheel wear. On the other hand the sharpening
by electro contact discharge (ECD) is characterized by less grain strain. Therefore the additional grain wear by ECD sharpening is minor. The real sharpening effect arises from the thermal erosion of the bonding material. This process removes the electrical conductive bond material (e.g. bronze) by a local thermal treatment [3,4]. sharpening methods abrasive sharpening
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Fig. 6: Principles for in-process sharpening By continuous in-process sharpening with an abrasive stick it is possible to reduce the effects of the microgeometrical wear mechanisms on the grinding forces (Fig. 7). Due to the removal of chip loading and bond material a constant grain protrusion is preserved. The sharpening intensity can be easily adjusted by the removal rate of the abrasive stick QdsB. An increase of QdsB leads to decreasing grinding forces. According to a higher sharpening intensity the macrogeometrical grinding wheel wear Ar, rises simultaneously. Hence, the inprocess sharpening changes the process behavior fundamentally. Stable grinding processes with constant grinding forces are possible and the sharpening intensity is responsible for the level of the grinding forces and the grinding wheel wear Ars.
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Fig. 7: Tool grinding with abrasive in-process sharpening The in-process sharpening by ECD in principle has the same positive influence on the grinding process. It is able to reduce the effects of the microgeometrical wear mechanisms on the process behavior. In ECD sharpening the intensity is easily adjustable by the sharpening voltage UdsO.
The increase of the sharpening voltage (D.C.) generates a higher grain protrusion in the grinding wheel and reduces the grinding forces. Higher voltages lead to a higher grain protrusion, because of an intensified removal of the chip loading and of bond material. Therefore, the macrogeometrical grinding wheel wear Ars increases [ 5 ] . 40
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Fig. 9: Comparison between abrasive and ECD in-process sharpening Fig. 9 shows the comparison between these two sharpening methods concerning the macrogeometrical grinding wheel wear Ars after a specific material removal of V', = 1000 mm3/mm. The three groups of columns represent different sharpening conditions with a variation of the sharpening intensity. Comparable grinding processes are adjusted by approximately equal grinding forces. The total grinding wheel wear Ars is divided into the wear caused by the grinding process Arsw (Fig. 9, hatched area) and the additional wear through sharpening Arsds(B)(Fig. 9 ,white area). Low sharpening intensities cause higher grinding forces and low additional wheel wear Ars. When using higher sharpening intensities the wheel wear increases. The differences between abrasive and ECD-sharpening grow with higher intensities. This shows the technological advantage of the electro contact discharge process in sharpening. The additional grain treatment caused by sharpening is very low and therefore the macrogeometrical grinding wheel wear Ars is reduced.
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PROPERTIES OF THE CUTTING TOOLS AFTER GRINDING In tool manufacturing different diamond grain sizes can be used in grinding. As a result of a variation of the process parameters the contact conditions in grinding can be changed. The effects of the process parameters on the contact condition in grinding are characterized by the average chip thickness ku.In face grinding the chip thickness mainly depends on the four parameters: diadiamond concentration C, cutting mond grain size 6, speed v, and axial infeed speed vfa [3]. The influence of these parameters on the chip thickness is expressed in formula 1. The chip thickness describes the mechanical load in the material removal and is predominantly influenced by the diamond grain size.
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In addition to the mechanical load it is important to take the thermal conditions in grinding into account. In grinding processes typically a high amount of energy is generated and transformed into heat which penetrates the workpiece. Hence the grinding power is an important value to focus on (Fig. 11). With decreasing chip thickness the grinding power related to the contact area of a single grain Pll,~and the thermal load at the ground surfaces increases. The workpiece properties are influenced by both mechanical and thermal - loads during the grinding process. These loads are responsible to cause a certain state of surface integrity. Decreasing chip thickness in grinding leads to lower mechanical loads and less excavations at the cutting edges. Simultaneously the thermal effect in the contact zone between the diamond grains and the workpiece surface is intensified. As a result of these mechanisms the material characteristics after grinding are closely connected to the grinding conditions which can easily be summarized by the calculated chip thickness kU. The surface integrity after grinding can be described by the residual stresses, which are measured with X-ray diffraction. The residual stresses are depending on the grinding conditions. Fig. 12 shows the close correlation between the residual stresses in the surface of the ground workpieces (cutting tools) and the chip thickness in grinding. +lo0 MPB
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Fig. 11 summarizes the effect of the chip thickness on the cutting edge quality. The measurement of the cutting edge roughness is based on the measuring and evaluating principle of the maximum peak-to-valley height The cutting edge roughness is directly depending on the chip thickness. The mechanical load during the grinding process is responsible for the formation of the ground cutting edge. In order to reach cutting edges with less excavations it is necessary to reduce the chip thickness in grinding.
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Fig. 11: Effects of mechanical and thermal load in grinding
Fig. 10: Potentials of using fine grained grinding wheels with in-process sharpening The improvement by in-process sharpening can be realized with different diamond grain sizes of the grinding wheels. Fig. 10 shows the process behavior of fine grained grinding wheels with the additional use of ECD in-process sharpening. It is possible to use diamond grain sizes down to D10 in efficient grinding processes with constant forces. This is only possible because the inprocess sharpening preserves a stable level of the grain protrusion by the removal of bond material and of chip loading. These results show, that grinding with in-process sharpening is able to combine high performance grinding processes and the use of small grained grinding wheels in order to reach a high cutting edge quality.
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Fig. 12: Residual stresses depending on the chip thickness in grinding Grinding processes with a large chip thickness generate high compressive residual stresses in the subsurface of the ground cutting tool. If the chip thickness is reduced, especially to create high precision cutting edges (Fig. l l ) , lower compressive stresses are effected. The residual stress state influences the fraction and wear re-
sistance of the ground tool flank. The mechanical and thermal loads of the turning process are responsible for additional stresses in the cutting tool. The capability of the cutting tool surface to resist those stresses depends on the surface integrity after grinding [6, 7, 81. In order to examine the significance of the residual stress state after the tool grinding process, indentation tests with a Vickers pyramid are carried out. By this mechanical load a stress state in the subsurface of the workpiece is generated. The superposition of the residual stresses caused by grinding and the additional stresses generated by indentation are responsible for the material behavior. Fig. 13 shows four different ground surfaces after similar indentation tests. The surfaces which are ground with a coarse grained grinding wheel are characterized by high compressive residual stresses and no failures occur after the indentation test.
- 1
Fig. 13: Surface damage after Vickers indentation (F = 590 N) By using fine grained grinding wheels and small chip thicknesses in grinding tensile residual stresses are created in the subsurface of the ground workpiece. The damages of the ground surface and subsurface is stronger in grinding with a small chip thickness. These process conditions lead to a less beneficial surface integrity and reduce the fraction resistance of the ground surface. According to these properties grinding with a high chip thickness should be preferred. The results of the tool grinding process can be summarized as follows. There are two main influences of the grinding process. On the one hand there is the mechanical influence which is responsible for the created quality of the cutting edge. In order to reach the continuously increasing demands on the cutting edge quality, it is necessary to reduce the mechanical treatment of the cutting edge by decreasing the average chip thickness in grinding. On the other hand the thermal conditions in grinding have to be taken into account. Improvements in grinding require a reduction of thermal loads to achieve a better surface integrity. These two directions in the design of tool grinding processes are contrary. Therefore, it is necessary to investigate and acquire the best compromise in process design. In the following an analysis of the effects of the process design in tool grinding on the wear behavior of Ti(C,N) reinforced ceramic cutting tools is carried out in hard turning.
EFFECTS OF THE TOOL GRINDING PROCESS ON THE WEAR BEHAVIOR The main wear mechanisms in hard turning are characterized by the abrasive flank and crater wear. Especially the width of flank wear land VBc is used as the tool life criterion. The properties of the tool flank concerning surface integrity and residual stresses are generated in the tool grinding process and influenced by the chip thickness in grinding. Therefore the effects of the conditions in grinding on the wear resistance of the tool flank in hard turning are investigated. Different ground tools are used in hard turning (Fig. 14). Typically there are two phases in the wear behavior of cutting tools in hard turning. In the first phase the flank wear suddenly increases up to values of VBc = 60 - 80 pm. This value depends mainly on the first contact condition between the sharp cutting edge and the workpiece material. Afterwards the second wear phase begins. In this phase the flank wear increases continuously with the cutting length. The reached tool life depends on the growth of the flank wear in both phases. Whereas the first phase is more or less a random effect in the wear behavior, it is necessary to look closer at the continuous wear in the second phase. The small diagram in Fig. 14 shows the rate of the growth of flank wear in one minute cutting time and is called wear rate VVB. 250,
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Fig. 14: Effect of grain size in tool grinding on the wear behavior of the cutting tool in hard turning The results in Fig. 14 show the influence of the grain size of the diamond grinding wheels which are used for tool grinding on the average wear rate determined by four replications. Tool grinding with small grain sizes causes a higher wear rate in hard turning and therefore a worse tool life is reached. The increase of the grain size in tool grinding effects lower wear rates. The relationship between the diamond grain size and the wear rate is not linear. Therefore it is necessary to have a closer look on the interactions. It has to be taken into account, that these tests are carried out with the same chamfer size and shape of the cutting tool. If smaller grain sizes in the grinding wheel are used, the chamfer at the cutting edge can be minimized and leads to lower passive forces and enlarged tool life. Hence, the effects in grinding correlate precisely with the chip thickness h,,, also an interaction between the
415
chip thickness and the wear rate should be analyzed. Fig. 15 shows the wear rate in hard turning as a h c t i o n of the chip thickness in grinding the radius of the cutting tool buR. The curve can be divided into two different = 0.27 pm the areas. If the chip thickness is less then buR wear rate increases by a W h e r reduction of the chip thickness. Consequently with a smaller chip thickness the wear behavior is worse and the tool life decreases. If the = 0.27 pm the wear chip thickness is bigger than buR characteristic changes. The chip thickness has less influence on the wear rate. Up to a chip thickness of hcuR= 0.35 pm the wear rate is constant. These results show the decisive effects of the grinding conditions in tool grinding and the created surface integrity as well as the influences on the cutting tool performance, wear rate and tool life.
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Fig. 15: Influence of the chip thickness in tool grinding on the wear rate in hard turning
CONCLUSION The cutting edge quality, surface integrity and residual stresses are decisively influenced by the tool grinding process. In order to achieve a high cutting edge quality with less excavations it is necessary to reduce the mechanical load during grinding by using fine grained grinding wheels. The grinding processes are strongly influenced by the microgeometrical wear mechanisms especially in using small grained grinding wheels. Inprocess sharpening technologies enable stable and productive grinding processes by reducing the influence of the microgeometrical wear mechanisms taking place in the grinding layer. The performance of cutting tools depends on the properties of the cutting tool material. These properties are influenced by the tool grinding process. Increasing thermal loads in grinding lead to damages in the subsurface of the ground cutting tool. If the chip thickness in grinding remains below a critical level the reachable tool life decreases. However, with smaller excavations after grinding smaller chamfers can be realized and the tool life increases. The knowledge of these contrary effects enables the optimization of the process design in tool grinding.
416
ACKNOWLEDGEMENT The work described in this paper was carried out with support of the German Federal Ministry of Education, Science, Research and Technology (BMBF) in a special research program on new materials (MaTech).
REFERENCES (1) Wobker, H.-G.: Schleifen keramischer Schneidstoffe, Dr.-Ing. Dissertation Universitiit Hannover, 1991 (2) Marshall, D.B.; Evans, A.G.: Nature of Machining Damage in Brittle Materials, Proc. R. SOC.Laond. A 385 (1983), pp. 461-475 (3) Friemuth, T.: Schleifen hartstoffverstiirkter keramischer Werkzeuge, Dr.-Ing. Dissertation, Universitiit Hannover, 1999 (4) Karpuschewski, B.; Friemuth, T.; Fripan, M.: InProzeD-Scharfen - abrasiv oder kontakterosiv, Gesucht: ((Scharfe Scheibenn zum Werkzeugschleifen, Technica 5/98, pp. 14-17, 1998 ( 5 ) Tbnshoff, H. K.; Friemuth, T.: Electro Contact Discharge Dressing (ECDD) of Diamond Wheels for Tool Grinding, International Conference on Precision Engineering, 20.-22. November 1997, Proceedings Volume 2, pp. 565 - 570, Taipei, Taiwan, 1997 (6) Tbnshoff, H. K.; Brinksmeier, E.: Determination of the Mechanical and Thermal Influences on Machined Surfaces by Micro-hardness and Residual Stress Analysis, Annals of the CIRP 20 (1980) 2, 1980 (7) Lucca, D.A.; Brinksmeier, E.; Goch, G.: Process in Assessing Surface and Subsurface Integrity, Annals of the CIRP 47 (1998), Vol. 2
ULTRASONIC ASSISTED FACE GRINDING AND CROSS-PERIPHERAL GRINDING OF CERAMICS Eckart Uhlmann and Nikolai-Alexander Daus* Institute for Machine Tools and Factory Management, TU Berlin, Germany
ABSTRACT The examples of ultrasonic assisted face grinding and cross-peripheral grinding demonstrate a possibility to achieve high material removal rates in machining brittle hard materials, while keeping the workpiece quality constant. Apart from an increase in active speeds, the kinematics lead to constantly altering nominal engagement angles of the abrasive grains at a frequencyf;ls= 20 kHz and an amplitude Ails of several micrometers. The material removal and wear mechanisms of ultrasonic assisted grinding will be explained on the basis of the description of the changing kinematic conditions compared to conventional grinding processes.
In few process variants on the other hand, the ultrasonic oscillation was effected by the tool. Figure 1 displays the processes ultrasonic assisted face grinding and cross peripheral grinding, which are both suitable for die-sinking and for machining grooves of brittle hard materials such as ceramics.
Fig. 1: Ultrasonic face grinding and cross-peripheralgrinding
INTRODUCTION An obstacle to the introduction of functional components of brittle hard materials such as advanced ceramics are still the considerable finishing costs. Finishing is mainly conducted by means of material removing processes with a geometrically undefined cutting edge, grinding with diamond grain being the most important among these processes. Complex contours are often manufactured using ultrasonic lapping. Another possibility for the machining of ceramics is the technological combination of the processes ultrasonic lapping and grinding, creating the hybrid process ultrasonic assisted grinding. As early as 1956, bound abrasive grains were used in ultrasonic assisted machining since the obtainable machining results in ultrasonic lapping are very small (1). Up to now, a number of grinding processes has been superimposed with ultrasonic vibrations. Many authors agreed that a reduction of process forces leads to an increase in material removal rate (2-7). However, the influence on the surface quality of the workpiece and on the tool wear is still controversial. Whilst the range of investigated process variants is rather wide, the effective correlations leading to an improvement compared to conventional grinding were neglected in most cases.
PROCESS VARIANTS Since the defined transfer of ultrasonic vibrations to grinding wheels, e.g. in surface grinding, is rather difficult, the ultrasonic vibration is introduced in the contact zone by means of a stimulation of the workpiece in many process variants.
Presuming an ideal stiff machine system and an ideal cylindrical tool, material removal in face grinding is conducted by the engaging diamond grains on the front of the tool. In cross-peripheral grinding, material removal mainly takes place with the help of abrasive grains on the periphery. Due to axial ultrasonic oscillation however, the diamond grains at the face side are responsible for the surface formation in the grounds of the machining track. The Institute for Machine Tools and Factory Management has an ultrasonic machining centre of the company Sauer, Stipshausen, Germany, at their disposal, which is specifically laid out for the loads occurring in these grinding processes. Grinding tools with a steel sintered bond and abrasive diamond grains were used for the technological investigations. The tools are designed as hollow bodies through which the cooling lubricant is lead to the contact zone. In face grinding it is thus guaranteed that the contact zone is supplied with a sufficient amount of cooling lubricant even at high depths of cut.
KINEMATICAL CORRELATIONS The superposition of the kinematics of conventional grinding processes with a longitudinal ultrasonic oscillation results in a fundamental change in the resultant motion. Irrespective of the position of active partners, tool and workpiece and the direction of the oscillation, there is a recurring change in active speed Ave and the frequencyhIs. depending on the amplitude Consequently, there is the acceleration ails reaching its maximum on den extreme positions of the oscillation. Compared to the conventional process, the result in this case are different maximum engagement angles aell,~,mm 417
on the points of inflection of the oscillation perpendicular to the workpiece surface or different cross grinding angles Q,, parallel to the workpiece surface. In ultrasonic assisted face grinding the resultant cutting speed results of the concurrence of the cutting speed v, resulting from the peripheral speed of the tool, the axial feed speed vfa and the axial ultrasonic speed vfius.The ultrasonic speed is calculated as follows: vfol/S = 2 ' A [ J S ' I r * h / S
*sin(2*n'hlS 't)-
(eq* l )
Depending on the kinematic parameters, the duration of the oscillation and the amplitude can be varied. Figure 4 illustrates the dependency of the maximum engeagement angle or cross grinding angle, respectively, on the cutting speed and the ultrasonic amplitude. It becomes distinct that an increase of the amplitude or of the frequency results in steeper angles. Related to this the ultrasonic accelerations and presumably the mechanic stresses of the active partners also rise.
Figure 2 displays the movement path using the example of a diamond grain.
0 Fig. 2: Active movement in ultrasonic assisted grinding and engagement angels
If the axial ultrasonic speed runs parallel to the machined surface, the abrasive grains engage in the form of a sinus oscillation in the workpiece surface. This way, the contact between the grain edge and the workpiece surface is continuously maintained. However, there is a change in cross grinding angles of the engaging diamond grain. This is the case in ultrasonic assisted cross-peripheral grinding. Compared to conventional grinding, the contact length is much higher. Figure 3 (left) illustrates such a grinding trace on aluminium oxide.
Fig. 3: Sinusoidal continual engagement in aluminium oxide and pulsed discontinual engagement in zirconium oxide
With axial oscillation however, the engagement of diamond grains on the face is characterised by instantaneous local interruptions of the contact between abrasive grain and workpiece that recur with the work frequency. If the fracture toughness of the machined material is sufficiently high, it is possible to obtain a surface structure with trough shaped engagements. Figure 3 (right) displays a surface produced by ultrasonic assisted cross-peripheral grinding of zirconium oxide. The engagement was effected through diamond grains on the face of the tool. The maximum engagement angle or cross grinding angle, respectively, is the characteristic parameter for the motion path resulting from kinematics of ultrasonic assisted grinding. From the size of this angle it is possible to derive the impact of the ultrasonic oscillation on the process.
418
1
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Due to the current technical restrictions concerning the range of parameter values, no statements can yet be made as to what concrete effects a rise in amplitudes, frequencies and cutting speeds would have on bonding system, abrasive grains and sub-surface damage on the workpiece.
MATERIAL REMOVAL AND WEAR MECHANISMS A result of the pulsed grain engagement in ultrasonic assisted grinding are mechanical loads on workpiece and tool that are higher than those in conventional grinding. The thermal loads on the other hand decrease (9). This is caused by a microsplintering of the diamond grains due to high mechanical loads. Contrary to the conventional process, flattening of the grains were rarely noted. The splintering causes a perpetual generation of sharp edges on the grains that maintain the cuttability of the tool and reduce the friction. With the reduction of friction and contact times temperatures in the sub-surface of ceramic workpieces also decreased. Figure 5 displays the front diamond grains of a grinding tool of the specification D126 St50 C90 after conventional and ultrasonic assisted machining of silicon carbide with v, = 2,3 d s and V J ~ 4 mdmin. The investigated process was face grinding with an amplitude A(,,s= 14 pm and a frequency ftIs= 20 W z . Only after a surface related material removal of v,""= 55 mm3/mm2 could a considerable flattening of the diamond grains be discerned in the conventional process, which lead to high process forces. Ultrasonic assisted engaging grains on the other
7
hand are characterised by splintered, sharp cutting edges.
Fig. 5 : Wear on a diamond grain D46 after conventional (right) and ultrasonic assisted (left) grinding
In conventional grinding, bonding material stay behind the abrasive grains. This give additional stability against chipping to the grain. In ultrasonic assisted grinding, this process do not appear, the reason being completely different engagement conditions and higher engagement angles. In accordance with the complete change in engagement conditions in ultrasonic assisted grinding, the removal chips are also significantly different. Figure 6 uses the example of zirconium oxide to illustrate this. In conventional cross-peripheral grinding, the removed chips are marked by plastic deformation processes. Their shape resembles that of conventional removed chips chips from machining of metals. On the other hand, the removal particles in the ultrasonic assisted process are smaller in size. Traces of deformation are hardly discernible. Rapid accelerations and changes of direction of the tool motion prevent continuous separation. Low process temperatures and higher engagement depths resulting from the oscillation reduce the affinity for plastic deformation.
Fig. 6: Zirconium oxide removal chips after ultrasonic (left) and conventional (right) cross-peripheral grinding
The mechanical loads which are concluded from the wear behaviour of the tools lead to a change of properties even in the sub-surface of the workpieces. In the case of surface grinding with axial oscillation it was concluded that there are no disadvantages arising for the residual stresses and strength characteristics (8). After comparing ultrasonic assisted cross-peripheral grinding with plane parallel lapped ceramic workpieces, the conclusion was confirmed (9).
TECHNOLOGICAL INVESTIGATION The focal point of the technological investigations was the machining of advanced ceramics and glass by ultrasonic assisted face grinding and cross-peripheral grinding. It was investigated to what extent the obtainable material removal rates can be increased by ultrasonic assisted of the grinding process. Subject of the investigations was further the question what effect the increased mechanical loads and the high frequency oscillation have on the surface formation and the properties of the machined component. The process forces were measured with a force measuring system by the company Kistler. A subsequent evaluation software was applied to analyse the values. Since the frequency of the oscillation of fils = 20 kHz is very high, a dynamic force measurement was not conducted. Therefore, the force values given below represent a mean stationary value. Concerning economic efficiency, it was ascertained that it is possible to considerably increase the material removal rates compared to former prototypical installations. This is especially the case in ultrasonic assisted face grinding, in which the integrated flushing box permits the supply of cooling lubricant and the transport of removed particles from the contact zone even at higher depths of cut. In contrast to results gained in previous investigations, there was no above-average increase in process forces with increasing depth of cut. Even with poor accessibility to the contact zone from outside, a stable process course is possible. The following figure 7 demonstrates that with increasing feed speed v J ~
the stationary surface related axial forces also rise. Since the effective front of the tool A,$ corresponds to the workpiece surface to be removed A,, the feed speed is equal to the surface-related material removal rate Q,"".The axial forces continue rising during machining, irrespective of the feed speed.
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30
s u f m related material removal rate Q",
reduction of Q ~ U S . , (AUS ,~ = 14 pm) = 32" to G ~ (A(,S / ~= 10 , pm) ~ ~= 24". Another result of these kinematical changes is a reduction of ultrasonic accelerations and consequently a reduction of the pulsed mechanical loads. In this context, the example of machining of aluminium oxide is used in Figure 9 to demonstrate the resultant process course dependent on the material removal rate. Higher amplitudes permit an almost stationary process course after an initial rise of process force which cannot be obtained with smaller amplitudes. This is probably due to the fact that a insufficient amount of sharp edges is generated by grain splintering and the increase of frjction. In the case of ultrasonic assisted face grinding of aluminium oxide, the maximum surface related material removal rates were Qw " > 25 mm3/mm2min. The reduction of process forces appearing in ultrasonic assisted face grinding as opposed to the conventional process, can also be detected in cross peripheral grinding.
Fig. 8: Comparison of the axial force in conventional and ultrasonic assisted face grinding with different amplitudes
Figure 8 displays the influence of the ultrasonic oscillation on the process forces during face grinding. It is discernible that in conventional grinding, high process forces already occur at smaller surface-related material removal rates and smaller feed speeds and do not permit an increase in material removal rates. Due to the superposition with axial ultrasonic vibrations, the process forces are significantly reduced in such a way that surface related material removal rates of Q," ' ' > 15 mm3/mm2minare obtained during a secure process.
4
i p -N-
38
mm*
I
I
I
1I
t
t
t
ts 2 f l iw l 0:-
U
-*
1 workpi
IT7
+ I
1
0
420
l
i1
L:
I 1
D126 D151 grain size Fig. 10: Material removal rate of ultrasonic and conventional crossperipheral grinding as a function of the grain size
046
4
a
I
Figure 10 illustrates the obtainable material removal rate as a function of the cross-peripheral grinding modus and the size of the diamond grain. In this case, the feed speed was controlled by determining a maximum process force in feed direction. It becomes distinct that the difference between ultrasonic assisted and conventional grinding is much more subtle that of die-sinking. This is due to the fact that at the periphery, the abrasive grains that mainly exert the material removal are subject to different loads and engagement conditions than on the front. For surface grinding with axial ultrasonic vibration of the workpiece, ZAPP
describes that the grain is exposed to a continual change in load direction as a result of the oscillation. In connection with those turns during the tool engagement he concludes that these changes lead to a more intense material removal (6). The obtainable surface qualities are greatly dependent on the characteristics of the machined materials and of the machining process. This can be proven by comparing the surface qualities of ceramics, which were machined with ultrasonic assisted grinding and plane parallel lapping. Excluding the material silicon carbide, the arithmetical mean deviation of those samples ground with ultrasonic assistance are slightly above the value established for that of plane parallel lapping.
t iron
12 I
ultrasonic assisted grinding plane parallel lapping
SUMMARY The superposition of grinding kinematics with ultrasonic vibrations opens up new prospects for the machining of ceramics. This article shows considerable advantages for processes with a large workpiece-toolcontact. This was put down to the change in resultant cutting speed ratio during ultrasonic assisted grinding as opposed to conventional grinding. This change also leads to an improved process behaviour. As a result of the high mechanical loads, the diamond grains are subject to micro splintering on the tool and hence fresh sharp edges are formed in the course of the machining process. These wear characteristics entail a stable process behaviour and small process forces, which enable distinctly higher material removal rates than in the conventional process. A significant deterioration of the surface quality or of the sub-surface characteristics was not detected. It was possible to show that smaller process forces can be obtained by increasing the ultrasonic amplitudes. Moreover, these forces show an almost stationary course dependent on the removal rate. In ultrasonic assisted face grinding, the supply of the contact zone with cooling lubricant through the tool and the ultrasonic oscillation guarantee a reproducible process course even at higher depths of cut. The transport of the removed particlesout of the contact zone is enhanced by the high frequent lifting motion of the tool.
ACKNOWLEDGEMENTS Special thanks to the Bundesministerium fir Bildung und Forschung (Federal Ministry of Education and Research) for supporting the research project "Ultrasonic assisted grinding".
REFERENCES Fig. 1 1 : Arithmetical mean deviation depending on the material and the machining process (10)
For the materials silicon carbide and zirconium oxide (ZN 40) the best surface values were measured during ultrasonic assisted grinding. Summarising the above, the obtainable surface qualities of ceramic materials machined with ultrasonic assisted grinding can be established at the level of conventional finishing procedures. Despite higher mechanical loads, surface qualities comparable to those of lapping can be achieved. For the machining parameters which are displayed in figure 11 bending strengths of ceramic workpieces were investigated. It is obtainable that ultrasonic assisted grinding causes bending strengths similar to or higher than those of plan-parallel lapping. Moreover ultrasonic assisted grinding leads to insignificantly higher deviations of bending strengths. These are caused by the alternating load on the workpiece-subsurface (9).
Colwell, L.: The Effects of High-Frequency Vibrations in Grinding. Transactions of ASME, May 1956, p. 124-131. Dam, H.; et al.: Surface Characterization of Ultrasonic Machined Ceramics with Diamond Impregnated Sonotrode. NIST-Special Publication 847 (1993), p. 125-133. Hoia An, V.: Erkundung der Effekte im Arbeitsergebnis beim Ultraschallschwingschleifen. Dissertation, TU Dresden 198 1. Markov, A. I.: Ultrasonic Drilling of Hard NonMetallic Materials with Diamond Tools. Stanki I Instrument, Vol. 48; Issue 9, 1977. Nerubai, M.: Leistungssteigerung beim Schleifen mit Diamant unter Ultraschall (russisch). Stanki I Instrument Vol. 48; Issue 2 (1 977). Zapp, M.: Ultraschallunterstutztes Schleifen von Hochleistungskeramik- Ein Beitrag zur gezielten Beeinflussung der Eigenschaften von Bauteilen
42 1
durch eine ganzheitliche Prozefikettenbetrachtung. Dissertation Univ. Kaiserslautern, 1998. (7) Pei, Z.: Rotary Ultrasonic Machining of Ceramics: Characterization and Extensions. Thesis University of Illinois, USA, 1995. (8) Uhlmann, E.: Surface Formation in Creep Feed Grinding of Advanced Ceramics with and without Ultrasonic Assistance. Annals of the CIRP, Vol. 47/1/1998, p. 249-252.
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(9) Uhlmann, E.; Holl, S.-E.; Daus, N.-A.: Bearbei tungsbedingte Ranchonenbeeinflussung von keramischen Werkstoffen durch das ultraschallunterstUtzte Schleifen. 59. Jahrbuch Schleifen, Honen, Liippen und Polieren, p. 45-57,2000. (10) Spur, G.; et al..: Machining of Com lex Contours by Ultrasonic Assisted Grinding. 3!* International Machining & Grinding Conference October 4-7, 1999, Cincinnati, Ohio, USA, p. 6 1 1-625.
V. Material Design and Process Development
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ADVANCING IN MECHANICAL PROPERTIES OF SILICON NITRIDE: THE ROLES OF STARTING POWDERS AND PROCESSING A. Bellosi”, G.N. Babini
CNR-IRTEC, Research Institute for Ceramics Technology, 48018 Faenza, Italy
ABSTRACT The characteristics of raw Si3N4powders and of the mixtures containing sintering aids were proved to be the key factors in controlling microstructure and properties of dense silicon nitride. The relationships between the surface properties of powders and the interaction with the additives, their homogeneity and distribution were investigated, using two routes to introduce sintering aids: ultrasonication and chemical coprecipitation. Sintering aids (La24+Y203) that originate highly refractory grain boundary phases were employed. Depending on the raw powders, strength varies in the range 500 to 1100 Mpa at R.T, in the range 300 to 800 MPa at 140O0C, toughness varies from 4 to 6 Mpam”. The factors influencing microstructure and properties are discussed.
INTRODUCTION In order to improve the properties of silicon nitride, the key features are the characteristicsof raw powders, the development of controlled powder processing and densscation procedures [ 1-61. The mixing/milling process to introduce additives modifies the powder surface, determines the distributionof the sintering aids and the inter-particle interactions. If the additive is a powder, the finite size of the particles cause difficulties for intimate mixing. Advantages can be gained using chemical methods to produce additives in a finely divided form or to distribute them homogeneously on a nanoscale on the silicon nitride particles, however these methods need to be optimized. The present study aims to asses the role of the powder characteristics on microstructure and properties of silicon nitride, starting from four different Si3N4 powders, and two different powder processing routes: i) ultrasonic mixing: a clean, fast, cheap, reliable method; ii) chemical coprecipitation from nitrates: a complex procedure which offers the potential to control distribution and microchemistry of intergranular phases. As additives, La203+Y203having high solidus temperature and viscosity were used; their amount was kept as low as possible to obtain full dense materials.
EXPERIMENTAL PROCEDURES The four commercial Si3N4 powders were selected on the basis of their production process: nitridation of
silicon (powders M and P), chemical synthesis by liquidhapour phase (powders U and B). Two compositions of powder mixtures were prepared (amounts in wt0/0): Si3N4+2La2Q+2Y203(composition 1) Si3N4+3La203+3Y203 (composition 2) The techniques used to add the sintering aids were the followings (further details in Ref. [3, 41): -ultrasonic mixing (u): separate ultrasonic stirring of additive and Si3N4powders in water, than mixing the two dispersions under magnetic stimng at pH =lo. Ultrasonic stirring by pulsed cycles (10 s) for a total time of 10 min. -chemical coprecipitation from nitrate solutions (c) separate ultrasonic stirring in water of Si3N4powder and of Y- and La-nitrates, mixing of the two dispersion under controlled pH (9.5-10.5) in order to allow metal hydroxides precipitation on S13N4 particles surface. After three stages of sedimentation and washing, the mix was freeze dried and calcined at 400°C for lh under flowing nitrogen. Characteristics of the starting Si3N4powders and dried powder mixtures are summarized, respectively, in Tables I and 11. In order to optimize the homogeneity of the mixtures, the colloidal behaviour is of great importance for further processing steps, therefore Z-potential values and p H ~ p(pH of the isoelectric point) of the starting powders and powder mixtures was determined by Electrokinetic Sonic Amplitude (ESA) measurements. The ESA potentiometric titrations were conducted from natural pH to pH= 11 and from pH= 11 to pH=3, the procedures are described elsewhere [7,8]. The distribution of the additives and crystalline phases of the mixtures were estimated by SEM and Xray diffraction. On raw powders and on the mixture of composition 2, atomic ratios were measured by X P S analyses, details are reported in Ref. [7, 81. All the mixtures were hot pressed under vacuum in an induction heated graphite die using a pressure of 30 MPa and a temperature of 1850°C. The microstructure of the dense materials was analyzed by SEM, EDS and XRD. Mechanical properties were measured with the following methods: -Young’s modulus (E): frequency resonance method on a 0 . 8 ~ 8 . 0 ~ 2mm 8 sample; -the flexural strength (0)in a 4-pt bending fixture, on 2 . 0 ~ 2 . 5 ~ 2mm 5 bars, up to 1400°C in air; -the fracture toughness (KIc): Direct Crack Measurement (DCM) method, with a load of 98.1 N in a hardness tester; -the Vickers microhardness (HV1.0): on polished surfaces with a load of 9.81 N.
425
surface silica. XPS analyses confirm the results of the chemical analysis of oxygen shown in Table I. The presence of different surface species on Si3N4 particles are confirmed by the values of the pHIEp measured by ESA (Table I). The shift of the pHEP is due to differences of surface oxygen content: a more acid surface derived from the presence of surface silica (or a mixture of Si02 and Si3N4)in powder P and B, a more basic surface due to oxynitride film coating in powder U, a surface mainly covered by basic slylamine groups in powder M [7, 8, 10, 111. The chemistry of coating layer drives the particles reactivity: oxynitrides seems to be more reactive than silanoYsilylamineones. The surface characteristics of silicon nitride powders, undergo modification depending on processing route, as shown in Ref. [7]. Also powder U, which has a more homogeneous and slightly oxidized surface, after attrition milling, ultrasonic mixing and ball milling, reveal increasing amount of SiO2, in a direct relationship with processing time. Among the mentioned processing routes, ultrasonication changes surface characteristics less than other routes [7] Therefore, in the present study, it was chosen as an alternative procedure to chemical methods for the addition of sintering aids to silicon nitride powders. The addition of sintering aids involves several interactions between silicon nitride particles and additives like heterocoagulation, segregation, precipitation; in this respect the surface properties of the starting powders are of great importance. These characteristics influence the dispersing behaviour of the suspensions during the addition of the sintering aids, as the specific adsorption of the metal ions influences the solid-liquid interface properties. The two metal oxides ( Y 3 0 and La& ) used as sintering aids
RESULTS AND DISCUSSION Raw Si3N4powders and powder mixes Among the tested Si3N4powders, three of them (U, B, M) have similar specific surface area, higher than that of powder P (Table I). Powders M and P have relatively wide particle size distribution and irregularly shaped (mainly acicular) particles. Powders U and B have approximately round particle shape. Type and amount of impurities depend on the synthesis route of the powders: apart oxygen, which amount is low in powder M, the content of the remaining impurities is about two orders of magnitude higher in powders (P and M) produced by nitridation of silicon than in powders (U and B) produced by chemical synthesis. The presence of surface oxygen is due to a partd surface oxidation of powder particles, because of the thermodynamic instability in air of silicon nitride. Therefore, particle surface is covered by a layer containing mainly amphoteric silanol (Si-OH), basic secondary ammine groups (SiZ=NH) and Si2N20 in different amounts, depending on powder preparation and subsequent treatments. According to the atomic ratios measured by XPS, powder U presents a homogeneous and complete surface layer composed by Si2N20. In powders B, P and M two distinct XPS silicon signal were detected and attributed to Si3N4and Si02 surface sites, that accounts for the presence of a (probably patchwise) oxygen-rich surface [7]. Presumably, as reported specifically for silica samples [9], several kinds of Si-0 species are present at the particle surface, these including also Si-OH bonds. Powder M contains the lower amount of silica. As XPS revealed the presence of fluorine, probably the raw powder was chemically treated in order to reduce
Table I. Characteristics of raw Si3N4powders: s.s.a.: BET, d: equivalent spherical diameter, D: aggregates diameter : from grain size distribution; oxygen amount measured by LECO, impurities from the powder supplier, ~ H E Pisoelectric point (ESA), surface atmic ratios ( X P S ) . dp Si3N4ratio is about 94% in U, B, P and 92% in M.
I s.s.a. - . I
d IDs0 (mLg-') (pm)
U 11.5 B 12.2 P 7.3 M 11.6
I Impurities(wt??) 0
C
C1
lpHIEp Fe
Ca
0.16 0.51 1.09 0.01 0.01 0.005 0.15 0.57 1.44 0.1 0.005 0.0003 0.26 0.77 1.35 0.29 0.06 0.01 0.16 0.2 0.81 0.2 0.01 0.01
Al 0.005 0.0003 0.03 0.08
6.6 4.0 5.1 7.1
1
atomicratio* /Estimated N/Si Si/O N/O Surface sites YO 1.0 2.2 2.2 100 Si2N20 0.7 1.1 0.8 52 Si3N4;48 SO, 1.5 0.5 0.8 60 Si3N4;40 SO, 1.1 3.9 4.2 87 Si3N4; 13 SiO,
Table IL Characteristics of the powder mixtures: s.s.a. (BET); oxygen content (wt%); pHEp: isoelectric point; ApH,: variation of the isoelectricpoint in the coprecipitated mixtures in respect with the starting powders; ApHEpto 11.1: difference in the isoelectric points of the mixture and of the additives (11.1); atomic ratios from X P S analyses. s.s.a.(m2/g) 0 (WtYO)
2uu
2uc
1uu
1uc
11.9
10.9
11.0
12.54
3.4 4.2 +2.31 +3.11 8.1 8.7 (+1.5 (+2.2)
1.76 +0.67
2.23 +1.14 10.9 (+4.2) 0.22
PHIEP ApH=p ApH=p to1 1.1 0.02 Y/Si LdSi 0.007
426
0.30 0.20
present basic surfaces @HIEPare 10.4 and 11.8 respectively), their minimal solubilitieswere calculated to correspond to pH=10.2 and 9.8 respectively [7]. On these basis, regardless of the mixing process, experimental pH for the addition of the oxides was strictly maintained in the range 9.5-10.5, in order to have minimal solubility of Y-. and La-oxides and to fall into the pH range of opposite z potential values. Electroamtic analyses [7, 81 showed that a good “coating” of the silicon nitride particle surface with additive species can be obtained with ultrasonication, the experimental atomic ratios LdSi and Y/Si approximate the theoretical values [7]. The preferential surface enrichment of yttria in respect to lanthania is due to its higher positive z-potential at the selected pH values leading to heterocoagulation phenomena with the particle surface. SEM analyses indicated the presence of randomly distributed aggregates of additives which have low dimensions (<3 pm). Ultrasonication has the advantages to avoid contamination that usually derives from milling media and to perfom the homogenization in a very short time. Regarding the chemical coprecipitation, ESA titration curves of all the mixtures lc (composition 1) highllght the following features: i) the pHmpshifts towards basic values, i.e. towards those of the additive oxides; pHm values of the mixtures close to those of the additives indicate a homogeneous coating without segregation of precipitated hydroxides: the Si3N4 particle surface behaves as a surface coated with Y(OH)3 and La(Om3 ii) the difference in the pHmP values between the starting silicon nitride powder and the relative mixtures (Table 11) gives an estimation of the heterocoagulation and coprecipitation processes and a relative comparison of the raw powder surface reactivity. The experimental atomic ratios by X P S analyses (Y/Si and LdSi) (Table 2 mixture 2U) appear to be at least one order of magnitude larger in coprecipitated powders than in ultrasonicated ones. As in the previous case, yttrium appears more enriched at the surface in comparison to lanthanum due to better specific adsorption (heterocoagulation) of Y203than La203for the reasons above explained. It is possible to establish a merit rating concerning the surface reactivity of each silicon nitride powder, in respect with an ideal configuration: -powder B exhibits a surface composed mainly by silica that favours a large number of scarcely reactive nucleation sites. It implies a weak but homogenous growth, which induces the formation of a film with a high screening effect in respect to the substrate, as suggested by the high ApHmp.-Powder U: the homogeneous surface layer of Si2N20on the particles gives rise to a rapid nucleation of Y and La oxides with few nuclei but a large growth. It provides a smooth coating and slight change in the surface area. The morphology and smoothness of this coating was observed by TEM analyses [12]. -Powder P: features like high surface concentration of silanol groups with low initial reactivity, large grain size distribution, low specific surface area and irregular particle shape favour inhomogeneous heterocoagulation of the additive species on particle surface that results in the worst screening effect in
respect to the substrate.-Powder M the low content of silanol groups and the presence of sylilamine groups (on the basis of the pHm values) should imply a scarce surface reactivity, that results in a minimal heterocoagulation phenomenon of the additive species and inhomogeneous surface coating of Si3N4particles. Concerning powders P and M, it can be hypothesized the formation of clusters of amorphous additive precipitates, consisting of clumps of nanosized, near round shaped particles, which are supposed to be the first cause of the observed increase in specific surface area (Table I and 11). As crystalline phases, peaks due to additive phases (La203,Y203and some La(OW3) were detected in the mixed powders. In addition, the coprecipitated powders present an higher oxygen content than the ultrasonicated ones, due to calcination (Table 11). The results confirm that morphology and distribution of the additive coating on the surface of the core particles depend on precipitation and heterocoagulation mechanisms, which are related to surface chemistry, size and distribution of core and precipitated parbcles, amount of the additive precursors, nature of the other ion species in solution, kinetics of precipitation, Ph, viscosity): their influence is not yet completely clear and the optimization of the chemical processes needs to be further investigated. Microstructure of hot pressed materials SEM micrographs on plasma etched surfaces reveal the microstructure of materials hot pressed with different Si3N4nitride powders (Fig. la-d) and different powder processing route (an example in Fig 2a,b). The values of the mean grain size and of aspect ratio are reported in Table 111. The microstructure presents in any case elongated grains mixed with smaller and equiaxed grains, although the features of these morphologies are different in the various materials. Only materials from powder U reached full density, porosity and defects are absent, instead for all the other samples some pores, microcracks, porous areas are observed (examples are shown in Fig. lc,d). Morphological features, porosity, defects, type and amount of grain boundary phases depend on starting powders characteristics and are related to the mechanisms involved in densificatiodgraingrowth. The use of powder U, both processed by ultrasonication or chemical coprecipitation, offers the best chance to give origin to fine and uniform “in situ” composite microstructure, the highest aspect ratio and the lower grain size, absence of defects. The increased amount of additives in composition 2 in comparison to composition 1 favours the development of very fine grains with aspect ratios up to about -10, particularly in the sample obtained by chemical coprecipitation. Materials from powder B, although the high h a l density, reveal low aspect ratio and some inhomogeneities, which features depend on the powder processing route. E n h a n d grain growth is observed in sample obtained by coprecipitated powders (Fig. 2 b), while microcracks associated to inhomogeneities are present in samples from ultrasonicated powders.
427
Figure 1. Microstructure of hot pressed materials, on the surface perpendicular to the applied pressure, from ultrasonicated powder mixtures, with the Si3N4powders U, B, M, P.
Figure 2. Microstructure of silicon nitride hot pressed from the same powder (M), processed by : a) ultrasonication, b) chemical coprecipitation
42 8
Powder P showed a very different behaviour depending on the processing route. As above described, the coprecipitation did not cause a homogeneous coating of particles with additives. This factor, associated to the high content of impurities and to the large particle size distribution in the initial powder, resulted in a relatively low liquidus temperature, which favoured the growth of grains with low aspect ratio and an inhomogeneous final microstructure (Fig. 2c). The final density is rather low, particularly with ultrasonicated powders: the wide particle size distribution hindered the homogeneous distribution of the additives and, consequently of the liquid phase during sintering. Pores and areas with scarce amount of liquid phase were found. Samples obtained from powder M, particularly the coprecipitated mixtures, reached a high final density. The microstructure is quite coarse (grains up to 5-8 pm were observed) but no evident defects were detected (Fig. 2d). Nevertheless, in the correspondence of coarse grains, some cracks were observed, originated by stress accumulated at their neigborough, due to the impingement of adjacent growing grains. All the above described features depend on the additive system and the variation of the liquidus temperature (Table 111) originated by the presence of impurities and oxygen during sintering. In fact the liquidus temperature (estimated by densification curves [12] is up to 2OOOC lower with powder U and P than with powder U. In all the cases the relatively high oxygen content in chemically processed mixture causes a relatively lower eutectic temperature in comparison to those of ultrasonicated mixtures. As the viscosity of the liquid phase at the sintering temperature is higher for compositions with high liquidus temperature, the diffusion-controlled processes: i.e. densification and grain growth, are different for the various samples because the diffusion coefficients are inversely related to the viscosity of the liquid phase. An inverse relationship between the liquidus temperature (i.e. viscosity) and the mean grain size can be drawn. Moreover, generally, an higher viscosity limits mass transport [l, 131, this results in retarded grain growth and in the evolution of finer and fibrous microstructure: it was experimentally confirmed by the
Tl1q "C
Pf %
d pm
a
E GPa
Hv GPa
behaviour of samples from powder U containing two different amounts of additives. Composition 2 in fact had the higher liquidus temperature and the lower mean grain size, associated to the highest aspect ratio. In the hot pressed materials, grain boundary phase is amorphous except than in samples from powder B, where traces of Y-La-silicates were identified. This derives from to the higher amount of oxygen in the starting powders. After annealing tests at 1400°C for 6 hours, amounts from 2 to 6 vol% of crystalline phases (La-Y-Si-0-N and La-Si-0 phases with various stoichiometries) were revealed. Two devitrification humps generally occurred during thermal expansion tests: the former in the temperature range 1000-1200OC and the latter 1250-14OO0C, being attributable respectively to the devitrification of silicates and of oxynitrides. It is well known [14] that the substitution of nitrogen for oxygen in the glass network produces a more rigid network increasing the viscosity of the glass. The transition temperature of these glasses increases with nitrogen content and with the type of the network modifier: glasses in the system La-Y-Si-0-N have transition temperatures higher of about 200°C in respect with other oxynitrides [14, 151.
Mechanical properties (Table 111) Young's modulus. It is mainly influenced by porosity, as confirmed by the low value measured on samples P. Values (325-329 GPa) measured in dense samples are higher than those found for materials produced with other sintering aids due to the presence of stiff grain boundary phases [4]. In sample B the presence of silicates in the intergranular phases lowers E. Hardness. Residual porosity, large mean gain size and the presence of microstructural defects are responsible for the low hardness measured on samples P and M. Toughness and crack propagation. The microstructure with elongated and textured P-Si3N4grains allows a toughening effect from a combination of factors like amount of elongated grams, aspect ratio and grain size. In addition, an influence of porosity has not to be excluded. The effect of microstructure is evidenced from tests in two dlrections (N and I to applied pressure during hot pressing, Table III) [4].
KIC
r3
(ma)
madm
429
Crack propagation partially follows the grain boundaries but, to a greater extent, cracks run through the grains. Crack deflection is more frequently observed in samples with aspect ratios higher than 7 andor is highlighted in materials with the largest grain size. Room temperature flexural strength. Striking differences were found: excellent strength values (9001200 MPa) were measured only in samples produced with powder U. No remarkable difference is due to the powder processing route. It means that, when using a using a fine and pure Si3N4powder, suitable to favour homogeneous dispersion of additives, the defect population is not so influenced by processing. Also with the use of the other silicon nitride powders, the feature which mainly affects the presence of critical defects is associated to the characteristics of the starting silicon nitride powders, because.they strongly determine the homogeneity of the additive distribution and the microstructural evolution. Among the nearly fully dense samples, the lowest strength regards powder B, that, although “fine and pure” has particles with acid surfaces, which impedes the uniform surface coating with the additives. The observed defects, in fact, are due to the poor homogeneity of the grain boundary phases with consequent pores and cracks. In samples from powders P and M, the presence of large particles in the starting powder is the main reason for the development of microstructural defects like exaggerated grains associated to poorly sintered areas and for the scarce homogeneity in the distribution of the grain boundary phase. High temperature flexural strength. In addition to phenomena such as preexisting defects, also relaxation of machining stress and other effects take place, the most important are the softening of the grain boundary phases and the introduction of oxidation pits. Materials from powder U, particularly with composition 2, showed remarkable performances up to 140OOC (up to about 800 MPa). Constant strength was observed from 1000 to 140OOC for sample 1Uc and from R.T to 12OO0C for samples 2Uc and 2Uu, followed by a decrease at 140OOC of only 18%. It confirms that the selected additive system originated highly refractory grain boundary phases, as above mentioned. Samples from powder B maintain good strength level up to 1400OC. In materials produced with powder P, and particularly M, a relevant strength decrease is detected at 12OO0C, where the values are about half of those measured in samples from powder U. The strength decrease is related to the relatively low refractoriness of grain boundary phase, that is affected by impurities in the starting powder. Another important factor is the distribution of grain boundary phases. The coarser is the microstructure, as in samples P and M, the larger are the glassy pockets at triple grains junctions and thicker is the glassy phase between the grains: failure can originate there and than advance along the grain boundary layer. The results make it clear that the “quality” of the raw Si3N4powder is the main key factor for the production of high performance silicon nitride, but, in addition, a 430
severe control of the powder treatments and the design of sintering aid systems are prerequisites.
REFERENCES (1) G. Wotting and G. Ziegler, Influence of Powder Properties and Processing Conditions on Microstructure and Mechanical Properties of Sintered Si3N4,Ceram. Intern., 10, (1984) 18-22. (2) H.J. Kleebe and G. Ziegler, Influence of Crystalline Secondary Phases on Densification of Reaction Bonded Silicon Nitride During Post Sintering Under Increasing Nitrogen Pressure, J. Am. C-. SOC.,72, (1989) 2314-2317. (3) C. Galassi, V. Biasini and A. Bellosi, Effects of Powder Characteristics and Mixing Processes on the Microstructure and Properties of Silicon Nitride, Proces. of Adv. Mater., 3, (1993) 153-161. (4) A. Bellosi, F. Monteverde, G.N. Babini, Influence of Powder Treatment Methods on Sintering, Microshucture And Properties of Si3N4-Based Materials, in Engineering Ceramics ’96, ,NATO AS1 Series Vol. 25, Kluwer Acad. Publ., Dordrecth, 1997, 197-212. ( 5 ) T. M. Shaw and B. A. Pethica, Preparation and Sintering of Homogeneous Silicon Nitride Green Compacts, J. Am. Ceram. Soc,62, (1986) 88-93. (6) H. J. Kleebe, W. Braue, H. Schmidt, G. Pezzotti, G. Ziegler, Transmission Electron Microscopy of Microstructure In Ceramic Materials, J. Europ. Ceram Soc.,16, (19%) 339-351. (7) F. Bertoni, C. Galassi, S. Ardizzone, C.L. Bianchi “Water-based Si3N4 Suspensions: Effect of Processing Routes on The Raw Materials, J. Mat. Res. 15, (2000) 155-163. (8) F. Bertoni, C. Galassi, S. Ardizzone, C.L. Bianchi, Surface Modification of Si3N4 Powders by Coprecipitation of Sintering Aids, J. Am. G r a m SOC.82, (1999) 2653-2659. (9) J. F. Moulder, W. F. Stickle, K. D. Bompen, Handbook of X-Ray Photoelectron Spectroscopy, Perkin Elmer, Eden Praire (1991). (10) S.M. Castanho, J.L.G. Fierro, R. Moreno, Surface Oxidation of Si3N4 Green Compacts Effect of Sintering Aids, J. Europ. Ceram SOC. 17, (1997) 383-392. (1l)E. Lyden, L. Bergstrom, M. Persson, R. Carlsson, Surface Modification and Dispersion of Silicon Nitride and Silicon Carbide Powders, J. Europ. Ceram. Soc.7, (1991), 361-368. (12)A. Bellosi, A. Bondanini, Influence of Powder Characteristics and Processing on Properties of Silicon Nitride, submitted to Mat. Sci. Eng. (13)s. Hampshire, Nitride Ceramics, in Materials Science and Technology, Vol. 11, VCH Weinheim, 1994, pp.119-172. (14)J. Chen, P. Wei, Y. Huang, Properties of La-Y-N0 glasses, J. Mat. Sci Lett. (1977) 1486-1488. (15)s. Hampshire, R. A. L. Drew, K. H. Jack, Viscosities, Glass Transition temperatures and Microhardness of Y-Si-Al-0-N glasses. Comm. Am. Ceram. Soc.73, (1990) 117-135.
MATERIALS DESIGN OF COMPOSITE MATERIALS - COMPATIBILITY OF SELF-DAMAGE MONITORING AND STRENGTHENING H. Yanagida
Research Institute, Japan Fine Ceramics Center 2-4-1 Mutsuno, Atsuta-ku, Nagoya-city, 456-8587 Japan
ABSTRACT Brevity in technology is the keyword of next centurylmillennium. We have to avoid spaghetti syndromes of modern technology. Authorities have to find novel concepts to make technology more easily understood to avoid unnecessary confusion. Integration is another important key word. Improvement of structural reliability and capability of self-damage monitoring may be achieved if we design so beforehand. This opens a new field of materials design in composite materials. Typical case is a concrete structure reinforced by CFGFRP (carbon fiber and glass fiber co-reinforced plastic) bars. This may be extended to diagnose damage of fragile materials such as ceramics used for heat engine component and concrete piles applying fiber reinforcement and percolation phenomena of conducting powders. R & D is conducted by Kenmaterials Research Consortium led by Professor Yanagida. Technology must be cooperated by willing citizens. Yanagida recently established NGO, Port for Techno-Democracy.
by INAX where humidity and temperature adjusting capability like soil and mechanical strength like ceramics are achieved to save energy during fabrication and during actual duties, electric cable trolleys for Shin K m e n able to monitor damage without additional sensors by JR Tokai, etc.
KEN-MATERIALS RESEARCH CONSORTIUM Researches co-operated between industries and academia are made accelerated if knowledge is spread easily among groups through agreements. Y anagida lab of University of Tokyo used to have several liaison contracts, among them Yanagida considered often it is necessary to transfer knowledge obtained by a contract A to another contract B to solve problem laid in the contract. Sometimes it is required to transfer knowledge from B to A. Thus a research consortium is established and expanded. The main concepts of the consortium are to pursue simplicity, to integrate structural reliability and useful functions, to make people willingly commit and conscious to environmental issues. The logo-mark is shown in Figure 1.
INTRODUCTION R & D of so-called intelligent materials is not only performed upon academic basis but also proceeded by industrial sector. In February 1994 some industrial companies gathered to form research consortium for R & D of so-called intelligent materials. H.Yanagida was appointed as the chairman of the consortium. Through discussion about the concept and naming of the consortium, the consortium has been called as "Ken-materials Research Consortium". Objectives of R & D of kenmaterials are intensively concerned with safety, environment and friendliness to people. Ken stems from the pronunciation of some Chinese characters meaning wisdom (self-control), sensing (functional capability), structure (reliability), integration, simplicity, soundness (supported by people) and ecology. Typical examples are CFGFRP=carbonfiber and glass fiber co-reinforced plastics with improvement of mechanical performance and capability of self-monitoring of damage developed by Yanagida et al, soil-ceramics or earth-ceramics
Structure
Soundness
-
Sensing
Brevity
Simplicity
<
Philosophy
Method
Figure 1. Logo-mark of Ken-materials Research Consortium
SIGNIFICANCE
OF
ENTERING
NEW
MILLENNIUM We discuss the future based upon knowledge about the past. It is significant to discuss the future considering a term of thousand years.
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The Ceramic Society of Japan and The American Ceramic Society recently celebrated the centennial anniversary. Establishment of the professional societies means birth and unification of ceramists. Professional groups like ceramists, alchemists and scientists are recognized in a term of hundred years. The term of hundred years means the birth of professional groups suffixed -ists. Back to the order of 1000 years, we had a very excellent lady named Murasuki-shikibu to write The tales of Genji'. She was not a professional novelist. The oldest earthenware ever found was excavated from Kanita-town of Aomori prefecture, Japan and it is considered only around 15 thousand years ago. Among them we can still notice some earthen ware is reinforced by natural fibers. History of civilization is the history of exploitation.Imgation caused enrichment of salinity in soil. Rapid economic growth was made possible by rapid consumption of underground resources. So-called advanced technology made general public alienated. Upon reconciliation of these facts we have to redesign scheme of technology of millennium scale. Ten thousand years scale corresponds to regression from earthenware to stoneware which is an important ancestor of ceramics. Age of the oldest stoneware by excavation recently goes back rapidly to the origin of human beings, presently supposed 2.6 million years ago in Ethiopia. Yanagida proposes a hypothesis that ape became humanbeings with development of stoneware. By the era scale of millennium ancestors of modem ceramics, stoneware and earthenware at the first appeared together. Most of the people fabricated stoneware and earthenware for themselves without being specialized as professionals. They are supposed to have committed themselves willingly and enjoyed actions of fabrication and the nature surroundmg them. No serious environmental troubles did not appear in the millennium scale years ago. Important issues of R & D for the future are, therefore, paying attention to environmental impact, people's willingness to technology, recognition of roles of ancestors of ceramics.
DIRECTIONS OF TECHNOLOGY, MINIATURIZATION, ENLARGEMENT, INTEGRATION AND BREVITY Y anagida proposes that directions of R & D in technology are miniaturization, enlargement of structures, integration and pursuing brevity. It may look that they conflict each other. However, they are based upon common concepts. Miniaturization of devices saves space, resources and energy. So-called advanced technology has been and still is being
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developed along this direction. Yanagida has been directing a National Project 'Frontier Ceramics' sponsored by STA Japan. Large scale structures require much energy and resources. Development of technology related to large scale structure, therefore, is very important for saving energy and resources. Maintenance technology has also to be developed to save resources. Key-word to technologies related to large scale structures is simplification and plainness. Yanagida and hs group have developed technology where strengthening and damage monitoring capability are fulfilled simultaneously by only one action. Simultaneous fulfillment of more than two requirements by one action leads to concept of integration. Integration makes technology simpler and plainer. Ken-materials research consortium presided by Yanagida is trying to develop plain technology whch is well accepted by general public. As the logo-mark shows in Figure 1, key words are 7 Chinese characters pronounced as 'ken' in Japan. Integration of structural reliability and capability of self-damage monitoring is said to have opened new dimension in R & D of composite materials.
EXAMPLES OF KEN-MATERIALS Yanagida has proposed a typical case for selfdamage monitoring capability and improvement of mechanical performance in CFGFRP, carbon fibers and glass fibers co-reinforced plastics[l,2,3]. Shmidzu Co. Ltd. and SOK Co. Ltd collaborated hlm. If we start with FRP, glass fiber reinforced plastics, an action of mixing with carbon fibers improves elastic modulus and gives rise to capability of damage monitoring by checking change in electric conductance. If we start, on the other hand, with CFRP, carbon fiber reinforced plastics, the action of mixing glass fibers improves stiffness to avoid sudden fracture. Figure 2 shows relation between load and deformation. CFGFRP has been developed as an alternative of steel bars to reinforce concrete structure to avoid troubles arising from erosion of steel. Yanagida has reached the material from a completely different view point to design strengthening and capability of self-damage monitoring simultaneously. In CFRP w e can measure electric conductance. However, the loss of electrical conductance there means fatal fracture of the material. The change in electric conductance before the carbon fibers undergo fracture is very small to detect. Mixture with glass fibers makes possible the material stand further behind the point the carbon fibers suffer fracture. A remarkable change is observed whle the material remains still unfractured. The decrease in conductance corresponds the portion of fractured carbon fibers. Life detection is possible by measuring the change. Latent damage after an earthquake of large scale structures such as bridges, highways, buildings
carbon powder has the capability to estimate small crack formation and loading hstory in concrete materials. Micro cracks of the order of 0.01 CTO in concrete structure is easily detected. This method can be applied to the cases to moiiitor damage of the structures under high temperature, in deep water or ground, or exposure to strong irradiation. 700
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Figure 2. LoadlStrain Relation (Steel, Carbon Fiber, Glass Fiber, CFGF Composite) can be monitored by measuring the change in conductance includmg carbon fibers. Damage by invaders through the shell walls of safety chamber can be delayed and monitored by applying the CFGFRP to reinforce the walls. Mixture of carbon fibers with FRP enables not only improvement of mechanical performance of the material but also monitoring of damage. No additional sensors to monitor damage is added here. This is the typical case of integration and integrated material. Breakdown of carbon fiber usually takes place when the fiber is elongated around 1.0%.In order to check damage due to strain less than 1.0%, change in percolation of carbon powders around the fibers to reinforce ceramic matrix is developed by Kenmaterials group of Japan Fine Ceramics Center[4,5]. The group successfully has designed and fabricated the GFRP (Glass fiber Reinforced Plastics ) and CMC (Ceramic Matrix Composites ) materials which posses the function of fracture detection by introducing electrical conductive phases in ceramic or plastic matrix composites. The percolation structure of conductive particles (carbon and TIN) in matrix of the composites enables very sensitive detection in a small stain range. The composite materials with fracture detection by electrical conductivity can bring many industrial applications concerning self diagnosis for deformation or damage in materials. Figure 3 is load - and AR - displacement curves of the concrete specimen reinforced by the GFRP with carbon powder. It can be seen that the change of AR is corresponding the load change owing to the crack formation in the concrete specimen. the residual electrical resistance was observed in the specimen after the loading unloading test, which means that the GPRP with
Displacement I rnm
Figure 3. Load and electrical resistance change as function of displacementin the bending test of concrete specimen reinforced by GFRP-C powder
ENVIRONMENT RELATED MATERIALS Soil-ceramicswas first developed to use industrial waste from tile manufacturing factories. If the waste is heated up to high temperature, energy consumption becomes too much. Ishida of INAS treated the waste with natural fibers hydrothermally around 150 'C to acheve enough strength as wall or floor materials. Mixture of natural fibers used to be popular in earthen ware of Jomon-era. The materials hydrothermally treated still have enough porosity to adsorb and desorb water. The materials can adjust humidity and temperature mildly. Air conditioners usually require help of humidifier when used for warming or dehumidifier when used for cooling. The soil-ceramics make technology simple and comfortable. Technology used for the material is very old and very new. Ceramic materials used to defined as inorganic materials fabricated by heat treatments. This inay not be true hereafter. Unheated ceramics such as soil with lime and MgO-C bricks may become popular. Electric trolley for Shin Kcrirsrn consists of core and ring separated by insulator. \.Tihen erosion of ring reaches the core, collector blade connects electrical circuit between core and ring. This is very simple atid reliable evidence of erosion of trolley.
AVOIDING SPAGHETTI SYNDROME OF TECHNOLOGY Modern technology is sometiines too much complicated to understand. Yariagida defines such ii complicated technology as suffering spaghetti syndrome. There are five symptoms in the synclroine.
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The first is a complicated appearance dfficult to understand. The second is a method to apply more components to solve problems. The third is one's values to consider the more complicated the more advanced. The fourth is paying more attention to trivial things overlooking essential matters. Finally technology becomes degraded by making more complicated. This in not only true for technology itself but also leads to distrust from people. To estimate degree of intelligence Yanagida proposes an equation. He defines 'Wisdom index' as: WI = ( Necessary Merits ) I ( The number of Components )n where n is greater than 1. In the case of CFGFRP, merits in numerator are strength as steel, chemical stability better than steel, two-step wise fracture as steel, lighter weight than steel, damage monitoring capability. Only carbon fibers and glass fibers embedded in plastics are used in dominator. Sometimes one suffering from spaghetti syndrome boasts of large value of dominator instead. It is possible to design compatibility of damage monitoring and strengthening in ceramics if the materials are reinforced by stiff fibers where some of them are surrounded by conductive powders.
REFERENCES (1) N. Muto, H. Yanagida, T. Nakatsuji, M. Sugita, Y. Ohtsuka and Y. Arai, Design of intelligent materials with self-diagnosing function for preventing fatal fracture, Smart Mater. Struct. 1, (1992) 324-329. (2) H. Yanagida, N. Muto, T. Nakatsuji and M. Sugita, Materials design for explicit warning of fracturing, JSME International Journal Series A, 36, (1993). (3) N. Muto, H. Yanagida, T. Nakatsuji, M. Sugita and Y. Ohtsuka, J. Am. Ceram. SOC. 76, (1993)875-79. (4) M. Takada, S.G. Shin, H. Matsubara and H. Y auasda, Fracture Detection of Fiber Reinforced Composites Using Electrical Conductivity. UKJapan Seminar on Intelligent Materials, March, 1996. (5) H. Yanagida, H. Matsubara and N. Muto, Compatibility of toughening and micro-crack detection in brittle materials, Composites at Lake Louise '99, October-November, 1999.
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THE DESIGN OF COMPOSITION AND MECHANICAL PROPERTIES OF a-p SiAlON CERAMICS DENSIFIED WITH HIGHER ATOMIC NUMBER RARE EARTHS E. Dolekqekiq*, H. Mandal*, R. Oberacker** & M. J. Hoffmann** (*) Anadolu University, Department of Ceramic Engineering, Eskisehir, Turkey (**) University of Karlsruhe, Institute for Ceramics in Mechanical Engineering, Karlsruhe,
Germany.
ABSTRACT In addition to the conventionally used yttrium oxide, many studies have been carried out in recent years using rare earth oxides as densification additives for a-P sialon ceramics. Although densification is much easier with these additives, it is difficult to achieve the designed composition and therefore mechanical properties after sintering. This behavior has been explained in terms of the stability region of a-sialon, which decreases as the cation size of rare earth increases. As a result, the use of these oxides as sintering additive is limited. In the present work, a-P sialon starting composition with m=0.8 and n=1.7, which was designed to produce %80 a- and 20% p-sialon, has been densified by capsul free sinter-HIPing using Ln,O,, where Ln= Nd, Yb and equimolar mixture of Nd-Yb. The effects of different rare earth oxide additives to achieve the designed composition and properties have been investigated. It has been found that the designed compositions could not be achieved with single dopants as expected. However, by using multi cation doped a-sialon materials, the designed composition and properties were easily achieved.
INTRODUCTION Because of many excellent inherent properties, silicon nitride based materials have obtained great attention as an engineering ceramics. One commercially succesfull application has been the use of sialon ceramics as a cutting tool material for machining metals, in particular when cutting cast iron or nickel based alloys. The high toughness and hardness values of silicon nitride based materials makes them especially suitable for intermittent cutting at high speeds, depth of cuts or feed rates. Mixed hvo-phase ( a and p) sialon ceramics offer possibilities of tailoring the microstructure designing by a starting composition, for example equiaxed a-sialon
grains can be matched with elongated P-sialon grains to form a toughened composite, and consequently the properties of the final product can be optimized combining the high fracture toughness of the p-sialon with the good hardness of a-sialon [ 1,2]. A distinguishing feature of the a-p sialon system from the silicon nitride system is that a+ sialon phase transformation is fully reversible. Because two phases have characteristic morphologies, the mechanical properties of these materials, therefore, can be controlled by the a : p sialon ratio and , in turn, by production conditions. The phase composition is controlled by sintering additives. Because the a-sialon phase can accommodate metal oxide additives, whereas, p sialon phase does not, metal oxides are rejected to the intergranular regions during the a+ sialon phase transformation. Because of the a+$ sialon phase transformation, some properties of the sialons may deviate from the designed value. Optimization of phase content and the microstructure of sialon ceramics is possible by heat treatments after sintering. A range of hardness, strength, and toughness values can be obtained from a single starting composition by heat treatment method [3]. Although this method may open new routes for tailoring the microstructures and controlling mechanical properties for the applications below -1000 C", many applications for a-p sialon ceramics are in the range of 1000-1300 C" and a+p sialon phase transformation can continue during use of the materials with consequent continuous change in properties. Design of the composition and properties of a-p sialon ceramics is difficult due to this transformation. New compositional and processing techniques are needed. In the present work, a-p sialon starting composition with m=O.8 and n=1.7, which was designed to produce %80 a- and 20% p-sialon, has been densified by capsul free sinter-HIPing using Ln,O,, where Ln= Nd, Yb and equimolar mixture of Nd-Yb. The effects of different rare earth oxide additives to achieve the designed composition and properties have been investigated.
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Sintering
EXPERIMENTAL PROCEDURE
Density
Additive Starting powder mixtures were prepared by using Si,N, (UBE-ElO), A1N (HC Starck- Berlin, Grade C), A1,0, (Alcoa, Grade A1 6SG) together with M203where (M= Nd, Yb and Nd-Yb) (99.9 % Aldrich Rare Earth Products) as the densifying additives to give a-/3sialon composition where m=0.8 and n= 1.7 (see Figure 1). The added metal oxides were calcined at 800°C for 4 hours before use, to remove the absorbed water. When calculating the compositions, 1.4% 0 and 1.6% 0 (according to the manufacturer's specifications) present on the surfaces of Si,N, and AIN respectively were taken in account. Powder mixtures were prepared by wet milling under ethanol in a planetory ball mill for 4 hours using a sialon jar and silicon nitride balls. The size of the prepared batches were 50 grams. The milled slurries were dried at 40°C in rotary evaporator, resulting in weak agglomerates which easily could be passed through a 125 pm sieve. Specimens were compacted into pellets by pressing uniaxially with 30 MPa and then isostatically pressed under 400 MPa. The green pellets were capsul free sinter-HIPed in an ASEA QIH6 HIP-unit at 1800 "C (1 hour at 1 MPa and 2 hours at 50 MPa argon gas pressure). Product phases were characterized by X- ray diffraction using a diffractometer (Rigaku Rint-2000 series). Polished surfaces of sintered samples were examined in Scanning Electron Microscope (CAMSCAN S4) after gold coating by using back scattered mode with an Energy Dispersive X-ray analyzer. Hardness measurements were carried out at room temperature using a Vickers diamond indenter with 9.8 N (1 kg) load.
\ W."
Figure 1. a-sialon plane showing the starting composition used in this study. The a-sialon forming region for Y as mapped out by Sun et a1 [6] is marked by the grey line.
I
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Yb203-Nd203
3.34
I
32% a-SiAlON 68% B-SiAlON 76% a-SiAlON 24% P-SiAlON 84% a-SiAlON 16% P-SiAlON
Table 1. Density measurements and XRD results of each samples densified with sinter-HIPing technique. In all samples, a and p sialons are the predominant phases and no other crystalline phase(s) appeared in the grain boundaries while cooling. The relative amounts of a and P-sialon phases after sintering were established by XRD and the results are also presented in Table 1. Although starting compositions for all samples aimed to produce 80% a-sialon and 20% p-sialon, Yb203doped sample gave results close to the designed composition, while Nd20, doped sample gave only -30% a-sialon and -65% a-sialon. Sample, which contains equimolar mixture of Nd203and Yb203as the sintering additives also reached the designed value. These differences in the content of a-sialon phase may be explained by a++ sialon phase transformation. There is a a+ sialon phase transformation during the cooling step of sintering schedule. The amount of a+ phase transformation is directly related with the radius and the valancy value of the doped cations. There is a relationship between the a-sialon contents of the samples and cation radius of rare earths. So, the tendency for a-sialon formation is decreased with the increasing ionic radius of rare earth cation. The stability region of a-sialon ceramics is related with the size of the dopant and the temperature. Increasing the size of the rare earth and decreasing the temperature leads to an reduced solubility of A1 and 0 in the asialon [4] Therefore, under the same sintering conditions, the stability region of Yb a-sialon is larger than Nd-a-sialon. The ratio of a:(a+p) is plotted against the cation radius (see Figure 2).
RESULTS AND DISCUSSION Table 1 gives densities of samples after sinterHIPing. Examination of polished cross-sections of the sinter-HIPed sialon ceramics by SEM showed only very few micro-pores to be present, indicating that the samples had reached virtually theoretical density after sinter-HIPing.
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0.85
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Figure 2. Linear relationship between the cation radius and a-sialon content of the samples (the cation size of equimolar mixture of Nd-Yb is calculated by taking the average of Nd and Yb cations).
The Vickers hardness (HV1) of sintered samples is shown in Figure 3. The hardness increases in proportion to the amount of a-sialon present, therefore, the hardness of neodymium sample was low, while the Yb203 and (Nd, Yb),03 doped systems showed relatively higher hardness values. As same as the asialon content, the sample, which contains equimolar mixture of Nd,O,-Yb,O, as sintering additive, gave higher hardness value than the expected. 2050
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Figure 3. Linear relationship between the cation radius and hardness (HV1) of the samples. SEM micrographs of the samples also support these XRD results. Typical back-scattered SEM images of the sinter-HIPed and polished surfaces are given in Figure 4. Because the contrast depends on mean atomic number, micrographs very clearly distinguish between phases included: p grains (which contain no sintering additive cation) are black and more needlelike, whereas the a-sialon grains (which contain a small amount of sintering additive cations) are grey and more equiaxed whilst the cation rich glassy phases appear white, because of the high cation content. The microstructure is homogeneous everywhere in the cross-section of the sintered samples. Most of the densifying cations (except Nd”) are incorporated into a-sialon structure and therefore almost glass free grain boundaries are shown. It can be clearly seen from the micrographs that the amount of intergranular phase decreases with increasing a-sialon content. Therefore, intergranular phase in samples densified with Yb,03 and (Nd, Yb),03 additives are considerably less than Nd203 doped composition. SEM micrographs of sintered samples also show that, Nd,O, and equimolar mixture of (Nd, Yb),O, doped samples both had p-sialon grains which had high aspect ratios. These type of grain microstructures increase the toughness value of ceramics. Many of the results obtained in the present study are consistent with conclusions previously deduced for rare earth stabilized a-sialon materials except the sample sintered with equimolar mixtures of Nd,O,Yb,O,, which gave higher amount of a-sialon and therefore higher hardness value than expected. This behavior can be explained in detail in reference [ 5 ] that the increased size of the single phase a-sialon phase field in the ytterbium sialon system.
(c>
Figure 4 Back-scattered SEM micrographs of the (a)Nd,O,, (b)n203, (c)(Nd-Yb),O,, added samples after sinter-HIPing.
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CONCLUSION Dublex a-p sialon materials using three different rare earth oxides have been explored. According to the results presented here, it was very difficult to obtain the desired composition, more so with certain rare earths than with others. This can be explained in terms of the ionic radius of the rare earth cations and the increased tendency for a-sialon formation as the ionic radius of the rare earth cation decreases. Therefore, for Yb,O, addition, nearly all the ytterbium atoms go in to the asialon structure while for Nd,03 addition, very few neodymium atoms go in to the a-sialon structure. However, by adding Yb3+ cation into Nd a-sialon system, the amount of a-sialon phase and the hardness value were increased.
REFERENCES T. Ekstrom & I. Ingelstrom, “Chracterization and Properties of Sialon Ceramics”, In Proceeding of the International Conference on Non-Oxide Technical and Engineering Ceramics, ed. S. Hampshire. Elsevier Applied Science Publishers, London, 1986, pp. 23 1. G. Z. Cao, R. Metselaar & G. Ziegler, “Microstructure and Properties of Mixed a-P Salons”, In the International Symposium on Ceramic Materials and Components for Engines, ed. E. Carlston, T. Johansson & L. Kahlman, Elsevier Applied Science Publishers, London, 1992, p. 188. T. Ekstriim &M. Nygren, “SiAlON Ceramics”, J. Am. Ceram. Soc.,75 [2], pp. 259-276,1992. A. Rosenflanz and 1.W.Chen; J. Am. Ceram. SOC.82 (1999) 1025-1036 H. Mandal, D. P. Thompson, and K. H. Jack, “a@ Phase Transformations in Silicon Nitride and Salons”. Key Engineering Materials, 159-160, 1- 10, 1999. W. Y Sun, T. Y. Tien, and T. -S. Yen, “Solubility Limits of a‘Sialon Solid Solutions in the System Si, Al, Y/N, 0’.J. Am. Ceram. SOC.,74,2547-50, 1991.
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GRAIN-BOUNDARY PHASE CONTROL OF SILICON NITRIDE MATERIALS Takero Fukudorne" and Masahiro Sat0 Kyocera Corporation, R&D Center Kagoshirna 1-4 Yamashita-cho Kokubu, Kagoshima Japan
ABSTRACT
EXPERJMENTAL PROCEDURE
A new silicon nitride material has been developed. The strong points of this material are high strength up to 1 lOO"C,and it is possible to be sintered under normal pressure. These strong points were obtained by controlling the grain-boundary phase. Composition for the grain-boundary phase of this material was selected so that it could be sintered at low temperature. And the grain-boundary phase of this material was able to be crystallized easily. The microstructure of the new material was observed by using Transmission Electron Microscopy (TEM). The crystallization effect on high temperature strength and stress rupture strength were characterized through the TEM observation.
The direct nitridation powder was used. The yttrium oxide and aluminum oxide was selected for the sintering additives. These powders were mixed by ball milling. The mixed powder was pressed under 80MPa by Cold Isostatic Press (CIP). This green body was sintered at 175OOC for 5 hours in normal pressure nitrogen. The microstructure was observed by using SEM. The crystallization of the grain-boundary was confirmed by using X-ray Diffraction Pattern (XRD) and TEM. The flexural strength and the stress rupture strength were measured both at room temperature and 1000"C.
INTRODUCTION Silicon nitride materials have been expected as good candidates structural materials for high temperature use. They have been investigated by many researchers. As a result, the properties of silicon nitride materials have improved over the years. Recently, however, the demands of the market are in the direction of more reliable and more cost effective materials. To satisfy these demands, we have developed a new silicon nitride material. [l] Silicon nitride materials are known to require sintering additives in the process of sintering. These additives form liquid during sintering, and this liquid remains as a grain-boundary phase. Characteristics of silicon nitride materials are varied by controlling the grain-boundary phase. The new silicon nitride material can be sintered at low temperature under normal pressure. And the grain-boundary phase of this material was able to be crystallized easily. There are many reports about the crystallization of the grain-boundary phase. [2-31 By crystallizing the grain-boundary phase, the mechanical properties at high temperature were improved. [4-71 In this study, the microstructure of the new material was analyzed by using TEM to confirm the effect of crystallization of the grain-boundary phase on high temperature strength. The crystallization of grain-boundary phase during the measurement of the mechanical properties at high temperature was also observed. It was confirmed that the crystallization of the grain-boundary phase improved the mechanical properties at high temperature.
RESULT Many kinds of silicon nitride materials are available in the market. Fig.] shows the flexural strength of the new material and the conventional silicon nitride materials produced by Kyocera Corporation. The flexural strength was measured by the four point bending test in the ambient atmosphere. The new material has high strength both at room temperature and high temperature in spite of the normal pressure sintering. The yttria-alumina system was selected for sintering additives, since this system is inexpensive compared with the other additives. Composition of these additives was controlled so that this silicon nitride could be sintered at relatively low temperature. For this reason, the flexural strength decreased over 1200"C dramatically.
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Fig.2 shows the stress rupture strength of the new material and the conventional material at 1000°C. The new material has superior stress rupture strength compared with the conventional material. The slope of the new material is small compared with the conventional material. These results indicate that the expected life time of the new material is longer than that of the conventional material. Namely, the new material is more reliable than the conventional material.
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Fig.2 The stress rupture strength of the new material and the conventional material at 1000°C.
DISCUSSION HIGH STRENGTH AT ROOM TEMPERATURE The cross section surface was observed after 4 point bending test at room temperature by using Scanning Electron Microscopy (SEM). Fig.3 shows the microstructures of the new material and the conventional material. The new material has no large and elongated grain. This microstructure results in high strength at room temperature.
Fig.4 shows the X-Ray Diffraction (XRD) patterns of specimens with and without the heat-treatment. The non-heat-treated specimen revealed only beta silicon nitride pattern. A secondary crystalline phase from the grain-boundary phase appeared in specimens heat-treated over one hour. From these results, it was confirmed that the grain-boundary phase of this new material was crystallized easily. Fig.5 shows flexural strength before and after crystallization. The non-heat-treatment specimens have the amorphous grain-boundary phase. The heat-treated specimen for one hour has a partially crystallized grain-boundary phase. The heat-treated specimen for 100 hours has an almost crystallized grain-boundary phase. The average strength wasn't improved, but the Weibull coefficient was improved by crystallization of the grain-boundary. The low strength values disappeared after crystallization of the grain-boundary phase.
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Conventional Material New Material Fig.3 The microstructures of the new material and the conventional material. EFFECTS OF THE CRYSTALLIZATION The new material has superior flexural strength and stress rupture strength at high temperature. The crystallization of the grain-boundary phase may have contributed to the improvement of these properties. To confirm this consideration, flexural strength and stress rupture strength at high temperature after crystallization of the grain-boundary phase were measured. The 440
Fig.6 shows the stress rupture strength of the new material before and after crystallization at 1000°C. The grain-boundary phase of the specimens after crystallization was almost crystallized. The stress rupture strength was improved by the crystallization of the grain-boundary phase. The slope of the new material after crystallization became small compared with the new material before crystallization.
The above two results suggest that the superior mechanical properties of the new material at high temperature were the result of the crystallization of the grain-boundary phase. The grain-boundary phase of the new material was crystallized easily during measuring of mechanical properties at high temperature.
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1 10 Failure Time (h)
100
Fig.6 The stress rupture strength of the new material before and after the crystallization at 1000°C.
Fig.8 shows the TEM micrographs of the new material specimen after the stress rupture test. This specimen ruptured after 16 minutes of applying 700MPa. The portion of the grain-boundary phase was crystallized as can be seen in Fig.S(c). It was confirmed that the grain-boundary phase of the new material was crystallized easily during the measurement of mechanical properties at high temperature.
MECHANISM OF THE FRACTURE The two specimens after stress rupture test (625MPa-1OOh) were analyzed by TEM. Before the test one was heat-treated at 1000°C to crystallize the grain-boundary phase, the other was not heat-treated. Both the specimens survived 100 hours rupture test without failure. Fig.9 shows the cavities created by the stress rupture test. The sizes of the cavities of the non-crystallized specimen were larger than that of the crystallized specimen. There were many cavities in the non-crystallized specimen. On the other hand, it is difficult to find these cavities in the crystallized specimen. These results suggest that the crystallization of the grain-boundary phase suppressed the cavity formation.
THE TEM OBSERVATION OF THE GRAIN-BOUNDARY PHASE The new material was observed by using TEM (JEM201OF: JEOL) to confirm a state of the grain-boundary phase. Fig.7 shows the TEM micrograph of the as-sintered new material specimen. (b) Crystallized at 1000% for lOOh (a) Non crystallized Fig.9 The TEM micrographs of the specimen after the stress rupture test (625MPa-1OOh).
The grain-boundary phase was not crystallized. (a) (b) (c) Fig.7 The TEM micrographs of the as-fired new material. Photo (a) shows bright field. Photo (b) shows dark field. Photo (c) shows high-resolution lattice image.
Early stage of micro cracks creation was also observed. Fig.10 shows the early stage of the micro crack between the silicon nitride grain and the grain-boundary phase. A portion of the grain-boundary phase was crystallized. There existed non-crystallized grain-boundary phase between the silicon nitride grain and the crystallized grain-boundary phase. The micro crack appeared between the silicon nitride grain and non-crystallized grain-boundary phase. The width of the micro crack was about 5nm. This micro crack was not completely continuous.
Fig. 10 Micro crack between the silicon nitride grain and the grain-boundary phase after the stress rupture test at 1000°C (625MPa-1OOh) (a) Bright field (b) Dark field (c) High-resolution Fig.8 The TEM micrographs of the specimen after stress rupture test at 1000°C (700MPa-16min.).
Fig. 1 1 shows the micro crack between the two silicon nitride grains. These photographs may also show the early stage of the micro crack. The width of the micro
44 1
crack was also about 5nm. The micro crack was not also completely continuous. There seems to exist something between the two silicon nitride grains. This could possibly be considered as amorphous film. [8]
phase improves the mechanical properties at high temperature. After the stress rupture test at lOOO'C, the early stage of micro cracks was also observed by TEM.
REFERENCES (1) M. Sato, K. Sakaue, T. Fukudome and S. Wakida,
I I Fig. 11 Micro crack between the two silicon nitride grains after the stress rupture test at 1000°C (625MPa-1OOh) Fig.12 shows the grain-boundary between the two silicon nitride grains before and after the stress rupture test. The amorphous film was not observed between the two silicon nitride grains before the stress rupture test. However, the existence of the yttrium concentration was confirmed before the stress rupture test between the two silicon nitride grains by Electron Probe Micro Analyzer (EPMA). This result suggests the existence of the amorphous film between the two silicon nitride grains. On the other hand no corresponding micro crack was observed in the specimen heat-treated at 1000°C for 100h. As mentioned above, the sizes of the cavities in the crystallized specimen were very small. These results indicate that the crystallization of the grain-boundary phase could improve the stress rupture strength.
(a) After stress rupture test (b) As-sintered Fig. 12 Grain-boundary between the two silicon nitride grains. The new material has high stress rupture strength, because the grain-boundary phase of the new material is easily crystallized. The reason for superior stress rupture strength of the new material can be also addressed to less amount of amorphous film between silicon nitride grains.
CONCLUSION A new silicon nitride material has been developed. This new material has superior mechanical properties both at room temperature and at high temperature. By SEM and TEM analysis, we confirmed that the new material has a fine grain microstructure. The grain-boundary phase of this material is easily crystallized. The crystallization of the grain-boundary
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(2)
(3)
(4)
(5)
(6)
Development of New Silicon Nitride Materials for Automotive Applications. 6" International Symposium on Ceramic Materials and Components for Engines ( 1997) 173-176 R. Ramesh, E. Nestor, M. J. Pomeroy and Hampshire, Classical and Differential Thermal Analysis Study of the Glass-ceramic Transformation in a YSiAlON Glass. J. Am. Ceram. SOC.,81, [ 5 ] (1998) 1285-1295 D. V. Szabo, G. H. Campbell, J. Brulely, M. J. Hof&nann and M. Ruehle, Heterogeneous Nucleation in the Intergranular Phase of a SiAlON. J. Am. Ceram. SOC.,75, [I] (1992) 249-252 T. Fukudome, M. Sato, K. Sakaue, and T. Fujimoto, Mechanical Properties of Silicon Nitride at High Temperature and Crystallization of the Grain-boundary Phase. ll* Fall Meeting of the Ceramic Society of Japan. (1 998) 126 A. Yamakawa, M. Miyake and K. Ishizaki, Change of YAP-Phase Content and Some Properties by the Annealing of Si3N4-AI2O3-Y2O3-A1N. J. Ceram. SOC. Japan, 101 [9] (1993) 996-1000 A. Tsuge, H. Inoue and K. Komeya, Grain-boundary Phase Crystallization of Silicon Nitride with Material Loss During Heat Treatment. J. Am. Ceram.
Japan, 105, [6] (1997) 453-475
SINTERING AND MICROSTRUCTURE OF SILICON NITRIDE WITH MAGNESIA AND CERIUM ADDITIVES Haitao Yang, Ling Gao and Runzhang Yuan State Key Lab of Advanced Technology for Materials Synthesis and Processing, Wuhan University of Technology, Wuhan 430070, P. R. China
ABSTRACT This paper deals with the pressure less sintering and microstructure of Si3N4 -MgO-Ce02 ceramics. The sintered Si3N4-MgO-Ce02ceramics achieved a relative density of 98,5% and a bendino strength of 950 Mpa. XRD, TEM and EDAX analysis indicated that the main composition of the glassy phase in the sintered Si3N4MgO-Ce02 ceramics was cerium silicate and hardly contained any MgO. Abnormal grain growth occurred at 1850 "C, which led to micro cracks and dislocation. Some whiskers made of silicon were growing in fiactal patterns during TEM observation, which provided a good experimental example for further study of fiactal growth.
The TEM specimens were prepared in the usual was by cutting, grinding and finally ion beam thinning.
RESULTS AND DISCUSSION Effect of additive compositions The selected compositions are listed in Table 2. Figure 1 shows the relationship of MgOICeO, as a sintering aid for Si3N4.The relative density of the sintered Si3N4with 10% MgO is 97%, in comparison with the result of only 91% for that with 10% CeO,. But the combination of MgO and CeOz is better than either one. A highest relative density of 98.5% is obtained at a MgO/CeO, weight ratio of 5/5.
Keywords: Silicon, Nitride, Sintering, Microstructure Bending Strength
INTRODUCTION It is difficult to densify pure silicon nitride into useful high strength ceramics due to its covalent nature of bonding. Metal oxides such as Mg0['221,A1203[3*41, and rare-earth oxides[5971 have been found to be effective sintering and microstructure of silicon nitride with a combination if ceria and magnesia additives - although the effect of these two cations on sintering have both been extensively studied separately in combination with silica.
EXPERIMENTAL Starting materials and selected compositions The characteristics of the Si3N4powder used in the present study are listed in Table 1. The selected compositions are listed in Table 2. Experiment The selected compositions were mixed and ball cemented milled in alcohol for 24 h with WC-~YOCO carbide medium. The powder mixtures were dry-pressed into bars in a steel die at 120 Mpa. The compacts were embedded within a Si3N4+50wt. %BN mixed-powder bed in a molybdenum crucible and pressure less sintered in a 1 atm N2 atmosphere. The bulk density of the sintered specimens was measured using Archimede's Principle. Without being grinding, the as-sintered bar specimens, 5 x 5 x 30 mm in size, were used for 3-point bending strength measurements. Phase identification was made by X-ray diffraction using CuKa radiation. An H-800 transmission electron Microscope fitted with an EDAX was used for TEM work.
The as- sintered bars were used directly for 3-point bending strength vs. MgO/Ce02 weight ratio. The strength is very poor if MgO or Ce02 is added alone. But the strength is excellent if MgO and Ce02 are added in optimal ratio. The highest strength of 950 Mpa for the composition with a MgO/Ce02 weight ration of 5/5 indicates that the combination of MgO and Ce02 sintering aids is very effective. Microstructures Glassy phase: Fig. 3 shows a typical microstructure of the Si3N4+5%Ce02 ceramics sintered at 1800 "C for 60 min. The glassy phase remains at multigrain junctions as well as p - p Si3N4grain boundaries. The glassy phase can be confirmed directly by its electron diffraction, which appears to be an ambiguous circle because of the absence of Bragg diffraction. The result of the EDAX analysis of a glassy phase shows that Si, Ce are rich in the glassy phase but the amount of Mg is very poor (Mg: 0.64 at%, Si: 65.34 at%, Ce: 9.23 at%, Ce: 9.23 at%, Cu: 21.31 at%, W: 3.48 at% and Cu was introduced by specimen holder). This reveals that after sintering at 1800 "C, the main composition of the glassy phase in the sintered Si3N4+5%MgO+5%Ce02ceramics is cerium silicate and hardly contains any MgO. This result is confirmed by X-ray diffraction (XRD)analysis. Fig. 4 shows the XRD pattern for the Si3N4+5%MgO+5%Ce02ceramics sintered at 1800 "C for 60 min. There are no traces of Ce02, but MgO is found. This suggests that Ce02 has entered the glassy phase, and most of the MgO did not enter the glassy phase.
