Materials Aspects in Automotive Catalytic Converters

Material Aspects in Automotive Catalytic Converters Edited by Hans Bode

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

Further Titles of Interest: B. Cornils, W. A. Herrmann, R. Schlägl, C.-H. Wong (Eds.) Catalysis from A-Z ISBN 3-527-29855-X S. Hagen, S. Hawkins Industrial Catalysis ISBN 3-527-29528-3 S. M. Thomas, W. J. Thomas Principles and Practice of Heterogenous Catalysis ISBN 0-471-29239-X G. Ertl, H. Knözinger, S. Weitkamp Handbook of Heterogenous Catalysis ISBN 0-471-29212-8

Material Aspects in Automotive Catalytic Converters

Edited by Hans Bode

Deutsche Gesellschaft für Materialkunde e.V.

Prof. Dr. Ing. Hans Bode Bergische Universität GH Wuppertal FG Werkstofftechnik Gaußstr. 20 D-42097 Wuppertal Germany

International Congress „Material Aspects in Automotive Catalytic Converters“, held from 03–04 October 2001 in Munich, Germany Organizer: DGM · Deutsche Gesellschaft für Materialkunde e.V.

This book was carefully produced. Nevertheless, authors, editor 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.

Library of Congress Card No. applied for. A catalogue record for this book is available from the British Library Deutsche Bibliothek Cataloguing-in-Publication Data: A catalogue record for this book is available from Die Deutsche Bibliothek ISBN 3-527-30491-6 © WILEY-VCH Verlag GmbH, Weinheim (Federal Republic of Germany), 2002 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. Satz: W.G.V. Verlagsdienstleistungen GmbH, Weinheim Druck: betz-druck GmbH, Darmstadt Bindung: J. Schäffer GmbH + Co. KG, Grünstadt Printed in the Federal Republic of Germany

Preface Based on increased ecological demands, the car and car-supplying industries strive to meet the challenging requirements for higher performance and extended service life of future vehicle generations. Maintaining good performance is mandatory particularly in the view of thinner supports, higher cell densities and higher temperatures. Performance and service life predictions, based on tests or on modelling and simulation techniques, will depend on reliable materials data. Only very close cooperation between researchers and producers will help to meet these requirements. It was the aim of MACC, the second international conference on Materials Aspects in Automotive Catalytic Converters, to foster this cooperation. It refered to papers from both industry and research institutes which concentrate on the high-temperature mechanical and oxidation behaviour of both metal-supported and ceramic supported automotive catalysts. The metalsupported catalyst is based on a ferritic steel with 5–8% aluminum, 17–22% chromium and small additions of reactive elements. More than 11,000,000 units were produced in the year 2000. The ceramic-supported catalytic converter is based on corderite. The production rate is much higher. Both materials have specific advantages and disadvantages which determine the application for a given car model. In addition to these two basic groups of catalytic carriers, coating and canning aspects were also addressed by the conference programme. Especially the influence of coating thickness and composition is becoming more and more important when going to thinner supports and higher cell densities. I am very obliged to the authors for their valuable contribution to a comprehensive programme that covers the whole chain of product development and application, beginning with the melting process and ending with recycling aspects. Munich, October 2001 Prof. Dr.-Ing. Hans Bode Conference Chairman

I Introduction Contribution of Automotive Catalytic Converters R. Searles, Association for Emissions Control by Catalyst, Brussels (B) ...................................3

II Metals Development Status of Metal Substrate Catalysts R. Brück, Emitec GmbH, Lohmar .............................................................................................19 Materials Issues Relevant to the Development of Future Metal Foil Automotive Catalytic Converters J. Nicholls, Cranfield University, Cranfield (GB); W. Quadakkers, Forschungszentrum Juelich (D) ................................................................................................................................31 High Temperature Corrosion of FeCrAlY / Aluchrom YHf in Environments Relevant to Exhaust Gas Systems A. Kolb-Telieps, Krupp VDM GmbH, Altena (D); R. Newton, Cranfield University, SIMS, Cranfield (GB); G. Strehl, TU Clausthal, Institut für Allgemeine Metallurgie Clausthal-Zellerfeld (D); D. Naumenko, W. Quadakkers, Forschungszentrum Jülich, IWV-2, Jülich (D) .....................................................................................................................49 Improved High Temperature Oxidation Resistance of REM Added Fe-20%Cr-5%Al Alloy by Pre-Annealing Treatment K. Fukuda, K. Takao, T. Hoshi, O. Furukimi, Technical Research Laboratories, Kawasaki Steel Corp., Chiba (Japan) ......................................................................................59 Oxidation Induced Length Change of Thin Gauge Fe-Cr-Al Alloys C. Chang, L. Chen, B. Jha, Engineered Materials Solutions, Inc., Attleboro (USA) ................69 Improvement in the Oxidation Resistance of Al-deposited Fe-Cr-Al Foil by Pre-oxidation S. Taniguchi, T. Shibata, Department of Materials Science and Processing, Graduate School of Engineering, Osaka University, Osaka (J); A. Andoh, Steel and Technology Development Laboratories, Nisshin Steel, Osaka (J) ...........................................83 Factors Affecting Oxide Growth Rates and Lifetime of FeCrAl Alloys W. Quadakkers, L. Singheiser, D. Naumenko, Forschungszentrum Jülich (D); J. Nicholls, J. Wilber, Cranfield University, School of Industrial and Manufacturing Science, Cranfield (GB) ...........................................................................................................93 On Deviations from Parabolic Growth Kinetics in High Temperature Oxidation G. Borchardt, G. Strehl, Institut für Metallurgie, TU Clausthal (D) ......................................106

VIII Effect of Reactive Elements and of Increased Aluminum Contents on the Oxide Scale Formation on Fe-Cr-Al Alloys V. Kolarik, M. del Mar Juez-Lorenzo, H. Fietzek, Fraunhofer-Institut für Chemische Technologie, Pfinztal (D); A. Kolb-Telieps, H. Hattendorf, R. Hojda, Krupp VDM GmbH, Werdohl (D) .....................................................................................................117 High Temperature Strength of Metal Foil Materials M. Cedergren, K. Göransson, R&D, AB Sandvik Steel, Sandviken (S) ..................................126 Lifetime Predictions of Uncoated Metal-Supported Catalysts via Modeling and Simulation, based on Reliable Material Data H. Bode, University of Wuppertal (D); C. Guist, BMW AG, Munich (D) ..............................134 Elastic-Plastic Thermal Stress Analysis for Metal Substrates for Catalytic Converters S. Konya, A. Kikuchi, Nippon Steel Corporation, Futtsu (J) ..................................................144 A New Type of Metallic Substrate R. Lylykangas, H. Tuomola, Kemira Metalkat Oy (SF) ..........................................................152