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Abnormal grain growth: Prevention of abnormal grain growth (AGG) is important if one chooses sintering to obtain high-strength ceramics. In the present study, AGG occurred at sintering temperature of above 1850 "C. AGG increases the grain boundary stress and leads to micro cracks [Fig. 51 and dislocations [Fig. 61, which is harmhl to the mechanical properties. So, for Si3N4MgO-CeOZ system, the sintering temperature shall not exceed 1800 "C. Fractal growth during TEM observation: Since B. B. Mandelbrot first provided his idea of fractal in 1970's, there has been much theoretical and experimental progress in the understanding of fractal phenomena [8-101. In order to have a better understanding of this complicated process, more direct observation of fractal in real world is required. In our experiment, we found some whiskers were growing in fractal patterns on Si3N4 grains during TEM observations (Fig. 7a, b). EDAX showed that these whiskers were made of Si. Under the radiation of 200 kv transmission electron beams in TEM, Si3N4,decomposed into Si (gas) and N2.Si (gas) then condensed to Si (solid) at the region that the radiant intensity was poor. The fractal growth process may be initiated by this evaporation-condensationprocess of Si, which provided a good experiment example for further study of fractal growth.
CONCLUSIONS
Si3N4 with Different Amounts of Sintering Aids. J. Mater. Res., 11, (1996) 120-126 (5) J. T. Smith and C. L. Quakenbush, Phase Effects in Si3N4 Containing Y2O3 and Ce02:I Strenght. Am. Ceram. SOC.Bull., 59 (1980), 529-532 (6) C. M. Wang, X. Pan and M. J. Hoffmann. Grain Boundary Films in Rare-Earth Glass-Based Silicon Nitride. J. Am. Ceram. SOC.79, (1996) (7) W. A. Sanders and D. M. Mieskowski, Strength an Microstructures of Sintered Si3N4 with Rare-Earth Oxide Additions. J. Am. Ceram. SOC.64 (1985) 304309 (8) B. B.Mandelbrot, D. E. Passoja and A. J. Paully, Fractal Character of Fracture Surfaces of Metals. Nature 308 (1984) 721-722 (9) David Avnir and Dina Farin, Molecular Fractal Surfaces. Nature 308 (1984) 261-263 (10) L. J. Huang, J. R. Ding and H. D. Li, Growth of the Fractal Patterns in Ni-Zr Thin Films During Ion-Solid Interaction. J. Appl. Phys. 83 (1988) 2879-2881 Table 1: Characteristics of Si3N4Powder* Phase
a 88
*
N 31,3
0 2,O
Si 58,O
Fe < 0.4
Particle Size
(cc)
1.2
Purchased from Zhuzhou Cemented Carbide Works Zhuzhou, 4 12000 China
Table 2: The Selected Compositions (wt.%) Samples 1 2 3 4 5 6 7
MgO-Ce02 ceramics achieved a relative density of 98.5% CeOz ceramics achieved a relative density of 98.5% and a bending strength of 950 Mpa. The main composition of the glassy phase in the sintered Si3N4MgO-CeOz ceramics was cerium silicate and hardly contained any MgO. Abnormal grain growth occurred at sintering temperatures of above 1850 "C, which led to micro cracks and dislocations. During TEM observation, under the radiation of transmission electron beams in TEM, some whiskers made of silicon were growing in fractal patterns, which provided a good experimental example for further study of fractal growth.
Content of elements (wt.%)
(wt.%)
MgO+ 10 8 6 4 2 0 5
CeOz++ 0 2 4 6 8 10
Si3N4 90 90 90 90 90 90
5
90
+
Purchased from Tianjing Chemicals, Tianjing 3000000 China ++ Purchased from Hunan Rate Earth Institute Changsha 4 10000 China
ACKNOWLEDGEMENTS The authors greatly acknowledge the support of K. C. WONG EDUCATION FOUNDATION; HONG KONG.
REFERENCES (1) G. Ziegler, J. Heinrich and G. Wotting, Relationships Between Processing, Microstructure and Properties of Dense and Reaction-bonded Silicon Nitride. J. Mat. Sci., 22, (1987) 3041-3086. (2) A. J. Pyzik and D. F. Carroll, Technology of SelfRainforced Silicon Nitride. Annu. Rev. Mater. Sci (1994) 189-212 (3) Y. Goto and G. Thomas, Phase Tranformation and Microstructural Changes of Si3N4 During Sintering. J. Mater. Sci., 30 (1995), 2194-2200 (4) S. Y. Yoon, T. Akatsu & E. Uasuda, The Microstructural and Creep Deformation of Hot-pressed 444
Fig. 1: Relative density vs. MgO/Ce02 (10% MgOCeO2, sintered at 1800 "C for 60 min)
Ce02 (wt.%)
Fig. 2: Bending strength vs. MgO/Ce02(10% MgOCe02, sintered at 1800 "C for 60 min)
Fig. 5: Micro cracks by a abnormal large grain (sintered at 1850 "C for 60 min)
&F 3: Typical microstructure of the Si3N4+5%MgO+ 5% Ce02 ceramics A
I
*
a
2 04 Deg) Fig. 4: XRD pattern for the Si3N4+5%MgO+ 5% Ce02 ceramics sintered at 1800 "C for 60 min, 0, a-Si3N4;A, 8- Si3N4;0 MgO; A,CeOz;
0,WC
Fin. 6: Dislocations in a abnormal large grain (sintered at 1850 "C for 60 min)
a) b) Fin. 7: Fractal growth on Si3N4grains during TEM observation a) Observing 2 minutes b) Observing 6 minutes
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CHARACTERISATION OF MULTI-CATION STABILISED ALPHA SIALON MATERIALS M.RTerner*, S.P.Swenser, Y.-B.Cheng, Department of Materials Engineering, Monash University, Australia
ABSTRACT Si3N4 based ceramics have many desirable engineering properties including high hardness, corrosion resistance, and strength retention at elevated temperatures. Reaction sintering to produce these materials requires high purity, synthesised starting powders that result in a costly but homogenous product best suited for specialised high performance applications. Presently, a carbothermal reductionnitridation process is being investigated as a low-cost alternative for a-sialon production. One of the main advantages of this is the ability to utilise lower grade, impure starting minerals, whereby impurities present may be incorporated into the a ’ structure. The effect of various potential stabilising cations present in this system is investigated via electron microscopy and EDXS techniques.
occumng aluminosilicate minerals with significant Mg or Ca contents may be potential candidates for a’ formation. Furthermore, other impurity elements in the raw minerals may potentially be incorporated into the a‘ structure, thus producing multi-cation stabilised asialons. The impurities may however result in a high quantity of residual glass and thus have a deleterious effect on mechanical properties, especially at high temperatures. Hence this material may be best suited to low-medium temperature applications that can benefit from the intrinsically high hardness and good corrosion resistance of a-sialon. It is apparent that impurities in the starting minerals play a critical role in the development of the final material microstructure and mechanical properties. This paper presents the results of a characterisation study of an a’material produced by CRN of low-grade minerals.
INTRODUCTION
METHOD
Alpha sialons (a’) possess many desirable engineering properties such as high hardness, corrosion resistance, wear resistance, and high strength retention at elevated temperatures. Despite these attractive properties, the high cost of production coupled with low fracture toughness has limited the use of a ’ materials to specialised, high performance areas such as metalcutting tools. One of the main production costs is that of the raw materials - the reaction sintering process utilises costly, high-purity, synthesised nitride and oxide powders, and high firing temperatures (1600- 1SOOOC). Sialons are formed from the substitution of Al and 0 into Si3N4,and in the a ’ case there is extra substitution of A13+beyond that of the 0’content, which results in an overall excess of negative charge. Charge compensation occurs by adding various stabilising cations such as Li+, Mg2+,Ca”, Nd3+,Sm3+,Y” to the system, usually via oxide powders. These cations can inhabit two large interstices, or ‘cages’ located within the a ’ unit cell. An alternative processing route is that of simultaneous carbothermal reduction and nitridation (CRN) of oxide starting powders in the presence of a carbonaceous reductant and nitrogen atmosphere. This process has been successhlly used to manufacture Si3N4from silica [l], and f3- and 0-sialons from clays [Z-31,but has received less attention for the production of a-sialons [4-61. A major attraction of this process is the possibility of utilising lower-grade and thus relatively inexpensive starting powders. Naturally
The material to be investigated, designated ‘Cl’, was produced via simultaneous carbothennal-reduction of a mixture of two locally available aluminosilicate based minerals. The composition contained CaO in significant quantity, and minor amounts of MgO, Ti02 and F@O3 were present in addition to a few other trace elements. This composition was chosen because it contained several possible stabilising elements - Ca, Mg, Ti, and Fe, though only Ca and Mg are currently known to act as a ’ stabilisers. Furthermore, it was known from previous work that the CRN of this composition would produce a predominantly a-sialon powder. To facilitate hrther analysis, the a’powder produced was hot pressed to form a dense pellet. X-ray diffraction was performed using a Rigaku-Geigerflex BraggBrentano difiactomer with Ni-filtered Cu k a radiation. The microstructure was viewed after etching the sample in molten NaOH using a Jeol JSM 6300F FEG SEM, and a Philips CM20 TEM equipped with an Oxford Pentafet EDXS detector was used to analyse an ionbeam thinned foil.
RESULTS / DISCUSSION X-ray Diffraction The XRD trace in Figure 1. shows that a-sialon is the dominant phase formed, along with a minor amount of AlN and glass. Several unidentified peaks remain in the spectrum.
447
a’
a’
a’
a’
1 *fa‘ a’
10
1s
?
20
I
a’
a’
25
a’
30
2-*
4a
a’
a’ a’ 45
so
a’
55
a’
(10
Figure 1. X-Ray Diffraction Trace of C 1
Microstructural Overview (SEM) Micrographs of the etched surface of C 1 are given in Figure 2. The majority of grains display a roughly equiaxed morphology, with an average grain size of 0.51pm. Large, irregular 3-5pm inclusions are also present (Fig. 2b).
Figure 2. Microstructural overview
448
Transmission Electron Microscopy (TEM) Overviews of two typical regions within C1 are given in Figures 3 a) and b). The majority of grains are sub-micron in diameter, with occasional elongation of the a ’ also evident. The small -250pm overlapping particles in a) show distinct hexagonal facets often characteristic of a-sialon. In addition, small, 40-100nm inclusions (indicated by arrows) can be seen within the a’grains. A typical CBEDP and EDS spectrum from an asialon grain is given in Figures 3d) and e). The elemental distribution clearly shows the sialon elements Si, Al, 0 and N, and Ca. The carbon signal is an artefact from the conductive sample coating. This sialon is relatively Al rich, and on average has an Si:Al:Ca peak height ratio of -4:2:1. Neither Fe nor Ti were detected within the sialon, and interestingly neither was Mg, the only other known a‘ stabilising cation in the system. From these results it is clear that only a single-cation stabilised Caa-sialon was produced in this system. From recent work by Wang et al. [7-81 it would seem that there is simply insufficient Mg for multi-cation (Mg,Ca)-a-sialon formation, where it was shown that large Mg:Ca ratios of 5050 did not result in significant Mg levels within the a’, and even small Mg additions to Ca-a’ systems tended to promote an increase in the Ca content of a’. The bright field image in Figure 3c) shows two elliptical, -lpm AlN grains, both containing small particles. The EDS spectrum (Fig. 3g) shows a slight trace of Si and Mg in addition to the strong A1 peak, however this phase is not a polytypoid phase and the CBEDP was clearly indexed as A N . These small traces are the result of a small substitution of A13+by Mgz+and an equivalent amount of Si4+to balance the valency. Small 40-1OOnm inclusions previously seen within the a ’ grains were also detected within AlN grains, as shown in Figure 3c). A CBEDP and EDS spectrum are given in Figures 3h) and i). These inclusions have been identified as cubic TiN, with a small inclusion of Si4+
00
n
22
3
@kaV5
6
??
88
Figure 3. a-c) BF overviews showing a', AlN and TiN particle morphology; d) CBEDP of a'; de) EDS spectrum from a'; f ) CBEDP of AlN, g) EDS spectrum of AIN; h) CBEDP of TiN; i) EDS spectrum from TIN. evident in the EDS spectrum. The surrounding AlN grain is the most likely source of the Al signal. Nitridation of TiOz to form TiN in Si3N4-ceramic systems has been reported to be thermodynamically viable at similar temperatures to those used in this process [9],and has been deliberately used to form TiNsialon composite materials with enhanced electrical conductivity [lo].
TEM analysis of the large, dark particles seen in the SEM overviews is given in Figure 4a). The particles have been identified as cubic FeSi, with a significant substitution of Ti4+for Si4+.The larger size of the Ti4+ cation produces a slight but detectable increase of >1.5% in the crystal d-spacings as determined from diffraction patterns. Fe is one of the main impurities in many commercial silicon powders used for Si3N4 production and clay minerals used to produce p-sialons,
449
and its effect on nitridation reactions has been studied previously [ll-121. It was found that Fe has a catalytic effect on nitridation reactions through the formation of FeSi, transient liquid phases, however the effect of residual FeSi, particles on mechanical properties has received little attention. The morphology as seen in Figure 2b) is indicative of solidification of a transient liquid, and the poor dispersion and large size of the agglomerates is likely to be detrimental to the mechanical properties of the material. This issue must be addressed in hrther work.
indicates that aside from the small amount found in the FeSi phase, the Ti was completely nitrided to TiN.
00
I'I
22
3
MkeVS
0
7
88
Figure 5. a) BF image of a glassy pocket; b) EDS spectrum taken at x
Conclusions Carbothermal reduction nitridation was successfidly used to produce a predominantly a-sialon material. However this analysis has shown that despite the potential to produce a multi-cation stabilised a', only a single-cation Ca-a' phase was formed. Ti and Fe impurities in the starting composition formed discrete secondary phases, which in the case of FeSi were large agglomerates likely to be detrimental to mechanical properties. Mg did not enter the a-sialon, but was predominantly tied up in the residual glassy phase, as were many of the other trace impurities. Due to the interchangeability of many of the cations present due to similarity in size andor valency, a small extent of substitution was evident all phases found. Figure 4. a) BF image of FeSi grain (dark); b) CBEDP of grain; c) EDS spectrum Analysis of the residual glass is given in Figure 5. Given that the Ca content of this system is far in excess of that required for sialon formation, it was no surprise to find the Si02-based glass also very rich in Ca. The Mg is also accounted for in the glass, as are many of the other trace impurities. The Mo signal is an artifact from the ion-beam thinning process. The lack of Ti detected
450
References (1) Y.W.Cho, J.A.,Charles,
Synthesis of Nitrogen Ceramic Powders by Carbothermal Reduction and Nitridation. Part 1: Silicon Nitride. Mat. Sci and Tech., 7, (1991) 289-98. (2) C.Bishop, A.Hendry, Thermal Analysis of Formation of Sialons by Carbothermal Reduction of Clays. J.Tht~m.Anal.,42, (1994) 697-711.
(3) C.G.Barris, et al., Reaction Bonded 0-Sialon and 0 Sialon-Silicon Carbide. J. Aust. Cer. Soc.,33, (1997) 15-20. (4) M.Mitomo, M.Takeuchi, M.Ohmasa, Preparation of a-Sialon Powders by Carbothermal Reduction and Nitridation. Cer. Int., 14, (1988) 43-8. ( 5 ) J.W.T. van Rutten, et al., Carbothemal Preparation and Characterisationof Ca-a-sialon. J.Eur. Cer. SOC., 15, (1995) 1-6. (6)R,Metselaar, et al., The Synthesis of a- and pSialons from Fly Ash. Key Eng. Mat., Proc. PacRim 2, Cairns, Australia, (1996). (7)P.L.Wang, C.Zhang, W.Y.Sun, D.S.Yan, Formation Behaviour of Multi-Cation a-sialons Containing Calcium and Magnesium. Mat. Let., 38, (1999) 178185. (8) P.L.Wang, Private communication. (9) M.B.Trigg, E.R.McCartney, Comparison of the Reaction Systems Zr02-Si3N4 and Ti02-Si3N4. Comm.Am.Cer.Soc.,Nov, (198 1) C-15 1-152. (1O)F.Hong, R.J.Lumby, M.H.Lewis, TiN/Sialon Composites via In-Situ Reaction Sintering. J. Eur. Cer. SOC.,11, (1993) 237-239. (1 l)S.M.Boyer, A.J.Moulson, A Mechanism for the Nitidation of Fe-Contaminated Silicon. J.Mat. Sci., 14, (1978) 1637-1646 (12)A.D.Mazzoni, E.F.Aglietti, E.Pereira, p'-Sialon Preparation from Kaolinitic Clays. Appl. Clay Sci., 7, (1993) 407-420.
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GRAIN BOUNDARIES OF SiC-Si02COMPOSITE Haihui Ye*, Georg Rixecker, and Fritz Aldinger Max-Planck-Institut fur Metallforschung and Institut fur Nichtmetallische Anorganische Materialien, Universitat Stuttgart, PulvermetallurgischesLaboratorium, 70569 Stuttgart, Germany
ABSTRACT
EXPERIMENTAL PROCEDURE
Silicon carbide was liquid phase sintered using Si02 as the only sintering additive. HREM was operated to observe the grain boundaries in the Sic-Si02 composite. AES depth profile and energy selected TEM were used to identify the existence of carbon in the interface of SiC-Si02. The significance of carbon in the intergranular amorphous films in the Sic-Si02 composite is discussed in detail basing on Clarke’s equilibrium thickness model.
Sic-based samples were liquid phase sintered with SiOz as the only sintering in order to obtain the intergranular phase without extra elements such as Al, N, Y, which would complicate the chemical microanalysis and make theoretical calculations difficult. a-Sic (UF 15, H.C.Starck) was used as raw material. The amount of silica was 20 wt%. CO atmosphere was used in order to prevent the following reaction during sintering: (1) S i c + 2Si02 + 3SiO + CO By sintering at 1860°C for 30 minutes, quite high densities (up to 94.2% of the theoretical density) were achieved[101. Specimens for TEM investigation were prepared by a standard slicing, polishing, dimpling procedure and finally by Ar-ion milling until perforation. In order to observe the grain boundary and measure the thickness of amorphous intergranular phase, H E M was carried out at a voltage of 400kV, using a JEOL JEM 4000FX instrument. Another model experiment was performed to identify the composition of the interface between silica and Sic. A high-purity single crystal of Sic was polished optically flat and then heated to 1100°C in air atmosphere for 3 hours. A thin film (650 nm) of S O 2 was thus formed on the surface of Sic. To crystallize the Si02 film and to approach thermodynamic equilibrium, the sample was heated to 1450°C (the softening temperature of silica) and held for 12 hours. Auger Electron Spectroscopy (AES) depth profiling of the surface Si02 film was performed to investigate the composition of Si02-SiC interface. The primary sputtering ion energy was 3keV. For TEM observation, the sample was cut to two equal pieces and epoxied with the film sides facing each other. The sandwich was cross-sectioned into -1 mm thick slices, which were polished dimpled, and ion milled to electron transparency. Energy selected imaging (ESTEMI) using an in-column energy filter on the Zeiss EM 9 12 Omega TEM (100kV) with slow scan CCD camera image acquisition was performed to provide the chemical distribution information for the Si02-SiC interface.
INTRODUCTION Silicon carbide can be sintered to high density using many additive systems: B-C, AI2O3-Y2O3,or AIN-Y203. These additives either promote solid state sintering or form liquid phases at elevated temperature and fill the intergranular spaces. The distribution of these liquid or glass phases in the microstructure influences both the sintering behavior and the final properties of the ceramic and have therefore been the subject of intensive investigation. A general model for the intergranular glass phases in ceramic materials has been advanced by Clarke[ I]: his calculation of the force balance across the interface which suggested that nearly all the intergranular phases exhibit a stable equilibrium thickness. This model successfully explains the existence of equilibrium-thicknessglass films (-1 -2nm) in Si3N4-basedceramics[2], the ZnO-Bi203 system[3], and the TiOz-Si02 system[4]. However, when this model is considered in connection with Sic-based ceramics, some uncertainties arise. According to Clarke’s calculations, intergranular Si02 films are not stable at Sic-Sic grain boundaries; the calculated equilibrium thickness is zero, because the attractive van der Waals force between Sic particles is larger at all distances than the steric repulsive force which is related with the ordering of the silica tetrahedra in the intergranular film. This theoretical prediction was in agreement with that early TEM work on liquid phase sintered Sic, where no siliceous intergranular phase was formed[5], but was doubted later by the other scientists[6-91 who had clearly observed thin intergranular films at Sic-Sic boundaries in S i c based ceramics. Thus, further research is necessary to check the validity of Clarke’s model with respect to the intergranular phases in Sic-based system.
RESULTS AND DISCUSSION Fig.] is a low-magnification TEM image of the SiCSi02 composite. In the triple junctions, there is a lot of amorphous phase because the content of Si02 additive
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Fig. 1. Low-magnificationTEM image ofSiC-Si02 composite
intergranular film is the most important one. For example, a small addition of CaO (80 ppm) into a-Si3N4 causes the intergranular film thickness to decrease, while larger CaO concentrations(220ppm and 450 ppm) cause it to increase[ll]. So, what is the film composition in SiC-Si02 system? Is it only amorphous silica? If not, how does it affect the equilibrium film thickness? To answer these questions, a second model experiment was designed and performed as described in the ‘Experimental’ section. Since both experiments were conducted at high temperature for an extended time, the measured composition of the amorphous phase between crystallized SiOz and single crystalline S i c in the second model experiment, is considered to be representative of that of the intergranular phase in the first experiment. In the second model experiment, AES was used because it is highly surface sensitive. Only the top lnm of the surface of the solid is analyzed due to the low escape depth of Auger electrons. The depth profiling based on AES can determine the variation in type and concentration of species as a function of distance beneath the surface. The technique involves the removal of successive layers from the material under study by sputtering and recording the Auger spectrum of the freshly exposed surface. Fig. 3 shows the results of the measurement at a sputtering rate of 20nm depth per minute. At a depth of about 650nm, the interface between Si02 and S i c was found with a thickness of
Fig. 2. High-resolution TEM images of intergranular amorphous phase in liquid phase sintered Sic-SO2 composite
in this composite is quite high (20 wt%). Quartz particles can be found in some large triple junctions, which is in agreement with the X-ray diffraction results. Both selected area electron diffraction and X-ray diffraction show that S i c remained in the a-Sic modification during sintering. An intergranular amorphous phase film can always be found by H E M , as shown in Fig. 2, where the film thicknesses are measured to be 1.2nm and 1.5nm. On the whole, ten intergranular films have been analyzed in detail; their thickness ranged from 0.7nm to 6nm, and the mean thickness was 1.5nm. This result is in agreement with many experimental observations in recent year’s papers [6-91, but at variance with the theoretical prediction by Clarke’s model. Since in the present model experiment, SiOz was selected as the only sintering additive in order to exclude any disturbances produced by extra elements, it corresponds to the model ceramic system (SiC-Si02-SiC), for which the model calculations were performed, to the highest possible degree, and therefore most clearly manifest the inconsistency existing in the Clarke’s model. However, there always remains some uncertainty in sintering experiments, as to whether thermodynamic equilibrium has been reached. Among the many factors which can affect the equilibrium film thickness, the composition of the
454
I00
600
700
800
900
Sputter Depth (nm)
Fig. 3. AES depth profile of heat treated SiO2 surface layer/SiC single crystal specimen
b a Fig.4. Bright field image(a) and energy selected TEM image for carbon(b) ofa SiO2 layer / Sic single crystal specimen
65nm. Along the transition from SiOz to Sic, the intensity of the oxygen signal goes dohwn to zero while the intensities of silicon and carbon gradually rise up to their maximum. Although substrate roughness would also possibly present this kind of a gradual transition even for an atomically sharp interface[ 121, this effect is very small in this case because Sic substrate was polished and proved by TEM micrographs to be rather smooth (Fig. 4(a)). Thus, the only plausible reason of the graduating carbon concentration along the interface is that C from the S i c substrate dissolves into the SiOz at high temperature and forms a SiO,C, interface layer. This amorphous phase is clearly visible by using energy selected TEM Imaging for carbon as shown in Fig. 4(b). The thickness of the Carbon-containing interface layer is 60nm, which corresponds to the chemical width measured by AES depth profiling. Returning to the liquid phase sintered SiC-Si02 system, where the thickness of the intergranular amorphous phase (0.7-6 nm) is much less than the thickness of the interface layer in the single crystal-oxidation experiment, it is reasonable to consider that carbon from Sic grains can also dissolve into the intergranular amorphous phase at high sintering temperatures, replacing part of the Si-0 bonds by Si-C bonds to form silicon oxycarbide(SiO,C,). Silicon oxycarbide is an amorphous metastable phase [ 131 whose possible configurations are [SiO$], [Si02C2],and [SiOC3]. The
(a) U
\ Energy barrier 50.2
0.61
r y ;.
,
. , . ,,.,
. . ..
, ,,
h(ea) . . I
Film Thickness. h. (nm)
F i g 5 Calculated Interaction Energy E as a hnction of film thickness h for (a) Si3NAi02-Si3Ns and (b) SiC-SiO4iC systems, assuming that aq”’=100MPa, {=2.5A. (a) Hapa=76~10-21J, and (b) I-I.,~~=233xlO~~’J. The resulting equilibrium thicknesses are (a) I .8nm and (b) Onm, respectively.
significance of this kind of microstructure will be discussed as following after the important theoretical calculation based on the Clarke’s model. According to Clarke’s model, the total interaction energy across an intergranular liquid film due to van der Waals attraction and steric repulsion is given by
L+ J where H N ~is the Hamaker constant for grains of material ‘a’separated by a film of material ‘p’ , h is the film thickness, cq; is a constant called ‘ordering force’ which determines the strength of the repulsion, 6 is the structural correlation length which is assumed to be of molecular dimensions. For the Si3N4-SiOz system, where the Hamaker constant is 76~1O-~’J [I], aq; = 1 OOMPa, and E,=2.5A (approximately the size of a Si04 tetrahedron), the interaction energy E(h) can be plotted as a function of film thickness h (Fig.S(a)). It is apparent that along with the decreasing of the film thickness there maybe two local energy minima, namely the ‘secondary minimum’ at the so-called ‘equilibrium film thickness’ (for Si3N4-Si02, 1.8nm), and the ‘primary minimum’ at varnishing film thickness (crystal-crystal grain boundary). In sintering experiments, the intergranular film thickness decreases under the influence of the attractive Van-der-Waals force, starting from a large initial separation of the grains. Unless the Van-der-Waals interaction is very large, the thinning of the intergranular film will stop as force equilibrium with the steric repulsive force is reached, thus establishing an equilibrium film thickness. Since the energy barrier between the primary and secondary minimum is high, the resulting intergranular amorphous films are thermodynamically meta-stable. Same situation exists in S&N4 and most other liquid-phase-sintered ceramics, their Hamaker constant being comparable or smaller then in the case of Si3N4.The only exception is the SiCSi02 system, where the Hamaker constant is exceptionally large, reaching 233xlO-”J [I]. The energy-thickness diagram of this system is plotted in Fig. 5(b) (cq; = IOOMPa and E,=2.5A). Here, no energy bamer exists and the equilibrium thickness of the amorphous film is zero. However, in the actual sintered SiC-Si02 ceramic, the
SOz
SiO,C,
Fig. 6. Schematic illustration of the structure of silica and silicon oxycarbide. The correlation length is increased as carbon dissolves into silica.
455
--
Another model experiment was designed to identify the existence of carbon in the amorphous interface of SicSOz. Carbon is shown to change the structural correlation length of the intergranular phase, increasing the steric repulsion and finally stabilizing the intergranular amorphous films. If this influence of carbon in the intergranular phase is taken into consideration in the calculation of the equilibrium film thickness, Clarke’s model is found to describe the observed microstructures of LPS-Sic materials correctly.
I
10
0.1
1W
Flm Thickness. h. (nm)
REFERENCES (1) D.Clarke, J. Am. Ceram. SOC.,70[ 13, (1987) 15-22.
0.1
10
100
Film ThiCkM88, h. (nm)
Fig. 7. Calculated Interaction Energy E as knction of film thickness h for the SiC-SiO2-SiC system, assuming that H+=233xIO-”J, cqo’ = I OOMPa, (a) <=3.OA, and (b) <+.OA. The equilibrium thicknesses are (a) 1.3nm and (b) 2.4nm. respectively.
intergranular phase is not pure silica, but silicon oxycarbide, as shown in the above experiments. The effect of carbon is to bridge four Si-based tetrahedra, instead of only two tetrahedra in the case of oxygen, as shown in the schematic illustration(Fig. 6), and thus increase the size of the structural correlation length 5. The intergranular film thickness has been demonstrated to be relatively insensitive to the magnitude of the ordering force aq;, but very sensitive to the correlation length[l4]. A small increase of the correlation length will greatly enhance steric repulsion and can thus cause dramatic change in the energy-thickness diagram. Fig 7(a) and Fig.7(b) show the calculation results for the Sic-Si02 system with 6=3.0A and 6=4.OA, respectively. With increasing correlation length, the secondary energy minimum appears, and the equilibrium thickness increases from 1.3nm for 5=3.0A to 2.4nm for 4.OA. These revised calculations based on Clarke’s model are then consistent with the experimental observation, that intergranular films are characteristic for all liquid phase sintered materials based on Sic, even in the special case of the pure SicSi02 system.
CONCLUSION A model system Sic-SiO2 has been liquid phase
sintered. HREM shows that intergranular amorphous films always exist in the final product, which is contrary to the theoretical calculation based on Clarke’s model.
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(2) H.J.Kleebe, J. Ceram. SOC.Japan, 105[6], (1997) 453-475. (3) H.Wang and Y.M.Chiang, J. Am. Ceram. SOC. 81[1], (1998) 89-96. (4) H.D.Ackler and Y.M.Chiang, J. Am. Ceram. SOC. 82[1], (1999) 183-189. ( 5 ) R.W.Carpenter and W.Braue, J.Mater.Res. 6[9], (1991) 1937-1944. (6) S.Turan and K.M.Knowles, J. Microscopy, 177, (1995) 287-304. (7) L.S.Sigl and H.J.Kleebe, J. Am. Ceram. SOC.76, (1993) 773-776. (8) J.J .Cao, W.J .MoberlyChan, L .C.DeJonghe, C.J.Gilbert, R.O.Ritchie, J. Am. Ceram. SOC.79[2], (1996) 461-469. (9) L.K.L.Falk, J. Eur. Ceram. SOC.17, (1997) 983-994 (10) H.Ye, G.Rixecker and F.Aldinger, Liquid phase sintered S i c with SiOzadditive, Proc. EUROMAT 1999, in press. ( 1 1) I.Tanaka, H.J.Kleebe, M.K.Cinibulk, J.Bruley, D.Clarke and M.Ruhle, J. Am. Ceram. SOC.77[4] (1994) 91 1-914. ( 1 2) M.Thompson, M.Baker, A.Christie and J.Tyson, Auger Electron Spectroscopy, John Wiley & Sons, New York, (1985). (13) C.Onneby and C.G.Pantano, J. Vac. Sci. Technol. A 15(3), ( 1 997) 1597-1602. (14) H.D.Ackler, Thermodynamic Calculations and Model Experiments on Thin Intergranular Amorphous Films in Ceramics, Ph.D. Thesis, Massachusetts Institute of Technology, Cambridge MA, (1 997).
MILLIMETER WAVE SINTERING OF CERAMICS G. Link, S. Wee*, L. Feher, M. Thumm' Forschungszentrum Karlsruhe GmbH, IHM, D-76021 Karlsruhe, Germany # and Universitat Karlsruhe, IHE,D-76128 Karlsruhe, Germany
ABSTRACT At the Forschungszentrum Karlsruhe, Germany, a
compact 30 GHz gyrotron system has been established in order to investigate technological applications in the field of high temperature materials processing by means of microwave radiation in the millimeter wave (mmwave) range. Besides the improvement of the system design, research activities are mainly engaged in studies on debindering and sintering of various types of advanced structural and functional ceramics. The use of microwaves at 30 GHz distinguishes from the widespread industrial microwaves with frequency of 2.45 GHz and 0.915 GHz with respect to field and therefore heating homogeneity. Furthermore dielectric absorp-tion and therefore heating at higher frequencies is stronger, so that a larger variety of materials especially low loss materials can be heated. Due to volumetric heating and enhanced sintering kinetics the application of microwaves allows to shorten the processing time and therefore to reduce the energy consumption. Besides these effects microwave technology gives the unique possibility to influence microstructure and therefore physical properties of the ceramic materials. This paper will give an impression about the benefits of the mm-wave technology with respect to sintering of various advanced ceramics.
INTRODUCTION Microwave heating is a relative new technology in the field of materials processing although it was already conceived about 50 years ago. The utilization of microwaves for heat generation was discovered accidentally during testing of magnetrons at the Microwave & Power Tube Division of Raytheon in 1950 [l]. Since the early 1950s when the first industrial microwave ovens for heating of food were fabricated [2] there is growing evidence to support the use of microwaves in certain industrial processes such as food preparation or rubber vulcanization. Investigations of Levinson in the 60s [3] have shown that sintering of ceramics appeared on microwave heated supports. Sutton found that in addition to water removal from high alumina castables microwave processing leads to heating of the ceramic itself to temperatures well above 1400 "C [4]. Based on these results several research
activities in different research groups have been conducted. The heating process for debindering or sintering is one of the most crucial steps in the processing of modern high-performance ceramics. Due to the low penetration depth of IR-radiation in a standard resistance heated or gas fired sintering furnace, temperature gradients are induced in the ceramic parts. This results in a hot surface and a colder interior whiich are leveled out during a time consuming thermal conduction process. Such temperature gradients inherent in conventional heating techniques lead to thermal stresses within the ceramic body and therefore to its possible destruction, if not an optimized timetemperature program with low heating rates is applied. The use of microwaves allows to transfer energy directly into the volume of the materials which results in the possibility to apply high heating rates, that means significant shortening the processing time. Furthermore, the densification process of ceramic bodies seems to be enhanced by sintering in a microwave field, which has been demonstrated by several authors [5,6] allowing a reduction of the sintering temperatures and dwelling times compared to the processing in a standard sintering furnace. This makes the use of microwave technolclgy very attractive because of specific economical benelits such as energy conservation, reduced cycle time, reduced operating space and improved environment. In addition microwave technology gives the unique possibility to influence microstructure (e.g. grain size) and physical properties of the ceramic materials.
GYROTRON INSTALLATION Based on the number of installed units and power, microwaves with a frequency of 2.45 GHz and 915 MHz found the broadest application up to now. With the foundation of research on nuclear fusion and ithe development of plasma heating technologies the idea of microwave sintering experienced a renaissance, due to the availability of high power millimeter wave sources, so-called gyrotrons. At the Oak Ridge National Laboratory this technology was demonstrated first for experiments on sintering of ceramics [7]. They used a 200 kW 28 GHz gyrotron from Varian Associates Microwave Tube Division, now called Communications & Power Industries, Palo Alto; California.
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Fig. 1: Experimental setup for sole mm-wave heating (left) and hybrid heating (right).
In 1994 at the Forschungszentrum Karlsruhe a compact prototype gyrotron-system was installed for materials processing. It includes a 30 GHz, 10 kW gyrotron from the Institute of Applied Physics, Nizhny Novgorod, Russia [8]. The power can be computer controlled in that way, that the temperature measured with a thermocouple at the sample surface is following any preset time-temperature program. A calibration of the thermocouples by measuring the melting point of gold yields an error of less than 1%.
Advantages compared to Magnetrons The use of mm-waves distinguishes from the application of widespread industrial microwaves in the dm-range with frequencies of 0.915 GHz or 2.45 GHz with wavelengths of 32 cm and 12.24 cm respectively. First of all the absorbed power P inside the material, is proportional to the angular frequency o and the dielectric loss factor &", which itself is a function not only of frequency but also of temperature. This loss factor is increasing with temperature as well as with frequency in the microwave range for many low loss advanced ceramics.
L
This means, in case that the installed microwave power for a certain heating experiment with low loss ceramics is not sufficient to reach sintering temperatures there are only three things which can be done to overcome this problem. Either the ceramic mass that means the load of the microwave applicator has to be increased if possible in order to increase the part of the
458
power absorbed in the ceramic load compared to the power absorbed in the thermal insulation and resonator walls. That means increasing the heating efficiency by increasing the ceramic batch. But this unavoidably leads to restrictions in maximum heating rates due to limited microwave power. Another thing that can be done is preheating the ceramic by infrared radiation to temperatures where coupling of the ceramic samples itself, that means E " ( W , T ) is sufficient high for microwave heating. This can be realized either by hating the resonator walls with conventional heating technology or by arranging microwave susceptors such as graphite or silicon-carbide for indirect preheating of the ceramic samples. But the only solution without any further restrictions to the arrangements inside the oven is to change to higher frequencies since heating with higher kequencies such as 30 GHz is, corresponding to equation (l),by a factor of 10 to 100 stronger compared to the heating with magnetron frequencies. This means that a larger variety of materials, especially advanced technical ceramics with low dielectric losses, such as Al203 or Si3N4, can be heated more easily with gyrotron compared to magnetron radiation. Another difference between magnetron and gyrotron frequencies can be seen in the heating patterns evolving inside the ceramic samples. Each sample represents a dielectric resonator so that interference patterns of the electremagnetic field may evolve, depending on frequency, dielectric properties, resonator geometry and physical size of the samples itself. Since heat is created by interaction of the electromagnetic energy with the materials, inhomogeneous fields invariably lead to hot spots and non uniform processing. Due to the temperature sensitivity of the microwave absorption coefficient E l of many ceramics, rapid heating may result in local thermal runaway, melting or cracking. if
such hot spot cannot be avoided. Therefore with short wavelengths this problem is less critical. The distance between hot spots is so short that they can be leveled out by thermal conduction. Finally, it is much easier to establish a homogeneous field distribution inside an applicator with shorter wavelengths provided that an optimized applicator geometry is used. This leads to a more homogeneous temperature field within the parts to be processed [ 101.
EXPERIMENTAL RESULTS OF MILLIMETER WAVE SINTERING Experimental procedure Although many dielectric ceramics are ideal candidates for a volumetric and fast heating with microwaves, in practice some problems appear with pure microwave heating. Because heating with microwaves is always selective materials with high dielectric losses absorb more energy than materials with lower losses. Ceramic samples achieve a higher temperature than their environment, depending on the type of material. This results in energy losses from sample surfaces by conduction, convection and radiation processes leading to temperature gradients opposite to the gradients existing with conventional heating. Several methods have been developed to overcome this problem. One method uses thermal insulation material, transparent to microwaves but opaque to thermal radiation, which helps to avoid energy losses from sample surface (Fig.1). Materials typically used for this purpose are highly porous silica, alumina or mullite fiber boards, depending on the temperature range they are exposed to. In case that this method is not sufficient enough microwave susceptors can be used additionally. Such susceptors typically made from lossy Sic ceramic are able to compensate the energy losses from sample surfaces by their own thennal radiation, provided that such susceptors are arranged at the optimal positions. Once the thermal insulation and the arrangement of susceptors is fixed the temperature evolution during the process cannot be influenced any more. The use of these methods allows an essential reduction of temperature gradients but these methods will never be able to avoid temperature gradients totally as long as the samples are irradiated by microwaves and penetration of microwave is sufficiently high. In thermodynamic equilib-rium of sample surfaces with their environ-ment, volumetric heating always results in a inverse temperature profile along the samples interior [91. As long as single and small ceramic samples were used for investigations of mm-wave effects the problem of inverse temperature profile was solved by using ceramic fiber boards for thermal insulation. As soon as sample or batch dimensions passed a certain limit the
thermal insulation box had to be equipped by mm-wave susceptors. Since such a packing of samples will never be feasible in a industrial production process the problem of energy losses form sample surfaces was faced by the installation of a hybrid heating system [ 111. This system consists of a box made f?om mullite ceramic fiber boards and eight 300 Watts MoSizheating elements, distributed along the side walls (Fig.1). This is a small design with a usable volume of 1 liter only, but it allows the insertion in the existing mm-wave applicator without large construc-tive changes. The maximum working tempera-ture of this oven is 1650 "C at various gas atmospheres. To investigate the sintering behavior in the mmwave field more closely a dilatometer was designed and built into the mm-wave applicator. Therefore a commercially available dilatometer from the company Linseis, Germany with an inductive displacement transducer and an Alumina sensor rod was modified in an appropriate way to avoid microwave leaking out of the applicator. It allows in-situ measurements of Ihe sample's extension and shrinkage during complete heating cycles up to temperatures of 1600°C.
Nanoceramics The development of new techniques for the production of nanocrystalline oxide ceramic powders with average grain sizes smaller than 50 nm provides the opportunity for the production of ceramic materials with novel mechanical properties, such as low temperature creep and plastic deformation at temperatures well below the melting point [12]. But such novel material properties can only be exploited if it is possible to overcome the problem of grain growth during the sintering process. With conventional sintering techniques it is practically impossible to avoid strong grain growth due to the need of long processing time. The application of high power mm-waves for sintering seems to be a promising technical approach to high-density nanocrystallineceramics. The material which have been investigated is Yttriastabilized Zirconia with an average grain sizes of 36 38 nm (determined by BET surface measurement). A comparison of the dilatometer curves obtained with equal processing parameters by conventional, resistant heating to the curves obtained with mm-wave sintered samples shows an enhanced densification in the mmwave field. The shrinkage of the sample sintered in the mm-wave field starts earlier, about 150°C lower than the conventionally sintered one (Fig. 2).
459
1
Piezoceramics ._. .... conventional
-3OGHz
4
1
Shrinkage dL/b [%]
One of the most interesting class of functional ceramics are piezoceramics based on the system leiidzirconate-titanate(PZT). As lead oxide has a rather high vapor pressure and tends to evaporate during the sintering process shortening of the processing time and a possible reduction of the sintering tempera-ture will help to reduce this difficulty. As can be seen from Figure 4, the comparison of dilatometer curves conventional and at 30 GHz shows a clear reduction of sintering temperature of 150°C.
-20 -2 4
I 600 750 900 1050 1200 Temperature ["C]
Fig. 2: Dilatometer curves of zirconia, 3OGHz and conventional,heating rate 5 "C/min Densities of more than 90 %TD were obtained at 1100°C and at 1200°C densities of 97 - 99 %TD were achieved with millimeter-waves. The average grain sizes in the finished samples which were determined by peak broadening of X-ray diffraction patterns are less than 100 nm, in spite of several hundred nanometers in conventional sintered samples. This was also confumed by characterization of the microstructure as can be seen following (Fig.3). With the mm-wave sintering technique it is possible to densify Zirconia samples fully with a nanoscale diameters of grains.
The advantage is even more effective for sintering samples with a higher surface to volume ratio. An example is a microstructured part as shown in Figure 5. The microstructure part was manufactured by hot moulding technique. This PZT piezoarray consists of 75 x 75 columns with a single column cross-section of 100 x 100 pm and 400 pheight. Such arrays are used for the production of 1-3-composites which x e ultrasonic sensor elements designed for isostriain thickness vibration at 3.5 MHz.
....... ~
conventional -3OGHz
1
0 -4
Shrinkage
-8
dLIb [%]
-12 -1 6
-20 400
600
800
1000 1200
Temperature ["C]
Fig. 4: Dilatometer curves of PZT, 3OGHz and conventional, heating rate 5 "C/min
Fig. 3: Zirconia, sintered 1200"C, no dwell
460
First experiments revealed that using mm-waves the sintering temperature of the conventional process can be reduced from 1200 "C down to 1100 "C and the dwelling time from 60 min. down to 10 min. only, in order to get similar densities. This leads to a reduction of PbO losses as low as 0.5 %. The performance of 1-3composites revealed a coupling coefficient of 74% compared to 70 % of those equipped with conventional
sintered arrays. This coupling behaviour and the low impedance enable a very good performance.
MRS Symp. Proc., Vol. 430, Microwave Processing of Materials V, ed. by M.F. Iskander J.O. Kiggans, J.-Ch. Bolomey, Pittsburgh, PA, (1996) 157-162.
7. H.D.
Kimrey, M.A. Janney, P.F. Becher; Techniques for ceramic sintering using microwave energy; Conf. Digest 12" Intern. Conf. on Infrared and Millimeter Waves, ed. R.J. Temkin, Orlando, Florida, (1987)136-137.
8. Yu. Bykov, A. Eremeev, V. Flyagin, V.Kaurov, A. Kuftin, A. Luchinin, 0. Malygin, I. Plomikov, V. Zapevalov, L. Feher, M. Kuntze, G. Link, M. Thumm, Gyrotron Installation for millimeter-wave processing of materials, Vakuumelektronik Displays form VDE-Verlag GmbH, Berlin Offenbach, Vol. 132, (1995) 103-108.
9. L. Feher, G. Link, M. Thumm, Eletrothermal Fig. 5: PZT array, sintered 120O0C, 20 min
SUMMARY AND CONCLUSIONS Sintering tests in a gyrotron installation at the Forschungszentrum Karlsruhe, IHM, have been performed in order to study the densification behavior of several types of structural and functional ceramics under mm-wave radiation. For the investigated types of ceramics it was shown that mm-wave sintering is superior to conventional heat treatment. In order to reach the desired densities with the mm-wave technique the necessary processing times are markedly shorter and the sintering temperatures are lower.
Heating Model for Microwave/Hybrid-Processed Materials, Proceedings of 24'h Infrared and Millimeter Waves Conference, Monterey, (1999), F-B5
10. L. Feher; G. Link, M. Thumm; Optimized Design of an Industrial Millimeter Wave Applicator for Homogeneous Processing of Ceramic Charges; Conf. Roc. of Microwave and High Frequency Heating 1997, Fermo, Italy; ed. by A Breccia, R. De Leo, A.C. Metaxas, Bologna, Italy; (1997) 442446.
11. G. Link, L. Feher, M. Thumm, Hot Wall 30 GHz Cavity for Homogeneous High Temperature Heating; Roc. of 7'h Int. Conf on Microwave and High Frequency Heating, Valencia, ed. J.M. CatalACivera et al., (1999) 165-168.
12. J. Karch, R. Birringer, H. Gleiter, Ceramics ductile
REFERENCES
at low temperature, Nature, Vol. 330 (lo), (1987) 556.
1. I.J. Chabinsky; Application of microwave energy past present and future, in Microwave Processing of Materials MRS Symp. Roc.Vol. 124; (1988) 1730.
2. J.M. Osepchuk; Trans IEEE on Microwave Theory and Technique, Vol. MlT-32[9], (1984) 26.
3. M.L. Levinson US Patent No. 3585258 (1971). 4. W.H. Sutton, BM.H.Brooks, 1.J.Chabinsky; Microwave Processing of Materials. MRS Symp. Roc. Vol. 124 Pittsburgh. (1998)287-295.
5 . M.A. Janney, H.D. Kimrey; Diffusion Controlled Processes in Microwave-Fired Oxide Ceramics; MRS Symp. Proc., Vol. 189, Microwave Processing of Materials 11, ed. by W.B. Snyder, W.H. Sutton, Pittsburgh, PA, (1991) 215-227.
6 . G. Link, V. Ivanov, S. Paranin, V. Khrusov, R. Bohme, G. Miiller, G. Schumacher, M. Thumm; A Comparison of mm-Wave Sintering and Fast Conventional Sintering of Nanocrystalline Al203;
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INTERNALLY COOLED MONOLITHIC SILICON NITRIDE AEROSPACE COMPONENTS Jonathan E. Best*, James D. Cawley*, Ramakrishna T, Bhatt** and Dennis S. Fox*** (*) Materials Science and Engineering, Case Western Reserve University
10900 Euclid Avenue, Cleveland, Ohio 44106 (**)Army Research Laboratory Vehicle Tech. Dir., 21000 Brookpark Rd. Cleveland, Ohio 44135 (***) NASA Glenn Research Center, 21000 Brookpark Rd. Cleveland, Ohio 44135
ABSTRACT A set of rapid prototyping (RP) processes have been combined with gelcasting to make ceramic aerospace components that contain internal cooling geometry. A mold and core combination is made using a MM6Pro (Sanders Prototyping, Inc.) and SLA-250/40 (3Dsystems, Inc.). The MM6Pro produces cores from ProtoBuildTM wax that are dissolved in room temperature ethanol following gelcasting. The SLA250/40 yields epoxy/acrylate reusable molds. Parts produced by this method include two types of specimens containing a high density of thin long cooling channels, thin-walled cylinders and plates, as well as a model hollow airfoil shape that can be used for burner rig evaluation of coatings. Both uncoated and mullitecoated hollow airfoils has been tested in a Mach 0.3 burner rig with cooling air demonstrating internal cooling and comfirming the effectiveness of mullite coatings.
INTRODUCTION The efficiency of a gas turbine engine is increased by raising the turbine inlet temperature and the maximum temperature is ultimately limited by material properties. In order to approach service near 90% of their melting point, superalloy blades and vanes have become elegantly engineered’. In part this is microstructural, for example, the progression from equiaxed to directionally solidified to single crystal and the gamma gamma-prime strengthening. But of equal importance is the development of coring technologies that permit internal cooling and the application of environmentally resistant coatings. The former allows the surface temperature of the component to kept to a maximum that is significantly below the ambient whereas the latter mitigates the effects of oxidation and other corrosion processes as well as providing a thermal barrier. Much of the work on producing ceramic materials has focused of the development of material properties such as resistance to thermal shock and oxidation, high temperature strength and stiffness, and creep resistance. And the results have been singularly successful. In particular, both commercial and research grades silicon nitrides (usually alloys containing rare earth oxides in
combination with aluminosilicates) are available which can tolerate sudden quenches in excess of 1000°C, have strengths greater than 800 MPa, and which tolerate service operating temperatures up to 13OOOC. In-situ reinforced silicon nitride formulations have been developed with fracture toughnesses in excess of 8 MPadm. The combination of such properties are immediately suggestive of application in gas turbine (and rocket) applications’. Furthermore, the effectiveness of internal cooling has been demonstrated by Tsuchiya et al’ who have shown that the it is reasonable to for the surface temperature of silicon nitride components in a 150OOC environment to be kept below 1300°C through the use of internal air cooling, which require modest flows. Engineering the surface is suggested by work that has been done demonstrating the use of plasma-sprayed mullite to produce thermal and environmental barrier coatings on silicon carbide and silicon nitride4. The work reported herein represents an effort to develop a process strategy to prototype silicon nitride parts capable of internal cooling that is robust and easy to implement. Conceptually, the process is quite straightforward: molds and cores are produced by two different commercial RP processes and assembled; a gelcasting slurry is prepared and forced into the cavity under low pressure; after gelling, the part is liberated through a combination of dissolution to remove the cores and diassembly to remove the reusable molds. The resultant wet part is dried and fired via conventional means. Details of the process are given elsewhere’. Gelcasting is a process developed at Oak Ridge National Laboratory6 and it is particularly amenable to use with soft tooling, such as that produced by polymerbased RP processes because only very low pressures are employed. In fact, Jamalabad et al. suggested the use of gelcasting in fugitive molds sometime ago’. The approach taken in this work differs in that a combination of RP methods is used, the design of the molding system is different, and that a process for setting the gel, immersion is a poly(ethy1ene glycol) (PEG), is used incorporated. The particular implementation of gelcasting used in this work uses water-soluble organic materials. Two monomers are employed: methacrylamide (MAM) which polymerizes as a linear chain and N,N’-
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methylenebisacrylamide (MBAM) which is a crosslinker. Just prior to casting an initiator, N,N,N’,N’- tetramethylethylenediamine (TEMED), and a catalyst, ammonium peroxydisulfate (APS), are mixed into the slurry to commence the gelation process. Standard stereolithography was used to produce reusable molds. A SLA250/40 was used to build epoxy/acrylate parts that survived molding and subsequent immersion in ethanol and PEG without noticeable damage or distortion. Dozens of casts were made using the same epoxy/acrylate parts. The Sanders Prototyping process was used to produce cores and mold lines which were dissolved, after casting, by soaking in room temperature ethanol for several hours. A ModelMaker MM6Pro was used. The Sanders process employs two materials to !kilt a part. The actual part is built out of a “green wax” (ProtoBuildTM)and the surrounding supports are built out of a “red wax” (ProtoSupportTM). The waxes are melted and each layer of the part is deposited using two separate ink jets to spray droplets. A mechanical grinder levels off each layer to the correct thickness. These process steps continue until the entire part is built. Upon completion the red supporting wax can be selectively removed by heating to 40-50°C in a VS-0 bath. Beyond the fact that the green wax can be dissolved with ethanol at room temperature, this method is particularly suited to the production of cores or small characteristic dimension or requiring fine details. All CAD work was carried out using Rhino3D (Robert McNeel and Associates) a NURBS package that runs on PentiumTMclass PCs.
PROCEDURE Three types of parts were built. The first was a reverse-engineered model airfoil shape based on archival superalloy parts left over from the original thermal barrier coating development program at NASA Lewis Research Center (now NASA Glenn Research Laboratory). The part allows internal cooling and was readily coated with plasma sprayed mullite. The other two were demonstration parts to permit the a critical evaluation of the possibility for using thin cooling channels of long length within the walls of thin-walled plates and cylindrical sections. In initial experiments, molds were entirely epoxy/acrylate and a mold release was used to effect separation of the gelled part from the mold. Although often successful, distortion due to demolding stresses was enough of a nuisance that soluble mold linings were developed and used with great success. Mold linings were made by the exact same process as the cores and so were simultaneously removed. The net result is that the space filled by the slurry was entirely bounded by surfaces made up of the Protobuild material. This was true for all of the parts discussed. Figure 1 shows the CAD representation of the molding system used to make the thin-walled internally cooled hollow cylinder. Similar mold systems are used for other components. In practice, the joints of the assembled mold were sealed with a soft patching wax (Kindt-Collins) and held together with teflon tape as
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needed. After the casting has gelled, the entire mold is immersed in ethanol to dissolve the wax components of the molds. Leachable paths are designed such that all the wax parts are connected to the external surface. The wax linings and cores dissolve away in a few hours. The part is then submerged in PEG for drying. As the gating structure would restrict drying shrinkage creating undesirable stresses within the part, it is trimmed from the part prior to drying. After the part is dried, typical ceramic firing procedures are employed. The process steps are illustrated in Fig. 2 using a different mold/core set to illustrate each.
FABRICATED PARTS The results of the molding process as well represented by the results shown in Fig. 3 for the cylinder. The cast part contains some flashing that must be manually removed, but the overall shape and size is accurate and the high density of small diameter holes produced by the cores are excellent. The pathway of the cooling holes is visible in the fired piece because of the locally thin walls. Very similar results were obtained with the flat plate. The simulated airfoil specimen is documented in Figs. 4 and 5 . Figure 4 illustrates the results of the reverse engineering process and Fig. 5 shows the resultant silicon nitride version, which was an excellent geometrical match for the original casting. Mach 0.3 jet fuel burner rig testing was conducted on superalloy, uncoated silicon nitride and mullite-coated silicon nitride specimens. A typical test condition is illustrated in the photograph appearing in Fig. 6. Quantitatively results indicate that cooling air produced the expected comparable reductions in surface temperature, up to 130°C for flowrates of approximately 1.5 lit/s.
CONCLUSIONS The fabrication of internally cooled monolithic silicon nitride aerospace components has been demonstrated. Gelcasting coupled with rapid prototyping technology allows for the production of these components which range from thin-walled shapes with cooling channels to simpler hollow shapes. The internal geometry of these parts is made possible by wax cores produced by the Sanders Prototyping process, which can be dissolved from wet gelcast parts with ethanol. This technique brings the same flexibility of shape forming associated with metal casting to the realm of ceramics processing. A hollow silicon nitride airfoil produced with this process has been tested with cooling air in a Mach 0.3 burner rig. Cooling air reduced the surface temperature up to 130°C below the uncooled condition of 1185OC.
ACKNOWLEDGEMENTS Z. Liu, CWRU, was of enormous help in general ceramic processing and rapid prototyping. R. Babuder, CWRU/NASA, was similarly invaluable in his help with binder burnout and firing. Mike Cuy, NASA, assisted with the burner rig testing. The work was funded under NASNCWRU Cooperative Agreement on Ceramic Processing Grant # NAS3-404.
REFERENCES 1.
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C. T. Sims, “A History of Superalloys for Superalloy Metallurgists,” pg. 399 in Superalloys 1984, edited by M. Gell, C. S. Kortovich, R. H. Bricknell, W. B. Kent, and J. F. Radavich, AIME, 1984. Pollinger, J.P., Progress in fabrication of silicon nitride structural components for turbomachinery applications. ASME Paper 96-GT-347. Tsuchiya, T., Furuse, Y., Yoshino, S., Chikami, R., Tsukagoshi, K., Mori, M., Development of aircooled ceramic nozzles for a power generating gas turbine. ASME Paper 95-GT-105. K. N. Lee and R. A. Miller, “Development and Environmental Durability of Mullite and MulliteNSZ Dual Layer Coatings for S i c and Si3N4Ceramics,” Surface and Coatings Technology, 86-87 [ 1-31 142-8 (1996). J. E. Best, J. D. Cawley, R. Bhatt, and D. Fox, “Fabrication and Testing of Internally Cooled Monolithic Silicon Nitride Aerospace Components, J. Am. Ceram. Soc., submitted. 0.0.Omatete, M. A. Janney, S. D. Nunn, “Gelcasting: from Laboratory Development toward Industrial Production. Journal of the European Ceramic Society 17 (1997) 407-413. Jamalabad, V.R., Whalen, P.J., Pollinger, J., “Gelcast molding with rapid prototyped fugitive molds,” , pp. 71-78 in Proceedings of the Solid Freeform Fabrication Symposium, ed. D.L. Bourell, J.J. Beaman, H.L. Marcus, R. H. Crawford, J. W. Barlow. The University of Texas at Austin, Austin, Texas, 1996.
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Figure 1 shows the CAD rendition of the assembly of the thin-walled internally cooled cylinder mold system. The individual parts can be seen in the center overview from right to left: inner gating, inner lining, core half, outer lining half, outer gating half. The design feature are highlighted numerically: 1) registration features, 2) slurry casting tube, 3) bottom filling tray, 4)flow path slats in core, 5 ) overflow space, 6) gate removal is accomplished by cutting along this plane, 7) uppermost of three slots for dissolution of outer lining, and 8) leachable ethanol path for core dissolution (square cutouts are obsolete flow paths).
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Figure 2 shows the processing steps of the soluble core gelcasting process: a) filling the mold, b) dissolving ProtoBuild mold components in ethanol, and c) the final green part separated from, but still within the enoxv/ acrvlate mold.
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Figure 3 shows various views of the thin-walled cooled cylinder: a) untrimmed green state, b) trimmed green state, and c) fired component. Diameter of coin is 19.05 mm.
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Figure 4 Reverse engineering of a internally cooled model airfoil test specimen. a) Deduced CAD file from measurements on original casting. b) and c) Two views of the original casting and a stereolithography model made for verification purposes.
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Figure 5 Two views of the green hollow simulated airfoil. Diameter of coin is 19.05 mm.
Figure 6 Silicon nitride airfoil in Mach 0.3 jet fuel burner flame with.3 lids of cooling air. Superalloy casting and mullite-coated airfoils also were tested.
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PREPARATION AND HIGH TEMPERATURE STRENGTH OF G&A1209/MgO COMPOSITES M. Shimada, T. Sakamoto and H. Yamane Institute for Advanced Materials Processing, Tohoku University Sendai 980-8577, Japan
ABSTRACT Gd4Al~OdMg0 composite (MgO: 85 ~01%) was prepared from the mixed powders of Gd203, A1203and MgO at 1600 "C for 5 h in air. The relative density of the composite was 99.1%. The X-ray powder diffraction pattern of the composite was identified by the mixture of Gd4A1209and MgO. No other peaks were observed. The composite had a uniform microstructure with grain sizes of 1 = 2 pm for Gd4AlZOYandof about 4 pm for MgO. A bending strength of 400 MPa measured at 1100 "C on cooling from 1400 "C was 1.6 times higher than that measured at 1100 "C on heating. This could be explained by the transformation toughening mechanism of the high temperature phase Gd4Al2aat 110 "C cooling.
INTRODUCTION Recently, there has been a great effort to overcome the intrinsically brittle nature of ceramics in order to apply ceramic materials to engineering purposes. Since the fracture toughness of ceramic materials is not high enough, their use under high stress is limited. It is highly desirable to improve the mechanical properties of ceramic materials. Transformation toughening is one effective approach to improving fracture toughness and fracture strength of brittle ceramics. Transformation toughening requires the stress-induced martensitic transformation of particles, as well known in tetragonal zirconia ceramics (1). The main purpose of our study is to find new oxide ceramics with stress induced phase transformation. Rare-earth aluminate, %A1209 ( R rare-earth element) is monoclinic with space group P2Jc at room temperature (2). Yamane et. Al. (3) reported that YdA11Q showed a reversible thermal phase transformation from low temperature monoclinic to high temperature monoclinic phases at 1377 "C. The results of hightemperature X-ray powder diffraction showed that the unit cell volume of the high temperature phase Y4A1209 measured at 1400 "C on cooling from 1450 "C was 0.5% smaller than that of the low-temperature phase at the same temperature of 1400 "C on heating from room temperature (3). This volume change was in accordance with the results of dilatometry. The phase transformation of YdA1109 was concluded to be diffusionless by comparison of the crystal structures of the high- and lowtemperature phases determined by high-temperature neutron diffraction (4). Gd~A1209showed a reversible thermal phase transformation and thermal hysteresis at
around 1100 "C is about 200 "C lower than that of y4AlZa. The present study attempts to stabilize the hightemperature phase of Gd4A1209 by preparing composites with MgO and examines the effect of phase transformation toughening by measuring the high temperature fracture strength.
EXPERIMENTAL Powders of Gd203, A l Z 0 3 and MgO were used as starting materials. These powders were weighed in the appropriate proportion of G&A1209 and 15 vol% Gd4Al209/85 vol% MgO and mixed in a plastic vessel for 24 h by a wet ball milling with ethyl alcohol and alumina balls. After drying at 80 "C for 5 h in air, the mixed powders were uniaxially pressed at 30 MPa, and then isostatically pressed at 200 MPa to form pellets (5x30~50mm). The compact pellets were sintered at 1600 "C for 5 h in air. The sintered pellets were cut into rectangular coupons 2 x 4 ~ 2 0mm. The crystal structure was studied by X-ray diffraction (XRD). The microstructure was observed by scanning electron microscopy (SEM). The thermal behaviour of Gd4A1109and Gd4AlZO&ig0 composites was characterized from room temperature to 1500 "C by high temperature XRD and thermal mechanical analysis (TMA). The bending strength of Gd4A120&lg0 composites was measured from room temperature to 1400 "C by a three-point bending test with a crosshead sped of 0.5 mm/min and a span length of 10 mm.
RESULTS AND DISCUSSION The X-ray powder diffraction pattern of Gd4A120&ig0 composite (85 vol% MgO) was identified by the mixtures of Gd4A12OYand MgO phases. No other peaks were observed. The Scanning electron micrograph of polished surface of composite is shown in Fig. 1. In this figure, with area is G d 4 A l 2 a and black area is MgO. As seen in this figure, the composite shows a uniform microstructure with grain sizes of 1 = 2 pm for Gd4A12O9and of about 3 - 4 pm for MgO. The relative density of composite was 99.l%, which indicates that composite was well densified by heating at 1600 "C for 5 in air. The results of high temperature XRD showed that the lattice parameters of Gd4Al2O9shrank0.2 = 0.3% along 469
the c axis at the phase transformation fiom the lowtemperature phase. Since the XRD patterns of hightemperature phase was similar with those of lowtemperature phase, it is considered that the crystal structure of high-temperature phase belongs to a monoclinic structure. Anisotropy of thermal expansion was observed below the phase transition temperature. The average thermal expansion of the a axis was about twice as large as that of the c axis.
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Fig. 3 Temperature dependence of bending strength of Gd4A120&lg0 composite
Fig 1 Polished surface of G&A120&lg0 composite Temperature dependence of linear thermal expansion of MgO, Gd4A1209/Mg0composite is shown in Fig. 2. A large thermal hysteresis was found in Gd4A1209.From the results of linear thermal expansion of Gd4A1209,the volume contraction started at 1100 "C and finished at 1230 "C on heating process, and on cooling process the volume expansion started at 1000 "C and finished at 880 "C. The thermal linear expansion of the composite was between those of G&A1209and MgO. No thermal hysteresis like Gd4A1209 was found in composites Gd4A1209/Mg0due to the small amount of Gd4A1209 content.
The temperature dependence of bending strength measured for the Gd4A1209/Mg0composites is shown in Fig. 3. The bending strengths were 190 MPa at roomtemperature and about 280 MPa at 700 "C on both heating and cooling processes. At 1100 "C, a bending strength of 230 MPa was measured on heating, and a higher strength of 400 MPa was measured on cooling. The bending strength measured at 1 100 "C on cooling was 1.6 times as much as the bending strength measured at the same temperature on heating. This improvement could be explained by the transformation toughening mechanism of the stabilized hightemperature phase of G&A1209at this temperature.
ACKNOWLEDGMENTS This work is partly supported by NED0 under the Synergy Ceramics Project of The Industrial Science and Technology Frontier Program promoted by AIST, MITI, Japan.
REFERENCES
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(1) R. C. Garvie, R. H. Hannink and R. T. Pascoe, Ceramic Steel?, Nature 258 (1975) 703-704 (2) J. W. Reed and A. B. Chase, The Unit Cell and Space Group of 2Y203.A1203, Acta Crystallogr. 15 (1962) 812 H. Yamane, M. Omori, A. Okubo and T. Hirai, HighTemperature Phase Transition of Y4A120g,J. Amer. Ceram. SOC.76 (1993) 2382-2384 H.Yamane, M. Shimada and B. A. Hunter, HighTemperature Neutron Diffraction Study of Y4A1209, J. Solid State Chem. 141 (1998) 466-474
CHARACTERIZATION OF IN SITU SIC-BN COMPOSITES Guo-Jun Zhang* and Tatsuki Ohji Synergy Ceramics Laboratory, National Industrial Research Institute of Nagoya, Nagoya, Aichi 463-8687, Japan ABSTRACT Sic-BN composites were prepared by in situ reaction using Si,N4, B4C and C as reactants. Adding Sic powder into the reactants controlled the content of BN in the composite. For comparison, monolithic S i c and SIC-BN composites with the same phase compositions were produced by conventional process. The in situ process was advantageous for obtaining better composites with fine grain size, homogeneously and isotropically distributed microstructures. The elastic modulus decreased markedly with increasing BN content and obeyed well both to the exponential and polynomial rules. The bending strength increased slightly with increasing BN content up to 15 vol%, and then gradually decreased; the decrease was accelerated above 35 ~01%.The results were explained by a percolation network formation of the BN phase. The fracture toughness decreased with increasing BN content. It can be excepted that these in situ Sic-BN composites have excellent mechanical strain tolerance and thermal shock resistance.
INTRODUCTION Silicon carbide (Sic) is a potential candidate material for structural applications at high temperatures. In many cases of application ceramic components are assembled with other components made of metals to form a hybrid structural part. The failure of such hybrid part often occurs in the ceramic components at the connecting section due to the mismatch of thermal expansion coeffi-cients and elastic strains in different materials. To avoid these fractures, it is required that the elastic modulus of Sic ceramic material is reduced. BN ceramics, which shows very low elastic modulus among ceramic materials, have been widely used as the second phase to reduce the elastic modulus and improve the thermal shock resistance of monolithic ceramic materials, such as in Si3N4-BN[l-31 and SIC-BN [4-61 composites. Unfortunately, the bending strength of BN-containing composites
generally decreased with increasing the BN content compared to the monolithic ceramics. One of the main reasons is that the poor sinteribility of BN affects densification behavior of BN composite. The other is that because BN is weak, BN agglomerates or large BN particles/platelets may act as fracture flaws. This means that BN dispersoids with fine particle size and homogeneous distribution are the key factors to obtain high strength composites. Reactions are also used to synthesize BN-based composites. For example, Coblenz and Lewis [l] used the in situ reactions of Si3N4+ B203,A1N + B203or Si3N4+ A1N + B,O, to produce SO2-BN, A1203-BN and mullite-BN composites, respectively. Kusunose et a1 [3] prepared Si3N4/BN nanocomposites by in situ formation of nanosized h-BN coated on a-Si3N4.particleswith a mixture of boric acid and urea in hydrogen gas. In addition, in the Si-B-C-N system, Riedel et a1 [7,8] developed a silicoboron carbonitride ceramic material, which is stable to 2000 OC, from polymeric boron-containing polysilylcarbidi-imideprecursor. By in situ process, one can generally obtain composites with finer and more homogeneous microstructures, higher chemical and microstructural stability at high temperature, and better mechanical properties than the conventionally processed ones. Recently the authors proposed the following reaction to synthesize SIC-BN composite: Si3N4+ B4C + 2C = 3SiC + 4BN.. ....(1) The weight percents of Sic and BN in the obtained composite according to reaction (1) are 54.79 % and 45.21 %, respectively, and the volume percents are 46.29 % and 53.71 %, respectively. By adding Sic powder into the reactants, the content of S i c in the composite can be adjusted. Here the preparation of the in situ composites will be firstly introduced, then the mechanical properties of the obtained composites will be characterized.
EXPERIMENTAL PROCEDURE The starting powders were Si3N4 (E-10 grade, Ube Industries, Japan), B4C (F1 grade,
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particle size lpm, Denki Kagaku Kogyo Co. Ltd., Japan), C (2600# grade, particle size 13 nm, Mitsubishi Chemical Co., Japan), A1,0, (particle size O.lpm, purity 99.99%, Daimei Chemical Co., Japan), Y203 (particle size 1.06pm, purity 99.9%, Shin-etsu Chemical Co., Japan), a-Sic (A2 grade, mean particle size 0.63pm, purity 99.1%, Showa Denko Co., Japan) and h-BN (SP-2 grade, particle size 4pm, Denki Kagaku Kogyo Co.Ltd., Japan). For improving the sinterability in the reaction process, 10 wt% A1,03-Y203 additives (7:3 mixture of Al,03 and Y203)related to the Sic contents in the composites were used. Composites with no A1203-Y203 additives were also produced for companson. About sample designation, for example A-B25 means that the BN content in the composite is 25 vol% and with A1203-Y,03additives. Whereas, N-B25 means that the BN content in the composite is 25 vol% and without Al,03-Y,03 additives. However, the sample only from reaction (l), i.e. without Sic addition, is denoted as B55. The powders were ball-milled for 24 h in ethanol using Zr02(Y203)balls and subsequently dried. Hot pressing of the mixed powders was conducted in an argon atmosphere under 30 MPa in a graphite die with BN coating. For studying the phase formation mechanism, hot press at different temperatures for 60 min was conducted and then the phase composition was determined by X-ray diffraction (XRD) using CuKa radiation. The final composites were hot pressed at 2000 "C for 60 min. The linear dimensional change of the specimens during hot pressing was measured by a displacement gauge. The obtained data were used to calculate the densification behavior, which is represented by the relative density to the final phase composition of the composite. Reaction behavior was investigated by differential thermal analysis (DTA) up to 1700 "C using a heating rate of 10 Wmin. Three-point bending strength was tested on bars of 2.5 mm x 3 mm x 20 mm using a span of 16 mm and a crosshead speed of 0.5 mdmin. Fracture toughness was measured by SENB method. The strength and toughness data were average of 5 measurements. Young's modulus, E, parallel to the hot pressing direction, was measured by the pulse echo method. The microstructure of the composite was observed by scanning electron microscopy (SEM). RESULTS AND DISCUSSION
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Effect of AI2O3-Y203Addition on Reaction Process The possible reactions in the Si3N4-B4C-C system during the hot pressing process are: Si3N4+ B,C + 2C = 3SiC + 4BN.. ....(1) Si3N4+ 3C 3SiC + 2N2............(2) Si3N4+ 3B4C= 3SiC + 2N2+ 12B.. . (3) Si3N, + 3B4C= 3SiC + 4BN + 8B.. . (4) 4B + C = B4C......... .................. (5) Si3N4+ 3C +4B = 3SiC + 4BN.. .... (6) Figure 1 shows the calculated enthalpies AHo and free enthalpies AG: of these reactions using the thermodynamic data from reference [9]. It can be seen that the total reaction (1) is exothermic and can take place thermodynamically in common temperature range (AG; I 0 when T I 9944 "C). In the case of reaction (2), it will occur at temperatures higher than 1488 "C and give out nitrogen gas (N,). However, upon thenno-dynamics the coexistence of B4C will inhibit the reaction (2) since the reaction (4) can take place much easily (i.e. much lower free enthalpy) in the experimental temperature range. The byproduct B in the reaction (4) is not released easily compared to the by-product N, gas in reaction (2). This should be the thermodynamic basis for preparing the SiCBN in situ composite according to the proposed reaction (1). Furthermore the byproduct B in the reaction (4) can act as reactant either in the reactions (5) or (6). +
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Fig. 1 Free enthalpy curves of related reactions The XRD patterns for the reaction (1) with (A-B specimen) or without (N-B specimen) A1203-Y203additives taking place at 1000 "C, 1200"C, 1450"C, 1550"C, 1700"C, and2000 "C as well as for the green body are shown in Fig. 2. It can be seen that the reaction did not take place below 1200 "C in spite of the addition of A1203-Y@3 additives. For A-B
specimen at temperatures higher than 1450 "C, the reaction took place gradually, but did not complete until the temperature reached 1700 "C. For N-B specimen, the reaction almost did not take place at temperatures below 1500 "C and occurred gradually from 1700 "C. It can be concluded that the addition of Al,O,-Y,O, additives improved the reaction process and the crystallization of BN phase. The final products of these two specimens only showed Sic and BN phases. In all the XRD patterns, no other main diffraction peaks can be seen except the starting reactants a-Si,N4 and B4C and the products p-Sic and h-BN. There also
existed low peaks indicated as a-Sic (6H and 4H types). The starting reactant C is in amorphous and there is no peak in the patterns. Accordingly, the reaction sequence was suggested as follows: Si,N, + 3B4C = 3SiC + 4BN + 8B 4B + C = B4C or Si,N, + 3C +4B = 3SiC + 4BN Some small peaks appeared when the sample was hot pressed at 2000 "C, comparing the XRD pattern with that of 1700 "C. These peaks were difficult to identify, however the content of the related impurity phase should be low.
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2-theta (deg) Fig.2 XRD patterns of the reaction (1) with or without Al,O,-Y,O, additives DTA has been performed for A-B55 specimen using heating rate of 10 "C/min in Ar atmosphere; the curve is shown in Fig. 3. It can be concluded fi-om this curve that the reaction begins to take place at 1400 "C and this result is coincident with the XRD analysis
discussed in the above section. On the other hand, it can be seen that the reaction took place gradually from the broad exothermic peak on the curve. According to the calculation method of adiabatic temperature (T,) of combustion [ 101, the Tad of the
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exothermic reaction (1) is 1799 K. However, because the reaction took place gradually during the temperature rising, and all of the reactants or products of this reaction have high melting or evaporating points, the exothermic character of this reaction is considered to be not harmful but helpful to the progress of the reaction. It is necessary to note that because of the temperature limitation of the DTA apparatus, the analysis was stopped at 1700 "C. The densification curves of the SIC-BN in situ composites A-BO, A-B15, A-B35 and AB55 are shown in Fig.4. The pressure and temperature during hot pressing process are 120 r
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c
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contribute to the densification. B4C and C are well known as additives for improving the sinterability of Sic [11,12]. Also, in the present reaction synthesis system, the reactants B4C and C should contribute to the densification process. This is very likely the reason why the densification behavior was improved for the specimen with low in situ BN content (B15 specimen). Because of the poor sinterability of BN, the densification behavior became worse for more BN content in the in situ composites. In addition, it can also be seen from Fig.4 that the densification process almost finished after 30 min at 2000 "C. This suggests that the hot pressing time can be shorter to obtain a dense in situ Sic-BN composite with fine microstructure.
1500 1000
500
0
Fig.4 Densification behaviors also shown in this figure. Compared with the monolithic Sic (BOYi.e. BN content 0 vol%), the densification behavior of A-B 15 specimen was apparently improved. However, with increasing the content of in situ BN, the densification decreased. Nevertheless, the final relative density after heating for 60 min at 2000 "C was higher that 95% for all the specimens. As shown in Fig. 3 the synthesis reaction has already finished after heating for 60 min at 1700 "C. That is, hot pressing at higher temperatures than 1700 "C only
474
Microstructures and Mechanical Property Evaluation The SEM micrograph of the fracture surface A-B15 are shown in Fig.5. The in situ formed hexagonal BN is in flake shape and the Sic phase is in equiaxial shape. The hot press direction was vertical in these pictures. It should be noted that the in situ formed BN flakes are homogeneously and isotropically distributed. These BN flakes are about 1 pm in length and 0.1 pm in thickness and are located at boundaries between Sic grains. Such microstructure is illustrated schematically in Fig. 6. The size and shape of the BN flakes do not depend on the BN content, nor on the initial particle size of reactants of the in situ reaction. For S i c phase, there is not obvious change in particle size (2 to 3 pm) when the BN content is lower than 15 ~01%.However, as the BN content increased in a range above 25 vol%, the grain size of Sic decreased and reached about 1 pm in B55 composite. That is, above a certain value of BN content, the grain boundary BN flakes inhibit grain growth of Sic. This phenomenon is plausibly attributable to a percolation network formation of the BN phase. It has been reported that the percolation threshold is different for materials with different microstructure, but that is around 20 vol% for randomly distributed two-phase system. In the present case, even in B5 specimen with only 5 vol% BN, two-flake pairs of BN can be seen. In B15 specimen, more BN flake pairs appeared, however percolation network has not formed yet. In the specimens with BN content higher than 35 vol%, percolation network can be clearly seen. Consequently, the percolation threshold Vc for BN phase in this in situ SIC-BN composites can be reasonably taken as 25 ~01%.It can be considered that the grain growth of Sic is inhibited by a percolation network of the BN
-
phase, which presumable forms at volumes above the percolation threshold.
or, by a polynominal function as: E = E,(1-a,V + a,V2)..................... (3) where E,, b, a, and a, are constants and V is the volume fraction of the second phase. For the present SIC-BN in situ composites, the exponential and polynomial approximate curves are given by: E = 432.39exp(-2.6139V)................(4) E = 421.61(1 - 2.1674V + 1.4271V2)...(5) The respective reliabilities (R2) are very high, 0.9955 and 0.9976, indicating excellent fitness of both the approximations. 1000
I
5.5
Fig.5 Fracture surface of A-B 15 specimen /Grain boundary
I 0
20 40 BN content (~01%)
60
Fig.7 Mechanical properties
"-
Fig.6 Schematic illustration of the microstructure Figure 7 illustrates the mechanical properties of the Sic-BN in situ composites as a function of BN content. The upper bound (Voigt bound), E, (E, = 1 EiVi, i stands for the component phase, V the volume percent of i phase and the lower bound (Reuss bound), EL (E; = ZEi-'Vi) of Young's modulus are also shown in this figure. Despite the porosity deviation, it can be seen that the Young's modulus is reduced with increasing the BN content and the values are plotted in between E, and EL. For second-phase composite, variation of Young's modulus with volume fraction is expressed either by an exponential function as [13]: E = E,exp(-bV).. .......................... (2)
Generally, as a second phase with low Young's modulus suchs as pores and BN is incorporated, the strength is degraded. For instance, in the case of the BN nanoparticledispersed Si,N, composites [3], while the strength increased slightly (less than 4%) as BN content increased up to 5 vol%, further increase of the BN content resulted in remarkable degradation of the strength. The Sic-BN in situ composites, however, exhibited 9% strength increase at 5 vol% BN, and 4.4% at 15 vol% BN, compared with the Sic monolith, as shown in Fig.7. Moreover, the strength decreased only about 6% when the BN content was less than 35 ~01%.Due to the thermal expansion mismatches of BN and Sic, the BN percolation network can facilitate defect formation at the boundaries, degrading the material strength. As discussed in the above section, because that the BN flake pairs appeared and increased with increasing BN content before the formation of BN percolation network, the strength was degraded slightly with increasing BN content below percolation threshold. Although the BN incorporation above the percolation threshold leads to the boundary defect development, it also plays a role to refine the matrix grain as discussed in
475
the previous section and improves the material strength. These two contradictory factors may cause the gradual strength decrease in the BN content range from 15 to 35 vol% , as shown in Fig. 7. The fracture toughness increased a little for B5 and then decreased gradually and the details of the mechanism are under investigation. The strain to failure of the in situ Sic-BN composites calculated from & = 0 E-' is also shown in Fig.6. For improving the reliability of structural ceramics, particularly when they are used in conjunction with different materials like metals, large strain-to-failure, or strain tolerance, is an essentially important property. In the present case, the strain-tofailure increased remarkably with increasing the BN content and reached the maximum, which was about 2.5 times larger than that of the S i c monolith, at 45 vol% BN content.
of B,03 with AIN and/or Si3N4to Form BNToughened Composites, J. Am. Ceram. SOC., 7 1, (1988) 1080-85. (2) E. H. Lutz and M. V. Swain, Fracture
CONCLUSIONS
83-C-84. (7) R. Riedel, A. Kienzle, W. Dressler, L. Ruwisch, J. Bill and F. Aldinger, A Silicoboron Carbonitride Ceramic Stable to 2000 "C, Nature (London), 382, (1996) 796-798. (8) B. Baufeld, H. Gu, J. Bill, F. Wakai and F. Aldinger, High Temperature Deformation of Precursor-derived Amorphous Si-B-C-N Ceramics, J. Eur. Ceram. SOC.,19, (1999) 2797-2814. (9) I. Barin and 0. Knacke, Thermochemical Properties of Inorganic Substances, Spinger-Verlag, BerlidHeidelberg and Verlag Stahleisen m.b.H., Dusseldorf, Germany, 1973. Z. A. Munir, Synthesis of High (10) Temperature Materials by Self-propagating Combustion Methods, Amer. Ceram. SOC. Bull., 67, (1988) 342-49. Y. Zhou, H. Tanaka, S. Otani and Y. (1 1) Bando, Low-Temeprature Pressureless Sintering of a-Sic with A14C3-B,C-C Additions, J. Am. Ceram. SOC.,82, (1999)
SIC-BN in situ composites were prepared based on the in situ reaction of Si3N4,B4C and C by hot pressing at 2000 "C for 60 min under 30 MPa. The effect of BN content on the densification behavior, microstructure, elastic modulus, bending strength and fracture toughness of the SIC-BN in situ composites was studied. The change of elastic modulus with BN content obeyed both of the exponential and polynomial rules. The densification behavior of this in situ system was excellent even at high BN content. The grain growth of S i c was inhibited by the in situ formed BN flakes at high BN contents above 25 ~01%. The bending strength increased slightly with increasing BN content up to 15 vol% and then decreased gradually. Above 35 vol%, the strength decrease was accelerated. These results were explained according the BN percolation network formation. The markedly low elastic modulus as well as the relatively high strength were indicative of excellent strain tolerance of this material. ACKNOWLEDGMENTS This work has been supported by AIST, MITI, Japan, as part of the Synergy Ceramics Project. REFERENCES (1) W. S. Coblenz and D. Lewis , In Situ Reaction
476
Toughness and Thermal Shock Behavior of Silicon Nitride-Boron Nitride Ceramics, J. Am. Ceram. SOC.,75, (1992) 67-70. (3) T. Kusunose, Y. H. Choa, T. Sekino and K. Niihara, Mechanical Properties of Si,N,/BN Composites by Chemical Processing, Key Engineering Materials, 161- 163, ( 1999) 47580. (4) P. G. Valentine, A. N. Palazotto, R. Ruh and D. C. Larsen, Thermal Shock Resistance of SiCBN Composites, Adv. Ceram. Mater., 1, (1986) 81-87. (5) R. Ruh, A. Zangvil and R. R. Wills, Phase and Property Studies of Sic-BN Composites, Adv. Ceram. Mater., 3, (1988) 411-15. (6) R. Ruh, L. D. Bentsen and D. P. H. Hasselman, Thermal Difisivity Anisotropy of SiC/Bhr Composites, J. Am. Ceram. SOC.,67, (1984) C-
1959-64.
(12) H. Gu, Y. Shinoda and F. Wakai, Detection of Boron Segregation to Grain Boundaries in Silicon Carbide by Spatially Resolved Electron Energy-Loss Spectroscopy, J. Am. Ceram. SOC., 82, (1999) 469-72. W. D. Kingery, H. K. Bowen and D. (13) R. Uhlmann, Introduction to Ceramics, Second Edition, John Wiley & Sons, Inc., New York, USA, (1976) 768-777.
COATING EXPERIMENTS ON CARBON FIBERS USING A CONTINUOUS LIQUID COATING PROCESS N. Doslik*, R Gadow
Institute for Manufacturing Technologies of Ceramic Components and Composites, University of Stuttgart, Allmandring 5b, D-70569 Stuttgart, Germany
ABSTRACT Carbon fibers are found today in a wide range of applications as ceramic and metal matrix composite component with characteristic features as high tenacity. Commercially available fibers aren't appropriate for application at high temperatures because of the decay of tenacity caused by oxidative degradation. The best prospects for new high temperature stable coating materials show the use of SiBCN- and SiCN-systems. The realization is carried out by a liquid polymer impregnation with polyorgano-(boro)silazanes in a continuous
fiber coating process. These precursors are optimized in viscosity by the addtion of solvents and further compounds after detailed rheological investigations and adapted in the coating process. In a continuous t h e m treatment and fiber trzlnsport the precursor coatings are dried, polymerized and ceramized leading to systems which are resistive against oxidation and chemical attack e.g. in metal melts at high temperatures. The oxidation resistance of the coated composites was evaluated using thermogravimetricanalysis in atmosphere.
INTRODUCTION
Fig. 1 shows the schematic view of the whole pilot plant for continuous liquid fiber coating. The coiler system allows to use different kind of commercially available carbon fibers with 3000 up to 12000monofilaments. The pilot plant consists of two separate sections. Within the first, the fibers are led through a cleaning bath for desizing and then are dried (Fl). Commercial-grade carbon fibers normally have epoxy resin or PVA as sizing. This layer may interact with the precursors during impregnation, causing undesired reactions dunng drymg and pyrolysis. Furthermore it can deteriorate or reduce the adhesion of the ceramic layer to the carbon fiber surface. The second section of the LPI plant (F2, F3) is made as a vacuum chamber, where inert argon atmosphere protects the precursor polymers against contamination with oxygen or water. Before running the pilot plant a dipcoating process is used for primary investigations and optimization of the coating properties. The sequence of the preliminary investigations is divided in the following steps:
The high potential of carbon fibers as refractoq reinforcing component in ceramic and metal matrix composites"] strongly depends on the chemical and physical compatibility to the surrounding matrix. Therefore, new precursor systems based on polyorgano-(boro)silazanes are applied for coatings. One of the best prospects show the use of SiBCN-systems, especially to achieve high temperature stable materials. For carbon fiber coating, the knowledge of the rheological and wetting behavior of the precursor solvents is one of the most important requirements to get coatings without defects.[*]The precursors are optimized in viscosity by addition of solvents and W e r compounds and after detailed rheological investigations they are adapted in the coating process. By optimization of the rheological properties homogeneous and crack-free coatings on carbon monofilaments are obtaine~i[~] The realization of the continuous fiber coating is camed out by a liquid polymer impregnation. The liquid coating is followed by a two step thermal treatment process which is carried out in inert atmosphere. The first step allows drymg, curing and polymerization of the precursors and the second performs controlled pyrolysis to ceramize the intermediate inorganic polymer coating leading to systems which are resistive against oxidation and chemical attack e.g. in metal melts at high temperatures. r I
L
Fig.1:
Liquid phase coating and impregnation @PI)
1. Rheological characterization and optimization of the compounds 2. Dip-Coating process under argon atmosphere 3. Polymerisationunder argon atmosphere 4. Pyrolysis under argon atmosphere
EXPERIMENTAL RESULTS For fiber coating, the knowledge of the rheological and wetting behavior of the diluted precursors is one of the most important requirements to get coatings without defects and flaws. The flow and wetting behavior depends on the pseudoplastic flow type, which allows a superior wetting than the newtonian flow type.1413[51 The pseudoplastic flow type shows a characteristic decrease of viscosity with increasing shear stress or shear rate. Pseudoplasticity can be achieved by using additives which build up temporarily chain structures. These structures are de477
stroyed in a reversible process with increasing shear stress. A new measurement method to determine the rheological behavior of the precursors is the rotative oscillation method, which analyzes the viscuelastic properties. Viscoelasticity stands for the ratio of the elastic and plastic (viscous) part of flow properties. For measurin&a sample is poured in a gap between a cone and a plate. The cone makes a rotative oscillation. Torque and phase displacement are measured, rheological data are calculated. The method of determination used for the compounds is the amplitude sweep measuring method (AMS, frequency = constant, amplitude = variable). The amplitude sweep is able to determine the linear viscoelastic behavior, the flow point, the point of change of predominance of viscoelastic moduli and the stability of the additive. The storage modulus G stands for the elastic part of the viscoelasticity, the loss modulus G ' for the plastic part. With increasing storage modulus the sample shows a solid state like behavior, with increasing loss modulus the sample shows a fluid like." In Fig2 the SiBCN-precursor Rt7], diluted in the aprotic solvent tetrahydrofiuane, shows newtonian flow behavior with a constant low viscosity lq*l= 0,016 Pas. The loss modulus is six orders of magnitude higher than the storage modulus, thus the sample is very thin fluid and has non appropriate adhesion on the substrate. A liquid with a high loss modulus and a low viscosity at low shear rates is not able to stick to the surface of the substrate.
T
-
1E-4
B 0.02
1E4 c
s,
0.01
2
41
b
i 3
1N1E-4-
h
It
t
1M. 1Ed
-3
1E-7.
. .
1Ed
14014
a1
0.M
shear stresa % pa]
Fig.4: Newtonian flow behavior of SiCN-precursor PCS without additive
The SiCN-precursor PCS['], also diluted in tetmhydrohrane, shows a comparable rheological behavior (Fig.4) to sample P2 without PVB additive (Fig.2) with a non appropriate wetting characteristic. The additive effect of PVB on the flow behavior of PCS is very significant (compound mixture: 7,7 x 10" moVl PCS, 7,O x 10" mom PVB, 17,74 x lo4 moVl TJXIF). It induces in addition to the pseudoplasticty (Fig.5).
5
1EJ a
r
Fig.3 shows the precursor P2 with the additive polyvinylbutyral (PVB). This sample shows pseudoplastic behavior due to the reversible interaction with the additive. From textile ingineering and sizing it is known that pseudoplasticity enhances the wetting and penetration of fiber strands by liquid coatings. The higher viscosity at low shear stresses Iq*l = 0,037 Pas enables the precursor to wet the substrate and to stick on the surface (compound mixture: 15,25 x lo-' mom P2, 3,O x lo4 moVl PVB, 17,74 x lo4 mom THF).
-
-*--.-* -*--H--* * a-*-+W
.
HE
1
030
-
*-a* --L
0.01
0.1
shear stress 7 [pal
lE-3
Fig.2 Newtonian flow behavior of pFecursor P2 without additive
025
- 0.20 g
0.01
I
- 0.15 rF
mpwraool*.-lkm*~
1E-41Eb-
0*1°%
0.05 0.00 0.01
0.1
1
shear stress 7 pa]
Fig. 5: Hardening behavior of SiCN-precursor PCS with additive PVB
10,015
Fig. 3: Structural viscous flow behavior of precursor P2 with additive PVB
478
The increasing storage modulus depends on the observed hardening process of PVB. Dumg the period of shear stress variation (totally ca. 12 mia, measurement each 30 sec) an irreversible crosslinking of the precursor / additive mixture occurs simultaneously. The expected pseudoplasticity, expressed by decreasing viscosity under rising shear stress, is overcompensated by the viscosity increase due to the crosslinking. The storage modulus increases
from 1E-7 Pa to 0,9 Pa. The loss modulus increases smoothly from 0,l to 0,8 Pa. The sample makes a transition to a slightly solid state characteristics (transition point: 0,l Pa shear stress). The viscosity moves from 0,05 to 0,25 Pas after hardening. The sample has a good wetting behavior due to the pseudoplasticity and a superior adhesion due to the hardening process. The SiCN-precursor HPSr8] is unlike the other examined precursors a highly viscous liquid. Nevertheless the rheological data show an newtonian flow behavior (Fig.6). PVB addition was not expected to be sumssful because of the critically high viscosity caused by crosslinking. To point out if the wetting properties of this precursor are appropriate for a satisfying coating, the multimode frequency sweep ( M F S ) method has to be used because it provides evaluable data even for newtonian fluid coatings. But the mathematical model is only valid under the condition of linear viscoelastic behavior (G' parallel to G" ). The MFS method6] pig. 7) shows first if the precursor H P S is suitable for the coating process (loss modulus G" data form a light grey plane in 3D plot) and second in which area the linear viscoelastic properties are applicable thus proving linear viscoelastic behavior. The dark grey plane, showing the storage modulus G', demonstrates sigmiicantly less intensive solid like elastic behavior with lower values in [pa].
The oxidation resistance of the coated composites was evaluated using thermogravimetricanalysis in atmosphere (Netzsch STA 409C, Al2G-crucib1e, heating rate 10 Wmin). Fig. 9 and 10 shows the mass loss of the SiBCNprecursor P2 with PVB and the SiCN-precursors PCS with PVB and HPS, G1 without additive, as used in the coating experiments. All graphs show similar curves: first a mass loss through polymerisation reactions and second a stabilisation of mass with a transition to constant values (P2PVB: from 755OC; PCSPVB: fiom 730OC; HPS: from 700°C; G1: from 710°C).
rw
z
m
f
1" 40
Fig. 8: Thermogravimetric analysis of SiBCNprecursor P2 with PVB
0.1
1
shear mess r[Pa]
Fig.6: Newtonian flow behavior of the SiCNprecursor HPS without additive
Fig. 9: Thermogravimetric analysis of SiCNprecursor PCS with PVB,G1 and HPS
The constant run of curves at higher temperatures indicate that no oxidation occurs.
COATING RESULTS
Fig. 7: MFS: range of linear viscoelasticity of EIPS without additive. The fluctuation of the measured data in the lower range of Gand G" is caused by the rheometer setup.
After the rheological optimization the fluid coating of the carbon fiber filaments is carried out. The SEM micrographs show impressively some of the appropriate coating results after polymerisation and r l y s i s , too, of the SBCN- and SiCN-precursors.[71,[8 [ Fig. 10 shows the result of the optimized compound mixture P2 with addtive PVB on a carbon fiber monofilament after polymerisation The SEM micrograph of the fracture shows clearly the very good bonding between the coating and the carbon fiber. There are no sticking areas between single filaments.
479
Fig. 11: SEM micrograph of SiCN-precursor PCS with additive PVB afetr polymerization at 215OC under argon atmosphere The polymerisation, to fix and cure the coating on the carbon fiber monofilaments, is followed by the pyrolysis at higher temperaturesto build up the ceramic structure of the coating. Fig. 12 shows the SiBCN-precursor p2 after pyrolysis at 1100°C under argon atmosphere. The result is a homogeneous and crack-fkee ceramic coating with nonstickingmonofilaments.
Fig. 10: SEM micrograph of SiBCN-precursor P2 with additive PVB after polymerization at 305OC under argon atmosphere Similar results are reached for the polymerisation of the optimized compound mixture of the SiCN-precursor PCS (Fig. 11).
Fig. 12: SEM micrograph of SiBCN-precursorP2 with additive PVB after pyrolysis at llOO°C under argon atmosphere (monofilament)
480
CONCLUSION The wetting and flow properties of ceramic precursors on carbon fibers depends on the viscoelastic flow behavior of the selected coating polymers or blends. Therefore one can influence the wetting properties by adjustment of the viscoelastic properties introducing chain forming additives like PVB and matching the viscosity by dilution. The ratio between loss modulus and storage modulus is the indicating parameter for the viscoelastic behavior and the formation of an adherent coating. The value of this new measurement method using a rotative oscillation in rheometry has been proved by experimental results. Based on the experimental data results on SiCN- and SiBCNprecursor coatings are introduced and optimized crackfree and homogeneous monofilament coatings of carbon fiber filaments have been obtained Fig. 13: SEM micrograph of SiBCN-precursor P2 with additive PVB after pyrolysis at llOO°C under argon atmosphere non-sticking (monofilaments)
The pyrolysis of the SiCN-precursor G1 under NH3 atmospherer'O1 reaches an visually similar result, as in Fig. 14.
REFERENCES R. Gadow, S. Kneip, G. ScMer, Ceram. Trans. 103 (2000) 15ff N. Doslik, R. Fischer, R. Gadow, 24th Annual Cocoa Beach Conference 2000, Cocoa Beach, USA, Transactions of the Am. Ceram. Soc.2000, in print N. Doslik, R. Fischer, R. Gadow, 102"' Annual Meeting AcerS 2000, St. Louis, USA, Transactions of the Am. Ceram. Soc. 2000, in print Th. Metzger, S. Neuber, Messung des Fliefi- und Deformationsverhaltens von Stoffen, Chemietechnik 9 (1991) 50ff, Dr. Alfred Huthig Verlag GmbH, Heidelberg H.-G. Fritz, Einfikung in die Rheometrie der Kunststoffe, Editor: Technische Akademie Esslingen, 1996 H. Giesekus, FWnomenologischeRheologie, Springer-Verlag Berlin 1994 R. Riedel, private communication, precursor P2, supplied by University of Darmstadt, Fachbereich Materialwissenschaft,D-64287 Darmstadt, Ge-Y G. Ziegler, private communication, precursor PCS and HPS, supplied by University of Bayreuth, Institut fiir Materialforschung D-95440 Bayreuth, Germany U. Klingebiel, private communication, precursor G1, supplied by University of GiSttingen, Institut fiir Anorganische Chemie, D-37077 Gottingen, Ge-Y N. Doslik, R. Gadow, B. Jaschke, U. Klingebiel, R. Riedel, Appl. Organomet. Chem., Wiley & Sons, New York, issue 2000, in print
Fig. 14: SEM micrograph of SiCN-precursor G1 after pyrolysis at 9OOOC under NE13 atmosphere
48 1
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EFFECT OF A1 COMPONENT ON MECHANICAL PROPERTIES IN Al-PENETRATED ALUMINA Sang-Woo Hong* and Sung-Churl Choi Department of Ceramic Engineering, Hanyang University, Seoul 133-791 South Korea
ABSTRACT AIlO;/AI composites were fabricated by reacting mullite preforms with aluminum, and their mechanical properties were investigated according to the content of metal phase in the composite and the direction of metal penetration. We controlled the metal content in the composite by varying sintering temperature of mullite performs(l700C, 1650C, 1625'C and 1600°C). Two kinds of composites were prepared for the tests to investigate the effect of penetration direction. One is parallel, and the other is perpendicular to the direction of Al penetration. The reaction of Al with mullite preform provided the enhanced mechanical properties to the A1203/AI composite, such as hardness and fracture toughness, showing mutually interconnected microstructure. The microstructure of the composite and the mechanical properties were affected by the penetration direction. The mechanical properties were also affected by Al contents in the composite.
INTRODUCTION Metal reinforced ceramics have many applications due to the improved toughness of the inherently brittle material. But application of ceramic-metal composites for industrial fields has been limited in the point of machinability and near-net-shape capability, resulting in high manufacturing costs. Recently, in an attempt to reduce manufacturing costs, in-situ synthesis techniques for composites such as directed metal oxidation (DIMOXTM),'-4 pressureless metal infiltration (PRIMEXTM),' reaction bonding,s4 and liquid-solid displacement have received considerable attention. Reactive metal penetration (RMP) processing is a kind of liquid-solid displacement and has much merit in the near-net-shaping of composites." Reactions between molten aluminum and oxides have been known for many years particularly in relation to aluminum refining and these aren't used as a technique to synthesize composites, especially reaction synthesis of composites in the amorphous Si02/AI system. However, it is appropriate to fabricate ceramic-metal composites because of good physical, and mechanical properties of products. The study of this kind is recently made by Breslin et a1,* Loehman et al ,9 and Matsuo and Inabe.' For the application of ceramic-metal composites under real condition, it is necessary not only to optimize
'
fabrication processing but also to control the mechanical properties ofthe composite.'."-'2 In this study 4 kinds of mullite performs were prepared by varying sintering temperature to control A1 content in the composites. Two kinds of composites were prepared for the tests to investigate the effect of Al penetration direction. One is parallel and the other is perpendicular to the direction of Al penetration. Air atmosphere was also used for its convenient.
EXPERIMENTAL PROCEDURE Preparation of AVAlzOJComposites Starting materials were commercial mullite powder and aluminum powder(S. P. C. I, Korea). Mullite powder was attrition-milled for 24hours and uniaxially pressed. The pressed compacts were cold isostatically pressed at 200MPa and sintered at 16OO"C, 1625"C, 1650°C and 1700C for 4 hours to vary the densities of the performs. Table 1 lists their density, apparent porosity and reactionability. Mullite preforms were reacted with Al bar, which was prepared by pressing A1 powder. A box furnace was used for the reaction. The reaction between ' for 5 hours in mullite and Al was carried out at 1 100C an air atmosphere. An air atmosphere was chosen for its ' convenience. The mullite perform sintered at 1600C didn't react with Al, since an air atmosphere and A1 bar used in this study restricted the wetting of A1 to mullite perform. To investigate the effect of A1 penetration direction, two kinds of composites sintered at 1700"C were prepared for the tests. One is parallel and the other is perpendicular to the direction of Al penetration. Sintering Temp(C)
Density (dcm')
Apparent porosity ("/)
Reaction (Yes//No)
A1700
3.004
0.293
Yes
B1650
I
2.897
I
0.529
I
Yes
I
C1625
I
2.803
I
2.922
I
Yes
I
D1600
2.637
12.749
No
To compare these composites, we made another sample via conventional powder metallurgy. Alumina (A1203,99.9% purity, Alcoa A16SG, The Netherlands) having 0.3-0.4 W particles was ball-milled for 24 hours
483
and then dried. The dried powder was sintered at 1600°C for 2 hours. Characterization Scanning electronic microscopy(SEM) was used to examine the microstructures of AI/A1203 composites. A1/Al2O3composites were etched by dissolving A1 with NaOH solution. X-ray diffractometry (XRD) analysis was carried out on representative specimens to determine the phases which were present. Bulk densities of fabricated specimens were measured by using Archimedes immersion method. Four point bending tests were conducted with universal testing machine (UTM) in which the specimens were placed in a four-point bend fixture with inner span 10 mm and outer span 20 inm and a crosshead speed of O.Smm/min. Fracture toughness was measured, based on single edge notched beam (SENB) method, and Vickers indentations with 10 kg load using pyramidal diamond indenter were made on all materials to measure microhardness.
microstructure of the fabricated AVA1203composites. In Fig. 2, the microstructure of fully reacted composites shows mutually-interconnectedceramic and metal phases, and consists of a brighter phase indicating 3-dimensional A1203and darker phases occupied by Al. (a) is the upper part of reacted surface, which is the plane normal to the direction of A1 penetration. (b) is the side part of reaction surface on SEM photograph. Alumina has a longly connected channel shapes like pipes and the rest is occupied with aluminum. During total reaction, aluminum's penetration occurs very well, specially to the one direction.
RESULTS AND DISCUSSIONS Phase Analysis and Microstructure Fig. I showed mullite, A1 and AI2O3peaks detected by X-ray diffractometry. These results indicate that mullite preforms reacted with the A1 bar at 1 IOO'C and created A1203. At this temperature, it is sufficient for reactive metal to penetrate into mullite preform because aluminum's wetting angle to mullite is less than 90 degree'.' and Gibbs free energy is negative.' But the mullite perform sintered at 1600°C just partially reacted at the contact surface with A1 because Al bar and an air atmosphere restricted the coiiiact area between mullite and AI." Contrary to other works, small Si peaks were detected by XRD. It is supposed that Al bar used in this study restricted transport of Si to the Al bar. Residual Si did harm on AI/A1203composites since the K c of silicon is z 1 MPa.m"'.''
I
20kV lkx 1 O . O G
I
Fig 2.SEM photographs of AI/A1203 composites (a) the upper part of reaction (b) the side part of reaction surface 60
55 A
pl
50
45
40
35
A1700
B1650
C1625
Fig 3. Bending strength and fracture toughness of Al/A1203composites 20
40
60
80
213
Fig 1 . XRD patterns of mullite preform and A1203/AI composites The metallic phase was leached out with sodium hydroxide (NaOH) solution to observe the surface
484
Mechanical Properties Fig 3. shows the bending strength and fracture toughness of AI/A1203 composites. A 1700 sample had the best bending strength and the lowest fracture toughness. The strength of the composites increased with decreasing Al contents in the composites. Fracture toughness of the composites was better than that of
E %
2
alumina monolith prepared by powder metallurgy. The reason for this increase in fracture toughness can be seen in the fracture surfaces from the bending tests. Fig 4. shows SEM photographs of the fracture surfaces of the composites. The composites show many projection-like ridges and some basins which result from grain pull-put. It can be inferred from this result that plastic stretch prohibited cracks and resulted in the increase of fracture toughness composite. Fig 5 shows the Vickers hardness. Hardness value increased with the decrease of ductile Al phase in the composites.
toughness and bending strength. It is supposed that the growth of AI/Al2O3is nearly parallel to the direction of Al penetration, so this orientation can make differences in the mechanical properties in the composites.
CONCLUSION The reaction of aluminum with the oxide preform provides the enhanced fracture toughness to the resulting composites compared with alumina. By varying the sintering temperature of mullite, A1203/AI composites with different Al contents can be fabricated. 1. The reaction between mullite and A1 depends on the density of the mullite preform. Air atmosphere restricted the wetting of A1 to rnullite. So, rnullite preform with higher density is needed for the reactive metal penetration. 2. A1203/AIcomposites with different Al contents were fabricated and showed mechanical properties depend on the Al contents. 3. The direction of Al penetration affected the mechanical properties of Al2O3/AI composites. The normal direction of the composite showed better bending strength and fracture toughness.
Composites
3.43 - 3.64 253 - 280 5.14 - 5.61
Fig 4. SEM photographs of the fracture surfaces (a) B I650 (b) C 1625
0
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-
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$
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Fig 5. Vickers hardness of AI/AI2O3composites. The direction of Al penetration slightly affected the mechanical properties of the composites. The samples parallel to the penetration showed better fracture
able 2. The comparision of .properties o f , , A1203and . AI/AI2O3composites. REFERENCES (1) B.R. Lawn, Fracture of Brittle Solid, 2nd Ed., Cambridge Solid State Science Series, Cambridge Univ. Press (1993). (2) A. W. Urquart, Novel Reinforced Ceramics and Metals: A Review of Lanxide’s Composite Technologies, Mat. Sci. and Eng., A144, (1991) 75-82. (3) C.R. Kennedy, “Reinforced Ceramics via Oxidation of Molten Metals”, Ceram. Ind., 12, (1991) 26-30. (4) Saburo Hori, Net Shape Manufacturing of CMC and MMC by Lanxide Proocesses, Jpn. Ceramics, 32[2], (1997) 93-97. (5) N. Claussen, T. Le and S. Wu, Low-Shrinkage Reaction-Bonded Alumina, J. Europe. Ceram. SOC.,5, (1989) 29-35. (6) S. Wu, A.J. Gesing, N.A. Travitzky and N. Claussen, “Fabrication and Properties of Al-infiltrated RBAO-based Composites”, J. Europe. Ceram. SOC.,7, (1991) 277-81.
485
(7) S. Matsuo and T. Inabe. Fabrication of AI-Alz03 composites by Substitutional Reaction in Fused Aluminum, Jpn. Ceramics, (1991) 222-23. (8) M.C. Breslin, J. Ringnalda, J. Seeger, A.L. Marasco, G.S. Daehn and H.L. Fraser, Alumina/Aluminum Co-continuous Ceramic composite(C4) Materials Produced by Solid/Liquid Displacement Reactions: Processing Kinetics and Microstructure, Ceram. Eng. Sci. P~oc.,15[4], (1994) 104-12. (9) R.E. Loehman, K.G. Ewsuk and A.P. Tomsia, Synthesis of A1203/AI Composites by Reactive Metal Penetration, J . Am. Ceram. SOC.,79[ I], (1996) 27-32. (10) W.G. Fahrenholtz, K.G. Ewsuk, D.T. Ellerby, and R.E. Loehman, Near-Net-Shape Processing of Metal-Ceramic Composites by Reactive Metal Penetration, J . Am. Ceram. SOC.,79[9], (1996) 2497-99. ( I 1) Hiroaki Makino and Shigetaka Wada, Effects of Microstructure of Ceramics on Microstructure Induced by Round Indenter, Jpn. Ceramics, 27[ 101, (1 992) 943-47. (12) K.S. Lee, A Study on the Contact Damage in Silicon Nitride Bilayer, Ph. Ds thesis, Dept. of Mat. Sci. and Eng., KAIST, Taejon, Korea (1998). ( 1 3) V. Laurent, D. Chatain and N.Eustathopoulos, Wettability of Si02 and Oxidized Si c by Aluminum, Mater. Sci & Eng., A135, (1991) 89-94 ( 1 4) T. Watari, T. Torikai, W.-P. Tai and 0. Matsuda, of Fabrication and Mechanical Properties c(-AI20~/P-Al20~/AI/Si Composites by Liquid Displacement Reaction, J. Mat. Sci., 35, (2000) 5 15-520. (15) I. J . McColm, Ceramic Hardness, Plenum Press, New York ( 1990)
INFORMATION Department of Ceramic Engineering Hanyang University, Seoul 133-791 sung dong-gu haeng dang-dong, South Korea. TEL : 82 + 2 + 2290 0505 FAX : 82 + 2 + 2291 6767 e-mail :
[email protected]
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486
CAVITATION CREEP IN THE NEXT GENERATION SILICON NITRIDE
’,
P. Lofaj# S.M. Wiederhorn? G.G. Long: P.R Jemian? M.K. Ferber4
(*)
(l) Institute of Materials Research of SAS, 043 53 Kosice, Slovakia National Institute of Standards and Technology, Gaithersburg, MD 20899, USA (3) University of Illinois at Urbana-Champaign/ Argonne National Laboratory IL 60439, USA (4) Oak Ridge National Laboratory, Oak Ridge, TN 37831, USA
ABSTRACT The tensile creep behavior of a commercial grade of silicon nitride, SN 281, has been studied in the temperaturerange 1350°C - 1550”C, within the periods of up to 10,000 h. The corresponding creep damage has been investigated by electron microscopy and by anomalous ultra small-angle X-ray scattering (AUSAXS). Failure strains were around 0.5 %. The creep resistance of this grade was found to be several orders d magnitude greater than that of previously examined silicon nitrides. The corresponding stress exponents were higher than 6. Creep damage consisted CE multigrain junction cavities. Strain rate data. fit the cavitation creep model of Luecke and Wiederhom, which is based on a dilatation of the secondary phase pockets. Evolution of the pockets containing secondary phases was confiied by A-USAXS. High creep resistance in the current material was attributed to the suppression of cavitation via restriction of the redistribution of the secondary phases (crystalline secondary phases and residual glass) with higher effective viscosity. High effective viscosity seems to result fiom new type of sintering additives containing Lu. Lutetium-based secondary phases may be essential for the next generation of silicon nitrides with the operating temperatures up to 1500°C.
reduction in the level of pollutant emissions. However, the long-term reliability of ceramic gas turbines was not proven. One of the principal limitations which ultimately controls the lifetime, creep resistance, is sufficient in the current generation of silicon nitride ceramics for operation periods over 10,000 h only to temperatures of below 1350OC (Fig. 1) [3]. However, the maximum TIT in the advanced high power metallic gas turbines with thermal barrier coatings and special cooling of the blades already exceeded this temperature [4]. Thus, the potential advantage provided by ceramics was significantly undermined, especially when the reliability and price of the ceramics were considered. Ceramics can be competitive only if they provide significantly higher TIT which superalloys cannot principally sustain. The aim of the current work is to investigate tensile creep behavior and creep mechanisms in a silicon nitride grade with creep resistance of the magnitude needed 6 s the next generation of materials for structural applications. L W
10-7
10-8
INTRODUCTION The development of more efficient gas turbines is ultimately constrained by the maximum operating turbine inlet temperature (TIT) and by the availability of the structural materials with sufficient creep and oxidation resistance. Over the last decade, structural ceramics, particularly silicon nitride, were considered to be the most promising materials for the structural parts in gas turbines subjected to high stresses and temperatures up to 1400°C [I]. The tests of the CGT 302 developed within the Japanese national project “Research and Development of Ceramic Gas Turbines (300 kW class)” demonstrated a feasibility of a lowpower ceramic gas turbine operating at 1400°C with the thermal efficiency of 42% [2], which is twice that of the corresponding metallic turbines with lower TIT. Additionally, CGT 302 exhibited a significant
10-9
-.,
1 n-10
1100
1200
1300
1400
Temperature, “C Fig. 1. The comparison of the creep performance in different generations of silicon nitride ceramics at 150 MPa [3].
EXPERIMENTALPROCEDURE Material Characterization The material studied is the hot isostatically pressed silicon nitride designated as SN 281* (Kyocera Corp.,
#/ Current address: NIST, Gaithersburg, MD 20899, USA; e-mail:
[email protected].
* The use of commercial designations is for identification only and does not indicate endorsement by NIST. 487
Kyoto, Japan). It consists of SisN4 matrix grains with the mean diameter of about 1 pm and aspect ratio 2 to 4, and a small number of the large grains of length up to 30 pm and diameter of 3 -6 pm, often containing porelike defects and nuclei fiom seeding (Fig. 2). Secondary phases in the pockets are crystalline and they contain Lu, which was also reported in the gas-pressure sintered version of this grade, SN 282 [5]. The mean value of the four-point bending strength (3040 mm) at room temperature is 687d26 MPa; at 1200°C it is approximately 680 MPa, and 580 MPa at 1400°C [6-71.
Fig. 2. The microstructure of the silicon nitride studied with typical bimodal grain size distribution. Tensile Creep Testing The tensile tests were performed on flat, dog-bone specimens (SR51 type [8]). The gage size of the specimens is 2 mm x 2.5 mm and the gage length is 15 mm. They were loaded via single pin S i c pull rods, lever arm and dead weight. Tensile strain was measured in situ using a pair of silicon carbide flags suspended by their own weight on the specimen and laserextensometry system [8]. The raw data were recorded by PC in the interval of 5 min or 15 min and averaged over a time period corresponding to 3 to 7 data points. Details of the testing procedure are described elsewhere [8- 101. Twenty-one specimens were tested in air at 1350"C, 1400"C, 1450"C, 1500°C and 1550°C at the stresses ranging fiom 150 MPa to 380 MPa and periods up to 10 000 h. Two specimens broke during loading. Three specimens broke prematurely and the measurement system failed in three of the tests. Data fiom these specimens were excluded from consideration. MicrostructureCharacterization The phase composition of the specimens was investigated by X-ray difhction at 40 kV (Cu K,, h = 1.54046 A) fiom different zones of the bulk samples. The interior of the specimens, revealed by grinding away half of the specimen thickness, was used as a representative composition for the evaluation of the changes in the bulk. The samples for the transmission electron microscopy (TEM) studies were prepared h m f m creep tested specimens. They were cut from the core of gage section parallel to the direction of the applied stress. After hand grinding and polishing to -100 pm
488
thickness and dimpling, the foils were thinned by ion milling at 5 kV using Ar gas until foil perforation. Transmission electron microscopy investigations on carbon coated specimens were carried out at 200 kV and 400 kV. Creep damage was investigated by a scanning electron microscope (SEM) on the secondary fmcttm surfacesproduced at room temperature after creep and on the polished and plasma etched cross sections. Anomalous Ultra SmalCAngle X-ray Scattering Cavity size distribution and the evolution of the secondary phase pockets were investigated using anomalous ultra small-angle X-ray scattering on the beam line 33ID-D at the Advanced Photon Source at Argonne National Laboratory. The beam incident on the sample was monitored by an ionization chamber. The beam scattered by the sample was detected by a silicon PIN photodiode operating in an unbiased mode. The details of the USAXS instrument were reported elsewhere [ 1 I]. The samples with the dimensions of 2.5 mm x 4 mm x 0.15 mm were prepared fiom the as-received material and h m the gage and grip zones of three specimens after creep. The grip-gage pairs of specimens were used to eliminate the effects of initial porosity and possible changes in the microstructure during prolonged heat treatment on cavity and pocket size distributions. Data collection fiom USAXS scan for each specimen required approximately 20 min with a 5 s counting time per data point. The iterative method of Lake was used to desmear the data [12]. The anomalous USAXS technique involves measurement at -250 eV, -100 eV, -40 eV and -10 eV below the LIIIabsorption edge energy, which is 9244 eV for Lu. The presence of LuaSi207 as the secondary phase was assumed in the calculations. The details of the AUSAXS technique and evaluation method were reported elsewhere [131. RESULTS Creep Behavior Fig. 3 shows the long-term tensile creep behavior of the material at 1400°C under the stress of 200 MPa. The test was interrupted prior to failure after 10,200 h. During this period, the test was interrupted once due to the power outage. The total strain is around 0.5 %, however, due to the several failures in the data collection, an uncertainty of > 0.1% is included. Despite that, the transient stage exceeds 6000 h azd the strain rate after this period is less than 1 x 10- I/s, close to the limit of the laser extensometer resolution. A conservative estimate of the minimum strain rate at these conditions is 7 x lo-" l/s. Similar transient creep was observed at 1450°C. The short-term tests at 1550°C indicated conventional creep behavior with the primary and secondary stages. However, these data were affected by the oxidation of the S i c pull rods which form glass bubbles which interact with the laser beam. Because of prolonged lifetime, the number of tests performed until fmal failure was limited. Stress dependence of the minimum strain rate fiom the tests unaffected by premature failure is summarized in Fig. 4. It can be seen that the stress exponent exceeds 6.
t 7
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14OO0C/ 200 MPa
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0
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2000
4000
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,
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8000
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Time, [h] Fig. 3. Tensile creep curve obtained at 140OOC under the stress of 200 MPa in air indicate prolonged transient stage and the minimum rates below 1 x 10” Us.
IU -
-
. . . . .
10-7
as-received material. The same phases were detected by X-ray d i W o n &er creep tests at 1400°C and 1550°C after 10000 h and 520 h, respectively, indicating that the lutetium silicates are very stable secondary phase.
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200 300 Applied Stress, [MPa] Fig. 4. Stress dependence of the minimum strain rate in traditional log-log representation (* indicates data not used in the analysis). Creep hcture surfaces were similar to those in the other silicon nitride grades with relatively large creep damage zone and rough morphology typical for creep rupture regime. Fig. 5 illustrates intergranular !?actme occurring during secondary hcture at room temperature after creep. The crack propagated along the interhce between the matrix grains and large reinforcing grain. The shape of smaller matrix grains and the locations CE the multigrain junctions are visible due to the residual secondary phase left on the facet of the large grain. TEM observation of the creep tested specimens revealed very few cavities compared to the TEM studies on earlier silicon nitrides which were deformed to larger strains [7, 10, 14-15]. The cavities observed were exclusively the isolated multigrain junction cavities (Fig. 6). Cavity size is h m 100 nm to 400 nm which is equal to the size of the secondary phase pockets. Energy dispersive X-ray analysis confmed the presence of Lu in secondary phase pockets and subsequent X-ray difE-action analysis revealed the presence of Lu2Si207 (JCPDS cards No. 35-326, 34-509 or 3 1-777) and LuSi2N207 (JCPDS card No. 33-847) as the dominant crystalline secondary phases in the
Fig. 5. Intergranular fracture observed on the large grains on the secondary fracture surfaces after creep at 1400OC.
Fig. 6 . TEM micrograph of the multigrain junction cavity found in the studied material after creep.
489
The results of A-USAXS study for the grip and gage of the specimen tested at 1400°C are illustrated in Table 1 and Fig. 7. The size distributions of voids in the corresponding grips (pores) and gages (cavities + pores) are very wide. The volume tiaction of cavities, fv, is very low and close to the lower limit of the resolution. The size distributions of the secondary phase pockets have distinct peaks between 300 nm and 400 nm, which approximately correspond to those of the void size distributions. The total volume of each phase, which is calculated as a surface area below the corresponding distribution, and the mean size corresponding to the maximum volume of the both phases with certain size, are shown in Table 1. The resulting volume fiaction d cavities, was calculated as a difference between the gage and grip data. volume
M o n
mea n SPP 351 nm 421
size
f
SPP I voids I voids cavities asI 378 0.065 I 0.0027 nm received 0.040 0.0050 363 #1 grip #1 0.033 0.0064 389 360 0.0014 I I I I I gage #2 I 0.057 I 0.0029 I 350 I 430 I grip 0.046 0.0032 365 377 0.0003 #2 gage I I I I I Table 1. Summary of the U-SAXSdata from the specimens tested #1 - at 1400"C, 150 M a for 2804 h, and #2 - at 14OO0C,200 MPa for 215 h ( "SPP" is used for the secondary phase pockets, "voids" are related to the sum of the initial porosity in the as-received material and cavitation after creep).
DISCUSSION Direct comparison of the creep data from the material studied with those from the earlier grades of silicon nitride is not possible because of higher temperatures and stresses required to produce measurable creep. Therefore, strain rates were extrapolated to 150 MF'a and plotted in Fig. 8 using a temperature dependence similar to that in Fig. 1. This comparison indicates that at the strain rate of -3 x 10"' Us under the stress of 150 MPa, the strain of 1% is reached after 10000 h at the temperature of approximately 1500°C. This temperature capability is at least by 180°C higher than that of NT 154. A classification of the materials based on their mq resistance is implied from Fig. 8. The first generation d silicon nitride materials containing MgO additives exhibited sufficient creep resistance (normalized to 150 MPa) at temperatures below 1000°C. Second generation of silicon nitrides containing YAG secondary phases was resistant up to approximately 1100" - 1150°C. Introduction of pure rare-earth oxide addltives such as Y 2 0 3 and Y b 2 0 3 , improved creep resistance and increased the correspondingtemperature up to 1300°C 1325°C.
490
-1
h
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0
I
0
Fig. 7. The volume fraction size distribution, f (D), of the secondary phase pockets and voids in the grip -(A) and in the gage part -(B) of the specimen crept at 1400°C for 215 h. The current silicon nitride which can operate up to 1500"C, can be conditionally considered as the material of fourth generation. Apparently, the main reason for the improvement in this case is related to the Lu containing secondary phases originating from sintering additives (probably LuzO~).However, an understanding of the mechanisms controlling creep and of the Wmca between different materials are necessary to confirm such assumption. A number of recent creep studies on difkent silicon nitrides emphasized the role of cavitation among the process contributing to tensile creep strain [6-7, 9-10, 13-18]. A linear dependence between the strain and volume density change (volume hction of cavities) [7, 9-10, 15, 181, sound velocity change [6], Young's modulus change [ 171, a method based on the difference between tensile strain and compressive strain [ 161 and direct measurement of cavitation by USAXS in NT 154 [15] and SN 88 [13, 181 confirmed that cavitation at multigrain junctions are responsible for more than 90% of the tensile strain. Thus, cavitation was concluded to be the main creep mechanism in silicon nitride and possibly in other two-phase materials [7,9-10, 161. High stress exponents in the current material (Fig. 4) support a possibility of cavitation creep. TEM observation indirectly implies a correlation between strain and cavitation: small strains correspond to very low concentration of cavities. Similarly to other materials, only multigrain junction cavities were found Their size is comparable to the size of multigrain
junctions filled with secondary phase. Thus, despite limited data imposed by low strains, it can be concluded that cavitation is the main creep mechanism in the studied material, similar to that of materials with lower creep resistance [7]. This conclusion is supported by the A-USAXS data (Fig. 7) where the volume weighted cavity size overlaps with the size distribution ofthe pockets. The A-USAXS data (Table 1) indicates that cavitation may occur via redistribution of the secondary phase between the secondary phase pockets, since a correlation between the volume fraction of cavities and secondary phase is found an increase in cavity content results in a corresponding decrease of the secondary phase content (Table 1). Such mechanism was assumed in the recent cavitation creep model i?r silicon nitride proposed by Luecke and Wiederhorn [9]. Luecke and Wiederhorn cavitation creep model i?r two-phase materials is based on the effect of dilatation d the primary phase (silicon nitride) when subjected to deformation. It occurs via grain boundary sliding and results in expansion of the pockets with s o h secondary phase. Hydrostatic tensile stresses generated in some d the pockets are high enough for cavity nucleation. Consequent cavity growth occurs via redistribution of the secondary phase fiom the pockets with cavities into neighboring pockets and it is driven by the stress gradients between those pockets. The minimum strain rate, M6t, calculated for such mechanism is [9] &/at = B (l/q) (T exp (-Q/RT)exp(ao)
the term “secondary phase” is related to crystalline secondary phases and low amount of residual glass, both solution-precipitation and viscous flow can be considered. However, it is difficult to determine which of the processes is rate controlling dissolution of the crystalline phase, its diffusion via amorphous boundary phase, its precipitation or viscous flow of the glass. Thus, the term “effitive viscosity” is related to the solubility and diffusivity of the crystalline secondary phases and viscosity of the residual glass as well. Stepwise increase in the effective viscosity of the secondary phases based on different types of additives provides an explanation for the variations in creep resistance between different generations of silicon nitride while the other parameters, such as activation energy and stress dependence, remain the same (Fig. 9). Despite the absence of the experimental data on the effective viscosity of different rareearth based crystalline secondary phases, it can be concluded that lutetium based secondary phases exhibit higher effective viscosity than other rare-earth based oxides.
-
(1)
where B and 01 are the coefficients of proportionality, q is the effective viscosity and Q is the activation energy. This equation assumes exponential dependence on stress which is different fiom conventional power law. Such dependence corresponds to the stress exponents n 1 at low stresses and n >> 1 at high stresses [9] which is indicated in Fig. 4. Experimental data h m this Figure are replotted according to the Eq. (1) in Fig. 9. Because of limited stress range and number of data, both exponential and power law fit the data equally well. However, the stress exponents of 6 cannot be explained based on conventional d i h i o n creep models while they are directly involved in cavitation creep model. The principal processes involved in cavitation creep are schematically shown in Fig. 10 (A) and Fig. 10 (B). In idealized case of cavitation at multigrain junctions without dilatation (Fig. 10 (A)), cavity formation would correspond to the removing of the secondary phase h m the pockets and out of the solid. However, matter is conserved in a real situation and the secondary phase fiom the cavities is redistributed between the remaining uncavitated pockets (Fig. 10 (B)). This inevitably leads to the dilatation of the uncavitated pockets which is accompanied by a sliding apart of the hard grains cf primary phase. Thus, the redistribution of the secondary phase h m cavitated to uncavitated pockets is the ultimate controlling process of tensile creep deformation. Redistribution rate is determined by the effective viscosity of the secondary phases, q, (see Eq. (1)). Physical meaning of the effective viscosity depends on the possible mechanisms for the redistribution. Since
-
1100
1200
1300
1400
1500
1600
Temperature, “C Fig. 8. A comparison of the creep resistance for different silicon nitrides with that of material studied depending on the temperature and sintering additives.
3 10-9 b)
G 10-l0
lo-” 150
250 300 Applied Stress, MPa
200
350
Fig. 9. Stress dependence of the minimum strain rate in SN 281 described according to the cavitation creep model of Luecke and Wiederhorn [9].
49 1
The support tbr the stay of F. Lofaj at NIST provided by Fulbright Commission is gratefilly acknowledged.
REPERENCES
Fig. 10. Schematic representation of the cavitation at multigrain junctions without dilatation - (A) and cavitation creep resulting from redistribution of the secondary phase between the uncavitated pockets - (B).
CONCLUSIONS The improved creep resistance of a new generation d silicon nitride provides a potential for the operating temperatures up to 1500°C over prolonged periods of time and tensile stresses exceeding 100 MPa. Despite considerably lower strain rates, the principal a e q ~ mechanism is cavitation at multigrain junctions, identical to the earlier silicon nitride grades. Cavitation occurs via rehstribution of the secondary phase between the cavitated and uncavitated pockets. This process is ultimately controlled by the effectve viscosity of the secondary phases which is determined as the solubility and hffusivity of the crystalline secondary phases and viscous flow of the residual glass. Lutetium-based secondary phases seem to be crucial for the increase of the effective viscosity resulting in subsequent suppression of cavitation and increase in creep performance in the next generation of silicon nitride ceramics.
ACKNOWLEDGMENT APS at ANL is supported by NIST, DOE, ORNL, University of Illinois at Urbana-Champaign and others.
492
(1) T. Kamei, Overview of 300 kW Class CGT Project, Proc. of 1995 Yokohama Int. Gas Turbine Congress, Yokohama, Japan (1995) I- 143-146. (2) M. Yoshida, K. Tanaka, S. Tsuruzono and T.Tatsumi, Development of Silicon Nitride Components for Ceramic Gas Turbine Engine (CGT 302), Ind. Ceramics, 19 (1999) 188-192. (3) S.M. Wiederhorn, High Temperature Deformation d Silicon Nitride, Z. Metallkd. 9 (2000) 1053-58. (4) S. Amagasa, K. Shimomura, M. Kadowaki, K Takeishi, H. Kawai, S. Aoki and K. Aoyama, Study on the Turbine Vane and Blade for a 1500°C Class Industrial Gas Turbine, J. Eng. Gas Turbines Power Trans. ASME, 116 (1994) 597-603. ( 5 ) T. Ohji, Long-Term Tensile Creep Behaviors d Silicon Nitride for 1350°C Class Ceramic Gas Turbines, 102" Annual Meeting Am. Ceram. SOC., St. Louis, 2000. (6) F. Lofaj, J.-W. Cao, A. Okada and H. Kawamoto, Comparison of Creep Behavior and Creep Damage Mechanisms in the High Performance Silicon Nitrides, Proc. 6th Int. Symp. Ceramic Materials & Components for Engines, Arita, Japan (1998) 7 13718. (7) F. Lofaj, A. Okada, Y. Ikeda and H. Kawamoto, Creep Processes in the Advanced Silicon Nitride Ceramics, Key Eng. Materials, 171- 174 (2000) 747754. (8) J.D. French and S.M. Wiederhom, Tensile Specimens from Ceramic Components, J. Am. Ceram. SOC.,79 (1996) 550-552. (9) W.E. Luecke and S.M. Wiederhom, A New Model for Tensile Creep of Silicon Nitride, J. Am. Ceram. SOC.,82 (1999) 2769-78. (10) R.F. Krause, Jr., W.E. Luecke, J.D. French, B.J. Hockey and S.M. Wiederhom, Tensile Creep and Rupture of Silicon Nitride, J. Am. Ceram. Soc., 82 (1999) 1233-41. (11) G.G. Long, A.J. Allen, J. Ilavsky, P.R. Jemian, P. Zschack, The Ultra-Small-Angle X-ray Scattering Instrument on UNICAT at the APS, Proc. 1lIh U.S. Synchrotron Radiation Instrumentation Conf. 1999, (SRI"99), American Institute of Physics, in press, 2000. (12) J.A. Lake, An Iterative Method of Slit-Correcting Small-Angle X-ray Data, Acta Cryst., 23 (1967) 191- 194.
(13) P.R. Jemian, G.G. Long, F. Lofaj and S.M. Wiederhorn, Anomalous Ultra-Small-Angle X-ray Scattering ftom Evolving Microstructures During Creep, Proc. MRS Fall Meeting 1999, Boston, USA, (2000) in press. (14) M.K. Ferber, M.J. Jenkins, T.A. Nolan and R.L. Yeckley, Comparison of the Creep and Creep Rupture Perfomance of Two HIPed Silicon Nitride Ceramics, J. Am. Ceram. SOC.,77 (1994) 657-65. (15) W.E. Luecke, S.M. Wiederhom, B.J. Hockey, R.E. Krause, Jr., and G.G. Long, Cavitation Contributes Substantially to Tensile Creep in Silicon Nitride, J. Am. Ceram. SOC., 78 (1995) 2085-96.
(16) F. Lofaj, A. Okada and H. Kawamoto, Cavitational Strain Contribution to Tensile Creep in Vitreous Bonded Ceramics, J.Am.Ceram.Soc., 80 (1997) 1619-23. (17) J.-W. Cao, F. h f a j and A. Okada, Application of an Ultrasonic Technique to Creep Cavitation in Silicon Nitride, J. Mater. Sci., in press.
(18) F. h f a j , P.R. Jemian, J. Ilavsky, G.G. Long and S.M. Wiederhom, Evolution of Cavities During Creep in Silicon Nitride by Anomalous Ultra-Small Angle X-ray Scattering, 102"dAnnual Meeting Am. Ceram. SOC.,St. Louis, 2000. I
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THERMAL CONDUCTIVITY AND PHONON SCATTERING MECHANISMS OF B -Si3N4 CERAMICS Koji Watari, Kiyoshi Hirao, Takaaki Nagaoka, Motohiro Toriyama National Industrial Research Institute of Nagoya, Hirate-cho, Kita-ku, Nagoya 462-8510 Japan ABSTRACT B-Si3N4 ceramics with various grain size were fabricated in order to investigate effect of grain size on the thermal conductivity. Strong relationship between thermal conductivity and grain size is not found under the present work. Furthermore, calculation of phonon mean free path of B -Si3N4 ceramics demonstrates that phonon scattering occurs at an average interval of 1030 nm. It is, therefore, concluded that the thermal conductivity is controlled by the type and amount of crystal defects into the grains.
INTRODUCTION B -Si3N4 ceramics are widely used in industry because of its excellent mechanical properties, including strength, fracture toughness and wear resistance. Recently, it has been established through calculation of intrinsic thermal conductivity and thermal conductivity measurements that B-Si3N4 is a high thermal conductivity (> 100 W/m"C) material. Watari et al. estimated from the average mass of atoms, average volume occupied by one atom, Debye temperature and Gruneisen's constant that a maximum thermal conductivity at room temperature of B -Si3N4 is about 400 W/m"C [l]. On the other hand, Haggerty and Lightfoot focussed that Sic and Si3N4 are nearly identical except for the number (n) of atoms in each primitive cell, and reported that a p r d c t e d intrinsic 200 W/m"C for n = 14, and 320 W/m value is C for n = 7 [2]. Measurement of the conductivity of B-Si3N4 single crystal was carried out by Li et al.. They have focussed on an extremely large elongated single crystal grain (mean diameter: 17 w m and length: 100km), and reported that conductivity of a
-
-
B -Si3N4 single crystal was 180 and 69 W/m"C along the c- and a-axis, respectively [3]. As for ceramic materials, their conductivities have been reported to be 10-155W/m"C [4-111. Textured B-Si3N4 ceramics indicated higher conductivity compared to non-oriented ceramics [8]. The conductivity of textured B-Si3N4 obtained by tape-casting with seeds and HIPing was achieved to 155W/m"C in the direction of tape-casting [lo]. It is significant to clarify thermal conduction mechanism of sinkred B -Si3N4. Controlling factors on the conductivity of B-Si3N4 ceramics have been proposed by numerous researchers. The factors are divided into two groups: (a) microstructure effect and (b) crystalline perfection. As a microstructure effect, characters of grain boundary phase, i.e. amount, thickness, distribution and chemical composition, and characters of grain, i.e. size, aspect ratio and orientation, would be suggested. Many authors focussed on thickness of grain boundary phase and grain size (number of two grain-grain junctions), because typical thermal conductivity of silicate-based glass, which is almost similar to grain boundary glassy phase, is known to be as low as 1 W/m"C. Hirosaki et al. concluded that the conductivity increased with increasing grain size (decreasing number of two gmngrain junctions) [7]. However, larger grain size lead to lower mechanical strength of the sinkred materials, because the flaw size relating to fracture becomes larger. Thus, a combination of high thermal conductivity and mechanical strength can not be achieved by fabricating large-grained Si3N4 ceramics. Recent work by Kitayama et al. reported that the conductivity of B -Si3N4 ceramics quickly decreased as the grain boundary film thickness increases in a range of a few tenths of a nanometer 1121. According to their conclusion, it is important to decrease the grain boundary film thickness in order to enhance the
495
conductivity. Grain-boundary phase is formed by an eutectic reaction between oxide on Si3N4 parhcles and sintering aid, which promotes densification. It is, therefore, difficult to obtain high thermal conductivity and dense material without reducing the amount of sintering aids. In this work, effect of grain size on the conductivity of B -Si3N4 ceramics with various sintering aids was investigated experimentally. Furthermore, the thermal conduction and phonon scattering mechanisms of sintered B-Si3N4 are discussed based on experimental results and calculation of phonon scattering distance.
EXPERIMENTAL PROCEDURE High-purity a -Si3N4 raw powders (Ube industries, E10) with single addition of Y203 or A1203, and concurrent additions of A1203-Y203 or Y203-Mg0 were mixed by ball-milling, using a 2-propanol solvent. After drymg, the mixed powders were pressed into pellets using a stainless steel die, and were then CIPed under 400 MPa The CIPed specimens were sintered at 1800, 1850 and 2000 “c under N2 gas pressures of 1 and 200 MPa. Also the mixed powders were then placed into a graphite die and hot-pressed at 1800 “c under a uniaxial pressure of 40 MPa in flowing N2 atmosphere. The sintered specimens were embedded in a Si3N4BN-SiO;? powder mixture with additives in a graphite crucible, and annealed at 1850-2500 “c for 2-10 h under N2 gas pressures of 1 and 200 MPa. After the samples were sintered and annealed, their bulk density was measured by a method involving displacement in water. The oxygen and yttrium contents of the specimens were measured using hot-gas extraction analyzer and inductioncoupled plasmaatomic emission spectrometry, respectively. The microstructure was evaluated by SEM observation. The thermal conductivity at room temperature was evaluated by a laser flash technique
[a.
RESULTS As representative data, the characteristics of Y203doped specimens hot-pressed at 1800 “c and annealed at 2500 “c are shown in Table 1. The bulk densities of the hot-pressed and the annealed specimen are close to theoretical density of B -Si3N4, indicating their fully
496
Table 1. Characterstics of B -Si3N4 cemnics after hot-pressing and subsequent annealing. Hot-pressed Si3N4 Density (kglm3) 3236 Oxygen content 2.1 (mass.%) Yttrium content 2.7 (mass.%)
Hot-pressed and annealed Si3N4 32 14 2.0 2.7
densification. No significant variations in oxygen and yttrium contents were found between the hot-pressed and the annealed materials. On the other hand, enhancement of p n size due to subsequent annealing was recognized as shown in Fig. 1. It is, therefore, possible to discuss a relationship between the thermal conductivity and grain size in sintered B-Si3N4 without reducing the change of amount of sintering aids. Fig. 2 indicates a relationship between thermal conductivity of B-Si3N4 with various sintering aids and mean gain size. Mean grain size was determined using linear intercept method. For Y 2 0 3 and Y 203MgO additions, their conductivities increase as the grain size increases up to about 5 j ~ mand , thereafter are almost constant. In the case of additions of A1203 and Y203-Al203, a significant change of the conductivity depending on the grain size is not Observed.
DISCUSSION The thermal conduction of B -Si3N4 is mainly due to phonons. The phonon mean free path, L ,is described bY G = 3 KlVC (1) where K is the thermal conductivity, V the group velocity of the phonons, and C the specific heat 1131. Watari et al. calculated the phonon mean free path of B -Si3N4 ceramics with values of 80-150 W/m“c, and reported that their phonon mean free path is 10-30 nm. It indicates that phonon scattering occurs at an average interval of 10-30 nm [11,14]. Considering that conventional -Si3N4 ceramics have grains as a major
b
Figure 1. SEM photographs of plasma-etched surface of B -Si3N4 specimens fabricated by (a) hot-pressing at 1800 "c and (b) subsequent annealing at 2500 "c.
3
40-
L.l
2
--c+
20-
6 mol% A1203
-
b
0
phase (> 95 vol.%) and grain boundary phase (< 5 vol.%) as a minor phase, number of phonon scattering sites in the grains is much larger than that in grain boundary phase [141. The reason that there exists many phonon scattering sites into the grains is also explained by comparing with intrinsic thermal conductivity and measured conductivity of the grain. The conductivity of a B -Si3N4 single crystal grain in a sintered material with thermal conductivity of 155 W/m"C was 180 W/m"c along the c-axis. This value is 56 % for the intrinsic value (320 W/m"c [2]), presenting a large
I
amount of crystal defects into the grains. In this work, we produced B -Si3N4 ceramics with diverse grain sizes of 1.5 to 14 a m . If the conductivity had been influenced strongly by grain size, the speamens of diverse grain sizes would have exhibited large differences in conductivity. However, significant effect of grain size on the conductivity of sintered B-Si3N4 is not found as shown in Fig. 2. The conductivity of B -Si3N4 ceramics therefore must not be controlled by the grain size. Point defects, dislocation, solid solution are given as crystalline perfection to reduce the conductivity of
497
Point defects, dislocation, solid solution are given as crystalline perfection to reduce the conductivity of the grains. Presence of dislocations in the grruns has been reported by Lee Brito [16], and Munakata [17]. They mentioned that the major grains have misfit dislocation network and dislocations tangle. On the other hand, solid solutions have been observed into the B-Si3N4 grains. A1203 addition results in formation of SIALON. Aluminum and oxygen in Si3N4 turn out to be substitutional impurities in the crystal. Oxygen repiaces nitrogen, and aluminum replaces silicon (183. Then, mass differences and vacancies form in lattices, and act as phonon scattering sites [6]. Therefore, the conductivity of A12CX-doped specimens is low. Hirosaki et al. found crystal particles, 10-30 nm in diameter of Y-Ndapatite in the p u n s of Y2CX-Nd203 added specimens. The formation is due to solution-reprecipitation process for liquid-phase sintering [19]. Recently, Kitayama et al. have succeeded in measurement of the oxygen content in the B-Si3N4 grains by the hot-gas extraction method, and found that the grains have oxygen contents of 0.150.25 % [20]. The oxygen atoms in the grains will become point defects. These crystal defects will be factors influencing on the conductivity of sintered B-Si3N4. At present, it is, however, difficult to determine significant phonon scattering sites. Further work must focus on detailed type and amount of crystal defects into the grains.
[la,
CONCLUSION The thermal conductivity d B-Si3N4 ceramics was strongly influenced by the type and amount of crystal defects into the grains. To develop high thermal conductivity B -Si3N4, removal of crystal defects into the grains is significant.
REFERENCES [l] K.Watari, B-C.Li, L.Pottier, D.Fournier, and
M.Toriyama, Key Eng. Mater., in press. [Z] J.S.Haggerty and A.Lightfmt, Ceram. Eng. Sci. Proc., 16,47587(1995). [3] B-C. Li, L.Pottier, J.P.Roger, D.Fournier,
K.Watari, and KHirao. J. Euro. Ceram. Soc., 19, 1631-40 (1999). [4] M.Kuriyama, Y .Inomata, T.Kijima and Y.Hasegawa, Am. Ceram. Soc. Bull., 57, 1119-22 (1978).
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[5] K.Tsukuma, MShimada and M.Koizumi, Am. Ceram. Soc.Bull., 60, 910-12 (1981). [6] K.Watari, Y.Seki and K.Ishtzaki, J. Ceram. Soc. Japan, 97,5642 (1989) N . H i d , Y.Okamot0, M.Ando, F . M d and Y.Akimune, J. Am. Ceram. Soc., 79,2978-82 (1%). [8] K.Hirao, KWatari, M.EBrito, M.Toriyama and S.Kanzaki, J. Am. ceram.soc.,79,2485-88 (19%). 191 Y.Okamoto, N . H i W , M.Ando, EMunakata and Y.Akimune, J. Mater. Res., 13,3473-77(1998). [101 K.Watari, KHirao, M.E.Brito, M.Toriyama, and S.Kanzaki, J. Mater. Res., 14, 1538-41 (1999). [111 K. Watari, K.Hirao, M.Toriyama, and K.Ishizaki, J. Am. Ceram. Soc.,82, 777-79 (1999). [12] M.K~tayama,K.Hirao, M.Toriyama, and S.Kanzaki,J. Am. Ceram. Soc., 82,310512 (1999). [13] J.M.Ziman, in "Electrons and Phonons", Oxford University Press, London, 1%0. Chapter VIII. [141K.Watari, K.Hirao, M.E.Brito, and M.Toriyama, Abstract of 1997 the 4th IUMRS International Conference in Asia, pp. 477, Paper #. L 4.2, Chiba, Japan, Sep. 16-18 (1997). [15l W.ELee and G.E.Hilmas, J. Am. Ceram. Soc.,72, 1931-37(1989). [16] M.E.Brito, K.Hirao,K.Watari, and S.Kanzalu, presented at 1996Annual Meeting & Exposition of Am. Ceram. Soc.(Paper No.,SVII-40-%), Indianapolis, U.S.A., April 1417, 1996. [17] F.Munakata, C.Sato, N.Hirosaki, M.Tanimura, Y.Akimune, Y.Okamoto and Y.Inoue, J. Ceram. Soc.Japan, 105, 858-61 (1997). [18] A. Tsuge and K.Nishida, Am. Ceram. SOC. Bull., 57,424-26 (1978). [19] N.Hirosaki, T.Saito, EMunakata, Y.Akimune, Y.Ikuhara, J. Mater. Res., 14,2959-65 (1999). W] M.Kitayama, K.Hirao, A.Tsuge, KWatari, M.Toriyama, and S.Kanzaki, J. Am. Ceram. Soc. in press.
A HIGH THERMAL CONDUCTIVE -SILICON NITRIDE SUBSTRATE FOR POWE MODULES
P
Hiroshi Yokota and Masahiro Ibukiyama Research Center, Denki Kagaku Kogyo K.K. 3-5-1 Asahimachi, Machida-shi, 194 Tokyo, Japan
ABSTRACT We have newly developed a silicon nitride substrate with a thermal conductivity of 100 Wm-'K-'. A high thermal conductivity with sufficient mechanical properties was achieved by processing a high purity raw powder of p-silicon nitride, with ytterbium oxide as an additive and developing the homogeneous microstructure of a sintered material. This material was densified to over 99 % of the theoretical density at the sintering temperature of 1800 "C for 8 h under a nitrogen pressure of 0.9 MPa. In this material, the flexural strength was between 550 and 600 MPa, and fracture toughness was 6 MPa*m"2. It is thus considered that this material could be an attractive substrate for power modules.
INTRODUCTION Recently, power modules using aluminum nitride substrates have been more attractive than aluminum oxide substrates due to the high thermal conductivities of the former. Aluminum nitride substrate used in typical power modules is sandwiched by copper plates. The main objective of aluminum nitride substrates for this application is to prevent cracks and fracture caused by thermal expansion mismatch stress between the aluminum nitride substrate and the copper sheets during thermal cycles. On the other hand, very recently, silicon nitride substrates have begun to be used due to their excellent mechanical properties. We already developed a silicon nitride substrate with a thermal conductivity of 70 Wm-'K-', in 1997. If silicon nitride substrate with higher thermal conductivity can be developed, this substrate will be more attractive for power modules. Hirosaki et al. fabricated p-silicon nitride ceramics with a room temperature thermal conductivity of 122 Wm'K" by gas pressure sintering at a temperature of 2200 "C[l]. They reported a significant improvement of thermal conductivity due to the decrease in two grain junctions accompanying grain growth. Hirao et al. fabricated p-silicon nitride ceramics with a highly anisotropic microstructure a i d exhibiting a high thermal conductivity of up to 120 Wm 'K1 by using rodlike psilicon nitride seed particles and annealing at 1850 "C for up to 66 h under a nitrogen pressure of 1 MPa [2]. Watari et al. further annealed this material at 2500 "C under a nitrogen pressure of 200 m a . This material reached the thermal conductivity of 155 Wm 'K-' [3]. In those processes, grain growth of p-silicon nitride promoted by prolonged annealing time or extremely high annealing temperature was required to
obtain high thermal conductivity. It is well known that smaller grain size results in higher flexural strength of sintered materials. The presence of such fine and homogeneous grain is considered to result in higher flexural strength. If the grain size increase, however, the mechanical strength of the sintered materials will decrease. Silicon nitride ceramics with high thermal conductivity thus tend not to be strong mechanically. In fact, there have been no studies obtaining both high thermal conductivity and high mechanical strength. Additionally in these studies, silicon nitride raw powders have been sintered with Y,O, or Y,O, - based as additives [l-31. In this study, we concentrate on investigating the silicon nitride ceramics with both high thermal conductivity and mechanical properties sufficient for the need of a substrate produced in a low temperature sintering. The experiments have been designed to determine whether silicon nitride that has been sintered with a rare-earth additive (lanthanide oxide, Ln,O,, where Ln is a lanthanide element) exhibit higher thermal conductivity than silicon nitride with Y203.
EXPERIMENTAL PROCEDURES Table 1. Properties of raw powder of silicon nitride Powder (a) (b) Impurities Oxygen 1 wt% 0.80 0.80 80 60 Aluminum / ppm 100 50 Iron 1 ppm 50 60 Calcium / ppm 32.5 30.5 a-phase content / wt% Specific surface area / m'g ' 12.5 15.0 0.8 0.5 Average particle size / pm All the samples were prepared from two types of high purity silicon nitride powders and high purity rareearth oxides. Table 1 shows the two types of silicon nitride powders used in this experiment. Two powders (a) and (b) were synthesized by direct nitridation of metallic silicon. Both were manufactured in our company.
(1) Determining the thermal conductivity with Ln,O, When comparing different rare-earth oxide additives Yb,O,, Er,O,, Dyz03,Y,O,, Sm,O,, Nd,03 and La,O,, the amount of sintering rare-earth oxide was 2.5 mol% . Powders of p-silicon nitride type (a) were ball milled with each rare-earth oxide additive and 1 wt% ZrO, using a solvent of methanol for 3 h. After drying, pellets of powder mixtures in each composition were
499
uniaxially pressed using a pressure of 10 MPa, which resulted in a green density that was about 55 9% of the theoretical density. The pellets were sintered at 1900 "C for 8 h under a nitrogen pressure of 0.9 MF'a. The samples were inserted in a BN tube with BN end caps to minimize decomposition and reduction reactions at the pellet surface.
(2) Preparation of silicon nitride substrate Two types of the raw powders of silicon nitride (a) and (b) were used to investigate the influence of microstructure of sintered materials on thermal conductivity and bending strength. The raw powders (a) and (b) were mixed with 10 wt% of Yb,O, and 2 wt% ZrO, as additives by ball milling for 3 h. After drying, the powder mixtures, an organic binder and water were ball milled for 3h. Green sheets were formed by the use of extruding method, adjusting the sheet thickness to about 0.85 pm. Subsequently, the green sheets were punched into a rectangular shape ( 7 5 x 4 5 mm). The sheets were then calcinated at the temperature of 550 "C for 3 h , which resulted in a green density that was about 55 5% of the theoretical density. The sheets were sintered at 1800, 1820, 1845,1870 and 1930 "C for 8 h under a nitrogen pressure of 0.9 MPa. The samples were contained inserted in a BN tube with BN end caps to minimize decomposition and reduction reactions at the pellet surface. 3) Characterization After sintering, the samples were characterized with regard to weight loss, density. The density (p) of sintered material was determined by Archimedes method immersion using deionized water. The microstructure of the sintered materials was examined by scanning electron microscopy (SEM, JSM-8404 JEOL, Japan) of polished and plasma-etched surfaces. In order to characterize quantitatively the microstructure, large grains having a diameter greater than 2 pm were identified in an area of 200 X 200 pm' for each samples, and the area fraction of the large grains was calculated using an image analyzer. Quantitative analyses of oxygen content of sintered materials were carried out with a hot gas extraction analyzer (TC-436, LECO, St. Joseph, MI). The phases that were present in the bulk were determined by XRD analysis. (MCP-3, Mac science, Japan ) The thermal diffusivity (a) and the specific heat (Cp) of the samples were measured at room temperature by a laser flush method using a thermal constant analyzer (TC-3000, ULVAC, Japan). The thermal conductivity (K) was calculated from the equation. K
=a xc px p
For only the sintered materials for use as the substrates, three pointed flexural strength was measured with a span of 30 mm, and a cross head speed of 0.5 mm/min at room temperature. Fracture toughness was determined by the IF method at the load of 196 N at room temperature.
500
RESULTS (1) Comparison of Y203with Ln,03 Table 2 gives the sintering characteristics of silicon nitride-Ln,03 composition pellets sintered at 1900°C for 8 h. The lanthanide oxides (Ln,O,) ,when added in equimolar amounts, were densified to over 99 9% of the theoretical density at the sintering temperature of 1900°C for 8 h. It was noticed that weight loss was largest when Yb203was used as an additive. Table 3 gives oxygen content, thermal conductivity and phases for silicon nitride - Ln203 compositions sintered at 1900°C for 8 h. The thermal conductivity varied, from 99 Wm 'K for silicon nitride-La,03 to 121 Wm-'K-' silicon nitride- Yb203. It was found that predominant secondary phase for each composition was K phase (2Ln,03- lSi0,. lSi,N, ), with the exception of silicon nitride-Ybz03, which revealed J phase (4Ln,O31SiO; 1Si3N4). With regarded to oxygen content of the sintered materials, the lowest oxygen content was obtained in the case of silicon nitride- Yb,O,. Table 2. Sintering characteristics of silicon nitride Lnz03compositions sintered at 1900 "C for 8 h Sintered %of Weight Density Theoretical loss 1 gcm I wt% density YZO, YbZO, ErZ03
Dy2°3 Sm203 Nd203
La,O,
3.27 3.38 3.40 3.35 3.37 3.34 3.32
100 99.0 100 99.4 100 100 100
2.02 6.16 2.78 2.27 4.17 2.79 3.01
Table 3. Thermal conductivity and phases for silicon nitride - Ln,03 compositions sintered at 1900 "C for 8 h Oxygen* * * Thermal Phases Content conductivity by XRD
ErZ03
DyZ03 Sm203
Nd,03 La,O,
0.11 0.59 0.37 0.41 0.57 0.65
121 109 114 114 108 99
SN>>J* * SN>>K SN>>K SN>>K SN>>K SN>>K
*K phase reveals 2Ln,O,- lSiO,-lSi,N, **J phase reveals 4Ln,0,-1Si0,-1Si,N4 ***Oxygen content reveals the difference total oxygen content of sintered material and initial contents of oxygen with regard to Ln,O,and ZrO,.
110 3
1201
t
L4
om
0
Sm
55
.M
>
100
n
I
Powder(a) Powder(b)
I
90
.cI
c
oLa
loot
a 2 80 0
-3 a
*
70
80
sintering temperature/OC Fig. 1Thermal conductivity of the samples as a function of ionic radius of rare-earth oxide.
Fig.3 Thermal conductivity as a function of sintering temperature.
Fig.l shows the effect of ionic radius of rare-earth oxide, when added to silicon nitride on the thermal conductivity. The thermal conductivity of silicon nitride increased as the ionic radius of rare-earth oxide decreased.
(3) Properties of silicon nitride substrates Fig.3 shows the thermal conductivity of silicon nitride-Yb,03 of substrates as a function of sintering temperature. All the samples were densified to over 99 % of the theoretical density. The thermal conductivities of substrates using powder (b) increased remarkably when sintering temperature increased. It is noticed that, however, the thermal conductivities of substrates using powder (a) were higher than those using powder @) at any sintering temperature, and reached to 96 Wm 'K at the sintering temperature of 1800 "C for 8 h.
'
Area fraction of large grains/area% Fig.2 Correlation between area fraction of large grains and thermal conductivity.
Fig.2 shows that the effect of area fraction of large grains on the thermal conductivity. Thermal conductivity of silicon nitride-Ln203 increased with area fraction of large grains. It was found that the growth rate of silicon nitride was fastest in the case of silicon nitride-
50 1
(A) 1800 oc x 8 h with powder (b), K=96Wm-'K-l
@) lS7O
"'
with powder (a)' K=99 wm
(c)1820 "c X 8 h with powder (b), ~ = 8 Wm 0 'K1
Yb,03 reached the highest thermal conductivity. Additionally it is noticed that the several remarkable
other rare-earth oxide related systems exists four nitrogen containing, pseudoternary phases. On the other hand, the Yb-system has one pseudoternary compound (J phase). Hoffmann et al. mentioned that the reason for in the instability of the H phase (lOLn,O3-9Si0,-1Si,N,), K (2Ln,O,. lSi0,. 1Si3N,) phase, M phase (lh,O,. lSi,N,) in the Yb-systems is attributed to the smaller cation radius of the Yb3'-ion in comparison to Y3'[4]. The secondary phase compositions are influenced by the amount of the rare-earth additive and the amount of SiO,, which resulted from oxygen content of the raw powders of silicon nitride. In this experiment, since both the amounts of the rare-earth oxides and Si0,were equimole, the existence of J phase was attributed to the smaller cation radius of the =3+-ion. The liquid phase in the silicon nitride-Yb,O, is maintains the Si0,rich liquid phase with difficulty due to the instability of the thermodynamics, as compared to K phase in other rare-earth oxides. Therefore, SiO, remaining in the liquid phase is considered to react with silicon nitride, and then evaporate to the atmosphere as follows; Si,N,
(D) 1930 "CX 8 h with powder (b), ~ = 9 5Wm-'K-' Fig.4 SEM photographs of polished and plasma -etched surface of samples. Fig.4 shows that SEM photographs of polished and plasma-etched surface of samples. The photographs denoted symbol (A), (B), (C) and (D) in Fig.4 correspond to the denoted plots in Fig.3. It found that the
502
+ 3Si0,
-
m i 0 + 2N,
(1)
Accordingly, the largest weight loss of silicon nitride-Yb,O, is attributed to the reaction (1). It is well established that oxygen as an impurity lowers the thermal conductivity of aluminum nitride. In the aluminum nitride-Y,O, systems, Virkar et al. mentioned that the equilibrium activity of Al,O, was considered to affect oxygen solubility in the aluminum nitride grains, the equilibrium activity of Al,O, decreasing with the increasing Y,OJAI,O, ratio as yttrium-rich aluminates are formed [5]. With regarded to silicon nitride-Ln,O, systems,
Yokota et al. reported that the thermal conductivity of silicon nitride increased as the impurities in the grains decreased by using a p-silicon nitride raw powder of high-purity [6]. Yokota also reported that since oxygen was the major impurity in the grains, the high thermal conductivity was thus achieved by reducing the oxygen content of the grains. Therefore, it can be said that the liquid phase (J phase) associated with high Yb,OJSiO, ratio, in otherwords ytterbium-rich silicates during the grain growth in comparison to K phase of other rare-earth oxides should have lower oxygen solubility in the silicon nitride grains. Additionally the high growth rate of silicon nitride-Yb,O, enhances the removal of oxygen from the grains. Hirosaki et al. reported that significant improvement of thermal conductivity is due to decrease in two grain junctions accompanying grain growth [l]. In this experiment, since the thermal conductivity of silicon nitride-Ln,O, increased with the area fraction of large grains, the high thermal conductivity is also here attributed to a decrease in two grain junctions as grain growth. Therefore, thermal conductivity of silicon nitrideLn,O, depends on the type of additive, the highest thermal conductivity of 121 Wm 'K' being achieved by the silicon nitride-Yb,O,.
(2) Compatiblity between high thermal conductivity and high mechanical strength
s3
600
Powder(b)
cj
500
It is appears from this data that it is difficult to obtain high thermal conductivity and good mechanical properties with a bimodal microstructure. This would be because large elongated grains are considered to be the origin of cracks, when flexural testing is done. Therefore, it can be said that the achievement of both high thermal conductivity and sufficient mechanical properties by using the powder (a) due to the homogeneous microstructure of the sintered material achieved thereby . Emoto et al. investigated that the grain growth behavior of fine-grained p-silicon nitride with varying amounts of nuclei, and they concluded that the grain growth driving force depended on the amount of nuclei in the silicon nitride raw powder [7]. In this experiment as well, the difference of the microstructure of the sintered materials between powder (a) and (b) is considered to be the amount of nuclei in the silicon nitride raw powders.
SUMMARY We have newly developed a silicon nitride substrate with a thermal conductivity of 100 Wm-'K'. A high thermal conductivity with sufficient mechanical properties was achieved by processing the high purity raw powder of p-silicon nitride with ytterbium oxide as an additive and developing the homogeneous microstructure of a sintered material. This material densified to over 99 % of the theoretical density at the sintering temperature of 1800 "C for 8 h under a nitrogen pressure of 0.9 MPa. In this material, the flexural strength was between 550 and 600 MPa, and fracture toughness was 6 MPa-m"'. It is thus considered that this material could be an attractive substrate for the power modules.
REFERENCES
F: c
m
3 400 ia?
'.
3
L
300
,
'0
,
80 85 90 95 100 105 -1
-1
Thermal conductivity/Wm K
Fig.5 Relation between thermal conductivity and flexural strength. Fig.5 shows that relation between thermal conductivity and flexural strength. The sintered materials grown from powder (a) have both thermal conductivity of approximately 100 Wm%' and flexural strength of approximately 550 MPa in striking contrast to those grown from powder (b). Fracture toughness of the sintered materials grown from the powder (a) was 6 MPa m1l2
(1) N.Hirosaki, Y.Okamoto, M.Ando, EMunakata, and Y.Akiume, Thermal Conductivity of Gas-PressureSintered Silicon Nitride, J.AmCeruni.Soc., 79 [ll] (1996) 2878-82 (2) K.Hirao, K.Watari, M.E.Brito, M.Toriyama, and S.Kanzaki, High Thermal Conductivity in Silicon Nitride with Anisotropic Microstructure, .JAnr. Ceruni.Soc.,79 [9] (1996) 2485-SS (3) K.Watari, K.Hirao, M.E.Brito, M.Toriyama, and S.Kanzaki, Hot Isostatic Pressing to Increasing Thermal Conductivity of Si,N, Ceramics, J.Muter.Res., 14 [4] (1999) 1538-1541 (4) L.J.Gaukler, H.Hohnke, and T.Y.Tien, The System Si,N, J.AmCerum.Soc.,63 [35] (1973) (5) A.V.Virkar, T.B.Jackson, and R.A.Culter, Thermodynamic and Kinetic Effects and Oxygen Removal on the Thermal Conductivity of Alumminum Nitride, J.AnrCeruni.Soc., 72 [ 111 (1989) 2031-42 (6) H.Yokota and M.Ibukiyama, to be published. (7) H.Emoto and M.Mitomo, Control and
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Characterization of Abnormally Grown Grains in Silicon Nitride Ceramics, J.Euro.CerumSoc., [17](1997) 797-804
5 04
CHARACTERISATIONOF THE PORE STRUCTURE OF BIOMORPHIC CELLULAR SILICON CARBIDE DERIVED FROM WOOD BY MERCURY POROSIMETRY A.Hemg*', U.Vogt', T.Graule', T.Zimmermann2J.Sel12 'Department of High Performance Ceramics 2Departmentof Wood Swiss Federal Laboratories for Materials Testing and Research (EMPA); 8600Duebendorf, Switzerland
INTRODUCTION In the last decade growing interest in converting biological structures like paper and fibre like morphologies (rice husks and flax) into ceramic materials has been reported by several research groups"3'. A particular field of interest was in reproducing wood-like structured ceramics. Due to the highly anisotropic cellular appearance, which cannot easily be copied by artificial means, such materials might be attractive for filter and catalyst applications. Different conversion routes of pyrolysed wood (biological derived templates), mainly into Sic-ceramics via infiltration of gaseous silicon14',gaseous SiO'" or liquid silicon'"'' and sol derived silica'8' are known. Mainly reaction mechanisms and changes in porosity are monitored. The following paper describes the pore channel development, reaction mechanisms and degree of conversion of wood from original to pyrolysed to ceramic state with particular focus to wood thickness. The conversion from pyrolysed to ceramic state is carried out by silicon gas infiltration, the aim being full conversion to Sic. The obtained porous ceramics are tested with the focus being on oxidation stability and influence of oxidation on porosity and pore size distribution.
EXPERIMENTAL PROCEDURE The wood-derived ceramics were produced from two morphologically different kinds of wood species: deciduous (beech) and conifer (pine) trees. They differ mainly in the degree of uniformity of structure. The conifer structure is very uniform and consists of up to 95% similar sized tracheides whereas deciduous wood is more diverse regarding functions and morphology of the wood cells (vessels and fibres)'". The wood samples were sliced perpendicular to the fiber axis into sections with thicknesses of I , 2, 4 and 8mm, dried 70"C/15h, afterwards pyrolysed in Nzatmosphere with a heating rate of lWmin up to 500°C and SWmin up to 1400°C followed by a four hours hold. The resulting carbonised preforms were infiltrated in a graphite furnace by gaseous silicon under an argon
atmosphere. The heating rate was IOWmin up to 1200°C and SWmin up to peak temperatures of 1600, 1800 or 2000°C followed by a hold of 4 or 8 hours. Cyclic oxidation was carried out at 1200°C and weight changes were recorded after removing the specimens from the furnace. Characterisation methods used were x-ray diffraction for phase identification and SEM for morphological analysis. To determine the free carbon content specimens were oxidised at 1O0O0C/1h. Under these conditions weight gains due to oxidation of silicon and Sic are almost negligible. Consequently weight loss corresponds to free carbon. Porosity was examined by mercury porosimetry.
RESULTS AND DISCUSSION i) Starting material The measured mean porosity of dried beech wood is 56% and therefore slightly less than the porosity of pine with 73%. The distribution of pore sizes strongly depends on the thickness in axial direction of the examined pine wood samples, fig.app. 1. This phenomenon is well known and attributed to the socalled bottleneck effect'"": With increasing specimen thickness narrow ends of tracheides progressively influence the results of pore size measurement as some lumens in the interior consequently are not filled by mercury at the pressure corresponding to their pore diameter. However, they may be filled at pressure corresponding to the narrow pits at the walls of each tracheid cell. Consequently an increasing specimen thickness results in a decrease in pore size. This result often misleads the interpretation of porosity data. The solution is to keep specimen size as small as possible. For beech no difference in pore size distribution between 2 and 8mm thickness is noticed. Total porosity of the dried wood samples is independent of specimen thickness. ii) Pyrolysis Pyrolysis of dried wood is accompanied by sharp weight loss above 200"C, mainly in a temperature range up to 400"C, fig.app. 2. During the main weight loss a
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process of disruption or decomposition of the (bio)polymeric (cellulose, lignin) substance is initiated. On pyrolysis of the cellulosic structure small molecules such as water, CO2, aliphatic acids, carbonyls and alcohol's are volatilised. As small molecular species are removed from the originally macromolecular network, the resultant chemically reactive lattice 'zips up' around the vacancies which are left by volatises and the new lattice becomes more carbonaceous towards a more stable graphite state"". However, the x-ray diffraction patterns in fig.app. 3 gives little evidence for preferred orientation of carbon as graphite as peak pyrolysis temperature is raised from 600 to 1800°C. Only a slight increase in sharpness and height of the peaks at 28 = 26.2" at 1400 and 1800°C is observed which is assumed to indicate to the formation of graphite. The degree of graphite formation can not be concluded from this data. After pyrolysis skeletal densities for beech and pine of 1,5 g/cm3 on average are far below the density of graphite (2.26g/cm3). Thus it is proposed that the pyrolysed wood consists of a carbon structure which is not fully graphitic, tab. 1. Due to volatilisation shrinkage occurs depending on the orientation of the wood cells in correlation to radial, tangential and axial direction refemng to fig. 1. tab. 1 shows clearly that shrinkage parallel to orientation of the trunk axis (axial orientation) and wood cells is less than perpendicular to trunk axis (radial and tangential orientation).
stemming from incomplete conversion at contact areas between specimen and supporting furnace furniture where no infiltration took place, the measured mass increase of specimens of different thicknesses were normalised to lmm where fully conversion to Sic was observed. fig. 2 shows a decrease in conversion rate as pyrolysed wood preforms become thicker.
fig. I main orientations of wood A-axial, T-tangential, R-radial"*'
tab. 1 Shrinkage and weight loss of dried wood during pyrolysis Pine Beech Shrinkage Tangential 37 40 (%) Radial 31 30 Axial 23 21 Weight loss 74 75 (%)
Density (g/cm') Porosity
Apparent Skeletal
0.417 1.604 74
0.478 1.405 66
(%)
iii) Infiltration by Si(ga5) During infiltration the carbon of the pyrolysed wood reacts with silicon vapour
+ SiC(s)
The theoretical weight increase for the specimen for full conversion is 233m%. To eliminate measurement errors 506
.
O i
Weight loss and shrinkage during pyrolysis work in the opposite direction to the change in total porosity from wood to pyrolysed state. Therefore only a small increase in porosity (10%) during conversion from wood to char for beech and no porosity increase for pine are observed. Differences in pore size distribution as a function of slice thickness for pine are again observed (ref. part i, starting pine material), fig.app. 4. fig.app. 5 shows micrographs of the pyrolysed woods. Differences in pore size and pore distribution between a deciduous and conifer wood are evident even in pyrolysed state.
C(s) + Si(g)
-g 0.4 1
0
2
4
6
-
1 8
thickness of wood (mm)
fig. 2 influence of thickness on conversion to Sic (beech, 1800"C/4h) The rate of conversion depends as well on the partial pressure of silicon gas as on rate of diffusion of carbon and silicon through Sic after first layer of Sic is formed. Both factors are mainly ruled by temperature. Consequently as psi and the diffusivity of carbon and silicon rise with temperature more Sic is formed. As conversion to Sic is enhanced, carbon content drops down. Bum off of residual carbon after silicon infiltration by oxidation at lO00"C shows a decreasing carbon content as the peak reaction temperature rises, tab. 2. The lower carbon content of pine compared to beech points to a faster and more thorough conversion of pine. It is suspected that the different pore size distribution, pore diameter and mean cell wall thickness
(pine 0.5-Ipm, beech 1-2pm) of wood cells are responsible for this difference, fig.app. 5 . In fig.app. 5 it can also be seen, that the wood cell structure is completely reproduced while pyrocarbon is converted to Sic during silicon vapour infiltration. The microstructure of the S i c is similar for both wood species. Size of disk shaped S i c crystals ranges between 0.5-2.5pm, fig.app. 6. tab. 2 influence of temperature on residual carbon content of wood (%) derived ceramics, (8mm thickness) Pine Beech Peak temperature ("C) I hours __ 4 40 1600 4 9 36 1800 2000 4 7 23
The microstructure is further influenced by the proximity of the pyrolysed wood within a specimen to the silicon vapour source, fig.app. 7. The closer an area of wood within the specimen is to the source the faster and more thorough the conversion to SIC. Crystallite size is rising parallel as degree of conversion improves. The top of a 8mm thick beech specimen (furthest away from silicon vapour source) is only partially converted to S i c and free carbon remains in the structure, while areas close to silicon source are fully converted, tab. 3. Increasing the holding time at peak temperature from 4 to 8 hours improves ceramic yield and nearly all carbon, even in the upper parts, is converted to Sic. In order to understand the limit of conversion to S i c in the upper parts of 8mm thick pyrolysed woods, SEM examinations were carried out to find barrier of Si c or silicon, which was thought to suppress or minimise transport of silicon vapour. Inside the wood cells no such layer was found. However a barrier layer was found to exist on the specimen surface closest to the silicon source. Since deposition rate of silicon vapour on the wood close to the vapour source is very high, formation of S i c in these areas is very fast until all carbon is converted to Sic. The silicon condensing thereafter can not be transported into the structure and a silicon surface layer develops. This causes a drop in silicon vapour transport rate before carbon far away from the source is fully converted to Sic. Only by increasing the reaction time sufficient amounts of silicon can be transported to the upper regions of the pyrolysed wood specimen, as shown in tab. 3. Measurement of the pore size distribution of S i c derived from wood was carried out after removing the silicon layer. fig.app. 8 reveals a slight increase in pore size for both wood species as their state changes from pyrolysed to ceramic. Pores below 1pm are closed. Depending on the thickness of specimen total porosity decreases on average by 10% relative to the pyrolysed wood, tab. 4. The difference in total porosity might correlate to the ceramic yield during conversion from pyrolysed wood to ceramics: The smaller a specimen slice the better the conversion to S i c and the higher the decrease in porosity. Porosity decreases due to expansive character of S i c formation (AV= +56% [Ctemplate density 1 .5g/cm', SiC=3.2g/cm3]).
Significant differences in pore size distribution as a function of specimen thickness seen in the wooden and pyrolysed state, are not found either for pine or beech. tab. 3 carbon content of different parts of wood-derived Sic-ceramics as a function of position relative to the silicon source (8mm thickness) Wood Peak temposition' Residual species perature carbon (wt%) ("C) hour beech 1800 4 upper 66 lower 5 beech 1800 8 upper 10 lower 0 pine 1600 4 upper 55 lower 28 pine 1800 8 upper 0 lower 0 tab. 4 influence of specimen thickness on total porosity of wood-derived Sic-ceramics Thickness (mm) beech Pine 2mm 56 48 58 53 4mm 8mm 60 63 iv) Oxidation of wood-derived Sic-ceramics
Cyclic oxidation puts stresses on the material due to sudden temperature change as the specimen is cooled in air. Consequently thermal tension in the specimen is high, however the wood derived ceramics showed no appearance of cracks after cycling. During cyclic oxidation a rapid weight loss of 10% is measured for both wood species, indicating the burnoff of residual carbon, fig.app. 9. Subsequently a weight gain due to formation of a silica layer is observed. The pore size distribution and total porosity of both wood species remains constant during oxidation, fig.app. 10.
CONCLUSIONS The experimental results showed: 0 pyrolysis of wood followed by silicon vapour infiltration forms a highly porous ceramic the internal structure of wood is transformed to Si c with remarkable accuracy conversion of wood to ceramic can be controlled by adjusting the temperature and time during infiltration to account for wood thickness 0 the measured pore size distribution in the original and pyrolysed states depend strongly on wood thickness for coniferous wood; no such dependence is observed for beech 0 during conversion to S i c the pore size distribution and porosity changes 0 wood-derived ceramics are oxidation resistant and maintain their pore size distribution and total porosity.
' Refers to position realtive to the silicon vapour source: upper
part - far from source, lower part - close to source
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[ 13 Friedrich H. et al: Herstellung von keramischen Laminatstrukturen durch Pvrolyse von PolymerFullstoff infiltriertem PaDier, Verbundwerkstoffe und Werkstoffverbunde, DGM/Wiley-VCH, Weinheim, Germany, 1999, p.399-404 [2] Kleber S., Hermel W.: Keramik aus Flachs ein konkurrenzfahiges Technoloeiekonzept, 5.Internationale Tagung Stoffliche Nutzung nachwachsender Rohstoffe ,Chemnitz 1998 [3] Krishnaro R., Mahajan Y.:Conversion of raw rice husk to Sic bv pvrolvsis in nitrogen atmosphere, J. Europ. Ceram. SOC.18, 1998, p. 147 [4] Sieber H. et al: Biomorphic cellular ceramics, Advanced Engineering Materials 3,2000, p. 105-109 [5] Mukerji J. et al: Conversion of oak to cellular silicon carbide ceramic bv vapour Dhase reaction with SiO, submitted to J. Amer. Ceram. Soc. 2000 [6] Shin D. et al: Silicon/silicon carbide comDosites fabricated by infiltration of a silicon melt into charcoal, J. Am. Ceram. Soc. 82, 1 1, 1999, p. 32513253
[7] Greil P. et al: Biomorphic cellular silicon carbide ceramics from wood, part I+II, J. Europ. Ceram. SW. 18, 1998, p. 1961-1983 Ota T. et al: Biomimetic Drocess for producing [8] Sic "wood", J. Amer. Ceram. Soc. 78, 12; 1995, p.3409-3411 [9] Wagenfuehr R.: Anatomie des Holzes, VEB FachbuchverlagLpz. 1989 [lo] Blankenhorn P. et al: Porositv and Dore size distribution of black cherry carbonized in an inert atmomhere, Wood science vol. 1 1,1 [ 1I ] Patrick, J.: Porositv in carbons, chp. 1, Hodder Headline Group, London 1995 [ 121 Sachsse H.: Einheimische Nutzholzer und ihre Bestimmung nach makroskoDischen Merkmalen, Parey Verlag Bin. 1984
1.8 1.6
1 ;:;
.- 0 1
3 0.8 5
0.6 0.4 0.2
0 0.0I
0.1
1
pore diamter (pm)
10
100
0.01
0.1
1
pore
(p)
fig.app. 1 influence of wood thickness on porosity of beech (left) and pine (right), dried at 70°C/15h
508
10
100
CJ-
c?
I
I
I
I
I
I
1
I
8
I
I
I
t
rn Ln m
0.01
1
0.1
pore diameter (p)
10
0.01
0.1
10
1
100
pore d w r ( C I I 1 1 )
fig.app. 4 influence of wood thickness on porosity of pyrolysed beech (left)and pine (right),
N,
1400OC
fig. app. 5 continued next page
509
fig.app. 5 SEM micrographs of pyrolysed and
infiltrated pine (left) and beech (right)
fig.app. 6 Morphology of S i c derived from pyrolysed pine (left) and beech (tight), close to silicon vapour source
fig.app. 7 influence of relative Si(,,-source
510
position on microstructure of S i c in beech, samples with 8mm thickness
0. I
0.01
10
1
pore d
10
0.0 1
0.1
10
1
100
pore dhm%e.r(rvn,
i m (pl)
fig.app. 8 influence of wood thickness on porosity of Si(gas)-infiltratedbeech (left) and pine (right), pyrolysis: N ,1400°c, infiltration: argon, 1800/4h
0
-12
0
2Ooo
lo00
3Ooo
4Ooo
time (min) fig.app. 9 oxidation ( I 200°C) of S i c derived from pine and beech via S&,) infiltration (4mm. 1800"C/4h), SEM of beech after oxidation 1200'C/64h
1.8
,
0.01
0.1
1
pore d m t e r (rUn,
10
11
0.01
0.1
I
10
loo
pore d m t e r (p)
fig.app. 10 pore size distribution of wood-derived ceramic from beech (left) and pine (right) after oxidation 1200"C/64h, 4mm thickness
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TBC CONSISTING OF NEW METAL-GLASS COMPOSITES M. Dietrich*, V. Verlotski, R. VaDen, D. Stover
Institut fur Werkstoffe und Verfahren der Energietechnik, Forschungszentrum Jiilich GmbH, D-52425 Jiilich, Germany.
ABSTRACT A new concept for thermal barrier coating (TBC) systems is presented, based on a metal- glass composite (MGC). In this TBC system the composition of the MGC is chosen in such a way that the thermal expansion coefficient of the composite is close to the one of the substrate. This leads to reduced thermal stresses and hence improved thermal cycling life times. Another advantage of the gas tight composite coatings is their ability to protect the bondcoat from severe oxidation. Correspondingly, longer life times have been found for these TBCs in oxidation tests. In this paper measurements of thermal properties as well as results of oxidation experiments at high temperatures and thermal cycling tests will be presented. Additionally some aspects of the evolution of the microstructure during thermal aging is shown.
one of the most important sources of failure in TBCsystems. A way to overcome this problem is to improve the thermal expansion match with the substrate material and to obtain a good adherence. This could be obtainable with a metal-glass composite (MGC). In this material it is possible to adjust the coefficient of thermal expansion in a wide range [7], by varying the ratio metayglass. In layer systems like TBCs the adaptation of its thermal expansion to that of the substrate material reduces the stresses, which are induced by thermal mismatch. An other advantage of the material is the absence of open porosity, permitting the protection of the substrate material and the bondcoat from corrosive gases8. Up to now TBCs are permeable for these gases, exposing the bond coat to oxidation and leading to the formation of a TGO. Corresponding to such a gas tightness, longer lifetimes for TBCs could be expected during thermal aging in oxidizing atmosphere.
INTRODUCTION TBCs find an increasing number of applications to protect high-temperature metallic components; for example, TBCs are deposited on components of gas turbines and diesel engines either to increase the inlet temperature with a consequent improvement of the efficiency or to reduce the requirements for a cooling system [ l ] or to enhance the lifetime. The selection of TBC materials is very specific [2] and some basic requirements are low thermal conductivity, no phase transformation accompanied by volume change during heating, chemical inertness, low thermal expansion mismatch with the metallic substrates and good adherence [3]. The best compromise among these requirements is presently offered by partially stabilized zirconia as 7-8 wt% Y,O,-ZrO, (YSZ) on top of a MCrAlY bond coat, deposited either by air plasma spraying (APS) or by electron beam physical vapor deposition (EB-PVD) [4]. However, the application of YSZ-TBCs is limited by two factors. At higher temperatures than 1200"C, phase transformations from tetragonal (t') to tetragonal and cubic (t + c) and then to monoclinic (m) occur, giving rise to the formation of cracks in the coating [5]. An other factor for failure is the insufficient oxidation resistance of the substrate which has to be protected by the gas tight MCrAlY-bondcoat (BC). This BC is forming a thermally grown oxide (TGO) layer [6]. The stresses associated to the growth of this Al,O,-layer are presumed to induce spalling of the coating and so to be
EXPERIMENTAL For the preparation of metal-glass-composit coatings (MGC), industrial available bondcoat powder and ordinary white container glass were chosen. In order to use plasma spraying for the deposition of metal-glass composites, the powders had to be adapted. Taking into account the metallic component of the composites, only vacuum plasma spraying (VPS) was a possible spraying technique for these powders. Data of raw materials for the preparation of metal glass powder mixtures are compiled in Table 1. The thermal expansion coefficient of composites with different metal-glass ratio has been determined by dilatometer measurements. The preparation of the powder starts by ball milling glass- and metal powder with zirconia balls for 24 hours in ethanol. After drying, the powder is homogenized by ball milling with steel balls and subsequently annealed under HJAr-atmosphere at 700" for 2h. The slightly densified and agglomerated powder is then dry coarse milled in a mortar mill and finally a powder fraction of 40-90 pm is sieved. For the deposition of the composite powder, a VPS equipment from Sulzer Metco with a F4 gun was used. The powder was sprayed on Inconel 617 and Inconel 783 substrates under 60 mbar Aratmosphere. The thickness of the MGC-TBC varied between 500 and 600 nun.