III Ceramics Development Status of Ceramic Supported Catalyst C. Vogt, E. Ohara, NGK Europe GmbH; M. Makino, NGK Insulators Ltd ...........................173 Evaluation of In-Service Properties and Life Time of Automotive Catalyst Support Materials U. Tröger, M. Lang, Zeuna Stärker GmbH & Co. KG, Augsburg (D) ...................................186 Loads, Design and Durability Evaluation of Mount Systems for Ceramic Monoliths G. Wirth, J. Eberspächer GmbH & Co., Esslingen (D) ..........................................................191 High Performance Packaging Materials M. Vermoehlen, D. Merry, S. Schmid, Corning GmbH, Wiesbaden (D) ................................202

IV Catalysts Three-way Catalyst Deactivation Associated With Oil-Derived Poisons J. Kubsh, Engelhard Corporation, Environmental Technologies Group Iselin (USA) ..........217 Catalytic Reduction of NOx in Oxygen-rich Gas Streams, Deactivation of NOx StorageRaduction Catalysts by Sulfur C. Sedlmair, K. Sehan, Technische Universität München, Institut für Technische Chemie II, Garching (D); J. Lercher, A. Jentys, University of Twente, Faculty of Chemical Technology, Enschede (NL) ...................................................................................223

IX Catalytic Reduction of NOx in Oxygen-rich Gas Streams: Progress and Challenges in Catalyst Development W. Grünert, Lehrstuhl Technische Chemie, Ruhr-Universität Bochum (D) ...........................229 Atomic Structure of Low-Index CeO2 Surfaces H. Nörenberg, University of Oxford, Department of Materials, Oxford (GB); J. Harding, University College London, Department of Physics and Astronomy, London (GB); S. Parker, University of Bath, Department of Chemistry, Bath (GB) .....................................237 Nanostructured Ceria-Zirconia as an Oxygen Storage Component in 3-way Catalytic Converters-Thermal Stability B. Djuricic, Austrian Research Centers, Seibersdorf (A), S. Pickering, Institute for Advanced Materials, Petten (NL) ...........................................................................................241

V Recycling Recycling Technology for Metallic Substrates: a Closed Cycle C. Hensel, Demet Deutsche Edelmetall Recycling AG & Co. KG, Alzenau (D) ....................251

VI Miscelleanous Hot-Corrosion of Metal and Ceramic Honeycombs by Alkaline Metals for NOx Adsorption M. Yamanaka, Nippon Steel Technoresearch, Futtsu (J); Y. Okazaki, Nippon Steel, Toukai, (J) ..............................................................................................................................263 The Effect of Trace Amounts of Mg in FeCrAl Alloys on the Microstructure of the Protective Alumina Surface Scales P. Untoro, M. Dani, National Nuclear Energy Agency, Kawasan PUSPIPTEK, Serpong (Ind); H. Klaar, J. Mayer, Gemeinschaftslabor für Elektronenmikroskopie, RWTH Aachen (D); D. Naumenko, J. Kuo, W. Quadakkers, Institut für Werkstoffe und Verfahren der Energietechnik (IWV-2), Forschungszentrum Jülich (D) ................................271

Subject Index* A Adsorption, NOx 263 Al deposition 83 Alkaline metals 263 Alloys 271 Aluchrom YHf 49 Alumina surface 271 Aluminum content 117 Annealing 59 Atomic structure 237 Automotive catalysts 3, 31, 186 C Catalyst 134 - ceramic 173 - three-way 217, 241 - deactivation 217 - development 19, 229 - support materials 186 Catalytic converters 3, 31, 144, 241 Catalytic reduction 223, 229 CeO2 surfaces 237 Ceramic components 241 Ceramic honeycomb 263 Ceramic monoliths 191 Ceramic supported catalyst 173 Ceria-Zirconia 241 Corrosion 49, 263

Elements, reactive 117 Emission limits 152, 202, 251 Environmental protection 251 Exhaust gas systems 49, 152 F Fe-Cr-Al alloy 59, 69, 83, 93, 117, 271 FeCrAlY/Aluchrom YHf 49 Foil 31, 83, 126 G Gas flow 152 Gas streams 223, 229 Gas systems 49 Growth kinetics 106 H High temperature - corrosion 49 - oxidation 59, 106 - strength 126 Honeycomb 263 Hot-corrosion 263 I Increased aluminum content 117 In-service properties 186 K

D

Kinetics 106

Data, reliable 134 Deactivation 217, 223 Development 173 - catalysts 19, 229 - converters 31 Durability evaluation 191

L Length change 69 Lifetime 93, 186, 251 Lifetime predictions 134 Loads 191 Low-index CeO2 surfaces 237

E Elastic thermal stress 144 *

The page numbers refer to the first page of the article

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

281 M Mat, vermiculite 202 Material - data, reliable 134 - issues 31 - packaging 202 - strength 126 Metal foil 31, 126 Metal honeycomb 263 Metallic substrate 152, 251 Metals, alkaline 263 Metal substrate catalysts 19, 144 Metal-supported catalysts 134 Mg 271 Microstructure 271 Mixed gas flow 152 Modeling 134 Monoliths, ceramic 191 Mount systems 191 N Nanostructure 241 NOx - adsorption 263 - reduction 223, 229 - storage 223 O Oil-derived poisons 217 Oxidation 69, 106 Oxidation resistance 59, 83 Oxide growth rates 93 Oxide scale formation 117 Oxygen-rich gas streams 223, 229 Oxygen storage 241 P Packaging materials 202

Parabolic growth kinetics 106 Plastic thermal stress 144 Poisons, oil-derived 217 Pre-annealing treatment 59 Pre-oxidation 83 Protective surface 271 R Reactive elements 117 Recycling technology 251 Reliable material data 134 REM 59 S Separation process 251 Simulation 134 Storage-reduction catalysts 223 Stress analysis 144 Structure, atomic 237 Substrate 19, 144, 152 Sulfur, NOx reduction 223 Surfaces 237, 271 T Thermal stability 241 Thermal stress analysis 144 Thin alloys 69 Three-way catalyst 217, 241 Trace amounts, alkaline 271 U Uncoated metal-supported catalysts 134 V Vermiculite mat 202