513
I NiCoCrAlY I powder Grain size
Content in theMGC [wt.%]
I
I
Silicate glass powder
25 - 40
5 - 10
63
37
I
I
Composition 29.7% Ni [MolYo] 31% Co 30% Cr 8% A1 0.7% si
71% S O 2 14% NqO + K20
10% CaO
RESULTS AND DISCUSION Dilatometer measurement showed the strong influence of the metal content on the thermal expansion coefficient of the MGC. Fore the pure glass it is 9,5* 10K-' in the temperature range of 20°C to 300°C. This value is rising with the metal content and reaches 12,3*10-6K-'in the temperature range of 20°C to 7OOOC at 68 wt% corresponding to 35 ~01%.Higher metal contents were not used in order to guaranty complete covering of the metal particles by the glass matrix. An optical micrograph of the cross-section of an as sprayed MGC-TBC shows the typical lamellae structure of plasma-sprayed coatings where the light component is metal, the gray one is glass and black represents pores (Figure 1).
Table 1. Composition of the MGC
Two types of oxidation tests were carried out in air isothermal exposure of samples in a furnace and oxidation in a thermal gradient across the sample thickness. The first test was carried out at 1000°C for 1000 or 1500 h. The samples were taken out of the furnace every 24h, natural cooled, controlled and then taken back in the furnace. The temperature was chosen in order to simulate the working conditions at the TBC/BC Interface. In the gradient test TBC surface temperatures of 1200°C were achieved with a stable substrate temperature level below 1000°C. Taking into account the thickness and the thermal conductivity of the substrate, the temperature gradient across the TBC was about 200K. With this setup the thermal conductivity of the MGCTBC were estimated using the equation h(T) = (dq/dA) * (dx/dT) The temperatures were measured directly on the TBC surface and the substrate back-side with thermocouples; dq was obtained by a measurement with a YSZ standard sample with known thermal conductivity. Further more a burner rig test was carried out using a methane flame. The sample was cycled on air by raising the surface temperature from room temperature to 1200°C within 30 seconds and with 5 minutes dwell time while the back surface was cooled by air. Subsequently cooling to room temperature for 2 minutes was enforced by cooling air jets from both sides. The microstructure of the TBC was found to change during thermal aging. This change was studied with seven samples prepared in an identical manner and exposed during an oxidation test to 1150°C in air. The exposure time was varied from 1 to 300 h. The changes were followed by scanning electron microscopy (SEM), energy dispersive X-ray microanalysis (EDX) and Xray powder diffraction (XRD).
5 14
I
1ooum Figure 1. As sprayed MGC-TBC
The uppermost layer does not show any microcracks, neither were inclusions or oxides found at the interface. The porosity is a closed porosity of about 5% During annealing at 1000°C which is the BC temperature in technical applications, grain coarsening without formation of cracks or spalling was observed Figure 2.
1
* '.
d
'
s
** *r
1Ohm
Figure 2. Plasmasprayed MGC-TBC after annealing at 1000°C for 300h
I
Furthermore a metal-deficient zone of 10-40 mm thickness is formed on the MGC-TBC surface, within the first hour of exposure (Figure 3). The thickness of this depletion zone does not increase with further annealing. Until now, there is no explanation for its formation.
Figure 3 Metal deficient layer at the surface of an annealed MGC-sample
Several XRD measurements, made during removing layer by layer in steps of 30pm, showed the growth of a-Al,O, mainly from the interface TBC/BC into the glass matrix. An other phase found in the glass matrix was Anorthite (CaAl,Si,O,) It was seen that A1 atoms diffuse from the BC and from the metal particles of the MGC into the glassy matrix. It is presumed that the difference of aluminum solubility in the glass and the alloy, combined with the oxidation of the aluminum in the glass are the driving force for this diffusion. A XRD study of the surface of a MGCTBC after lOOh oxidation at 1100°C in air shows reflections belonging to a-AlzO3. This indicates that at least a fraction of A1 is oxidized to a-Al203. The addition of A1203 to the glass corresponds to a shift in the composition leading to a preferential crystallization the two phases.. Plasma-spraying is similar to fast heating-up and cooling down. During processing there is no time for A1 diffusion into the glassy matrix. Diffusion, formation of alumina and crystallization of the glass can take place only during subsequent long-time aging at high temperature in air.
-
Figure 4. A1 K a X-mapping of annealed MGC-TBC: Evolution of tne interlayer between TBC and BC Quantitative EDX mappings of A1-K radiation (Figure 4) shows the evolution of this layer between TBC and BC depending on the annealing time. EDX measurements of the layer revealed that one major element of the layer is Al.
-
MGC and may contribute to the formation of the intermediate Al-rich layer at the interface MGC/bondcoat. The growth of this intermediate layer starts quickly and seems to stop within the first 100h. A reason could be the crystallization of the glass, slowing down diffusion. After calibration of the gradient furnace setup the specific thermal conductivity of a MGC-TBC has been determined to 1.4f0.2 W/mK. In comparison, the corresponding value for porous YSZ-TBC is 0.8 W/mK. It follows that in order to achieve equal thermal barrier properties the thickness of a MGC-TBC must be nearly twice that of a traditional YSZ-TBC. The results of oxidation and thermal cycling tests can be summarized as follows : Plasma-sprayed MGC TBC on a IN738 alloy substrate did not show spalling during isothermal exposure to air at 1000°C for 1500h. YSZ-TBC failed due to strong spalling after 900h under identical experimental conditions. Plasma-sprayed MGC-TBC withstand isothermal oxidation tests in air at about 1200°C for more after 300h before failing. Plasma-sprayed MGCTBC do not show spalling during oxidation test in a thermal gradient furnace, even above 300h at 1250/1O0O0Cwhile YSZ-TBC fail already after 200h. Similar to YSZ-TBC, plasma-sprayed MGC coating did not fail during ther&al cycling until 1000 cycles-The good cycling behavior even with 600 pm thick coatings can be explained with the high thermal expansion coefficient which is 12.3*10-6K-'. Burner rig tests without BC showed very poor results for YSZ as well as for MGC. This indicates, that the gas tightness of the MGC alone is not sufficient for the
515
oxidation protection of the substrate. The results are listed in Table 2.
YSZ MGC YSZ MGC
MGC YSZ without BC MGC without BC
1000 1000 1200 1200
900 >1500 85 300
1200/1000
>loo0
1ooooc
20
1ooooc
250
Table 2. Results from thermal tests
CONCLUSION Metal-Glass Composite (MGC) was also found to be very promising for the use as TBC material. During thermal aging tests, the gas tight MGC coating seems to protect the substrate material better from oxidation than classical YSZ TBC does, leading to longer life times in isothermal oxidation tests. Even with a thickness up to 600 pm the adherence of the coating and its spalling resistance in thermal cycling was demonstrated to be comparable to the corresponding results for YSZ coatings . A better understanding of the microstructural and phase evolution during thermal aging and measurement of thermal and mechanical properties are subject of ongoing work. Up to now no experiments have been done in excess of 1250OC. Passing to higher temperature is the aim of further development.
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(1) P.Hancock, M. Malik, Materials for Advanced Power Engineering. Part 1.D.Coutsouradis et al. (eds.), Kluwer Acad. Publishers, Dordrecht, (1994) 685 ff (2) G.A. Kool, Current and Future Materials in Advanced Gas Turbine Engines. J. Therm. Spray Tech., 5 [ l ] (1996) 31-34 (3) V. Arnault et al., Thermal Barrier Coating for Aircraft Turbine Airfoils: Thermal Challenge and Materials. Rev. MCtallurgie - CITI Science et GCnie des MatCriaux, 96 [5] (1999) 585-597 (4) S.Stecura, Optimisation of the Ni-Cr-Al-Y/ZrO,Y203Thermal Barrier System. Adv. Cer. Mat., 1 [ 11 (1986) 68-76 (5) R.A. Miller, J.L. Smialek, R.G. Garlick, Phase Stability in Plasma Sprayed Partially Stabilized Zirconia-Yttria. Science and Technology of Zirconia, Advances in Ceramics,Vol.3, A.H. Heuer and L.W. Hobbs (eds.), Am. Cer. SOC.,Columbus, OH, USA (1981) 241-251 (6) R. Miller, Oxidation-Based Model for Thermal Bamer Coating Life. J. Am.Cer.Soc., 67 [8] (1984) 517-521 (7) V. Verlotski, D. Stover, H.P. Buchkremer, R. VaBen, Warmedtimmende Glas-MetalVKeramikSchichten. German Patent No. 198 52 285, Date: May 3rd,2000 (8) R. VaBen, X.Q.Cao, V. Verlotski, H. Lehmann, M. Dietrich, D. Stover, Two new candidates for thermal barrier coatings. To be published in Surf. Coat. Tech.
ASPECTS ON SINTERING OF EB-PVD TBCS Klaus Fritscher, Uwe Schulz, Christoph Leyens and Manfred Peters
DLR German Aerospace Center, Institute of Materials Research, D-5 1170 Koln, Germany INTRODUCTION Thermal barrier coatings (TBCs) of partially yttria stabilized zirconia (PYSZ) on turbine airfoils are one of the most promising recent approaches to improve the overall economy of fossil fuel dependent energy conversion systems. The benefits result from reduced fuel consumption by higher efficiencies, prolonged service intervals as well as lifetimes. Today the most durable TBCs on rotating turbine parts are PYSZ coatings that are applied by electron-beam physical vapor deposition (EB-PVD). The EB-PVD process procures the most superior strain and thermoshock tolerant coatings due to their unique columnar microstructure. In order to integrate these TBCs into the design of modem aero engine parts the manufacture of so-called “designed-$ TBCs - a quantitative understanding of the correlation between processing, microstructures, and the resulting lifetime of the coatings is advised. Mechanisms which affect the microstructures have to be understood.
ing of TBCs on their substrates can be obtained by (electron) microscopic observations. Indnect evidence is feasible by specific surface analysis methods. Mercury intrusion, gas permeametry and gas adsorption techniques are commonly used analytical methods for the characterization of porous media, such as powders and green or sintered compacts. Their general capability to the specific surface area determination in adherent coatings on substrates, however, is limited. Gas adsorption is suited best and was used to characterize plasma sprayed coatings [51. High-rate EB-PVD PYSZ TBCs on alumina substrates were investigated in this study. The specific surface area of the respective TBCs was determined by the BET method via inert gas adsorption described in [6] as a function of sintering temperature and time.
EXPERIMENTAL
Sintering of the ceramic top layer is a mechanism that can impair vital thermal and mechanical properties in TBCs. Sintering contributes to a reduction in thermal isolation [l] and an increase of the in-plain modulus of elasticity [2]. Sintering phenomena were identified during thermocyclic exposures. During cycling oxide columns were observed to crack intracolumnarly, and sintering of the crack faces occurred [3]. Sintering needs to be defined in order to clanfy what can happen in PYSZ TBCs. In pure single components like PYSZ it takes place completely in the solid state. Besides solid state reactions also vapor formation and re-condensation as means of transport may occur. Sintering may be accompanied by shrinkage, leading to densification. But sintering and shrinkage are by no means identical. Apparently no definition of sintering exists that takes full account of all theoretical and practical aspects and the various stages of sintering. Sintering is a fairly complicated process that involves several mechanisms of material transport. A basic definition of sintering can be given like that: it is a thermally activated material transport in porous matter that aims at a reduction of the specific surface area.
6mm Qameter and 20mm long cylindrical recrystalized fully-dense alumina tubes were EB-PVD coated with PYSZ TBCs of approximately 220pm thickness. Recrystallized alumina was taken instead of metallic substrates to minimize diffusive contamination of the TBC and to preclude similar effects of thermally grown oxides on the heated metallic substrates from the inner surface of TBCs [7]. The TBCs were applied by high rate EB-PVD in a rotary mode, as is conventionally used for TBC deposition on turbine parts [8]. The denomination “TBCs” in t h ~ spaper will exclusively indicate this type of EB-PVD TBCs. The total weight of the TBCs on the alumina substrate was typically 0.25g. The specific surface of the samples was measured in the as-coated condition. They were then heated isothermally in air at 700, 800, 900, 1000 and 1100°C. They were taken from the furnace after a definite time for specific surface measurements and redeposited into the furnace. This procedure was repeated several times to obtain a consistent plot of data for each temperature.
Macroscopic shrinkage was observed by thermoanalytical dilatometry on free-standing plasmasprayed PYSZ TBCs at >9OO0C [2] and at 1038°C [4] but will scarcely happen on fully adherent TBCs in this temperature range. Direct evidence of sinter-
The specific surface area of as-deposited and of airannealed samples was determined by the volumetric BET method (ASAP 2000, Micrometrics, Georgia). Nitrogen or krypton as adsorbent gas was used depending on the range of surface area expected. Spe-
517
cific surface area values were obtained by multipoint determination (five points) from the adsorption isotherm. The pressure range was 0.1 to 0.3 atm.
RESULTS AND DISCUSSION Mechanisms The variation of the specific surface area of PYSZ TBCs versus time at temperatures between 700 and 1100°Cis given in Fig. 1.
concluded that on sinteringbigger pores will grow at the expense of the thinner pores to lower the inner specific surface area. And the resultant bigger pores will be attracted more and more by the surface of the columns for the same reason of overall surface reduction. There they become converted into open porosity and contribute to some increase of the specific surface area. The effect of open porosity formation on sintering follows classical sintering sequences but was not quantitativelyaddressed in this study.
700"
PYSZ TBCs on alumina substrate
1
0,l
10
Time in hours
100
1000
Fig. 1: Specific surface area of EB-PVD processed PYSZ TBCs on recrystallized alumina substrate between 700 and 1100°Cversus time It shows that at higher temperatures a lower specific surface area is attained, and the time required for fading down to lower specific surface areas is reduced. Vice versa, at lower temperatures, the specific surface areas remain higher and the times to come up to low values are longer. After short annealing times at lower temperatures of 700 and 800OC a transient increase in the specific surface area was observed. This effect seems to be in some agreement to initial sintering stages of some powder compacts that increase in volume due to rearranging of grains. A rearrangement of grains or a shift in the crystallographic planes of the columns appears unlikely as they are essentially monocrystals. They are still in a low-energetic state, and therefore no driving force is likely to alter the pertinent stage of energy. The following explanation is more likely. Microscopic observations of sintered columns have shown the presence of numerous sub-micron open pores in the surface. Similar open pores were not observed on columns of as-received columns, so they obviously formed on sintering. Instead, the formation of a finely dispersed closed porosity mainly along the center of the columns as well as in line with the joints between the feather arms (in the following the feather arms will be denominated "secondary columns") is confirmed by transmission electron-microscopic observations [9]. The closed pores are typically between only 5 and 50nm. It is
518
1
10 100 Time in hours
1000
Fig. 2: Specific surface area recession (So - S)/So of EB-PVD processed PYSZ TBCs on recrystallized alumina substrate between 700 and 1100°Cversus time
In Fig. 2 the values for the specific surface S (denomination of the actual speclfic surface is S , the specific surface of the starting TBC material is So) of Fig. 1 are used to calculate the (So- S)/So ratio to bring about the specific surface area recession versus time. The values are fairly good representatives for the initial stages of sintering mainly at 700 and 800°C where shrinkage and densification cannot be noted. If the transport processes in sintering have stabilized straight lines in the double log graph are expected as will be briefly explained [ 101. In the early stages of sintering, preferably of small powders, the parameter (So - S)/So can be used to identhe sinteringmechanism
((So - S)/S"?
=
c
'
t
(1)
where C is a kinetic term that includes mass transport constants, t is the sintering time and v is an exponent that is close to n/2 for small powders, where n is the exponent in the initial-stage sintering equation (the sintering stage that is controlled by neck growth). n depends on the mechanism of mass transport. E. g. n is 7 for surface-diffusion controlled mechanisms, and other mechanisms provide lower n exponents. The inclination of the line plots at 1100°C in Fig. 2 is shallow and uniform. The number of data, however, is restricted, and their accuracy depends on the mass of the TBCs of 0.25g that is at least one order of magnitude too small if compared to recommended
mass weights for giving values with small standard deviations. Moreover, the accuracy drops more and more on repeated sintering of the samples due to their reduction of the specific surface area. In some of these cases of BET values below 1mVg the adsorbent had to be changed from nitrogen to krypton for the sake of better resolution. These values turned out to be of the order of 5 to 10% lower than with nitrogen, which, however, was estimated to be within the range of scatter of these data and was taken here to be of minor concern. But the low inclination of the plots in Fig. 2 still suggests that it is indicative of dominating surface diffusion control. The value for v = 2n will then give the best fit. It must be born in mind that the dependencies of equation (1) are applicable to initial sintering of small-grained powder compacts. The basic question in this context is whether the dependencies can be transferred to anisotropic columnar particles that are apart from each other having secondary minute columnar protrusions all over their surfaces. And only a few of these secondary columns are contacting others. This situation is really different from the one as for small particles. So the conclusion of predominant surface-dfision control is still speculative. During sintering the specific surface area declines rapidly from an initial value So to an intermediate value S. So is given as an average value in Fig. 1 and is more detailed by the values of the respective samples in Table 1. These data also show the scatter of So of individual BET measurements on identical samples from the same EB-PVD batch.
Annealing temperature
So
s
im mz/g
im mz/g
PYSZ PYSZ
1000
4.619
PYSZ
1100
Material
(s"-s)/s" after 100 h
0.914 0 92
The results on sintering show a dramatic reduction in surface area in Fig. 1 or, alternatively, an increase of the surface area reduction pamneter (So- S)/So in Fig. 2 at all temperatures. The values are most prominent for the highest temperature of 1100°C. No equivalent effect in dilatation analytical work was obtained on sintering in the temperature range 2 900°C [2]. These quantities for specific surface changes are unexpectedly high with regard to the minor macroscopic effects. The flux of matter during sintering must have been directed along superficial short-circuit diffusion paths because volumerelated effects were not observed at < 900°C. These
observations give some more support for the suggestion of dominating surface diffusion control.
.
................................................
.-
TBC on alumina
I
j
L
m
I
1 10 Time in hours
0,1
100
Fig. 3: Specific surface area recession (So - S)/So of EB-PVD processed PYSZ TBCs on recrystallized alumina substrate compared to free-standing identical microstructures at 1100°Cversus time Fig. 3 shows the difference of sintering of TBCs at 1100°C on alumina substrates and without substrates as free standing coatings. The differences are apparent during the first hour. Then the values become equivalent. It indicates that the straight positions of the columns on a substrate accelerate sintering only at the beginning. Later the columns show no different behavior. Some columns apparently exhibited an activated sintering mode probably due to forced contact with neighboring columns and have arranged afterwards. The same situation does not happen with free standing coatings where stress relief on contact points before sintering is feasible. Fig. 4 shows the impact of annealing in laboratory Pa on the degree of surface reduction. The vacuum atmosphere clearly retards sintering. The light gray color of the vacuum-annealed TBC is indicative for loss of oxygen &om the zirconia lattice.
air and vacuum of approximately
1000°C, 100 h
c
air
vacuum
Annealing atmosphere
Fig. 4: The influence of the atmosphere air and vacuum during annealing at 1000°Cfor 100 h on the specific surface area recession indicating less sintering in vacuum
519
The dependency of the migration of matter on the oxygen partial pressure needs some comment. The predominant point defect in partially stabilized zirmnia are double charged oxygen vacancies in a Schottky-type disordered lattice. Oxygen is quickly transported through zirconia as is well known. But migration of matter is controlled by the slowest moving species. These are the cations. Creep experiments with single-phase zirconia show, however, that the creep rate is independent on the type and concentration of the stabilizing element. But it depends on the oxygen pressure [1ll. From these observations it was concluded that the flux of matter is controlled by neutral cation vacancies, e. g. VZr*, that are, of course, a minority defect species. Their concentration can be described - in full analogy to equations with charged ions - by the mass law with reference to the two main elements according [V,,* 3 [V0*]2 = const.
0
0
EI
p-“4
0
(2)
(3)
0
of ionic conductionvia oxygen holes in PYSZ that is also experimentally confirmed by [121. Low oxygen pressures will therefore increase the amount of migration-retarding neutral Zr vacancies according to (2) and hence retard sintering in a low oxygen environment. This explanation is concluded from indirect and sometimes speculative information in literature and has to be taken with restraint.
Favorite sites in TBCs for sintering: Information from microscopic observations on the (scanning) electron microscope Microstructural investigations are necessary to iden@ the location and to understand the kinetics of
sintering.. Formation of TBC microstructure via the EB-PVD process and service life can be understood as continuous dynamic processes in high temperature environments, which have fundamental similarities. An imagination of what is happening in the course of vapor deposition will make the investigator sensitive for the mechanisms that rule the progressive sintering on high temperature exposure. Three observations are reported from electron microscopic examinations: 0
520
Inner pores of typically 5 to 50nm diameter are found within the main columns that have approximately 3 to 30pm diameter. Many of the pores are “in line” with the open boundaries between the secondary columns [9] as just mentioned in the last chapter.
The “young” secondary columns in close neighborhood to the top with the habit planes on it are the most peaky having the highest aspect ratio [131. They seem to be usually smoother at lower parts of the columns.
The observationssuggest the following conclusions:
The constant depends on temperature and on the oxygen pressure according to p-ln. This leads to the dependency from the oxygen pressure [V0*]
Sub-micron thick secondary columns are observed all over the primary columns but not at the prominent top of the columns.
0
The pores in line with the interstices closely resemble the well known simulation experiments of pore formation according to the three- or multi-particle interaction models. It suggests that neighboring high-aspect ratio secondary columns come in contact to each other at several contact points (probably in the course of ongoing deposition of vapor along the flanks of the secondary columns). The contact points will then exhibit progressive “neck formation”. Pores will form. This mechanism is suggested to be also probable on sintering. The formation of peaky secondary columns is a dynamic growth process due to shadowing and deposition and simultaneous smoothing of the peaks by the tendency for surface reduction. The peaks have grown the quickest on top of the columns and slower apart from the top due to shadowing by the prominent top of the columns. The interstices between the secondary columns are concluded to be most effectively filled up at the top of the columns by the tremendous vapor flux on high-rate vapor deposition. Thus they form habit planes instead. The time of formation of most peaky structures at high temperatures is restricted as they become smoothened by diffusional loss. This fact is related to general observations that surface atoms have their highest mobility on low radius concave surfaces by enhanced vacancy injection. These locations of extensive surface diffusion continuously attract matter from surrounding surface arm and mainly from the peaks. Thus the peaks of the secondary columns are the most prominent points to be susceptible to sintering on contact with other surfaces. The more peaky they are the more effective they promote sintering. The secondary columns at the lower part of the main columns are “older” than those near the top, and due to that they were smoothened by off-diffusion over a longer time during the processing step.
Modeling of the TBC microstructure The microstructure of EB-PVD TBCs is commonly assumed to be composed of parallel cylinder shapes as is explicitly suggested by the denomination "columnar". If the close paclung of identical cylinders in TBCs is assumed, similar to cigarettes closely matched in a box, a paclung density of 78.5% is attained. Ths value agrees reasonably well with the real packing density in TBCs. But if a representative value for the specific area in TBCs of 5mz/gis taken the diameter of the parallel cylinder shapes will be only 0 . 2 ~This . value is one to two orders of magnitude too low. Thus this approach will be rejected. A better fitting conceptual design of a simple model
shape of an EB-PVD TBC is based on pyramidal forms, as given in Fig. 5b).
overall specific surface area of as-coated EB-PVD TBCs is twenty-fold. The model is used to demonstrate the main impact of secondary columns on the specific surface area. It addresses the cuboidal habit of the columns (see Fig. 5.c.) as well as their tapered formation on competitive growth from the vapor phase (see Fig. 5.a.). A specific surface area for the model shape, however, turns out to be 20 times less than the real specific surface areas measured. The divergence between calculated and measured data has to be attributed to the neglection of the surface area of secondary columns in this concept and must be taken into account. If the secondary columns are also assumed to be pyramids for simplicity, so a fair coincidence of real and calculated data is obtained if the intermediate aspect ratio of the secondary columns of the order of 1 : 20 is considered.
CONCLUSIONS FROM MODELING ON FAVORITE SINTERING SITES The model clearly suggests that the secondary columns and their aspect ratio cannot be ignored when discussing sintering of TBCs. They have the main impact on the overall BET surface. So a decline of the specific surface area on sintering can mainly/exclusivelybe attributed to a reduction of the aspect ratio of the secondary columns. The aspect ratios on prolonged sintering can be calculated from the BET values and be given as a function of temperature and time. So a decline from 5 to 0.5m2/gis taken equivalent to an aspect ratio reduction of the secondary columns from 1 : 20 to 1 : 2. There may be (minor) dfierences between real and calculated aspect ratios. Among others, they may result from deviating cross sections of the secondary columns from pyramidal to more lamellar forms in the asdeposited condition that transform to more circular shapes on high temperature exposure.
SUMMARY
Fig.5: Design of a pyramidal model shape (in 5b.) for calculating the specific surface area (the contribution of secondary columns excluded) of an EB-PVD TBC as in 5.a (cross-sectional view) and 5.c (top view). Proposed coating thickness is H = 400p.m consisting of six layers of the respective thickness h, N2, W4, N8, N16 and N32, the basic plane of approximately l O p m being considered solid without porosity. The contribution of the secondary columns on the
Sintering of high-rate EB-PVD TBCs on alumina substrates was investigated through specific surface recession measurements between 700 and 1100°C by gas adsorption methods. The specific surface recession is dependent on time and temperature and was highest at 1100°C. At 700 and 800°C an intermittent increase before subsequent surface reduction was observed. Less specific surface recession occurs in vacuum. Specific surface calculations suggest that the specific surface recession is dominated by the secondary columns that reduce their aspect ratio on sintering.
52 1
Acknowledgments BET measurements were carried out by Dr. L.-M. Berger, Fraunhofer Institute of Ceramic Technologies and Sintered Materials, Dresden. Careful sample preparation by J. Brien, C. Kroder, H. Mangers and H. Schurmann is highly appreciated.
References 1. R.B. Dinwiddle, S.C. Beecher, W.D. Porter,
2.
3. 4. 5.
522
B.A. Nagaraj: The effect of thermal aging on the thermal conductivity of plasma sprayed and EB-PVD thermal barrier coatings, ASME paper 96-GT-282 (1996) F. Sziics: Thermomechanische Analyse und Modellierung plasmagespritzter und EB-PVD aufgedampfter Wi4rmedWmschichtsystemefiir Gasturbinen, Fortschr.-Ber. VDI Reihe 5 Nr. 5 18. Dusseldorf. VDI Verlag 1998 J.A. Haynes, M.K. Ferber, W.D. Porter, E.D. Rigney, Oxid. Metals 52 (1999) 31-76 H.E. Eaton, R.C. Novak, Surf.Coat. Technol. 32 (1987) 227-236 L.M. Berger: Powder and compact porous materials characterizationby adsorption, Roc. 1992 Powder Metallurgy World Congress, Vo1.8, J.M. Capus and R.M. German eds.,
MPIF/APMl 1992,235-247 6. S. Brunauer, P.H. Emmett, E. Teller, J. Amer. Chem. SOC.60 (1938)93 7. K. Fritscher, M. Schmucker, C. Leyens, U. Schulz, Mater. Sci. Forum 251-254 (1997) 965-970 8. U. Schulz, M. Schmucker, Mater. Sci. Eng. A276 (2OOO)l-8 9. L. Lelait: "Etude microstructurale fine de rev& tements ceramiques de type baniere thermique", Ph. D. thesis, Universite d Orsay, France 1991 10. R.M. German,Powder metallurgy science 2nd edition, MPIF Princeton, New Jersey 1994, p.250-252 11. M. S. Smelzer, P.K. Talty, J. Amer. Cer. SOC.58 (1975) 124-130 12. A. Knopp, N. Htife, W. Weppner, P. Kountouros, H. Schubert, Science and Technology of Zirconia V, Technomic Publication Comp., Lancaster, Pennsylvania 1993, p. 567-575 13. U. Schulz, K. Fritscher, C. Leyens: Two-source jumping beam evaporationfor advanced EB-PVD TBC systems, ICMCTF 2000, San Diego, accepted for publication
CRACK PROPAGATION IN A THERMAL BARRIER COATING SYSTEM G. Blandin*, R.W. Steinbrech, L. Singheiser Forschungzentrum Julich, Institut fur Werkstoffe und Verfahren der Energietechnik 52425 Julich, Germany
ABSTRACT Yttria stabilized zirconia (YSZ) coatings with a corrosion resistant MCrAlY bond coat provide an effective thermal protection for advanced gas turbine components of Ni-based superalloys. The crack propagation trough the top YSZ layer and the intermediate metallic bond coat of a thermal barrier coating system has been studied. Three-layer specimens - IN 792 substrate / NiCoCrAlY bond coat / air plasma sprayed YSZ thermal barrier coating (TBC) - were mounted on a four-point bending device with the ceramic layer on the tensile surface. Isothermal experiments were conducted between room temperature and 950°C. During deformation, the load/deflection curve was monitored and the propagation of cracks was observed in situ. At room temperature, cracks first appeared in the zirconia TBC at about 0.2 % strain. Under further deformation, the cracks grew through the whole ceramic layer, finally entering and causing fracture of the metallic bond coat. Residual stresses in ceramic and bond coats were determined by comparing strength values from literature with the measured bending stress at which a crack appeared in each layer. At high temperature, cracks stopped at the interface between ceramic TBC and metallic NiCoCrAlY layer. This behavior reflects the high ductility of the bond coat material. The change from brittle fracture to non-fracture of the bond coat was utilized to estimate the ductile-to-brittle transition temperature (DBTT) of the material. The stress/strain considerations used in the present approach and the residual stresses in the thermal barrier system are discussed.
INTRODUCTION Thermal barrier coatings (TBCs) are increasingly applied to gas turbine components to improve service life and efficiency [l]. The combination of a plasma sprayed yttria stabilized zirconia (YSZ) top coating and an oxidation resistant MCrAlY bond coat (M = Ni, Co) is one of the most commonly used TBC systems [2-31. The YSZ coating protects the turbine materials from hot spots and reduces the average temperature of the coated component, e.g. a 300 pm TBC reduces the temperature by 80°C [3]. The primary failure mode of TBCs is spallation, typically occurring after extended cyclic thermal exposure during cooling down from operation temperature [4]. As one of several failure mechanisms, spallation
results from cracks which grow along the interface between top and bond coats when the ceramic layer is under compressive stress. The determination of residual stresses in the YSZ top coat resulting from spraying process and thermal expansion misfit between the bonded layers [ 5 ] is essential to estimate the critical failure stresses of TBCs. In recent years, much work focussed on the characterization of residual stresses in TBC systems, using the established sin2€'-' and hole drilling methods, the layer removal technique [6] or the changes in curvature of TBC specimens [7-81. In the present study, tensile strain is applied to the three-layer composite in bending experiments. Using a four-point bending device, the specimens are strained between room temperature and 950°C. The cracking in ceramic TBC and metallic BC is observed in situ with a high temperature telescope system. The work focuses on the initiation and propagation of cracks in the top coat and their behavior at the interface with the bond coat. The residual stresses in top and bond coat are deduced from first appearance of cracking and strength values reported in the literature. The experimentally determined residual stresses are compared to a theoretical stress distribution in TBCs, calculated with a linear thermoelastic approach. Finally the behavior of the NiCoCrAlY material is discussed with respect to temperature and ductility.
THEORETICAL ASPECTS During bending, the stress at a position y of the composite material is given by: O ( Y )= ( y ) + O'f (Y)+ O B (Y) (1) where b0 is the residual stress, oT the thermal stress (when the experiment is performed at high temperature) and oEthe bending stress. Fig. 1 shows a three-layer specimen strip and the adopted geometric notations. The surface of the YSZ top coating is chosen as origin ( y = O ) .
Fig. 1: Geometry of a three-layer TBC composite
523
In this approach, materials are considered to be isotropic and homogeneous with a pure elastic response to external forces. As long as the three materials behave elastically, the three-layer system is representative of the TBC system (k=3). Due to thermal expansion mismatch, the TBC experiences tensile stresses at elevated temperatures, causing cracking of the TBC. Once the top coat is cracked, only the two metallic layers have to be taken into account ( k = 2 ) . If cracking also occurs in the MCrAlY bond coat, the bending behavior of the TBC specimen is dictated by the substrate only.
THERMAL STRESS We assume that the substrate layer is thick enough to maintain the composite specimen straight during all manufacturing steps and during temperature exposure without mechanical loading. In this case, the thermal stress at a position y can be calculated as a function of the thermoelastic and geometric parameters of each layer [8]:
E)dj ~
j=l
(2) where F iis the biaxial elastic modulus of each layer i ( i = l ...k with k = l ...3), ai the thermal expansion coefficient and di the thickness. To and T define the temperature before and after heating.
BENDING STRESS Following bending theory, the stress at a position y is proportional to the distance from the neutral axis and inversely proportional to the radius of curvature at the neutral axis R:
Fig. 2:Deflection and curvature of a bending specimen
EXPERIMENTAL 4 mm-thick plates of the Ni-based superalloy IN 792 were coated (IWV1, Forschungszentrum Jiilich) with an industrial NiCoCrAlY variant by vacuum plasma spraying. After a two-step heat treatment (2 h at 112OOC and 24 h at 850°C), a ceramic coating (ZrOz + 7-8 wt % Yz03) was air plasma sprayed onto the substrate / bond coat composite. The thickness of metallic bond coat and ceramic top coat was 100 and 340 pm respectively. Fig. 3 shows a cross-section micrograph of a plasma sprayed TBC system. The morphology of the ceramic layer is typical of plasma sprayed materials, consisting of splats formed during the deposition of the molten particles. Fig. 3 also reveals a high concentration of defects (characteristic pores and microcracks). Image analysis indicated for this material a porosity of about 15%.
with Ei the elastic modulus of the layer i (i= 1 ...k with k = 1 . .. 3 ) and fn the position of the neutral axis [9]. From the equilibrium equation of bending stresses
i=l
ti
the position of the neutral axis can be determined as a function of the elastic and geometric parameters:
A E i d i [ 2 g d j+ d i \ I, =
i=I
\
j=1
I
(5)
2 c Eidi i=l
The radius of curvature at the neutral axis is directly related to the deflection of the beamf (Fig. 2):
Substrale
I
.
-. .
Fig. 3: Cross-section micrograph of a plasma sprayed TBC system
where L is the span length of the four-point bending device. Using this geometric relation, equation 3 becomes:
524
Some relevant thermoelastic data of the three materials are given in Tables 1 and 2 as a function of temperature [lo-131.
Top coat * Bond coat * Substrate [ 101
9.1 + 1 . 8 ~ 1 0T ' ~- 1 . 6 ~ 1 0 -T2 ' ~1 . 7 ~ 1 0 -T3 ~' 12.4 + 3.7~10"T - 2 . 7 ~ 1 0T2 - ~+ 4 . 8 ~ 1 0T3 -~ 9.8 + 1 . 3 ~ 1 0T- ~- 1.9x10-' T2 + 1 . 2 1~O-* T3
Table 1: Temperature dependenceof the thermal expansion coefficient for the three materials. * Thermal expansion coefficient measured with free standing materials
Temperature Topcoat [ l l ] Elastic modulus Bondcoat* (GPa) Substrate * 4-point tensile Top coat [ 121 strength (MPa) .
I
(MPa) . , *
RT 600°C 950°C 46 44 40 122 109 99 136 114 90 35
strength Bond coat [ 131 980
700
30
Table 2: Temperature dependent elastic properties. Elastic modulus measured in bending on thick materials
Specimen strips (47 x 4 mm2) were prepared from the composite material. The thickness of the substrate was mechanically reduced to produce a specimen thickness of 3 mm. Finally one side face was polished using a final 1 pm diamond paste. Isothermal four-point bending tests were performed with the three-layer specimen strips from room temperature to 950°C. The bending tests were controlled by the deflection of the specimen. A total deflection of 1 mm was applied in 30 minutes. For high temperature tests, the specimens were mounted in the four-point device under a 5 N pre-load and heated up with a rate of 8"C/min. The bending tests were started after a twohour holding time at the chosen temperature. All bending strips were positioned in the bending device so that tensile stress developed in the top coat (Fig. 4). Since the bending device was mounted in a furnace with a silica window, in situ observation of the polished side face could be carried out by using a high temperature telescope system.
Specimen
43L2 -4a2) AP (8) 4bd3 Af ' with L = 40 mm the span length, a = 10 mm the distance from the support to the load applicator, b = 4 mm the specimen width and d=3 mm the specimen thickness. E=
RESULTS LOAD/DEFLECTION CURVES Fig. 5 shows the applied loads as a function of deflection at room temperature (RT), 600 and 950°C. At RT and 600"C, the loaddeflection curves can be approximated by straight lines indicating that the composite material behaves elastically during the entire experiment. At 950"C, the elastic behavior is limited to deflections below 300 pm. 1000
,
800
600 400
200 0 0
200
600
400
800
Deflection (pm) Fig. 5: Loaddeflection experimental curves
The equivalent elastic modulus of the composite is proportional to the slope of the straight part of the loaddeflection curve (Equation 8). The obtained moduli are presented in Table 3 as a function of temperature and compared to theoretical values ETh
=
Eldl + E2d2 + E3d3
.
(9)
d
Temperature Slope of the linear part (Npm-') Experimental equivalent modulus (GPa) Theoretical equivalent modulus (GPa)
RT
600°C
950°C
1.180
0.909
0.585
120
93
60
125
106
85
Table 3: Comparison between experimental and theoretical equivalent elastic modulus in bending of the three-layer specimen
Fig. 4:Four-point bending device
Load P and deflection f were recorded and used to calculate the equivalent modulus of elasticity in bending E of the composite material using [ 141:
With increasing temperature, the elastic modulus decreases (Table 2). At RT and 6OO0C, both experimental and theoretical moduli are similar. At 950"C, the theoretical modulus is about 30% higher than the measured one, probably resulting from the larger error in slope determination.
525
IN SITU OBSERVATIONS Fig. 6, 7 and 8 illustrate the crack propagation in the TBC system at room temperature. Cracks were first observed at deflections between 250 and 300 pm (Fig. 6). Bending cracks do not start at the ceramic surface but seem to initiate Dreferablv at
penetrate the BC, e.g., pores located near the interface of TBC and BC play the role of crack stoppers. Cracks enter into the pore and do not cross the interface (Fig. 7b). Additional strains are necessary to induce further propagation.
Fig. 6: Crack initiation at two locations of the YSZ top coat at RT for deflections between 250 and 300 pm
Fig. 7 shows the crack pattern at two locations of the top coat at the end of the bending experiment (f==950 pm).
Fig. 9: In situ observation at 600°C: crack a) initiation at 185 pm deflection and b) propagation at 500 pm deflection in the top coat
Fig. 7: Crack pattern at two locations of the TBC system a1 RT for a 950 pm deflection
Once the interface is reached, cracks grow into the brittle bond coat following a straight trajectory within only a few micrometers of deflection (Fig. 7-a). They stop finally in the tougher substrate material, dissipating the crack driving force into the plastic zone in front of the crack tip (Fig. 8). However, cracks not always
526
Fig. 10: Crack pattern at two locations of the TBC system at 600°C for a 970 p m deflection
In the ceramic top coat, the first crack is observed at a deflection of about 185 pm (Fig. 9-a). As room temperature behavior, the cracks initiate preferably at bigger defects near the top coat surface. Once initiated, the
cracks propagate again towards the interface between top coat and bond coat. At 500 pm deflection, they cover about two thirds of the total ceramic thickness (Fig. 9-b) and reach the interface with the NiCoCrAlY bond coat at 700 pm deflection. However, none of the cracks enters the bond coat layer, even when additional bending strain is applied. The crack stop is also independent from the presence of pores near the interface (Fig. lo). Fig. 10-a and b show two cracks in the ceramic top coat for a 970 pm deflection. At 950°C testing temperature, cracks are also observed in the ceramic layer, without any mechanical loading. Again pores near the surface are the origin (Fig. 11). The formation of cracks probably results from the tensile thermal mismatch stress, which develops in the ceramic layer once the initial compressive residual stress is relaxed. With additional bending strain, the thermal cracks of the YSZ-layer grow again towards the interface with the bond coat and arrest there. Even higher strains do not induce crack propagation. During cooling, the bond coat contracts faster than the top coat so that compressive stresses develop in the ceramic and the crack opening diminishes (Fig. 12).
DISCUSSION The bending stress as a function of deflection in a multilayer specimen is given by equation 7. At room temperature, major cracks are first observed in the ceramic top coat at deflections between 250 and 300 pm. Below 250 pm, the YSZ layer is intact and bending strain and stress can be calculated by using a three-layer model (k = 3). The obtained strain and stress distributions are shown in Fig. 13. Cracks first appear in the YSZ top coating at a bending strain and stress of about EB=0.2 % and o B = 9 0 MPa, respectively. Since the tensile strength of the porous TBC material is about (TC = 35 MPa (room temperature and as sprayed conditions, Table 2, [ 12]), we can deduce that the ceramic top coat is under a relatively low compressive residual stress ~ O = o c- o g = - 55 MPa. Bond coat
0.20
0.00
p T YS? I
Substrate
b
II 1 I
-0.20 250
ii 1
I 1
Fig. 11: In situ observation of the ceramic top coat cracks before bending test a) at RT and b) at 950°C -250
Fig. 13: a) Bending strain and b) bending stress in a three layer specimen for a 250 pm deflection
Fig. 12: Crack in the YSZ top coating a) at 950°C for 1 nun deflection and b) after cooling to RT
Deflections for crack initiation in the ceramic top coat and propagation through the bond coat are resumed in Table 4. Temperature Crack in the top coat
RT 250 pm
Crack in the bond coat 950 pm
600°C
950°C Thermal 185pm cracks Cracks don't enter
Table 4: Deflection at observation of cracks in top coat and bond coat
The bending cracks reach the interface between top and bond coat at deflections of about 950 pm. At this stage of the bending experiment, the ceramic coating is cracked over its entire length (segmentation) and does not influence further the behavior of the composite material. A two-layer model has been adopted to calculate the strain and stress (k = 2). At 950 pm deflection, the bond coat faces bending strains of about 0.6 % and tensile stresses between 7 10 and 765 MPa. These values are much lower than the strength of MCrAlY materials oc=980 MPa (Table 2, [13]), indicating that the bond coat is under a tensile residual stress of about 0 0 ~ 2 1 5 MPa. Independently, equation 2 can be used to estimate the theoretical residual stress distribution in three-layer specimen strips. Considering the thermal and mechanical history of specimen preparation, the following steps have to be considered: - Heat treatment of bond coat / substrate composite 24h at 850°C. We assume that the composite is stress free.
527
Cooling down of the two-layer specimen to RT and heating up to 300°C for the plasma spraying of the ceramic TBC. Spraying of the ceramic top coat: the molten YSZ particles are rapidly cooled down. The high tensile stresses, which should result in the TBC are considered to relax through the formation of microcracks. Thus the tensile strength of the YSZ-TBC provides an upper boundary value for the remaining residual TBC stress. Cooling down of the three-layer system to RT. As shown in Table 4, the residual stresses in the YSZ top coat and the NiCoCrAlY bond coat, which are calculated on these assumptions, are quite similar to the values obtained from the bending experiments.
Bending experiments Thermoelastic calculation
Top coat -55 MPa -63 MPa
Bond coat 215 MPa 200 MPa
Table 5: Residual stresses in top and bond coat
In the case of the high temperature tests, thermal stresses have to be added to the residual stresses to obtain the stress distribution in the TBC system before bending. Fig. 14 shows the respective stress distributions at RT, 600 and 950°C. Bond coat
v -4 200
b
Substrate
-
ii
I1
!! 1
ii
RT -0-
600°C
flection of the specimen reaches 1 mm, which corresponds to bending stresses at the interface with the top coat of 720 and 650 MPa at 600°C and 950°C respectively. In any case, the resulting stresses of 780 and 610 MPa should be sufficient to induce cracking in the metallic bond coat. However, the cracks stop at the interface between top coat and bond coat (Fig. 10 and 12), which proves the high ductility of this metallic layer. If we assume that the BC only fractures in a brittle mode below the ductile-to-brittle transition temperature (DBTT), then 600°C seems to be above this temperature and it is unlikely that the calculated stresses can be maintained. The conducted experiments lead to the conclude that the DBTT of the NiCoCrAlY layer is below 600°C.
CONCLUSION Using in situ observations of cracking in a TBC system during bending, the critical stresses of ceramic top coat and metallic bond coat were determined. Residual stresses of about -55 MPa and 215 MPa in top and bond coat were deduced respectively. These values are in a good agreement with residual stresses predicted from linear elastic calculations. With increasing temperature, tensile stresses develop in the YSZ top coat as a result of thermal expansion mismatch. They are sufficient to induce thermoelastic cracking at 950°C. At high temperature, the cracks do not penetrate the metallic bond coat, reflecting the high ductility of the NiCoCrAlY material. The change from brittle fracture at RT to non-fracture at 600 and 95OOC indicates that the ductile-to-brittle transition temperature (DBTT) of the material is below 600°C. Complementary bending tests at temperature between RT and 600°C would allow a more precise evaluation.
-0- 950°C
ACKNOWLEDGEMENT The authors are grateful to Ms. Funke, IWV1, Forschungszentrum Jiilich, for providing the coated specimens.
Fig. 14: Stress situation in a TBC specimen at RT. 600 and 950°C
With increasing temperature, the compressive residual stress in the top coat relaxes and tensile stress develop. Thus the ceramic layer expands at a lower rate than the two metallic materials (Table 1). At 950"C, the calculated thermal stress ( T ~ ~ MPa ~ ~ ~is =already ~ O significantly above the bending strength of the YSZ layer and cracks are formed thermally before any mechanical loading (Fig. 11). Following the same thermoelastic considerations, nominal bond coat stresses can be derived. In the bond coat, the tensile residual stress relaxes (60 MPa at 600°C) and compressive stress develops (-40 MPa at 950°C). At the end of the bending experiment, the de-
528
REFERENCES 1
2 3
4
W.J. Brindley, Thermal Barrier Coatings of the Future, J. of Thermal Spray Technol., 6[1], (1997) 3-4. S. Stecura, Optimization of the NiCrA1Y/Zr02Yz03 Thermal Barrier System, Adv. Ceram. Mat., I[ 11, (1986) 68-76. P. Fauchais, A. Vardelle and M. Vardelle, Recent Developments in Plasma Sprayed Thermal Barrier Coatings, Workshop on Thermal Barrier Coatings, 85" Meeting of AGARD, Structures and Materials Panel, 13-17 October 1997, Aalborg, Denmark. R.V. Hillery, B.H. Pilsner, R.L. McKnight, T.S. Cook and M.S. Hartle, Thermal Barrier Coating
5
6
7
8
Life Prediction Model Development, NASA Contractor Report 180807, (1988). M.K. Hobbs and H. Reiter, Residual Stress in ZrOz8%Y203 Plasma Sprayed Thermal Barrier Coatings, Thermal Spray: Advances in Coatings Technology, ASM International, (1988) 285-290. P. Bengtsson and C. Persson, Modelled and measured residual stresses in plasma sprayed thermal barrier coatings, Surf. Coat. Technol., 92, (1997) 78-86. S.C. Gill and T.W. Clyne, Investigation of residual stress generation during thermal spraying by continuous curvature measurement, Thin Solid Films, 250, (1994) 172-180. G. Blandin, R.W. Steinbrech and L. Singheiser, Response of Thermal Barrier Composites to Temperature Exposure, EUROMAT'99, 27.-30.9.1999, Munich, Germany.
9 G.H. Ryder, Strength of materials (1970). 10 COST 501, Advanced Blading for Gas Turbines, 3rd Annual Report (1992). 1 1 D. Basu, G. Blandin, L. Singheiser and R.W. Steinbrech, Elastic Behavior of Thermal Coatings: A Comparison of Mechanical and Thermoelastic Tests, to be published. 12 W. Mannsmann, Keramische Warmedammschichtsysteme: Eigenschaften und Verhalten unter mechanischer, thermischer und thermomechanischer Beanspruchung, PhD. Thesis, Karlsruhe University (1995). 13 M.G. Hebsur and R.V. Miner, High Temperature Tensile and Creep Behavior of Low Pressure Plasma-Sprayed Ni-Co-Cr-Al-Y Coating Alloy, Mat. Sci. Eng., 83, (1986) 239-245. 14 P. Dadras, Metals Handbook, 9th Edition, Vol. 8: Mechanical Testing, (1985).
529
This Page Intentionally Left Blank
POROSITY GRADED SILICON CARBIDE EVAPORATOR TUBES FOR GASTURBINES WITH PREMIX BURNERS M. Droschel", R. Oberacker", M. J. Hoffmann", W. Schaller**, D. Munz** (*) Institut fur Keramik im Maschinenbau, Universitat Karlsruhe (TH)
D-76131 Karlsruhe, Germany (**) Institut fur Materialforschung, Forschungszentrum Karlsruhe
D-76021 Karlsruhe, Germany
ABSTRACT Porous silicon carbide (Sic) ceramics are promising materials for liquid fuel evaporator tubes in gas turbine combustors. For evaporator tubes with four different design variations (two-layer concept and three continuous porosity gradients) finite element method calculations were made to determinate the thermal stresses due to the temperature gradient during steady state and transient operation. By comparing the local stress distribution with the local strength it could be shown that a tailored porosity gradient is necessary to meet the local stress/ strength requirements. For tubes with such continuous porosity gradients a new processing route based on pressure filtration has to be developed. The formation of one and more dimensional wax concentration gradients in the filter cake (corresponding to porosity and pore size gradients after sintering) is adjusted by controlling the composition of a mixture of Sic- and Sic-wax slurries and the filtration pressure. The design concept of a laboratory casting apparatus which uses this new concept and an experimental/ numerical method for the derivation of the process parameters is also described in this paper.
INTRODUCTION A concept for a gas turbine combustor with premix burner is shown in Fig. 1. The evaporation of the liquid fuel and the burning zone are decoupled [ 11. The fuel is sprayed on the outer surface of the Sic-evaporator tube.
Due to the flowing air from the compressor, a thin film of fuel is formed on the outer surface of the evaporator, which vaporizes completely. The homogeneous mix of air and fuel vapor enters the reaction zone inside the tube where the combustion takes place. It is known, that a porous evaporator surface allows much higher evaporation rates compared to a dense material. This could result in advantages with respect to the design of the combustion system, if its reliability is maintained despite of the porosity. However, the evaporizer tube has to be gas tight on the inner side and it should be porous at the outer surface, where the evaporation takes place [2]. Between this porous outer surface and the dense inner surface a porosity gradient should be introduced.
POROSITY GRADIENT DESIGN BY FINITE ELEMENT CALCULATIONS GEOMETRICAL BOUNDARY CONDITIONS The evaporator is an axial symmetric component with an inner radius of ri = 22.5 mm. The inner surface consists of a dense S i c layer with 2.0 mm thickness which is followed by a porous Sic-layer, serving as an evaporator (Fig. 2).
radius r
nozzle
porosity P. "/o
Fig. 2: Design ofthe evaporator tube.
Fig. I : Concept o f a gas turbine combustor with premix burner [I].
For the porosity gradient different design variants were studied. According to a previous work [2] four design variants of the evaporator tube fulfill the de-
53 1
mands regarding the fuel evaporation. The first design variation is a homogeneous layer (non-graded) with the outer radius ra = 38.8 mm and a constant porosity of 50 %. Three alternative porosity gradients with a constant porosity of 50 % at the outer radius of r, = 38.8 mm are considered. The investigated three porosity gradients are defined by Eq. 1.
Eq. 1 with: rp 5 r 5 ra, rp= 24.5 mm, r, P(r,)
=0
%, P(r,)
=
= 38.8 mm
50 %
With n = 1.O a linear gradient, with n = 0.5 a convex gradient and with n = 2.0 a concave gradient is obtained. The resulting porosity gradients are shown in Fig. 3 in comparison to the non-graded two layer concept.
mode is reached after 2.2 to 13.0 seconds, depending on the porosity gradient. Then the temperature of the inner wall is 15OO0C, the temperature of the outer wall is 550°C.
FINITE ELEMENT METHOD STRESS CALCULATIONS Transient temperature calculations until steady state and the corresponding stress calculations were made for the different design variations shown in Fig. 3. The maximum stresses through this period appeared in zdirection (0,). Figures 4 - 7 reveal the stresses (T, along the r-axis for steady state conditions and for the time, where the maximum stress occurs during the heating. The stresses are compared with the local strength o0 determined in [2]. The strength (T, is the porosity dependent Weibull stress, which was determined for S i c specimens with homogeneous porosity by ring-on-ring tests. stress, MPa
porosity P. %
strength, MPa
I 3-
400
I400
50
1-200
"%.,
-400
20 I
0
_.--
20
25
,...
,./
,I.'
. -.
...'
. . concaie
-600
n=?
30 35 tube radius. inin Fig. 3: Porosity gradients chosen for comparison by finite elemet method.
40
MATERIAL DATA, THERMAL LOAD ANDSTRESS BOUNDARY CONDITIONS Because of its advantageous features in high temperature application, S i c has been chosen for the present study. The material data (density, Young's modulus, Poisson's ratio, strength, thermal expansion coefficient, thermal conductivity and specific heat) used in finite element calculations are based on data from homogeneous porous S i c specimens, which were experimentally determined in a former study [2]. Due to the independency of the thermal expansion coefficient on porosity, it is assumed €or the stress analysis that the component is stress free after sintering. The component regarded is axially symmetric and has an infinite length in z-direction. The component is allowed to expand freely in all directions. For this component the thermal expansion coefficient is independent of the location. Thermal stresses occur due to the inhomogeneous temperature distribution, which appear both in stationary and transient mode. A component temperature of 150°C in accordance with the temperature of the incoming compressor air that at the beginning of the heating state is assumed. After ignition of the flame in the combustor the inner surface of the evaporator tube is heated. The stationary
532
20
25
30 tuberadius, mm
40
35
Fig. 4: Two-layer component:6,and 0,)versus r-axis.
In the two-layers component shown in Fig. 4, the stresses (T, created under transient conditions in the porous layer are much higher than the strength oo.The component would fail with a very high probability in the beginning of the first heating process. All design variants with a graded porous layer (Fig. 5 - Fig. 7) indicate compressive stresses at the inner and outer surface during heating. Along the r-axis a maximum of tensile stress close to the middle of the porous layer is found. For the convex and the linear porosity gradient the maximum stress exceeds the local strength of the material and the failure probability of the component is very high, similar to a two-layers component. strength, MPa
stress, MPa 400
I
400
--4---
200 0 -200 -400
I
-400
transient state
-600
-600 20
25
30 tube radius, m m
35
Fig. 5: Convex porosity gradient: o,and 0,)versus r-axis.
1
Under steady-state conditions no maximum is found for tensile stresses along the r-axis inside the component. The results of Fig. 7 show that it is possible to keep the local stresses lower than the local strength by using a concave porosity gradient. Only for this geometrical boundary condition the stress 0, does not exceed the local strength o,,. strength, MPa
stress, MPa 400
I
at-/....\ ;
200 0
/
-200
400
-..-.
200
*%_
steady stale
--_. -....__
-200
i
i
-400
0
-400
transient state
-600
-600
With non of these methods a more dimensional concentration gradient could be adjusted. For evaporator tubes with defined porosity gradients in radial and axial direction as shown in Fig. 2 a new processing route has been developed. This process uses continuous pressure filtration of aqueous slurries containing S i c and wax particles. The latter act as pore formers which are burned out of the cake prior to sintering. By controlling and adjusting the local filtration pressures and the wax concentration in the slurry, defined gradients of wax content in the filter cake are realized in and perpendicular to the cake forming direction. The local porosity and pore size in the sintered material correspond to the wax concentration and wax particle size in the filter cake. This process requires a new type of pressure casting apparatus and a numerical process control.
I
20
25
30 tube radius, nini
35
40
Fig. 6: Linear porosity gradient: 0,and (T,, versus r-axis. stress, MPa
strength, MPa I400
400
EQUIPMENT FOR PRESSURE FILTRATION OF TWO-DIMENSIONAL GRADED COMPONENTS To realize axial-symmetrical porosity or pore size gradients in tubes a new pressure filtration apparatus (Fig. 8) has been constructed.
0 -200 -400 -600
; transient state
software controlled gas pmsure (0 I 0 M Pa)
-
-400 -600
Sic-wax slurries with different wax conccntrdtions
The use of a tailored transition fimction for porosity seems to be advantageous for the investigated component. The concave porosity gradient provides the best relation between stresses and local strength at any time. However, manufacturing of evaporator tubes with tailored porosity gradients requires a new processing routes such as pressure filtration.
PRESSURE FILTRATION PROCESS FOR THE PREPARATION OF GRADED COMPONENTS
flow nietcr and throttlc valve for etch tilts
pressure filtration in segniented metal filters rotating cylinder with passing bores
I I , filtrate
Fig. 8: Pressure filtration apparatus for two dimensional concentra-
tion gradients in plates and tubes .
Several papers concerning the production of different hnctionally graded materials by using slip casting technique have been published earlier [3 - 101. One dimensional gradients in cake forming direction were occured by sequential or continuous changing the slurry composition during the casting process. The filtration during the slip casting process was realised by using porous molds [3 - 71 or by vacuum pumps [8 - 101. Gradient formation in the filter cake was only influenced by changing the slurry composition, but not by a variation of the filtration pressure during the process.
It is designed for processing of cylindrical tubes (outer diameter 65 mm, thickness 15 mm and height 55 mm). The bottom of its pressure chamber can be replaced for the fabrication of parts with other geometries, e.g. plates. A maximum gas pressure of 10 MPa can be applied. The Sic- and Sic-wax slurries are stored and mixed inside the pressure chamber. The wax concentration of the Sic-wax slurry at the cakesuspension interface can be continuously changed. To manipulate the local cake thickness, segmented metal
533
filters are used. The filtrate volume of each metal filter, corresponding to the local cake thickness, is measured by flow meters and controlled by throttle valves. Tube shaped components with compositional gradients in radial and in axial direction can be obtained by adjusting the chamber pressure, the wax concentration at the filtration front and the local flow resistance in accordance to the desired local wax concentration in the tube at any given time of the process. However, there is no general solution for the correlation between the process variables. Therefore, a numerical simulation model has been developed for process control.
MODELLING OF THE CAKE FILTRATION PROCESS FOR HOMOGENEOUS SLIPS The growth of a filter cake during pressure filtration is schematically shown in Fig. 9.
ent experiments. The slurries are electrostatically stabilized, their viscosity is about 10 mPas at a pH of 7.5. The S i c particles have a mean diameter of d50 = 0.8 pm, the waxes have a d50 of 150 pm. For pressure filtration of homogeneous Sic-wax slumes an apparatus was used, as shown in Fig. 10. Constant filtration pressures of 0.5 to 4 MPa were chosen for the experiments. The filter cakes had a diameter of 60 mm and a thickness of 10 mm. The filtrate volume and the filtrate volume flow were measured by a balance. For pressure filtration with constant pressure and a constant solids loading the permeability D of the filter cake and the flow resistance R of the metal filter can be derived by plotting (t/V) versus the filtrate volume V [ 121. This plot results in a linear dependency according to Eq. 3. The permeability D and the flow resistance R can then be calculated from this linear fit of the experimental data by Eq. 4 and Eq. 5. gas pressure
m
h
j
,pressure
chamber
ltercake
filtrate I
filtrat
metal filter
Fig. 9: Schematic diagram of the cake filtration process.
The cake forming rate during pressure filtration is described [ 1 I] by the basic equation of cake filtration:
-dh_ dt
-
K
. Ap
Eq. 2
q.(k+R)
Cake growth can be controlled through the process parameters time t and pressure Ap. The solids loading c and the filtrate viscosity q are known variables, which are determined by the slurry. The flow resistance R of the metal filter can be measured easily in a calibration experiment. The porosity E and the permeability D depend on the microstructure of the filter cake which itself is determined by the filtration process. The filtrate volume V in Fig. 9 is related to the cake thickness h and the cross section of the metal filter A (V=(h.A)k). The parameter K is a constant determined by the solids loading c and the porosity E ( ~ = c / ( l - ~ - c )The ) . cake structure parameters can be derived from experiments with homogeneous suspensions [ 121. Knowing all parameters, Eq. 2 can be integrated numerically.
EXPERIMENTAL DETERMINATION OF CAKE STRUCTURE PARAMETERS Aqueous S i c and Sic-wax slurries with a constant solids loading of c = 32.5 vol% were used for the pres-
534
\ metal filter
balance
Fig. 10: Laboratory casting equipment for pressure filtration of homogeneous cakes.
-t= a + b . V V
D=
R=
q*K 2 . b .A .Ap
a.A.Ap
Eq. 3 Eq. 4
Eq. 5
rl
The concentration constant K is determined by E and c as already explained. In the investigated range of filtration pressures (0.5 to 4 MPa), the cake porosity E and permeability D were practically pressure independent. This is an indication for incompressible cake structures. The Sic/ wax ratio, however, has significant influence for these parameters as shown in Fig. 11. The cake porosity E decreases up to a wax concentration of 40 ~01%.This can be explained by the increased packing eficiency of the bimodal particle mixtures. The permeability D increases by a factor of 4 (from 1.10-'' m2 to 4.10-'' m2) when the wax volume fraction of the solids reaches 60 vol%, despite of the reduction in porosity. This is an indication for increased pore channel diameters of the cakes with bimodal particles.
14
35t
:----'
--- -------
0
10
20
30
40
50 wax concentration in solids, vol%
lo
60
70
Fig. I I : Dependence of cake structure parameters on the Sic/ wax- ratio.
The strong influence of the wax content on the cake structure parameters E and D can not be neglected in the solution procedure for Eq. 2 and makes numerical integration necessary.
gradient in the cake can be simulated. The solution provides either the course of the filtration pressure for a given function of the cake forming rate, or the course of the cake forming rate for a given function of the filtration pressure. With respect to a low level of casting defects, casting at constant cake forming rates seems to be the better way, which was therefore chosen for the subsequent calculations. In the process calculations, the porosity fimctions investigated in the finite element calculations and additionally porosity fimctions with exponents of n = 0.1, 0.25, 4 and 10 (see Eq. 1) were investigated. All these porosity gradients are compared inFig. 13. porosity P, % 60 two-laver ioint
50 40
30
MODELLING OF PRESSURE CASTING FOR ONE DIMENSIONAL GRADIENTS One dimensional concentration gradients in the filter cake are formed by changing the wax concentration in the solid phase during pressure filtration as shown in Fig. 12. 100% 0 % (3
X %
100-x %
0%
100 Yo
Fig. 12: Processing of one dimensional wax concentration gradients.
This filtration process is regarded for short time increments Ati. For this purpose, Eq. 2 is modified and solved in an incremental form as shown in Eq. 6. As shown before, the cake porosity E and thus the concentration constant K depends on the wax content of the slurries. During filtration the total flow resistance Rges, i (which contains the flow resistances of the metal filter and that of the previously formed filter cake) increases according to Eq. 7.
20 10 0
20
25
30 tube radius, mm
40
35
Fig. 13: Porosity gradients chosen for simulation of filtration.
For the porosity functions with one dimensional porosity gradients the pressure filtration process was simulated with the experimentally derived cake structure parameters shown in Fig. 11. The results for a constant cake forming rate of dh/dt = 0.1 mm/s are presented in Fig. 14 and Fig. 15. Due to the shrinkage during drying and sintering, the wet cake has a thickness of h = 20.9 mm. From cake thickness and constant cake forming rate the total filtration time can be calculated to 209 seconds. wax concentration in slurry, vol% 70
1
two-laver ioint
-___-_.__-.
'
0
50
I00 I50 filtration time, s
I
'
250
200
Fig. 14: Simulation of the filtration process for the porosity gradients of Fig. 13: wax concentration in slurry.
Eq. 6
Eq. 7
By integration of Eq. 6 and Eq. 7 the filtration process for any desired one dimensional wax concentration
The course of the wax concentration in the slurry shows a linear dependency on the desired gradient fbnction (Fig. 14). This more or less trivial finding is valid for the special case of constant cake forming rate. Rather complicated functions are calculated for the course of the filtration pressure (Fig. 15). The maximum filtration pressure Apma has to be applied at the end of the filtration process (t = 209 s). The kinetic of filtration pressure depends strongly on the exponent n
535
filtration pressure, MPa
REFERENCES
6
5 4
3
2 I
0
- .- joint
two-layer
0
50
I00 I50 filtration time, s
200
250
Fig. 15: Simulation ofthe filtration process for the porosity gradients of Fig. 13: filtration pressure.
in Eq. 1. The two-layer joint (n + 0) requires by far the lowest maximum filtration pressure (Apmm= 1.37 MPa), due to its high amount of cross section with high permeability. With increasing exponents n > 0 the value of the maximum filtration pressure is raised up to a maximum of Apmax = 5.57 MPa for n = 1.25. For exponents n > 1.25 the maximum filtration pressure is reduced again and reaches a value of Apmar= 3.61 MPa for n -+ co. The kinetic of pressure filtration cannot be understood only on the basis of permeability D, as it depends strongly on the concentration parameter E. The results of the process simulations shown in Fig. 14 - 15 apply not only for the experimentally derived filtration parameters (Fig. 11). The method can also be used for any other binary slurry mixture if their cake structure parameters D and E are known. By using SiCwax slurry mixtures with different wax concentrations and/ or waxes with different particle sizes defined gradients in the pore structure (porosity a n d or pore size) of the sintered S i c evaporator tube can be adjusted.
SUMMARY Silicon carbide evaporator tubes with porosity gradients are promising materials for liquid he1 in gas turbine combustors. Finite element method calculations show that for these devices a tailored porosity gradient is necessary to meet the local stress/ strength requirements. To realize such components a new processing route based on the continuous pressure filtration has been developed. With this process it is possible to produce defined one- and more dimensional concentration gradients of a second particle type in a filter cake. The filtration process parameters can be derived by a numerical integration of the basic filtration equation. The required cake structure parameters can easily be measured in constant pressure experiments with homogeneous filter cakes.
ACKNOWLEDGEMENT The financial support of the Deutsche Forschungsgemeinschaft (Ob 104/6, Mu 466/26) is gratehlly acknowledged.
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M. Brandauer, A. Schulz, S. Wittig, Optimization of fuel prevaporization on porous ceramic surfaces for low NO,-combustion, Combustion technology for a clean environment, Lisboa, Portugal, (1995) M. Droschel, Grundlegende Untersuchungen zur Eignung poroser Keramiken als Verdampferbauteile, doctoral thesis, University of Karlsruhe, IKM 022, ISSN 1436-3488, (1998) J. Requena, R. Moreno, J. S. Moya, Alumina and AlumindZirconia Multilayer Composites Obtained by Slip Casting, Journal of the American Ceramic Society, 72 [8], pp. 1511-1513, (1989) J. S. Moya, A. J. Sanchez-Herencia, J. Requena, R. Moreno, Functionally gradient ceramics by sequential slip casting, Materials Letter 14 [5, 61, pp. 333-335, (1992) J. Chu, H. Ishibashi, K. Hayashi, H. Takebe, K. Morinaga, Slip Casting of Continuous Functionally Gradient Material, Journal of the Japanese Ceramic Society, 101 [7], pp. 84 1-844, ( 1993) B. R. Marple, J. Boulanger, Graded Casting of Materials with Continuous Gradients, Journal of the American Ceramic Society, 77 [lo], pp. 2747-2750, ( 1994) B. R. Marple, J. Boulanger, Slip Casting Process and Apparatus for Producing Graded Materials, United States Patent, No. 5498383, (1994) H. Mori, Y. Sakurai, M. Nakamura, S. Toyama, Formation of Gradient Composites Using the Filtration Mechanism of Binary Particulate Mixtures, Proceedings of 3rd International Symposium on Functionally Graded Materials, FGM-3, Lausanne, Switzerland, 1994, ISBN 288074-290-0, pp. 173-178, ( 1995) K. Taka, Y. Murakami, T. Ishikura, N. Hayashi, S. Watanabe, Y. Uchida, S. Higa, T. Imura, D. Dykes, Development of stainless Steel / PSZ hnctionally graded materials by means of an expression operation, Proceedings of 4th International Symposium on Functionally Graded Materials, FGM-4, Tsukuba, Japan, 1996, ISBN 0-444-82548-7, pp. 343-348, (1997) S. Watanabe, T. Ishikura, A. Tokumura, Y. Kim, N. Hayashi, Y. Uchida,’S. Higa, D. Dykes, G. Touchard, The Use of a Functionally Graded Material in the Manufacture of a Graded Permittivity Element, Proceedings of 4th International Symposium on Functionally Graded Materials, FGM-4, Tsukuba, Japan, 1996, ISBN 0444-82548-7, pp. 373-378, (1997) H. Gasper, Handbuch der industriellen F e d Flussig-Filtration, ISBN 3-7785- 1784-8, ( 1990) Verein Deutscher Ingenieure, Filtrierbarkeit von Suspensionen: Bestimmung des Filterkuchenwiderstandes, VDI-Richtlinie 2762, (1997)
POROUS CERAMICS FUNCTIONAL CAVITIES FOR SYSTEM INNOVATION Horst R. Maier Rheinisch-Westfalische Technische Hochschule (RWTH), Aachen, Germany
Abstract This IKKM-topic is concerned with the basics and processing of closed and open porous ceramics leading to monolythic structures with application oriented ,,functional cavities". The computer-aided modelling approach includes material tayloring, multimaterialoriented design and joining, quality assured processing as well as application-oriented prototype testing. Three running developments are presented: -Diesel soot filters -Exhaust sound absorbers and -Thin steel slab-casting. All three development examples are supported by patent application. In addition, an outlook is given on started activities concerning biomedical bone replacement and biotechnological immobilization of microorganisms.
1. Introduction Our development of monolythic ceramics with functional cavities is governed by the leading idea to combine aspects of quality, ecology and economy in production as well as during application at the earliest possible stage. In our opinion only the wedding of most suitable materials and processes will lead to a successful operation of functional units and systems. The history is rich in settled applications based on ceramic structures with closed and/or open pores, which we call ,,functional cavities". Porous building materials are of advantage as far as weight, thermal conductivity, sound insulation and room climate are concerned. Homogeniously distributed fine closed pores in Alumina and Chromoxid ligthweight bricks are used in contact with aggressive melts of metals and glasses in order to cope with heat losses, corrosion effects and thermally induced stresses. Open porous coarse foamstructures made of Alumina and Silicon Carbide have been successfully applied in the filtering of molten metals for about 25 years. Hot gas filtering of coal combustion atmosphere at about 850°C based on SiC-tubes has reached an acceptable standard and continue the course to nanostructures as membranes for different applications. Open porous honeycomb structures made of Cordierit have been in mass production for over ten years as automotive and industrial catalyst carriers. In the field of thermal insulation, sound attenuation and gas filtering at high temperatures ceramic fiber structures still play a dominant role, but they are confronted with a risk on health due to fibres lost and whiskers created in service. Looking ahead, there are challenging application profiles in sight with prosperous market potential.
2. Diesel Soot Filters Based on Direct Electrical Regeneration The basic development started with a multiclient project of the european automotive industry (ERATS, 19911994). As a result Silicon Carbide has been evaluated as most favourable in comparison to Cordierit, Mullite and Glass Ceramics. The work continued in cooperation with Thomas Josef Heimbach GmbH, Diiren, in public sponcered projects (1 990- 1997) and led to the system specified below.
2.1 Concept Approach Targeting on stationary and quasistationary diesel engines with accessable electrical power supply, emphasis has been put on direct electrical regeneration (without additives at about 600 "C) of honeycomb elements or electrically connected modules in a subsequent way in order to achieve a narrowed cycling in backpressure. As a result of the ERATS-Project the wall thickness of the honeycoumb structure has been set to the range between 0.75 and 1.5 mm, which gives advantage to thermal shock (increased thermal capacity), to processing (quality and yield) as well as to application (reliability). The disadvantage of packing density (ratio of filtering surface per element voulume) in comparison to market standards had to be compensated by the positive effect of subsequent regeneration and by taylored pore size distribution.
2.2 Processing Features The most suitable shaping process for honeycomb structures is extrusion. Ecological reasons led to a watersoluble raw material mix. In order to reduce wear during extrusion the boarder line case without primary SiCpowder (Si and C is included plus organic binder) in the range of 10 to 70 pm. The organic binder is converted into C-raw material by a coking process in an inert-gas atmosphere. Reaction bonding offers the advantage of zero-shrinkage and of stearing the electrical conductivity in a wide range with support of suitable accitives, e.g. Boron with a particle size not greater than 10 pm and an amount of 0.05 % to 1.O % by weight. The final microstructure is settled in a combined reactiodrecristallizationheat treatment. Reaction to Sic in nitrogen or nitrogenous inert gas is performed between 1400 and 1900 "C and followed by recristallization between 2 100°C and 2300°C. The Boron additives influence the microstructure as well, e.g. pore size distribution. More details are given in US patent number 6,017,473 Ill.
2.3 Material and System Characterization The standard cross sections of the extruded honeycomb filter elements are 42 cpsi (cells per square inch) and 11
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cpsi. The cells are reciprocally closed, force the raw gas to run through the filtering walls and clean gas is leaves the filter element. The microstructure in Fig. 1 illustrates a high level of uniformity. The sharp setting of the median pore diameter to 70 pm in Fig. 3 corresponds to the wall thickness of about 1.5 mm and the allowable backpressure.
unloaded honeycomb element versus the gas flow rate is characterized in Fig. 3.
Fig. 4: Filter unit with 25 filter elements of type 11 cpsi
Fig. 1: SEM microstructure of honeycomb standard Porosity level P,=60% 121 0.25 0.20 0.15 0.10 0.05 0.00 1000
100
10
1
pore size (vm)
Fig. 2: Pore size distribution of honeycomb standard, Porosity level P,=60% /2/ A p (mbar)
The electrical connection for direct electrical regeneration is placed on both ends of each filter element. It consisted in the first stage of metallic flats brazed along the filter element. A further improvement concerning longterm reliability has been achieved with ceramic bonded ceramic endcaps with force fitted wire connectors. A serial connection of certain filter elements to a filter module is also possible in order to optimize the regeneration cycle and to reduce the overall wiring effort. The electrical units (elements or modules) are electrically insulated by suitable interlayers. A filter unit with 25 filter elements of type 11 cpsi , illustrated in Fig.4, contains a filtration area of about 2 m2 related to diesel engines in the range of 50 to 75 kW and depending on soot load and allowable backpressure. The filtration time is estimated with about 2.5 h related to a regeneration time of about 4 min. at a power level of 1.5 kW for each element. That corresponds with about 1.5 to 2 % of the nominal power output. The specific filter system volume (without housing connectors) ranges between 0.23 and 0.34 IkW. Further information is given in 121.
120
2.4 Conclusion
100
The principle of direct electrical regeneration has been proven by field tests in power heat couplings, fork lifts and buses. There is still room for improvements Concerning pressure drop during filtration, soot storage capacity per filter volume, filtering-regeneration cycle as well as regeneration power and system volume (weight) related to nominal power output. This will be done in parallel with setting up a production line. The material and processing basis developed is not limited to diesel soot filtering or to direct electrical regeneration and opens the way to new applications.
80 60 40 20
0 0
40
80 120 gas flow rate (d/ h)
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Fig. 3: Pressure drop of a single unloaded filter element, (1 1cpsi, A=0,08m2,T=2OoC) The specific electrical resistance is usually between 3 ohm*cm at room temperature and shows a NTC behaviour. With a linear coefficient of expansion of 3.3*10" /K (20...600"C), a thermal conductivity of approx. 5 W/(mK) and a maximum operating temperature of 2000°C (depending on atmosphere), the honeycomb standard is most suitable in comparison with common material alternatives. The pressure drop of an
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3. Exhaust Sound Absorbers for Automotive and Aircraft Engines The development started in cooperation with Deutsche Basalt Steinwolle, Bovenden, and Arvin Cheswick, Roermond (1994 - 1997) with the motivation to replace rock wool fibres by monolithic ceramics for automotive
exhaust silencers at compariable (lower) system cost and weight with favourable sound attenuation. This standard has been improved by a new modelling approach and applied to a BMBF-project for aircraft engines.
and is allowing a lightweight sandwich design between metal canning and ceramic modules.
3.1 Concept Approach Following the strict cost, weight and sound attenuation guidelines first of all a cheap material base in link with a costeffective processing route with high flexibility concerning porosity level in the range of 60 to 85 % and pore size distribution in the range of 0.5 to 5 mm had to be defined. The combined effects of microstructure, module shape, arrangement and permeability on exhaust flow,, backpressure and sound attenuation have been studied by benchtests and computer aided modelling. In addition reliability considerations have been done including joining, mechanical, thermal and chemical loads.
. .
-
.
Granulate
25 --
vdume fraction pi]
/A\
3.2 Processing Features In order to avoid environmental pollution organic preforms and fillers have been replaced by inorganic pore formers with extreme high porosity of up to 95 %. After melting in a suitable ceramic matrix the pore former distribution turns into a pore distribution of slightly reduced size. Thus, a high flexibility is gained concerning the level of porosity and pore size distribution. A favourable combination as well in terms of raw material costs (below 3 ECUkg in small batches), is detailed in 131. Perlite as a pore former begins softening at 1190°C and is compatible with a Ti02 matrix build during subsequent sintering in the range of 1450°C to 1600°C. The shaping process is similar to that of the slip casting technique. The water soluble raw material mix contains 0.3 % organic dispersal only, and the highly thixotropical mass is casted in moulds without open porosity to raw shaps in a wide size ranges (bars cylinders, tubes, flats, foils). Green, white and final machining gives a high flexibility to any demand. In total we call this approach internally ,,oecopor-ceramic". At high porosity level (above 50 %) the increasing glass content tends to build closed pores and creates plastic deformation during sintering. These effects can be reduced, if neccessary, by replacing a part of the inorganic pore formers by natural pore formers, that are burned out during firing. Details are led down in patent application 131.
dB 125
120
3.3 Material and System Characterization The microstructure of the inorganic pore former perlit and the resulting oecopor-Ti02 matrix is shown in Fig. 5. There are no new phases formed between them. Due to large pores, necessary for sound attenuation, Hgintrusion porosimetry is not possible and the open pore size distribution in Fig. 6 is estimated from size distribution of perlit granules and image analysis of microstructure in terms of cross section fraction. The total porosity level of 75 % splits into 41 % open porosity and 34 % closed porosity. The high proportion of closed porosity caused by the glassy phase is inactive to sound attenuation but adds to the mechanical strength
115 1000
ston wool pack backpr %urn 22 kPa
2000
3000
4000
5000
6000 rpm
Fig. 7: Front silencer of Ford Mondeo (66 kW) with stone wool pack and IKKM oecopor-Ti02 Although cost and weight guidelines have been met, the principle has not been applied yet due to increase of backpressure and change of pipe arrangement. The knowledge gained with automotive silencers has been transferred to an aircraft engine Lycoming 0-360 of 134 kW and a volume of 5.9 1. The end muffler remained
539
packed with steel wool and the front muffler bas been changed according to Fig. 8. In addition to three IKKh4-oecopor-Ti02 insert the chamber and flow arrangement has been changed as well. The improvements shown in Fig. 9 are verified in bench tests and include a considerable reduction of sound level and backpressure from 120 to 80 mbar. Another version of the front silencer has been successfully tested in a 50 h flight profile without any signs of degradation. Further information is given in (4).
casting by reducing the thickness of the slabs and decreasing the number of rolls. The crucial ceramic components for the continuous near net shape casting guiding and sealing the steel between tundish and mould - are the so called submerged entry nozzles (SEN) presently based on graphite - alumina -zirconia materials. IKKM is bilaterally involved since 1994 concerning - computer aided design of the geometries and the thermal insulation of traditional graphite nozzles, - new joining techniques for hybrid nozzles and - simulation of the corrosion attack. The gained knowledge has led to a new material and processing approach based on pure zirconia materials.
4.1 Concept Approach
Fig. 8: Front silencer in reference layout (a) and with oecopor-Ti02 ceramic modules (b)
105 100
95 90 85 80 75
The submerged entry nozzles have to - present excellent corrosion resistance against steeuslag attack, - provide excellent erosion resistance inside against steel, - prevent clogging effects and gas penetration and - survive thermal shock at casting start. The graphite based materials provide a sufficient service life due to corrosion attack only by a wall- thickness of about 20 mm. This is of disadvantage for thin slab casting and quality of steel. Low porous (total porosity < 5 Vol.%) fine grained zirconia materials - well known from the advanced ceramics - with less than 0.5 wt% Si02 provide an excellent corrosion and erosion resistance against slag/steel attacks and secure the purity and quality of the casted steel. Unfortunately these materials suffer under thermal shock attack. The main task is to improve the thermal shock behaviour with dispersed phases and microcrack patterns in the zirconia matrix and keep the corrosion rate at least at the same level as of pure zirconia.
4.2 Processing features
Fig. 9: Front silencer of aircraft engine (134 kw) with reference (a) and IKKM oecopor-Ti02 (b)
3.4 Conclusion It has been verified that a structure like oecopor-Ti02 is attractive Concerning cost level, is effective for sound attenuation, is suitable for state of the art joining and is stable in short and medium time service. It is recommended to keep the total porosity below about 75 '3'0. There is still room for improvement concerning level of open porosity, pore size distribution and especially for computer aided tayloring of module shape, size and overall arrangement.
4. Thin Steel Slab-Casting with Advanced Ceramic Components A new generation of continuous casting developed by SMS-Demag is based on the principle of thin slab
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The slip casting technique has been selected as the most suitable processing in order to disperse additives and phases homogeneously and manufacture thin walled components. Electrofbsed Mg-PSZ with 3.5 wt.% MgO and a grain size range of 0 to 12 pm has been used. Without additives the structure shows an open porosity level of 16% (density 4.8 g/cm3) after sintering at 1600°C for 2 h. The best improvement so far has been achieved with additions of 1 wt.% A1203(median grain size 1.0 pm) and TiOz (median grain size 0.2 pm). During sintering the Ti02 incorporates in the neighbourhood of the zirconia lattice, weakens the zirconia cell due to repulsion between the dopant cations and the vacancies; a part of the MgO stabilising agent removes from the zirconia cell and reacts with the A1203 to form MgA1204. The in-situ formation of MgAlz04 spinel creates a microcrack pattern due to a volume expansion of appx. 5%. In addition due to the destabilisation of the zirconia (loss of the MgO stabilising agent) a martensitic phase transformation (tetragonal to monoclinic with approx. 5 Vol.% expansion) takes place followed by a second microcrack pattern. By these means the open porosity level increases
slightly, the permeability remains stable, but the thermal shock resistance is improved considerably. Also the corrosion resistance is improved due to the higher monoclinic amount. This promising microstructure has been transferred into complex nozzle shapes of 350 mm height, 350 mm width and 7.5 mm wall thickness. These sensitive components required specially developed supporting techniques during slip casting, drying, handling and sintering in order to keep the shape, tolerances and quality within acceptable limits. More information about the effect of in-situ formed and externally added spinel phases on thermal shock improvement are given in the patent application 151.
mostly disappear above the transformation temperature of monoclinic to tetragonal. 4I . ”Q
,
I
“1.6-
E 5 1.4-
-c 1.2e 18- 0.8.
i06. 0.40.2 -
4.3 Material and system characterisation
0 $0
250
450
650
850
1050
1250
1450
temperature [‘C]
The microstructure of Mg-PSZ with in situ formed spinel (black phase) is shown in Fig.10. The proportion of monoclinic/tetragonaVcubic phases of zirconia is measured with 55/8/37 Vol.%. The monoclinic phase is represented by the long sharp twins and the cubic one by the grains. The interlinked microcracks due to zirconia destabilisation and spinel formation increase the basic open porosity of 16% only by 2%. The corresponding open pore size distribution is given in Fig. 11. The level of closed porosity remains stable at about 3%.
Fig.13: Thermal expansion of Mg-PSZ + spinel The main properties are compared in Table 1 with outstanding results for the created Mg-PSZ+spinel material: The thermal shock resistance - reflected by the remaining 4-point bending strength after water spalling test from 600°C - to RT has been improved from 28 to 72 MPa. With tribute to the high monoclinic amount of 55 Vol.% the corrosion rate has been further reduced from 0.6 m m h to 0.3 mmh. This means, nozzles made of Mg-PSZ+spinel with a wall thickness of 7.5 mm show a higher durability than nozzles made of conventional graphite-alumina-zirconia materials with a wall thickness of 20 mm and a corrosion rate of 4 mm/h. Fig.14 shows the endpiece of a hybrid submerged entry nozzle and demonstrates that the new material standard can be transferred into large complex and thin walled shapes by slip casting. Further informations are given in 161.
I
Table 1: Mechanical, thermal and chemical properties Properties Mg-PSZ+ Mg-PSZ
I
I
I
1
Fig. 11: Microstructure of Mg-PSZ+spinel 0.016
:\ Mg-PSZ, reference
“M 0.014
-
0.012
‘3
0.010
L
E 8
i **: i55OoC in stedacid-slagbath 100
10
1 0.1 0.01 open pore diameter [pm]
Fig. 12: Pore size distribution of Mg-PSZ + spinel The effect of the microcrack patterns is also reflected in the thermal expansion characteristic in Fig.13. The level is clearly reduced but the hysteresis widens and has to be taken into account during temperature cycling applications. The microcracks due to destabilisation
Fig. 14: Submerged entry nozzle, endpiece based on Mg-PSZ+spinel 541
4.4 Conclusion The material and processing approach presented offers thin walled corrosion resistant structures of complex shapes that give vital support to the system development of thin steel slab casting. The considerably improved thermal shock resistance allows preheating techniques presently used in the steel industry. In order to combine the achievements with proven materials the development of hybrid designs and corresponding joining techniques are of outstanding importance. The principle results can be transferred to other applications in steel or non steel industry by carefidly computer aided matching of application temperature profiles with thermal expansion hysteresis adjustments.
5. Started Activities and Outlook Each of the given examples of functional cavity applications shows ist own behaviour concerning material and processing selection, level of porosity, proportion of closed and open pores, size distribution of pores that needs to be grasped as an interlinked system in terms of computer aided modelling and simulation /7/. New activities have been started in the field of immobilization of microorganisms and bone replacement. Fig. 15 shows microorganism of about 1 pm settled onto ceramic cavities. They are used to generate a preproduct of vitamin C. At a porosity level of about 50 % as suitable pore size in the range of 150 pm has been estimated by analytical simulation of the immobilization process.
0.1
1
10
100
1000
PoresizeI)lml
Fig. 17: Spectrum of running and future activities for functional cavities with ceramics
6. Acknowledgements Thanks are expressed for the fruitful cooperation with the industrial partners, Thomas Joseph Heimbach GmbH, Arwin Chesswick and SMS-Demag as well as with the Institut fiir Luft- und Raumfahrt, RWTH Aachen. Respect is payed to the public fundings of the MWMT Nordrhein-Westfalen and the Deutsche Bundesstihg Umwelt.
7. References /1/ Maier, H.R. et al. "Method for Producing porous Molded Body." United State Patent No.: 6,017,473; Jan25,2000 /2/ Maier, H.R.; Best, W.; Schuhmacher, U.;Schafer,W. "Charakteristicsand Design of Diesel Soot Filters Based on Direct Electrical Regeneration" in: CFV Ber. DKG (1998), Nr. 5, S. 25-29 /3/ Pfaff, E;Aneziris, C.; Maier, R.M. "Verfahren zur Herstellung poroser keramischer Strukturen", Offenlegungsschrift DE 19605149A1 Datum: 14.08.97 /4/ Drobietz, R. ;Krusch, C. ;Neuwerth, G. ; Aneziris,C.; Jacob, D. ;Maier, H.R.: "Keramische Schalldtimpferstrukturen fiir Pkw- und Flugmotoren". In: Werkstoff und Automobilantrieb. Diisseldorf 1999, S. 155-172. (VDIBericht 1472) ISBN 3-18-091472-6
Fig. 15: Immobilization of coryne bacterium glutamicum Fig. 16 illustrates a bone replacement structure made of oecopore - Ti02 - ceramics copying the gradual pore size distribution of natural bone. At the ingrowth interface a pore size range of 150 to 400 pm at a porosity level of about 35 % seems to be favourable.
Fig 16: Dummy structure of bone replacement The spectrum of running and future activities is summerized in Fig. 17. It only presents the tip of the iceberg of the fascinating innovative potential of hctional cavities with ceramics. 542
151 Maier, H.R.; Aneziris, C.; Pfaff, E. "Verfahren zur Herstellung eines Keramikwerkstoffs auf Basis Zirkonoxid" Patentanmeldung DE 19938752.4, Datum: 16.08.99 /6/ Maier, H.R.; Aneziris, C.; Pfaff, E.
"MgO partial stabilisized zirconia materials with titanium, alumina or spinel additives for the near net shape processing" Journal of the European Ceramic Society, accepted and to be printed, Sept. 2000 /7/ Maier, H.R. "Porose Keramik, der Natur abgeschaute Funktionshohlraumeals Anreiz fiir neue Anwendungen" In: Horizonte: die RWTH auf dem Wege ins 21. Jahrhundert / Roland Walter (Hrsg.). Berlin 1999 : Springer, S. 127-136. ISBN 3-540-66373-8
FEASIBILITY STUDIES ON APPLYING IN-SITU SINGLE CRYSTAL OXIDE CERAMIC EUTECTIC COMPOSITES IN NONCOOLED HIGH EFFICIENCY TURBINE SYSTEM K. Hirano, T. Suzuki, A. Sasamoto Mechanical Engineering Laboratory, Agency of Industrial Science and Technology, MITI Namiki 1-2 Tsukuba-shi, Ibaraki-ken 305-8564, Japan
ABSTRACT MGC materials, which are in-situ single crystal oxide ceramic eutectic composites have many potentialities, such as high strength, high creep resistance, and high oxidation resistance at ultra-high temperature. In the New Sunshine Program of the Agency of Industrial Science and Technology, a leading research and development has started since FY1998 for three years to apply MGC materials as ultra-high temperature structural materials for gas turbine systems. In this paper, an outline of the leading research and development is given, and the results established by the Mechanical Engineering Laboratory are introduced.
INTRODUCTION Recently, MGC(Melt Growth Composite) materials have newly been researched and developed[l]. MGC materials are in-situ single crystal oxide ceramic eutectic as A1,O3N,A1,O,,(YAG), composites such Al,O,/GdAIO, (GAP)and A1,03/Er,AI,0,,(EAG). MGC materials are expected to be one of the most interesting and attractive ultra-high temperature structural materials in the field of power generator industry, aeronautics, and aerospace, because they have many potentialities that they keep high strength, high creep resistance, and high oxidation resistance at ultra-high temperature. In addition they can deform at ultra-high temperature and they have possibility to fabricate complex shape of components. However, they do not have enough fracture toughness and thermal shock resistance for structural materials, because they are made of oxideloxide ceramics. Also there have been few researches of fatigue strength and fatigue crack growth resistance. Then in order to apply MGC materials to ultra-high temperature structural materials, it is necessary to improve MGC materials in these points. The New Sunshine Program of Agency of Industrial Science and Technology has started a leading research and development for “MGC Ultra High-Eficiency Turbine System” aiming to utilize MGC materials for ultra-high temperature structural materials for a gas turbine system. In this paper, an outline of the leading research and development is briefly given and the current research activities by the Mechanical Engineering Laboratory are introduced.
OUTLINE OF LEADING RESEARCH AND DEVELOPMENT The leading research and develcpment for “MGC Ultra High-Eficiency Turbine System” is one of the New Sunshine Programs(Energy Saving Technology and Development Programs) under the Agency of Industrial Science and Technology. The period qf the leading research and development is for ,hree years from FY1998 and the following items are investigated. (1) Fracture mechanism and improvements For materials performance in severe environments (2) Possibility of low-cost process technology for large complex, near-net shaped components (3) Aero-mech. design methodology for turbine components based on computational fluid dynamics (4) Turbine cycle analysis and system integrated technologies Under the Agency of Industrial Science and Technology, the leading research and development is directly administered by NEDO(New Energy and Industrial Technology Development Organization) with Technology Research Association of Gas Turbine for Practical Ability Progress, Ishikawajima-Harima Heavy Industries Co. Ltd., Kawasaki-Heavy Industries, Ltd., UBE Industries, Ltd., and Japan Ultra-high Temperature Materials Research Institute. As a national laboratory in the Agency of Industrial Science and Technology, the Mechanical Engineering Laboratory joins the leading research and development, and leads and supports it for its own knowledge base. The Mechanical Engineering Laboratory also performs the research especially in the field of evaluation of mechanical properties to improve higher performance MGC materials[2-4] and in the field of component design to optimize MGC gas turbine blades by computational fluid dynamics
CURRENT RESEARCH ACTIVITIES IN MECHANICAL ENGINEERING LABORATORY In this paper, the current research activities in the field of evaluation of mechanical properties of MGC materials are indicated. Evaluation of fracture toughness and fatigue crack growth resistance is discussed and
543
development of a materials testing system in simulated gas turbine environments is shown. MATERIALS AND EXPERIMENTAL PROCE DURE The MGC materials used in this study are binary insitu single crystal oxide ceramic eutectic composites Al,03N,A15012 (YAG) and ternary in-situ ceramic eutectic composites AI,O3N,AISO,, (YAG)/ZrO,. The composition of the binary eutectic composites is Alz0,iYAG=67/33wt%. They were fabricated by the unidirectional solidification method at the melting temperature of 1900°C by lowering a crucible at a speed of 5mm/h. The composition of the ternary eutectic composites is A1,03NAG/Zr02=54/28/18wtO/o. They were also fabricated by the unidirectional solidification method at the melting temperature of 1900 "C by lowering a crucible at a speed of 3Omm/h. The microstructure for the AI,O,/YAG binary eutectic composites is shown in Fig.1. Threedimensional networks of single crystal Al,O, and single crystal YAG are observed. Fracture toughness tests were conducted by the indentation fracture method at a room temperature. In the binary eutectic composites, in order to evaluate each phase, the indentation load of 75gf was kept for 10 seconds. In the ternary eutectic composites, the indentation load of 200gf was kept for 10 seconds. The value of fracture toughness K, was evaluated by the following equation.
K,c=O.O1 1E04P 6a4'( l/a)d Where; E: Young modulus 21: Crack length P: Indentation load 2a:Indentation size Cyclic fatigue crack growth tests were conducted at a room temperature in a load-controlled mode at a stress ratio R=O.l and at frequency f=l lOHz with a sinusoidal wave using a closed-loop electro-hydraulic materials testing machine.
-
I
1 I
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20bm
L Fig. 1 Microstructure for AI,O,/YAG binary eutectic composites
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Fig2
Specimen configuration and dimensions for fatigue crack growth test
Specimen configuration and dimensions are shown in Fig.2. hior to the fatigue crack growth test, a precrack about 1-2mm long was introduced by the bridge indentation method. After the fatigue crack growth tests, fractographic observation was performed on fracture surfaces using a scanning electron microscope and a laser scanning microscope. EXPERIMENTAL RESULTS AND DISCUSSION FOR ORIENTATION DEPENDENCE OF FRACTURE TOUGHNESS The micrographs of typical crack growth behaviour for the binary in-situ single crystal oxide ceramic eutectic composites Al,03/ YAG are shown in Fig. 3. It is found that in the YAG phase the crack easily propagates, and in the A1203phase the crack length is smaller than that in the YAG phase and the crack is inclined to propagate along the phase boundary. From the indentation hcture test result, it is found that fracture toughness for tne YAG phase is the smallest, and the sequence of fracture toughness is the YAG phase, phase boundary, and the A1,03 phase. The mean value of the fracture toughness on the C-R plane is 2.11~Pam", the mean value of the fracture toughness on the C-L plane is 2.17MPamIR, and the mean value on the L-R plane is 2.4SMPam". Hence, there is an orientation dependence of fracture toughness. The C-R plane, which is the smallest fracture toughness orientation, is the perpendicular plane to the solidification direction, and in this case the Palmqvist cracks are parallel to the solidification direction. 2-parameter Weibull plots of fracture toughness for the A1,03/YAG binary eutectic composites are shown in Fig.4. In each plane, a mixed Weihull distribution is indicated and each part of the mixed Weibull distribution corresponds to the YAG phase, phase boundary, and the Al,03 phase. The C-l. plane, which has the smallest fracture toughness, has thc largest shape parameter and has the smallest scatter of fiacture toughness. The C-L and L-R planes, which are perpendicular to the solidification direction, have larger mean values and larger scatter of fracture toughness than the C-R plane.
(a) YAGphase
(b) Phase boundary
(c) A1203phase
Fig.3 Micrographs of typical crack growth behaviour for A1203NAG binary eutectic composites
80.0 ......................
....................
50.0 % .........
7
-CC -
................
..................
30.0
LL
..........................
10.0
L
-4
2
1
3
L 2.0
4
5
KIC
YAG Phase
K c MPaJm
Fig.4 2-Parameter Weibull plots of fracture toughness for A1,0, NAG binary eutectic composites 0 A120,/YAG
P i-
.-99.9
A Al20,/YAG/Zr0, \
0
..................................................
1
.-
f 80.0
........................................ 0'-
Lt
' -1 7 7
C -C
1.............................
- 50.0 g
...................
-2 t ..................................................................
-
-3 ................ e O -4
- 30.0 10.0
A .................................................
w
A
.A
k 2.0
2-parameter Weibull plots of fracture toughness for the AI,O,NAG binary and AI,O,NAG/ZrO, ternary eutectic composites for the C-R plane iJe shown in Fig.5. The mean value of fracture toughness for the AI,O,/YAG binary eutectic composite$ in the C-R plane is 2.11MPam1', and that for the AI-0,NAG/Zr02 ternary eutectic composites in th.: C-R plane is 2.4MPam''. And at the same time the szatier of fracture toughness in the AI,03NAG/Zr02 ternary eutectic composites is smaller than that in the AI20,/YAGbinary eutectic composites. The AI2O3/k'AG/ZrO2ternary composites have higher straigth than the AI,O,/YAG binary eutectic composites[5]. It is concluded that higher strength and higher fracture toughness are accomplished with the ternary eutectic composites.
545
EXPERIMENTAL RESULTS AND DISCUSSION ON FATIGUE CRACK GROWTH RESISTANCE Fatigue crack growth curves for the Al,O,/ YAG binary eutectic composites are shown in Fig.6. In the RL orientation, in which the crack propagation direction was parallel to the solidification direction, the fatigue crack propagated continuously. In the L-C orientation, in which the crack propagation direction was perpendicular to the solidification direction, the fatigue crack propagated discontinuously, and the crack
.
fi . I
, * "L-C' Orientation " ' i " " t T ,
3.5
m
;
3 _
E E
!
.+
L
0 ;
_ ........ ......... .......... .......... ..........i........i :
:
:
:
:
.........:........._
0 ;
i
........; .........4..........:..........4..........:!........8..4:..........:j .........-
i
i
8 i
.mi
:
*.
+*;w
sometimes stopped at the phase bomdary, and propagated along the phase boundary. The effect of orientation on fatigue clack growth resistance for the Al,03/ YAG binary eutectic composites is shown in Fig.7. In the R-L orientation, fatigue crack growth rate dddN can be uniquely related with maximum stress intensity factor and can also be Fatigue expressed by the power law(da/dN=C Lax"'). crack growth resistance for sintered Al,03 ceramics is also shown in Fig.7. It is found the gradient of the da/dN-IC, relationship in the R-L orientation for the Al,O,/ YAG binary eutectic composites is smaller than that for sintered A1,0, ceramics. The maximum stress for the Al,03/YAG intensity factor threshold value LKfi binary eutectic composites is about 1SMPam". The maximum stress intensity factor threshold value for sintered AI,03 ceramics is 1.9MPam". Hence, the fatigue crack growth resistance in the R-L orientation for the Al,O,/YAG binary eutectic composites is the same as or smaller than that for A1,0, ceramics. 0
.
i 5 4 0 ~ ~ 3 5350N+360N 4
2.5
Fig.6
Fatigue crack growth curves for AI,O,/YAG binary eutectic composites
II
Al2O3 ceramics
-Sintered
-A1203/YAG
loa
Eutectic Cornp. R-L
I
(a) Fatigue fractured region(da/dN*
m/cycle)
Al 0 /YAG Eutectic Comp. L - C L 2 3
1 o - ~ ............................................................................
...-
2
10-l1 1 0-l2 1
Fig.7
546
2 Kmax MPaJm
3
Effect of orientation on fatigue crack growth resistance for Al,O,/YAG binary eutectic composites
Fig.8
(b) Fast fractured region Fractographs of fracture surface for R-L and L-C orientation of Al,O,/YAG binary eutectic composites
High Pressure Vessel+ Materals Testing Machine
Actuator
~-.n
I
Fig.9
Controller+ computer
I
Schematic representation of a materials testing system in simulated gas turbine environment
In the L-C orientation, in which discontinuous fatigue crack growth behaviour was observed, it has not been clarified whether the crack growth rate dddN is uniquely expressed by the maximum stress intensity factor Lax. However, fatigue crack growth rates in the L-C orientation for the AI,O,/YAG binary eutectic composites are much smaller than those for sintered AI,O, ceramics at any Lax. Hence, it is concluded that fatigue crack growth resistance in the L-C orientation for the AI,O,/YAG binary eutectic composites is larger than that for sintered A1,03 ceramics. Fig.8 shows fractographs of fracture surfaces in the R-L and L-C orientations of the A1,03/YAG binary eutectic composites. In both the fatigue fractured region and the fast fractured region in the R-L and L-C orientations, only intergranular fracture is observed. In fatigue fracture of sintered Al,O, ceramics, it is well known that transgranular fracture is usually observed. Hence, the fatigue crack growth mechanism for the AI,O,/YAG binary eutectic composites is quite different from that for sintered AI,O, ceramics.
DEVELOPMENT OF A MATERIALS TESTING SYSTEM IN SIMULATED GAS TURBINE ENVIRONMENTS Gas turbine structural materials are actually used in ultra-high temperature, high-pressure water vapor environments. Then, it is necessary to develop a materials testing system in simulated gas turbine environments. The schematic representation for the system is shown in Fig.9. The system consists of a testing machine with a high-pressure vessel, an environmental simulator, and a computer and a contoller.
The materials testing machine is controlled by a closed-loop electro-mechanical servomotor. The crosshead speed is 0.0001-1.0mm/min and the maximum load is 20kN. The high-pressure vessel consists of two concentric chambers. The environment is only coritrolled in the inside chamber. The environmental sirnulator consists of a water vapor circuit device and a gas supplier, and it can control water vapor, N, gas, and O2 gas environments up to 0.98MPa. The endurance tests will be conducted by the materials testing system.
FUTURE PROSPECT In order to improve the thermal efficiency of gas turbine systems, it is necessiuy to develop a new heat resistant material to improve the turbine inlet temperature without cooling the turbine structural components. From the results of the leading research and development[5], the thermal efficiency increases about 15% if MGC materials are applied to the turbine structural components such as turbine blade and combustion liner. Now the data on the mechanical properties and fabricating technique for MGC materials are increasing. Hence, the planning of a technology development program for a MGC ultra high-efficiency turbine system is needed.
CONCLUSION An outline of the leading research and development for “MGC Ultra High-Efficiency Turbine System” is given, and the current activities by the Mechanical Engineering Laboratory have been introduced.
547
(1) For the AI,O,/YAG binary eutectic composites, fracture toughness for the YAG phase is the smallest, and it decreases in the sequence of the YAG phase, phase boundary and the Al,O, Phase. From 2parameter Weibull plots, a mixed Weibull distribution is indicated and each part of the mixed Weibull distribution corresponds to the YAG phase, phase boundary, and the AI,O, phase. (2) Fracture toughness for the AI2O3/YAG/ZrO2ternary eutectic composites show larger values than that for the binary eutectic composites in the C-R plane. Hence, higher strength and higher hcture toughness is accomplished with ternary eutectic composites. (3) In the R-L orientation for the AI,O,/YAG binary eutectic composites, which has the smallest fatigue crack growth resistance, fatigue crack growth resistance is the same as or smaller than that of sintered Al,O, ceramics. In the L-C orientation for the A1,OJYAG binary eutectic composites fatigue crack growth resistance is larger than that for sintered AI,O, ceramics. (4) A new materials testing system which can perform mechanical tests in simulated gas turbine environments is newly developed.
ACKNOWLEDGEMENT This leading research and development is part of the New Sunshine Program, Agency of Industrial Science and Technology. The authors wish to thank New Energy and lndustrial Technology Development Organization.
548
The authors also wish to thank Technology Research Gas Turbine for Practical Association of Ability Progress and the cooperating contractor’s companies (Ishikawajima-Harima Heavy Industries Co., Ltd. Kawasaki-Heavy Industries, Ltd., UBE Industries, Ltd., Japan Ultra-high Temperature Materials Research Institute)
REFERENCES (1) Y. Waku, N. Nakagawa, T. Wakamoto, H. Ohtsubo, KShimizu and Y. Kohtoku, A Ductile Ceramic Eutectic Composite Nith High Strength at 1873K3, Nature, 389-6646,(1997) 49-52. (2) K. Hirano, A. Kamei, and T. Suzuki, Fracture Toughness Variation for In-sitb Single Crystal Ceramic Eutectic Composites, Abstract for American Ceramics Society lOlst Annllal Meeting & Exposition,( 1999) 191. (3) K. Hirano, Future Prospects of R&D on Ultra-high Temperature Structural Materials, Proc. of 7th International Fatigue Congress-Fatigue’99, Vo1.4, (1999), 2 119-2126. (4) K. Hirano, T. Suzuki, A. Kamei and F. Tamai, R&D on In-situ Single Crystal Oxides Ceramics Eutectic Composite, Proc. of 7th Euro-Japanese Symposium on Composite Materials and Tranwportation, (1999), 183-184. ( 5 ) Progress Report for the Results of Leading Research and Development for MGC Ultra High-Efficiency Turbine System(1999), Technology Research Association of Gas Turbine for Practical Ability Progress.
SYNTHESIS AND PROPERTY TAILORING OF REACTION-BASED COMPOSITES: THE RBAO AND THE 3A PROCESS Sven Scheppokat*, Mark Roeger*, Peter Beyer* *, M. Leverkohne* *, Rolf Janssen*’* *, and Nils Claussen*’* * (*) Materials Engineering Hamburg GmbH Nartenstr. 4a, D-21079 Hamburg, Germany (**) Technical University Hamburg-Harburg Advanced Ceramics Group D-2 1071 Hamburg, Germany
ABSTRACT Ceramics and metal/ceramic composites are promising materials for a variety of applications in engines due to their unique property profile including excellent high-temperature behavior, good tribological properties, oxidation resistance, and low specific weight. In order to make economic use of these materials, it is necessary to manufacture them by a low-cost process that can be adapted to varying demands. Two related processes, the RBAO (Reaction Bonding of Aluminum Oxide) and the 3A (Alumina Aluminide Alloys) process combine low cost with broad compositional variability and therefore with a wide range of microstructures and properties. RBAO and 3A materials can be green machined and provide reduced sintering shrinkage. While the RBAO process yields purely ceramic materials (alumina composites), the 3A process results in metal/ceramic composites with interpenetrating networks.
THE RBAO PROCESS In the RBAO (Reaction Bonding of Aluminum Oxide) process, mixtures of A1 and A120; are intensively milled, compacted to high green density, and fired in oxidizing atmosphere.[ 1-61 During this heat-treatment, A1 is oxidized completely to A120;. The newly formed A1203 particles sinter and bond the originally added A120;. ZrOs is normally added to the precursor powder in order to improve the microstructure and the mechanical properties of the final product. The milling is performed in organic liquids such as ethanol or acetone in order to avoid excessive hydrolyzation of the Al. The milled material is dried and then compacted. Due to the presence of ductile Al which forms metallic ligaments in the green bodies, RBAO green compacts achieve exceptionally high green strengths that allow extensive green machining. This provides an economical advantage by saving costs (e.g. elimination of need to manufacture dies for pressing samples of specific
geometry) and is also suitable for fast manufacturing of small numbers of components or spare parts.
Fig.1: Green machined RBAO compacts
Typical green bodies compacted unidirectionally at 50 MPa and subsequently isostatically pressed at 300MPa reach strength values of about 20 MPa.The green bodies are then heated in air slowly to 1000°C to allow the A1 oxidation to take place, after that they can be fired fast with a typical sintering temperature of 1550°C. Sintered RBAO components achieve typical strength values > 650 MPa and a fracture toughness of 3.5 MPadm. Increasing the zirconia content results in fracture toughness values > 5 MPadm. The RBAO process can be modified by adding other phases, e.g. additions of S i c to the precursor lead to mullite materials and make zero shrinkage possible.
3A MATERIALS The Alumina Aluminide Alloys (3A) process is a novel route for low-cost manufacturing of metaVceramic composites with interpenetrating networks.[7- 101 While the densification of metal/ceramic composites usually requires the application of pressure, [11,12] the 3A process allows pressureless densification. The process
549
consists of a redox reaction between Al and a metal oxide M,Ob A ] + MaOb + M i-AI2o3 (1) followed by the formation of an intermetallic phase between the metal and excess A1 in the material: + M,Al, (2) AI+M Due to its very high affinity to oxygen, A1 reduces the oxides of many other metals, these will be called the reactive oxides here. Examples of reactive metal oxides that have been used in this process include Ti02, Fe203, Cr103, NiO, Zr02, and Nb205. The composition of the aluminide formed is determined by the ratio of residual A1 to the metal M in the material after completion of the redox reaction. For example, in the system Ti02/Al, materials containing TiAl, Ti3A1, and Ti have been made. Alumina is always present in the final product as a product of the redox reaction. Depending on their composition, aluminides can be very refractory (T, NbA13 = 1657"C, T, Nb3AI = 1960"C), as well as low specific weight (p TiAI3 = 3.4 g/cm3). Depending on the type and volume fraction of the metal in the final product, 3A materials reach strength values > 600 MPa and fracture toughness values >10 MPadm [8]. There are two basic variants of the 3A process: One of them is a powder metallurgical, or sintered, route (s-~A),the other one involves metal infiltration of a porous ceramic preform (i-3A). The basic reactions in both cases are the same. The s-3A process starts from intensively milled precursor mixtures of Al, A1203, and a reactive metal oxide MaOb. Due to the presence of a ductile metal phase (Al), the powder compacts made from these mixtures exhibit high strength and can be green machined. Densification is achieved by pressureless sintering in inert atmosphere or vacuum. At temperatures > 5OO0C, the redox reaction between A1 and the reactive metal oxide takes place. This reaction proceeds mainly as a solid state reaction below the melting point of A1 (660°C). Because the redox reaction is generally highly exothermal, carefkl process control (slow heating ramps around the reaction temperature) is needed in order to avoid thermal runaway. The densification of the materials typically takes place at temperatures between 1400 and 1550°C. A variant of the s-3A process is the so called 3AMC (3A Metal Composites) process. In this case, the precursor mixtures contain a metal, such as Fe or Cr instead of the reactive oxide, and only a small amount of Al. The primary objective of the A1 addition here is to improve the sintering behavior of the material by reducing the native oxide layer on the metal. Because only a small amount of Al is present, and the redox reaction takes place only on a very small scale, much less energy is released, and thermal runaway is not likely to occur. Therefore, this version of the process requires less careful control of process parameters, and faster heating rates can be employed. In the case of Fe/AI2O3 composites, the usehlness of A1 as a sintering aid has been established [S], in the case of Cr/A1203 and Nb/AI2O3composites, dense materials have also been attained without the addition of A1 [13,14]. However,
550
recent experiments indicate that A1 also promotes densification in the case of Cr.
Fig. 2:SEM micrograph of a 30 Vo1.-YOCr(A1) / 70 Vo1.-YOA1203composite Fig. 2 shows an SEM micrograph of a 3AMC composite containing 30 Vo1.-% Cr(AI) and 70 VoL-% A1203.The material is dense with few pores. The metal phase shows a high degree of interconnectivity, and the material is electrically conductive. Grain sizes are in the range of up to 5 pm. The metal content has a strong influence on the mechanical properties of the final product. In general, strength decreases, but toughness increases with increasing metal content. In the i-3A process, a porous preform containing A1203and the reactive oxide is infiltrated with liquid A1 and subsequently annealed at higher temperature in order to allow the reaction to take place. The i-3A process is a net shape process that allows low temperature synthesis of materials that are suitable for high temperature applications due to the formation of the refractory intermetallic phase. It is also particularly suitable for local reinforcement of components and can be used in existing industrial processes such as die casting or squeeze casting. With its different versions, the 3A process and its derivates are very versatile. They allow the cost-efficient manufacturing of materials covering a wide range of compositionswith good properties.
ACKNOWLEDGEMENT The authors wish to thank NED0 and the Innovation Foundation Hamburg (Innovationsstihng der Stadt Hamburg) for their support.
REFERENCES: (1) N. Claussen, T. Le und S. Wu: ,,Low-Shrinkage Reaction-BondedAlumina", J. Eur. Ceram. SOC.5 29-35 (1989)
(2) N. Claussen, N.A. Travitzky und S. Wu: ,,Tailoring of Reaction-Bonded A120; (RBAO) Ceramics", Ceram. Eng. Sci. Proc. 1 1 806-820 (1990) (3) S. Wu, D. Holz und N. Claussen: ,,Mechanisms and Kinetics of Reaction-Bonded Aluminum Oxide Ceramics"; J. Am. Ceram. SOC. 76 [4] 970-980 (1993) (4) D. Holz, S. Wu, S. Scheppokat und N. Claussen: ,,Effect of Processing Parameters on Phase and Microstructure Evolution in RBAO Ceramics"; J. Am. Ceram. SOC.77 [lo] 2509-25 17 (1994) ( 5 ) D. Holz: "Herstellung und Charakterisierung von reaktionsgebundenen A1203-Keramiken (RBAOVerfahren) am Beispiel des Systems A1203/Zr02", PhD Thesis (in German), VDI Verlag FortschrittBerichte Reihe 5 Nr.367 (1994)
(6) M. Roger: "Aufbereitung und Verarbeitung von RBAO-Precursormischungen zur Herstellung hochfester Bauteile", PhD Thesis (in German), VDIVerlag Fortschritt-Berichte Reihe 5 Nr.524, (1 998) (7) N. Claussen, D.E. Garcia, and R. Janssen, "Reaction sintering of alumina-aluminide alloys (3A)", J. Mat. Res. 11 [ 1 11 2884-2888 (1996) (8) S. Schicker, T. Emy, D.E: Garcia, R. Janssen, and N. Claussen, "Microstructure and Mechanical Properties of Al-assisted Sintered Fe/A120; Cennets", J. Eur. Ceram. SOC.19 [13/14], 24552463 (1999) (9) S. Schicker, D.E. Garcia, J. Bruhn, R. Janssen, and N. Claussen, "Reaction Processing of A1202 Composites Containing Iron and Iron Aluminides" J. Am Ceram. SOC.80 [9], 2294-2300 (1997) (1 0) D.E. Garcia, S. Schicker, J. Bruhn, R. Janssen, and
N. Claussen, "Synthesis of Novel Niobium Aluminide-Based Composites", J. Am Ceram. SOC. 80 [9], 2294-2300 (1997) (1 1) X. Sun, J.A. Yeomans, "Microstructure and Fracture Toughness of Nickel Particle Toughened Alumina Matrix Composites", J. Mat. Sci. 31, 875880 (1996)
(12) P.D. Djali and K.R. Linger, "The Fabrication and Properties of Nickel-Alumina Cermets", Proc. Br. Ceram. SOC.26 113-127 (1978) (13) D.E. Garcia, S. Schicker, J. Bruhn, R. Janssen, and N. Claussen, "Processing and Mechanical Properties of Pressureless-Sintered Niobium-Alumina-Matrix Composites", J. Am Ceram. SOC.80 [9], 2248-2252 (1997) (14) D.E. Garcia, S. Schicker, R. Janssen, and N. Claussen, "Nb- and Cr-A120; Composites with Interpenetrating Networks", J. Eur. Ceram. SOC.18 601-605 (1998)
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PROPERTIES OF SILICON NITRIDE/CARBIDE NANOMICROCOMPOSITES ROLE OF S i c NANOINCLUSIONS AND GRAIN BOUNDARY CHEMISTRY
-
P. sajgalik, M. Hnatko, Z. Len&% Institute of Inorganic Chemistry, Slovak Academy of Sciences, Diibravskh cesta 9, SK-84236 Bratislava, Slovak Republic inclusions on the mechanical behavior of composite.
ABSTRACT Two different types of S i c nano-inclusions were identified within the Si3N4/SiC nano/micro composite. Sic with the oxygen rich interface is created by the reduction of Si02 present in the silica melt and the S i c inclusion with ,,clean" interface is created by the precipitation of SiCN amorphous powder. The first type of nano-inclusion influences the high-temperature properties while the second one affected the roomtemperature properties of the composite.
INTRODUCTION SiC/Si3N4nano/micro composites are intensively studied during the last decade [ 1-61. Good room as well as high temperature properties are reported for these materials. The background of their excellent mechanical properties is widely discussed from their introduction by Niihara, [I]. Nano/micro composites consist of Si3N4micro-grains and S i c nano-inclusions, which are located within the host Si3N4grains and on their grain boundaries. As it was reported in the works of Pan et al. and Sajgalik et al. [7, 81, not only position of nano-grains is different (grain boundary and grain inside) but also the surface chemistry. The S i c nanoinclusions with a "clean" surface and the inclusions containing the oxygen rich surface layer were reported. Sic nano-inclusion origin and the role with respect to the mechanical behavior of composites is not clear yet in spite of the fact, that Pan suggested the hypothesis of their formation, [7]. Hypothesis of Pan was based on the size effect of the nano-inclusions. Based on a model experiment, the present paper shows the origin of inclusions and discusses the effect of both types of
the
EXPERIMENTAL SiC/Si3N4 nano/micro-composites were prepared by seeding of the starting powder with amorphous SiCN fine powder, Table 1. Amorphous SiCN powder, with calculated composition given in the Table 2, was added together with the alumina (Alcoa, USA, grade A16) and/or yttria (H.C. Starck, Germany, grade Fine) in the amount listed in Table 1 to the crystalline Si3N4 powder (UBE Industries, Ltd., Japan, grade SN-E10). The powders were attrition milled in isopropanol for 4 h. Dried powder mixtures SNYA and SNYAIO, respectively (Table 1) were cold isostatically pressed (CIP) at a pressure of 750 MPa. The CIPed samples were gas pressure sintered in BN crucible at 1900 "C for 2 h in a nitrogen atmosphere of 10 MPa. The density of samples measured by the water immersion method was > 99 % theoretical density (TD). Samples SNY and SNY20, respectively (Table 1) were cold pressed in the steel die with the pressure of 100 MPa, then embedded into the BN and hot-pressed by pressure of 30 MPa at 1750 "C and 0.2 bar over-pressure of nitrogen. Density of hot-pressed samples were measured by mercury immersion method and for all samples was > 98 % TD. The bars of (45 x 4 x 3) m for mechanical testing were cut from densified prisms (50 x 15 x 10) mm after gas pressure sintering and discs (50 mm diameter and 5 nun height) after hot pressing, respectively. The bars with tensile surface polished to 15 pm finish were used for the 4-point bending tests (20/40 innedouter span, cross-head speed O.Smm/min) at room temperature.
Table 1 Composition and sintering method used in processing of samples for RT and HT measurements. Properties
Sample
RT
SNYA
RT
SNYAlO
HT
SNY
HT
SNYZO
Sintering method GPS 1900 "C12 h GPS 1900 "C/2 h HP 1750 "C/2 h HP 1750 "C/2 h
Si3N4 wt% 92
A1203 wt. % 3.4
Y203 wt. % 4.6
amorph SiCN wt. %
82
3.4
4.6
10
95
5
75
5
20
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Table 2 Calculated composition of the amorphous Si-N-C powder after crystallization. Si3N4 Sic Free carbon wt. %
wt. %
wt. %
66.36
27.21
6.43
Fracture toughness was measured by the indentation method and calculated using the formula of Shetty, [9]. The creep tests in bending (20/40 innedouter span) were performed in air from 1200 to 1400 "C, with step 50 "C. Stepwise loading regime was selected at each temperature. In order to study the S i c inclusion origin the simple model experiments were performed. Samples were prepared by mixing of oxide powders Y2O3 (PID, 600 101-024, 99.99%), amorphous SiOz (50 m'g-', Aerosil OX-50, Degussa, Germany) and pigment grade carbon black (SBET= lo00 m'g-'). The chemical compositions of samples are given in Table 3. The powders were ball milled for 24 hours in isopropyl alcohol, dried and subsequently cold pressed in the steel die with a pressure of 100 MPa. Sample SY was heated at temperature 1750°C with heating rate of 10"C/min. Sample SYC20 was heated at temperature 1940°C with heating rate of 5O"C/min in a graphite crucible with powder bed of YzO3-SiO2-BN composition. The lower temperature (1750°C) was selected to have a similar conditions like during sintering of Si3NdSiC composites. A slight overpressure of nitrogen (0.15 MPa) was used during heat treatment. The compacts were heat treated with intention to observe the chemical changes in the melt. Surfaces of melted specimens were polished to l p n finish and plasma etched with CF, and 0 2 gas to highlight the grain structures. The etched surfaces were examined by SEM (JEOL JSM-35), and the elemental analysis was conducted using energy disperse spectrometry (EDX, Cambridge, U.K). The crystalline phases present in the ground samples were identified using X-ray diffractometry (STOE powder diffraction system with C O Gradiation).
RESULTS AND DISCUSSION Origin of the S i c Inclusions SIC inclusions with oxygen rich surface layer SiCN amorphous powder contains free carbon, Table 2. That is why the model experiments were concentrated on the possible chemical reactions of carbon with the oxide constituents, i.e. Si02 and YzO3. Composition of model glasses, which simulate the composition of grain boundary films in the SNY and SNY20 composites is listed in Table 3. SY is a glass of eutectic composition in the phase diagram Si02-Yz03 (TE = 1660 "C). To this basic composition 20 wt% of carbon was added and heated up to 1940 "C. The high carbon content in sample SYC20 (20wt%, C/Si02 = 3.09) caused carbothennic reduction of SiOz and formation of Sic, and consequently a shift of starting composition (TE = 1660 "C) to the region with higher eutectic temperature according to the phase diagram
554
(TE(YzSiz07-Y4Si301z) = 1900 "C ). The XRD analysis showed that the major phases in this specimen were YzO3 and Sic, Table 4. The matrix of this sample contains areas with different colour. It indicates the presence of phase with different composition as pure YzO3. EDX analysis of the matrix of SYC20 specimen proved the presence of yttrium and oxygen in great amount and a small amount of silicon and carbon. The following reaction took place: SiOz(s,l) + 3C(s) 3 SiC(s) + 2CO(g)
(1)
KI(1750°C) =96.4, AH1(1750°C) = -573.8 kJ/mol K1(1900"C)= l.01x103, AH1(1900"C)=-569.03kJ/mol
SIC FORMATION
SURFACE OXIDE LAYER FORMATION
Fig. 1Schematic of S i c formation by reduction of SiOz within SiOz + YzO3 melt The schematic of the present reaction within the silica melt is shown in Fig. 1. The experimental confirmation of the schematic was partially proved by the model experiment and the EDX analysis shown in Fig. 2. The intimate contact of amorphous carbon and S i c inclusion within the host Si3N4 grain, reported in previous works [8,10-131 and the S i c produced by reaction (1) shown in Fig. 2 are considered a confirmation of the processes leading to the formation of S i c nano-inclusions with an oxide rich surface layer. This proof is supported by the previous work of authors [8], where the EDX analysis confirmed that the surface of inclusions with an oxygen reach layer is poor on Si. This observation also excellently confirms the proposed chemical origin of these inclusions.
High Temperature Creep The high temperature creep measurements show an improvement of creep resistance of the nano/micro SNY20 composite comparing to the reference SNY. The improvement of the creep resistance is usually
Fig. 2 S i c formation within the silica-yttria glass by addition of amorphous carbon Table 3: Composition of glass samples Sample Carbon content [wt. Yo] SY SYC20 20 attributed to the distribution of the fine S i c nano-grains along the grain boundaries [ 141. Recent publication of Rendtel et al. [5] shows that there exists an optimum of Sic nano-grains content which lies within the range of 10-15 wt. %. The amorphous SiCN powder added in the amount of 20 wt. % to the starting mixture yields 5.5 wt. % of S i c present in the SNY20 nanolmicro composite. This value is far below 10-15 wt. % optimum. In spite of this fact the creep strain of SNY20 is approx. one order of magnitude lower comparing to the reference SNY ceramics, Fig. 3. The model experiment described above can be used for explanation of creep behaviour of nano-micro SNY20 composite. The excessive carbon introduced along with SiCN amorphous powder reduces SiOl content similarly as it was observed in the model experiment. This fact has two consequences: decrease amount of oxide grain boundary phase, shift of eutectic temperature higher, according to the phase diagram the next eutectic point (at decreased content of SOz) is at 1900 “C. That means the contribution to one order increase of creep resistance of nano-micro SNY20 composite comparing to the monolithic SNY is not caused only by distribution of S i c submicrometer grains at the grain boundary but mainly by changing the chemistry of grain boundary phase. This explains rapid increase of creep resistance in presented case, even the content of Sic nano-grains is rather low, 5.5 wt% comparing to the optimum of 10-15 wt% suggested by other authors [51.
SIC inclusions with “clean” surface SiCN amorphous powder is added to the starting mixture as an microstructure forming additive. SiCN is solving in the silicate melt at the temperature of sintering. After solving the original amorphous powder
y203
[wt. Yo] 59.37 47.53
Si02 [wt. Yo] 40.63 32.47
precipitate in three different chemical forms, as it is schematically shown in Fig. 4. These are as follows, crystalline p-Si3N4, S i c and amorphous carbon, Fig 4b. S i c precipitated from amorphous SiCN a source of inclusions with “clean” interface.
Room Temperature Properties The strength and fracture toughness of composites prepared in the present study are listed in the Table 5. Microstructure of materials, SNYA and SNYA10 is morphologically similar, except large elongated grains (up to 25 -30 pm) in the monolithic SNYA. High Weibull modulus of 19 of microlnano-composite give an evidence that the size of technological defects is undercritical. Fine S i c precipitates in the microstructure serve as the grain growth inhibitors, as it was pointed out in [8] and these are responsible for the homogeneous particle size distribution. Strength of 1.2 GPa of nano-composite (20% higher comparing the strength of monolith) cannot be explained only on the base of microstructure. The inclusions with a “clean” surface are causing the stresses within the grains because of different thermal expansion coefficients of S i c and Si3N4, [8, 131. These stresses influences the overall status of the composite. The simple calculation taking into account the difference in thermal expansion coefficient between S i c nano-inclusion and Si3N4host micro-grain and using Hook’s law shows that the tensile stress of 1.3 GPa is produced around the Sic inclusion by cooling the sample from the sintering temperature when the inclusion is incorporated into the host grain by grain growth mechanisms, Fig. 5. This stress is large and deform the host Si3N4 lattice, as it was pointed out in [8]. Si3N4grain with S i c inclusions containing a glassy surface layer is supposed to be stress free because of relaxation of thermal stresses within the oxygen layer.
555
Fig. 3 Creep rate of SNY20 nano/micro composite and the refering Si3N4based ceramics SNY Table 4: Results of the phase analysis of the samples Sample Firing cycles Heating Molar ratio rates CbJSiO2 SY 1750°C / 0.5h / N2 10"C/min SYCZO 1940°C / 0.5h / N2 5O"C/min 3.09
Phase detected by XRD Y2Si207-orthoromb.,SiOz, Si2N20 Y2O3, S i c
Fig. 4 Schematic of microstructure formation. a) a-Si3N4and amorphous SiCN are solving in the Si02+ Y203 melt, p- Si3N4remains stable; b) b) p-Si3N4and S i c ("clean" surface) precipitate from the melt, carbon is reacted with silica and S i c (with oxygen rich layer) are formed by chemical reaction (marked by arrows in Fig. 4c); c) c) final microstructure consisting of p-Si3N4micrograins containing both types of nano/inclusions.
Presence of residual stresses of above reported quantity which are concentrated into the Si3N4 micro-grain containing S i c inclusions with glassy free surface layer must play an important role. Tensile stresses within the Si3N4 grain lead to the creation of the sub-grain boundaries, as it is schematically shown in Fig. 6. If the refinement of the microstructure by presence of stress caused sub-grain boundaries is the reason for increase of strength or if the other phenomena take part should be subjected to the further systematic investigation. The energy-filtering TEM investigation showed that the SiC/Si3N4 nano-composite consists of Si3N4 micro-grains, S i c microhano-grains, Si3N4 micro-grains with S i c nano-inclusions. Different grains the microstructure of material SNYAIO consists of cause also the other changes in mechanical behavior.
556
TI=Tsinterinu
T2=7room
-
AVNn = l/(l+Ps~dT) 141+PscAn p = [E/3(1-2p)] AVWT~= 1.3 GPa Fig. 5 Schematic of residual stresses formation within the Si3N4grain with S i c nano-inclusion
Strength of particular grain boundary must be different because of stress caused by neighbourhood of various kinds of grains. Presence of S i c micrograins with higher coefficient of thermal expansion comparing to Si3N4grains causes tension at their grain boundaries. Similarly the boundary between Si3N4and Si3N4grain with oxygen free S i c inclusions should be under tension because of tensile stress caused by their presence. These statements lead to the logical conclusion that SiC/Si3N4 nano/micro-composite contains “stronger” and “weaker” grain boundaries between the micro-grains as a results of their thermal stresses. The strength of the micro-grain boundaries plays an important role in influencing the fracture behavior of the composite and consequently the Table 5 RT Mechanical properties. Sample Fracture toughness/h4Pa.m’n 7.4 SNYA SNYAlO 6.9 SNY 7.1 SNY20 6.4
fracture toughness. Weak boundaries will promote the inter-granular fracture in the nano/micro composite, Fig. 7. Their occurrence will contribute to the prolongation of the crack path. From this point of view a number of weak boundaries (volume fraction) is an important parameter. Their volume fraction will be depended on the volume fraction of Si3N4micro-grains with S i c inclusions without any glassy interlayer. It means, the fracture toughness would be also indirectly influenced by the volume fraction of the S i c inclusion with a “clean” interface to the host Si3N4grain. As can be seen from Fig. 8, the S i c inclusion within the large grains can serve as a fracture origin. In this sense is their role strongly negative.
Bending StrengtMPa
Weibull modulus
990 1203 870 710
7 19
Fig. 6 Schematic of the sub-grain boundary formation in SiC/Si3N4nano/micro composite
Fig. 7 Intergranular crack path within the Si3NdSiC nano/micro-composite, bar 2 pm
Fig. 8 Fracture surface of large Si3N4 grain with S i c nano-inclusions
557
CONCLUSIONS The present paper showed that the origin and the amount of S i c inclusions within the Si3N4 based composites are important factors influencing their room as well as high temperature properties. The S i c nano-inclusions with the “clean” interface are originated by the percipitation of an amorphous SiCN powder present in the starting powder mixture. These are considered to affect the room temperature fracture toughness and strength of the composite. The S i c nano-inclusions with the oxygen rich interface are originated by the chemical reaction of the free carbon (present in the SiCN powder) with the silica rich melt, which creates during sintering. This reaction is responsible for the decrease of the overall volume fraction of the residual glass phase in the composite and for its chemical modification. These nano-inclusions affect the high temperature properties of the composite. The hypothesis of Pan, [7] on formation of S i c inclusion with oxygen rich layer depending on the inclusion size was not proved.
ACKNOWLEDGEMENT Present work was partly supported by Slovak Grant Agency VEGA, project 215 118/00. P. Sajgalik acknowledged the Alexander von Humboldt Foundation for the finacial support during his stay at University of Karlsruhe, Germany where a part of this work was carried out.
REFERENCES K. Niihara, New design Concept of Structural Ceramics - Ceramic Nanocomposites, J. Jpn. Cer. SOC. 99[10] (1991) 974-982. K. Niihara, K. Suganuma,, A. Nakahira, K. Izaki, Interfaces in Si3N4 Nano-Composites, J. Muter. Sci. Lett. 9 (1990) 598-599. G. Sasaki, K. Suganuma, T. Fujita, K. Hiraga, K. Niihara, Interface Structure of Si3N4 Matrix Composite with Nano-Meter Scale S i c Particles, Mar. Res. Symp. Proc. 287 (1993) 335-340. M. Herrmann, C. Schubert, A. Rendtel, and H. Hubner, Silicon Nitride/Silicon Carbide
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Nanocomposite Materials: I, Fabrication and Mechanical Properties at Room Temperature, J. Am. Ceram. SOC. 1998,81(5), 1094-1108. 5. A. Rendtel, H. Hubner, M. Henmann, and C. Schubert, Silicon Nitride/Silicon Carbide Nanocomposite Materials: 11, Hot Strength, Creep, and Oxidation Resistance, J. Am. Ceram. SOC.81[5] (1998) 1109-1120. 6. P. Rendtel, A. Rendtel, H. Hubner, H. Klemm, and M. Henmann, Effect of Long Term Oxidation on Creep and Failure of Si3N4 and Si3N4/SiC Nanocomposites, J. Eur. Ceram. SOC. 19[2] (1999) 217-226. 7. X. Pan, J. Mayer, M. Ruhle, J. Am. Ceram. SOC. 79 [39] (1996) 585-90. 8. P. Sajgalik, M. Hnatko, F. Lofaj, P. HvizdoS, J. Dusza, P. Warbichler, F. Hofer, R. Riedel, E. Lecomte, M.J. Hoffmann: SiC/Si3N4NanoMicroComposites - Processing, RT and HT Mechanical Properties, Journal of the European Ceramic Society, 20 [4] (2000) 453-62. 9. D. K. Shetty, I.G. Wright, P.M. Mincer, A.H. Clauer, Indentation Fracture of WC-Co Cermets, J. Mater. Sci. 20 (1985) 1873-1882. 1O.P. Sajgalik,, J. Dusza, F. Hofer, P. Warbichler, M. Reece, G. Boden, J. Kozinkovi, Structural Development and Properties of SiC-Si3N4 Nano/Micro Composites, J. Muter. Sci. Left., 15 (1993) 72-76. 1l.P. Sajgalik, K. Rajan, P. Warbichler, F. Hofer and 3. Dusza: Silicon Nitride Based Nano- and MicroComposites with Enhanced Mechanical Properties, Key Engineering Materials 159-160(1999) 405- 10. 12.P. Sajgalik, K. Rajan, R. Riedel: Sub Grain Boundary Formation in Si3N4Based Ceramics, Key Engineering Materials 161-163(1999) 229-34. 13.P. Sajgalik, M. Hnatko, Z. LencBS: Silicon nitridelcarbide nano/micro composites for room as well as high temperature applications, Key Engineering Materials 175-176(2000) 289-300. 14.K. Niihara, T. Hirano, K. Izaki, F. Wakai, High Temperature Creep/ Deformation of Si3NdSiC Nanocomposites, in Silicon-Based Structural Ceramics, B.W. Sheldon and S. C. Danforth Eds., Ceramic Transactions Vol. 42, The American Ceramic Society, Ohio 1994, pp. 208-219.
LAYERED Si3NJ(SiAlON+TiN) COMPOSITES WITH SELF-DIAGNOSTIC ABILITY Z. LenEeS"', P. Sajglik', M. Balog', K. Frohlich', E. Roncari3 (1) Institute of Inorganic Chemistry, SAS, 842 36 Bratislava, Slovakia (2) Institute of Electrical Engineering, SAS, 842 36 Bratislava, Slovakia (3) Research Institute for Ceramics Technology, CNR, 480 18 Faenza, Italy ABSTRACT Multilayered Si3Nd(P-SiA10N+TiN) composites with functional properties were prepared by reactive hot pressing. Improvement of bending strength and fracture toughness of layered materials was observed in comparison to the bulk P-SiAlON+TiN composite. High anisotropy was achieved for the electrical resistance of the functionally graded layered materials, in which the TiN content stepwise increased from zero to 33~01%. The laminated structure and the knowledge of stress status of individual layers allowed the design of functional material with self-diagnostic ability. Monitoring the change of electrical conductivity of pSiAlON+TiN composite under tensile load seems to be a promising method for the prediction of crack generation and preventing fatal fractures.
INTRODUCTION In the last decade the preparation of ceramic layered materials has gained increasing attention not only due to the possibility to combine materials with different physical properties but also due to their lower sensitivity to defects [ 1-31. Remarkable improvement in strength and fracture toughness was achieved for this kind of ceramic materials. Nevertheless, the catastrophic failure of ceramic components hinders their wider application in the industry. Fracture prediction in engineering ceramics is a promising method to increase the reliability and practical application of ceramic materials. One of the possible methods for forecasting the fracture is the measurement of electrical resistivity of conductive layer as generally the resistivity increases with cracks generation. Layered material design is a very useful tool for the preparation of ceramic composites with built in selfdiagnosis function [4-61. For non-oxide ceramic composites as the electro conductive phase TiN might be used due to its low resistivity (3.10-'SZcm). On the other hand, TiN has a poor oxidation resistance, which hinders its high temperature application in air. This can be solved by the preparation of Si3N&3-SL410N+ TiN) layered composite. All the components of this laminate are separately used for special engineering applications. In layered form their advantageous properties can be combined: Si3N4 layers have a high strength and (pSiAlON + TiN) layers have higher hardness and low thermal Conductivity. The P-SiAlON (Si6,AlzOzN8~z,
where 0 < z 4.2) matrix has a good oxidation resistance [7], and can protect TiN against corrosion. A number of papers were published on Si3N4/TiN particulate composites and the preparation of laminated damage resistant Si3NdTiN trilayer composites was also reported [8,9]. In this work except of improved mechanical properties anisotropic electrical conductivity of layered material is expected, because Si3N4is an insulator with extremely high electrical resistivity 1013 Rcm, while TiN is conductive [lo]. The combination of this kind of materials is advantageous for self-detection fimction, because composites consisting of an insulating matrix (e.g. SiAlON) with electroconducitve inclusions are known to show a large variation in electrical resistivity with changes of composition. A large decrease in electrical resistivity occurs above the threshold volume fraction of conductive phase through the formation of conductive paths. In this composites fracture can be predicted by measuring the change of electrical resistivity under an external load. Present paper deals with the design and preparation of the Si3Nd(P-SiA10N+TiN) layered composite with enhanced mechanical properties and modified electrical conductivity. The role of the residual stresses with respect to the electrical conductivity is also discussed. EXPERIMENTAL SN and SNT is the designation of two basic powder systems used. The SN starting powder mixture consisted of a-Si3N4powder (E-1 0, Ube Industries, Japan), 5 wt% Y203 and 2 wt% A1203 (both Hokko Chemicals, Japan) sintering additives. The SNT mixture consisted of aSi3N4,AlN (type F, Tokuyama Co., Japan) and Ti02 (<5 pn, Wako Pure Chemical Industries, Japan). The compositions were adjusted to have a P-SiAlON matrix with 10 - 33 vol% TiN in the final product after the following combined reaction: 4(6-z) Si3N4+ 62 TiOz + 122 AlN = 12 Si6.,Al,0zN8., + 62 TiN + 2 N2 (1) In the case of composites with higher than 21 vol% TiN in the final product a part of this conductive phase was added as TiN powder (TTN15, Tioxide Chem., UK), because it's preparation by in-situ reaction would produce A1203exceeding the limiting solubility in Si3N4 to form P-SiAlON (ZS 4.2). The starting powders were milled with Si3N4balls in Turbula homogeniser for 2 hours. Methyl-ethyl-keton
559
contained also 2 or 5 vol% J3-Si3N4 seeds (SN-2 and SN-5). The suspensions were homogenised on rollers for 35 hours. Green ceramic sheets were formed fiom the slurries by tape casting with final thickness of -100 pm. The SMON+TiN tapes used in L4 samples (Fig. 1a) were covered &om one side with BN spots by screen coating. The diameter of BN spots was 250 pm [l 13. The ceramic sheets were punched, stacked and finally warm pressed at 120°C. The order of layers is schematically shown in Fig. 1. The binder was burned out at 60OoC/3h. Samples were sintered in the hot pressing graphite resistance fiunace. For the in situ preparation of TiN a slow heating rate 2"Clmin was applied in the range 1350-1550°C under vacuum. Afterwards nitrogen gas was introduced and 3 0 MPa mechanical load was applied. The samples were hot pressed at 1800°C for 2 hours in 0.15 MPa nitrogen atmosphere.
electrical resistance was registered.
RESULTS AND DISCUSSION
The density of sintered samples was close to t theoretical density (> 98%). Only a few pores we observed near the surface of samples with TiN, optimised conditions were used for densification [ 12 The optical micrograph of dense L4-20 sample is show in Fig. 2. The microstructure of STN-20 (SiAION 20~01%TiN) layer and SN-5 (Si3N4+ 5% seeds) lay of this sample are shown in Figs. 3a,b. A certa alignment of TiN particles and elongated p-Si3N4grai grown on seeds can be observed in direction normal the applied load due to the hot pressing and previo tape casting.
4 SN-0 STN SN-5
I
I
SN-2 SN-0
Fig. 2. SEM micrograph of L4-20 sample.
SN-0 STN-15 STN-25 STN-33 STN-25 STN-15 I
I
SN-0
Fig. 1 Schematic of layer order in: a) L4 samples (numbers indicate the P-Si,N, seed content) and b) FGM samples (numbers indicate the TiN content).
Bulk densities were measured by the Archimedes method in water. Three-point flexural strength measurements were carried out at room temperature on 3 x 4 x 36 rnm bars with a span of 30 mm, and a crosshead speed of 0.5 mm/min. The fiacture toughness was measured by Single-Edge-V-Notched Beam (SEVNB) method.
560
The layer sequence of L4 samples is complex a one "layer unit" is built fiom three different Si3N4laye and one P-SiAlON+TiN (STN) layer. The p-Si3N4se content increases towards the STN layer, Fig. 2. T aim of this design was to get gradually increasi compressive internal stresses in the layers aft sintering. The layer units are repeated several tim across the sample. The detailed theoretical backgrou is given in [13,14]. With increasing seed content t sintering rate and shrinkage of SN layers decreased, the elongated grains hindered densification. It result in residual stresses. The residual stress of adjacent Si3 layers was partly influenced also by the preferentia oriented elongated p-Si3N4 particles (there is anisotropy for the CTE of p-Si3N4in a and c direction where %= 3.3.104 K ' and ac=3.8.10-6 K ' , [15]). T presence of residual stresses was identified by Vicke indentation, because at all layer interfaces cra deflections were observed. The layered structure of sintered FGM sample shown in Fig. 4. With increasing TiN content al increases the tensile stress in the layer. Due to t stepwise increase of TiN content the cracking of STN 33 layer [12] was successfully solved.
The 3-point bending strength of sintered samples is shown in Fig. 6. A similar trend is observed as for the fracture toughness. The strength of layered composites 11 I 10
I
-
I
A
v
9 8 -
layered 10% TIN layered 20% TiN layered 33% TiN
4 3
M-I 0
L4-10
L4-20
FGM
Fig. 5. Fracture toughness ( S E W ) of monolithic and layered samples
was much higher compared to the monolithic M-10 sample due to the presence of strong Si3N4layers. With increasing TiN content the strength of L4 samples decreases due to the increasing residual stresses. Higher residual tensile stresses and lower volume fraction of Si3N4layers caused the decrease of strength of FGM material. Fig. 3. Microstructure of (a) STN-20 and (b) SN-5layers A
1100
v
a
5 5
F
.
1000 900 -
p!
L
u)
o, 800
C .U
5
m
-
700 600 -
-
M-10
L4-10
L4-20
FGM
Fig. 6. Three-point bending strength of monolithic and layered samples
Layer sequence
Comparing all samples the L4-10 layered material has the highest strength and fkacture toughness.
Fig. 4. Optical micrograph of FGM material.
The mechanical properties of monolithic and layered samples are shown in Figs. 5,6. The fracture toughness of layered L4-10 material is markedly higher in comparison to monolithic M-10 sample, Fig. 5. The high increase of fkacture toughness (107 %) can be explained by the laminated structure, by strong alignment of elongated P-Si3N4 particles, and by reasonable thermal expansion mismatch between Si3N4 and SiAlON + TiN layers. Moreover, the BN spots significantly deflected or entrapped the propagating cracks.
Physical properties The measured electrical resistivity of monolithic M10 sample and layered samples with increasing TiN content are listed in Table 1. Anisotropic electrical conductivity was measured for the layered L4-20 and FGM samples in parallel (x) and perpendicular @) direction to the layers, Table 1. In parallel direction the electroconductive TiN increased the conductivity. Although, the TiN content was only 20 vol% in the (p-SiAlON + TIN) layer, good electrical conductivity was obtained for these layers due to the network-like distribution of TiN.
561
Table 1.Electrical resistivity p of samples
TiN
Sample
Px
A
1 : ::3 1 in situ + powder
0.5
M-10
10
5.106
5.10'
L4-10
10
8. lo'
2.1012
L4-20
20
6.10-2
5.10"
FGM
33
4.1 O4
8.1 O'O
0.0
t
a A
10
19
I 20
21
22
a
23
TiN content I vol%
The electrical conductivity of layered material was much lower in perpendicular direction to the layer alignment, because the presence of insulating Si3N4 layers. Very good conductivity was obtained for the FGM material with 33 vol% TiN, as it was expected. Self-diagnosis However, good mechanical properties with reasonable Weibul modulus (m~(15;30))and anisotropic electric conductivity were obtained for the layered material, their application in the industrial scale is still hindered due to the low reliability. From this reason a selfmonitoring layer was joined to the layered composite, as it is schematically shown in Fig. 7, layer STN-X.
I
SN-0
1
I I
SN-0
1
Fig. 8. Change of electrical resistivity p with TiN content
for conductance is 20.9 vol%, if the particles are spherical [16]. The alignment of TiN particles was increased also by the presence of elongated Si3N4 grains. The electrical resistance of samples with 22 and 23 vol% TiN is little bit higher in comparison to the samples with lower TiN content. In these samples only 18 vol% TiN was prepared by in situ reaction, the remaining part was directly added as a powder. From this reason the distribution of TiN particles was not so homogeneous as in the samples with lower TiN content. On the base of these results the composition pSiAlON+19%TiN was selected for the self-monitoring layer. This functional layer will be between two Si3N4 layers in practical use (Fig. 7, X=19). The outer layer is under compressive stress and the inner conductive STNX under tensile stress. The level of residual tensile stresses is a very important factor, because exceeding certain level (-300 IvlPa) crack formation can be observed in direction normal to the layer interface, as it is shown in Fig. 9.
I
Fig. 7. Schematic of FGM material with self-monitoring STN-X layer (SN-0: pure Si3N4layer).
The first step was the optimisation of the TiN volume fraction in the self-monitoring layer. The variation of electrical resistivity with TiN content is shown in Fig. 8. This chart shows that the threshold TiN volume fraction is about 19 ~01%.These results suggest that the in situ formed TiN particles are well dispersed and aligned, because the critical volumetric fiaction calculated on the base of percolation theory
562
Fig. 9. SEM micrograph of crack formation normal to the layer alignment due to the presence of high residual tensile stresses.
The residual tensile stresses in three-layered composite can be calculated according to the following equation [ 171:
where v is the Poisson’s ratio, E the Young’s modulus, d the layer thickness, AT the temperature difference, /?, and fi parameters include the coefficients of thermal expansion (CTE) and shrinkage of adjacent layers. Except of layer thickness the other parameters are fixed with the composition of layers. The calculations showed (Eq. 2) that the ratio of innerlouter layers should be about 4 to have residual tensile stresses in the inner conductive layer less than 150 MPa. The thickness of tape casted ceramic sheets was 100 pm and fiom this reason the thickness of three-layered testing bars was 600 pm (100-400-100). The results of electrical resistivity measurements are shown in Fig. 8 by uptriangle symbols. The resistance of layered materials decreased in comparison to bulk materials, what was unexpected and is not clear yet. With Vickers indentation tensile stresses were observed in the middle layer, because the cracks normal to the layer interface were by 15% longer in comparison to the parallel cracks. Probably also another method, e.g. piezospectroscopic determination of residual stresses should be used [18]. Partly clarify these results additional measurements were conducted on bulk bars with 19 vol% TiN under load. The testing bars were positioned on the 3-point bend test holder and the load was increased stepwise. The results of electrical resistivity measurements under tensile load are shown in Fig. 10.
c
to zero (dp/dF = 0 ) means that the probability of crack formation increased to the serious level. In the practical use it means that the ceramic component should be changed. The results showed that there is a simple but useful tool for the self-diagnostic of brittle ceramic materials, although additional measurements under continuous or cyclic load are necessary. CONCLUSIONS Dense Si3NJ(p-SiA10N + TiN) layered composites were prepared by tape casting and reactive hot pressing. The bending strength and fracture toughness of layered L4 and FGM materials were remarkably higher in comparison to the conventional P-SiAlON+TiN composite. Anisotropic electrical properties were obtained for the layered materials, if the TiN content exceeded 19 ~01%.The experimental results confirmed that the layered material design is suitable for the preparation of multifunctional ceramic materials. Layered material with possible self-diagnostic ability was also designed. The monitoring layer contained 19 vol% TiN and the change its electrical resistance was rather sensitive to the applied load. The results suggest a possibility for producing layered ceramic materials with “health-monitoring” function and prevention against fatal cracks. ACKNOWLEDGMENT Work was supported by the Slovak Grant Agency VEGA, contract No. 2/5118/99 and by the bilateral project between SAS Slovakia and CNR Italy, No. 1 1/1.
3,26--1
3.24
.818
REFERENCES
8 8
‘9 3.18 8
3.16 8
.
0
20
3.14
1
.
40
60
80
100
120
Load I N
Fig. 10. Change of electrical resistivity of STN-19sample with applied load.
At the beginning the electrical resistivity slightly decreased and fiom 60 N load it started to increase. The increase of electrical resistivity can be explained by the formation of microcracks in the conductive layer and by the reduction of the conductive paths (cracks between TiN particles). The change of electrical resistance was in the order of 0.1 Qcm, which can be recorded with available measuring devices. The course of extrapolated curve on the measured points with local minimum is advantageous for the self-diagnostic function. The output of measuring device can be simply recorded and if the derivation of electrical resistance by load is equal
Harmer, M.P., Chan, H.M. and Miller, G.A., Unique opportunities for microstructural engineering with duplex and laminar ceramic composites. J. Am. Ceram. SOC., 1992, 75, 1715-28. Sajgalik, P., LenWS, Z. and Dusza, J., Si;N, based composite with layered microstructure. In 5’ Int. Symp. Ceramic Materials and Composites for Engines, ed. D.S. Yan, X.R. Fu and S.X. Shi, World Scientific, Singapore-New Jersey-London-Hong Kong, 1995, pp. 198-201. Sajgalik, P., LenECS, Z., Silicon nitride layered composites - relationship between residual stresses and fiacture behaviour. In 6‘h Ceramic Materials and Components for Engines, ed. K. Niihara et. al. Technoplaza Co., Japan, 1998, pp. 675-78. Ishida A., Miyayama M., and Yanagida H., Prediction of fracture and detection of fatigue in ceramic composites fiom electrical resistivity measurements, J. Am. Ceram. SOC.,77 (1994) 1057-1061. Nakamura M., Saitoh K., Ikeyama M., Kozuka T. and Shigematsu I., A study of the utilization
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of an electroconductive surface layer on a Si3N4 ceramic formed by ion implantation as a crack gage and a stain gage, J. Ceram. SOC. Jpn., 101 (1993) 139-142. Akiyama M., Nonaka K., Watanabe T. and Miyazaki K., Self-detection of fi-acture in PZT and MoSi2-Mo2B5layered composite ceramics, J. Ceram. SOC.Jpn., 102 (1994) 686-687. Jack, K.H., Review: SiAlON and related nitrogen ceramics. J. Mater. Sci., 11 (1976) 1135-58. Huang, J.L., Chou, F.C. and Lu, H.H., Investigation of Si3N4-TiN/Si3N4-Si3N4 trilayer composites with residual surface compression. J. Mater. Res., 1997, 12(9), 2357-65. Huang, J.L., Lee, M.T., Lu, H.H. and Lii, D.F., Microstructure, fi-acture behaviour and mechanical properties of TiN/Si3N4composites. Mater. Chem. Physics, 1996,45,203-10. Bellosi, A., Guicciardi, S. and Tampieri, A., of Development and characterization electroconductive Si3N4-TiN composites. J. Eur. Ceram. SOC.,1992,9,83-93. LenECS, Z., Hirao, K., Brito, M.E., Toriyama, M., Kanzaki, S., Si3N4-based layered composites with discontinuous BN interlayer, accepted to the J. Am. Ceram. SOC.,1999. LenECS, Z., Sajgalik, P., Roncari, E., Hirao, K., Design of Si3N4Based Layered Composites for Multifimctional Application, Key Engineering Materials, 175-176 (2000) 173-182.
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MULTILAYER CBIC COMPOSITES FOR AUTOMOTIVE BRAKE SYSTEMS R. Gadow, M. Speicher University of Stuttgart Institute for Manufacturing Technologies of Ceramic Components and Composites Allmandring 5b D-70569 Stuttgart GERMANY
ABSTRACT
ingredients into a sheet product that is usually about 4 mm thick and 1.500 mm wide. The SMC process is illustrated in Fig. 1.
Reaction bonded carbon fiber reinforced Sicceramics are promising light weight materials for tribological applications('). The current state of the art concerning the manufacturing technology is the reaction bonding process via silicon infiltration in carbon preforms'233'.Designing the porosity of these preforms is a possibility to control and optimize the Investigations on different designed resin derived and carbonized preforms have been carried out by characterizing pore volume, pore size distribution and melt infiltration experiments. The microstructural morphology of the preforms was modified by mixing ceramic fillers in the applied resin system thus varying the shrinkage and pore formation during pyrolysis. Using mercury porosimetry the influence of different fillers on the morphology was studied. The results describe the dependence of filler type and content on the porosity and evidence that not only the pore volume fraction but also the pore size distribution influences the melt infiltration. To produce carbon fiber reinforced preforms the optimized resin-powder mixture is applied in a SMC (Sheet Moulding Compound) manufacturing process. This manufacturing technology offers a wide range of possibilities to combine short and continuous fibers with matrix material. The fiber length and content influence not only the mechanical and tribological properties but also the porosity and in conclusion the silicon infiltration during the reaction bonding process.
MANUFACTURING TECHNOLOGY Generally, the manufacture of fiber reinforced preforms involves two stages. Incorporation of fibers into the unconsolidated matrix material, followed by consolidation of the matrix'@. The current state of the art are the RTM (Resin Transfer Moulding), the liquid phase infiltration of fibers and compounding of short fibers"'. The SMC (Sheet Moulding Compound) is transformed from its liquid, fiber and powder filler
fig. 1
SMC manufacturing process (on top) computer controlled impregnation process
All ingredients, except the fibers, are mixed together to form a resin paste, which is transferred to a
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doctor box where it is deposited on a moving carrier film. The doctor box controls the amount of resin paste applied. Simultaneously,the carbon fiber rovings are fed into a rotary cutter above the resin covered carrier film. The fibers are chopped to length from 12-25 mm and dropped onto the resin paste. The deposition of fibers is random, but generally oriented slightly parallel to the direction of the film travel. The amount of fibers is controlled by the speed of the carrier film. A second carrier film is coated with resin paste and is laid resin side down on the top of the chopped fibers (fig. 1). This stage of the process creates a resin paste and carbon fiber 'sandwich' which is subsequently sent through a series of compaction rollers where the carbon fibers are wetted with the resin paste and trapped excess air is squeezed out of the sheet. At the end of the compaction rollers, the SMC sheets are taken up on a storage roll(*). Endless fibers can be. obtained as well by SMC technology if no cutter blade is used. The viscosity of SMC resin paste in the doctor box is very low. Before the SMC can be used for molding it must age or mature. This maturation time is required to allow the relatively low-viscosity resin to thicken chemically, so it is easier to handle and more appropriate in industrial forming, especially if large shape components are required Before molding the SMC sheets have to be cut into pieces of predetermined size and shape. The cut pieces are subsequently laminated and assembled into a charge pattern. The combination of short and long fiber reinforced sheets opens up the possibility to realize a great spectrum of suitable structures for different applications. The charge is placed in a preheated mold and compressed. Typical mold pressures are in the range of 150-250 bar. Under influence of heat and pressure the SMC is transformed from its leather-like quality to a very low-viscosity liquid behavior. So net-shape forming process is possible. After this forming and curing process the component is pyrolized under nitrogen atmosphere and then siliconized under vacuum. During both processes the shape of the component is mostly unchanged, so that the finishing effort can be minimized.
properties the infiltration can only be optimized by an appropriate morphological design. To realize a dense, oxidation resistant CMC component it is necessary to influence the pore volume and pore radii distribution obtained during pyrolysis through selective addition of fillers with well defined grain size, specific surface and chemical reactivity. An increasing powder filler content (Sic) leads to an isotropic decrease of shrinking factor which results in an increasing pore volume('). This increasing pore volume is based on the enlargement of pore radii (fig 2). A part of the former small pores is maintained, so a bimodal pore size distribution is obtained. 30 phenolc resin 9691 MI (Bakeld)
25
z
20
--4Owt%Sic03
15
I
fig. 2
'
I
, t
pore radii distribution depending on S i c filler content
The measurements by mercury porosimetry show no linear dependence of the pore volume on the filler volume content (fig 3). -c bulk density
--- calcul.density
50
+- app. density
40a a 30 00.
u"
20z
MORPHOLOGICAL STUDIES The complete transformation of a carbon fiber reinforced carbon preform to a reaction bonded SiCCMC material is based on two simultaneously competing processes: the silicon melt volume diffusion and the heterogeneous chemical reaction between silicon and carbon. For a large component with considerable wall thickness a controlled and effective silicon infiltration is of greatest importance. With respect to the requirement of realizing a complete chemical conversion (ceramization), a time and cost effective manufacturing the required final material
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10
fig.3
20
30
40
50
70
porosity and density depending on S i c filler content A strong rise of the porosity occurs below 30 wt. % Sic, due to the limitation that only open pores can be determined by mercury porosimetry. Beginning at 30 wt. % S i c filler, the open and the closed porosity are joined and therefore the complete pore volume is
measured. The addition of 40 wt % S i c leads to a homogeneously distributed porosity and above 60 wt. % a small, bimodal pore radii distribution is obtained (fig. 2 and 4). In any case the porosity strongly depends on the shrinking behavior of the applied resin during wwo
pyrolysis and, as mentioned above, on the particle size distribution of the powder filler. In conclusion it is possible to control the porosity via powder filler volume content. However the amount of filler is limited by the viscosity of the resin paste.
sic
10
fig 4 pyrolyzed preform morphology using various S i c powder content left: overview; right: detail pyrolyzed carbon matrix (grey), Sic-filler (white), porosity (dark)
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An increase in fiber length (25 mm) and content (44vol. %) leads to lower porosities and a plane pore radii distribution (fig 6).
MELT INFILTRATION EXPERIMENTS The main goal of this porosity optimization is to influence the melt infiltration during the reaction bonding process. A compounding with a high amount of filler, a minimum binder content and only 29 vol. % carbon short fibers (3mm) leads to a high porosity of the preform with a bimodal pore size distribution (fig. 5). Small pores guarantee the capillary force and wide pores provide a rapid silicon mass flow. In this case silicon melt infiltration leads to a dense ceramic body keeping constant shape of the porous green body"'.
30 25
180 Total cumulativevolume (mnP/g): 166,88 Total porosity (X): 24,s Bulk density (glcw): i,46 Apparent density (g/cnP): 1,93
I 1 *
n
20
15 C-fiben (25mm) 44 ~ 0 1 %
i a
10-
160 140 0
120
5
/
5-
4M3
0-
350
2 3 250 ; 300
Gfiben(3mm) 29~01%
J
I .
pore radii [A]
0
200 7 150
- Spec.Vol. O I Og. u
50
0 0 0 0 @ @ @
8 86, &I 49
@, 8 80 44 % 4
0 0 0 6% 9 b#'?
8
26
% '
"
+J
pore radii [A]
fig 6 pore radii distribution of the preform and microstructure of CMC siliconized at 1600 "C, 5 h
fig. 5 pore radii distribution of the preform and microstructure of CMC, siliconized at 1600 "C,
30 min
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The use of UD-fibers with a high volume content results in a low open porosity of the preform with a similar plane pore radii distribution as well, but in addition a high amount of small pores in the range of 10-100 A is detected (fig 7) due to the alignment of long fibers. These small pores are difficult to infiltrate and therefore a higher closed porosity and free carbon content of the siliconized ceramic is obtained.
30
~
25 c1
20
~
E a> 1 5 -
p
140
Preform (pyro~ped)
Total cumulative volume (mmVg): 129.23 Total porosity (X): 18.78 Bulk density (g/cm3): 1 53 Apparent densty (glcm’): 1.90 C-fiben(UD) 50~01%
io-
fig 7 pore radii distribution of the preform and microstructure of CMC silconized at 1600 “C for 10 h
SUMMARY AND DISCUSSION The morphology of this special SMC is strongly influenced by its formulation, concerning the type of phenolic resin, additives and the amount of carbon fibers and powder fillers. High contents of long, unidirectional (UD) fibers result in maximum reinforcement in the required direction, but the infiltration process is complicated due to the resulting low porosity of the preform. The remaining carbon content is high. Consequently the oxidation stability is limited. The use of insulated short fibers with lower volume content and simultaneously higher volume fraction of powder fillers in the resin mixture promise a dense ceramic body with a high corrosion stability. But the mechanical properties are currently insufficient, especially the fracture toughness due to a high residual silicon content. The SMC technology offers a wide range of possibilities to combine various fiber length and filler contents. By enveloping UD-fiber layers with short fibers at the surface a design was realized which combines the advantages of the short fiber CMC and the UD-fiber reinforced component (fig 8). The inner UDlayers guarantee fracture toughness and the short fiber layers at the outside provide corrosion resistance and, for tribological applications, a steady friction coefficient combined with minimum wear This opens up the possibility for a ‘life time’ brake disk. Short and reproducible production cycles as well as a net shape forming process, due to the high mould filling capacity of the prepreg sheets, can be realized by the SMC compounding process. The fulfillment of the automotive industry requirements concerning cost effectiveness and reproducibility opens up the possibility for a serial manufacturing method for ceramic brake disks.
fig 8 structural design of high performance friction material left: cross section schematic cross section of infiltrated 280 mm brake disk right:
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REFERENCES R. Gadow, M. Speicher: Herstellung faserverstiirkter, reaktionsgebundener Siliciumcarbidkeramiken unter Verwendung intermetallischer Siliciumlegierungen, Mat.-wiss. u. Werkstofftech. 30, No. 8, pp. 480-486 (1999), WILEY-VCH Verlag C. W. Forrest et al.: Special Chemics 5, pp.99-123 ( 1970) W. B. Hillig et al.: GEC Tech. Inform. Ser. 74 RD 182 (1974) M. Singh, D. R. Behrendt: Microstructure and mechanical properties of reaction-formed silicon (RFSC) ceramics, Materials Science and Engineering, A1 87, pp. 183-1 87 Elsevier Sequoia (19941, R. Gadow, M. Speicher: Manufacturing and CMCComponent development for brake disks in automotive applications, Ceramic Engineering and Science Proceedings, Vol. 20, No. 4, The American Ceramic Society, 1999, ISSN 019662 19, pp. 55 1-558 D.C. Phillips: Fibre reinforced ceramics, Handbook of Composites, Vol. 4, ed. by A. Kelly and S.T. Mileiko, Elsevier Science Publishers B.V., 1993, ISBN 0 444 864474 Gadow, R.; Kienzle, A.: Processing and Manufacturing of C-Fibre Reinforced SiCComposites for Disk Brakes, Proc. 6" Int. Symp. On Ceramic Mat. and Components for Engines, Arita, Japan, K. Niihara et al. eds., pp. 412-418 (1997), ISBN 4-9980630-0-6 SMC/BMC - Design for Success!, European Alliance for SMC, WDW Werbedruck Winter ( 1997) R. Gadow: Dissertation, University of Karlsruhe (1 986)
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SINTER ADDITIVE OPTIMIZATION IN PROCESSING OF ALUMINUM NITRIDE FOR HEAT EXCHANGER COMPONENTS R. Fischer*, R. Gadow and M. Lischka University of Stuttgart, Institute for Manufacuring Technologies of Ceramic Components and Composites, Stuttgart, Germany
ABSTRACT AIN ceramics are known for their wide range of applications in electronics mainly as substrate materials for semiconductor ICs in computer applications, optical components and high performance heat exchanger in electrical devices. There is a potential for use in hot gas I water heat exchangers because of the advantageous thermophysical properties of AlN. Conventional domestic heat exchangers made of steel and copper (widely used in thermal value heating systems) are less efficient and less reliable than polycrystalline AlN heat exchangers due to their long term corrosion behaviour problems based on dilute acid condensation products from the vapour phase. AIN as a polycrystalline ceramic material has superior thermal conductivity properties with room temperature values up to 250 W/mK. Such high performance is not necessarily required for a commercial product, but allows up to a certain extent the variation of sinter additive phases, enabling to design the best combination of mechanical and corrosion resistance properties in combination with a high thermal conductivity. The influence of different additives on the sintering process and the corrosion resistance was investigated for injection molded components. The optimization was verified analyzing the physical properties, the oxidation resistance and the crystal optical analysis of ceramic thin sections. The optical properties of the edge zone of the AlN crystal depend on their chemical properties. Thus the optical properties are a promising measure for the efficiency of the additives.
INTRODUCTION Aluminum nitride A1N ceramics are known for their wide range of applications in electronics and electrical devices, mainly as substrate materials for semiconductor IC in computer technology [I], optical components and high performance heat exchangers, e.g. thyristor coolings in electrical traffic technology, because of their advantageous thermophysical properties (Tab. 1) and their thermal expansion coefficient matching with silicon.
Table 1: material properties of A1N ceramics [2] material properties density p [g/rm']
thermal expansion coefficient CXK [106Kq electric resistivity [h] at RT thermal conductivity h [WhnKl e- modulus E [kN/nd] fracturetoughness I
bending strength 06 [~/mm']
3.81 k0.02 5.6 (RT-lO00"C) > 10l2 140 - 250 310 3.35 f 0.2 350 f40
Conventional heat exchangers made of steel and copper widely used in high performance thermal value heating systems show corrosion problems caused by condensation products of low temperature exhaust gases. An alternative is an all ceramic heat exchanger, based on aluminum nitride (AIN). A1N is a polycrystalline ceramic material, which has superior thermal conductivity properties with room temperature values up to 250 W/mK. This high value is not necessarily required for a commercial product. A1N allows by variation of sinter additives the design of the best combination of mechanical, thermophysical and corrosion resistance parameters. The conventional manufacturing method for AIN components is axial or cold isostatic pressing. The press forming method requires an additional green body machining operation step, which is expensive. The variation of the component design is limited on less complex shapes. The fast and economic ceramic injection molding method enables to produce complex components in industrial quantities with near net shape and high surface quality. The CIM production of such heat exchangers requires an efficient and optimized manufacturing process to achieve the desired net shape tolerances and reproducable high surface quality. For serial CIM production, the knowledge of the sinter processings is an essential. The effect of three sintering additives, CaO (1 wt.% due to chemical incompatibility to AIN in higher concentrations), Ce@ (3 wt.%) and Y 2O (3 wt.o/) on the chemical properties of the AIN materialwas investigated and analyzed with the crystal optical analysis of thin sections of the samples.
57 1
AlN COMPOUNDING A1N ceramics are most commonly manufactured via axial and cold isostatic pressing. The expected advantage of injection molding is based on the high efficiency, process stability and materials economy resulting in low specific cost for mass production of an advanced ceramic component with complex net shape and performing surface quality obtained by this CIM processing [4]. For injection molding of ceramic components, the knowledge of the rheological behaviour [ 5 ] of the compounds is one of the most important requirements to obtain components without defects. The flow behaviour depends on the grain size distribution, grain fraction mix, grain morphology and the binder formulation. Low flow resistance due to optimized viscosity leads to a good form filling behaviour, no warping and allows netshape tolerances. It leads furthermore to reliable mechanical properties, low mold abrasion and supports the optimization of machine parameters (p, T, v). Depending on the rheological behaviour of the compound, a highly performing and stable processing can be achieved.
debinding method is very fast and gentle to the green body. The sintering process works at very high temperatures up to 1900°C and in the case of A1N under inert gas atmosphere, Ar or N with optional addition of NH3.
1-
I
__
--
CanFOt cracks pores
A special ceramic feedstock together with a well defined and optimized rheological behaviour has been developed [3]. The optimized feedstock (78 wt.% AlN, 22% BASF binder and additives) recipe is made by comilling or mixing sinter additive oxides (CaO, Y 20 and Ce203)for superior corrosion resistance properties together with commercial AlN powder with monomodal grain size distribution. After mixing with organic binders and plastifiers in a heatable doublesigma-blade header, the compounding is performed by continous extrusion and rotation cutting to pellets.
RHEOLOGY AND INJECTION MOLDING The CIM process (Fig.1) is related to the polymer and metal injection molding but has some special features. Due to the high powder filler content of the compound a reliable compoundation (about 80 wt.%), homogenisation is very important to achieve defect fiee green body components with high density and performing mechanical properties. The abrasive properties of the ceramic powder require a wear resistant design of the injection molding machine, especially the plastification unit and the injection nozzle. Highly wear resistant and polished steel is also required for the mold. Before sintering, the temporary organic binder has to be removed with a well controlled pyrolysis or chemical degradation. Otherwise the component will be destroyed during the sintering process. Especially for the CIM and MIM process the BASF catalytical debinding method is applied. To remove the binder, mineral or organic acid cracks the polymer binder into gaseous monomers, which are combusted in an afterburning process. This
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I---.(
Figure 1: CIM process flow chart
DEBINDED AND SINTERED AlN COMPONENT
-
The AlN green bodies were catalyticaly debinded with the BASF / IFKB method (6 hours, 125"C, NZatmosphere and organic acid ( 5 vol.% oxalic acid, 55 vol.% acetic acid, 40 vol.% formic acid) for a smooth debinding of the non-oxide AlN powder) and sintered in a graphite kiln at 1850°C with three hours dwell in N2 atmosphere (1 bar). The oxide dope additives of the compound lead to a superior densification due to a reduction of the oxygen concentration on the AlN powder surface and accelerates the densification during the sinter process [6].The liquid phase coating also improves the corrosion resistance due to the superior chemical resistivity than the undoped AlN. The liquid sintering process is shown in Fig. 2. AlN is soluted in the solution range A by the the liquid phase and transported by diffusion to the range of enrichment B and deposited. This process leads to the densification of the material [7]. Due to the complex sintering behaviour of AlN an improved sintering process was developed. The sintering curve is shown in Fig.3. The SEM-micrographs of the sintered body with Ce02 sintering additive (Fig. 4 and 5) show a typical ceramic AlN sinter texture with some defects which are caused by debinding problems. The specimen with this
additive develops the best texture. Fig. 2 shows the typically liquid sintering phase, which coates the AlN powder and leads to a densification of the material. The pore volume is 5 %. The samples with CaO and YO3 additives show inferior results. First long time corrosion tests (30 days in concentrated burner condensate) showed an explicite improvement in corrosion resistance . Further tests under working conditions will be done in the near future.
Figure 5: SEM micrograph of a typical sintered CIM specimen with debinding caused defects AlN specimen from CIM processing
OPTICAL CHARACTERIZATION WITH POLARIZED LIGHT MICROSCOPY Figure 2: Liquid phase sintering mechanism of A1N (liquid phase: black; AlN grain: grey); A: range of solution; B: range of enrichment
r *.Oh
! ,.Oh
! 2.0h
! ,,Oh
! .Lh
! %Oh
! 'Lh
mgM1
!
ILh
! *Oh
,.Oh
Figure 3: Sinter program for A1N components
Figure 4: SEM micrograph of a sintered AlN
!
r
ro..n
Due to its crystal structure, AlN shows optical birefrigeance. The polarized light microscopy is a method to make birefrigeance visible and measureable. For the examination of thin sections of the samples, the thickness is reduced to 8 - 10 p m [8, 91. The observed interferece colour depends on the chemistry and concentration of the sintering additive. The higher the birefrigeance, the more oxygen diffusion into the AlN structure takes place. The best thermal conductivity has the pure AlN. With increasing oxygen impurities the thermal conductivity decreases but the corrosion resistance increases. The result of our examinations is that the value of birefrigeance gives information about the quality of thermal conductive components (Fig. 6 , 7 and 8). The higher the birefrigeance, the lower the thermal conductivity.
Figure 6: AlN sample with 1 wt.% CaO . Interference colours: orange 1S t order.
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Accurate process monitoring via simultaneious process control (SPC) of the flow behaviour during injection and form filling results in an increase of the quality and netshape size of the sintered ceramic component with low texture induced damages and allows to decrease the production cost together with higher lifetime of the mold. This results allow in a first step a small scale manufacturing of heat exchanger components.
Figure 7: A1N sample with 3 wt.% C e Q . Interference colours: Yellow 1St order
Figure 9: Conventional steel/copper domestic heat exchanger (courtesy of Bosch-Junkers, Germany)
Figure 8: AlN sample with 3 wt.% additives. Interference colours: Blue 2"dorder The C e G additive shows a moderate ogygen concentration in the sinter phase with a interference colour of yellow first order. CaO with lower concentration shows higher interference colours, YO3 shows blue second order and the lowest thermal conductivity
Figure 10: Form filling study for the new designed ribbed AlN heat exchanger component, demonstrated with a fan rotor as reference shape
COMPONENT DESIGN The shape of the sintered component is checked by using a coordinate measurement machine allowing to verify the dimensional stability after debinding and sintering procedure. Due to the advantageous thermophysical properties of AlN, the heat exchanger elements can be designed more compact or smaller than conventional models (Fig.9 and 11). A segmented ribbed design (Fig.10 and 12) enables to achieve an efficient convective flow of the hot burner exhaust gas around the exchanger. The design of a pressed thyristor cooler (Fig. 11) is less effective for the use in domestic heat exchangers. Because of the completely different temperature gradients, the flow rate on the water side is inferior to the ribbed design. The production cost for this design is significantly higher because the required amount of ceramic powder material is increased at about 50%.
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Figure 11: AlN thyristor heat exchanger, acc. ANCeram, Bayreuth, Germany [2]
$rib
[4]
pas side
Figure 12: Schematic illustration of a ribbed A1N heat exchanger component
[5]
[6] [7]
CONCLUSION Due to its advantageous thermophysical properties, AlN has been recently established. The material is known in a wide range of applications in electronics, computer applications and optical components. A new application is a high performance hot gas / water heat exchanger with superior performance for domestic thermal value heating systems up to 15 kW capacity, which can substitute conventional steel/copper heat exchangers. AlN shows superior thermal conductivity and an improved long term corrosion behaviour. A first compact design for a ribbed heat exchanger was developed. The CIM manufacturing of AlN heat exchangers was shown to be competitive for a production in near net shape, high surface quality and for large lot sizes. The CIM process must be optimized for ceramic requirements in regard to molding parameters and mold design.
[8]
[9]
engineering and science Droceedines 10 [4] (1999), pp.595-602 Gadow, R., SuBmuth, G., SpritzgieBen in der technischen Keramik, Technische Keramik, 1988, pp.206-2 10, Fritz, H.-G., Einfuhrung in die Rheologie und Rheometrie der Kunststofle, Technische Akademie Esslingen, Esslingen, 1996. German R. M. (1996): Sintering Theory And Practice.- J. Wiley & Sons, inc. Hundere, A. M. (1995): Sintering Of Aluminium Nitride (AlN) Ceramics.- Institutt for uorganisk Kjemi Norges Tekniske Hargskole, Universitetet i Trondheim. Fischer, R., Franz, E.-D. und Telle, R. (1992): Gefigeuntersuchungen an SpritzguRkeramiken mit der Methode der DurchlichtPolarisations-Mikroskopie. Sprechsaal fir Betriebsmanagement & Technologie International Magazine Traditio-nal and Advanced Ceramics ,Glass and Materials Technologie, 125. (1992) ; 12 S. 797-804. Wahlstrom, E. E. (1979): Optical Crystallography.- John Wily and Sons, New York.
The quality of CIM formed ceramic components depends on the rheological properties of the feedstock and the manufacturing parameters like flow resistance, form filling behaviour and deforming / ejection. Netshape tolerances, reproducable high surface quality and minimized mold abrasion are the result of optimized processing parameters. The chemistry of the sintering additives is very important for the components quality. C e Q seems to be the best additive for optimized AlN components. Due to the minmized oxygen exchange with the A1N structure, the expected thermal conductivity and the corrosion resistance is superior. Polarized light microsopy confirms this fact.
REFERENCES [ 11
[2] [3]
Werdecker, W., Aldinger, F., Aluminum nitride- an alternative ceramic substrate for high power applications in microcircuits, IEEE Trans. CornDon.. Hvbrids and Manuf. Technoloey, 7 [4] 1984, pp. 399-404 ANCeram, Product Information, ANCeram GmbH & Co., Germany, 1996 Fischer, R., Gadow, R., Schafer, G., Fabrication of AlN Heat exchangers by ceramic injection molding, Ceramiq
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NEW OPPORTUNITY FOR BIMODAL MICROSTRUCTURE CONTROL IN SILICON NITRIDE Hai-Doo Kim, Byung-Dong Han, Dong-Soo Park, Dong-Ho Han*, and 11-So0 Kim* Ceramic Materials Group, Korea Institute of Machinery and Materials, 66 Sangnam-Dong, Changwon, Kyungnam, 641-010, Korea * Dept. of Advanced Materials Eng., Dong-eui Univ., San 24, Kaya-Dong, Pusanjin-Ku, Pusan, 614-714, Korea
ABSTRACT Two step sintering process comprising of carbothermal reduction process of the compacts of Si3N4, Y2O3, A1203, and C mixture and the following gas pressure sintering process was investigated in order to produce the bimodal microstructure in silicon nitride ceramics. Carbothermal reduction treatment of powder compacts of silicon nitride, sintering additives and carbon mixture takes the oxygen out of Y-Si-Al-0-N glass, which leads to the precipitation of crystalline oxynitride phase such as apatite due to the increase in N/O ratio in the glass. Precipitation of apatite phase from Y-Si-A1-0-N glass leads to the decrease in glass content. Precipitation of apatite phase and disappearance due to its melting during sintering process result in the retardation of densification process up to 1750“c , followed by the rapid increase in density due to the recovery of liquid content, which provides the situation to maximize the difference between the growing grains and the matrix grains, which leads to the distinct bimodal microstructure. Since grain growth and densification processes are thermally activated processes, holding the densification process to a high temperature followed by the reactivation of sintering process promotes the development of bimodal microstructure. An analogous ‘transient solid phase sintering’ induced by the carbothermal reduction treatment of powder compacts of silicon nitride, sintering additives, and carbon mixture may provide a new opportunity to control the microstructureinto bimodal way.
INTRODUCTION One of the methods to improve the strength and h t u r e toughness of silicon nitride is to control the microstructure into bimodal way, in which the large elongated grains in a fine matrix act as reinforcing grains, mostly, by the crack bridging mechanism(1-2). In order to have an extreme bimodal microstructure, fine matrix grains are prerequisite, which can be obtained from fine starting a-Si3N4 particles(3-5). Intensive milling is an effective method for reducing the particle size. However, it might give rise to an increased oxygen content due to the increase in surface silica content, which affects both phase development and microstructures(6). It is, therefore, necessary to reduce the
oxygen content while keeping the milled a-Si3N4particles as fine as possible. Possible ways to reduce the oxygen content after an intensive milling are chemical treatment such as HF rinse(7) or carbothermal reduction treatment(& 14). Most of the works on carbothermal reduction process are concerned with the reduction of SiOz layer on Si3N4 particles(8-14). However, when the carbon is added to Si3N4powder mixture with sintering additives, the result would be different from the above case, in that the carbon may take oxygen not only from Si02 but also fiom oxide glass or oxynitride glass generated from Si3N4(+Si02)and the sintering additives because the carbothermal reduction temperature employed in most cases(8-14) is 1450°C or higher, which is well above the eutectic temperature(1370C ) in the system Y203-A1203-Si02( 15) and, at 1450C, a- to p-Si3N4 phase transformation is possible(6, 16). When the carbothermal reduction treatment of the powder compacts consisting of silicon nitride, sintering additives and carbon takes the oxygen out of the glass formed, this changes the N/O ratio in oxynitride glassy phase during sintering compared to the general case without carbothermal reduction treatment. Since the oxynitride glasses with different N/O ratio show different properties such as viscosity, glass transition temperature, etc.(l7) and might result in a different precipitation behaviour due to the limited solubility of nitrogen in oxynitride glass(l8), this would leads to a different sintering behaviour and , thus, different microstructural development. In this work, the carbothermal reduction process of the compacts of Si3N4, Y2O3, Al2O3, and C mixtures and the following gas pressure sintering process were investigated. Special attention was paid to the second phases precipitated after carbothermal reduction process, the phase development with increasing temperature and the microstructural development.