I Introduction

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

Contribution of Automotive Catalytic Converters Robert A Searles Association for Emissions Control by Catalyst, Brussels, Belgium

1

Abstract

Catalyst-equipped cars were first introduced in the USA in 1974 but only appeared on European roads from 1985. In 1993 the European Union set new car emission standards that effectively mandated the installation of emission control catalysts on gasoline-fuelled cars. Now more than 300 million of the world’s over 500 million cars and over 85% of all new cars produced worldwide are equipped with autocatalysts. Catalytic converters are also increasingly fitted on heavy-duty vehicles, motorcycles and off-road engines and vehicles. The paper will review the technologies available to meet the exhaust emission regulations for cars, light-duty and heavy-duty vehicles and motorcycles adopted by the European Union for implementation during the new century. This includes low light-off catalysts, more thermally durable catalysts, improved substrate technology, hydrocarbon adsorbers, electrically heated catalysts, DeNOx catalysts and adsorbers, selective catalytic reduction and diesel particulate traps. The challenge is to abate the remaining pollutants emitted while enabling fuel-efficient engine technologies to flourish. This is paramount to the achievement of air quality and greenhouse gas targets given the large increase in the number of vehicles on European roads since 1970 and the projections for further increases in vehicle numbers and greater distances driven each year in future.

2

Introduction

AECC is an international association of European companies making the technologies for automobile exhaust emissions control: autocatalysts, ceramic and metallic substrates, specialty materials incorporated into the catalytic converter and catalyst, adsorber and filter based systems for the control of gaseous and particulate emissions from diesel and other lean burn engines. 2.1

European Emission and Fuel Legislation

The European Union (EU) emission limits for passenger cars set from 1993 have already been lowered from 1996 and again from 2000. For passenger cars and light commercial vehicles the emission standards and fuel composition, including sulfur levels, have been agreed for 2000 and 2005. [1] New test cycles (ESC and ETC) and tougher emission standards for heavy-duty diesel vehicles have been finalized for 2000 and 2005. The limits for Enhanced Environmentally.Friendly Vehicles (EEV) are set and can serve as a basis for fiscal incentives by EU Member States. A

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KgaA ISBN: 3-527-30491-6

4 further reduction in limit values for nitrogen oxides (NOx) in 2008 is subject to a review by the European Commission in 2002 on technical progress. [2] The Working Party on Pollution and Energy (GRPE), an expert group of the World Forum for Harmonization of Vehicle Regulations (WP.29) at UN-ECE in Geneva, is developing a Worldwide Harmonized Heavy Duty Certification procedure and is looking in new measurement protocols in order to ensure that ultra fine particles are controlled by future emission legislation to minimize the health effects of diesel particle emissions. A proposal by the European Commission to set tougher, catalyst-requiring emission limits for motorcycles is being ratified by the European Parliament and Council. Tighter emission limits from 2003 for new types of motorcycles are agreed and correspond to a reduction of 60% for hydrocarbons and carbon monoxide for four-stroke motorcycles, and 70% for hydrocarbons and 30% for carbon monoxide for two-stroke motorcycles. A second stage with new mandatory emission limits for 2006 are expected to be based on the new World Motorcycle Test Cycle (WMTC) which is also being developed by the UN-ECE in Geneva. In the final report of the European Auto Oil II Programme [3], it was concluded that some air quality problems, such as atmospheric levels of particulate matter and ozone, are not yet solved. The challenge is to abate the remaining pollutants emitted while enabling the development of fuelefficient engine technologies. This is paramount to the achievement of air quality and greenhouse gas targets given the large increase in the number of vehicles on European roads and the projections for further increases in vehicle numbers and greater distances driven each year in future. 2.2

Exhaust Emissions from Internal Combustion Engines

Exhaust emissions can be lowered by: · Reducing engine-out emissions by improving the combustion process and fuel management, or by changes to the type of fuel or its composition · Retrofitting catalytic converters and associated engine and fuel management systems if they are not original equipment · Decreasing the time required for the catalytic converter to reach its full efficiency · Increasing the conversion efficiency of catalysts · Storing pollutants during the cold start for release when the catalyst is working · Using catalysts and adsorbers to destroy nitrogen oxides under lean operation · Using particulate filters with efficient regeneration technology · Increasing the operating life of autocatalysts and supporting systems. This paper reviews all the above opportunities, except the first, from the standpoint of material requirements and will also look back into the history of the materials developed for catalytic converters with these requirements in mind.

3

A Brief History of Automotive Catalysts

The first reference to a catalytic converter known to AECC is a patent [4] published to a French chemist, Michel Frenkel, in 1909. The device uses a kaolin (china clay) “honeycomb” with 30 grams of platinum as the active catalytic material. (Figure 1) The patent describes

5 “deodorizing” the exhaust using air blown in by a fan. As far as is known the device was not put into commercial production at the time. No doubt the high loadings of platinum were a deterrent and there was no air pollution concerns in those early, carefree days of motoring.

Figure 1: 1909 Catalytic Converter invented by Michel Frenkel

Figure 2: Eugene Houdry in 1953 with a small prototype catalytic converter

The next report [5] of the concept of catalytic converters was in the 1920s. Another European invention, this time German, was taken to General Motors in the US and was described as a collection of wires and beads, again coated with platinum. The tests were at first a success with the device glowing red, but within seconds the catalyst had failed. This was because tetraethyl lead had been recently introduced in the US as an octane booster, but was at that time unknown in Europe. Lead poisons catalytic converters. The French engineer Eugene Houdry can be considered as the father of the modern catalytic converter. Born in France he moved to the US and invented a revolutionary method for cracking low-grade crude to high-octane gasoline – the “cat cracking” process. After the 1939–1945 war he set up the Oxy-Catalyst company and turned his attention to the health risks from the increasing volumes of automobile and industrial exhausts. In 1962, the year of his death, he patented the first modern catalytic converter (Figure 2). The modern history of the catalytic converter started with the developments that lead to the 1970 US Clean Air Act and the rate of invention has accelerated greatly. Excellent histories of the industry [6, 7, 8] have been published so only a summary will be covered here. The modern catalytic converter, based on platinum group metals deposited on a ceramic honeycomb base or monolith, was first patented in 1965 [9]. However the industrial use of catalysts was then dominated by catalysts deposited on pellet or bead supports. In the first years after 1974 when catalytic converters were used in the US and Japanese markets, both pellet (Figure 3) and monolithic converters (Figure 4) were used. The loss of catalyst material by attrition in pellet converters was largely overcome by reactor design. Early prototypes of ceramic honeycombs were made by two approaches: 1. Dipping paper in ceramic slurries, corrugating them and laying up a unitary structure, firing the composite and shaping 2. Calendaring a plastic material containing ceramic powders between grooved and plain rollers, rolling up into a unitary structure, firing the composite and shaping.