EXPERIMENTAL PROCEDURE 2 batches of 93 wt?! a-Si3N4 (E-10, Ube Industries Ltd., Tokyo, Japan), 6 wt% Y203 (Fine, H. C. Starck Co., Berlin, Germany), and 1 wt% A1203 (AKP-30, Sumitomo Chemical Co., Osaka, Japan) with and without 0.6 wt% carbon(Activated carbon, Aldrich Chemicals Co.,
577
Milwaukee, WI) were prepared by planetary ball milling for 20 h using Si3N4 balls of 5 mm in diameter and methanol in a Si3N4jar. Hereafter, these batches are named as C-0 and C-0.6. Slurries were dried on a hot plate while stirring, sieved to -100+150 mesh, pressed into compacts of 30 mm in diameter by CIP at 250MPa The compacts were carbothermally processed in a tube furnace at 1450°C for 10 h in a flowing N2 atmosphere(0.2 literlmin). In order to examine the phase development, the compacts after carbothermal reduction treatment were sintered at 1550 to 1850°C for 0.5 h under 1 MPa N2 pressure by 100°C interval. Compacts were also sintered at 1850°C for 3 h and 6 h under 2 MPaN2 pressure to see the densification and microstructural development. After the planetary ball milling, particle size distribution and specific surface area were measured by laser diffraction method (LS-130, Coulter Co., Hialeah, FL) and BET method (ASAP-2010, Micromeritics, Norcross, GA), respectively. Oxygen and carbon contents before and after carbothermal reduction treatment were measured using Oxygen Analyzer (TC-136, LECO Co., St. Joseph, MI) and Carbon Analyzer (CS-344, LECO Co.). Density was measured by Archimedes method. Phase development was investigated using X-ray d i h t o m e t e r (Rigaku Co., Tokyo, Japan) on the cross section of the specimen. CuKa radiation, 35kV-25mA condition and O.O67O/s scan speed were employed for this investigation. The d p ratio was calculated using Gazzara's equation.(19). The cross section of the specimen was polished to 1 pn with diamond paste and plasma-etched (96 % CF4 + 4 % 02) for microstructural analysis using SEM (JSM-5800, Jeol, Tokyo, Japan).
calculation considering that oxygen content in a-Si3N4 powder is 1.16 wt Yo and that in the sintering additives is 1.74 wt %. Oxygen content of the same batch after 4 h planetiuy ball mill is 3.5 wt% according to the measurement using the oxygen analyzer. The oxygen content of 2.9 - 3.5 wt % seems, therefore, to be a reference value because the oxygen contents of the batches which were prepared without any special treatments such as a prolonged milling or an oxidation may fall into this range. Oxygen content of C-0 batch after 20 h planetary ball mill is 4.0 wt% and that of C-0.6 batch is 4.2 wt%, while those after carbothermal reduction treatment are 4.0 wt% and 3.6 wt%, respectively. The remaining carbon content after carbothermal reduction is in 0.1-0.14 wt% which is comparable to the carbon content of a-Si3N4 starting powder(O.1 wt%). Fig. 1 shows the XRD patterns of C-0.6 with increasing temperature. Just after the carbothermal reduction treatment, a-Si3N4as major peaks and p-Si3N4and apatite as minor phases. Appearance of apatite phase suggests that oxygen-deficient glass resulting from the carbothermal reduction treatment leads to the Y-Si-Al-0-N glassy phase with high N/O ratio. Increase in N10 ratio of the oxynitride glass gives rise to the precipitation of apatite phase. With increasing temperature, a to p phase transformation occurs with apatite second phase until 1650°C.At 1750"C,a to p phase transformation is complete and apatite phase disappeared. The melting point of apatite phase is known as 1750"C(20) or 1700-1800"C(21), which suggests that the disappearance of apatite phase is due to melting. The precipitation of apatite phase is expected from the pseudoquaternary phase equilibrium diagram for the system Si-Y0-N(22), in that the composition point shifted downwards due to the oxygen loss while the yttrium content is nearly constant. RESULTS AND DISCUSSION Density increases with increasing temperature and time for the specimens of C-0 and C-Odis shown in Fig. 2. C-0 The particle size distribution after 20 h planetary ball specimens without carboth-al reduction treatment shows milling shows very narrow particle size distribution the typical densification curve in which the rapid density ( D ~ : 0 . 3 8D50:0.22p, ~, Dlo:O.13p). increase at a relatively low temperature followed by a Oxygen Content of a batch of 93 wt% 6 wt% slower densification by the diffusion-controlledor reactiony 2 0 3 and 1 wt% A 2 0 3 without milling is 2-9 &? bY controlled mechanisms(23). C-0.6 specimens show the
0 l85O0C
0
9 1
r x c m
hh.A
A
= -
P)
0
L
1750°C n
2 ?J
1650'
.
A
A.
,
-
4-
I 15
.
, 20
.
I 25
.
, 30
.
, 35
.
, 40
.
I 45
.
I 50
2Theta (28)
Fig. 1. XRD peaks for the specimen C-0.6 with sintering temperature 578
Fig. 2. Relative densities of the specimens with increasing sintering temperature
Fig. 3. SEM micrographs for the specimen C-0 and C-0.6 sintered at 1850“Cfor(a) 0.5h and (b) 6h delayed densification until 1750°C which is the melting point of apatite phase, followed by the rapid increase in densification due to the increase in liquid content. Note that the density of C-0.6 specimen at 1750°C is still &5%TD while that of C-0 is nearly 90%TD. C-0.6 specimens have the apatite phase as second phase after carbothermal reduction. The precipitation of crystalline phase such as apatite from glassy phase leads to the decrease in liquid content. This leads to the retardation of sintering process, especially, for early stage of liquid phase sintering. Just after the melting of ap atite at 1750“c, the amount of liquid increases, which triggers the densification process at a relatively high temperature such as 1750“c . Since the grain growth and densification processes are thermally activated process, the occurrence of initial stage sintering at a relatively high temperature such as 1750°C compared to a general case at 1450°C may give rise to a significant difference. Fig. 3 shows the SEM micrographs for C-0 and C-0.6 specimens sintered at 1850°C for 0.5 h and 6 h. After sintering for 0.5 h, C-0 specimens showed dense microstructures with fine grains which correspond to 97% TD, while (2-0.6 specimen showed some pores with 80% TD(Fig.3-a). It is noticeable that large elongated grains are already formed even at 80% TD (C-0.6, 0.5 h), which can act as the seeds for the growth of large elongated grains. With further increase in sintering time to 6 h, width of the large elongated grains is getting larger and larger while the matrix grains remain fine(Fig.3-b). These microstructures are due to the formation of apatite phase that acts as a
transient phase. The precipitation of apatite phase retards the densification process at low temperature, and with increasing temperature, just after the melting of apatite transient phase, it triggers the densification process with increased amount of liquid phase. This suggests that suppressing the densification process to a certain high temperature followed by the reactivation of sintering process seems to provide a beneficial effect on the development of a bimodal microstructure.
CONCLUSIONS 1. Carbothermal reduction treatment of powder compacts of silicon nitride, sintering additives and carbon mixture takes the oxygen out of Y-Si-Al-0-N glass. 2. Loss of oxygen in Y-Si-Al-0-N glass leads to the precipitation of crystalline oxynitride phases such as apatite due to the increase in the N/O ratio. 3. Precipitation of apatite phase and its disappearance due to melting during sintering process result in the retardation of densification process up to 1750“C, followed by the rapid densification due to the increase in liquid content. 4. Occurrence of rapid densification process at a high temperature such as 1750“C maximizes the size difference between the growing grains and the matrix grains, and this leads to a distinct bimodal microstructure.
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REFERENCES 1. P. F. Becher, “Microstructural Design of Toughened Ceramics”,J. Am. Ceram. Sw-,W2L255-69, (1991)2. M. Mitomo, “Design and Processing for Silicon Nitride h-situ composites,” 6th International spposium On Ceramic Materials and Components for Engines, pp.85-90, (1997). 3. Y- J. Park, “Ab~Ormalgrowth Of faceted grainS in liquid matrix,” Ph.D thesis, KAIST, Korea, 1994. 4. Y. J. Park, N. M. Hwang and D. Y. YOO& “Abnormal growth of faceted (WC) grains in a (Co) liquid matrix,” Metall. Mater. Trans., 27A, 1-11, (1996) 5 . C. J. Lee, Deug J. Kim, and E. S . Kang, “Effect Of aSi3N4particle size on the microstructural evolution of Si3N4 ceramics”, J. Am. Ceram. Soc.,82[3], 753-56, (1999) 6. Hai-Doo Kim, Ellen Y. Sun, Paul F. Becher, Hyo-Jong Kim, Byung-Dong Han, and Dong-Soo Park, “Effect of increased oxygen content due to intensive milling on phase and microstructural development of silicon nitride, submitted to J. Am. Ceram. SOC. 7* Natansohn, Eand J‘ Rourke, “Effect of powder modifications on the properties of silicon nitride ceramics”, J. Am. Ceram. Soc., 76[9], 2273-84, (1993) 8. I. H. Kang, K. Komeya, T. Megum, M. Naito, and 0. Hayakawa, “Effect of Silicon Nitride Seeds Addition on the Particle Size and the Crystal Form of Resulting Powder in Carbothermal Reduction-Nitridation of Silica,” J. Ceram. SOC.Jpn., 104[6], 471475, (1996). 9. Magnus Ekelund and Bertil Forslund, “Carbothermal preparation of silicon nitride: Influence of starting’ material and synthesis parameters”, J. Am. Ceram. soc., 75[3], 53239, (1992) 10. Rasit Koc and Swaroop Kaza, “Synthesis of alpha silicon nitride from carbon coated silica by carbothermal reduction and nitridation”, J. E m . Ceram. Soc., 18, 147177, (1 998) 11. Simon J. P. Durham, Karti Shanker, and Robin A. L. Drew, ‘‘ Carbothermal Synthesis of Silicon Nitride: Effect of Reaction Conditions,” J. Am. Ceram. Soc.,74[1], 31-37, (1991 .) 12. Alan W. Weimer, Daniel F. Carroll, David W. Susnitzky, and Donald R. Beaman, “Morphology and sinterability of thermally treated carbothemally synthesized silicon nitride powders”, J. Am. Ceram. Soc.,82[6], 1635-38, (1999) 13. G. Woetting and G. Ziegler, “Powder characteristics and sintering behaviour of Si3N4powders”, Powder. Metall. Int., 18[1], 25-32, (1986) 14. Koji Watari, Mitsuru Kawamoto and Kozo Ishizaka, “Carbon behaviour in sintered silicon nitride grain boundaries:, Mater. Sci. Eng., A109,89-95, (1989) 15. U. Kolitsch, H.J. Seifert, T. Ludwig, and F. Aldinger, “Phase equilibria and crystal chemistry in Y203-A1203SiOz system”, J. Mater. Res., 14[2], 447-55, (1999). 16. G. Woetting, H. Heuer, and E. Gugel, “The influence of powders and processing methods on microstructure and properties of dense silicon nitride”, Mat. Res. Soc. Symp. Proc., Vol. 287, 133-46, (1993) 17. Stuart Hampshire, Robin A. L. Drew, and Kenneth H. Jack, ccViscosities, glass transition temperatures, and microhardness of Y-Si-Al-0-N glasses”, J. Am. Ceram. 9- ’
~~
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SOC.,67[3], (2-46-C-47, (1984) 18. R. E. Loehman. “Structure, formation, and stability of oxynitride gla~ses”,~pp. 167-73,& “Tailoringof M e c G c a l properties of Si3N4 ceramics”, eds M.J.Hof3ka.m and E: the lied science-v01.276, G.petzow, NATO As1 Kluwer Academic ~blishers,(1994) n a d D. P. Messier, “Determination of 19. C. p. G halysis~~, phase content of si3N4 by x-ray Am. C-. SOC.Bull., 56[9], 777-81, (1977). 20. ~ans-Joachim KlWbe and GWnter Zieglp, “Influence of crystallie secondary phases on the densification behavior of redon-bonded silicon during postsinkhg increased pressure79,J. Am. Ceram. SOC.,72[12], 2314-17, (1989) nsialOns: A study in ~ ~ t e r i a l ~ 21. K. H. Developmentlt, In Non-mde and Engineering Ceramics, ed. S. Hampshire. Elsevier Applied Science, London, pp.l -30, (1986) 22. L. J. Gauckler, H. Hohnke, and T. Y. Tien, “The system Si3N4-SiO2-Yz0T,J. Am. Ceram. Soc., 63[1-21, 35-37, (1980) 23. M. Mitomo, and M. Uenosono, “Microstructural development during gas-pressure s i n e g of alpha nitride”, J. Am. Ceram. Soc.,75[1], 103-08, (1992)
DESIGN OF SiCN - PRECURSORS FOR VARIOUS APPLICATIONS Gunter Motz*, Jurgen Hacker and Gunter Ziegler University of Bayreuth, Institute for Materials Research, D-95440 Bayreuth, Germany
ABSTRACT Main applications for preceramic polymers are ceramic fibres, matrix composites and coatings. The requirements for the polymer are determined by the intended application and differ from a cross-linkable liquid to meltable and curable or unmeltable but soluble solids. For applying these precursors on a larger scale, they must be processable by conventional polymerprocessing techniques. Moreover, the starting material should be reasonably cheap and readily available, and the synthesis should be uncomplicated. Therefore, novel tailored Si-C-N precursors were synthesised for various applications only by using the same two chlorosilanes but by different reaction types. Applying the liquid phase infiltration (LP1)-process for producing CMCs coammonolysis of CH3C12SiH and CH3C12SiCH=CH2 leads to a liquid cross-linkable precursor with a high ceramic yield (75 %) after pyrolysis at 1400 "C. Hydrosilylation of the chlorosilanes as the first step and subsequent ammonolysis result in a meltable solid (= 110 "C) with a very good stretchability because of specific ethylene-bridges. This polymer was also produced in a pilot plant and continuously meltspun to fibres with small diameters (= 20 pm). Subsequent curing and pyrolysis in a furnace at 1000 "C (N2)leads to non-oxidic SiCN fibres. The liquid SiCN precursor mentioned above was modified by an organometallic compound (e.g. Ti(NR2)4) and resulted in unmeltable but soluble solids. These polymers are suitable for applications as inorganic polymer- and at higher temperatures as ceramic-like-coatings (CLC). With corresponding precursor solutions different substrates (metals, ceramics, glass, plastics) with complex-shaped geometry can be coated by using simple dip- and spray-coating techniques. In all cases adhesion to the substrate is very strong. Selected results from the research areas mentioned above are presented in order to illustrate that a precursor must be tailor-made for successful use in different applications.
INTRODUCTION The main application areas for preceramic polymers are ceramic fibres, matrices in composites and coatings. The requirements for the polymer are given by the appropriate application and differ from a cross-linkable liquid to meltable and curable or unmeltable but soluble solids. For the broader distribution of these precursors,
they should be processable by conventional polymer processing techniques. Furthermore, the starting material should be readily available and relatively cheap and the synthesis should be uncomplicated. Polysilazanes are well suited to fulfil the requirements as preceramic polymers. Both educts, chlorosilanes and ammonia, are very cheap products of the chemistry industry. In most cases only a one or a two step reaction is necessary to get the desired polysilazane in a high yield. The degree of cross-linking (state of aggregation) and reactivity of the polymer is adjustable with different functional groups. In this work we report on three tailored new polysilazanes and their application as matrix in CMCs and graphite, ceramic fibres as well as ceramic-likecoatings.
EXPERIMENTAL PROCEDURE All operations for synthesis and processing of the precursors were carried out in an inert gas atmosphere because of sensitivity to air and moisture. The liquid silazane HPS was synthesised by coammonolysis of dichloromethylsilane and dichloromethylvinylsilane in toluene as described elsewhere [I]. Because of its low viscosity and high ceramic yield (Tab. 1) this precursor is suitable to infiltrate porous substrates like graphite, and for processing CMCs by the polymer infiltration route 121. Hydrosilylation and subsequent ammonolysis of the chlorosilanes yields quantitatively the ethylene-bridged colourless, brittle and meltable solid (ABSE). The melting point strongly depends on the processing conditions. This precursor was continuously meltspun to endless fibres with diameters of about 20 ym in a pilot plant (Fraunhofer ISC, Wiirzburg, Germany) by drawing speeds up to 500 d m i n . Subsequent curing of the green fibres by electron beam irradiation and subsequent pyrolysis in an innert gas furnace at 1000 "C leads to non-oxide ceramic SiCN fibres with very good oxidation resistance at temperatures up to 1500 "C. Both the synthesis of the HPS and the ABSE precursor can be transferred easily to a pilot plant (batch size 40 1). The reaction between transition metal compounds of the type M(NR2)4 (here M = Ti, R = alkyl) and the nitrogen function (NH-group) of the silazane backbone (here HPS precursor) is the basis for the synthesis of the unmeltable but soluble TIP precursor. The very moisture sensitive solid can easily be dissolved in solvents like cyclohexane, octane, toluene as well as xylene.
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Solutions of this polymer were used for coating of different substrates (steel, titanium, porous ceramics, glass, plastics) with complex geometry. The molecular weights of the polymers were determined cryoscopically in cyclohexane. The pyrolysis process was characterised by thermogravimetric measurements (TG), FTIR spectroscopy, MASNMR spectroscopy, X-ray diffraction (XRD) and elemental analysis. Gaseous products formed during heat treatment were identified by TG-coupled FTIR spectroscopy. All CMCs, fibres and coatings in the polymer as well as in the ceramic state were examined by scanning electron microscopy (SEM).
RESULTS AND DISCUSSION The properties of the various (Tab. 1) polymers resulted from the structural elements. The very liquid HPS precursor consists of a Si-N backbone with alternating Si-H and Si-vinyl groups. These functional groups allow a subsequent cross-linking into an unmeltable solid via hydrosilylation and polymerisation reactions. Using a radicalic initiator the cross-linking is possible at about 150 "C. In contrast to the HPS precursor the solid ABSE polymer consists of connected ethylene-bridged cycles without Si-H and Si-vinyl groups. The ethylene-bridges are the reason for the excellent stretchability, and the 5-membered cycles involve thermal stability of the molten polymer up to 170 "C for several hours. Further cross-linking only takes place via condensation of ammonia at temperatures higher than 200 "C. In the TiP precursor oligomeric HPS units are connected via a Ti atom. The degree of cross-linking is adjustable by the reaction temperature and the amount of the Ti-organic compound. For the synthesis of an unmeltable but soluble solid the optimal Si/Ti atom ratio ranges between 4:1 and 10:1. Heating the polymer up to temperatures higher than 200 "C causes further release of the corresponding amine and leads to a triple substitution at the Ti atom. The resulting polymer is insoluble. All precursors show a high ceramic yield up to 1000 "C (> 75 %). The molecular weights are in good agreement with the degree of cross-linking (see Tab. I).
precursor HPS ABSE
TiP
5 82
molecular state of synthesis weight aggregation yield 440 g/mol liquid, 0.05 Pas 85 % (at 20 "C) 1100 g/mol solid, meltable 70 % at ca. 100 "C 950 g/mol solid 95 % unmeltable
Infiltration of woven fabrics and porous carbon pistons with the HPS precursor After infiltration of woven carbon fibre fabrics (36 % fibre volume, KDL5002, SGL Carbon, coated with
150 nm pyro-C) applying the RTM technique [2] the liquid precursor was converted at 150 "C into an unmeltable fibre-reinforced thermoset by using an initiator. Subsequent pyrolysis in nitrogen atmosphere at lo00 "C leads to a very porous CMC. Therefore, the infiltration and pyrolysis steps were repeated up to 7 times (PI-P7). Because of the low viscosity and the low surface tension it was possible to get a minimal residual open porosity of 3 % in the P7 state (Fig. 1). For these specimens flexural strength values of about 330 MPa were achieved.
Fig. 1: CMC in the P7 state with minimal open porosity
Fig. 2: Graphite based piston (Motoren GmbH Greiner MGG) infiltrated with the HPS precursor and tempered at 1000 "C in nitrogen
ceramic yield (up to 1000 "C) 75 %
composition after 1000 "C [mass %] Si C N H O 51.4 26.5 21.9 < 0.1 0.4
75 %
49.9
27.2
20.9
< 0.2
1.9
77 %
39.7 16.5 Ti = 17.5
26.5
< 0.2
1.0
Graphite based pistons (Fig. 2) offer a great potential to reduce fuel and oil consumption as well as the exhaust fumes emission. However, commercially available graphite shows an open porosity of about 10 % and a still too low strength. By the use of the HPS precursor it was possible to close the open porosity of the piston and to increase the strength of the graphite by 10 percent. Nevertheless, further optimisation is necessary.
Processing ceramic fibres by using the ABSE precursor The brittle and meltable ABSE polymer fulfil the most important requirements for the melt spinning process (high melting point, thermal stability, and stretchability, suitable viscosity in the molten state). The precursor was continuously meltspun to green fibres with a diameter of 20 - 30 pm (withdrawal speed 500 d m i n ) in a pilot plant at the Fraunhofer ISC Wiirzburg. The fibre surface is smooth and the fracture surfaces show no pores or cracks. The main problem is to cure the green fibre in a very short time before pyrolysis for making the fibres infusible. For silazanes the ABSE polymer shows an astonishing resistance against moisture. Therefore, handling in air is possible for a short time. However, the low reactivity is disadvantageous for the chemical curing process. In contrast, irradiation with electron beam leads in a very short time to further cross-linking of the polymer backbone. The cured unmeltable fibres show an acceptable flexibility. After curing the fibres were heated up to lo00 "C for I h in nitrogen atmosphere. The surface remained smooth and no pores or cracks could be observed. To investigate the oxidation behaviour the ceramized fibres were heated for 12 h in air at 1500 "C. Figures 3 and 4 illustrate the oxidation resistance even at such high temperatures, because of the formation of a pore-free, passivise surface.
Fig. 3: SiCN fibre (ABSE) after oxidation in air at 1500 "C for 12 h
Fig. 4: Pore- and crackfree oxide coating on the fibre surface
Processing of polymer and ceramic-like layers by using the TiP Precursor The TiP precursor is a moisture-sensitive solid. Thus, coating can only be carried out in a protective atmosphere. The polymer can easily be dissolved in solvents such as cyclohexane, octane, toluene and xylene. The adhesion on various substrates, i.e., on metals, Si-wafer, ceramics, glass and plastics, is very strong. The thickness of the polymer coating depends on the concentration of the solution and the coating technique used. The layer thickness ranges from 0.5 (for dip-coating) up to 5 pm (for spraying). The polymer coatings on steel-sheets and silicon-wafer was pyrolysed in a furnace at various temperatures up to 1000 "C both in air and nitrogen or recently by laser radiation. This method enables the creation of ceramic layers even on aluminium or plastics. By annealing in air, the nitrogen in the layer will be partly replaced by oxygen, thus the system changes to SiTi0,C. Therefore, oxygen free SiTiCN as well as SiTi0,C coatings can be produced. The possibility of handling in air is of significance for the ceramisation process, because pyrolysis can be also carried out in air, which leads to a considerable simplification of the process (Fig. 5). The chemical stability of the coating against water, almost all organic solvents and acids is already given after hardening in air, whereas the corrosion stability of the coating in an alkaline environment strongly depends on the heat treatment. In most cases, good resistance to bases is reached after annealing at 300 "C. The hardness of the layers also strongly depends from the annealing temperature. To measure the hardness values of thin coatings and not the substrate hardness, a microhardness measurement with a nanoindenter (max. load 5 mN) was used. After cross-linking at 130 "C the TIP layer shows a hardness value of about 0.43 GPa (universal hardness [3]). With rising annealing temperature up to 1000 "C the hardness increases to 8.0 GPa. That is nearly the hardness of monolithic SiTiCN ceramic (8.5 GPa).
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Handling i n k Oxide Coating
Handling in Nitrogen: Non-Oxide Coating
Exchange of Nitrogen
Separation of Ammonia and ceramization
I
Ceramic-Like SiTi0,C Layer
-
\ T > 700 OC
J
Ceramic-Like
-
&[ SiTiCN Layer
Fig 5 : Processing of oxide and non-oxide polymer and ceramic-like layer
Latest results show the possibility to pyrolyse the polymer layer with laser radiation. If the silazane layer is thinner than 5 pm it is almost transparent for all kind of radiation. To improve the absorption of laser radiation in the thin coating ultrafine flake graphite powder was mixed into the solution of the polymer. After coating, the layer contains 10 vol% graphite powder. With this mixture, coatings with a thickness from 3 to 5 pm were produced on steel and aluminium substrates. With a Nd:YAG-laser and a laser power of about 600 W the hardness values of the layer remarkably increases up to 3.5 GPa and corresponds with a furnace pyrolysis at 600 "C. Because of shrinkage during pyrolysis layers thicker than 10 pm show many cracks. However, for precise, accurate ceramic-like structures (Fig. 6) this technique will be suitable (i.e. for microelectronic applications).
Figure 6 shows a coated A1203 substrate. The polysilazane layer was partly pyrolysed with an ArgonIon-Beam Laser. After irradiation the rest of the polymer was removed with an organic solvent.
CONCLUSIONS By varying the reaction type it is possible to produce novel tailored precursors in the systems Si-C-N(Ti) on the basis of only two inexpensive chlorosilanes. The nature and the number of the functional groups as well as the degree of cross-linking are the critical factors for the following applications. The tailored precursors can be used for infiltration of fibre preforms or porous carbon (HPS), for processing fibres (ABSE) and coatings (TIP). CMCs and infiltrated carbon manufactured with precursor HPS show a very low residual porosity and good mechanical properties. The ABSE polymer is an excellent precursor for the melt spinning process. Curing at low temperatures with electron beam was successfully tested and the subsequent pyrolysis leads to ceramic fibres with a high yield (70 wt%) and a very good oxidation resistance. The TiP precursor gives the possibility for producing polymer- and ceramic-likecoatings via simple painting techniques. The transition polymer - ceramic is adjustable by the annealing temperature. Therefore, the hardness and the corrosion resistance of the coatings depends to a great extent on the pyrolysis temperature. The hardening via laser radiation offers the possibility even to protect aluminium components or plastics as well and to generate precise ceramic-like structures.
ACKNOWLEDGEMENTS Fig. 6: With laser pyrolysed SiCN structure on an A1203 substrate
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This work has been supported bei the Bayerische Forschungsstiftung, Bayerisches Staatsministerium fur Wissenschaft, Forschung und Kunst (Bavaria,
Germany). In addition, F. Berndt and M. Ordung, University of Bayreuth, as well as H.-J. Krauss, University of Erlangen (LFT), Dr. B. Clauss, University of Stuttgart (ICF) and Fraunhofer ISC Wurzburg, are acknowledged for their support.
REFERENCES (1) J. Lucke, J. Hacker and G. Ziegler, Appl. Organomet. Chern., 11 181 (1997) (2) G. Ziegler, I. Richter, D. Suttor, Composites Part A, 30 41 1 (1999) (3) H.-H. Behncke, Hiirterei-Technische-MitteilungenHTM, 5 304 (1993)
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GELATIN CASTING AND STARCH CONSOLIDATION OF ALUMINA CERAMICS W. Pabst, E. Gregorovh, J. Havrda, E. Tfnovk
Department of Glass and Ceramics, Institute of Chemical Technology (ICT) Prague, CZ 166 28 Prague 6, Czech Republic
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ABSTRACT Two variants of slip casting of ceramic suspensions with organic additives into impermeable molds are studied: gelatin casting and starch consolidation. These forming techniques are applied to suspensions containing 75 and 80 wt.% of alumina powder of different purity and particle size. Gelatin casting is used to prepare ceramic bodies with approx. 95 % theoretical density. Starch consolidation results in a highly porous ceramics (with total porosities up to 35 YO)with a large pore size (tens of microns). The microstructure of the samples prepared is investigated by the Archimedes method (bulk density, porosity), mercury intrusion and optical image analysis.
INTRODUCTION Traditional slip casting exploits the capillary suction of porous (e.g. plaster) molds, possibly combined with external pressure, for dewatering of aqueous suspensions at the slip-mold interface. From a theoretical point of view this dewatering process can be modeled as a diffusion process (Fick's law) or as a filtration process (Darcy's law). Both models result in a square-root-of time kinetics for body formation, so that time is the critical factor when bodies with large wall thickness are to be produced. Furthermore the varying body formation rate, and the very fact that demixing occurs at the slip-mold interface, give way to microstructural variations (e.g. density gradients or particle orientation in the case of anisometric particles) from the surface of a body to its interior. Slip casting of ceramic suspensions into impermeable molds ("impermeable mold Casting", in the following abbreviated IMC) is a family of shaping techniques for the near net-shape forming of small- and large-sized ceramic parts. Some IMC methods are based on purely organic vehicles (e.g. the hot molding process [ l ] or some gelcasting processes [2]), while others use essentially water-based suspensions (e.g. direct coagulation casting [3], protein forming [4], starch consolidation [4-51 and gelatin casting [6-71, but also some gelcasting processes [8-91. The common feature of all these processes is a phase transition of the organic phase (mostly in connection with water) which transforms the suspension into a solid green body after
casting into the mold. This phase transition (e.g. stiffening of polymer melts, polymerization reaction, sol-gel transition, swelling of starch globules) can be temperature-induced or chemically initiated. In any case the process step of body formation can be controlled in such a way that stiffening occurs more or less simultaneously throughout the whole volume. The kinetics of body formation is thus essentially a phasetransition or reaction kinetics (typically of exponential type), modified at worst by the kinetics of heat transfer. Body formation in these IMC methods is therefore substantially faster for larger-sized components than in classical slip casting, and the slip-mold interface has no (or at least minimal) influence on the microstructure of the green body, in contrast to classical slip casting (and also paste forming techniques like injection molding). Apart from the specific principal advantages mentioned, i.e. the possibility of obtaining better uniformity of microstructure and the possibility of time-efficient production of larger-sized bodies (unrestricted by the square-root-of-time law), these new IMC techniques have a number of practical advantages over classical slip casting techniques, e.g. the fact that mold materials are to a large extent arbitrary (e.g. glass, metal, polyethylene, wax), molds have a long lifetime and can be reused without complicated cleaning or drying. On the other hand it has to be noted that the preparation and rheology control of the suspension itself needs at least the same care as in classical slip casting, and may be complicated due to the lack of experience with the interplay of the organic components with the aqueous slip. The present study examines the possibility of preparing alumina ceramic bodies by IMC of aqueous suspensions containing gelatin and starch, respectively. Gelatin casting exploits the ability of aqueous gelatin sols to gel upon cooling after heating to approx. 50-60 "C, starch consolidation uses the fact that starch globules absorb large amounts of water and swell to several times their original volume when heated to approx. 60-80 OC [4-7, 10-121.
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Starch Consolidation Gelatin Casting Gelatin casting has been tested in this work with a number of commercial alumina powder types ( h m the Advanced Alumina "AA" series by Sumitomo Chemical Co., Ltd./ Japan and several grades by Alcoa Industrial Chemicals / USA and Europe), differing mainly in purity (99.7 - 99.99 YO)and in particle sue Table I gives an overview of the powders (0.4 - 5 p). applied, their purity, particle size and specific surface according to the suppliers' information. Table I. Characteristicsof the powders used in this work according to suppliers' data (* Sedigraph 5100, + Cilas 850) Powdertype
Purity
"1 AA-05 AA-07 AA-2 AA-5 A-1000 SG A-16 SG CT-3000 SG HVA FG
> 99.99 > 99.99 > 99.99 > 99.99 99.8 99.8 > 99.8 99.7
Median sizeDSo
Specific surface
0.4-0.6* 0.6-0.8* 1.8-2.2* 4-6* 0.4* 0.4* 0.7' 5.3'
2.8-3.3 2.0-2.6 0.8-1.1 0.3-0.5 8.8 9.5 7.0 0.65 ~~
Suspensions containing 75 and 80 wt.% of solids were prepared conventionally by mixing the powders with distilled water, adding an appropriate amount (0.5-0.6 wt.%) of deflocculant (an alkali-& carbonic acid mixture), shaking with alumina balls and sonication. The de-airing step had to be omitted, because it led to significant increases of viscosity, which made the suspensions uncastable. Three different variants of gelatin addition have been tested In variant I dry gelatin is added to the asprepared suspension before heating, in variant I1 a preheated 5 wt.% aqueous solution of gelatin is added to the as-prepared preheated suspension, and in variant I11 dry gelatin is added to the suspension before homogenization by shaking with alumina balls in a polyethylene bottle in a water-bath (temperature approx. 60 "C). After casting the hot (approx. 60 "C) gelatincontaining suspensions into cylindrical polyethylene molds with diameter of approx. 30 mm, the samples cooled down to room-temperature, and gelling started. After removal from the molds the samples were dried in a laboratory drier at 110 "C before firing to 1550 "C (for HVA FG also 1600 "C) using a usual firing regime (2 "C/min, 2 h dwell at maximum temperature, free cooling in the furnace).
588
The technique of starch consolidation was in this work applied to a high-purity (99.99 YO)submicron (median size 0.5 pm) alumina powder (grade AA-05, Sumitomo / Japan). The preparation of the basic suspensions (75 and 80 wt.% of solids) was similar to that used for gelatin casting. Native potato starch (Naturamyl / Czech Republic) in the dry state was added in amounts of 5 , 7.5, 10, 15, and 30 vol.% (related to the oxide powder phase) and the whole mixture was carehlly stirred in a polyethylene vessel. De-airing was performed for 5 min, followed by 10 min sonication and again 5 min evacuation. Since the specifically lighter starch globules (density 1.45 g/cm3) showed a tendency to segregation within the suspension (density 2.29-2.50 g/cm3), the whole mixture was again carehlly stirred before casting. Cylindrical polyethylene molds with a diameter of approx. 30 mm were used for casting. Immediately after casting the molds were closed and put into a water bath for heating to approx. 70-80 OC for 1 h. Subsequently, the covers were opened and the green bodies were allowed to dry at ambient conditions (air at atmospheric pressure at room temperature) for 24 h in the mold. After removing the as-formed green bodies flom the mold, they were dried at 120 "C for several hours. The bodies were then fired to 1530 "C (some also to 1600 "C). In order to guarantee defect-flee burnout of the rather large (with respect to volume) amount of organic matter, the firing regime proposed by Lyckfeldt & Ferreira [5] has been used in the case of starchconsolidated bodies (1 "C/min up to 200 "C with 1 h dwell, 1 "C/min up to 300 "C with 1 h dwell, 1 "C/min up to 500 "C without dwell, 1 OC/min up to 1530 or 1600 "C with 2 h dwell). The particle size distribution of all alumina powders used in this work (and of the starch particles) has been measured by LALLS (low-angle laser light scattering, Analysette 22, Fritsch / Germany), the specific surface of the alumina powders by nitrogen adsorption (Nelsen-Eggertsen method). The rheological behavior of the starch-containing suspensions was measured by rotational viscometry (Rheotest' 2, Medingen / Germany) and pH of basic suspensions by the direct potentiometric method (Ionometer MS-3 1, Lavat / Czech Republic) using a glass and a calomel electrode. DTA and TG measurements (TG-750 Stanton-Rederofi / USA) were performed to characterize the burnout behavior of starch and gelatin. All alumina samples in the as-fired state were subjected to shrinkage measurements as well as to bulk density and porosity measurements using the Archimedes principle. Additionally, the samples prepared by the starch consolidation method were characterized by mercury intrusion (Poresizer 9320, Micrometritics I USA) and by optical image analysis (Lucia, Laboratory Imaging / Czech Republic).
RESULTS AND DISCUSSION Characterization of Powders, Suspensions and Organic Components Table I1 lists for the alumina powders used the average particle size (median equivalent diameter) measured by LALLS (using Fraunhofer theory for data evaluation), the specific surface calculated via the harmonic mean determined fiom the LALLS data (assuming spherical particle shape) and the specific surface measured by the Nelsen-Eggertsen method.
measurement is performed in water). This hypothesis is confirmed by the particle size distributions of these powders (not shown here), which are clearly bimodal with a second mode at sizes > 10 pm. Table I11 shows pH values of basic suspensions containing 75 wt.% of alumina in distilled water (pH approx. 6-7) with and without deflocculant (0.58 wt.% based on solids, pH approx. 7), respectively. Table 111. Measured pH values for aqueous suspensions of all alumina powders used in this work (powders indexed * did not form a liquid suspension in water without deflocculant)
Table 11. Measured particle size and specific surface of all alumina powders used in this work
I
Powdertype
AA-05
I
Median-1 Sppctz size DSo (LALLS) (LALLS) [pml b2kl 0.8 2.3
Powder type
Specific surface (NE)
[m2k1 3.4
A comparison of the particle size distribution curves (not shown here) shows that powders of the AA series (Sumitomo) have a narrow distribution, while the particle size distribution of Alcoa powders is generally broader. A comparison with Table I shows, that for all submicron powders the LALLS median diameters measured are larger than according to the suppliers' data. This has to be attributed to the fact, that Fraunhofer theory - although routinely used in this way - is not adequate for submicron particles and in order to obtain more reliable results Mie theory should be used for LALLS data evaluation. Nevertheless there is a clear correspondence of the median size results. The same can be said about the specific surface results obtained by the Nelsen-Eggertsen method compared to the suppliers' data (BET method) listed in Table I. Also the coincidence of the specific surface values calculated fiom LALLS data (via the harmonic mean size) with those measured by the Nelsen-Eggertsen method is with regard to the simplified measuring principle (compared to classical BET) reasonably good for all powders of the AA-series (Sumitomo) and for HVA FG (Alcoa). Systematic, however, is the large difference (by a factor of approx. 4) of the calculated and measured specific surface of the powder types A1000 SG, A-16 SG, and CT-3000 SG. Since differences in particle shape alone cannot account for such a large difference, the phenomenon must be attributed to the occurrence of agglomerates when the submicron Alcoa powders are suspended in water without deflocculant (note that the LALLS
I
AA-05 AA-07 AA-2 AA-5 A- 1000 SG A-16 SG CT-3000 SG
pH without deflocculant 4.4 4.2 4.8 7.3
---*
---* ---*
pH with deflocculant 7.9 8.0 8.1 7.9 9.0 9.1 9.2
Interestingly, the powder types A- 1000 SG, A- 16 SG, and CT-3000 SG did not form a liquid suspension in distilled water without deflocculant, while all other powders do. This may be due to the aforementioned fact that these three powder types show a large degree of agglomeration in water. With deflocculant, 75 wt.% suspensions of Sumitomo powders (AA-series) exhibit a pH of approx. 8, while for the Alcoa powders pH ir 9. Figure 1 shows the particle size distribution of the native potato starch used in this work measured by LALLS and Figure 2 shows the particle size distribution of this same starch as measured by automatic image analysis. The dry starch globules are of ellipsoidal shape (with only a small deviation fiom sphericity) with a median diameter of approx. 45-53 pm (45 pm has been determined by LALLS, 53 pm by image analysis) and a mode between 55 pm (LALLS) and 60 pm (image analysis). Figures 3 and 4 show thermogravimetric and DTA curves for the starch and the gelatin used in this work. Both exhibit strongly exothermic DTA peaks between approx. 350 and 650 OC. The last burnout products of starch vanish at temperatures higher than 600 OC, residual gelatin can remain up to temperatures higher than 700 OC. Starch burnout is most intensive between 400 and 600 "C. Therefore the application of the special firing regime, up to 500 "C at least, is useful. Since for gelatin casting the total gelatin content in the system is comparably low, no such care is necessary in this case.
589
0.1
10
1
100
1000
Equivalent diameter (um)
Figure 1. Particle size distribution of native potato starch measured by LALLS
0.1
1 10 100 Equivalent diameter (um)
1000
Figure 2. Particle size distribution of native potato starch measured by optical image analysis
the Sumitomo supermicron types (AA-2 and AA-5), but also unsuccessful for the submicron types (AA-05 and AA-07). Variant 111 was successfully tested for CT-3000 SG and seems to be the best variant for fast gelation (body formation within seconds - few minutes). While samples from 75 wt.% suspensions usually show deformation, and sometimes cracking, during drying, samples from 80 wt.% suspensions dry without problems to the desired shape and predictable dimensions. When using relatively coarse powders (AA-2, AA-5 and HVA-FG) casting itself is greatly facilitated (easier for HVA FG than for AA-5 and AA2), but these systems are prone to phase separation before gelation is completed. From the submicron powders investigated the Alcoa types A-1000 SG, A16 SG, and CT-3000 SG show evidently the best casting behavior in connection with gelatin. Many characteristic differences between these three powder types and the AA-types have been mentioned (cf. Tables I, 11, 111) but the definite reason is not yet clear and is a subject of current research. Table IV lists linear shrinkage, bulk density, and apparent (i.e. open) porosity of gelatin-cast alumina bodies made from a suspension with 80 wt.% of CT3000 SG after firing to 1550 "C. Table IV. Linear shrinkage A, bulk density p, apparent (i.e. open) porosity Po of gelatin-cast bodies (index ' means 80 wt.%, otherwise 75 wt.% solids content)
l6
Powdertype
0
200
400
600
800
Firing temp.["C]
A ["/.I
P PO [g/cm3] ["/.I
-2 1000
Temperature ( OC )
Figure 3. DTA and TG curves for starch
6 5 4 ln ln
2
3 , 2 3 1 % 0
40
20
-1
0
0
200
400
600 800 1000
Temperature ('C)
Figure 4. DTA and TG curves for gelatin
Gelatin Casting Independently of the solids concentration in the suspension (75 or 80 wt.%), gelatin casting according to variant I could be used for powders A-1000 SG, A16 SG, CT-3000 SG and HVA FG (Alcoa), but turned out to be impossible with all powders of Sumitomo's AA series (AA-05, AA-07, AA-2, AA-5). Gelatin casting according to variant I1 was more successful for
590
As expected, shrinkage, bulk density and porosity values clearly show that the bodies prepared from the supermicron powders AA-2, AA-5 and HVA FG (which are easily cast with gelatin-containing suspensions) are not yet sintered at the temperatures used. It was confirmed, however, that no defects occurred during firing, so that no special firing regime is needed for gelatin burnout. Bodies prepared by gelatin casting from the Alcoa submicron powders exhibited a linear firing shrinkage of approx. 16-18 % and a bulk density of approx. 93.5-95.0 'YO of theoretical density. After due optimization of the process bulk densities close to theoretical density can be expected for these powders.
Starch Consolidation Table V lists shrinkage A, bulk density p, open and total porosity (Po and P,) for as-fired alumina bodies prepared by starch consolidation with 5,7.5, 10, 15 and 30 vol.% of dry starch (based on solids) in a suspension containing 75 wt.% and 80 wt.% of AA-05 (after firing to 1530 OC and 1600 OC,respectively). Table V. Linear shrinkage A, bulk density p, apparent (i.e. open) porosity Po and total porosity Pt of starchconsolidated bodies (index means 80 wt.%, otherwise 75 wt.% solids content)
'
[vol.%]
[OC]
1530
15.7
cm31 3.48
1.4
13.0
It is evident that, when firing is performed under atmospheric pressure, starch-consolidated bodies cannot be densely sintered (i.e. to densities approaching theoretical density TD). Lyckfeldt [4] used 2-5 vol.% of starch in Si3N4 suspensions and obtained 89.1 % of TD for starch-consolidated Si3N4 bodies sintered by gas pressure sintering and 99.7 % of TD by hot isostatic pressing. For starch contents as high as ours, starch acts at the same time as a consolidating medium and as a pore-forming agent. Porosities as high as 35 % have been obtained and it seems that even higher porosities might be realizable. Table VI lists the bulk density p and open porosity Po measured by mercury porosimetry and the total porosity Pt determined by optical image analysis on polished sections of fired (1530 "C)alumina bodies prepared fkom 75 wt.% AA-05 suspensions with 7.5 and 15 vol.% of starch, respectively. Table VI. Bulk density p and open porosity Po(as measured by mercury porosimetry) and total porosity P, (as measured by image analysis) of starch-consolidated bodies Starch content [vol.o/,] 7.5 15
p
wm31 3.02 2.75
Po"] 15.7 24.8
error of 13-23 %), but even within this error the agreement is unsatisfactory for the highly porous sample. The discrepancy is due to the insufficient quality of the polished section (the preparation of which is difficult in practice for such highly porous samples). Table VII lists some of the pore size characteristics determined by image analysis for starch and specimens of fired (1530 "C) alumina bodies prepared from 75 wt.% AA-05 suspensions with 7.5 and 15 vol.% of starch, respectively. Table VII. Pore size characteristics determined by image analysis for starch and specimens of fired (1530 "C)alumina bodies prepared fkom 75 wt.% AA-05 suspensions with 7.5 and 15 vol.% of starch Starch Median [p] 53.2 Moderwnl 60
7.5 vol.% 55.8 70
15 vol.% 63.4 90
The pore sizes after firing (where the submicron alumina skeleton can be assumed to be sintered) are of the same order of magnitude as the particle size of dry starch. The absolute values are larger, which is not surprising, since the starch undergoes swelling during the heating step in water. The pore size distribution obtained by mercury intrusion shows a completely different picture (see Figs. 5 and 6). According to this method, the dominant part of the (volume-weighted) pore size distribution should be in the submicron range. The results of image analysis (which are rather close to that of dry starch, cf. Figs. 1 and 2) clearly prove that this is far fkom reality. The reason for this discrepancy is the very unrealistic model of a system of cylindrical pores with constant diameter, which forms the basis of the evaluation of the mercury intrusion data by the Washburn equation. Knowing the realistic pore size distribution from image analysis it is justified to interpret the minimum in the fkequency curve fkom mercury intrusion as the most fkequent size of interconnections between pores. For the starch-consolidated specimens (both with 7.5 and with 15 vol.% of starch in the suspension) this size is approx. 7.4 p,i.e. approx. one order of magnitude smaller than the pore size. 7 ,
PtWI 26.3f 6.7 45.2f 5.9 031
The bulk density and the open porosity determined by mercury intrusion are in good agreement with the values measured by the Archimedes principle (cf. Table V). The total porosity measured by image analysis contains a rather large error (absolute error in porosity approx. k 6 YO,which corresponds to a relative
1
10
100
1000
Equivalentdiameter (urn)
Figure 5 . Pore size distribution of as-fired alumina sample (prepared from a 75 wt.% suspension with 7.5 vol.% starch) according to mercury intrusion
591
25 n
0.1
1 10 100 Equivalentdiameter (um)
I 1000
Figure 6. Pore size distribution of as-fired alumina sample (prepared fiom a 75 wt.% suspension with 15 vol.% starch) according to mercury intrusion
CONCLUSION Two variants of slip casting of ceramic suspensions with organic additives into impermeable molds have been studied gelatin casting and starch consolidation. Gelatin casting has been used to prepare ceramic bodies with approx. 95 % theoretical density, but highpurity, submicron powders turned out to be problematic in combination with gelatin. The reasons for these problems are currently being examined, and the modification of the process with catalysts will be the subject of fbture research. Starch consolidation has been successhlly applied to suspensions of submicron, high-purity powders ( a - 0 5 , Sumitomo) and yields highly porous ceramics (with total porosities up to 35 'YO) with pore sizes comparable to that of dry starch (tens of microns as determined by LALLS and image analysis). Mercury porosimetry yields another important information, viz. the size (equivalent diameter) of the interconnections between pores, which is approx. 7 pm for the starch-consolidated samples.
Acknowledgement: This study was part of the research project CEZ:MSM 2231 00002 "Chemistry and Technology of Materials for Technical Applications, Health and Environment Protection" and supported by grant MPO No. FB-CV/64/98.
REFERENCES [ l ] R. Lenk, Hot Moulding - An Interesting Forming Process, cfi / Ber. DKG, 72 (1995) 636-642.
[2] 0. 0. Omatete, M. A. Janney, R. A. Strehlow, Gelcasting - A New Ceramic Forming Process, Ceram. Bull., 70 (1991) 1642-1649, [3] T. J. Grade, F. H. Baader, L. J. Gauckler, Shaping of Ceramic Green Compacts fiom Suspensions by Enzyme Catalyzed Reactions, cfi / Ber. DKG, 71 (1994) 317-323. [4] 0. Lyckfeldt, Novel Water-Based Shaping of Ceramic Components, Sixth E.Cer.S Conference and Exhibition (Extended Abstracts Vol. 2), British Ceramic Proceedings No. 60, Institute of Materials, London (1999), 2 19-220.
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[5] 0. Lyckfeldt and J. M. F. Ferreira, Processing of Porous Ceramics by Starch Consolidation, J. Eur. Ceram. SOC.18 (1998) 131-140. [6] Y. L. Chen, Z. P. Xie, J. L. Yang, Y. Huang, Alumina Casting Based on Gelation of Gelatine, J. Ew. C-. SOC.19 (1999) 271-275. [7] Z. P. Xie, Y. L. Chen, Y. Huang, A Novel Casting Forming for Ceramics by Gelatine and Enzyme Catalysis, J. Eur.Ceram. SOC.20 (2000) 253-257. [8] K. Prabhakaran, C. Pavithran, Gelcasting of Alumina Using Urea-Formaldehyde - I. Preparation of Concentrated Aqueous Slurries by Particle Treatment with Hydrolysed Aluminium, Ceramics International 26 (2000) 63-66. [9] K. Prabhakaran, C. Pavithran, Gelcasting of Alumina Using Urea-Formaldehyde II. Gelation and Ceramic Forming, Ceramics International 26 (2000) 67-7 1. [lo] R. Macrae, R. K. Robinson, M.J. Sadler (eds.), Encyclopedia of Food Science, Food Technology and Nutrition, Academic Press, San Diego (1993), 2176-2181 and 4372-4389. [l 11 A. M. Stephen (ed.), Food Polysaccharides and Their Applications, Marcel Dekker, New York (1995), 19-67. [ 121 R. H. Walter (ed.), Polysaccharide Association Structures in Food, Marcel Dekker, New York (1998), 1-168 and 289-328.
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LIQUID -PHASE SINTERED SILICON CARBIDE BASED CERAMICS WITH A1N-Y2O3AND AIN-La203ADDITIVES V. A. Izhevskyi’, L. A. Genova, A. H. A. Bressiani, J. C. Bressiani”
Instituto de Pesquisas EnergCticas e Nucleares, IPEN - CNEN/SP C. P. 11049, Pinheiros, 05542-970, S. Paulo, SP, Brazil ABSTRACT Microstructure development and phase formation processes during sintering of silicon carbide based materials with 10 vol.% of AlN - Y2O3, or AlN - La203 sintering additives were investigated. Densification of the materials occurred by liquid-phase sintering mechanism. Ratio of a- and p-Sic powders in the initial mixtures was a variable parameter. Shrinkage behavior during sintering was shown to strongly depend both on the combination of sintering additives and on the sintering atmosphere used. Kinetics of p-Sic to aS i c phase transformation under conditions of postsintering heat treatment at 190O-195O0Cwas studied, the degree of transformation determined by quantitative xray analysis. Evolution of microstructure resulting from p to a phase transformation was followed up by scanning electron microscopy. Transformationcontrolled grain growth mechanism similar to the one observed for silicon nitride based ceramics was established. Possibility of in-situ platelet reinforced dense Sic-based ceramics fabrication by means of sintering was shown.
Key words: silicon carbide, ceramics, sintering, phase formation, post-sintering heat treatment, phase transformation, microstructure. INTRODUCTION Silicon carbide is considered to be an important structural ceramic material because of a promising combination of properties, such as high oxidation resistance, good mechanical properties retained to high temperatures, high wear resistance, good thermal shock resistance due to high thermal conductivity, etc. All these properties are inherent to silicon carbide due to highly covalent bonding. The latter, however, causes complications with sintering of Sic-based ceramics to high densities, which disadvantage is characteristic for all non-oxide covalent compounds, such as Si3N4and AlN. Unlike the aforementioned compounds, which tend to decompose severely at high temperatures, silicon carbide can be densified by solid-state sintering process at high temperatures of about 210OoC with the aid of B and C [l], which dramatically improve the shrinkage kinetics. However, thus sintered materials have poor or, I
at the best, moderate mechanical properties (flexural strength of 300-450 MPa and fracture toughness of 2.54 MPaem-’”). Liquid-phase sintering of Sic can be achieved at much lower temperatures ( 1800-19OO0C)[2, 31 with the aid of metal oxides, such as A1203,Y2O3, and other rare-earth oxides [4-81. The densification of S i c by liquid-phase sintering lately draws more attention because the materials processed by this method exhibit superior mechanical properties. The liquid phase sintering of S i c is somewhat similar to the same process for Si3N4.The oxide sintering aids react with Si02,which is always present at the surface of Sic particles, forming an oxide melt and enhancing densification. However, oxides interact with Sic with massive gaseous products formation leading to high weight loss and porosity [9, 101. It is known that alumina may interact with silicon carbide according to the following reactions [4] : S i c (s) + A1203(s) -+ A120(g) + SiO(g) + cO(g) 2SiC(s) + A1203(s) -+ Al2O(g) + 2Si(l) + 2CO(g) 3SiC(s) + A1203(s)+ 2Al(l) + 3Si(l) + 3CO(g) These reactions occur more actively with the increase of alumina content and the temperature of sintering. These reactions can be to some extent suppressed by the application of the high external partial pressure of CO and/or application of reactive powder beds [4]. However, such an approach is not completely effective and enables to achieve the final density not higher than 98% of theoretical density. Moreover, such techniques are costly and do not guarantee the reproducibility of materials properties. Additional weight loss during liquid phase sintering of S i c occurs due to the reaction between Sic and Si02 surface films: 2SiOZ(s) + SiC(s) + 3SiO(g) + CO(g) It is therefore obvious that an alternative combination of sintering additives, which will eliminate or reduce the above mentioned effects determined for densification, will be the best possible solution. As it was suggested by Chia et a1 [ 113, aluminum nitride, AlN, in combination with yttria, Y2O3, may present a solution. However, this possibility was not investigated in detail and only a limited amount of information is available [ l l , 121. Which is of additional interest, due to certain structural similarities S i c and AIN produce solid
invited scientist. on leave from the Institute for Problems of Materials Science. National Academv of Sciences of Ukraine. Kiev. Ukraine
593
solutions and formation of mixed crystalline structures was observed [13-171. By analogy with silicon nitride, it seems to be possible to tailor the microstructure and the structuresensitive properties of silicon carbide based ceramics by varying a- to p-Sic ratio in the initial mixtures. If a-Sic is used as a starting powder, the final material is characterized by fine homogeneous microstructure with uniaxial grains [ 5 , 181, which results in moderate fracture toughness. Additions of p-Sic in combination with some specially developed thermal treatment leads to in-situ platelet reinforced material formation with improved mechanical properties [5,7, 191. In the present work sintering behavior and microstructure development of @-SiC-A1N-Y2O3 and @-SiC-A1N-La203 materials under conditions of pressureless sintering were investigated. The Y2O3 for La203 substitution was investigated for possible sinterability improvement due to lower refractoriness of La203 - based oxynitride systems as well as lower eutectic formation in relevant oxide systems [20]. Influence of RE sintering additive type on sinterability, phase and structure formation were studied. EXPERIMENTAL Mixtures were prepared from high-purity powders of a-Sic (UF-15, H. C. Starck, Goslar, Germany), p-Sic (B10, H. C. Stark, Goslar, Germany), AlN (H. C. Starck, Germany, grade C ), Y203 ( >99.98 % purity, Aldrich Chemical Company , USA), or La203 ( >99.9% purity, Sigma-Aldrich Co., Switzerland) by attrition milling with alumina milling media in isopropyl alcohol for 4 hours at 500 RPM. The total amount of sintering aids was kept constant at 10 vol.%. The ratio of a-SiW-SiC was a variable parameter, the content of a-Sic chosen as 5wt% and 10 wF!! relative to the amount of p-Sic. The ratio AlN/Y203 was kept constant at 312, the proportion chosen from the AlN Y2O3 phase diagram [21]. The ratio A1N/La203was kept at the same value in anticipation of similarities between these two oxides behavior, the exact phase relationships in A1N-La2O3 system being unknown until present time. The slurry obtained after attrition milling was separated from the milling media by passing through a ASTM 325 sieve and subsequently dried in a vacuum rotaevaporator. Samples of the mixture were taken for granulometric analysis and specific surface determination. The former was performed by means of laser granulometry (Granulometer 1064, CILAS, France), while the latter was accomplished by BET (ASAP 2000, Micrometrics Instrument Corp., USA).The process of drying was then completed in a drying box (48 hours, 65'C). Finally, the powder was passed through a ASTM 100 sieve to crush soft conglomerates. Green bodies in the form of cylindrical pellets 20 mm in diameter and 25 mm height were prepared by consequent uniaxial pressing at 100 MPa, and cold
5 94
isostatic pressing at 200 m a . Dilatometric experiments were accomplished in a high temperature dilatometer (DIL 402 E/7, Netzsch GmbH, Germany) with a graphite resistance furnace. Sintering was accomplished in a gas-pressure furnace (Thermal Technologies, Santa Barbara, USA) with a graphite heating element. Post sintering heat treatment (annealing) was carried out in the same furnace that was used for sintering. The sintered samples were subsequently annealed at temperatures 1900'C and 195OOC in nitrogen under normal pressure for up to 16 hours in order to achieve maximal degree of p-Sic to a-Sic phase transformation. Kinetics of the phase transformation was followed up by quantitative XRD analysis. Sintered and annealed samples were characterized for weight loss, density, phase composition and microstructure. Density was determined by Archimedes method. Evolution of phase composition and phase transformation kinetics were studied by X-ray diffraction (XRD) on a Siemens D-6000 diffractometer (Ni-filtered CuKa radiation, range of detection 10-80' 2 0 ). Microstructure was studied by scanning electron microscopy (SEM) on a Phillips XL-30 and on a JEOLJXA-6400 electron microscopes with EDS, analyzing attachment. Samples for microstructure investigation were prepared by standard ceramographic procedure of multistep grinding and polishing with subsequent chemical etching with Murakami's reagent (10 g of NaOH and 10 g of K3Fe(CN)6 in 40 ml H 2 0 at 110°C) for structural elements revelation.
RESULTS AND DISCUSSION The nominal formulation of the prepared mixtures together with some of their granulometric characteristics are presented in Table 1. Table 1. Characteristics of the powder mixtures Material a-SiC P-Sic
Y2O3
La203 AlN
wt.% wt.% wt.% wt.% wt.%
Specific surface area
The prepared mixtures were of submicron fineness, and had a narrow grain size distribution (dlo= <0.1 pm, dso = 0.38-0.40 pm, dw = 0.99-1.05 pm), the values of these parameters being very close for all mixtures prepared. Therefore, a favorable sintering behavior was expected. The results of the dilatometric investigations of the densification behavior of the prepared compositions under different sintering atmosphere - flowing nitrogen and flowing argon - are presented in Fig. 1. As it can be
seen, the behavior of Sic and SiL groups of compositions is strikingly different both in terms of linear shrinkage and shrinkage rate depending on the sintering atmosphere. The influence of the a-Sic content in the initial mixture has only a slight effect on the overall shrinkage behavior causing a minor diminishing in shrinkage rate. However, if the Y2O3 is substituted for LazO3, the shrinkage behavior changes dramatically. The linear shrinkage in this case is rather low in Ar atmosphere(Fig. 1. (a)), and substantially higher in N2 (Fig. 1. (b)), compared with the one exhibited by Y2O3 doped mixtures, thus making the sintering behavior in regard of the atmosphere of sintering of these two groups of compositions inverse. It must be also noted that in Ar atmosphere the beginning of shrinkage for La203doped mixtures is shifted about a 100°C towards higher temperatures, while in N2 by atmosphere shrinkage for both groups of mixtures starts at about the same temperature of 160OOC. Even more differences have been observed with regard of shrinkage rate (Fig. 1. (c) and (d)) of the studied compositions. If while sintering in Ar the difference in shrinkage rate basically is limited to the
4.02
4.01 -
4.04
0.02
-
4.03
-
0.05
-
0.02 0.00
0.01
-
-
0.00
4.06 -
s
2
value of the rate for La203 doped compositions being sufficiently lower, and the maximum being shifted by 2OO0C towards higher temperatures, the sintering in N2 provides a much more complicated picture. It can be seen that the materials doped with La203 exhibit four peaks on the shrinkage rate curves, while the Y2O3 doped ones have only two. Such behavior suggest a rather complex sequence of phase formation on heating in the former group of materials because every minimum on the shrinkage rate curve corresponds to a crystalline phase formation, which hinders densification. Only after a temperature higher than the melting point of this phase is reached the densification proceeds. As it was mentioned before, the exact phase relationships in the AIN-La203 are not known, and the molar ratio of the components was kept identical to the one in the case of AlN-Y203 system in anticipation of close similarity of the phase relationships in both systems. However, it is now obvious that in the case of La203 the phase relations are rather different, and in order to optimize both composition of the additives and
4 . 3 4.010
4.08-
4.10:
If
4.07 4.06
4.18
ma
800
1 m
1200
1400
1600
1800
zoo0
8w
Temperature ("C)
8w
1000
1200
1400
1600
1800
2WO
1800
2000
Temperature ("C)
0.0002
-
4.0002 0.0000
6
2
0.0004:
1-
-0.0001
-
I-
53
0.WO6-
0
=
0.0000
o'oool
-o.ooo2
:
-x-Sic10
-N
'0 0 . 0 0 0 3 -
0.MM8-
P 0.0010
-
0.0012
-
0.0004
d , . , . , . , . , . , . , . , , 600
800
low
12w
1400
18w
Temperature ("C)
18W
-0.0005
Mw 600
800
1000
1200
1400
1600
Temperature ("C)
Fig. 1. Dilatometric measurements: (a) and (b) - linear shrinkage in Ar and Nz, respectively; (c) and (d) - shrinkage rate in Ar and N2, respectively.
595
the sintering procedure further research is necessary, which is the next step in the course of the present research. In the case of the Y2O3 doped materials, the first of the observed peaks on the shrinkage rate curve while sintering in N2 is tentatively attributed to initial and limited liquid phase formation with participation of surface oxides, Si02 and A1203present on S i c and AlN powder particles, respectively, with Y2O3. The phase precipitating at about 17OO0C and subsequently melting at about 185OoCwas not identified until now. It should be also mentioned that with the increase of the a-Sic content from 5 to 10 wt.% in the initial mixture the number of peaks on the shrinkage rate curve of La203 curve diminishes and approximates the curve of the Y203 doped material with the same a-Sic content. Based on the results of the dilatometric investigations the time-temperature sintering schedule for furnace sintering was developed as presented in Fig. 2. The results of the furnace sintering of the prepared compositions in nitrogen and argon are presented in Table 2.
. ISSd'C .
2ooo-
e
1500-
f
two-
t
1.5
-
1.0 M
$
- -z
5 c
0.5 'II
500-
o
f
.
0
.
,
.
, 1W
50
,
.
.
, ZW
150
I 250
Time (min.)
Fig. 2. Time-temperature-gaspressure sintering schedule. The results of post sintering heat treatment and of the p-Sic to a-Sic phase transformation kinetics are presented in Fig. 3. As it can be seen, the p-Sic to aSic phase transformation occurs more readily with the La203 additions, while Y2O3 containing materials exhibit a somewhat sluggish behavior. At present time, only a tentative explanation is possible attributing such behavior to better conditions for diffusion provided by the secondary phases formed in La203 doped materials. However, firher more detailed investigation of the processes involved is necessary. im
-
!
o
l0 0
2
4
:-r?
_.....-
#-a-----0
596
.
P
Table 2. Green density, sintered density and weight loss after furnace sintering of Sic5 and Sic10 compositions.
As it can be seen, although the weight loss in the case of SiL series is lower than for S i c series, the former materials densify poorly under conditions of sintering used. In order to optimize the La203 doped materials it is necessary to develop a special sintering schedule based on the further phase formation sequence investigation which are under way. Severe overlapping of the reflexes of p-Sic due to the polytypism of SIC makes the quantitative phase analysis extremely difficult, However, the existing three well defined a-Sic reflexes (6H-polytype) enable to quantitatively evaluate the extent of the phase transformation by normalizing the sum of the (1 1l)3c and (006)6H reflexes intensities. Using this approach a calibration curve can be obtained using the inner standard technique, i. e., a series of powder mixtures of p-Sic with fixed a-Sic additions was analyzed by XRD, and the ratio between the relevant reflexes intensities was determined. Thus obtained curve was subsequently used for the extent of phase transformation during post sintering heat treatment determination.
2.0
.
'
0
6
8
10
12
14
16
18
Annealing time, h
Fig. 3. Kinetics of p-Sic to a-Sic transformation under conditions of isothermal annealing in NZ. The microstructure of the as-sintered and thermally treated SiC-AlN-Y203 material are presented on Fig. 4. If the as-sintered sample has a fine uniaxial grain morphology, the annealed material exhibits microstructure with elongated grain morphology as a result of the p-Sic to a-Sic transformation. The grain growth occurred as well, thus providing evidence of transformation controlled grain growth mechanism.
ACKNOWLEDGEMENTS The authors would like to express their gratitude to CNPq for financial support of Dr. V. Izhevskyi's participation in this research. The present research is supported by PRONEWINEP. The authors would also like to thank FAPESP for financial support of Dr. Bressiani's participation in the Conference. REFERENCES
Fig. 4. Microstructuresof the (a) as-sintered and (b) annealed (8 hours, N2 atmosphere) SiC-A1N-Y203 chemically etched samples. CONCLUSIONS Silicon carbide can be sintered up to high densities by means of liquid phase sintering under low gas pressure with AlN-Yz03 and AlN-LazO3 sintering additives. As-sintered materials exhibited fine-grained homogeneous microstructure. Densification behavior of materials with Y2O3 and La203 additions is shown to be inverse: While the former materials densify better in Ar, the latter show better sinterability in NZ.Moreover, the densification of La203 doped materials has a complex nature most probably due to a complex phase formation sequence on heating. This phenomena needs further investigation. Transformation controlled grain growth during post sintering heat treatment of the developed materials was established. A high degree or even a complete pSic to a-Sic transformation was achieved by such techniques in reasonably short time intervals. The kinetics of the phase transformation was shown to depend on the composition of sintering additives, the use of less refractory RE dopants accelerating the process. Thus, the possibility of in-situ platelet reinforced Sic ceramics was shown. Further research aimed at the optimization of both composition and processing of the materials under development is necessary.
(1) S. Prochaska, The Role of Boron and Carbon in Sintering of Silicon Carbide. Special Ceramics 6. Edited by P. Popper. The British Ceramic Research Association, Stoke-on-Trent, U. K., (1975) 171181. (2) R. A. Cutler and T. B. Jackson. Liquid phase sintered silicon carbide Proc. Third International Symposium Ceramic Materials and Components for Engines,. Edited by V. J. Tennery. American Ceramic Society, Westerwille, OH, (1989) 309318. (3) M. A. Mulla and V. D. Krstic. Low-Temperature Pressureless Sintering of p-Silicon Carbide With Aluminum Oxide and Yttrium Oxide Additions. Am. Ceram. SOC.Bull., 70, (1991) 439-443. (4) M. A. Mulla and V. D. Krstic, Pressureless Sintering of p-Sic with A1203 Additions. J. Mater. Sci., 29, (1994) 934-938. (5) S. K. Lee and C. H. Kim. Effects of a- Sic Versus p-Sic Starting Powders on Microstructure and Fracture Toughness of S i c Sintered with A1203-Y203 Additives. J. Am. Ceram. SOC.,77, (1994) 1655-1658. (6) M. Omori and H. Takei. Pressureless Sintering of Sic. J. Am. Ceram. SOC.,65 (1982) C-92. (7) P. Padture. In-situ Toughening of Silicon Carbide. J. Am. Ceram. SOC.,77, (1994) 5 19-523. (8) W. J. Kim and Y.-W. Kim. Liquid-Phase Sintering of Silicon Carbide. J. Kor. Ceram. SOC.,32, (1995) 1162-1168. (9) W. Bocker, H. Landfermann, and H. Hausner. Sintering of a-Silicon Carbide with Additions of Aluminium. Powd. Metall. Intern., 11, (1979) 8385. (10)K. Negita. Effective Sintering Aids for Silicon Carbide Ceramics: Reactivities of Silicon Carbide with Various Additives. J. Am. Ceram. SOC.,69, (1986) C-308-C3 10. (11)K. Y. Chia, W. D. G. Boecker, R. S. Storm. Silicon Carbide Bodies Having High Toughness and Fracture Resistance and Method of Making Same. United States Patent, 5 . 298.470, USA, 1994. (12)M. Keppeler, H.-G. Reihert, J. M. Broadly, G. Turn. I. Wiedmann, and F. Aldinger. High Temperature Mechanical Behavior of Liquid Phase Sintered Silicon Carbide. J. Eur. Ceram. SOC.,18, (1998) 521-526.
597
(13) R. Ruh, A. Zandvil. Compositions and Properties of Hot-Pressed Sic-AlN Solid Solutions. J. Am. Ceram. SOC.,65, (1982) 260-265. (14)A. Zangvil, R. Ruh. The Si3A&N4C3 and Si3AI5N5C3 Compounds as SIC-AlN Solid Solutions. J. Mat. Sci. Lett., 3, (1984) 249-250. (15)A. Zangvil, R. Ruh. Solid Solutions and Compositions in the Sic-AlN and SIC-BN Systems. Mat. Sci. Eng., 71, (1985) 159-164. (16)A. Zangvil. R. Ruh. Phase Relationships in the Silicon Carbide-Aluminum Nitride System. J. Am. Ceram. SOC.,71, (1988) 884-890. (17)J.-K. Lee, H. Tanaka, S. Otami. Preparation of S ic - AlN Composites by Liquid-Phase Sintering and Their Microstructure. J. Ceram. SOC. Japan, 103, (1995) 873-877. (18)L. S. Singhal, H.-J. Kleebe. Corehim Structure of Liquid-Phase-Sintered Silicon Carbide. J. Am. Ceram. SOC.,76, (1993) 773-776. (19)L. S. K. Lee, Y. C. Kim, C. H. Kim. Microstructural Development and Mechanical Properties of Pressureless-Sintered Sic with PlateLike Grains Using A1203-Y203Additions. J. Mater. Sci., 29, (1994) 5321-5326. (20)U. Kolitsch, H. J. Seifert, and F. Aldinger. Phase Relationships in the Systems RE203-A1203-Si02 (RE= Rare Earth Element, Y, and Sc). J. Phase Equilibria, 19, (1998) 426-433. (2 1) A. Jeutter. Untersuchung der Phasenbezieung im Sysytem AluminiumnitridYttriumoxid. Diplomarbeit an Universitat Stuttgart, (1993).
598
SLIP CASTING OF ATZ CERAMICS E. Gregorova, J. Havrda, W. Pabst, K. Kunei Department of Glass and Ceramics, Institute of Chemical Technology (ICT) Prague, CZ - 166 28 Prague 6, Czech Republic
ABSTRACT The preparation of ATZ ceramics by slip casting (of two commercial powder types, labeled TZ-3Y20A and ATZ 80) is studied. Low-viscosity slips (20-60 mPas) with 70 wt.% and 76 wt.% solids content, respectively, are prepared and cast into plaster molds. The particle size is found to be lower for TZ-3Y20A than for ATZ 80. According to mercury porosimetry the average pore size in dried green bodies is 22-24 nm and the open porosity approx. 48 %. After firing, all ATZ bodies investigated exhibit bulk densities close to theoretical (approx. 5.45 g/cm3) and three-point bending strengths of up to 1030 MPa and 995 MPa for TZ-3Y20A and ATZ 80, respectively. Peak values of flexural strength are achieved at firing temperatures of 1520-1530 "C, while higher firing temperature are detrimental to mechanical strength as a consequence of excessive grain growth.
INTRODUCTION ATZ ceramics (alumina-toughened zirconia) is one type of ceramics in the alumina-zirconia composite system. The main component is zirconia, and small alumina particles (< 1 pm after firing) are finely dispersed in this zirconia matrix. Although the fracture toughness reported for this type of ceramics is not as high as that of zirconia-toughened alumina (ZTA) or monolithic yttria-doped zirconia, ATZ ceramics exhibit the highest bending strengths known for ceramics at room temperature (up to 1800-2400 MPa for hot isostatically pressed ceramics) and - in contrast to monolithic TZP (tetragonal zirconia polycrystals) they do not undergo surface degradation (i.e. t-m transformation of zirconia) at moderately elevated temperatures (1 50-250 "C), especially in humid conditions (water vapor) [ 1-61. Traditionally, peak strength values in the order of 2000 MPa have only been achieved by hot isostatic pressing. In recent years, however, also other routes, including slip casting have been successfully performed for ATZ powder systems, although it seems that the strength values achievable for these types of ceramic materials shaped by slip casting are significantly lower, viz. in the order of 1000 MPa [791. In this work the preparation of ATZ by slip casting is studied. Two commercially produced types of ATZ powder mixes (labeled TZ-3Y20A and ATZ 80, respectively) are considered, both containing
approx. 20 hm.% alumina in a yttria-doped (3 mol% related to zirconia) zirconia matrix. The powders are characterized (in addition to the suppliers' information), the suspensions are characterized with respect to rheology and the bodies prepared are characterized with respect to microstructure and mechanical strength.
EXPERIMENTAL Two types of commercial powder mixes were used in this work (grade TZ-3Y20A by TOSOH Corp. / Japan and grade ATZ 80 by DAIICHI Kigenso Kagaku Kogyo Co., Ltd. I Japan), both containing approx. 20 wt.% alumina and approx. 80 wt.% yttria-doped zirconia. The characteristics of the powder are listed in Table I. Additionally to the suppliers' information concerning average size [lo], the particle size distribution was characterized by dynamic light scattering (Nicomp 3 80 Submicron Particle Sizer, Nicomp Particle Sizing Systems / USA), since the particle size information, especially concerning the Tosoh powder TZ-3Y20A, is insufficient (for pure zirconia powders the same supplier mentions an average particle size of 0.4 pm [ 111, but other authors [7-81 mention 0.3 pm for the alumina and 0.1 pm for the zirconia fraction in the ATZ powder mix). Table I . Characteristics of the powders used in this work according to suppliers' data [ 10-111
Zr02 content [wt.YO] A1203cont. [wt.YO] Y2O3 cont. [wt.%] specific surface rm2/gl average particle size [pm]
TZ-3Y20A (TOSOH) 75.0 f 0.3
ATZ 80 (DAIICHI) 76.0 k 0.1
21.0 k 0.3
19.5 k 0.1
3.95 k 0.05
4.0 k 0.1
15+2
15+2
---
0.3
599
chosen for the powder TZ-3Y20A to yield a suspension of comparable fluidity, cf. the similar findings in [12]. The suspensions were mixed, homogenized in a planetary mill (Pulverisette 6, Fritsch / Germany) in a zirconia bowl with zirconia balls, deaired, sonicated and again deaired. The rheology of the suspensions was characterized by rotational viscometry (RVl, Haake / Germany) for shear rates between 0 and 1000 s-'. The measured flow curves were fitted by a power law equation, and apparent viscosities were determined for selected shear rates. The kinetics of body formation was determined by measuring the thickness of the body forming in a one-dimensional situation on a plaster support after certain time intervals. The as-prepared suspensions were cast into plaster molds to obtain cylindrical bodies of diameter 5 mm and length 60 mm. After demolding and drying the samples were subjected to firing in an electrical furnace at atmospheric pressure to different temperatures between 1500 and 1570 "C. The firing schedule was 2 "C / min (ramp up), dwell 120 and 180 min, respectively, followed by free cooling in the furnace. Dried green bodies were characterized by mercury porosimetry (Poresizer 9320, Micromeritics / USA) and sintered bodies by the Archimedes method. Flexural strength was measured on sintered bodies by three-point bending (span 40 mm). Polished sections were prepared from sintered specimens, thermally etched, and investigated by scanning electron microscopy.
3Y20A (58 nm and 232 nm) than for ATZ 80 (109 nm and 367 nm). Since the peak area ratios are 0.56 to 0.44 and 0.51 to 0.49, respectively, it is not possible to assign the two peaks separately to alumina and zirconia. The measurements, however, show the presence of nanosized fractions in both powders, and there is no indication for ATZ-80 to be finer than TZ3Y20A. Figures 1 and 2 show the flow curves and apparent viscosities of suspensions containing 70 wt.% and 76 wt.% of TZ- 3Y20A and ATZ 80, respectively.
-
36
10
200
0
400
Y
600
800
1000
[llsl
Figure 1. Flow curve and apparent viscocity p versus shear rate y of a suspension with 70 wt.% TZ-3Y20A 35
1
-
10
RESULTS AND DISCUSSION
CI
0
n
E
Characterization of the Powders and Suspensions The mentioned average particle size of 0.4 pm given by the producer [ 1 11 for Tosoh zirconia powders represents a median equivalent sphere diameter determined by centrifugal sedimentation on a Horiba Capa 700 photocentrifuge, for the Daiichi type ATZ 80 the measuring principle is not given. The average size is that of primary particles, which are (at least in the case of the zirconia fraction of these powders) composed of crystallites whose size is about one order of magnitude lower, i.e. tens of nanometers. Apart from that, spray-dried granules (agglomerates) of considerably larger size (approx. 50 pm) are known to be present in the Tosoh powder [ 1 11. It can be seen in Table I that the alumina content is higher in the Tosoh powder. This finding has been confirmed by X-ray phase analysis. Beside the particle size information given by the producer, the particle size distribution for these two powder types has been measured by dynamic light scattering (Nicomp 380 Submicron Particle Sizer, Santa Barbara / USA). The volume-weighted particle size distribution shows a bimodal frequency curve in both cases, with the modi being clearly lower for TZ-
600
I
I
0
200
400
800
800
1000
Y [Ilsl
Figure 2: Flow curve and apparent viscocity p versus shear rate y of a suspension with 76 wt.% ATZ 80 Both suspensions exhibit pseudoplastic (shearthinning) behavior without yield, in the first case with a slight hysteresis (i.e. thixotropy). The flow curves (shear stress r versus shear rate y ) can be fitted with Ostwald-DeWaele's "power law":
r
= Ky"
where the consistency coefficient K attains values of 0.0859 and 0.0686 and n is 0.772 and 0.886 for the suspensions with 70 wt.% TZ-3Y20A and 76 wt.% ATZ 80, respectively. Apparent viscosities are very
low, approx. 20-60 mPas, cf. Figs.1 and 2. It is evident that even higher solids contents would be appropriate for slip casting. In order to achieve comparable apparent viscosities for both types of suspensions, the powder type TZ-3Y20A requires, however, a higher deflocculant content (0.9 wt.% versus 0.6 wt.% for ATZ 80, based on solids).
For the interpretation of these curves (and the average pore size of 22 nm and 24 nm it has to be reminded that mercury porosimetry assumes a (rather unrealistic) pore system of cylindrical pores. When pores with varying cross-section are measured, the percentage of small-sized pores (according to the Washburn equation) is always overestimated in the volume-weighted distribution.
Fig. 3 shows the kinetics of body formation on plaster supports. The rate of body formation is clearly lower for the (less concentrated) TZ-3Y20A suspension.
6
.
I
I
t
. 10
I
.- .*---Tz-3y-m '
0
opol
am
1
Quivalent pore diameter [pm] 0
m
P
10
40
m
Time [min]
Figure 3. Body formation kinetics of the suspension with 70 wt.% TZ-3Y20A and 76 wt.% ATZ 80
Figure 4. Pore size distribution (cumulative curve) measured by mercury porosimetry for dried green bodies of TZ-3Y20A
Characterization of the Green Bodies The as-prepared green bodies have been characterized in the dried state with respect to bulk density, apparent (skeletal) density, (open) porosity P and average pore size (equivalent cylinder diameter) by mercury porosimetry. Table I1 lists the results.
Bulk density
P
Skeletal density
Average [%I poresize [nm 48.4 I 22 47.5 I 24
O.ooC
am
41
1
Quivalent pore diameter [pm]
1 Figure 5. Pore size distribution (cumulative curve) TZ-3Y20A ATZ80
I I
2.72 2.81
I I
5.27 5.36
I I
Since in the green state closed pores are practically absent, the skeletal density values correspond to the theoretical density of the powder mix before firing. Xray studies of the Tosoh powder TZ-3Y20A [13] have revealed that in the as-supplied powder the content of monoclinic zirconia is about 20 wt.% (related to the total zirconia content), which corresponds to a theoretical density of about 5.3 g/cm'. Figs.4 and 5 show the pore size distribution curves (cumulative curves), also measured by mercury porosimetry, of the dried green bodies.
measured by mercury porosimetry for dried green bodies of ATZ 80
Characterization of the Sintered Bodies The as-fired samples were characterized by the Archimedes method (bulk density and apparent porosity) and their flexural strength was determined in three-point bending tests. Tables 111 through VI list the results in dependence of the firing temperatures and the dwell times used for firing. The apparent porosity (not shown in the tables) was in all cases < 0.05 %, i.e. taking into account the statistical errors of measurement - all bodies can be considered as fully sintered for the temperatures in question.
60 1
Both types of bodies can be considered as densely sintered for all temperatures between 1500 and 1570 "C. Assuming all zirconia to be present as tetragonal phase, the theoretical density would be 5.52 g/cm3, while when admitting approx. 12.5 wt.% of monoclinic phase (related to the total zirconia content) it would be 5.45 g/cm3[13]. Table 111. Bulk density of fired bodies of TZ-3Y20A in dependence of the firing temperatures and the dwell times used Firing temperature
1520 1530 1550 1560
I
Firing temperature ["C]
I
Dwell time 120 min
Dwell time 180 rnin
5.44
5.45
5.45
5.45
I I
1500 1510 1520 1530 1540 1550 1560 1570
I
I
5 45
I
5.45
Bulk density [g/cm3] Dwell time 120 min
I
5.42
I
- . .-
5.44 5.45 5- .-46 .5.47 5.47 5.47
---
Dwell time 180 min
-_-
5.44 5.45 I
I
-_-
5.47
--5.47 5.47 ~~
Table V. Flexural strength of fired bodies of TZ-3Y20A in dependence of the firing temperatures and the dwell times used
602
Firing temperature
-
Flexural strength [MPa] Dwell time
Dwell time
1500 1510 1520
903 915 995
---
1550 1560 1570
708
~~
750
xo 1
Bulk density [g/cm']
Table IV. Bulk density of fired bodies of ATZ 80 in dependence of the firing temperatures and the dwell times used
I
Table VI. Flexural strength of fired bodies of ATZ 80 in dependence of the firing temperatures and the dwell times used
---
703 699
Prolonging the dwell time from 120 min to 180 min does not increase the bulk density and deteriorates flexural strength. The optimum firing temperature for maximum strength (for both types at approx. 15201530 "C) does not coincide with the temperature for maximum bulk density (>1550 "C in both cases). The strength decrease due to an excessively high firing temperature is especially pronounced for the ATZ 80 samples, evidently as a consequence of unwanted grain growth. Peak strength values are 1030 MPa and 995 MPa for TZ-3Y20A and ATZ 80, respectively. Interestingly, the strength values of TZ-3Y20A samples are slightly higher (than those of ATZ 80 samples), despite the fact that many of these (cylindrical) samples show a central hole, which seems to be a relict of lower initial solids content in the suspension and the lower rate of body formation. These central holes are visible on the polished sections perpendicular to the cylinder axis. The absolute strength values for fired bodies are for both types of initial comparable to those achieved by Salomoni, Esposito and coworkers [7-81, who published 858 MPa for ATZ prepared by cold pressing and 1057 MPa for ATZ ceramics made by pressure slip-casting. The bulk densities for corresponding temperatures are slightly higher for ATZ 80 bodies than for TZ-3Y20A bodies. Probably this is a consequence of the higher initial concentration of the suspensions with this powder type, but a definitive statement can only be made when the tetragonalmonoclinic ratio is known for these samples. It can be said, however, that - for the bodies investigated - the overall influence of the initial concentration of the casting suspension on the microstructural characteristics of the final bodies is rather small. The microscopic investigation by SEM (documented on the poster presented at this Conference) reveals the less homogeneous microstructure and the larger grain size (especially of the alumina) of the ATZ 80 samples in comparison to the TZ-3Y20A samples and confirms the aforementioned hypothesis that excessive grain growth
is responsible for the strong strength decrease for higher firing temperatures.
CONCLUSION The preparation of ATZ ceramics by slip casting (of two commercial ATZ powder types, TZ-3Y20A. TOSOH and ATZ 80, DAIICHI) has been studied. Low-viscosity slips with 70 wt.% and 76 wt.% solids content and 0.9 and 0.6 wt.% of deflocculant have been prepared and cast into plaster molds. In contrast to suppliers' data, according to the measurements presented here the particle size is clearly lower for the TZ-3Y20A powder than for ATZ 80. This correlates well with the higher amount of deflocculant needed to prepare a suspension of comparable apparent viscosity. This fact, aside with the lower solids content, the differences in chemical composition (and thus surface characteristics) of the powders and the slower body formation kinetics can be responsible for the occurrence of central holes in the cylindrical samples. According to mercury porosimetry the average pore size (equivalent cylindrical diameter) is 22-24 nm and the dried green bodies exhibit an open porosity of approx. 48 YO. After firing, all ATZ bodies investigated exhibited bulk densities close to theoretical (approx. 5.45-5.52 g/cm3) and three-point bending strengths of up to 1030 MPa and 995 MPa for TZ-3Y20A and ATZ 80, respectively. Peak values of flexural strength are achieved at firing temperatures of 1520-1530 "C, while higher firing temperatures - although leading to a further increase in bulk density - are detrimental to mechanical strength as a consequence of excessive grain growth.
Partially Stabilized Zirconia, Key Engineering Materials 161- 163 (1999) 307-3 10. [6] R. Chaim, Pressureless Sintered ATZ and ZTA Ceramic Composites, J. Mater. Sci. 27 (1992) 5597-5 602. [7] A. Salomoni, A. Tucci, L. Esposito, I. Stamenkovic, Forming and Sintering of Multiphase Bioceramics, J. Mater. Sci. (Materials in Medicine) 5 (1994) 651-653. [8] L. Esposito, A. Salomoni, I. Stamenkovic, A. Tucci, Processing of Zr02-A1203 Powders: Consolidation and Characterization of Final Products, Special Meeting on Biomaterials Rimini 1992 (I. Stamenkovic, J. Krawczinski, eds.), Publ. Forschungszentrum Jiilich (1994), 37-45. [9] I. Stamenkovic, A. Salomoni, Colloidal Shaping and Sintering of Mixed Alumina and Partially Stabilized Zirconia, Cer. Acta 10 (1998), 11-17. [ 101 Data sheet Zirconia Oxide Products, Daiichi Kigenso Kagaku Kogyo Co., Ltd., 1997. [ I I] Data sheet Zirconia Powder, Tosoh Corp., 1997. [12] K. KuneS, J. Havrda, K. Hronikova, E. Gregorova, W. Pabst, Stabilization of Bioceramic Suspensions Prepared from Alumina-Containing Zirconia Powders, Ceramics-Silikaty 44 (2000) 1-8. [I31 W. Pabst, J. Havrda, E. Gregorova, B. KrEmova, Alumina Toughened Zirconia Ceramics Made by Room Temperature Extrusion of Ceramic Pastes, Ceramics-Silikaty 44 (2000) (to appear).
Acknowledgement: This study was part of the research project CEZ:MSM 2231 00002 "Chemistry and Technology of Materials for Technical Applications, Health and Environment Protection" and supported by grant MPO No. FB-CV/64/98.
REFERENCES [ I ] W. E. Lee and W. M. Rainforth, Ceramic Microstructures, Chapman & Hall, London (1994), 534-538. [2] R. W. Cannon, Transformation Toughened Ceramics for Structural Applications, in: Structural Ceramics (J. B. Wachtman, ed.), Academic Press, Boston ( 1 989), 195-228. [3] M. M. Schwartz, Handbook of Structural Ceramics, McGraw Hill, New York (1992), 3.35-3.38,4.48. [4] Z. Panek, KonStrukEna keramika (Structural Ceramics, in Slovak), SAV Bratislava (1 992), 4 1 45, 124-125. [5] B. Zhang, T. Isobe, S. Satani, H. Tsubakino, The Effect of Alumina Addition on Phase Transformation and Mechanical Properties in
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The Preparation for Sintered Body of Ce02 Based Complex Oxide in Low Temperature Solid Oxide Fuel Cells Using Colloidal Surface Chemistry Yong-Sin Hwang* and Sung-Churl Choi Dept. of Ceramic Engineering, Hanyang University, Seoul 133-791 Korea ABSTRACT In this study. the dispersion stability of Ce02 based complex oxide was studied, and density, porosity, and microstructure of green body were investigated using colloid surface chemistry to manufacture the Gd2O3 doped CeOz solid electrolyte in an aqueous system. To prepare the stable slurry for slip casting, the dispersion stability was examined as a function of pH using ESA(e1ectrokinetic sonic amplitude) analysis. The dynamic mobility of particles was enhanced when anionic and cationic dispersant were added the amount of 0.5 wt% respectively, but slurries pH didn't move to below 6.0 because of the influence of dopants. This phenomenon also appeared in the Ce02-Y203 and Ce02-Sm203 systems, so it could be inferred that rare earth dopants such as Gd2O3, SrnzO3 and Y2O3 not only have the similar motion with changing pH in an aqueous system but also can be dissolved in the range of pH 6.0 -6.5. In Ce02-Gd203 system, when the anionic dispersant was added the amount of 0.5 wt% and the slurry pH was fixed at 9.5, the green body density was 4.07 g / d , and the relative density of sintered body was 95.2 YO. It could be inferred from XRD analysis that Gd3' substituted into Ce4+ site because there was no free Gd2O3 peak
INTRODUCTION Generally, ftiel cell can be classified into alkaline fuel cell(AFC), phosphoric acid fuel cell(AFC), molten carbonate fuel cell(MCFC), and solid oxide fuel cell(S0FC) according to a kind of electrolyte. SOFC is more economical than other fuel cells and has a merit of using solid electrolyte. Target SOFC lifetimes are the order of lo4 10' hours [ 1 1. In solid electrolyte for SOFC, 21-02 and CeO2 system had been researched. Especially CeO2
-
system had been expected to be solid electrolyte material for low operation temperature because it has a high ion conductivity in low temperature. But Ce02 system is hard to be sintered because of its low self diffusion coefficient, and despite resintering at 1800 "c 1900C , it is difficult to have more than 90% relative density. Furthermore it shows low density of sintered body as the size of dopant increases. In the past, nearly all materials research activities on the fabrication of SOFC have been devoted to optimizing electrochemical, thermal and microstructural properties, but comparatively little research has been established on the process variables. So there is a need to be investigated for advanced substitution between matrix and dopant materials and control of sintering characterization. To overcome above problems, the research for the interaction between process additives and ceramics particles is necessary, and it is important to have dispersion stability for densed green body from control of stress between CeO2 and rare earth dopants. In this study, to optimize the microstructure using colloidal surface chemistry, not simple control of starting material, Ce02 complex oxide system that had dispersion stability through the change of surface chemistry of CeO2 as the addition of rare earth dopants(Gd203, Sm203, and Gd2O3) was investigated in the aqueous system. And from the condition of dispersion stability of CeO2-GdzO3 system, the sintered body was prepared, and it was compared the phase and microstructure of sintered body prepared by aqueous slip casting to those of the sintered body prepared by uniaxial press.