6

Figure 3: Schematic of pelletized catalytic con- Figure 4: “Cutaway” ceramic monolith catalytic converter verter

Both of these developments were ultimately replaced by extruded honeycomb substrates. These are based on cordierite (2MgO .2Al2O3 .5SiO2) and made from natural raw materials and a plastic material that is extruded to form a unitary structure with parallel fine channels and then fired to the final shape. These materials have high thermal shock resistance and high melting and softening points with higher attrition resistance and lower pressure losses than pellet converters, which they ultimately replaced. In the 1970s new ferritic steels became available that could be made into ultra thin foils, corrugated and then laid up to form a honeycomb structure. One such steel was developed at the Atomic Energy Research Establishment in Harwell, UK for “canning” Uranium 235 and was called Fecralloy. This name reflects the components of the alloy - Iron (Fe), Chromium (Cr), Aluminum (Al) and Yttrium (Y). The formation of a self-healing protective “skin” of alumina (Al2O3) allows the ultra-thin steels to withstand the high temperatures and corrosive conditions in auto exhausts. These materials also have high thermal shock resistance and high melting and softening points and facilitated the development of high cell densities with very low pressure losses.

a) b) Figure 5: a: Metallic substrate converter, b: Ceramic substrate converter

7 Further development of metallic and ceramic substrates (Figure 5a & 5b) is described in the next section.

4

Current Catalyst Technology for Emissions Control Autocatalysts

Oxidation catalysts convert carbon monoxide (CO) and hydrocarbons (HC) to carbon dioxide (CO2) and water and decrease the mass of diesel particulate emissions but have little effect on nitrogen oxides (NOx). Three-way catalysts (TWC) operate in a closed loop system including a lambda- or oxygen sensor to regulate the air-fuel ratio. The catalyst can then simultaneously oxidize CO and HC to CO2 and water while reducing NOx to nitrogen. These simultaneous oxidizing and reducing reactions have the highest efficiency in the small air-to-fuel ratio window around the stoichiometric value, when air and fuel are in chemical balance. 4.1

Fast light off catalysts

The catalytic converter needs to work as fast as possible by decreasing the exhaust temperature required for operation so that untreated exhaust is curtailed at the start of the legislated emissions tests and on short journeys in the real world. Changes to the type and composition of the precious metal catalyst (Figure 6) and to the thermal capacity of substrates (figure 7) have together effected big reductions in the required operating temperature and light off times have been reduced from one to two minutes down to less than 20 seconds. [10]

Figure 6: Effect of catalyst technology on light off temperature

Figure 7: Effect of substrate cell density on light off time

The introduction of the new generation platinum/rhodium (Pt/Rh) technology for current and future emission standards is a technically and commercially attractive alternative for current palladium (Pd) based technologies for high demanding applications in close-coupled and under floor positions using different cold start strategies. [11] 4.2

More thermally durable catalysts

Increased stability at high temperature allows the catalytic converter to be mounted closer to the engine and increases the life of the converter, particularly during demanding driving. Pre-

8 cious metal catalysts with stabilized crystallites and washcoat materials that maintain high surface area at temperatures around 1000°C are needed. Improved oxygen storage components stabilize the surface area of the washcoat, maximize the air-fuel “window” for three-way operation and indicate the “health” of the catalytic converter for On Board.Diagnostic (OBD) systems. Figure 8 shows the progress made with mixed cerium and zirconium oxides. [12]

Figure 8: Improvements to thermal stability and oxygen storage capacity (OSC)

4.3

Substrate Technology

The technology of the substrates, on which the active catalyst is supported, has seen great progress. In 1974 ceramic substrates had a density of 200 cells per square inch of cross section (31 cells/square cm.) and a wall thickness of 0.012 inch or 12 mil (0.305 mm). By the end of the 1970’s the cell density had increased through 300 to 400 cpsi and wall thickness had been reduced by 50% to 6 mil. Now 400, 600, 900 and 1200 cpsi substrates are available and wall thickness can be reduced to 2 mil - almost 0.05 mm (Figure 9). [13, 14, 15, 16, 17] In the late 1970's substrates derived from ultra thin foils of corrosion resistant steels came onto the market. In the beginning the foils could be made from material only 0.05 mm thick allowing high cell densities to be achieved. Complex internal structures can be developed and today wall thickness is down to 0.025 mm and cell densities of 800, 1000 and 1200 cpsi are available (Figure 10). [18, 19] This progress in ceramic and metal substrate technology has major benefits. A larger catalyst surface area can be incorporated into a given converter volume and this allows better conversion efficiency and durability. The thin walls reduce thermal capacity and avoid the penalty of increased pressure losses. Alternatively the same performance can be incorporated into a smaller converter volume, making the catalyst easier to fit close to the engine, as cars get more compact. These improvements in substrate technology are now being applied in conjunction with heavy-duty diesel engines with catalysts placed as close as possible to the engine in order to.reduce the time to light off. To improve conversion behavior, catalysts are placed close to

9 the exhaust port before the turbocharger (Figure 11) and close-coupled catalysts using hybrid substrates are fitted (Figure 12). [20]

Figure 9: Progress in ceramic substrate design

Figure 10: Progress in metallic substrate design

Figure 11: Pre-turbo catalyst

Figure 12: Close-coupled hybrid catalyst

4.4

New Technology for Emissions Control Stoichiometric combustion

Conventional three-way catalysts are continually developed to improve high temperature stability and light off performance and to meet the demands of both the most challenging emission legislation in the world and new applications including motorcycles. Their performance can be further extended by the following additional technologies. 4.4.1 Hydrocarbon adsorbers Hydrocarbon adsorber systems incorporate special materials, such as zeolites, into or upstream of the catalyst. Hydrocarbon emissions are collected when exhaust temperatures are too low for effective catalyst operation. The hydrocarbons are then desorbed at higher temperatures when the catalyst has reached its operating temperature and is ready to receive and destroy the hydrocarbons. This technology has the potential to reduce hydrocarbons to less than half the levels emitted from a three-way catalytic converter (Figure 13). [21]

10

Figure 13: Influence of improved three-way catalyst and hydrocarbon adsorber on emissions (European cycle).