-
EXPERIMENTAL The average particle size and purity of CeO2 was 4pm and 99.9%(AMR, Canada)
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respectively. Gadolinium oxide(Gd203, 99.99%, NONE-POULENC), Samarium oxide(Sm103, 99.99% RHONE-POULENC) and Yttrium oxide(Y203, 99.99% RHONEPOULENC) was selected as cationic dopants for CeO2. Betzl190(5wt% diluted dimethyamine-epichlorohydrin solution, copolymer, Trevose Co., U.S.A) and D3019(5wt% diluted solution, ammonium salt, Rhom and Haas Co., Phliadelphia, U S A ) was used as the cationic dispersant and anionic dispersant respectively. Deionized(D1) water( 18.2MsZcm, manufactured by Millipore Milli-Q plus) was used as a solvent, and PVA(A1drich Chemical Company. Inc., U.S.A) and PEG(S1iinyo Pure Chemicals Co., Ltd, Japan) were introduced as a binder and plasticizer, respectively. The experimental procedure used for the preparation of CeOzGd203 electrolytes fabricated by aqueous slip casting is illustrated in Fig. 1. Solvent Powder +
+Dispersant
Green Body &
7---Sintered Body
Fig. 1. Schematic of the preparation of the GdlOJ doped CeOl electrolyte by aqueous slip casting.
Because Ce02-Gd203 system not only has the higher oxygen ion diffusion [ 2 3 but also showed better characterization than other Ce02 based complex oxide in the relative density, mechanical strength, and microstructure, Ce02-Gd203 electrolyte was chosen for slip casting. Initially, the dispersants and deionized (DI) water( 18.2h.IsZcm, Millipore Milli-Q plus)
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were mixed and then the ceramic powders were added to the solvent mixture. and subsequently ball milled for 12 hrs. After adding the binder(PVA) and the plasticizer(PEG), the sluny was further ball milled for 12 hrs to obtain a well-dispersed slurry. To evaluate the dispersion stability of the slurry for Ce02 based complex oxide according to the kind of dispersants and concentration of dispersants, electrokinetic sonic amplitude(ESA-8000, Matec Applied Science, U.S.A) was used to measure the dynamic mobility of the particles. The rheology characterization of the slurries was measured by viscometer(Haake RS75, Rheostress, Germany). The chemicals used to adjust the pH(l.ON N&OH, and HNO3) were supplied by Duksan Pure Chemical Co., Ltd. The pore size and distribution of green body were measured by porosimeter(Autoscan-25, 60, Quantachrome Corp, USA). The density of green body was obtained from the geometric volume and mass of green body, and the electrolyte was sintered at 1600°C for 4hrs. The apparent density of the sintered body was introduced by ASTM C-1120. The analysis of phase was executed X-ray diffractometerp-Max 2C, Rigaku, Japan) under the condition of 30 KVand 40mA using Cu Ka target in the 28 range of 10" -80" . The microstructure of the electrolyte was investigated through SEM(Hitachi, Japan). And to compare the stability of phase stability and microstructure, the sintered body was prepared by uniaxial press.
EXPERIMENTAL RESULTS Fig.2 shows the results of dynamic mobility change for Ce02 using ESA. The ESA utilizes the electroacoustic effect. If an alternating electric field is applied to a colloidal slurry, the particles move back and forth at a velocity that depends on their size, zeta potential and frequency of the applied field. As the particles move, they produce an acoustic wave and this effect is known as the electrokinetic sonic amplitude(ESA). If the
applied electric field and velocity of particle E and V, respectively, the dynamic mobility( p ) of spherical particle can be defined as belou- [ 3 1. are
V= p
E
orthofenites((RFe03, R=Yb, Y, Sm) in H3PO4 and HBr [ 8 ] were reported to be dissolved. Especially, the dissolved polyvalent Y” ion acts as a counter ion in the solution which decreases the range of the double layer repulsion, thus causing the slurry to flocculate [ 7 1.
Also if the zeta potnetid(< ) is low, the relationship between dynamic mobility and zeta potential is below [ 3 1.
is the permittitivity of slurry and G is the value calculated from the relationship between the thickness of electrical double layer and particle diameter. The obtainedlEP for the investigated Ce02 powder was equal to earlier reported[ 4 3 values(pH.i,p = 6.2), and the same results is showed with the change of solid loading from 0.05 to 2 ~01%. E
-2
‘
5
a 9 Suspension pH
io
7
6
11
(A)
m
-2
I
5
6
7
8
9
1
0
Suspension pH 1
I
4
0
s.rP..rIo.
7
a
)
l
o
l
l
pn
Fig. 2. The dynamic mobility change of particles as a function of suspension pH for CeO2.
The change of dynamic mobility as the function of slurry pH for Gd203, Sm2O3, and Y203 is shown in the Fig. 3 (A), (B), and (C). In these case, the cationic and anionic dispersant were added amount of 0.1 and 0.5 wt% and the solid loading on the sluny was 2 ~01%.It must be noted that the pH of GdzOj, SmlO;, and Y203 couldn’t be move below pH 6.5 and from this result, it can be concluded that Gd203- Sm2O3, and YzO3 are dissolved in the range of pH 5.0 6.0 in the aqueous system. In the previous work, Y203 [ 5-7 ] in the aqueous and
-
r
“E
11
1 Suspension PH
(C)
Fig. 3. The dynamic mobility change of particles as a function of suspension pH for various rare earth; (A) Gd203,(B)Sm203,and (C) Y 2 0 3system.
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Fig. 4 shows the change of dynamic mobility as the function of slurry pH for 20 mol% GdlO;, SmlO; and Y 2 0 3 doped Ce02 in the aqueous system. The dissolution plienoineiia of rare earth group was be also investigated, and overall change of dynamic mobility was similiar through the examined pH range. The dynamic mobility of the particles is also reflected in the rheological behaviorpig. 5 ) with a change of viscosity between pH 7.4 and 10.0. The green body density with the change of dispersant and solid loading is shown in the Table. 1. The green body density was proportionalto solid loading because at the lower solids loading, the higher fluidity of the slips promotes segregation phenomena and particles do not pack homogeneously [ 9 1. When the solid loading was 24 vol%, the green body density was 4.07 g/d.
5
6
7
8
9
10
5
6
7
6 9 1 Suspension pH
0
I
Fig. 4. The dynamic mobility change of particles -as a function of suspension pH for various CeOz complex oxide; (A) CeOz-Gd2O3,(B) CeO2-SmZO3 and (C)Ce02-Y203system.
When the solid loading was 26 vol%, the green body density was decreased slightly because of the interference of the electrical double layer and collisions of each particles 10 I.
J
11
Suspension pH I 10
100
x)
log(Shsar rate)
Fig. 5. Viscosity of 20moI% Gd2@ doped CeOZfor different pH value titrated before aging.
-2 ; 5
6
7
8
9
Suspension pH
10
11
I
Solid loading
Green density
(vole/)
(dcrn3
GDCl
20
3.22
GDC2
18
3.35
GDQ
20
3.77
GDC4 GDC5
I
24 26
1
4.07
1
3.99
Table 1. Solid Loading and Density of Green Body.
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!b;;
Fig. 6 shows the pore size distribution of * s e e n body using 0.5 wt?hD-3019 with the solid loading of 24, and 26 ~01%. a?
:
--Id%Qf&q-.-
'
1:
t*q*.5lt-
a$
z...:
1:
am
-
. .:_
.: z..
..
....
! i
a !..
...
j
!
0%
lo
1
a1
Qol
IC400.(rmcnetes) Fig. 6. Pore size distribution of green body added by D-3019 with different solid loading
The phase of solid solution is analysized to compare the phase stability to that of the electrolyte fabricated by uniaxial press in the Fig. 7. In the Fig. 7 (A), it seems not to be formed cubic flourite structure because there are several Gd203 peaks. In case of Fig. 7 (B), it can be inferred that Gd3' is substituted at the Ce4' site stably. Fig. 8 shows the microstructure of sintered body sintered at 1600 "c for 4 hrs. In the case of the electrolyte prepared by uniaxial press (Fig. 8 (A)), pores are distributed in the grain and grain boundary impartially, so it seemed that the sintering was happened the first at the regime where the agglomeration was intensed [ 11 ] contrary to the electrolyte prepared by aqueous slip casting (Fig. 8 (B)).
SUMMARY In this study, the surface characterization was investigated as the change of pH and the amount of dispersants to have the stable phase and control the sintering behavior by the homogeneous doping of the rare earth dopants.
Fig. 7. X-ray diffraction pattern for 20mol% Gd203doped CeOz manufactured by (A) uniaxial press (B) aqueous slip casting.
The porosity and density of the green body, phase analysis and microsturcture of the sintered body were examined. It is inferred from ESA analysis that rare earth elements can be dissolved in aqueous. When CeOzGdzO3 slurry was prepared with the 0.5 wtYo anionic dispersant(D-3019), pH 9.5, and 24 vol% loading, the density of green body was up to 4.07 g/cm3. Compared to the sintered body manufactured by uniaxial press, the sintered body manufactured by aqueous slip casting had more substitutional stability.
609
Joseph M. Steigerwald, Shyam P. Murarka, Ronald J. Gutmann, " Chemical Mechanical Planarization of the Microelectronk Materials," pp. 124125, John Wiley & Sons, Inc., New York, 1997. V. ( 5 ) A. Hackley, TJ. Pa& B. H. Kim and S. G. Malghan, " Aqueous Processing of Sintered Reaction-Bonded Silicon Nitride: I , Dispersion Properties of Silicon Powder," J. Am. Ceram. Suc., 80[7], pp. 1781-1788 (1997). (6) F. Y. HO and W. C. J. Wei, " Dissolution of Y ttrium Ions and Phase Transformation of 3Y-TZP Powder in Aqueous Solution," J Am. Ceram. SOC., 82[6], pp.1614-1616 (1999). ('7) J. C. Farhas. R. Moreno. J. Requena and J. S. Moya. " Acid Basic Stability of Y-TZP Ceramics," Mater. Sci. Eng., A109, pp.97-99 (1989). (8) E. Hartmann and E. Beregi, " Dissolution Forms of Rare-Earth Orthofemtes," J. Crystal Growth, 166, pp.109-111 (1996). (9) J. M. F. Ferreira and H. M. M. Diz, " Effect of Solid Loading on SlipCasting of Silicon Carbide Slurries," J. Am. Ceram. Soc., 82[8], pp.1993-2000 (1999). (10) J. M. F. Ferreira, " Role of the Clogging Effect in the Slip Casting IEWO. Ceram. SOC., 18, Process", . pp.1161-1169 (1998). (11) Randall M. German, " Sintering Theory and Practice," pp.155-161, John Wiley & Sons, Inc., New York, 1996. (4)
(B)
Fig. 8. The scanning electron micrographs image for 20mol% Gd103doped CeOt sintered at 1600c for Jhr manufactured by ,(A) uniaxial press and (B) aqueous slip casting.
REFERENCES A. Selcuk and A. Atkinson. " Elastic Properties of Ceramic Oxides Used in Solid Oxide Fuel Cells(SOFC)," J. Euro. Ceram. SOC., 17, pp.1523-1532 (1 997). M. Kamiya, E. Shimada, Y. h a , M. Komatsu, and H. Haneda, " Intrinsic and Extrinsic Oxygen Diffusion and Surface Exchange Reaction in Cerium Oxide," J Electrochem. SOC.,147[3], pp. 1222-1227 (2000). O'Brien, R. W. " Determination of Psrticle Size and Electric Charge," U.S. Patent 5,059,909, 1991.
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PROCESSING AND PROPERTIES OF Tic-Ni,Al COMPOSITES T. N. Tiegs*, J. L. Schroeder", P. A. Menchhofer", F. C. Montgomery*, D. L. Barker*, F. Goranson** and D. E. Wittmer** (*) Oak Ridge National Laboratory
Oak Ridge, TN 37831-6087 USA (**) Southern Illinois University Carbondale, IL 62901 USA
ABSTRACT Tic-Ni,AI composites have properties suitable for engine applications. Powder compacts with binder contents from 30-50 vol. % were fabricated by pressureless sintering under vacuum followed by low gas pressure isostatic pressing. The final microstructure consisted of a 'core-rim' structure with T i c cores surrounded by Ti(Cr,Mo,W)C rims. Room temperature flexural strengths ranged from S O 0 MPa up to 1150 MPa. These strengths were retained or actually increased up to temperatures of 500°C. The fracture toughness values were all >12 MPadm and as high as 27 MPadm at Ni3Al contents from 30 to 50 vol. %, respectively.
INTRODUCTION Previous studies have shown that Tic-Ni,AI composites have an excellent combination of strength, fracture toughness, hardness and corrosion resistance [ 1-111. As a result, there is interest in using these types of materials for wear applications in diesel engines. Materials of interest contain 30-50 vol. % of Ni,Al alloy as a binder phase because these levels have thermal expansion characteristics similar to the other steel components in the engines. Because of the ductile nature of the metal, extensive toughening is obtained by local plastic deformation in the aluminide phase, which means these materials can exhibit very high mechanical reliability (e.g., Weibull modulus > 20 in preliminary studies) [3]. In addition to the high toughness, high fracture strengths in excess of 1 GPa can be observed. Nickel aluminide (Ni,Al) is an ordered intermetallic and has several attributes that make it attractive for applications in ceramic-metal composites. Its mechanical behavior is unusual in comparison to other alloys, in that the yield strength increases with increasing temperature up to about 800°C [9,10]. Typically, ordered intermetallics are brittle and act more like ceramic materials. However, it was discovered that minor amounts of boron allow Ni,Al to become ductile. These properties encouraged the early research on Ni,Al as a binder phase for ceramic particulate composites. Because these are composites, the physical properties of these Ni,Al-based cermets can be modified somewhat by changing relative amounts of the constituent phases. For example, the thermal expansion coefficients of the TiCNi,Al system, can range from 7 up to 15 x 10-6/oCby manipulating the respective volume contents of the
different phases. The expansion can also be tailored even further by altering the binder phase composition. Initial work also showed the aluminide binder phases provide good oxidation and corrosion resistance [1,11]. In addition, the composites are normally non-magnetic, however, with the appropriate substitution of Fe into the Ni,AI structure (about 220 atom. %), the materials become soft magnetics which may have advantages in particular applications. Finally, these composites exhibit adequate electrical conductivity so they can be shaped by electrical discharge machining (EDM). This may be a significant benefit in the manufacture of complex shapes where much of the cost is involved in grinding operations. Much of the early work involved hot-pressing to obtain high density materials for mechanical property testing. Subsequent work showed that both direct pressurelesssintering of mixtures of carbide and Ni,Al mixtures and reaction sintering where NiAl + Ni mixtures are substituted for Ni,Al powders are viable fabrication processes [2,4,7]. Other results indicate a meltinfiltration-liquid phase sintering process is also a potential fabrication method [8]. In each process, the equipment requirements and processing conditions are identical to those presently used in the powder metallurgy industry so they can be manufactured costeffectively. Because the properties of the aluminide-bonded ceramics are attractive for diesel engine applications, development of these materials was started. Future development of these materials is involved in tailoring of the compositions and processing to optimize the properties in composite systems with Tic. The issue of most interest was the fabrication of parts using cost-effective processing. Consequently, a study was done to correlate the powder processing and sintering with the effect on properties of Ni,Al-Tic composites. This present study used different currently available commercial prealloyed Ni,AI powders with various alloying additions. For comparison, composites were also made with some other nickel-based alloys which are typically used to make Nibased cermets. As far as processing is concerned, the only difference is that the densities of these NiCr(Fe) binders are slightly higher than for the Ni,AI (8.1 versus 7.5 g/cm3). Most of the was done at binder contents of 40 vol. % because this composition more closely matches the thermal expansion of steel. However, in this study, binder contents ranging from 30-50 vol. % were also examined .
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EXPERIMENTAL PROCEDURE The powder characteristics used in fabrication of the composites are shown in Table 1. Several Ni,Al alloys were included and are referred to by the IC- designation given in alloy development studies. Note the large size of the Ni,Al and NiCr powders produced by inert gas atomization compared to the T i c powder. These size differences have an effect on the densification behavior as will be discussed later. The samples were fabricated by milling fine T i c powder, with the metal powders at appropriate levels to produce composites with 30 to 50 vol. % binder phase. The milling was done in isopropanol for 19 h using WC-Co milling media and 1 wt. % polyethylene glycol (Carbowax 8000, Union Carbide, New York) added as a binder. Media wear during milling contributed -0.5 wt. % to each of the compositions. The mixtures were dried and screened to 100 mesh. Specimens were uniaxially pressed in either 25 or 55 mm diameter steel dies at -100 MPa (15 ksi). Sintering was done in a graphite element furnace at temperatures from 1400°C to 1500°C. The heating schedule consisted of a ramp of 10"Clmin from room temperature to 1200"C, a 0.5 h hold at 1200°C to allow for any degassing, and another ramp at 10"Clmin to the
final sintering temperature all under vacuum. The temperature was maintained at the sintering temperature for 0.5 h under vacuum followed by an argon gas pressurization to 1 MPa (150 psi) in 10 minutes and a hold under pressure for 10 minutes. The total time at the sintering temperature was 50 minutes. For all of the test samples, densities were determined by the Archimedes' method. For mechanical property testing, selected samples of high density were machined into bend bar specimens with nominal dimensions of 3 mm x 4 mm x 50 mm. Flexural strength testing was done in four point bending with innerlouter spans of 20 m d 4 0 mm. Fracture toughness was determined by both an indentation and indentationlfracture method [ 12,131. Hardness testing was done with a Vickers diamond indenter at a load of 50 kg. Scanning electron microscopy (SEM) was done on polished sections using back-scattered electron (BSE) imaging and with energy dispersive x-ray analysis (EDAX). The corrosion resistance was determined by measuring the weight loss during immersion in a 1N acid solution at 25°C after 100 h.
Table 1. Physical characteristics of powders. Powder TYPe
SuppliedGrade
Ave. Particle Diameter
Tic NiCr NiCrFe Ni,Al
Kennametal (Latrobe, PA)/Grade 2000 Alfa Aesar (Ward Hill, MA)/ Ni:Cr, 80:20 wt.% Alfa Aesar (Ward Hill, MA)/ Ni:Cr:Fe, 72:14-17:6-10 wt.% Homogeneous Metals (Clayville, NY) IC-50 (Ni- 11.3 A1-0.6 Zr-0.02 B) IC-218 (Ni-8.7 A1-0.2 Zr-8.1 Cr-0.02 B) IC-264 (Ni-8.4 Al-1.7 Zr-7.8 Cr-0.02 B) IC-396M (Ni-8.0 A1-0.9 Zr-7.7 Cr-3.0 Mo-0.01 B)
(PI
RESULTS AND DISCUSSION DensiJication - The densification results are summarized in Fig. 1. As indicated, high densities were obtained for most compositions at temperatures of 21400°C. These types of composites are densified by liquid phase sintering (LPS) and the relatively high liquid contents allows particle rearrangement to proceed rather easily. The effect of different Ni,Al starting particle sizes on densification can be seen by comparing the 40 vol. % IC50 composites in Figs.la and lb. The Ni,Al powders had particle sizes of <44 pm (Fig. la) and <75 pm (Fig. 1b). The larger sizes inhibited densification at 1400"C, but at 1450°C the effect was negligible. Prior work has shown that as the Ni3A1 melts, the liquid is drawn into the surrounding T i c powder by capillary action. This results in the formation of a large void that is difficult to eliminate. With the larger metal particles, this effect is exaggerated. Both the IC-50 and IC-218 based composites exhibited high densities with the V-LPHIP firing schedule. On the other hand, the composites with IC-264 and IC-396M 612
1.3 44-90 <44
<44,<75 <44 <44 <65
alloys, exhibited decreased densification. Both these alloys contained higher Zr contents than the other alloys and, in the case of IC-396M, Mo was also present. These amounts of Zr should have only a minor effect on the solubility of T i c and diffusion in the liquid phase and thus should not seriously affect LPS. However, the Zr would act as a powerful oxygen getter in the system due to the high free energy of formation and stability of ZrOz. This oxygen gettering could significantly influence the wetting between the T i c and the liquid phase and thus affect the solution-reprecipitation kinetics. Small amounts of surface oxygen are well known to adversely affect the wetting between carbides and molten metals and in a similar fashion, minor additions of Mo to Ni have been shown to reduce the contact angle with TIC from 17" to 0". Previous results have suggested inhibited sintering with Zr additions to Ni,Al alloys [9,10]. In the case for the composite utilizing IC-396M as the binder phase, the larger size of this powder (<65 pm) may have also been a factor in the decreased densification behavior. The non-aluminide alloys (NiCr and NiCrFe) produced composites with high densities even at moderate temperatures.
100
95
90
85
80 1400
1450
1500
Sintering Temperature C)
1400°C 1450°C 1500°C Sitering Temperature (“C)
(b) Fig. 1. Summary of densification results: (a) composites at different Ni,Al contents using IC-50 (<44 pm); (b) composites containing 40 vol. % of various binders.
Microstructure - Representative microstructures of composites containing 40 vol. % binder are shown in Fig. 2 . The T i c particles exhibit a ‘core-and-rim’ morphology caused by reaction of the liquid metal with the T i c and the precipitation of secondary T i c during grain growth. Such ‘core-and-rim’ structures are typically observed in Tic-Ni cermets. Energy dispersive x-ray analysis of the core-rim structure revealed several observations that were similar to previous results on TiCNi-Mo cermets. The central core region consists of the original TIC with no detectable other components. In the rim region, other carbide formers are observed, such as Cr, Mo, and W. The W is from the wear of the milling media during processing, whereas the Cr and Mo come from the constituents in the various binders. Preferential accumulation of the carbide formers in the rim structure on the T i c has been observed in other studies [14]. As shown in Fig. 2b, rounding of the grains was observed with the sample containing the NiCrFe binder. However, such rounding behavior was also associated with all of the Ni,Al alloys with Cr or Mo additions A typical fracture surface of the composites is shown in Fig. 3. Necking and deformation of the Ni,Al around the T i c grains can be seen on the fracture surface. In addition, debonding along the TiC-Ni,Al interface was predominant.
(b) Fig. 2. BSE images of microstructures of (a) Tic-40 % Ni,AI (IC-50) or (b) Tic-40 % NiCrFe. Both sintered at 1500°C. T i c particles exhibit a ‘core-and-rim’ morphology.
Fig. 3. Fracture surface of Tic-40 % Ni,Al (IC-50) composite sintered at 1450°C.
Mechanical Properties - The mechanical properties were determined for composites that achieved a density 295% T. D. A summary of the room temperature flexural strength results is shown in Fig. 4. As indicated, the strength generally increases with increasing binder content. The IC-218 alloy sample exhibited slightly higher strength at comparable binder contents and densities probably due to the slightly finer grain size of these composites with some contribution from solid solution hardening of the matrix. The majority of the strength limiting flaws for all the composites in the present study was the occurrence of metal “pools” as shown in Fig. 5. These metal accumulations are most likely the result of the large particle size of the starting
613
raw materials. The high temperature strength of the composites containing IC-50 Ni3Al alloy are shown in Fig. 6. As indicated, the strengths are retained or actually increase up to 500°C. This behavior is a direct effect of the increase in yield strength of the Ni3Al binder. Like the strength, the fracture toughness generally increased with increasing Ni,Al volume content as shown in Fig. 6. Such behavior would be anticipated based on the ductility of the binder phase. In all cases, the fracture toughness was K,, > 12 MPadm and for the composite with 50 vol. % Ni3Al and fabricated with prealloyed Ni,AI powders was as high as 27 MPadm. These values are exceptional and were one of the reasons for choosing these types of composites for development for diesel engine applications. The high toughness values are a result of the plastic deformation and crack bridging effects of the Ni,Al binder. These effects are illustrated in Fig. 8 (and also shown in Fig. 3 previously). Fig. 8 shows the crack behavior at the tip of the an indent introduced during hardness and toughness testing. Crack bridging is readily evident in Fig. 8 and the crack length is surprisingly small considering the indent load was 50 kg. The indent hardness was measured on several specimens sintered at different temperatures and is summarized in Fig. 9. As shown, the hardness decreased with increasing Ni,Al content. The IC-218 alloy sample exhibited slightly higher hardness probably due to some contribution from solid solution hardening of the matrix. The corrosion behavior of selected composites in the study is summarized in Fig. 10. The NiCr-based binders are currently used in corrosion resistant cermets and, in general, their behavior was better than the Ni,Al-based composites in the nitric and sulfuric acid solutions. However, in the hydrochloric acid solution, Ni,Al-based composites showed slightly better resistance Thermal expansion behavior of different composites is shown in Fig 11. As expected, increasing the Ni,AI volume content increases the thermal expansion coefficient. Comparison with the NiCr sample, reveals that the Ni3Al-basedcomposite has a lower expansion coefficient for the same binder content. The expansion coefficients for the composites in the study are similar to those of steel which is -1 1.5X10-6.
1200
,
,
, ,
,
, ,
,
,
Fig. 5. Fracture surface of Tic-40 vol. % Ni,A1 composite sintered at 1450°C.
1300
,
,
,
,
,
...................................
........
900
................
800
......
700 0
200
400
600
8W
Temperature (“C)
Fig. 6. Summary of flexural strength results on TiCNi,Al composites sintered at 1450°C.
30 25
IC-5W40
IC-218140 NiCrl40
NiCrFd40
IC-5W30
IC-50150
Binder TypeNd. Content (%)
Fig. 7. Fracture toughness results on TIC-Ni,A 1 composites sintered at 1450°C and 1500°C.
.
., ..................................
1000 I
m
n
3
800
srn
E
5 -
600
400
2 200
0 IC-50140
IC-21W40
NCrFd40
IC50/30
Binder TypeNol. Conlent (%)
Fig. 4. Summary of flexural strength results on TiCNi,Al composites sintered at 1450°C.
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Fig. 8. SEM image of microstructure of Tic-50 vol. % Ni,Al sintered at 1500°C. Crack bridging by the binder is evident at indent tip (top of photograph).
MPa up to 1150 MPa for volume contents from 30 to 50%. These strengths were retained or actually increased up to temperatures of at least 500°C. The fracture toughness values were all >12 MPadm and as high as 27 MPadm at Ni,AI volume contents from 30 to 50 vol. %, respectively. The thermal expansion coefficient of the Tic-40 vol. % Ni,AI composite is similar to that for steel.
8
P
(I,
6
I
e
B
4
ACKNOWLEDGMENTS 2
0 IC-218
IC-50
NiCr
NCrFe
Binder Type
Fig. 9. Summary of indent hardness results on Tic-Ni,AI composites sintered at 1450°C and 1500°C.
100
10
1
........ I
I
,
,
............. ...............
REFERENCES
. . . . . .
. . .
..............
............. ~
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.-
b
I
0.1
__
n ni IC-50
IC-218
NiCr
NlCrFe
Binder Type
Fig. 10. Weight loss for Tic-40 vol. % binder composites immersed in 1N acid solutions.
30 vol. % IC-50
40 vol. Yo IC-50
Research sponsored by both the Propulsion System Materials Program, DOE Office of Transportation Technologies under contract DE-AC05-000R22725 with UT-Battelle. The research also used the ORNL SHaRE User Facility supported by the Division of Materials Sciences, U.S. Department of Energy.
40 vol. % NiCr
Birder Canlent and Type
Fig. 11. Thermal expansion coefficients for selected TIC composites sintered at 1450°C.
1. T. N. Tiegs, et al, Mater. Sci. Eng., Vol. A209, No. 12,243-47 (1996). 2. T. N. Tiegs, et al, pp. 211-218 in Internat. Conf. on Powd. Metall., Met. Powd. Indus. Fed., Princeton, NJ (1995). 3. K. P. Plucknett, et al, Ceram. Eng. Sci. Proc., 17[3]314-321 (1996). 4. T. N. Tiegs, et al, pp. 339-357 in Internat. Symp. Nickel and Iron Aluminides, ASM International, Metals Park, OH (1997). 5. J. H. Schneibel, et al, pp. 329-338 in Ref.4. 6 . R. Subramanian and J. H. Schneibel, J. Metals, 49[8] 50-54 (1997). 7. T. N. Tiegs, et al, Ceram. Eng. Sci. Proc.,19[3] 447455 (1999). 8. K. P. Plucknett, et al, J. Mater. Res., 12 [lo] 25152517 (1997). 9. T. N. Tiegs, et al, Ceram. Eng. Sci. Proc.,Am. Ceram. SOC.,21[3]0721-728 (2000). 10. T. N. Tiegs, et al, Adv. Powd. Met. Partic. Mater.1999, Metal Powder Industries Fed., Princeton, NJ (1999). 11. D. E. Wittmer, et al, pp. 237-248 in Adv. Powd. Met. Partic. Mater.- 1999, Metal Powder Industries Fed., Princeton, NJ (1999). 12. G. Anstis, et al, J. Am. Ceram. SOC.,64 [91 533538( 1981). 13. P. Chantikul, et al, J. Am. Ceram. SOC.,64 [9] 539543 (1981).
CONCLUSIONS Tic-Ni,AI composites are under development for application in diesel engines because of desirable mechanical and physical properties, in addition to good
corrosion and wear resistance. Liquid phase sintering of the composites with 30 to 50 vol. % Ni,AI results in high densities being achieved for sintering temperatures 21400°C. The final microstructures consist of T i c grains and a surrounding matrix of a (Ni,Ti),AI. The Tic grains are made up of a core region of pure TIC and a rim region of a Ti(Cr,Mo,W)C solid solution precipitated from solution during liquid phase sintering. Room temperature flexural strengths ranged from >800
615
This Page Intentionally Left Blank
CERAMIC JOINTS BETWEEN S i c BODIES: MICROSTRUCTURE, COMPOSITION, AND JOINING STRENGTH J. Woltersdorf, E. Pippel, P.Colombo*
Max-Planck-Institut flir Mikrostrukturpbysik, Halle, Germany, *Universita di Bologna, Italy
Introduction Sic-based ceramic composites are promising structural materials because of their excellent heat resistance and mechanical properties. Moreover, silicon carbide fiber (SiCf) reinforced composites are of specific interest for future thermonuclear fusion reactor applications as they exhibit a low radioactivity by neutron transmutations, a high stability after the reactor shut-down, good high- temperature properties, a low plasma contamination and a low specific weight I . In order to fabricate large pieces or complicated structures, suitable joining techniques have to be developed. While the joining of monolithic S i c has been accomplished by using different techniques as, e.g., active metal brazing, solid state joining and reaction bonding 2, the joining of SiC/SiCf composites has not been extensively reported in the literature 3,4. The used joining techniques often require conditions which are not compatible with the upper temperature of about 1200”C, which currently limits the operation of SiC/SiCf composites owing to the degradation of the S i c fibers used so far’ and with respect to the requirements of fusion technology. Brazing with metallic fillers is not applicable because of the heavy nuclear transmutability of the fillers. Therefore, we developed an alternative joining method using preceramic polymers like polycarbosilane or polysiloxane 6,7, which yield a S i c or SiOC ceramic, respectively, when heated in an inert atmosphere 8,9. The present paper reports the results of microstructural and microchemical investigations of the interlayers in such polymer derived joints between SiC/SiCf bodies as revealed by transmission (TEM) and analytical (AEM) electron microscopy in addition to some data which characterize their mechanical properties including those of joined monolithic The study was performed in order to clarify the interlayer mechanisms determining the joining process.
Experimental Monoliths of sintered S i c were provided by Associazione EURATOM-ENEA, Italy, and the twodimensionally reinforced SiC/SiCf composite material was from Societ6 Europ6enne de Propulsion, France. The surface structure of the latter material was rather complex (total roughness Ra -25 pm, average peak-tovalley height Rz = 130 pm). To improve the surface quality all specimens were mechanically polished on one side, using a 5 pm S i c granulation, and ultrasonically cleaned in acetone prior to the joining.
Another feature of the SiC/SiCf composite material is that it possesses a thick (> 100 pm) over-coat of CVD S i c which was not completely removed by the mechanical surface preparation. Thus the fibers were not exposed and did not participate in the joint formation. The preceramic polymers used for the joining experiments were a methyl-hydroxyl-siloxane (SR350, General Electric Silicone Products Div., Waterford, NY) and polycarbosilane (PCS, Dow Coming X9-6348) which were homogeneously applied to the surface of the specimens to be joined. Subsequently, some samples were loaded with an axial pressure of 9.8.10” to 4.10-2 MPa, and heated at 200°C for 2 hours in order to achieve the complete crosslinking of the silicone resin or to cure the polycarbosilane. Then, the samples were heated in an argon (99.99 YO)flux for 1 hour at temperatures between 900 and 1200°C. The heating and cooling rates were very slow (1Wmin) as to minimize possible residual stresses due to the thermal expansion mismatch. In some experiments, commercial titanium (Cerac Inc. T- 1241, average grain diameter 15 pm) or p-Sic (H. Starck B10, average grain diameter 0.3 - 0.5 pn) powders were added to the polymers, in a weight ratio of 0.5 or 0.7 to 1 to act as fillers to lower the polymer shrinkage during the pyrolysis. Microstructural and nanochemical investigations were performed on joints between sintered SiC/SiCf composites, processed with SR350 preceramic polymer and applying optimum fabrication conditions (cf. next section). The examinations of the interfaces between composites and joining material were carried out by using a Philips CM 20 FEG (TENSTEM) electron microscope operated at 200 kV and equipped with a Voyager EDX system, which enables an atomic lattice plane imaging (HEM-mode) and a chemical analysis (element mapping) with a spatial resolution of a few nanometers. The TEM cross section specimens were prepared by standard techniques with final Ar-ion milling to electron transparency. These specimens allow an observation of the composites, the joining layers, possible reaction layers and interfaces side by side in one and the same specimen. For details concerning the special problems of TEM investigations of interfaces in ceramics and Si-C-0 composite materials including the elucidation of structure/property relations, see, e. g., refs. 12-19.
Results All the images reported are to be considered typical and representative of the interfacial region of each
617
sample. To get an overview of the interfaces to be examined and of their special features as mentioned above, Fig. 1 shows a light microscopic image of joint SiC/SiCf composites. Clearly the fibre structure (above and below), the 80 nm thick over-coat of CVD Sic (light gray) and the almost flat interface joing layer can be seen.
Fig.l: Cross-section of a joint SiC/SiCy (light microscopic image).
The amorphous joining layer of SiC/SiCf composites investigated is free of pores and has a thickness of 5-6 pn. As shown in th TEM image of Fig. 2 the interface between layer and SIC matrix of the composite is very flat again, with a 25 to 40 nm thin zone of a destroyed Sic structure next to the interface.
Fig. 2: Flat interface between S i c (upper IeB) and the joining layer with a 25-40 nm zone of destroyed S i c structure next to the inte$ace
The latter is demonstrated in the high resolution lattice plane image of Fig. 3: while in the interior of Sic (left) the { 1 1I } atomic planes of Sic regularly propagate over large areas, a lot of crystal defects (stacking faults and/or microtwins, characterized by an interruption and a slight change in the direction of the atomic planes) have formed near the interface to the joining layer
618
(right). Contrary to this, the Sic is free of defects at a distance of 25 nm away from the interface. While the joining layer is amorphous in its interior, the high resolution image reveals a precipitation of small amounts of carbon near the interface (cf., the 0.34 nm graphite atomic planes in the almost amorphousjoining layer of Fig. 3, right).
'Fig. 3: HREM image of the interface between Sic (lej) and joining layer with crystal defects in the Sic, notice the small graphitic particles in the amorphous joint (right) In the present stage of investigation it is difficult to infer that the SiC/SiCfcomposite surface has any effect on the development of a weak graphitic boundary layer 13, as it is well known that silicon oxycarbide glasses contain some turbostratic carbon '. Together with the CM 20 electron microscope the Voyager I1 EDXS-equipment allows the recording of X-ray spectra and element line scans from selected specimen areas with a resolution of some nanometers depending on the specimen thickness. For a proper evaluation of the following X-ray measurements it should be considered, however, that fluorescence phenomena and beam spreading effects in thicker specimen regions or at contamination hillocks will deteriorate the real spatial resolution of the chemical analysis although the electron probe has an actual diameter of only 1.5 nm. A further uncertainty in recording the elemental profiles may be due to a minor overlapping of the materials at the interface. Regarding this, in all the specimens investigated the chemical composition of the joining interlayer proved to be nearly the same with only silicon, oxygen and a small amount of carbon, i.e., a silicon oxycarbide glass has formed. Fig. 4 demonstrates a linescan across the interface of a SiC/SiCf joint presenting a situation similar to that of Fig. 2: a small electron probe was stepped along a definite line of 40 discrete points, with each recording a complete EDX-spectrum. The corresponding X-ray emission of carbon, oxygen and silicon from Sic (left) to the joining layer (right) is shown, with an overall line length of 226 nm (in the linescan the position of the interface, as visible by TEM, is marked by a perpendicular line). While some part of the carbon signal results from contaminationsin the microscope, the silicon X-ray intensity roughly
represents the specimen thickness decreasing from Sic toward the interlayer owing to preparation. The oxygen linescan, however, clearly reveals almost no diffusion from the joining layer into Sic (expect within a range
E l -Si
Fig. 4: X-ray line scans of Si, C, 0 across the interface in SiC/SiC’, length: 226 nm, obviously no oxygen has difised into the matrix
of 5 to 10 nm). The small 0-signal in the Sic region (left) probably results from surface oxidation of S i c during the preparation procedure. The quality of the joints was determined by k t u r e shear stress tests following the ASTM D905-89 test procedure, using a crosshead speed of 1 mdmin. Even if a pure shear stress field does not necessarily occur in this testing technique, it is, however, a suitable means of comparative evaluation. Fig. 5 represents the fracture stresses obtained for the joining of a-Sic monoliths. 50,
€
g
-I SR3G
-
--c sR3s PCS
sR35ocn
pcS+SiC
PCS I
900
-
S i c monolith. PCS was not effective as a joining material. At every joining temperature tested, the specimens exhibited a mechanical strength almost not measurable, independently of the thicknesses of the joint layers (ranging from about 20 pn up to 200 pm after curing). Because of the mentioned considerable shrinkage and density change to which preceramic polymers are subject due to the polymer-to-ceramic conversion 68, we expected that the introduction of fillers would have helped to reduce this volumetric variation, which could have adverse effects on the joint and its strength. In fact, for pure SR350 resin, the joint thickness is reduced by a factor of five to seven during the heat treatment at 1200°C. However, using inert (Sic) or active (Ti) fillers degraded the mechanical strength of the joints. A possible explanation is that with the introduction of these powders the joint thickness (measured after crosslinking) increased from about 20 pn to about 32 pn (for SR350 + Sic powder) and to about 57 pn (for SR350 + Ti powder). For PCSjoined specimens, the limited increase in strength with the introduction of powders into the joining material could be due to a more homogeneous layer present there. The following table lists the fracture stresses (in MPa) obtained from the SiC/SiCfjoints: Joining conditions
1000°C 1200°C
SR350,no pressure, unpolished SR350, no pressure, polished SR350, pressure=200g/cm2 SR350,pressure=400g/cm2 SR350+Ti, polished SR35O+Ti, unpolished
1.4 9.9 13.5
0.7
0.2 6.3 18.0 21.6 3.5 3.1
The increase of the joint strength with temperature is again obvious. Moreover, applying pressure during the formation of the joint is certainly advantageous. As for monolithic a-Sic, PCS does not afford the formation of proper joints. Mechanically polishing the SiC/SiCf composite was proved to be necessary, as this probably helped to create flat areas where joining could take place ’. The introduction of powders improved the joint strength only for unpolished specimens, because of the more effective filling of the surface roughness by the joining material.
,
loo0 1100 luw) Joining Temperature (“C)
Fig. 5: Fracture shear stresss for a-SiC monoliths joined at diyerent temperatures
It is obvious that using solely pure silicone resin (SR 350) as joining material results in a sufficient mechanical strength of the structures obtained. The shear strength increased with the joining temperature, and,for the samples heat treated at 12OO0C,most of the failures did not occur in the joint region but in the a-
Discussion The microstructural and nanochemical characterizations show that the pyrolytic decomposition of silicone resin produces a rather homogeneous silicon oxycarbide glass layer, the thickness of which seems to depend on the type of substrate joined. The investigations of the interfaces between the SiOC glass and the various S i c bodies indicate that the joining material is absolutely non-reactive, at least at the processing temperatures used, as there are neither reaction layers visible nor a chemical attack of Sic near the joints. Furthemore, the mechanical polishing of the substrates produced very flat interfaces. All these
619
particulars allow one to conclude that here the joining process is similar to adhesive bonding lo**. The joining mechanisms do not seem to involve any mechanical interlocking. They rather imply a chemical bonding with the direct formation of bonds between the Sic bodies and the inorganic adhesive material. The most typical microstructural feature of the interfaces is a 30 to 100 nm extended zone of enhanced diffraction contrast in Sic near the interface resulting from a destruction of the Sic lattice by mechanical strain and/or crystal defects. According to results of micro-indentation tests at the joint interfaces we can exclude the presence of residual stresses in the materials so that this zone can be attributed to the mechanical grinding/polishing of the Sic parts prior to the joining. The lack of residual stresses is certainly related to the low coefficient of thermal expansion (CTE) mismatch between Sic and SiOC materials lo. It has to be pointed out that all studied interfaces appear smooth on the scale observed, but a smootheqing of the interfaces during the bonding process seems quite improbable, considering the comparatively low temperature used for joining. Furthermore, as clearly shown by the microchemical analyses, no evident compositional gradient was observed in the glassy layer and in the Sic material in the proximity of the interface. Moreover, no growth or dissolution of the Sic material in the glassy interlayer was observed. The difference of the mechanical properties between joints of a-Sic and joints of SiC/SiCf bodies obviously is primarily not due to different microstructural or microchemical interfacial features. It might rather be due to the variation in the thickness of the joining layers: the SiC/SiCfjoint is thicker than aSic joint. Since oxygen does not seem to play an important role in the joining mechanism, the difference in the bonding ability between silicone resin and polycarbosilane preceramic polymers might be explained by their different decomposition behaviour upon pyrolysis. First of all, the weight loss during pyrolysis for pure PCS is considerably higher than for SR350 (PCS ceramic yield -55%, 6; SR350 ceramic yield -75%, '), possibly resulting in the formation of voids in the joint.
Conclusions Joints between Sic bodies were successfully produced using a silicone resin. In the case of joining monolithic a-Sic parts a much higher strength was attained than for joints of SiC/SiCf composites. Microchemical investigations showed that the joining material consisted of a silicon oxycarbide glass, and that no oxygen diffusion occurred between the glass and the joined Sic bodies. The flat interface structure and the lack of any reaction layer suggest that the joining mechanism involves the direct formation of chemical bonds between the Sic bodies and the joining material. The interfaces between joint and Sic material were similar for a - S i c monoliths and SiC/SiCf
620
composites. The lower strength of the SiC/SiCf composite joints is probably due to the uneven surface morphology of the composites. Polycarbosilane was not effective as a joining material, possibly because of its different decomposition behaviour upon pyrolysis. The introduction of Sic or Ti powders did not improve the mechanical strengthof the joints.
References 1. A. Donato, R. Andreani Fusion Technology, 29 (1996) 58-72. 2. T.Iseki In Silicon Carbide Ceramics - 1, S . Somiya and Y.Inomata eds., Elsevier Applied Science, London, 1991, pp. 239-263. 3. B.H.Rabin J.Am.Ceram.Soc.,75 (1992) 131-35. 4. P. Lemoine, M. Salvo, M. Ferraris, M. Montorsi and H. Scholz J.Am.Ceram.Soc.,78 (1995) 1691-94. 5. G. Simon and A.R. Bunsell J.Mater.Sci., 19 (1984) 3658-70 6. A. Donato,P. Colombo, M.O. Abdirashid, in: Ceram. Transactions,57,43 1436,1995 7. P. Colombo, M.O. Abdirashid, G. Scarinci and A. Donato, in Fourth Euro-Ceramics, Coatings and Joinings, Vol. 9, B.S.Tranchina and A. Bellosi eds., Gruppo Editoriale Faenza Editrice S.p.A., Faenza, Italy, 1995, pp. 75-82. 8. E. Bouillon, F. Langlais, R. Piller, R. Naslain, F. Cruege, P.V. Huong, J.C. Sarthou, A.Delpuech, C. LaRon, P. Lagatde, M. Monthiouxand and A. Oberlin J.Mater.Sci., 26 (1991) 1333-45 9. G.M. Renlund, S. Prochaza and R.H. Doremus J.Mater.Res., 6 (1991) 2716-34 10. E. Pippel, J. Woltersdorf, P. Colombo, A. Donato J. Eur. Ceram. SOC.17,1259-1265 (1997) 11. P. Colombo, V. Sglavo, E. Pippel, J. Woltersdorf J. Mat. Sci. 33 2405-2412 (1998) 12. J. Woltersdorf, E. Pippel Colloq. Phys. C1, Suppl. au no 1, T. 51, 947-956 (1990) (Proc. Int. Congr. on Intergranular and Interphase Boundaries in Materials, Paris 1989) 13. G.Grathwohl, M. Kuntz, E. Pippel, J. Woltersdorf Phys. Stat. Sol. (a) 146,393414 (1994) 14. E. Pippel, J. Woltersdorf, A. Hilhnel, R. Schneider Cer. Transactions (American Ceramic Society) 57, 273-278 (1995) 15. A. Hilhnel, E. Pippel, J. Woltersdorf J. Microscow (London), 177 (3), 264-271 (1995) 16. A. Hihel, E. Pippel, R. Schneider, J. Woltersdorf, D. Suttor, Composites Part A 27A, 685-690 (!996) 17. J. Woltersdorf, E. Pippel, A. H i h e l 2. f- angew. Muth. u. Mech. (Wiley) 78, Suppl.1 S81S84 (1998) 18. E. Pippel, 0. Lichtenberger, J. Woltersdorf, J. Mat. Sci. letters, accepted 2000 19. A. Hiihnel, E. Pippel, J. Woltersdorf Cryst. Res. & Techn.,accepted 2000
FATIGUE BEHAVIOR OF CERAMICS STRESSED NEAR FATIGUE LIMIT UNDER ROTARY BENDING H.N.KO Nakanihon Automotive College Sakahogi-cho, Kamo-gun, Gifu-ken, JAPAN 505-0077
ABSTRACT As a preparatory experiment to examine the fatigue damage of ceramics, the fatigue strength of plain specimens, stressed near fatigue limit, was studied at room temperature under rotary bending. The main materials tested were sintered Si3N4 and gas pressure sintered Si3N4,which were used practically for rotors of turbochargers. These materials had the effect of cyclic loading, although the degree of cyclic effect was different. It was known that the fatigue strength was not decreased by the pre-loading near fatigue limit. On the contrary, the pre-loaded specimens seemed to be stronger than virgin specimens. Similar tendency was confirmed on Y-TZP plain specimens. The effect of pre-loading, such as understressing effect or coaxing effect, may appear in some ceramics which satisfy a certain condition relating to the microstructure and microcracking.
near fatigue limit was discussed.
EXPERIMENTAL PROCEDURE The materials used were sintered Si3N4 (SSN) obtained from Kyocera (Japan) and gas pressure sintered Si3N4 (GSSN) obtained from NTK (Japan). Each material was fabricated with YzO3 and A 1 2 0 3 as additives. The microstructure of each etched surface is shown in Fig.1. The materials had more rodshaped grains with various aspect ratios. The properties of the materials are shown in Table 1. Sintered tetragonal zirconia polycrystals containing 2.5 mol % YzOs (YTZP), obtained from Toray (Japan), was also used in comparison with above two kinds of Si3N4. The specimen had a cylindrical shape and the diameter of the middle part was 8 mm and each end of
INTRODUCTION It is important to know the fatigue behavior of ceramics when using them as structural components, since the fatigue strength is generally lower than the static strength. The fatigue strength, the fatigue limit in particular, must be studied on ceramics to make their reliability higher as structural materials. It also seems necessary to know the fatigue damage of ceramics for their reliability as structural materials [l]. However, the basic data on fatigue behavior, such as the effect of cyclic loading, are still not sufficient, and elementary data of ceramics relevant to fatigue damage are few. More basic data on fatigue behavior seem necessary to clarify the fracture mechanism. As for the aspect, the fatigue behavior of two kinds of Si3N4 plain specimens, whose materials were used practically for rotors of turbochargers, was studied at room temperature under rotary bending [2,3]. The results obtained were compared with those under static fatigue, and the effect of cyclic loading was discussed. The rotary bending test was performed for a fairly long period, and the presence of fatigue limit was suggested. Therefore, the fatigue strength of Si3N4plain specimens, stressed near fatigue limit, was studied to obtain the fundamental knowledge on the fatigue damage of ceramics. Because of the unusual behavior of Si3N4, the similar test was also performed on Y-TZP plain specimens, having transformation-toughening behavior. Based on the results obtained, the effect of pre-loading
Fig.1
Microstructures of materials (a: SSN material, b: GSSN material).
62 1
SSN
Average grain size pm 0.62
Table 1 Properties of two kinds of Si3N4. Average Fracture toughness Vickers hardness aspect ratio IF method M P a 6 kgf/mm2 P=20kgf 2.01 6.3-6.5 1400
GSSN
0.43
1.89
Material
1500
6.3-6.5
the specimen had a larger diameter of 12 mm for chucking; the straight length of the middle part was 15.4 mm. The specimen was ground perpendicularly to its axis to make the finished surface smooth; the maximum roughness of the middle part of the specimen was less than about lpm. The machine used was an Ono's rotary bending fatigue testing machine operating at 3420 cycles per minute [4]; the loading type of the machine was fourpoint bending, and the stress state of the specimen was reversed bending.
t)
450: 400
350
Rotary bending test was carried out mainly within the range lo4 to 10' stress-cycles. Some specimens were tested at cyclic numbers more than lo8 to examine the existence of a knee; 10' stresscycles is equivalent to about 20 days under the present test condition. Static fatigue test was carried out mainly for less than lo6 seconds, using the non-rotating fatigue machine. The test points for each material are plotted on a logarithmic graph shown in Figs.2-4, respectively. The arrowed points indicate those stopped testing. The straight lines in the figures were obtained by the least squares method from the data except the arrowed points. As seen in Fig.2, the S-N curve for SSN material can be approximately represented by a straight line up to about lO'stress+ycles. The knee seems to exist at cyclic numbers more than 10' , and the fatigue limit seems to be about 380 MPa. A similar S N curve, indicating the existence of a knee, was obtained on GSSN material, as shown Fig.3. The fatigue limit seems to be about 400 MPa. The static strength was about 630 MPa for the two materials. It is necessary for proper selection of materials to know the strength degradation due to cyclic loading. In general, the effect of cyclic loading, estimated by the ratio of fatigue limit to static strength, seems to be different because of the strength of material [5]. Although the similar tendency was obtained in this study, the degree of cyclic effect was different in two materials. This was due to the difference of the static fatigue behavior for SSN and GSSN materials, as shown in Figs.2 and 3. The values of n, expressed in o"N=constant and o"t=constant, were 51 and 90 for GSSN material under rotary bending and static fatigue, respectively, and were 49 and 180 for SSN material. The static fatigue limit seemed to be about 490 MPa for GSSN material, and relatively lower than 540 MPa for SSN material. Thus, the ratio of fatigue limit to static
622
A
550 i n
RESULTS AND DISCUSSION Effect of Cyclic Loading
Density g/cm3 3.21 3.22-3.23
10
103 lo4 los
lo6 107 lo8 l o 9 N
Fig.2
S-N and S-t curves for SSN material under rotary bending and (A) static fatigue.
(0)
t (sec)
c--
olor
350
lo3
1
, I
I
I
lo6 N
los
104
,
I
107
,
I
10'
,
109
S-N and S-t curves for GSSN material under (0) rotary bending and (A) static fatigue.
Fig.3
n
am400 -
oo
O
0
300 lo2 lo3 lo4 los lo6 lo7 lo8 109
N Fig.4
Fatigue strength for Y-TZP material under rotary bending and (A) static fatigue.
(0)
fatigue limit was 400/490=0.82 for GSSN material, and was higher than 0.70 for SSN material. Considering the effective volumes of specimens under rotary bending and static fatigue [2], the ratio was 0.86 for GSSN material, where Weibull modulus was taken as 15. Therefore, it can be said that GSSN material has smaller cyclic effect in comparison with SSN material. The quantitative analysis on the degree of cyclic effect must be studied in future, since it is reported [6,7] on the present GSSN material that the plain specimen has no cyclic effect. On the other hand, Y-TZP material had a clear cyclic effect, referring to the limited data shown in Fig.4. The fatigue limit seemed to be about 370 MPa.
-
500
0
0 W
A
A* 0
0
g450 b
400 350
lo2 lo3
Fig.5
Effect of Pre-loading near Fatigue Limit The fatigue strength of specimens, pre-loaded near fatigue limit, was compared with that of virgin specimens. The test points for each material are plotted on a logarithmic graph shown in Figs.5-7, respectively. For SSN material shown in Fig.5, seven specimens, pre-loaded at the stress of 392 MPa within the range 1 . 5 ~ 1 0to~ 4 . 7 ~ 1 0were ~ ~ prepared for testing; at the stress level of 392 MPa, two virgin specimens were fractured at 3 . 7 5 ~ 1 0and ~ 1 . 2 3 ~ 1 0stresscycles ~ which were equivalent to about 8 and 25 days, respectively. Three virgin specimens were compared with seven preloaded specimens at the stress level of 451 MPa where the failure seemed to occur in virgin specimens. These data are shown in the figure by symbols 0 and A, respectively. The fatigue life seems to shift to a longer life due to the pre-loading. Test points, expressed by symbol A,were obtained on the pre-loaded specimens (A) at the stress of 451 MPa within the range 2 . 6 ~ 1 0to~ pre-loaded specimens were tested again 3 . 6 ~ 1 0; three ~ at the stress level of 470 MPa. For GSSN material shown in Fig.6, five specimens, pre-loaded at the stress of 392 MPa within the range 1 . 1 ~ 1 0to~ 1.7x107, were prepared for testing. The testing condition was changed to examine the pre-loading behavior of SSN material. As shown in the figure, each specimen was tested within the range 1 . 0 ~ 1 0 7to 3.6x107, increasing every stress level of 20 MPa until it was fractured; one test point with an arrow, shown at the stress of 529 MPa, indicates the datum stopped testing. The similar test was performed on Y-TZP material, having transformation-toughening behavior [8], to examine the unusual behavior of two kinds of Si3N4. The similar tendency can be seen in the results shown in Fig.7. These results show that the pre-loading does not weaken the fatigue strength of virgin specimens. It is important from a practical standpoint that the loading near fatigue limit gives little influence on the fatigue damage. This indicates that a flaw generation, having some adverse effect on the strength characteristics, does not occur. On the contrary, the pre-loaded specimens seem to be stronger than virgin specimens. This phenomenon is well known in steels as understressing effect or coaxing effect. Contrary to the expectation, any change of specimen surface due to pre-loading near
t-
lo4 los N
lo6 lo7 lo8
lo9
Fatigue lives for (0) virgin specimens and (A, A) pre-loaded specimens under rotary bending (SSN material).
0
--
Am
0
0
0
450
Q-or
elk-
350'
*
lo3 104
I
lo5 lo6
lo8
107
I
'
lo9 i o 1 O
N
Fig.6
Fatigue lives for
(0) virgin specimens and (A,
A, 0 , H , V ) pre-loaded specimens under rotary bending: each symbol for five pre-loaded specimens indicates the data obtained at different stress levels (GSSN material).
350' 103
,
-
I
104
105
lo6
I
- I
107
lo8
lo9
N Fig.7
virgin specimens and (A, Fatigue lives for (0) A, 0 , a)pre-loaded specimens under rotary bending: each symbol for four pre4oaded specimens indicates the data obtained at different stress levels (Y-TZP material).
fatigue limit, such as the stress-induced transformation for Y-TZP specimen, could not be detected apparently by X-ray measurement. Nevertheless, there must be some very small change in materials, stressed near fatigue limit. Although the clear characteristic sign of pre-loaded specimen is not be presented, it is believed
623
that some very small change, such as microcracking, occurs in the surface microstructure. Some strengthening mechanism, such as a blunting of the microcrack tip, may occur during the pre-loading. The accumulation of data on the fatigue damage or crack initiation is needed for k t h e r explanation.
Fracture Features and Fracture Mechanism The fracture surface of GSSN specimen had a flat region and a rough region, as already pointed out on SSN specimen [2]. An optical macrophotograph of GSSN specimen is shown in Fig.8. The flat region of GSSN specimen had no clear fracture origin such as a pore, although the flat region of SSN specimen had fracture origins such as a pore or a porous region and inclusion in many cases. The fracture occurred mainly as an intergranular process for both materials, and the fracture features were not different in the flat region and the area surrounding the region. The roughness of GSSN fracture surface, measured by scanning laser microscope, is shown in Fig.9 for reference; each datum is the average value of arithmetical mean deviation of the profile Ra, measured continuously in the area of 40x30 pm along the fracture propagation.
As shown in Fig.8, the fracture propagated radially from the flat region, which seemed to be restricted to the subcritical crack growth. The flat region was approximately semicircular in shape. The size was examined for virgin and pre-loaded specimens. The relationship between depth and applied stress is shown in Fig. 10; each line was obtained from virgin specimens and the depth was taken as a size of flat region, determined by the optical macrophotograph. The solid line, obtained from whole data for virgin specimens, shows that the size of flat region increases as the applied stress decreases. It should be emphasized that the dashed line, indicating the data under rotary bending, seems to be different from the dotted line, indicating the data under both static fatigue and static test. It is also seen in the figure that the data for pre-loaded specimens are equivalent to those for virgin specimens. Since the fatigue life seems to be controlled by the subcritical crack growth from the initial flaw, the effect of preloading must be related to the initiation of the subcritical crack growth; the fraction of transformed monoclinic phase, measured on the fracture surface of Y-TZP by Raman spectroscopy, seemed to be higher in the fracture initiation part. 750 700 650
600 h
550
2.500 b
450 400 2cn I
_I"
0.1
I
I
0.2
0.3
I
I
0.4 0.5 0.6
b (mm)
Fig.8 Fracture surface of GSSN virgin specimen after rotary bending test (0=45 1 MPa, N=6.59x104 ).
1.5
I
O
0
A
0
Fig.9
624
200
Fig.10 Relationship between depth of flat region and applied stress for GSSN virgin specimen under (0) rotary bending, (A) static fatigue and ( 0) static test ( 0 : pre-loaded specimen under rotary bending).
600 5 (Pm)
400
I
A
800
1000
Roughness of fracture surface plotted against distance from fracture initiation part of GSSN virgin specimen for the inside and outside of flat region ( 0 : 0=451 MPa, N=6.59x104, A : 0=481 MPa, N=4.79x104).
sl b.1
0.2
0.3
0.4
0.5
0.6
b (mm)
Fig.11 Relationship between depth of flat region and stress intensity factor for GSSN virgin rotary bending, (A) specimen under (0) static fatigue and (17) static test ( 0 :preloaded specimen under rotary bending).
The flat region seems to be a characteristic sign, indicating the difference of the fatigue behavior under rotary bending and static fatigue. The stress intensity factor KI at the deepest point in the flat region was obtained, assuming the flat region as a semicircular surface crack, to examine the failure condition. The results for GSSN specimens are shown in Fig.11; each line, obtained fiom virgin specimens, corresponds to that shown in Fig.10. As seen in the figure, the value of KI seems to depend on the size of flat region, and the value for a pre-loaded specimen seems to be equivalent to that for a virgin specimen. It also seems that the value of KI is lower under rotary bending than under static fatigue. The mean values of KI, for virgin specimens, were 10.0 M P a G and 11.1 M P a G under rotary bending and static fatigue, respectively. The difference of Kl was confirmed on SSN virgin specimens; the mean values of KI were 8.9 MPa& and 9.7 M P a G under rotary bending and static fatigue, respectively. It should be emphasized that the failure condition is different under rotary bending and static fatigue, although the related mechanism such as a bridging effect must be studied. Similar tendency could be seen in the results obtained on sintered AlZO3 [9]; this material had no clear flat region, and then the region of the subcritical crack growth was estimated, using the decorated pore.
CONCLUSIONS The fatigue behavior under rotary bending was studied at room temperature on two kinds of Si3N4and Y-TZP plain specimens. The results obtained are summarized as follows: 1. The materials had the effect of cyclic loading, although the degree of cyclic effect was different. The failure condition seemed to be different under rotary bending and static fatigue. 2. The rotary fatigue strength was not decreased by the pre-loading near fatigue limit. It could be said generally that the loading of fatigue limit gave little influence on the fatigue damage. 3. On the contrary, the pre-loading seemed to increase the strength of virgin specimens. The effect of pre-loading, such as understressing effect or coaxing
effect, may appear in some ceramics which satisfy a certain condition relating to the microstructure and microcracking.
ACKNOWLEDGEMENT Sincere thanks should be presented to Dr. 0. Kamigaito of Toyota Central Research and Development Laboratories for valuable advice and suggestions. Thanks should be also presented to Kyocera Kagoshima Factory for their support.
REFERENCES (1)H.N.Ko et al., Effect of Pre-loading on Strength of Sintered Si3N4. Proc. 5th Japan Int. SAMPE Symp., (1997) 569-574. (2) H.N.Ko, Fatigue Behavior of Sintered Si3N4 under Rotary Bending and Static Fatigue. Fracture Mechanics of Ceramics, 9, Plenum Press, New York, (1992) 517-533. (3) H.N.Ko, Cyclic Fatigue Behavior of Gas Pressure Sintered Si3N4. Proc. 6th Japan Int. SAMPE Symp., (1999) 241244. (4) H.N.Ko, Fatigue Strength of Sintered A1203under Rotary Bending. J. Mat. Sci. Lett., 11, (1986) 464466. (5) H.N.Ko, Fatigue Behavior of Sintered A1203 under Rotary Bending and Static Fatigue. Fracture Mechanics of Ceramics, 12, Plenum Press, New York, (1996) 1529. (6) A.Otsuka, H.Sugawara and Y.Ishihara, Static and Cyclic Fatigue of Glass and Silicon Nitride under Tensile and TensionCompression Fatigue Tests. Eng. Fracture Mechanics, 40, (1991) 903-911. (7)T.Niwa et al., Effects of Crack Size on Fatigue Behavior in Silicon Nitride. J. Ceram. SOC.Japan, 99, (1991) 296299. (8) G.Katagiri et al., The Stress-Induced Transformation of Y-TZP Fractured by Uniaxial Tension. MRS Int. Mtg. on Adv. Mats., 5, (1989) 313318. (9) H.N.Ko and A.Ueno, Decorated Pore Detecting Fracture Origin of Sintered Al2O3. J. Mat. Sci. Lett., 15, (1996) 10974099.
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LASER CUTTING AND JOINING OF 2D-REINFORCED CMC Marc Ordung", Florian Berndt, Gunter Ziegler University of Bayreuth, Institute for Materials Research, D-95440 Bayreuth, Germany
ABSTRACT To investigate influence of several cutting methods on flexural strength of ceramic matrix composites (CMC) was the main topic of this work. Strength values vary between about 230 MPa (COz laser) and about 339 MPa (water jet cutting), and can be related to the (thermal) damage of the material caused during cutting. A further topic concerns joining of CMCs with a liquid preceramic polymer after cutting. Shear strength of the joint strongly depends on the porosity of the matrix material before joining and varies between 5 MPa (joining after two infiltration and pyrolysis steps) and 29 MPa (joining after one infiltration step and subsequent crosslinking).
INTRODUCTION Ceramic matrix composites were produced by infiltration of woven C-fibre fabrics with liquid metallorganic precursors (LPI) followed by pyrolysis at 1000 "C in nitrogen. Complex three-dimensional parts can be obtained by cutting and joining semifabricated sheets. The criteria for the cutting methods are high flexibility and precision, minimum tool wear and high cutting speed. In principle the laser technique meets these requirements due to the noncontacting removal of material [l]. In this work, several cutting methods (different types of lasers, diamond saw, water jet) of the CMCs were compared based on the results of four point bending strength and the observation of (thermal) material damage. Furthermore, the composite material was joined by infiltration and pyrolysis using the same organic precursor for joining as for infiltration. Shear strength of the joint was determined by asymmetric four point bending test.
EXPERIMENTAL PROCEDURE Sample Preparation The liquid preceramic polymers (polyzilazanes) used for the infiltration of ceramic fibre preforms were developed at the Institute for Materials Research, University of Bayreuth. They are produced in a pilot plant in batches of several kilograms. The viscosity and the chemical composition of the liquid polymers can be systematically varied by the choice of the raw products and the processing conditions [ 2 ] .
Carbon fibre fabrics (KDL 5002, SGL Carbon) were used as reinforcing components. Preparation of the CMC was performed by liquid polymer infiltration (LPI) of the fibre fabric [3]. After infiltration, the liquid precursor was first converted into an unmeltable crosslinked polysilazane by thermal treatment at 150 "C. In the next step, the polymer was converted into an amorphous SiCN ceramic by pyrolysis at 1000 "C in nitrogen. The ceramic yield was about 75 % and the density increases from 1.1 g/cm3 to 2.2 g/cm3 before and after pyrolysis, respectively, leading to a significant volume shrinkage of the matrix material. The resulting porosity of the CMC was reduced to 5 - 10 % by further infiltration and pyrolysis steps. In total, the number of infiltration and pyrolysis steps was seven. The sample size was 50 mm x 110 mm x 3.5 mm.
Cutting methods Several cutting methods with different cutting speeds
(v) were used to obtain samples for the four point bending tests from the semifabricated sheets. Diamond saw cutting The samples were cut by a sliding copper wire, covered with diamond particles (v = 0.5 d s ) . Laser cutting In comparison with the mechanical cutting, the samples were cut by two different types of lasers. The first type was a carbon dioxide laser (h = 10.6 ym, Pcw = 700 W, v = 3 d s ) and the second one a copper vapour laser (A, = 510.6 nm, h2 = 578.2 nm, PAV = 64 W, v = 0.3 d s ) . Water jet cutting The samples were cut by an abrasive water jet. The abrasive medium was alumina containing sand and the jet pressure was 300 MPa (v = 3.3 d s ) .
Joining The size of the joining samples (CMC) was 3 mm x 4 mm x 40 mm. Two front faces were joined with the same liquid precursor used for infiltration of the CMC. The joining process was performed after the first infiltration step, before pyrolysis or after the first or second pyrolysis step, respectively. Further infiltration and pyrolysis steps were carried out very similar to the preparation process of the CMCs.
627
Mechanical properties
400
Four point bending tests were performed to determine the influence of the cutting method on the strength of the CMCs. The machined area of the samples were under tensile stress. Shear strength of the joint was determined by asymmetric four point bending test 141. A schematic drawing of the experimental setup of the asymmetric four point bending test is given in Fig. 1 .
380
E.
]1 I
4
320
w
1
220 240
IF
I T , I
3404'
~ ~ , . , . , . , * , . , . , . , , , , , 0.29 0.30 0.31 0.32 0,33 0,34 0,35 0,36 0,37 0.38
Displacement at rupture [%I
Fig. 2: Flexural strength and displacement at rupture of the CMCs with the machined area under tensile stress In the case of laser cutting the decrease of flexural strength depends on the thickness of the thermal damage zone of the material due to the high energetic laser beam resulting in chemical decomposition and crystallisation of the CMCs. As shown by x-ray diffraction, the amorphous SiCN matrix was transformed to p-Sic and small amounts of silicon (Fig. 3). The carbon peaks result from the C-fibre reinforcement.
T'-
3F(X-y ) 2bh(x+y)
Fig. 1: Schematic drawing of the experimental setup of the asymmetric four point bending test
RESI LTS
Cu vapour laser C
Mechanical properties Flexural strength and displacement at rupture of the C-fibre reinforced CMCs are shown in Fig. 2. Flexural strength of the C-fibre ceramic matrix composites is significantly influenced by the cutting method. The strength decreases by about 19 % ( 0 s : 230 MPa) after cutting with a cw-carbon dioxide laser compared with mechanical cutting by diamond saw (og:283 MPa). The pulsed copper vapour laser leads to no significant difference of the flexural strength (cB: 292 MPa). On contrary, water jet cutting results in an increase of the mechanical properties of about 20 % (aB:339 MPa) compared with mechanical cutting.
628
I
o
.
I
10
'
I
20
.
I
30
.
I
40
.
I
50
.
l
60
'
I
70
.
I
ao
.
1
90
20"]
Fig. 3: XRD patterns of the thermal damage zone of the C-fibre reinforced CMC after cutting with C02 laser and Cu vapour laser Amorphous S O 2 was detected in the thermal damage zone by energy dispersive x-ray diffraction (EDX). The thickness of the thermal damage zone caused by carbon dioxide and copper vapour laser are shown in Figs. 4 and 5 .
30-
25-
-B 20E . . I--
5 15-
-P: g
!
.
10-
3
. 5-
0.00 0.05 0,lO 0,15 0.20 0.25 0,30 0,35
Fig. 4: Micrograph of the qutting edge and thermal damage zone of the^ CMC after cutting with COz laser (light microscope, Nomarski contrast)
Displacement [mm]
Fig. 6
Shear strength of the joints of CMCs in relation to the stage of processing at joining, determined by asymmetric four point bending test
Two samples were tested for every processing stage at joining. The relatively high shear strength of about 29 MPa for the ceramic joints is comparable with the values measured on mixtures of preceramic oligomers with some powders (Si, Al, B), described by other authors [ 5 ] . However, the high temperature strength of these joints is limited due to the metallic components. On contrary, the inherent precursor-based joint used in this work keep the values constant up to high temperatures.
Fig. 5: Micrograph of the clptting ground of the CMC after cutting with Cu vapour laser (light microscope) The thickness of the thermal damage zone after cutting with a COz laser at 700 W is about 300 pm. On contrary, the copper vapour Idser leads only to a thermal damage of several microns.~Thermal damage of the CMCs decreases the load ca acity of the cross section during the bending test re ulting in a decrease of flexural strength.
t
CONCLUSIONS Although flexural strength of the C-fibre ceramic matrix composites is significantly influenced by the cutting method, all cutting methods described in this work are generally suitable for cutting C-fibre reinforced CMCs. Every method has specific technical and economical advantages and disadvantages. The characteristics of the cutting methods are shown in Table 1.
Joining Shear strength of the jdining experiments of the CMC were tested by asymme ric four point bending test after seven infiltration an pyrolysis steps. Shear strength of the joint strongly depends on the amount of existing matrix before joining1 Shear strength of samples joined after the first infiltrqtion step (11) was about 29 MPa. However, shear strength decreases to about 19 MPa and 5 MPa for samples joined after the first (Pl) and second pyrolysis step (P2), respectively (Fig. 6).
I
+ flexible cutting line
+ no material damage Table 1: Characteristics methods
of
the
different
cutting
629
The liquid polysilazanes developed at the Institute for Materials Research, University of Bayreuth, are well suitable for joining CMCs in an early stage of processing. Shear strength of the joint is sufficient and enables the construction of complex shaped parts from semifabricated sheets.
ACKNOWLEDGEMENTS The authors acknowledge the following organisation and companies for the performance of the laser cutting experiments and the water jet cutting: 0 Handwerkskammer Oberfranken, Bayreuth 0 Fa. ATZ-EVUS, Vilseck Fa. Diinkel und Keller, Bayreuth
REFERENCES W. Konig, F. Klocke, Fertigungsverfahren 3: Abtragen und Generieren, Springer VDI, Berlin, (1997) 191 ff.
J. Lucke, J. Hacker, D. Suttor, G. Ziegler, Synthesis and Characterization of Silazane-based Polymers as Precursors for Ceramic Matrix Composites. Appl. Organometall. Chem., 11, (1997) 181-194 G. Ziegler, I. Richter, D. Suttor, Fibre-reinforced composites with Polymer-derived Matrix: Processing, Matrixformation and Properties. Composites A, 30A, (1999) 41 1-417
0. Unal, I. E. Anderson, S. I. Maghsoodi, ATest Method to Measure Shear Strength of Ceramic Joints at High Temperatures. J. Am. Ceram. SOC., 80, (1997) 1281-1284 0. Unal, I. E. Anderson, M. Nosrati, S. I. Maghsoodi, T. J. Barton, F. C. Laabs, Mechanical Properties and Microstructure of a Novel SiC/SiC Joint, Ceramic Transactions, 77, (1997) 185-194
630
ANALYSIS OF COMPOUNDING AND INJECTION MOULDING PROCESS OF CERAMIC POWDERS Eckart Uhlmann and Philip Elsner* Institute for Machine Tools and Factory Management, TU Berlin, Germany
ABSTRACT Injection moulding of ceramic powders requires additional binder systems to obtain a sufficient green part solidity and the necessary flow properties of the compound for mould filling. The first stage of the process chain, the processing of the compound, is of particular importance, since homogeneity of the injection moulding compound is a prerequisite for a high quality sintered product. The compounding process was examined on a shear roll compactor and a kneading extruder. The processing quality was evaluated by means of existing compound characteristics as well as properties of green part and sintered products. Injection moulding was the second stage investigated. The effects of the injection parameters and their variants on the component properties of the green part and the sintered product were analysed. During further investigations, the injection moulding behaviour and the flow properties were examined during moulding processes of the differently processed compounds.
INTRODUCTION Producing compounds for powder injection moulding includes the supply of the raw materials or of the individual components in a fixed mass ratio to the processing aggregate, further the blending, the homogeneous distribution of the powder in the binder and the formation of a pourable granulate. In the case of powder injection moulding a solid material, in this case the ceramic powder, is not only disseminated in a component which in this case is the binder. Moreover, the agglomerates in the powder also have to be destroyed. This kind of compound production is referred to as dispersive kneading. To break up all powder agglomerates and to coat entirely every single powder particle with binder, the compound has to be homogenised. This is achieved with a sufficient number of material rearrangements and the respective shearing forces. Every remaining agglomerate causes a inhomogeneity in the component which may lead to a diminution of quality or even can be the cause of an early rupture by loading the work piece. The powder volume contents of the ceramic injection moulding masses which approach the critical filler rates places great demands on the processing and kneading aggregates concerning the required power rate and wear resistance. The investigations on the compounding of ceramic injection moulding masses were carried out on two different machine systems, which are described in the following chapter.
The second chapter describes the shaping process of injection moulding and the influence of the varied moulding parameters on the change of the work piece properties on the examples of the green part and the sintered product.
COMPOUNDING OF POWDER INJECTION MOULDING MASSES Investigations in compounding injection moulding masses were conducted on aluminium oxide powder of the specification CT1200SG by the company Alcoa and the binder system TP EK 583 by the company Clariant. Dry ceramic powder is indispensable for the production of a homogenous injection mass. In compounding, the amount of bonded humidity which depends on the respective air humidity can lead to a formation of bubbles in cause of evaporation. Even if these bubbles are destroyed in the melted mass during injection moulding, the humidity that settles on the work piece causes inhomogeneities in the material. This humidity poses a similarly grave problem during the follow up processes, since the water evaporates during thermal debinding before a sufficiently open porosity is reached. The high vapour pressure within the work piece causes a damage or a destruction of the component. For this reason, a desiccation of the powder prior to compounding is of utmost importance. Thermo-gravimetric humidity measurements were carried out during which a mean humidity rate of the powder CT 1200 SG of w = 0.38 % was determined. Therefore, an unproblematic processing of the compound to sintered products requires prior desiccation of the powder to reach a humidity rate of < 0.1 %. To achieve this, materials were dried for 30 minutes in a sigma-kneader at 6 = 120 "C, later the binder was added. The preliminary mixing lead to a dry pourable compound in which the powder was bound on the binder. This way, a dry compound could be produced which has a much lower hydrophilic tendency since the active surface has been significantly reduced. The formation of dust could also be limited which has a positive effect on the subsequent compounding process. The injection moulding masses were processed and homogenised on two different machine systems. A shear roll compactor of the company Bellaform Extrusionstechnik GmbH, Ingelheim, was used, which is a machine that has already been widely applied in powder processing. Moreover, an extruder of DraisWerke GmbH, Mannheim, was applied.
63 1
Kneading extruder: A kneading extruder is a closed processing aggregate which can be separated into three areas. In the indentation zone, the ceramic powder and the binder granulate are transported to the kneading zone by means of a twin taper screw system. Distributive kneading takes place in several kneading chambers that are arranged axially. These chambers are formed with different combinations of the rotor and the stator. The surfaces of the rotors and the stators are especially designed for the respective dispersion step in such a way that the material is continuously rearranged, sheared, compressed and stretched. For preliminary disperging, a small number of grooves is attached on the wheels of the first combination of rotor and stator. With every following dispersion step the number of grooves increases. The material is transported from one kneading step to the next by means of snail segments which are arranged on the main shaft. Cooling channels are inserted in the stators in order to transport the friction heat produced in the compound and hence to prevent material damage. Fig. 1 displays the design of a kneading extruder. stator
compound discharge
screw
rotor I
fig. 1 : Schematic image of a kneading extruder
Shear-roll compactor: A shear-roll compactor is an open, continuously working processing apparatus that serves to knead and homogenise the injection moulding masses. The rolls which are grooved like a thread enable the processing of materials with middle or high viscosity. It is possible to expose the rolls in the loading and in the unloading zone to different temperatures. Due to the grooved rolls, the material which is supplied with a metering hopper is continuously transported in the direction of the granulator. Fig. 2 is a schematic display of the way a shear roll compactor works. The rolls that rotate in opposite directions at different rotational speeds produce a periodically changing pressure in the compound material.
Before entering the gap between the shear-rollers the material is unfolded and then homogenised using extremely high shearing flows. The processing parameters for compounding are listed in fig. 3. Speed ratio nilnz [llmin] 0.5; 1
60 / 70
I
rcl
100; 110;90;90
The compounding tests could not be carried out to the intended extent because the kneading extruder got stuck during almost all of the tests which required a complete'disassembly of the installation for cleaning purposes. Only one compound could be produced with the powder CT1200SG and an elevated binder rate, as this was the only way a sufficient viscosity could be obtained for compounding. Hence a direct comparison of this compound with those that had been processed on the shear roll compactor was not possible. Thermo-gravimetric measurements were carried out to check the powder volume rates of the compounded granulates. A binder loss of 0,Ol mass-% was determined for all compounding processes. The same way, a weight loss of 0.01 % was determined during injection moulding irrespective of the examined parameters. These binder mass losses are caused by high local thermal stress, by friction heat during compounding and by extremely high flow speeds and the resultant friction heat during injection moulding. The internal cavity pressure was measured in the proximity and further away from the gate (fig. 4). From the different pressure levels, conclusions can be drawn concerning the cavity filling process and the quality of flow characteristics of the compound. pressure measuring points
n
far gate
n
I
granulator
fig. 2: Schematic image of a shear roll compactor
injection' moulded workpiece fig 4 Cavity pressure measuring points
632
I
The material which has been pre-kneaded and dried in a Sigma kneading apparatus was transported to the shear roll compactor with a vibrating conveyor channel. This way, a continuous material inflow was guaranteed and a homogeneous material transport along the gap could form. The mass flow rate was on average ms= 280 g/min. The powder material CT1200 SG was processed with the mass ratios of powder and binder mdmb = 86/14 and 85,5114.5 which corresponded to a powder volume content of $ = 62 % and $=61%.
near gate
To
1;2;3
temperatures T1.i; Ti& Tz.1; Tz.2
fig. 3: Processing parameters of a shear roll compactor
I
T,,l
Compounding amount i
In each injection moulding process, the compounding homogeneity was investigated using the injection moulding behaviour, or the required injection work WE was determined using the value of the hydraulics pressure pH needed as follows: WE = A , .;,]pH(t)dt tl
Piston diameter Average injection speed Hydraulic pressure Iniection time
Id=] [bar] Isecl
This correlation investigated by Gissing [Giss83] was established based on investigations on the machining of thermoplastics. Fig. 5 illustrates the injection work required for the compound produced on the shear roll compactor with varying compounding amount. What is clearly distinct is a decrease of the injection power at the compound which has been processed twice or three times at a roll gap of b, = 0.5 mm.
close to the gate and further away, since the investigations had revealed considerable differences. On the green parts, the measurements in the zone further away from the gate revealed a bending stress of only 69 % to 80 % of the values measured in the proximity of the gate. This is due to the loss of pressure within the cavity that leads to a decrease in material compression during form filling. The highest bending stresses were measured for the cases of dual compounding, both on the green part and on the sintered product. 22
2 MPa
5 18 a 16 14 2 12 g
2
8 6
2
4
2
0 compoundingamount i 250
8 MPa 210 5,190 170 150 cm 130 C
9
.8 1
2
3
Compounding amount i
90 70 50
1
2
3
compoundingamount i
holdirQ pressure: pN= 200 bar
fig. 5: Injection moulding work depending on the compounding amount
A significant decrease of the injection power WE occurs in the case of dual compounding in contrast to single compounding. The injection power decreases from WE = 373 J to WE = 347 J and corresponds to a reduction by about 7 %. The third material passage causes a further reduction of the injection power required by another 2 %. Increasing the roll gap from b, = 0.5 mm to b, = 1.0 mm does not bring about a significant improvement of the flow characteristics, even after several material passages. However, compared to the materials that underwent single processing, a slightly higher injection power was required at a roll gap of 0.5 mm which suggests a better homogeneity.
fig. 6: Bending stress of green parts and sintered products as a function of compounding amount
The change of the roll gap on the shear roll compactor did not lead to an improvement of the component properties. An increase of the binder rate resulted in an increase of green part solidity, but this also enhanced the sinter shrinkage. All 4 - point bending tests of green and sintered parts were performed according to DIN 5 1 110.
Fig. 6 displays the effect of the compounding amqunt on the solidity of the green parts and sintered products. Bending strength tests were carried out in the area
633
between the bending stresses within the components at a level of flow-up pressure of pN = 300 bar.
I
grnn part
9 225
c MPe
Parameters
Reference variable
v, [an*/sI
Injection rate ~
r F
~
6: 8: 10
~
~
Dwell level PN [bar]
~
~
200;300;500
100 200 300 400 bar 6W
0,l; 0.5; 5
flowup pressure p,,
Holding pressure time pt [s]
175
E
125
3
100
100 200 300400 t w 600 flowup pressure pr
~~
fig.7:Variation of the process parameters in injection moulding
In all tests, the investigations of the impact of temperatures on the properties of the green part and the sintered product were kept constant, both in the extruder and in the injection mould. The tool was held in a temperature of TwKZ= 53 "C while the temperature of the injection nozzle was TD= 165 O C . The injection tool had no hot runner system. The cavity of the mould had a length of 80 mm, a width of 20 mm and a height of 3.5 mm. The entire injection volume, including the gate, was approximately Vs = 1 3 3 cm3. In the proximity of the gate, the bending stress reaches a maximum of 04B = 17 MPa at v, = 8 cm3/s, respective of the green part solidity (fig. 8). grnn part
20
p
a"m
Gm
p
16 14 12 10
-
8
5
%
n 5
rinterd work piece
250
d MPe
'@175
% 150 125
2
6 5
6
7
8
9 tWs11
injection rate v.
injedlon
T, = 16SC
parameters:
T-
=5 y c
To = 165'C p a w n : T,=53-C injection
v.
= B m/s
V, = 3.0 cm'
t'0.5S
fig.9: lnfluence of the level of flow-up pressure on the bending stress
As shown in fig. 10, the bending stress of the green parts increases with increasing holding pressure time. This has the same effect on the areas close to the gate as on those further away from it. The highest bending stress is obtained at a holding pressure time of N t =5 s at 0448 = 17 MPa in the proximity of the gate and 04B = 1 2 3 MPa further away from it. The bending stress of the sintered products on the other hand reacts differently and decreases with increasing holding pressure time. In this case, a bending stress of 04B = 205 MPa is obtained at N t = 0,l s and only a t = 5 s. In the case of the value of 04B = 176 MPa at N sintered products, prolonging the dwell pressure time lead to a considerable reduction of the bending stress within the work pieces from 32 MPa at N t = 0,l s to 6 MPa at N t = 5 s.
im 75
green patt
6n
5
6
7
8
9 m(s11
InJectlonrate v,
p,, = 200 bar V, = 3.0 cn? t, = 0.5 5
fig. 8: Influence of the injection rate on the bending stress
Even in this case, a distinctly lower bending stress is measured in the area further away from the gate, where a maximum of 04B = 1 2 3 MPa was obtained. Changing the injection rate does not bring about a decrease in green part solidity, contrary to the course of the bending stress in the sintered product. In this case, the difference between the zone close to the gate and the areas further away becomes more subtle as the injection rate increases. Generally an increase in bending stress of the sintered product can be observed with increasing injection rate. The bending stress of the green parts only changes very slightly in the investigated range of the level of flow-up pressure both in the areas close to the gate as well as in the remote areas (fig. 9). The difference of the bending stress within the work pieces remains significant, as the tools are not affected by the chance in level of flow-up pressure. With increasing level of flow-up pressure, the bending stress of the sintered products mainly showed a slight decrease in the more remote zones. What is remarkable is the similarity
634
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IMAGES OF TYPICAL
RUPTURES Fig. 11 displays typical bending images of the green part, both the debindered one and the sintered sample, as well as the sintered product. All samples are taken from the same charge and were manufactured under equal conditions. On the partially sintered component a formation of layers is discernible. Such a layer formation can also be found in the sintered product, it occurred in the case of all injection parameter variations, irrespective of the kind of compounding applied. Defects during debinding can be excluded, since the formation of layers only appeared in the proximity of the gate and
8
decreased with increasing distance to the gate. Therefore the reason for the layer formation could be residual stresses in the green part caused by excessive thermal gradients in the solidification phase.
green part ~
brown part
sintered part
_ I _ _
fig. 1 1: Light microscope images of bending stress surfaces
Fig. 12 displays a typical inclusion of air caused by the injection moulding process. This kind of defect occurred only rarely, no correlation with the investigated parameter variations can be detected. Thi defect can rather be considered as a consequence of a inhomogeneous cavity filling process.
__ ig. 12: SEM image of an inclusion of air and of binder accumulatio in a green part
SUMMARY For the investigations, AI203-compounds of different powder volume rates were produced on two different processing aggregates. Compounding on a kneading extruder posed a number of problems, the required mixture ratios could not be produced due to excessive frictional resistance witch leads to a recurrent standstill of the installation. Using the shear roll compactor it was possible to produce compounds of different powder volume rates at varying processing parameters. Hereby, investigations were carried out into the effect of the varied parameters on the injection moulding behaviour and on the properties of the green part and the sintered product. The injection moulding behaviour was investigated by analysing the pressure at the changeover point of the hydraulic cylinder and the internal pressure of the cavity. These were measured both in the proximity of the gate as well as distant to it. An improvement of the flow characteristics of the compound could be noted with increasing compoun-ding amount. Moreover, the
pressure within the cavity decreased between the two spot marks, hence a more favourable homogenising could be excluded. An increase of the binder rate lead to an improvement of the flow characteristics of the compound and therefore reduced the required injection pressure. The same way, an increase in bending stress of the green parts was observed. However, the shrinkage of green parts and the sintered products increased. Both in the case of the green parts and of the sintered products, an enlargement of the roll gap on the shear roll extruder did not bring about any significant change in required injection work and bending stress. The best results with regard to bending stress of the green part were measured at a powder volume rate of $ = 6 1 % during twice compounding and shear roll distance of b, = 0,s mm. A correlation between the bending stress and the distance to the gate could be established on the samples. The difference is also due to the loss of pressure within the cavity which is responsible for a lower compression of the compound. The kneading extruder is not suitable for investigative purposes, as it is extremely vulnerable with regard to highly viscose masses, also because it requires a lot of cleaning. Providing the required temperatures or cooling poses a problem in the individual zones and, due to a high frictional heat, often leads to thermal damage of the binder system. An additional problem is the high wear on rotors and stators which may have a negative effect on the properties of the sintered product. The shear roll extruder which has an open and simple construction is a more suitable processing aggregate for a more cost-efficient compounding of ceramic injection moulding masses. A variation of selected injection moulding parameters revealed a dependency of the bending stress for the green parts as well as for the sintered products. However, the increase in solidity observed in the case of the green parts did not necessarily occur on the sintered products. After reducing the dwell pressure duration, an increase of solidity was observed on the sintered product, whereas the solidity of the green parts decreased considerably. The difference of bending stresses within the samples of the areas close to the gate and those further away from it also decreased by up to 37 % in the case of the green parts and up to 24 % in the case of the sintered Hence an ideal setting was determined for each varied process parameter which however did not lead to the most favourable results for all required component characteristics. In the course of the investigations a formation of layers within the work pieces was demonstrated by means of rupture images. The layer formation occurred irrespective of the kind of compound and the injection moulding parameters. It was possible to detect a correlation between the distance to the gate and this typical rupture images which only appeared in proximity of the gate. This might be due to residual stresses in the injection moulding part which are caused by excessive thermal gradients in the tool and in the moulded part.
635
Acknowledgements Special thanks to the authors of the Deutsche Forschungsgemeinschaft (DFG) for supporting the research project "Identifikation und Analyse der qualitltsbestimmenden ProzeBschritte und - gr6kn beim Pulverspritzgiekn von Sinterwerkstoffen" (Identification and analysis of process steps and process variables in powder injection moulding of sintered materials).
References Gissing, K., Lampl, A.: Uberwachung des SpritzgieOprozessesdurch Messung ViskositUabhangiger KenngroBen. Plastverarbeiter, 34. Jahrgang, 1983, Nr.5, S. 427-432
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USE OF THE SOL-GEL METHOD IN THE EXTRUSION OF ALUMINA CERAMICS AND ATZ CERAMICS W. Pabst, E. Gregorovi, J. Havrda Department of Glass and Ceramics, Institute of Chemical Technology (ICT) Prague, CZ - 166 28 Prague 6, Czech Republic
ABSTRACT A new forming technique is presented which applies zirconia sols and alumina gels as binders for paste extrusion. Alumina and ATZ (80/20) ceramics are prepared. After firing, the alumina samples exhibit bulk densities > 94 % of TD, apparent porosities of 0.4-0.7 % and total porosities of approx. 4.6 %, while for the ATZ samples the achievable bulk densities of sintered specimens axe only c 93 % of TD. While fued alumina samples attain strength values > 390 MPa, the strength of fued ATZ samples is much lower than could be expected for such materials. The reasons of this discrepancy and possible remedies are discussed.