4.4.2 Electrically heated catalyst systems This uses a small catalyst ahead of the main catalyst. A metallic substrate, onto which the catalyst is deposited, allows an electric current to pass so it will heat up quickly. This brings the catalyst to its full operating temperature in a few seconds. [22] 4.5

Lean Combustion

With the development of lean burn direct injection gasoline engines and increased use of diesel engines, lean combustion is the challenge for automotive catalysis but is essential to reduce fuel consumption and limit CO2 emissions. New diesel technologies with electronic management and direct injection can achieve further fuel consumption improvements. Conventional three-way catalyst technology used on gasoline engines needs a richer environment with lower air-fuel ratios to reduce NOx so a radical new approach is required. DeNOx catalysts and NOx traps hold out the prospect of substantially reduced emissions of oxides of nitrogen. NOx conversion rates depend on exhaust temperature and availability of reducing agents. There are four systems under evaluation by industry: 1. Passive DeNOx Catalysts using reducing agents available in the exhaust stream 2. Active DeNOx Catalysts using added hydrocarbons as reducing agents 3. NOx traps or adsorbers used in conjunction with a three-way catalyst 4. Selective Catalytic Reduction using a selective reductant, such as ammonia from urea. Each of these systems offers different possibilities in the level of NOx control possible and the complexity of the system. Fuel parameters such as sulfur content can affect catalyst performance. 4.5.1 DeNOx (or Lean NOx) Catalysts DeNOx catalysts use advanced structural properties in the catalytic coating to create a rich "microclimate" where hydrocarbons from the exhaust can reduce the nitrogen oxides to nitrogen, while the overall exhaust remains lean. Further developments focus on increasing the operating temperature range and conversion efficiency. 4.5.2 NOx Adsorbers (or Lean NOx Traps) NOx traps are a promising development as results show that NOx adsorber systems are less constrained by operational temperatures than DeNOx catalysts. NOx traps adsorb and store NOx under lean conditions. A typical approach is to speed up the conversion of nitric oxide (NO) to nitrogen dioxide (NO2) using an oxidation or three-way catalyst mounted close to the

11 engine so that NO2 can be rapidly stored as nitrate. The function of the NOx storage element can be fulfilled by materials that are able to form sufficiently stable nitrates within the temperature range determined by lean operating points of a direct injection gasoline engine. Thus especially alkaline, alkaline earths and to a certain extent also rare-earth compounds can be used. When this storage media nears capacity it must be regenerated. This is accomplished in a NOx regeneration step. Unfortunately, alkaline and alkaline earth compounds have a strong affinity for sulfation. As a consequence alkaline and alkaline earth compounds are almost irreversibly poisoned by the sulfur contained in the fuel during the NOx storage operation mode, leading to a decrease in NOx adsorption efficiency during operation. The stored NOx is released by creating a rich atmosphere with injection of a small amount of fuel. The rich running portion is of short duration and can be accomplished in a number of.ways, but usually includes some combination of intake air throttling, exhaust gas recirculation, late ignition timing and post combustion fuel injection. The released NOx is quickly reduced to N2 by reaction with CO (the same reaction that occurs in the three-way catalyst for spark-ignited engines) on a rhodium catalyst site or another precious metal that is also incorporated into this unique single catalyst layer (Figure 14).

Figure 14: NOx adsorber working principle

Under oxygen rich conditions, the thermal dissociation of the alkaline and alkaline earth sulfates would require temperatures above 1000 °C. Such temperatures cannot be achieved under realistic driving conditions. However, it has been demonstrated in various publications [23, 24, 25] that it is in principle possible to decompose the corresponding alkaline earth sulfates under reducing exhaust gas conditions at elevated temperatures. In this way the NOx storage capacity can be restored. The heating of the catalyst, for example by late ignition timing, does however result in a considerable increase in fuel consumption, which is dependent upon the sulfur content. Therefore, reducing the sulfur concentration in the fuel must be regarded as the most effective way of using the full potential of modern direct injection gasoline engines with respect to fuel economy and CO2 reduction. One of the demands for a desulfation strategy must be to avoid any H2S emissions above the odor threshold during desulfation. [26, 27] Developments and optimization of NOx adsorber systems have been and are currently underway for diesel and gasoline engines. These technologies have demonstrated NOx conversion efficiencies ranging from 50 to in excess of 90 percent depending on the operating temperatures and system responsiveness, as well as fuel sulfur content. [28, 29] The system is in production with direct injection gasoline engines.

12 4.5.3 Selective Catalytic Reduction (SCR) SCR is a widespread technology to reduce nitrogen oxide emissions from coal, oil and gas fired power stations, marine vessels and stationary diesel engine applications. SCR technology has been used successfully for more than two decades. SCR technology for heavy-duty diesel vehicles has been developed to the commercialization stage and will be available as an option in the series production of several European truck-manufacturing companies in 2001. SCR technology permits the NOx reduction reaction to take place in an oxidizing atmosphere. It is called “selective” because the catalytic reduction of NOx with ammonia (NH3) as a reductant occurs preferentially to the oxidation of NH3 with oxygen. Several types of catalyst are used, the choice of which is determined by the temperature of the exhaust environment. For mobile source applications the reductant source is usually urea, which can be rapidly hydrolyzed to produce ammonia in the exhaust stream. SCR for heavy-duty vehicles reduces NOx emissions by circa 80%, HC emissions by circa 90% and PM emissions by circa 40% in the EU test cycles, using current diesel fuel (<350 ppm sulfur). [30, 31] Fleet tests with SCR technology show excellent NOx reduction performance over more than 500,000 km of truck operation, and the experience is based on over six million kilometers of accumulated commercial fleet operation. [32, 33, 34] Though real world durability has been proven, the real challenge for using SCR systems to reduce NOx with heavy-duty diesel vehicles and buses is the development of an infrastructure for the delivery of the preferred reductant; a urea/water solution. The combination of SCR with a pre-oxidation catalyst, a hydrolysis catalyst and an oxidation catalyst enables higher NOx reduction under low loads and low temperatures. For combination technologies using oxidation catalysts, catalyzed filters or any catalyst formulations including precious metals the use of diesel fuel with sulfur lower than 10 ppm is necessary to keep sulfate particulate formation below future legislated limits. [35, 36, 37] 4.5.4 Diesel particulate filters (DPF) DPF systems consist of a filter material positioned in the exhaust designed to collect solid and liquid particulate matter (PM) emissions while allowing the exhaust gases to pass through the system. One type of filter material based on a highly porous cordierite monolith with channels blocked at alternate ends is shown in Figure 15.

Figure 15: Wall flow diesel particulate filter

Figure 16: Flow-through diesel particulate trap

13

Figure 17: Performance of wall flow DPF on PM size and number

A number of filter materials are used, including ceramic monoliths, woven silica fiber coils, ceramic foam, wire mesh and sintered or shaped metals. A new flow-through particulate trap has recently been developed using metal foils (Figure 16). [38] Collection efficiencies of these various filters range from 30 percent to over 90 percent, but most DPFs achieve over 99% when expressed as numbers of ultra fine particles (Figure 17). This is very important since health experts believe that it is the fine particulate that is carried deep into the lungs and which is thought to be the most dangerous size of PM. Since the wall flow filter would readily become plugged with particulate material in a short time, it is necessary to “regenerate” the filtration properties of the filter by regularly burning off the collected PM. The most successful methods to initiate and sustain regeneration include: 1. Incorporating a catalytic coating on the DPF to lower the temperature at which particulate matter burns. [39] 2. Using very small quantities of fuel-borne catalyst, such as cerium oxide. The catalyst, when collected on the DPF as an intimate mixture with the particulate, allows the particulate to burn at normal exhaust temperatures to form carbon dioxide and water, while the solid residues of the catalyst are retained on the DPF. [40] 3. Incorporating an oxidation catalyst upstream of the DPF that, as well as operating as a conventional oxidation catalyst, also increases the ratio of NO2 to NO in the exhaust. Trapped particulates burn off at normal exhaust temperatures using the powerful oxidative properties of NO2. [41] 4. Electrical heating of the DPF either on or off the vehicle, which would allow simple regeneration but imposing a fuel penalty. 5. DPF systems and intelligent engine-management allow efficient regeneration under all operating conditions. Diesel passenger cars equipped with a DPF in conjunction with fuelborne catalyst and an oxidation catalyst are now in series production. Recent evaluations indicate a good and durable performance of the system. [42, 43] Continuously regenerating DPFs are very successful in retrofit applications of older heavy-duty diesel vehicles and buses in various regions over the world. Real world durability of these systems is proven every day in major cities in Europe and the US. [44, 45, 39]

14 4.5.5 Combined emission control systems The stringent emission limit values (NOx and PM) for heavy-duty diesel (HDD) engines in 2005 and 2008, the request by the transport sector for minimum fuel consumption of the engine, requiring an optimized combustion process and therefore minimum CO2 and engine-out particulate matter, and the political, public and health concerns on the emission of ultra fine particles into the atmosphere make a combined emission control system an attractive proposition. Urea-based SCR or NOx adsorber systems, in combination with DPFs and an appropriate regeneration strategy will, it is anticipated, be used on the new generation HDD engines and show significant reduction of both NOx and particulate matter. [46, 47, 48]

5

Conclusions

Gaseous and particulate emission controls using catalytic, adsorption and trapping technologies with advanced materials have been developed. The automotive emissions control industry working with its partners in the automotive industry will meet the challenge of future emission regulations. Advanced catalyst and trap systems, with optimized engine management, will aid the achievement of the future low emission standards deemed necessary to meet air quality goals. Retrofitting of catalyst systems and particulate traps is increasing in response to fiscal incentive schemes introduced by governments and to meet the requirements of environmental zones - particularly in cities.

6

References

[1]

Official Journal of the European Communities, L 350, Vol. 41, 28 December 1998; Directives 98/69/EC (emissions) and 98/70/EC (fuels) Directive 1999/96/EC COM (2000) 626 French Patent FR 402173 and British Patent GB 9364/1909 Scientific American R. M. Heck and R. J. Farrauto, “Catalytic Air Pollution Control – Commercial Technology”, Van Nostrand Reinhold, 1995. M. L. Church, B. J. Cooper and P. L. Wilson, “Catalyst Formulations 1960 to Present”, SAE 890815. J. R. Mondt, “Cleaner Cars – The History and Technology of Emission Control Since the 1960s”, SAE, 2000. Engelhard Industries, US 3441381. R. J. Brisley et al, "The Use of Palladium in Advanced Catalysts", SAE 950259. J. Schmidt et al, “Utilization of advanced Pt/Rh TWC technologies for advanced gasoline applications with different cold start strategies”, SAE 2001-01-0927, March 2001 J-P. Cuif et al, “(Ce, Zr)O2 Solid Solutions for Three Way Catalysts”, SAE 970463. S. T. Gulati, "Thin Wall Ceramic Catalyst Supports", SAE 1999-01-0269

[2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]