INTRODUCTION For several unique features (e.g. homogeneity, superfine grain size, chemical versatility) sol-gel methods have gained considerable attention in ceramic technology as preparation methods supplementing the classical forming techniques. While sol-gel methods are successfully used to prepare powders (OD bodies), fibers (lD), and films (2D), the preparation of 3D bodies by a sol-gel method alone has not yet undergone a decisive breakthrough. The main reasons are the large drying shrinkage and the high internal stresses (socalled microstresses "of the second kind"). The former is a consequence of the very low equivalent oxide concentrations in the precursors and makes it impossible to match prescribed dimension and shape requirements. The latter usually causes cracking and often leads to complete destruction of the ceramic bodies, in most cases already during processing. A totally different problem is encountered with respect to the extrusion of oxide ceramic pastes. Extrusion is a versatile forming technique for smalland large-sized translationally symmetric ceramic bodies. In contrast to traditional silicate ceramic technology, where the clay h i i o n in connection with water forms the binder (plasticizer) for the paste to be extruded at room temperature, pure oxide or non-oxide ceramics require special binder formulations. Most of these formulations for the extrusion of advanced ceramics have certain drawbacks (queous methylcellulose binder systems e.g. need to be extruded at elevated temperatures, and dibutylphthdate
systems are dangerous fkom a hygiene viewpoint, paraffine-wax formulations, on the other hand, require elevated extrusion temperatures, careful rheology control [l] and very slow debinding prior to firing [2] etc.), so that the design of new binder formulations is highly desirable. Recently, a new generic variant of extrusion ("ABC paste extrusion", i.e. extrusion of pastes with accommodating binder composition) has been developed at the ICT Prague for ATZ (aluminatoughened zirconia) ceramics [3]. Sols or gels of a chemical composition close to the solid oxide phase which is to be extruded (i.e. the filler powder or powder mix), are used as binders (plasticizers) for the extrusion of pastes at room temperature. The sols and gels used can be considered as precursors that transform into the pure oxide phase (and in this sense accommodate to the main solid phase, e.g. alumina or zirconia) after an appropriate heat treatment. For the preparation of ATZ ceramics by extrusion yttria-doped zirconia sols can been used. A small amount of submicron alumina powder (type AA-03, Sumitomo Chemical / Japan) has been applied in [3] to adjust the binder composition (equivalent oxide ratio) exactly to that of the final ATZ ceramics, i.e. 80 wt.% zirconia and 20 wt.% alumina. This work presents results on pure alumina ceramics, where a boehmite gel has been used as a binder for extrusion, and on ATZ ceramics (80120 composite), where a boehmite gel (instead of submicron alumina [3]) is used (together with the yttria-doped zirconia sol) to adjust the binder composition.
EXPERIMENTAL The basic oxide powders to be extruded were submicron alumina (type AA-04, Sumitomo Chemical Co., Ltd. / Japan) and ATZ (type TZ-3Y20AYTosoh Corporation I Japan), having a particle size (median) of approx. 0.4 and 0.6 p,respectively. The boehmite gel is prepared fkom y-AlO(0H) flocs (Disperal Sol P2, Condea Chemie / Germany) with distilled water and nitric acid. The final concentration of this gel is 14.9 wt.% y-AlO(OH), its density approx. 1.11 g/cm3. The zirconia sol is prepared by dissolving zirconyl nitrate hydrate (ZIO(NO~)~. x H20, Sigma-Aldrich / Germany, where x has been determined by
637
thermogravimetry) in ethanol and adding an appropriate amount of yttria (Lachema-Chemapol / Czech Republic) stock solution (yttria dissolved in nitric acid) to yield 3 mol% of yttria (related to the zirconia) in the final zirconia oxide fraction after firing. Concentration and viscosity of this zirconia sol can be controlled by slow evaporation of the solvent. The equivalent oxide concentration, for which the sol-gel transition occurs, i.e. gelation sets in, lies in the range 18-24 wt.% (4-6 vol.%). For equivalent oxide concentrations of below 20 wt.% the zirconia sols are highly fluid (apparent viscosities of the order-ofmagnitude 100 &as, at shear rates of approx. 100 s-I), for higher concentrations the viscosity increases rather steeply. In order to guarantee satisfactory homogenization of the ATZ powder (and the boehmite gel) in the zirconia sol by dispersive mixing, the viscosity should not be too low, but the evaporation step has to be controlled very carehlly in order to avoid polymerization and transition to a dry xerogel. For an equivalent oxide concentration of 20 wt.% the density of the zirconia sol is approx. 1.29 g/cm3. The binder for the extrusion of ATZ powder is prepared by mixing zirconia sol and boehmite gel in a ratio that results after calcination in a ratio of oxides of 80 wt.% zirconia and 20 wt.% alumina. The pastes are prepared by mixing the binders with their corresponding oxide powders (AA-04 and TZ-3Y20A, respectively). In order to avoid contamination, mixing was performed by hand in a polyethylene vessel with an alumina rod. The last phase of mixing is more or less a process of kneading the highly viscous paste. The as-prepared pastes were then extruded in a small (laboratory-scale) piston-driven batch extruder, cf. [ 1, 21. A stainless steel tube (capillary) with internal thread, an inner diameter of 4 mm, and a length of 80 mm, was used as an orifice. After extrusion, the bodies were slowly dried in air at room temperature and subsequently fired to various temperatures between 1530 and 1600 "C (ramp up 2 OC/min, dwell 120 min, followed by natural cooling in the furnace). The as-fired bodies were subjected to bulk density and porosity measurements (Archimedes method in water), their shrinkage was measured by a digital slide caliper. Flexural strength values were determined by three-point bending tests on the as-fired cylindrical bodies without additional surface polish (specimens with average diameter of approx. 3.2 mm, span 40 mm).
of the zirconia sol being measured before mixing). The true density of the ATZ powder was calculated using for the zirconia component the monoclinic-tetragonal ratio determined for this powder (TZ-3Y20A in the asreceived state) by quantitative X-rayanalysis [3]. Table I. Weight and volume concentrations of oxide powders in the pastes before extrusion Paste I wt.% I vol.% Alumina I 75.4-76.9 I 46.0-48.1 29.5 ATZ 1 63.4 I For both pastes the concentrations listed in this table represent concentrations close to the maximally achievable limits for guaranteed extrudability. Apart from the practical difficulty to ensure sufficient homogenization by mixing (kneading), pastes with higher solids contents are prone to irremovable air inclusions and exhibit dilatant behavior as well as locking phenomena during extrusion. The striking difference in the volume concentrations attainable for extrudable pastes may possibly be attributed to the fact that the (submicron) ATZ powders, in contrast to the alumina powders, are present in the form of (supermicron) agglomerates, which are too "hard" to be dispersed even during high-shear mixing. The presence of 50-p-size granules in these powders is known [4]. Nevertheless, we do not consider this question to be finally settled as long as an electron-optical study of the microstructure of the samples has not yet been performed. Apart fiom the initial and final phases of extrusion, the extrudates of alumina pastes are relatively defect-fiee and their surface is smooth, while the extrudates of ATZ pastes exhibit more macroscopic defects and a higher degree of surface roughness. Drying has been performed simply by leaving the as-extruded samples for several days at ambient conditions (air at room temperature). An extra drying step (in a drier at approx. 110 "C) can be performed, but is not necessary for defect-fiee firing. For the compositions and processing considered here, no cracking occurred during drying or firing. Table I1 lists the measured bulk densities of alumina and ATZ samples in dependence of the firing temperatures used. Table 11. Measured bulk densities [g/cm3] for samples fued at different temperatures
RESULTS AND DISCUSSION Concentrations of oxide powders (AA-04and TZ-3Y2OA) in the respective pastes (before extrusion) are listed in Table I. In order to calculate the volume concentration from the weight concentration in the case of ATZ pastes, the density of the sol-gel binder phase was calculated via its composition (the actual concentration
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157OOC 3.82fo.04 160OOC 3.82M.04
I
4.99N.08
I 5.04kO.07
According to results of quantitative X-ray analysis the relative concentration of monoclinic phase with respect to the total zirconia content is approx. 12.5 wt.%. The theoretical density of the ATZ ceramics can therefore
be estimated to 5.45 g/cm3. It is evident that, while the alumina ceramics exhibit bulk densities larger than 94 % of theoretical density (and bulk densities > 95 % can easily be achieved), for the ATZ bodies the bulk densities are < 93 % of theoretical for all firing temperatures tested. The average diameter of the cylindrical samples after firing is 3.27 mm and 3.12 mm for alumina and for ATZ, respectively. Thus for the samples investigated the linear shrinkage in radial direction is between approx. 18.3 % (alumina) and 22.0 % (ATZ). Tables I11 and IV list measured porosity values. For all firing temperatures tested the porosities are clearly lower for the alumina samples than for the ATZ samples. While open porosities for the alumina samples are practically zero in all cases (taking into account the scatter of approx. 0.4 YO),those of ATZ samples are clearly larger than expected for sintered ceramics (several percent). It seems, that by increasing the firing temperature from 1530 to 1600 OC the open porosity of ATZ samples cannot be reduced further. According to these findings, the optimum firing temperature (where grain growth is as low as possible) can be expected to lie around 1550 OC for both materials, similarly to that determined in [3]. Table 111. Measured apparent porosities [%I for samples fired at different temperatures
1570 OC 1600 OC
0.6 M.4 0.5 M.4
2.8 B . 8 3.4 M.9
The total porosities of ATZ samples after firing are approximately twice as high (8.7-1 1.5 %) compared to those of alumina samples (3.7-5.6 %). Table IV. Measured total porosities [YO] for samples fired at different temperatures
1
I Alumina I
ATZ
1
Table V lists flexural strength values determined in three-point bending for as-fired samples without surface polish. Table V. Measured flexural strength values [MPa] for samples fired at different temperatures (mean values, in brackets peak values)
I
I Alumina I
ATZ
1
It is evident that the strength values show considerable scatter. While the alumina samples exhibit strength values well comparable with those of alumina samples prepared by other forming methods, the strength values of the ATZ samples are at least one order of magnitude lower than those of ATZ ceramics prepared by cold pressing [5] and pressure slip casting [6], which are 858 MPa and 1057 MPa, respectively, not to speak of the peak values of up to 2400 MPa that are reported for hot-isostatically pressed ATZ ceramics or aluminacontaining zirconia [4,7]. The reasons for this extreme difference and the absolutely unsatisfactory strength values of the ATZ samples are a matter of current research. The initially low equivalent oxide concentration of the paste and the consequently low bulk density (high porosity) of the fired bodies is but one of the possible causes (and probably not the most significant one). Internal stresses, which are typical for gel-based systems and are a result of internal constraints exerted by the relatively rigid gel (or xerogel) skeleton hindering free shrinkage during drying, can also reduce strength, since they support the action of external loading. A quantitative study of internal stresses of ATZ ceramics by X-ray diffiction (line broadening) is in progress. Another critical point is the surface state of the samples: After extrusion, the ATZ samples exhibit clearly higher surface roughness (and more defects) than the alumina samples. Since unpolished specimens are used for the strength measurements, surface concavities can act as stress concentrators and crack initiators during the strength tests (three-point bending). A possible way to influence the surface state of the samples after extrusion (and possibly also the consistency of the pastes) is the addition of small amounts of organic additives (e.g. polypeptides or polysaccharides). These ways of modifying the paste composition to improve the sample surface are currently being investigated.
CONCLUSION Monophase alumina ceramics and ATZ composite ceramics have been prepared by a new forming technique (ABC paste extrusion), in which the sol-gel binder phase accommodates to the composition of the main phases (i.e. the oxide powders) after conventional heat treatment (i.e. drying and firing without special regimes). Linear shrinkage in the radial direction of cylindrical samples after firing is approx. 20 YO.Alumina samples can be easily sintered to more than 95 % theoretical density and almost zero open porosity, ATZ samples cannot (bulk density < 93 %, open porosity 3-6 %). Flexural strength values are satisfactory for the alumina samples (> 390 MPa) but totally unsatisfactory for the ATZ samples. Possible reasons are low equivalent oxide concentration in the paste, high internal stresses developed during drying, and high surface roughness after extrusion. Current research aims at a quantification of internal stresses in ATZ ceramics by X-ray diffiction (line broadening),
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an electron-microscopic investigation of the microstructure, and an empirical modification of the paste composition by addition of organic components to improve surface smoothness.
Acknowledgement: This study waspart of the research project CEZ:MSM 223100002 ”Chemistry and Technology of Materials for Technical Applications, Health and Environment Protection” and supported by grant MPO No. FB-CV/64/98.
REFERENCES [ l ] W. Pabst, J. Havrda, E. Gregorovi, Rheology of Ceramic Injection Molding Feedstocks. CeramicsSilikaty, 43 (1999) 1-11. W. Pabst, Influence of Extrusion and Injection Molding on the Microstructure of Ceramics (in Czech). PhD Thesis, ICT Prague 1998. W. Pabst, J. Havrda, E. Gregorovi, B. KrCmovi, Alumina Toughened Zirconia Made by Room Temperature Extrusion of Ceramics Pastes. Ceramics - Silikaty, 44 (2000) (to appear). [4] TOSOH Technical Bulletin No. 2-003: Properties of TOSOH Zirconia Ceramics. TOSOH Corporation / Fine Ceramics Department, Tokyo 1998. [5] L. Esposito, A. Salomoni, I. Stamenkovic, A. Tucci, Processing of Zr02-A12O3 Powders: Consolidation and Characterization of Final Products. Special Meeting on Biomaterials-mini 1992 (eds. 1. Stamenkovic, J. Krawcynski). Publ. Forschungszentnun Jiilich, Jiilich (1994) 37-45. [6] A. Salomoni, A. Tucci, L. Esposito, I. Stamenkovic, Forming and Sintering of Multiphase Bioceramics, J. Mater. Sci. Materials in Medicine, 5 (1994) 651-653. [7] R. W. Cannon, Transformation Toughened Ceramics for Structural Applications. Structural Ceramics (ed. J. B. Wachtman jr.) Academic Press, Boston (1989) 195-228.
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ADVANCED HOT-PRESSED CERAMIC MATRIX COMPOSITES (CMC) IN Sic,, SiCf, Cf / SiSN4 SYSTEMS I.Ju.Kelina, N.I.Ershova, L.A.Pljasunkova, E.I.Jakovenko The State Research Center of Russia Obninsk Research and Production Enterprise “Technologiya”
ABSTRACT Considerable recent attention has been focused on the development of CMC reinforced with Sic, whiskers, discrete and continuous Sicf and Cf fibers. Much attention is given to the microstructural aspects in hot-pressing of CMC with Si3N4matrix and various fibrous fillers. The peculiarities of highly dense matrix formation, filler distribution and orientation in matrix, retention of integrity of crystals and fibers in a composite material, the character of fibedmatrix interactions have been studied. Physical and mechanical properties of CMC in Si3N4-Sic,, Si3N4Sicf and Si3N4-Cfsystems have been compared. At this stage of investigations the effect of reinforcement in Si3N4-Sic, system manifests itself in the increase of high-temperature crack resistance up to 14 MPam” and in the retention of strength at matrix level up to 700-900 MPa at 15OOOC. Maximum values of resistance 8,2 and 9,s MPam” were obtained on combined specimens with short and continuous SiCf and Cffibers, respectively. In both cases the bending strength decreases by 10-20% as compared to monolithic Si3N4. The advantages of hot pressing technique in the manufacture of laminated CMC with discrete and continuous fibers are shown. The materials developed can be used for rotor blades of gas-turbine engines, combustion chamber segments, cutting plates, etc.
observed. Thanks to high durability, these materials are gradually displacing metal materials, including intermetallides and hardened superalloys (Fig. 1) [2]. At the present time the conventional methods of sintering including hot pressing are giving way to new more complex and power-intensive methods of CMC sintering which eliminate fiber damage as they do not require high pressure and temperature. However in practice the traditional methods of sintering remain efficient and easy to be used particularly when whiskers and short fibers are employed and laminated structures of composites are formed. At the same time a combination of suspension impregnation and subsequent hot pressing (SIHP) is one of the most commonly used methods of fiberreinforced CMC sintering. It can be suggested that the presence of two insulating atmospheres (nitrogen and graphite mold) eliminates fiber oxidation.
INTRODUCTION Reinforcement is efficient method of ceramics crack resistance increase. Among structural CMC, fiber-reinforced composites are of the greatest interest as the fibers possess a number of extreme characteristics (due to maximum anisotropy of the structure), and first and foremost high strength and high modulus of elasticity. The examples of ceramics crack resistance increase up to 30 -50 MPa.mln are known
1960
1980
Zoo0
2020
PI.
Owing to the combination of high strengths and crack resistance, CMC are considered to be advanced for engine manufacturing. Nowadays the tendency for hrther development of fiber-reinforced composites is
Fig. 1. Trends of development of advanced materials for engines
64 I
EXPERIMENTAL PROCEDURE
Table 1
Raw Materials Highly dense and high-strength monolithic materials based on composite [Si3N4-Y203, MgO] powders with 0,Ol-0,l pm particles obtained by plasma-chemical synthesis (PCS) were used as a matrix. The materials work in a wide range of temperatures up to 15OO0C, their properties have been investigated in detail [3]. A high activity of the powders on sintering under hot pressing makes possible the introduction of sufficient amount of fillers into them, generally the inhibitors of sintering. Various types of continuous SiCf and Cf fibers and dispersed fillers (short Cf fibers and Sic, whiskers) with different geometrical dimensions, chemical composition and properties were used as reinforcing elements (Fig.2, Table 1 [ 11): Sic, of AM 7 grade (Am.Matrix, US), TWS-200 and TWS-400 (Tocai Carbon.Co., Ltd., Japan) [4]; Sicf fibers of Hi@) and NLM-200 grade (Nippon Carbon, Japan), TM-S 1HOSPX (UBE Industries, Ltd., Japan) and the experimental home fibers VNIIPB; Cf fibers of UKN-P grade (Russia) and T300 and M60J fibers (Torayca, Japan). There is no optimum in every respect and universal reinforcing filler in nature. Each filler has its own advantages and disadvantages and its own fields of application.
Properties of fibrous fillers [11
I
C
I
I
I
I
Experimental procedure The studies were conducted with three types of specimens. These specimens were single-layer or multilayer structures with a given sequence of layers: matrix powder and whiskers of short fibers; prepreg based on fibers and matrix powder (Fig.3).
1 - Matrix 2 - Short fibres + matrix 3 - Prepreg Fig. 3. Types of specimensfor the investigation of composite materials
a
b
The compaction of layers were carried out by hot pressing in a graphite mold in nitrogen atmosphere at 1750-1850°C and at 20 MPa. The process lasted 1 hour. Particular attention was given to the microstructural aspects in developing CMC in Si3N4-SiC,, Si3N4-SiCf and Si3N4-Cfsystems and also to the effect of filler first of all on the properties determined by the conditions of crack initiation and propagation, namely, strength and crack resistance. The physical and mechanical properties were evaluated on the specimens measuring 3,5x5x70 mm within a wide range of temperatures (201500°C). The bending strength was estimated by the method of 3-point bending. The crack resistance was determined by the single edge notched beam method. The microstructure was studied by electron and optical microscopy and fractographic analysis.
RESULTS AND DISCUSSION C
-
Fig. 2. Morphology of fibrous fillers: a SiC4AM-7); b - SiCr(Hi(s));c Cfl300)
-
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Different properties of matrix structures and fillers make it possible to solve various material involving and technological problems in CMC synthesis. These are the increase of crack resistance, high-temperature strength, heat resistance, hardness and the development of thin complex-shaped components, multilayer structures, large items.
Composite materials were developed with due regard for the interrelation between chemical composition, microstructure, properties and technology. Physical and mechanical properties of composite materials depend on the realization of following characteristics: highly dense, defectless matrix, a uniform distribution of the filler at a given volume content, a given orientation of the filler in the matrix, optimum interfacial interaction and firm bond at tibedmatrix interface, minimum mechanical damage of the filler in the course of sintering. Complex studies of the microstructure of Sic,, Sicf, Cf /Si3N4 specimens have made it possible to reveal some mechanisms of its formation and the behaviour of various types of fillers in hot pressing: 0 Matrix structure with near-theoretical density is formed in all materials (Fig.4).
0 Large crystals AM-7, TWS-400 and short Cf fibers 6-10 pm in diameter and 2-50 pm in length are distributed uniformly without forming a framework [6]. There are yams and aggregations, which prevent their uniform distribution in matrix powder (Fig.6).
-
Fig.6. Distribution of short Cf fibers in SilNJ CdUKN-P) composite; a - x 400; b x 4000
-
Fig.4. Si3N4matrix microstructure in Si3NJ- SiCw(AM-7) composite material 0 In all cases the crystallographic axis "C" of Sic, whiskers and SiCf and Cffibers in the matrix is normal to the direction of forces in hot pressing (Fig.5) [ 5 ] .
C
II
1. -
-
-
-
-
F i g 5 Fibrous filler orientation in composite materials;a S&N4 SiCdTWS-400)~320;b SIN4 SiCdVNIIPV); c - SilN4 CdUKN-P), x320
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0 Sic, whiskers and Cf fibers remain intact and retain their geometrical dimensions under given hot pressing conditions; the experimental domestic SiCf fibers pass into polycrystal p-Sic (Fig.7).
Fig.8. Microstructure of Si3N4- C1(VNlIPV)composite material
Composite material properties are also determined by physico-mechanical and chemical properties, geometrical dimensions and volume content of the filler. It is taken into account in the development of CMC. The effect of filler characteristics on the final properties of CMC is shown below by way of specific examples. Components relationship A maximum density of the composite material in Si3N4-SiC, system is realized when the content of whiskers ranges from 10 to 30 vol.% and a maximum density and crack resistance are realized at 20 vol.% (Fig.10), and this is in good agreement with the data of home and foreign investigators [7]. The introduction of 10-20 vol.% of short SiCfand Cf fibers is an optimum amount for the realization of crack resistance. Geometrical dimensions Physical and mechanical properties of CMC depend to a great extent on so called scale factor of the filler on the dimensions of the whiskers introduced and lengthdiameter ratio. Large crystals AM 7 and TWS-400 are uniformly distributed in the matrix when the content is 10-40 vol. YO and in this case CMC possess high mechanical properties within the whole range of compositions. Characteristic of small crystals TWS-200 is the existence of local difficult- to-sinter areas as large as 60x60 - 200x300 pm in ceramics (Fig.9).
Fig.7. Fiberhatrix interface: a - SiC,(AM-7); b CAUKN-P); c SiCr (Hi($); d SiCr(VNI1PV)
-
-
-
0 In Si3N4-SiC, system a tight contact of whiskers with the matrix is achieved without any chemical interaction; in Si3N4-SiCfsystem a tight contact at the level of the chemical interaction of the components without the formation of the intermediate interfacial layer is observed; in Si3N4-Cfsystem on the contrary, the matrix and the fibers are not bonded to each other and a mechanical linkage probably occurs by virtue of surface roughness (Fig.7). In case of suspension impregnation, the continuous Sicf and Cffibers 6-22 pm in diameter in a monolayer up to 80 pm in thickness all the fibers are fblly separated from each other by a layer of the matrix material (Fig.8).
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Fig. 9. Distribution of short whiskers TWS-200 in matrix: a x320; b - x2000
-
following order from the viewpoint of increase of their activity with respect to Si3N4matrix: AM7 + TWS200+ TWS-400. In the process of hot pressing at temperatures higher than 1000-1200°C loss of SiCf fibers strength caused by S i c recrystalization is observed. It is accompanied by an increase in the size of crystals. At the same time and the same temperatures the interaction with oxygen increases with the resulting formation of unstable silicon monoxide SiC+02=SiO+CO? and this leads to fiber failure. A more sophisticated analysis of SiCf fibers of various grades in a composite material revealed a significant difference in their structure. Thus oxygenfree fibers Hi with stoichiometry approaching C/Si=1 possess high heat resistance and dense grained structure with the grain size 0,l-0,25 pm. Characteristic of the experimental domestic fibers under these conditions is a looser structure in which certain of the round grains of the same size are not connected with each other (fig.7). Among the fillers being studied TM and NLM fibers occupy an intermediate place and from the standpoint of their structure they are closer to Hi(s). To suppress high-temperature structure variations, especially recrystallization, 4 mass. YO of Ti is introduced into TM fibers composition.
These areas are aggregations of whiskers, which are practically not connected with the bulk. They act as defects inducing composite material failure at lower (660-690) stress and retain their morphology up to 1500°C. When the content of the filler is the same (20 YOof Sic,), a more pronounced (up to 5 pm) growth of Si3N4 matrix grains occurs in case of the reinforcement with large TWS-400 crystals than in case of the reinforcement with small TWS-200 crystals when the grain size do not exceed 2,5 pm (fig.10).
-
Physical and mechanical properties The effect of reinforcement of Sic, whiskers provides an increase of crack resistance of Si3N4-SiC, CMC up to 11 MPamIn and the microhardness - up to 26 GPa. The strength in this case is kept at the level of matrix strength (750-1000 h4Pa) (Tabl.2). The distinctive feature of the reinforced material is the retention of high strength over a wide range of temperatures and its increase up to 850 MPa at 1500°C [8]. Analogously the crack resistance increases up to 14 MPam'" with the rise of temperature up to 1500°C.
Fig. 10. Microstructure of Si,", - SIC, composite material: a - TWS-200 (20°C); b TWS-400 (1 SOOOC)
-
Chemical composition The difference in chemical composition of whiskers defines the difference in the character of the interfacial interaction of the components and the crystaYmatrix interface condition. As the content of the main component decreases and the content of 02, Si02 impurities increases the bonding strength increases and a thin film is formed on crystal surfaces. The whiskers investigated in this paper can be arranged in the
Table 2
Bending strength at various temperatures, MPa
Type of crystals Sic,
2O0C
1300°C
1500°C
Matrix
910
775
820
AM 7
1000
900
8501400°c
TWS-200
705
770
760
TWS-400
1000
775
775
I
I
I
I
Crack resistance at various temperatures, M P a d n
~~
~
Matrix
793
896
10,o
AM 7
ll,o
14,O
-
TWS-200
10.0
14.0
8.6
TWS-400
I
990
I
797
I
894
I
This is due to an optimum phase composition of composite ceramics and such a microstructure, which can be considered as the distribution of small grains of P-Si3N4 between large crystals of p-Sic (fig.10). The evaluation of physical and mechanical properties of CMC shows that a maximum value of the crack resistance in Si3N4-SiCfsystem is 8,2 in Si3N4-Cfsystem it is 9,s MPamln as compared to matrix value equal to 6,6 MPa.mln. In both cases the bending strength decreases by 10-20% and is 600-700 MPa. Such values of the strength are high enough for these systems and they make it possible to increase crack resistance at the expense of strength. These results have been obtained using pure fibers, which we consider as starting point. Thus the tibedmatrix bond for Si3N4-SiCfis very firm and the composite material becomes a brittle monolithic the strength of which decreases. In Si3N4-Cfsystem, on the contrary, the fiberlmatrix bond is weaker and the material from being failed. It should be expected that maximum values of properties will be characteristic of combined laminated specimens with short and continuous fibers (Fig.ll).
Fig. 1 1. Microstructure. of combined laminated ceramics with short and continuous fibers (Ct-UKN-P)
646
CONCLUSIONS
The potentialities of hot pressing have been demonstrated in the development of high-temperature composite materials reinforced with various dispersed and fibrous in Sic,,,, SiCf, Cfl Si3N4systems. Simultaneous increase of strength and crack resistance of composite materials is only possible in case of using SIC, whiskers as the reinforcing filler because they have the highest mechanical properties. The use of such fibrous fillers as Sicf and Cf offer essential advantages for the increase of CMC crack resistance with the decrease of strength by 1,5-2 times as compared to matrix level. The materials developed can be used for combustion chamber segments, rotor blades of gasturbine engines, cutting plates, etc. REFERENCES
(1) S.M.Barinov, V.Ya.Shevchenko, Strenght of engineering ceramics, Nauka, Moscow, (1996) 159. (2) K.Nishio, K.I.Igashira, K.Take, Development of a Combustor Liner Composed of Ceramic Matrix Composite (CMC). Journal of Engineering for Gas Turbines and Power, January, 121, (1999) 12-17. (3) I.Yu.Kelina, I.I.Tkacheva, A.V.Arakcheev, N.I.Ershova, and V.P.Paranosenkov, Hot-pressed constructional ceramic materials. Refractories, 3, (1992) 28-30. (4) Keith R.Karasec S i c whisker characterization: an update, Ceramic Bulletin, 70, [2], (1991) 224-228. ( 5 ) C.Olagnon, E.Bullock, Processing of high density sintered S i c whisker reinforced Si3N4 composites. Ceramics International. 17, (1991) 53-60. (6) M.D.Sacks, H.W.Lee, O.E.Rojas, Suspension processing of A1203/SiC whisker composites. J.Amer.Ceram.Soc., 71, [ 5 ] , (1989) 370-379. (7) Jitendra P.Singh, Kenneth C.Goretta, D.S.Kupperman, Fracture Toughness and strength of SiC-whisker-reinforced Si3N4 Composites. Advanced Ceramic Materials, 3, [4],(1988) 357360. (8)Hiroya Ishizuka and Yoichi Saida, Fabrication of S i c whisker reinforced Si3N4 ceramics turbine nozzles by slip casting, Proceedings of the International gas turbine congress, Kobe ( 1999) 29 1296.
HIGH TEMPERATURE Si3N4-BN COMPOSITE N.I.Ershova, 1.Yu.Kelina The State Research Center of Russia, Obninsk Research and Production Enterprise “Technologiya”, Obninsk, Russia
ABSTRACT Thermal and mechanical properties of Si3N4-BN composite and its behavior under severe thermal loads have been estimated. The characteristic feature of hotpressed Si3N4-BN material is the possibility of properties control over a wide range (bending strength 70-7OOMPa) by varying the ratio of starting components. This technology makes it possible to manufacture functional-gradient material and sandwich constructions in given systems and thus to extend the range of ceramic materials applications from engine members to heatresistant parts of aerospace equipment, bearings working in agressive medium. Results of development and testing of the articles for various applications are presented.
INTRODUCTION While the requirements to gas-turbine engines (GTE) become more exacting it is necessary to develop new materials capable of withstanding high temperatures and mechanical load. Promising materials from this point of view are ceramics of Si3N4-BN system which possess the optimal combination of properties - high strength, heat-resistance and good machinability [1,2]. In various elements of GTE, such as above-rotor seals (ARS), thermostable elements and also in separators of bearings it is necessary to provide moderated hardness and at the same time high strength and low oxidizability at temperatures up to 1300OC. In the Si3N4-BN system it is possible to produce composites with various designed properties and multilayer ceramics due to the feasibility to vary material properties with the variation of components ratio. So, prefabricated elements of ARS can be simplified by means of manufacturing of unbroken article with variable composition which is made of composite material.
EXPERIMENTAL PROCEDURE In this work the investigations were carried out on developing the composite ceramics in the Si3N4-BN system with the use of hot pressing process. Ultradispersed powder compositions Si3N4 - Y 2 0 3 (MgO) (particle size 0.05-0.12 p) made by joint-stock company “Neomat” (Riga city) were used for making
the matrix and coarser particles of hexagonal BN of 1.5-2.5 p in size were used as a filler. Mixing of ultradisperse composition Si3N4-Y203 (MgO) with boron nitride was carried out in combination with effective grinding for 100 hours. BN fraction of total mass was varied in the range of 1O-6O%. Hot pressing was carried out with the use of induction heating in graphite molds in nitrogen atmosphere at 1700-175OOC under pressure of 15-20 MPa. While manufacturing multilayer articles the additional stage is introduced on preliminary briquetting of powders of required composition and its laying in graphite mold in given sequence. The technology makes it possible to manufacture the articles with various schemes of arrangement of layers -horizontal and vertical [4]. In order to cany out a research into microstructure the equipment and methods of X-ray-phase and microstructural analysis, ultrasonic inspection and scanning electron microscopy were used. The bending strength has been estimated by methods of 3- and 4-point bending at 20 and 1300°C. For the Si3N4-BN system the experimental dependencies of change of thermo-physical properties up to 900-1300OC have been determined for the first time. Property indexes were estimated together in the course of one heating operation. Resistance of composite material to hightemperature oxidation was estimated under stationary heating at 1300°C for 50 hours and at 750-900°C for 250 hours. Service life of the specimens fi-om the material of this system was checked by 10-cycle tests in gas flow of rate 3000 W/(m2. K) at 153OOC. Heat-resistance of the material was determined by standard methods of water thermo-changes and by estimation of heat-resistance criteria R,and %. On the base of the investigated thermo-physical and physicaYmechanica1 characteristics the thermostrength calculations have been carried out for thermostressed state (TSS) for the specific articles from uniform and multilayer materials in designated operating conditions. Integration step over time in the problem of non-stationary heat-conductivity was taken equal lO-’s, over thickness - 100 nodes. Calculations were carried out in one-dimensional setup. The properties of matrix material and materials with various content of BN were given as dependencies on temperature. The accuracy of its evaluation defines the accuracy of calculations. During calculations the maximum values of tensile stresses om= were estimated which characterize the strength of articles to a greater degree. Comparative analysis was performed over values of thermal stresses that arose in articles during
647
heating and also over bending strength of uniform material. On this basis the safety margin of the material and its individual layers was estimated.
RESULTS AND DISCUSSION Microstructure The structure of hot-pressed specimens with a low content of BN represents the matrix, having the inserts of BN-plates, which are arranged perpendicularly to the direction of hot pressing. Si3N4 matrix has finely granular and compact structure with elongate morphology of grain and high degree of crystallinity. As the BN fraction of total mass increases the gradual substitution of Si3N4-matrixby BN-matrix occurs that results in decrease of packing tightness (fig.1). Besides, as the BN fraction of total mass increases to 40-60%, the boron-containing phases are distributed more uniformly (fig.2). The porosity of composite material remains practically at zero level even in specimens with a high content of BN. This is the advantage of the ceramics developed over foreign analogues.
The main phases in composite material are kSi3N4, BN. The intergranular phase is presented in various refractory compounds such as yttrium silicates, SiON2, 4Y203.Si02- Si3N4, YSi02N, Si3N4.Y203, 10 Y203*9Si02.Si3N4which increase high-temperature strength. The interlayer boundary is visible well in multilayer specimens. When BN content in composition increases to 50%, this boundary becomes more sharp. Stresses which occurred on boundary of layers due to considerable difference in their physicaYmechanical properties were eliminated by introduction of intermediate layer. Depending on the direction of property gradient the various schemes of layer arrangement were used in the article. Thus, while pressing ARS in the form of inserts or segments the briquettes were laid in vertical row, while pressing monolithic rim with internal wear layer - in one horizontal row. Petrographic analysis brought out the essential differences in the interface nature in both cases: in horizontal scheme of laying the smooth transition from one layer to another is watched independently of the BN content in them, in vertical scheme sharp interfaces between layers (fig.3)
-
10 % BN 30150 % BN
40 % BN
Fig. 1. Si3N4- BN composite material microstructure of various BN content
10% BN
40% BN
Fig. 2. Si3N4- BN composite material microstructure in characteristic beams of boron
648
0150 % BN Fig.3 - Interfaces between layers while laying briquettes according to various schemes This can be explained by migration of BN grains through the layer boundaries in the definite planes due to the presence of liquid phase in hot pressing and by low friction coefficient when the BN particles glide. Besides, maximum force which affects the powder articles in pressing is directed perpendicularly to the effort of hot pressing.
Mechanical characteristics The research into mechanical characteristics of composite material Si3N4-BNdemonstrated that strength value varies inversely as the BN content in composition from 660 to 75 MPa at room temperature and from 480 to 60 MPa at 1300°C.The approach of values of hightemperature- and room condition strength was also observed as the BN fraction of total mass increased (figAa). This is accounted for by the material progressively acquires the properties typical of pure BN - low strength and high resistance to temperature action - as the Si3N4matrix is substituted by BN matrix which possesses less density. For the same reasons the hardness index of material falls and wear index increases as the BN content increases (figAb). (Jb.
MPa
~~
~~
BN mass content, ./, 70
In mechanical properties the developed material 1.5 times surpasses composite materials in analogous
system that have been produced by means of slip casting and hot pressing [3,6,7]. Regularity of change of mechanical properties of multilayer material depending on BN fraction of total mass is the same as this regularity of uniform specimens. Specimens with horizontal arrangement of layers were failing mainly along a layer which possessed the elevated BN content or on the boundary between this layer and intermediate one. Thermophysical properties Research into thennophysical characteristics temperature conductivity (a), heat capacity (C), and heat conductivity (A) demonstrated that temperature conductivity (a) does not rigidly depend on addition mass of BN. This is accounted for by the nature of dependences a=f((r)for Si3N4and BN is the same, and they are close with each other in absolute values (fig.5). Heat capacity indexes increase as BN fraction of total mass in material increases ; in so doing the same nature of their dependence on temperature remains - as the temperature increases the heat capacity index increases too. The same temperature dependence is noted also concerning the heat conductivity but its indexes fall as BN content increases in material (fig.5,
0.07
0,OB
g
0.05
OPM# 0.03
20
0.02
10
0,Ol
0
CI)
b'
I0 20 30 40 50 BN mass content, %
Fig. 4. Mechanical properties of Si3N4-BN composite material a) bending strength dependence on BN content b) hardness and wear indexes dependence on BN content
Fig. 5 . Temperature dependencies of temperature conductivity u, heat conductivity 1 and heat capacity C, of Si3N4-BN material
649
Refractoriness Refractoriness tests showed high resistance of the material to oxidation - the change in mass was 0.0050.8% in compounds of 10- 60% BN in the fmt regime of heat treatment (T=1300°C; r =50 hours) and 0.20.4% in the second regime of heat treatment. When gas flow of rate 3000 W/(m2.K) at 1530°C affected the multilayer material the mass change was not more than 0.2%. It is related to the processes of BN erosional carry-over and oxidation which are accompanied by release of Bz03. Absolute values of mass changes depend directly on BN content in the layer and on this layer volume. Heat resistance Evaluation of heat resistance of composite materials by water thermo-changes standart methods did not bring out essential differences in their behavior - whatever the BN content all the materials withstood temperature difference over 1000-12OO0C, that exceeded the matrix material heat resistance of a minimum by 160-300 degrees. Heat resistance is affected most of all by ab, modulus of elasticity E and coefficient of heat conductivity a. Among known property complexes for the evaluation of heat resistance by criteria the criteria R, = G( 1-p)/Ea and & =RIA are most commonly used in nonstationary thermal action [ 5 ] . Calculation of heat resistance criteria R1 and & for composite material properties of the whole range of compounds at 20 and at 900-1300°C demonstrated that they did not increase monotonically (as from data reported in the paper [6]) but they have maximum values within the region 2 0 4 0 and 60% at room temperature and at BN 2040% at elevated temperature (table 2).
Material composi-
Criteria of R, and 4calculation at temperatures, 'c:
R,
I
Rc
The calculations were performed in engine starting and stopping regimes for ARS inserts and in regime of thermal shock for the washer. Thermal flow which affects the surface of ARS inserts and washer was given in terms of gas temperatures 1150 and 1700°C and heat-transfer coefficients which are equal to 3000 and 14000W/(m2K) and are constant in the course of heating. The results of more than 100 calculations are presented in the form of graphs and nomographs of stress time changes of omax and 0- over washer thickness and over individual layers of ARS inserts and graphs of distribution of temperatures and stresses in the moments of maximum stresses (fig.6). Q.
MPa
I
Fig. 6 Stress distribution in washer a) over time b) over thickness The TSS analysis of multilayer insert made it possible to explain some peculiarities - for any combinations of layers the introduction of layer with 10% BN resulted in the severe increase of om= in adjacent layers (fig.7). It is related to the sharp difference in linear expansion thermal coefficient of this layer from other compounds. The data obtained in the course of tests of multilayer insert and uniform washer conform well to calculation results. BN content \ T, "C SO%<<% Id% 0% us MPa 120
Estimation of thermo-stressed state (TSS) The optimization of geometry and composition of the uniform material of heat resistance articles and multilayer structure of above- rotor seals (ARS) was performed by numerical simulation method. The choice of articles is due to differences in their structure, operating conditions (washer is in considerably harder regime of thermal action), work time (hours for ARS, some seconds for washer) and others.
650
780
60
520
0
260
-60
-
-
Time 1600 sec Temperature Stress
-120
-
Fig.7. Temperature and stress distribution at i-th second of heating in 4-layer insert of ARS
Thus, the method proposed is integral part of article development route. It makes it possible to simulate the structure and geometry taking into account the minimization of thermal stresses. It considerably simplifies and accelerates the article development process.
The most important of them are the elements of gas-turbine engines - above-rotor seals of various configurations including uniform inserts 20x30~5mm in size, multilayer segments, integral rims 250 mm in diameter having internal wear layer, heat resistant washers, high- speed bearing details (fig.8).
REFERENCES
CONCLUSION The high-temperature composite material is developed in the Si3N4-BN system. It possesses high physicaVmechanica1characteristics and it is serviceable in the extreme thermal conditions, in short time at 1700OC. An important peculiarity of material is possibility of controlling its properties by means of varying ratio of main components. Manufacture technology of multilayer articles with property gradient over section is developed. The material has been widely evaluated in the structural application articles.
-
Fig.8. Articles from Si3N4-BN composite material
W. Sinclair, H. Simmons, Microstructure and Thermal Shock Behavior of BN Composites. J.Mater.Sci.Lett.,6,( 1987) 627-629. (2) H. Lutz Ekkehard, Michael V. Swain, Fracture Toughness and Thermal Shock Behavior of Silicon Nitride- Boron Nitride Ceramics. J. Amer. Ceram. SOC.,75,(1992) 67-70. (3). F. Toshihiko, I .Keiishiro, H. Akira and U. Ryoji, Mechanical Properties and Microstructure of Si3N4-BN Composite Ceramics. Ceram. Mater. And Compon. Engines: Proc. 3rd. Intern. Symp., Las-Vegas, Nov.27-30, 1988, Westerville (Ohio), (1989) 968-976. (4) N.I.Ershova, I.Yu.Kelina, Multilayer Ceramic Articles from Si3N4-BN-based Composite Material. Refractories and Technical Ceramics, 1, (1997) 6-10. ( 5 ) V. Dauknis, K. Kazakyavichus, G. Prantskyavichus and V. Yurenas, Research into Refractory ceramics Heat Resistance, Vilnius, Mintis, (1 97 1) 150. (6) I. Keiishiro, F. Toshiaki, 0. Kazuki, et al., Microstructure and Mechanical Properties of Si3N4-BNComposite Ceramics Obtained by Slip Casting. J. Iron and Steel Inst. Jap., 5 , N29 (1989), 1612-1619. (7) I. Keiishiro, F. Toshiaki and U. Ryoji, Development of Machinable Si3N4-BN Composite Ceramics. Kawasaki Steel Giho, 21, N24,(1989) 281-286.
(1)
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PROCESSING OF SICN-FIBRES PREPARED FROM POLYCARBOSILAZANES Jurgen Hacker, Gunter Mob*, Gunter Ziegler University of Bayreuth, Institute for Materials Research (IMA I), D-95440 Bayreuth, Germany
ABSTRACT
gaseous ammonia in toluene, and was described in detail elsewere [11. The number average molecular weights of the polymers were determined by cryoscopy in cyclohexane while the melt viscosities result from rheological measurements with a cone-plate-rheometer. The volatile species formed during pyrolysis and the ceramic yields of the precursors were measured by FTIR-TGA. The pyrolysis process was characterised in detail by MAS-NMR-and FTIR-spectrometry, XRD and elemental analysis of heat treated powder samples.
Polycafbosilazanes synthesised by treating bis(chlor0sily1)ethanes with gaseous ammonia in toluene (brittle, meltable solids, 80 wt% yield) were characterised by cryoscopy, viscosimetry, FTIRcoupled thermogravimetry (FI'IR-TGA), FTIR-spectrometry, liquidXRD and elemental and MAS---spectrometry, analysis. Melt spinning or drawing from the melt of optimised precursors resulted in polymer fibres of about 10 to 30 pn diameter. These fibres revealed a smooth surface, and no pores or cracks were detected in the fracture surface. Gas phase curing, causing an m e l t a b l e d a c e layer on the fibre, was chosen as process for making fibres infusible before pyrolysis. Different curing agents and various curing conditions were tested. Properties of the cured fibres varied with the curing agent concentration, the curing temperature and the curing time. Electron beam curing was tested as an alternative curing method. Pyrolysis of the cured fibres at 1000 and 1100 "C, respectively, in flowing nitrogen led to ceramic fibres with high ceramic yield (70 wt%).
Melt spinning of as received precursors was performed in a pilot plant (Fraunhofer ISC,Wiirzburg, Germany) and resulted in polymer fibres of about 10 to 30 pn diameter. Fibres with comparable diameters were drawn from polymer melts of thermally retreated precursors. The curing of the polymer fibres occurred by a chemical gas phase curing process (chlorosilanes and aluminium alkyles as curing agents, various curing conditions) or by electron beam (topic of present work). The cured fibres were pyrolysed in a flow of nitrogen up to 1000 "C (gas phase cured fibres) and 1100 "C (electron beam cured fibres), respectively, in a graphite tube furnace (1 hour annealing).
INTRODUCTION
RESULTS
During the last few years fibre reinforced CMCs have become more and more interesting for high temperature applications in oxidising and/or corrosive environments. For those applications as well as for tribological ones fibres have to M i l requirements like high tensile strength, high thermal stability, good oxidation and creep resistance and/or lowcost processing for use in such CMCs. Commercially available ceramic fibres with good high temperature properties are too expensive for economic industrial applications, and at present the quality of low cost fibres is not good enough. The aim of this work was to develop preceramic polymers for relatively lowcost SiCN-fibres which reveal sufficient properties.
Polymer Characterisation The polymers resulting in high yield (80 wt%) are brittle meltable solids with number average molecular weights range from 1000 to 2000 g/mole and different melting points depending on processing conditions. The melt viscosity at 120 "C ranges from 5 to 80 Pas. The melt exhibits Newtonian rheological behaviour. The mass loss of about 30 wt% up to 1500 "C determined by FTIR-TGAexperiments occurs in two steps between 250 and 500 "C (-15 wt0/0, distillation of low molecular oligomers and crosslinking by deamination) and between 500 and 800 "C (-15 wt%, methane and ethene as IR-active volatile species). From 800 to 1500 "C only about 2 wt% of mass are lost, presumably owing to the elimination of the hydrogen remained in the sample. Above 1500 "C the pyrolysis behaviour strongly depends on the furnace atmosphere. In argon the decomposition of the amorphous SiCN-matrix into siliconcarbide and molecular nitrogen occurs unhindered with a mass loss of another 15 wt%. Nitrogen atmosphere stabilises the matrix and only
EXPERIMENTAL PROCEDURE Polymer synthesis, characterisation of the precursors, fibre spinning, curing and pyrolysis of the cured fibres took place under inert gas conditions. The synthesis of the polycarbosilazanes was performed in a pilot plant by treating bis(chlorosily1)ethanes with
653
about 2 wt% mass loss were observed during a 1 hour annealing. The pyrolysis behaviour was examined in more detail by MAS-NMR- and FTIR-spectrometq on powder samples pyrolysed up to several temperatures with 5 hours annealing time in a nitrogen atmosphere. The 29Si-MAS-NMR-spectra give most information about the changes in the polymer and the matrix structure, respectively, during the pyrolysis process (Fig. 1). I
I
I
amorphous matrix which decompose during further heat treatment to siliconcarbide, free carbon and molecular nitrogen.
Fibre Spinninflrawing Melt spinning of optimised precursors results in polymer fibres with 10 to 30 pm diameter depending on the die diameter (175 to 300 pm) and the withdrawal speed (400 to 800 d s ) . The spinning pressure is generated by nitrogen gas pressure (up to 10 bar) and by a single screw extruder, respectively (up to 25 bar). Fibres drawn from the melt in a glovebox system have diameters in the same range. All polymer fibres have smooth surfaces and there are no cracks or pores in the fracture surfaces (Fig. 2).
Fig. 2 Extruder-meltspunpolycarbosilazane fibre
200
100
0
-100
-200
Chemical Shift [ppm] Fig. 1 %-MAS-NMR-spectra ofthermally treated polymer powder samples
The polymer structure remains unchanged up to 500 "C. At this temperature line broadening indicates the beginning of the organic-inorganictransition. At 700 "C the formation of the amorphous SEN-network is almost finished. Up to 1500 "C there are no sirnicant changes in the spectra. In this temperature range only local structural rearrangements seem to take place. The FTIRspectra from comparable powder samples support the NMR-results. At 1500 "C a narrow signal at about 4 9 ppm appears indicating the formation of a-Si3N4verified by XRD as the only cristalline phase at that temperature. The 1600 "C-spectnun shows only another narrow signal at -19 ppm characteristic for SiC4environments. The corresponding XRD-spectrum indicates crystalline Sic (mainly 6H-type). Between 1500 and 1600 "C first SiN4enriched domains are formed in a carbon-rich
654
Fibre Curing For the curing of the polymer fibres by a chemical gas phase curing process Merent curing agents like trichlorosilane (TCS) or trimethylaluminium (TMA) and different curing conditions were tested. The aim of chemical gas phase curing is to generate a infusible surface layer on the polymer fibres to stabilise the fibre shape during pyrolysis. TCS is known as an appropriate curing agent for polysilazane fibres [2]. The low boiling point (32-34 "C) enables a process temperature of 40 "C avoiding fibre softening. The SEM micrograph of a TCS-cured fibre (Fig. 3a) shows a thin surface layer formed by reaction of the TCS with the polycarbosilazane. The thickness of the surface layer strongly depends on the curing time and the curing agent concentration. Long curing times or high TCS concentrations result in surface layers of some microns in thickness. The handling of such fibres is very difficultbecause of their brittleness. A further useful agent for polysilazane crosslinking is TMA [3], which reacts with the NH-groups of the polymer forming methane as a volatile species. TMAcured fibres also show thin infusible surface layers (Fig. 3b). The dependence of the fibre shapes and properties on the curing conditions is the same as for TCS-cured fibres. In contrast to a gas phase curing process, curing by electron irradiation is a widespread bulk curing method in the field of Sic-fibre curing since a few years [4].
The fibres are cured homogeneously and no differences to the uncured state are observed (Fig. 3c). The relationship between irradiation doses and properties of the cured fibres is a topic of present research.
to strong, the fibres partly crack during pyrolysis, and the handling of the ceramic fibres is not possible because of their brittleness. Curing under the right conditions results in ceramic fibres with unchanged shapes and acceptable brittleness (see Fig. 4a).
b) TMA and c) electron beam
Fibre pyrolysis All curing methods discussed above result in crack-free amorphous ceramic fibres with relatively smooth surfaces after pyrolysis to 1000 and 1100 "C, respectively (Fig. 4), with ceramic yields of about 70 wt%. The shape of the pyrolysed fibres cured by TCS or TMA depends on the curing conditions: if the fibres are not cured enough, the infusible surface layer is not able to avoid fibre deformation (see Fig. 4b). If the curing is
Fig. 4 F'yolysed fibres (nitrogen atmosphere): a) TCS-cured (1000 "C); b) TMA-cured (1000 "C); c) cured by electron beam (1 100 "C)
Successful chemical gas phase curing is not possilbe on contact areas between the fibres and if the fibre diameter is too large. During pyrolysis the fibres will melt together at the contact areas because their surfaces are not cured there. Consequently fibre contact must be avoided during the curing process. Too large diameters
655
require strong curing conditions to avoid fibre deformation during pyrolysis. The consequence is an enhanced brittleness of the ceramic fibres. In contrast, ceramic fibres cured by electron beam do not change their shape (Fig. 4c) and do not fuse during ceramisation. The moderate brittleness of these fibres allows their handling without any di&culties.
CONCLUSIONS Reacting bis(chlorosily1)ethaneswith gaseous ammonia in a pilot plant scale results in polycarbosilazanes, which can be extruder-meltspunor drawn from the melt into polymer fibres with diameters of 10 to 30 p. Successful curing of these fibres is possible by a gas phase curing process (trichlorosilane or trimethylaluminiumas curing agents) or by electron beam. Pyrolysis of the cured fibres up to 1000 (gas phase cured fibres) and 1100 "C (electron beam cured fibres), respectively, in an inert gas atmosphere leads to crackfree amorphous ceramic fibres with a ceramic yield of 70 wt%.
Acknowledgements The authors would like to thank S. TraB1, University of Bayreuth, for recording of the MAS-=-spectra, the Fraunhofer ISC, Wiirzburg, for fibre spinning, D. Schawaller, University of Stuttgart (ICF) for fibre curing by electron beam and the Bayerische Forschungsstiftungfor the financial support.
REFERENCES (1) J. Hacker, G. Motz and G. Ziegler, Neuartige SiCN-Polymere als Vorstufen fiir keramische Fasern, Proc. Tagung ,Verbundwerkstoffe und Werkstofierbunde', Hamburg, Germany (1999), K. Schulte und K. U. Kainer (edts.), 459-463. (2) G. E. LeGrow, T. F. Lim, J. Lipowitz and R. S. Reaoch, Ceramics from Hydridopolysilazane. Am. Gem. SOC. Bull., 66 [2], (1987) 363-367. (3) D. Seyferth, G. Brodt, B. Boury, Polymeric Aluminasilazane Precursors for Muminosilicon Nitride. J. Mater. Sci. Lett., 15, (1996) 348-349. (4) K. Okamura and T. Seguchi, Application of Radiation Curing in the Preparation of Polycarbosilane-Derived Sic Fibers. J. Inorg. Organornet. Polym., 2, (1992) 171-179
656
DESIGN OF GRAIN BOUNDARY PHASES IN SILICON NITRIDE SILICON CARBIDE NANO-COMPOSITES
-
N. K. Schneider, M.J. Pomeroy and S. Hampshire Materials and Surface Science Institute University of Limerick, Ireland
ABSTRACT Glass ceramic-Sic nano-composites have been produced by introducing S i c into a Y-Si-AI-0-N glass. The results of these studies have been applied on Si3N4-SiC nano-composites to design the grain boundary phase composition of these composites. It has been shown that the choice of additives ratios and adequate heat treatment can lead to improved mechanical properties of the nano-composites such as wear resistance, elastic modulus, strength, hardness and toughness.
INTRODUCTION Although silicon nitride is a leading material for a number of applications, its use has been limited by the fact that the mechanical properties begin to degrade at temperatures above 1200°C. Due to the need for sintering aids, silicon nitride based ceramics contain oxynitride glass phases at the grain boundaries which can impair subsequent high temperature properties. Understanding their nature and behaviour is essential for the improvement of the Si3N4 matrix. Monitoring the amount and composition of these grain boundary phases can provide improved chemical and mechanical propertie~.['l[~1[~1 In this work an oxynitride glass of the atomic composition Y15.3Si,4.7Als.75052.~Ns.75has been studied which is known to produce 'B-phase' (Y2SiAI05N) as the major crystalline phase upon heat Subsequently the same ratios of Y, Si, and A1 as in the glass have been introduced into silicon nitride in various amounts i n the form of oxide powders as densification additives. B-phase has successfully been reproduced as a crystallisation product at the grain boundary of these silicon nitride matrix materials. Secondary phase additions such as S i c offer the potential for improvements in fracture toughness, creep behaviour and in flexural strength of Si3N4 at higher temperatures. Therefore nano-size S i c has been introduced into the silicon nitride matrix. Again the grain boundary compositions were
closely monitored. In addition, mechanical testing was carried out to characterise the resulting materials and their properties. Rate-controlled-sintering was carried out on the samples to optimise the densification conditions for pressureless sintering.
EXPERIMENTAL Oxynitride glass of the composition Y1~.~Si14.7A18.75052.5Ns.75 has been produced from raw oxide powders. The glass was studied and characterised by differential thermal analysis, dilatometry and X-ray diffraction. Optimum heat treatment schedules have been determined in order to maximise the formation of B-phase. This glass composition was crushed, milled and mixed with nano-size S i c (Mitsui Toatsu MSC 20, p-Sic, d50=0.15pm). Different heating cycles at temperatures between 1100 and 1300°C under nitrogen have been applied. Silicon nitride matrix and composite materials were produced using an aqueous process including attrition milling and freeze drying. Sintering studies were carried out using a high temperature dilatometer set-up. Matrix samples and composites were pressureless sintered under nitrogen atmosphere. Post sintering heat treatment have been applied in order to crystallise the grain boundary phases. Phase analysis has been carried out by XRD. Elastic modulus was assessed by ultrasonic measurements, micro-hardness was calculated from Vickers indentation. Flexural strength and fracture toughness was measured by four-point bend set-up on ground samples ( 2 0 ~ m ) .Fracture toughness was calculated from strength test samples which had been indented with different loads in the centre of the tensile surface. Microstructure of the samples has been assessed by scanning electron microscopy.
RESULTS Glass-Sic composites The densification behaviour of oxynitride glassS i c composites has been studied by dilatometry. Figure 1 shows the densification rate of the glass
657
with and without S i c addition. Densification starts at around 950°C, reaches a maximum at 1030°C and is complete at 1100°C. This is recorded for the glass matrix as for the glass-Sic composite. Glass-Sic composites have been sintered at various temperatures in order to identify the cycle which results in optimum densification. Figure 2 shows the result of these runs. The highest density was achieved after sintering at 1100°C for one hour.
ed by applying post sintering heat treatment known to give optimum crystallisation of B-phase. This schedule involves a 10h nucleation stage at 960°C followed by 5 hours at 1050°C for crystal growth. The microstructure of the heat treated glass and glass-Sic composite was studied by SEM. Figures 4 and 5 show the crystal structure of the glassceramic matrix and the glass composite after heat treatment.
Elastic modulus and micro-hardness have been assessed for the glass and the glass-composite, both as sintered and with a post-sintering heat treatment. As can be seen in Figure 3, the addition of S i c improved Young's modulus and microhardness. A further increase could be obtain-5 h
ij
s D W
I
0
700
600
800
900
1000 1100
1200
Temperature ("C)
I
Fig.4: Microstructure of glass-ceramic matrix
F i g . 1: Densification rate of glass-composite 100
3.9
h
h
2
98
03.8
Y
3
96
%
.-g3.7
n
L
94
0)
c)
3.6
.-* zi CT
92
-
1050 I100 1150 1200 Sintering Temperature ("C)
Fig.2: Densities of glass composite with 1Ovol-% Sic with different sintering temperatures Fig.5: Microstructure of glass-composite with IOvol-% Sic after crystallisation treatment
200
Judging by the grain sizes it can be seen that the sample with 1Ovol-% S i c has a larger number of smaller crystals. S i c seems to promote nucleation and inhibits crystal growth.
150 h
2 100
9,
50
0 Glass
Glass+lO% SIC
Glaes+l Ooh SIC, heat treated
Fig.3: Young's ( E ) and shear modulus ( G ) of g l a s s - S i c composites 658
Silicon nitride-Siliconcarbide composites Silicon nitride matrix and composite materials have been sintered with different additive amounts. For this, Y203 and A1203 were added in the weight ratio 4 : l in order to enable the formation of B-phase in the grain boundary phase. Sili-
con nitride matrix material was produced with 5 , 10 and 20wt.-% additives. The sintering behaviour was studied using a high temperature dilatometer in combination with a rate-controlled sintering program. Figure 6 shows the sintering curves of the matrices with the different additive amounts. Temperature
2000r t
-
5-83
-
-------
1.05
I
1750
1
1500
0.95
1250
0.9
1000
0.85
750
0.8
500
0.75
35
30
25
83- as sintered -83- sinteredand heat treated
Angle (2@)
Fig.7: XRD spectra of silicon nitride matrix with 10% additives
IGPa
7
0.7
250 0
60 lime[min]
Temperature
2000
20
15
T
-
120 %
10-83 l__l_
_ I _
7
1.05
10-
10-
B3+2.5
B3+5
10B3+10
10-
B3+20
Fig.8: Young's modulus of composites
1083
(Vo'.-")
0.8 0.75
0.7 0
60
limetmin] 120
."g. 6: Densification behaviour of silicon nitride matrix with 5, I0 and 20wt.-% additives The sample with 5% additives (5-B3)could not be sintered to full density at a maximum temperature of 1850°C. The composition with 10% additives (10-B3) reaches full density after 50 minutes at 1850°C. The sample with 20% additives can be sintered to full density at temperatures as low as 1600°C with a holding time of 30 minutes. Applying the same heat treatment schedule used for the glass composite revealed that B-phase could be crystal-
B3+ 2.5%
sic
B3+ 5%
sic
2083
20B3+ 2.5%
sic
20B3+ 5%
sic
Fig.9: Microhardness of composites with I0 and 20% sintering additives lised in the grain boundary phase. Beside B-phase also yttrium-disilicate (YzSiz07) developed in the grain boundaries as can be seen in the XRD spectra shown in Figure 7. Properties have been assessed on matrix samples and composites. Elastic modulus for the samples with 10% additives with increasing S i c content is shown in Figure 8 for the as-sintered and the sintered and heat treated samples. A maximum was obtained for the sintered and heat treated sample containing 5% Sic. Microhardness was measured on matrix and composite samples with 10 and 20% additives. Again, the composite containing 5% S i c has the highest value. Flexural strength was assessed by four-point bend test. Samples were tested after crystallisation heat treatment 659
Figure 11 shows a double logarithmic graph of the matrix material with 10% sintering additives. From the strength values obtained for the samples that had been indented with 25 and l00N the fracture toughness of 6.01MPa can be calculated.
CONCLUSIONS
10.83
lOB3+6%SiC
I
2083
2oBS+s% sic
Fig. 10: Flexural strength of matrix and Composites
on the machined bars. Flexural strength is shown in Figure 10 for the composites with 10 and 20% sintering additives. The highest strength value is reached for the matrix composition with 10% additives. Prior to testing strength, part of the samples were submitted to a Vickers indentation of various loads in the centre of the tensile surface. Knowing the indentation load, Young's modulus, hardness and strength, fracture toughness can be calculated after [Eq.l] [61.
p,
K,, = 77; ( E / H y 8(oP3
[Eq.1I where E is Young's modulus, H is Vickers hardness, cs
Oxynitride glass-ceramic-Sic and Si3N4-SiC nanocomposites have been produced successfully. Introducing S i c into glass ceramics improves mechanical properties such as hardness and Young's modulus. Heat treatment schedules lead to further increases of these values as a result of crystallisation of the matrix. Si3N4 matrix and composites have been produced with various amounts of Yz03 and A1203 as sintering additives in a constant cationic ratio. Post-sintering heat treatment can achieve the crystallisation of the grain boundary phase into a known crystalline phase based on the heat treatment applied to glass composites. Elastic modulus and hardness increase with the addition of up to 5% nano-Sic. Strength and fracture toughness decrease with increasing S i c addition. Maximum performance was measured for the samples with 10% sintering additives.
R
the as-indented strength, P the maximum load and q v = 0.52 is a geometrical, material-independent constant. This way fracture toughness was calculated for the composites with 10 and 20% sintering additives. Table 1 lists the toughness values obtained. The highest toughness was measured for the matrix containing 10% additives. With the addition of Sic, toughness decreased, similar to the strength values.
ACKNOWLEDGEMENTS The authors wish to thank Dr. F. Cambier, Belgian Ceramic Research Centre for his technical assistance. Part of this work was funded by the European Commission under Contract No. BRMA-CT97-5055.
REFERNCES 10-B3 6.01 4.36 4.32
0% S i c 2.5% S i c 5% S i c
20B3 5.27 4.56
1000
.
I
* n4
Fig. I I : Strength values depending on indentation load for the 10% additive matrix
660
111 Terwilliger, G.R. and Lange, F.F., J.Mut.Sci., 10 (1975) 1169. c21 Hampshire, S., Drew, R.A.L., Jack, K.H., J. Am. Ceram. SOC., 67 (1984) C46. [31 Ziegler, G., Heinrich, J. and Wotting, G., J.Mat. Sci., 22 (1987) 3041. r41 Lemercier, H., Rouxel, T., Fargeot, D. and Besson, J.-L.,J.Non-Cryst., Sol. 201 (1996) 128. [51 Ramesh, R., Nestor, E., Pomeroy, M.J. and Hampshire, S., Key Engin. Muter., 99-100 (1995) 21 1. 161 Chantikul, P., Anstis, G.R., Lawn, B.R. and Marshall, D.B., J. Am. Ceram. Soc., 64 [9] (1981) 539-43.
Yb-Si-Al-0-N GLASSES AND GLASS-CERAMICS AS GRAIN-BOUNDARY PHASES FOR SILICON NITRIDE MATERIALS Y. Menke, S. Hampshire, S. Newcomb, Materials and Surface Science Institute, University of Limerick, Ireland
ABSTRACT
EXPERIMENTAL PROCEDURES
The preparation of a glass and its devitrification at 1200°C in the Yb-Si-Al-0-N system with the composition in equivalent % of 35eloYb:45e/oSi:20e/oAl:83eloO: 17eIoN is reported. The properties of the glass and the glass-ceramic are examined. Changes observed in density, Young's modulus, hardness and fracture toughness during devitrification are represented. Oxidation studies have been carried out to examine the oxidation resistance of the glass and the glass ceramic under air and under water vapour. For the glass and the glass-ceramic no weight changes were observed up to 900°C. A weight gain of 1% was observed on heating under air during heating to 1200°C for the glass. For the glass-ceramic, no weight changes are observed up to 1200°C.
Yb203 (Rare-Earth Products Ltd), SiOz (FIuka Chemika), A 1 2 0 3 (Aldrich Chemicals) and Si3N4 (HC Starck) powders were mixed together in the right proportions with 2-Propanol (BDH) for 10 hours in a ball mill with sialon-balls as milling agent. The powders were then dried and a large pellet (-50 g) was compacted in a cold isostatic press. This was melted for 1 hour under a nitrogen atmosphere at 1715°C in a vertical tube furnace in a boron nitride lined graphite crucible. After rapidly removing from the hot zone, the glass was poured into a graphite mould, which had been preheated to 900°C. The samples were then subjected to an annealing treatment for Ih at 850°C prior to slow furnace cooling. A heat-treatment for 10 hours at 1200°C under a nitrogen atmosphere was carried out in a horizontal tube furnace.
INTRODUCTION The potential use of silicon nitride ceramics in high temperature engineering applications has created considerable interest over recent years. Because selfdiffusion processes are relatively slow in silicon nitride based materials, oxide sintering additives are required to provide conditions by which liquid phase sintering can occur [ I , 21. During sintering, the additive reacts with silica on the surface of the silicon nitride to form an oxynitride liquid, which, when cooled, remains as an intergranular glass [3, 41. The composition and volume fraction of such oxynitride glass phases determine the properties of the material, and, in particular, they have been shown to control its high-temperature mechanical behaviour [5-111. In this study a Yb-sialon glass has been prepared with the composition in equivalent % of 35e/oYb:45e/oSi:2Oe/oA1:83e/oO: 17eIoN. The glass was devitrified for 10 hours at 1200°C under nitrogen. The glass and the glass-ceramic properties were determined. DTA-studies have been carried out under air and water vapour to assess the oxidation resistance of the glass and the glass-ceramic.
Dilatometer measurements were carried out at a heating rate of 5"CImin to determine the glass transition temperature, the dilotemetric softening point and the thermal expansion coefficient. The Young's modulus of the glass and the glass-ceramic has been determined via an ultrasonic technique described previously [ 121. The resulting glass-ceramic phases were determined with a Philips X-ray powder diffractometer (Cu-Ka radiation). Combined DTA/TGA was carried out, using a Stanton Redcroft STA-780 series simultaneous thermogravimetric differential thermal analyser, on 90 mg cubes of the glass and the glass-ceramic samples. The samples were placed in a platinum crucible. Precalcined B.D.H. grade alumina powder was used as reference material. The samples were heated at 10"CImin up to 1200°C. In a second experiment a hold of 16 hours was followed by heating. Air and air bubbled through water were used as oxidation gases. A gas flow of 90 mYmin was employed in each experiment.
RESULTS In table 1 and 2 the thermal, physical and mechanical properties of the obtained Yb-Si-Al-0-N glass are given. The glass presents a moderately high Tg value of around 900"C, and good mechanical properties.
66 1
101
Table 1: Thermal properties of the Yb-Si-Al-0-N glass
15
100.8
-
h
ap 100.6 E
10
o, 100.4
5
c;;
i!? 100.2 0
100
-5
99.8 650
850
1050
1250
Temperature ("C)
Table 2: Mechanical glass and glass-ceramic properties Fig. 2: DTA of the glass under air at 10"C/min
In table 2 the mechanical glass-ceramic properties are compared to the corresponding glass properties. No changes in density during transformation from the glass to the glass-ceramic are observed. An increase of Young's modulus and fracture toughness with crystallisation is achieved while retaining a good hardness value.
Figure 3 shows the DTA/TGA traces for the glass-ceramic oxidised under air and under water vapour. The DTA curve for the experiment run under air + water vapour presents a lower signal to background ratio due to the sudden release of the air bubbles through the water.
In figure 1 the resulting glass-ceramic structure after crystallisation at 1200°C is presented. A fine grained, multiphase microstructure is obtained. The main phase at this temperature, identified by XRD,is Yb3Al5OI2(YbAG, JCPDS 23-1476) accompanied by pYb2Si207 (JCPDS 37-0458), X2-Yb2Si207(JCPDS 361476) and a small amount of Yb-B-phase (Yb2SiAI05N).
The glass-ceramic is almost fully crystallised after heat-treatment at 1200"C, as no further glass transition is visible in the DTA traces (fig. 3). In both experiments (air and water vapour) only a small weight gain is observed (fig. 3), which is consistent with the presence of only a small white oxidation layer, as observed at the surface of the samples (fig. 4). A hold of 16 hours at 1200°C did not result in additional oxidation (fig. 3).
Fig. 1: SEM image of the glass-ceramic
97
! 0
Figure 2 shows the DTA/TGA traces of the oxidised glass. The glass starts to oxidise at 1090°C with the oxidation product P-Yb2Si207 appearing. The maximum of the crystallisation peak is observed at 1127°C. A second phase crystallises at around 1200"C, which was identified as YbAG.
662
f -10 ! -15 250
500
750
1000
1250
Temperature ("C)
Fig. 3: DTA of the glass-ceramic under air and water vapour at 1O"C/min
Fig. 7: TEM image of the oxidised surface layer
Fig. 4: Image of oxidised sample
Different types of crystals are observed in the samples: (1) large, elongated (fig. 5 and 6) and (2) small rounded (fig. 7). Two types of cracks are observed: (1) elongated (fig. 5), which appear during heating and (2) curved (fig. 6), due to stress relaxation between two neighboured grains. No cracks are observed between the small rounded crystals (fig. 7).
CONCLUSION Oxidation experiments under a flowing air atmosphere have been carried out for the glass and for glass-ceramic materials prepared at 1200°C. Up to 900°C no weight changes were observed for both samples. For the glass, a weight gain of 1% was observed on heating under air during heating to 1200°C. Assuming that all nitrogen is replaced by oxygen during the experiment (weight gain of 2%), this value lies well within experimental error. For the glass-ceramic, no weight changes are observed up to 1200°C. Due to its good mechanical properties and its stability under air, this glass-ceramic would thus appear to have useful potential as a grain-boundary phase in silicon nitride based materials.
ACKNOWLEDGEMENT
Fig. 5: TEM image of the oxidised surface layer
This project was funded by the EU as part of a TMR network Project (Contract No. FMRX-CT960038). We are grateful to our colleagues Prof. D.P.Thompson, University of Newcastle upon Tyne, Prof. P. Goursat, University of Limoges, Dr. L. Falk, Chalmers University of Technology, Gothenburg, Prof. J.-L. Besson ENSCI, Limoges, Prof. R. Harris, University of Durham, Prof. J.P. Descamps Ecole Polytechnique de Mons and Dr. F. Cambier, Belgium Ceramic Research Centre, , University of Limoges for useful discussions.
REFERENCES [l] S. Hampshire, “Nitride Ceramics”, pp 119-171 in
Fig. 6: TEM image of the oxidised surface layer
Materials Science and Technology, ed. R.W. Cahn, P. Haasen, E.J. Kramer, Structure and Properties of Ceramics, Weinheim, Germany, 1994 [2] S. Hampshire, K.H. Jack, ‘“The Kinetics of Densification and Phase Transformation of Nitrogen Ceramics”, Special Ceramics 7, eds. D.E. Taylor & P. Popper, Brit. Ceram. Proc., 31,37-49, 1981
663
[3] D.R. Clarke D.R., Thomas G., “Microstructure of Y2O3 Fluxed Hot Pressed Silicon Nitride”, J. Am. Ceram. Soc.,61[3-41, 114-118, 1978 [4] C.C. Ahn, G. Thomas, “Microstructure and Grain Boundary Chemistry of Hot Pressed Silicon Nitride with Yttria and Alumina”, J. Am. Ceram. SOC.,66 [l], 14-19, 1983 [5] M. Cinibulk, G. Thomas and S . Johnson, “GrainBoundary Crystallisation and Strength of Silicon Nitride Sintered with Y-Sialon Glass’’, J. Am. Ceram. SOC.,61 [6], 1606-1612, 1990 [6] F.F. Lange, “Phase Relations in the System Si3N4Si02-MgO and their Interrelation with Strength and Oxidation”, J. Am. Ceram. SOC., 61 [l-21, 53-56, 1978 [7] G.E. Gazza, “Effect of Yttria Additions on Hot Pressed Si3Ni’. Am. Ceram. SOC.Bul., 54 [9}, 778781,1975 [8] D.R. Clarke, F.F. Lange and G.D. Schnittgrund, “Strengthening of Sintered Silicon Nitride by Post Fabrication Heat Treatment”, J. Am. Ceram. SOC.,65 [4], C51-C54, 1982 [9] A.W.J.M. Rae, D.P. Thompson and K.H. Jack, “The Structure of Yttrium Silicon Oxynitride and its Role in the Hot Pressing of Silicon Nitride with Yttria Additions”, Special Ceramics 6, ed. P. Popper, British Ceramic Society, Stoke en Trent, UK, 347360,1975 [ 101M. Chadwick, R.S. Jupp, D.S. Wilkinson, “Creep behaviour of a sintered Silicon Nitride”, J. Am. Ceram. SOC., 76 [2], 385-396, 1993 [ 111R. Ramesh, E. Nestor, M.J. Pomeroy, S. Hampshire, “Classical and Differential Thermal Analysis Studies of the Glass-Ceramic Transformation in a Y-Sialon Glass”, J. Am. Ceram. SOC.,81 141, 1-13, 1998 [ 121Standard practice for measuring ultrasonic velocity in materials, E494-92 (1992) 177-187 ’
664
Index
abrasive wear 161 accuracy 119 acoustic sensor 182 additives 425,435,571 adherence 74 adhesion 57,61 aerospace components 464 affordable materials 3 , 6 alloying additions 61 1 alumina 73, 163,239,247,279, 347,405, 483,587 -gels 637 -parts 387 aluminium nitride 347,571,593 aluminium oxide 37 1 aluminium hydroxide gel 36.5 anode 8,33,51ff aqueous tape casting 5.5 articulated configuration 223 articulating fixtures 223ff, 228 aspect ratio 123 autoclave 15 automotive application 383, 393 batteries 33 battery safety 34, 36 bending strength 391 bending tests 197 bimodal microstructure 577 binder 86, 88 biological 505 blade design 3 I 1 blank 399 bondcoat roughness 309 boron nitride 70 brake disks 569 brake technology 13 Brazilian disc test 211, 215 brazing 347 -technology 347 burnerrig 156 -tests 154
C/C-Sic 64 - composites 63,67 calcination 34 carbon fibers 477 carbon fibre-reinforced plastics 4
carbothermal reduction 577 casting 365 cathode 8 , 5 1 -materials 33 cell technology 7 ceramic nozzles 97 ceramic-to-metaljoining 350 cermets 85, 88 characteristic strengths 225 chemical stability 21f clay bonded 74,77 - silicon carbide 77 --filter 74f CMC 581,627,641 -brakes 63 coal fired power stations 255 coarse foamstructures 537 coating 477 coefficient of friction 63,65f, 67 coefficient of thermal expansion 64 cofiring 51ff co-generation technology 45,49 combuster liners 266 combustion 31, 199,203 combustor 47 complex machining processes 410 complex shapes 353,359 composite 431,611 - component structure 46 - component technology 46 -materials 3 compositions 61 1 compounding 63 1 compressive residual stress 157ff, 161 compressive stress 167,21lff, 239 computer aided design 255 computer simulations 267 conduction mechanism 495 consistency 119 contact 193 -fatigue 193 - pressure 205ff core-rim 85f - structure 86, 89 corrosion 69,153f, 156,199f, 205 -resistance 13, 17,260 cost 6,549 - effective 14f - - manufacturing 4,383
- - material 397 - - production 399 -efficient 65 - - manufacturing 67 -savings 10 crack aspect ratio 123ff crack free joints 391 crack geometry 122, 127 crack growth 109ff, 122, 124f, 129ff, 139, 143f, 151, 169ff, 175ff, 179,309, 319 crack initiation 169ff, 174 crack interlinking 163f, 167 crack propagation 175,233,523 -model 305 crack shape 121ff cracking 193, 197 cracks 17.5 creep 109ff, 139 - deformation 221 -models 297 -resistance 74f, 77,217,221f -rupture 109, 111, 113 -strain 297 -testing 291 critical stress 174 cross-peripheral grinding 4 17 crystallization 439, 657 -velocity 61 curing 653 cutting edge 377,411 cutting methods 627 cutting tools 21, 26,411 CVD process 9 1,93 cyclic contact loading 161 cyclic fatigue 279 -test 182
damage mechanism 285,319 damage mode 193, 197 damping 217,219ff, 222 -peak 221 density 4 design methodology 543 design standards 273 devitrification 661 dielectric ceramics 229 dielectric strength 230 diesel particulate filter 27
665
diesel version 35 diffusion barrier 41 disc radius 213f dispersion stability 605 dome 267 dood linearity 72 durability 69ff dynamic fatigue 99f, 109, 139ff -tests 141
EB-PVD process 517 elastic 217,222 electric 35 -vehicles 33 electrical measurements 175 electrical properties 175 electro-ceramic 212,214 - components 325 electrode material 33 electrolyte 8, 33ff, 51ff -material 33 - vapor deposition 8 electromotive force 69 electron beam 285 emission control 39 engine blocks 199 - materials 199 engine valves 181, 399 environment 27, 199 equilibrium thickness model 453 erosion resistance 70 evaporator 53 1 excellent 69 exhaust gas 27f, 31f, 41 exhaust pollution 27 external conductive path 8 extrusion 637 failure mechanisms 305 failure mode 21 1 fatigue 193 - behaviour 247 -crack 233 - damage 289,621 -design 247 - lifetimes 279 -limit 283 fiber filter 28 fiber-reinforcedceramic composite 5, 26 1 fibres 653 fibrous preform 15 filter 79, 83f -ceramic 80 - elements 73,76 -materials 73 - membranes 83f filtration 79ff - efficiency 74 -temperatures 73 fine finishing process 375 finite element method 21 lf, 305, 31 1, 53 1 finite Element System (FE) 339 fixing 8
666
flexural strength 115, l69,223,227f, 283,627 flexural test 174 flow pattern 267 flow properties 63 1 fracture 24,34, 193 -mechanics l09f, 117 - mechanism 113,543 -mode 167 -resistance 126 -toughness 70, 103ff, 117, 121f, 139ff, 149ff, 217,222 friction 104, 169ff, 187f, l91,199f, 202f, 205ff fuel cell 605 fuel consumption 39 function 208 -groups 40 functional cavities 537 gas pressure sintering .85,90. 97 gas sensing device 9 1 gas sensitivity 92 gas sensor 91f gas tight joining 61 gas tightness 58,60 gas tight seal 8 gas turbine 45,48f, 97, 100,261,311, 319 - combustor 53 1 gelatin casting 587 gelcasting 463 geometry factor 121ff, 124f glass 661 - ceramic sealants 57 -ceramics 57 glass phase 147 glass sealants 57 glass types 57,61 grain boundary l43,163ff, 167,439,453 -chemistry 139 -composition 139, 142, 155 -phase 142f, 154, 165, 167, 179,663 grain pull out 167, 193, 197 grain size 495 green density 52,54f green machining 37 1 grinding kinematics 42 1 grinding parameters 393 grinding process 399,411 G-value 182 hard turning 4 11 hardness 199ff, 205 heat exchanger 255,571 heat treated 91 heat treatment 92 highcreep 4 high-speed precision turning 410 high-temperatureoxidation 139 Hi-NicalonTM fibers 233 honeycomb structures 537 hot gas filter element 76 hot gas filtration 75,79 hot gas test 153
hot hardness 21 hot press techniques 67 hot pressing 85,641 hot-pressed silicon nitride 291 hybrid electric vehicles 33 hybrid gas turbines 49 hybridtype 47 hybrid vehicles 33 hydrolysis assisted solidification 365 impact damage 3 11 impedance spectroscopy 175ff impregnation 14 impulse excitation 222 - technique 217 -test 217 impurities 447 in situ synthesis 483 inclusions 553 indentation crack 121ff indentation fracture toughness 147 indentation load 121ff indentation strength 121ff indentation-quench 127f, 131ff, 147, 151 infiltration 14 injection moulding 631 inserts 21,26 insulating atmospheres 641 interacting mechanisms 108 interconnection 8 intergranular 147 - glass phase 147f -phase 167,22lf, 456 interlayer mechanisms 617 internal friction 217ff, 222 Japanese Industrial Standard 278 joining 617,627 - mechanism 387 -technology 347 kinetic damage 289 La,O, 593 lambda sensors 39 laminated ceramics 27 laminating seats 27 lamination 5 1,55 laser welding 387 laser-assisted turning 377 layer 79 layer structure ceramics 71 layered materials 559 lean bum 39,41ff -engines 44 lifetime 109ff, 157, 161, 175, 182, 193, 41 1 -model 305 - prediction 285,319,325 light metal components 383 lightweight 63 -brake systems 67 -design 4 - engineering 4 limit 621
linearity 69 liquid phase 85f, 88 - sintering 453, 593 load 187ff loading analysis 325 long-term reliability 487 low-cost cell production 11 low-cost process technology 543 lubrication 199,205 machanical properties anisotropic electrical conductivity 559 machine system 399 machining 25, 181f, 193,417 - behaviour 371 -damage 193 -techniques 377 manufacturing 15 -technologies 6 martensitic transformation 222 material removal 417 materials 23, 108, 153,212,537 -design 431 mathematical model of microhardness 299 matrix composite 13,63,233 mechanical behaviour 103, 199,553 mechanical failure 325 mechanical properties lSff, 74, 85, 88ff, 115,201,471,647 mechanical resistance 40 mechanical stability 40 mechanical strength 70 mechanisms 487 melt spinning 653 melting behavior 58ff melting point 69ff membrane 79ff mercury porosimetry 599 metalbased 85 metal-ceramic composites 85 metal cutting 21 metal forming 133, 137 metal-glass composite 513 metal joints 333 metal machining 85, 89 metal matrix composites 347,383 metallic binders 89 metallic phases (i.e. binder) 85 MgO-ZrO, 239 micro gas turbine 359 microcrack initiation 167 microcracking 193, 195ff microfracture 163,208,210 microgeometrical wear mechanisms 416 microhardness 299 micro-plastic deformation 157, 161 microstructural assessment 115 microstructural design 103, 108 microstructural parameters 105 microstructural stability 153, 156 microstructure 53, 85, 97ff, 104, 106, 118, 148, 154, 159, 163, 165, 189, 193, 201,206,439,443,457,593,605 microwave 457
miniaturization 359 mm-wave 457 modified compounding techniques 15 Mold Shape Deposition manufacturing 359 monolithic 71 mullite 483 multiaxial strength 239 multiaxial stress 291 multilayer 79,81,92f - filter ceramic 81 -structure 84 nanoceramics 459 nano-composites 657 near net shape 13,393 near-surface characteristics 157 net 76 nitridation 447 non-destructive 217 - evaluation 182 nonoxides 153, 156 novel concept 15 NOx emission 45,47,49 nozzles 97 numerical field simulations 175 optical properties 57 1 overcharge protection 34 overdischarge protection 34 oxid based candles 77 oxidation 153ff, 187,205,227 -resistance 77,661 oxidative crack healing 227f oxide 469 - based filter elements 75 - concentration 639 oxygen partial pressure 42 oxygen sensors 39ff oxynitride 657 particle removal 77 paste flow 267 pastes 353 performance 67 - applications 6 phase formation 593 phase transformation 175ff, 193,435 phonon scattering 496 physical properties 457 piezo-active structures 6 planar solid oxide fuel cells 57 planar technology 40 plasma spraying 187,305 plasma-sprayed Cr,O, coating 299 plasma-sprayed stripe 8 plastic deformation 193, 196f, 208,210 polycarbosilazanes 653 polycrystalline aumina 163, 167 polymer processing techniques 58 1 polymerisation 477 polysilazanes 581 porous 383,537,587 -preform 15 - silicon carbide 531
powder characteristics 425 power modules 499 power time dependence 333 preceramic polymer 617,627 precision 70f - machining 405 precursor 477,653 preparation 47 1 pressed-mould 267 pressure forming process 15 pressureless sintering 657 prevent gas leakage 8 prevention 70 process chain 399 process design 399 processing 61 1 propagation 167 - of a crack 176 prototype 353, 377 -testing 255 pseudoplastic rheological behavior 52 PTC-switching device 325 purpose-designed 115 pyrolysis 477,653 rapid prototyping 363,463 rareearth 435 rare-earth additive 499 RBAO (Reaction Bonding of Aluminium Oxide) 549 R-curve 121ff reaction 505 -bonded silicon nitride 393 -bonding 13ff, 75,565 - sintered silicon nitride 70 reactive metal penetration 483 recleaning 79ff, 84 reliability 6,273,621 reproducibility 119 residual crack length 283 residual stress 132, 157ff, 161, 163, 165, 199ff, 416 -factor 130 -parameter 121 residual tensile stress 97, 100 resin transfer moulding 15 resistive sensors 42 resonance vibrations 3 11 response 69ff rheological behaviour 356 rig test 153 rotor 359 roughing process 375 round robin 223,227f RSiC candles 77 RSiC-filter elements 75 safety concept 247 scratch 193 -damage 193, 195ff -testing 193, 195 sealed structure 70 secondaq phase 97ff, 577 selectivity 93 self bonding 75
667
selfmated 187 self-mated plasmasprayed 191 semiconductingmetal oxides 39 semiconductingoxides 39 sensitivity 91ff -analysis 175, 178 sensor 69,72 shaping process 631 sheathed thermocouple 69 Sheet Moulding Compound (SMC) Technology 16 short circuit 34 shrinkage 54f shrinkage rate 54 Si,N,/(fi-SiAION+TiN)composites 559 Si,NJSiC nano/micro Composite 553 sialon 147, 151,435,447 Sic composite 233 SiC/SiC, composites 617 SiC,N,-BN comosite 647 Si-C-N precursors 581 Sic-SiO, composite 453 silicon 109 -carbide 73, 139,347,453,493,505, 617,641,657 - - filter elements 74 - gas infiltration 505 -nitride 70ff, 109, 112, 114, 122, 181, 193,205ff, 2 10,217,219ff, 223,228, 247,261,311,359,359,365,377,425, 435,439,447,463,487,495,499,577, 621,641 --blades 265 - - nozzles 100 - - nozzles 97,265 sintered silicon carbide 76 sintering 53ff, 70, Slf, 84f, 443,457, 517, SiSiC components 260 sliding speed 187ff, 205f, 208ff sliding velocities 187 slip casting 15, 599, 605 slurry forming 365 SMC (Sheet Moulding Compound) 17, 565 SOFC 51,55 Solar Centaur-5OS engine 261 solid electrolyte 605 solid oxide fuel cell 7 , 5 1 spallation 523 specimen-loadingconfigurations 297 spray-dryed 34 stable crack 122, 126 starch consolidation 587 static fatigue 109, 111,279 stereolithography 353 stiffness 217,220f stiffness 222 storage of electricity 33 strength 22,24,26,77, 103, 115, 139ff, 193, 195, 197,20lf, 205,225,621 stress intensity factor K 121 stress pattern 132 stresses 164 stress-free elastic supporting structure 46
668
stress-free independent supporting structure 46 stress-lifetime curves 279 structural design 273 structural reliability 109 structure 85 sub-surface characteristics 421 surface finishing 229 surface flaws 229 surface properties 425 surface quality 421 surface tension 61 suspension 353,365 - impregnation 641 system 133, 135 tangential traction 171, 173f tape casting 40,51ff technologies 3 TEM 439,443 temperature 199ff, 203, 217ff - fracture strength 469 -resistant 45 - quenching 151 tensile creep 3 19,487 tensile strain 523 tensile strength 21 If, 239 tensile stress 21 Iff test method 116 testing 217 testing procedures 115 thermal barrier coating 285,305,513, 517,523 thermal conductivity 495,499 thermal efficiency 27,29,32,45 thermal expansion 57f, 58, 167 - coefficient 8,22,46,57,59,61,70, 165 - mismatch 166 - - stresses 167 thermal properties 83 thermal shock 22,24,26,48,69,71, 119. 127ff, 147f, 187,191,199,260 - behaviour 147 - properties 147, 151 -resistance 40,63,70,85, 127ff, 132, 147, 150f thermal spraying 187f, 199,299 thermal stability 40,63 thermal stress 27, 127ff thermally grown oxide 305 thermally sprayed 199,201 - coating 199,203 thermally stable 13 thermo-compression 5 1,55 thermocouple 69ff thermoelectro motive force 72 -response 70 thermomechanicalbehavior 14 thermomechanicalproperties 97 thermophysical material 201 thermophysical properties 17, 199f, 203, 649 thick film 43 -sensor 43
thin film 93 - gas sensing 91 - gas sensor 91
titania ceramics 229 tool 133, 135, 137 -ceramics 405 -life 133, 137 -wear 133,375 toughness 22,26,85, 115, 193 transformation toughening 469 transportation 6 transverse thermal conductivity 65, 67 tribological 64, 103, 157, 199ff, 210 - applications 63 -behaviour 65,67, 104, 136 -characteristics 107, 133,206 - conditions 133 -performance 103,106,108 - properties 187,205 -system 133,135, 137, 157 -tests 65f turbine blade cooling technologies 45 turbine components 3 I 1 turbine inlet temperature 45 turbine nozzle 48 turbine rotors 3 11 turbo compound engine 29ff turning 25 TZP 175, 178,180 ultrasonic analysis 182 ultrasonic assisted face grinding 4 17 ultrasonic welding process 333 universal oxygen sensor 41 valve 181ff, 247 varistor 325 viscosity 57, 59, 61,477, 600 voltage screening 229
wear 21ff, 103ff, 119, 133ff, 157ff, 163f, 167,169, 171, 173f, 187ff, 193,197, 199f, 202f, 205ff -damage 193 -mechanism 163,417 - properties 193 -rates 63 -resistance 22,26,63,64, 159, 163, 167 -stability 67 Weibull characteristic strengths 225f Weibull function 339 Weibull modulus 24,53,224ff Weibull statistics 223,228, 319, 333, 339 wetting angle 58ff Y,O, 499,593 Yb-sialon 661 Y-Si-A1-0-N glass 657 Y-TZP 621
zirconia 193, 217, 517,523 -SOIS 637 ZrO, planar technology 39,44