15 [14] J. Schmitt et al, "The Impact of High Cell Density Ceramic Substrates and Washcoat Properties on the Catalytic Activity of Three Way Catalysts", SAE 1999-01-0272 [15] S. Kikuchi et al, “Technology for Reducing Exhaust Gas Emissions in Zero Emission Level Vehicles”, SAE 1999-01-0772 [16] S. T. Gulati, "Design Considerations for Advanced Ceramic Catalyst Supports", SAE 2000-01-0493 [17] K. Nishizawa et al, “New Technologies Targeting Zero Emissions for Gasoline Engines”, SAE 2000-01-0890 [18] R. Brück et al, "The Necessity of Optimizing the Interactions of Advanced PostTreatment Components in Order to Obtain Compliance with SULEV-Legislation", SAE 1999-01-0770 [19] W. Maus, R. Brück and G. Holy, “Zukünftige Abgasnachbehandlungstechnologien für Otto-Motoren; Die nächste Generation Niedrigstemissionsfahrzeuge”, AVL Congress, Graz, Sept. 1999 [20] M. Reizig, R. Brück, R. Konieczny, P. Treiber, “New approaches to Catalyst Substrate Application for Diesel Engines”, SAE 2001-01-0189, March 2001 [21] N. Noda, A. Takahashi, Y. Shibagaki and H Mizuno, “In-line Hydrocarbon Adsorber for Cold Start Emissions – Part II, SAE 980423, Feb 1998 [22] F. J. Hanel, E. Otto, R. Brück, T. Nagel and N. Bergaul, “Practical Experience with the EHC System in the BMW ALPINA B12”, SAE 970263 [23] W. Strehlau, J. Leyrer, E.S. Lox, T. Kreuzer, M. Hori and M. Hoffmann: “New Developments in Lean NOx Catalysis for Gasoline Fuelled Passenger Cars in Europe”, SAE 962047 [24] M.S. Brogan, R.J. Brisley, A.P. Walker, D.E. Webster, W. Boegner, N.P. Fekete, M. Kraemer, B. Krutzsch and D. Voigtlaender, “Evaluation of NOx Storage Catalysts as an Effective System for NOx Removal from the Exhaust Gas of Lean burn Gasoline Engines”, SAE 952490 [25] U. Göbel, T. Kreuzer and E. S. Lox, “Moderne NOx-Adsorber-Technologien, Grundlagen, Voraussetzungen, Erfahrungen”, Proceedings of the VDA-conference, Frankfurt (1999) [26] K.-H. Glück, U. Göbel, H. Hahn, J. Höhne, R. Krebs, T. Kreuzer and E. Pott: “Die Abgasreinigung der FSI-Motoren von Volkswagen”, MTZ Motortechnische Zeitschrift, 6, 402 (2000) [27] U. Göbel, J. Höhne, E.S. Lox, W. Müller, A. Okumura, W. Strehlau and M. Hori, “Durability Aspects of NOx-Storage Catalysts for Direct Injection Gasoline Vehicles”, SAE 99FL-103 [28] H. Lüders, P. Stommel and S. Geckler, “Diesel Exhaust Treatment – New Approaches to Ultra Low Emission Diesel Vehicles”, SAE 1999-01-0108, Mar 1999.29 N. Ruzicka and T. Liebscher, “Possible Aftertreatment Concepts for Passenger Car Diesel Engines with Sulphur-free Fuel”, SAE1999-01-1328, Mar 1999 [29] S. Fischer, L. Hofmann and W. Mathes, “The development of the SINOx system for commercial vehicles for serial applications”, 20 th Vienna Motor Symposium, May 6-4 1999, VDI Fortschrittsberichte Reihe 12, Nr. 376, 267–282 [30] B. Amon, S. Fischer, L. Hofmann and J. Zuerbig, “The SINOx system for trucks to fulfil the future emission regulations”, CAPoC 5 Brussels, April 14–16 2000

16 [31] Siemens AG and TÜV Automotive, “Investigation on long-term stability of diesel DeNOx catalyst exhaust gas aftertreatment systems on 3 MAN and 10 DaimlerChrysler trucks – results of the 2nd Bavarian Road Test” – Final Report [32] N. Fritz, W. Mathes, J. Zürbig and R. Mueller, “On-road demonstration of NOx emission control for diesel trucks with SINOx urea SCR system”, SAE paper 1999-01-0111 [33] W. Miller, J. Klein, R. Mueller, W. Doelling and J. Zürbig, “The development of ureaSCR technology for US Heavy-Duty Trucks”, SAE paper 2000-01-0190 [34] E. Jacob, G. Emmerling, A. Döring, U. Graf, M. Harris, J. van den Tillart and B. Hupfeld, “Reduction of NOx from HD diesel Engines with urea SCR compact systems”, 19th Vienna Motor Symposium, May 3-5 1998, VDI Fortschrittsberichte Reihe 12, Nr. 348, 366–386 [35] E. Jacob and A. Döring, “GD-Kat: Exhaust Treatment System for Simultaneous Carbon Particle Oxidation and NOx Reduction for Euro 4/5 Diesel HD Engines”, 21st Vienna Motor Symposium, May 4–5 2000 [36] J. Gieshoff, A. Schäfer-Sindlinger, P.C. Spurk, J.A.A. van den Tillaart and G. Garr, “Improved SCR Systems for Heavy Duty Applications”, SAE 2000-01-0189 [37] R. Brück, P. Hirth, M. Reizig, P. Treiber and J. Breuer, “Metal Supported Flow-Through Particulate Trap; a Non-Blocking Solution”, SAE 2001-01-1950. [38] K. Voss et al, “Engelhard’s DPX catalysed soot filter technology for emissions reduction from Heavy-Duty Diesel engines with passive regeneration”, presentation given by R. Kakwani, SAE TOPTEC, Gothenburg, September 2000 [39] P. Zelenka et al, "Towards Securing the Particulate Trap Regeneration: A System Combining a Sintered Metal Filter and Cerium Addition", SAE 982598 [40] P. N. Hawker et al, "Effect of a Continuously Regenerating Diesel Particulate Filter on Non-Regulated Emissions and Particle Size Distribution", SAE 980189 [41] O. Salvat, P. Marez, G. Belot, “Passenger car serial application of a particulate filter system on a common rail direct injection diesel engine”, SAE 2000-01-0473, March 2000 [42] ADAC and UBA Press Release, 28 August 2001. [43] R. Allansson, B. J. Cooper, J. E. Thoss, A. Uusimäki , A. P. Walker and J. P. Warren, “European experience of high mileage durability of continuously regenerating diesel particulate filter technology”, SAE 2000-01-0480, March 2000 [44] T. Lanni, S. Chatterjee, R. Conway, H Windawi, et al, “Performance and durability evaluation of continuously regenerating particulate filters on diesel powered urban buses at NY city transit”, SAE 2001-01- 0511, March 2001 [45] M. Khair, J. Lemaire and S. Fischer, “Achieving Heavy-Duty Diesel NOx/PM levels below the EPA 2002 standards – an integrated solution”, SAE 2000-01-0187 [46] M. Khair, J. Lemaire and S. Fischer, “Integration of EGR, SCR, DPF and fuel-borne catalyst for NOx/PM reduction”, SAE 2000-01-1933 [47] G.R. Chandler, B.J. Cooper, J.P. Harris, J.E. Thoss, A. Uusimäki, A.P. Walker and J.P. Warren, “An Integrated SCR and Continuously Regenerating Trap System to Meet Future NOx and PM Legislation”, SAE 2000-01-0188

II Metals

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

Development Status of Metal Substrate Catalysts Rolf Brück Emitec GmbH, Lohmar

1

Introduction

The need for mobility will increase in future, particularly in the fast-developing nations. The event of this will permit division of labour, i.e. increased productivity, which in turn will bring the necessary increase in national income. (see figure 1).

Figure 1: The relationship between vehicle volume and gross national product

Growing world prosperity will also increase the volume of cars and motorcycles. A disproportionately large increase in mobility is therefore to be expected as compared to population growth. With a prospective population growth of 15 percent by the year 2010, this could lead to a 45 percent increase in car numbers and even a 50 percent increase in the volume of two and three wheel vehicles [1]. But this economically necessary development should not come at the expense of the environment. For this reason, emissions limits are being introduced around the world and continually being made stricter. The rapid further development in exhaust gas treatment systems caused by this is supported by new engine technologies, with reduced fuel consumption, low untreated emissions, and better engine control systems. Due to its geographical position and the associated climactic conditions of the Los Angeles valley, which have lead to high concentrations of exhaust contaminants, California has assumed a leading role in legislating for vehicle emissions. The remaining US states, Europe, Japan and other nations are following suit with similar emissions limits.

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KgaA ISBN: 3-527-30491-6

20 Figure 2 shows the development of the Californian emissions legislation for required hydrocarbon (HC) conversion rates. The conversion rates were based on untreated emissions of 2.0 g/mile.

Figure 2: The required conversion rates for the whole FTP test depending on untreated emissions and emissions limits

Figure 3: Current and future propulsion systems

The increase in conversion rates from 96.25 percent for low emission vehicles to 99.5 percent for super ultra low emission vehicles (with reference to the untreated emission mentioned

21 above of 2.0 g/mile) only amounts to around 3 percent in absolute terms, but compared to the residual emissions means a 7.5- fold reduction. The aim of all these efforts is to make the car so „clean“ in terms of its limited, contaminant emissions, that compared to environmental immissions, it emits equally low or even lower concentrations. This means that is not polluting the environment, but can even clean it. As well as improving traditional engines (Otto and Diesel) and catalyst technology, new propulsion concepts such as hybrid vehicles and fuel cell systems are also being discussed in order to reach these goals. However, when one compares the states of development and above all the costs of these systems, it is clear that at least over the next 15–20 years, the „normal“ combustion engine will continue to be the leading propulsion concept.

2

Metal Substrate Catalyst Technologies for Otto Engines

Besides ceramic catalysts, Metallit metal substrate catalysts have also been increasingly used over the last 15 years for exhaust gas treatment in mass-produced vehicles. The substrates comprise thin (0.05 mm–0.02 mm), smooth and corrugated metal foils that are connected to the monolith in a high-temperature brazing process. Figure 4 shows an example of a close-coupled catalyst.

Figure 4: Metal substrate catalyst in a close-coupled arrangement; metal microstructure with catalytic active coating

A large catalytic surface area is required for optimal catalytic effectiveness. Metal substrates are produced today in cell densities of up to 1600 cells/inch 2 (cpsi) and provide a specific geometric surface area (GSA) of up 6.8 m 2 /l (figure 5).

22

Figure 5: Geometric surface area (GSA) of uncoated Metalit substrates; influence of cell density

The catalytic reaction of hydrocarbon (HC), carbon monoxide (CO) and nitrogen oxide (NOx) only occurs at temperatures above 250 °C, so that for a cold start, before the catalyst has been heated by the exhaust gas, emissions are not converted. To be able to meet the required total conversion rates shown in figure 2, the cold start phase must be reduced to a few seconds. The heat capacity of the catalyst as an inert mass is the most important factor here. With consistent further development of the foil material, it is possible to reduce the foil thickness to 0.02 mm. And it is thus possible to produce catalyst substrates with 1600 cpsi, which have heat capacities 20 % lower than traditional 400 cpsi substrates (figure 6). Clearly improved cold start behaviour can thus be attained.

Figure 6: Heat capacity of coated Metalit substrate; Influence of foil thickness and cell density

The advantages of a high cell density catalyst with respect to HC results for a medium range car are shown in figure 7. The 1600 cpsi catalyst shows an improvement of 50 percent compared to a 600 cpsi catalyst. Uniform flow distribution is necessary to guarantee optimal utilisation of the catalyst volume. Often it is not possible to achieve a good flow in the car, par-

23 ticularly in the case of the close- coupled arrangement because of the spatial conditions. Due to the design freedom afforded by metal substrates, it is possible to produce cone-shaped catalysts (figure 8). The cone-shaped channels also guide the flow to the peripheral areas of the second substrate installed behind the „ConiCat®“, allowing better flow distribution and therefore a cost-optimised use of the catalyst system.

Figure 7: Influence of the substrate‘s cell density on HC emissions in the FTP emission test cycle; Æ 118 × 150 mm

Figure 8: Cone-shaped catalyst, „ConiCat®“; improved flow distribution and better space utilisation

24

3

Metal Substrate Catalyst Technologies for Diesel Engines

Following the experience of tailor fitted solutions for metal substrates for gasoline engines, research in automotive catalysts was carried out to develop components, which substantially improve emissions control of Diesel engines. As a result of the very efficient combustion process, Diesel engine exhaust gas temperatures are relatively low, especially under real driving conditions at partial load and speed. While for gasoline engines the main emissions control problem is how to reach light-off temperature in the catalyst as quickly as possible after starting the engine, Diesel engines have an additional problem. They tend to fall back below lightoff temperature during deceleration and low idling modes e.g. in the European driving cycle (figure 9).

Figure 9: Temperature comparison of a Diesel and a spark ignition engine in the European driving cycle

Figure 10 shows an engine layout with possible locations for installing catalysts. The closest position to the engine is with the catalyst in front of the turbo charger. This location offers the advantage of higher exhaust temperature compared to fitting it after turbine because this device works as a heat sink during cold start and takes away energy during operation. Further possibilities are the close-coupled position after the exhaust turbocharger and the standard underfloor position.

25

Figure 10: Possible catalytic converter positions

3.1

The pre-turbo catalyst

The most important criterion for the design of this type of catalyst is the very small space which is available in the connection of an exhaust manifold to the turbocharger. Most of the pre-turbo catalysts have substrate volumes of distinctly less than 100 cm3 . Figure 11 shows the pre-turbo charger catalyst in its fitted position. To show the mode of operation of the pre-turbo charger catalyst, emission tests were carried out on a 3.0 l 6-cylinder Diesel engine in the European test cycle. Figure 12 shows the advantages in the HC conversion rates.

Figure 11: Pre-turbo catalyst

It can be seen that the conversion rates were clearly improved over the whole test cycle. An average improvement of 20 percent was achieved in the first 500 phases of the test. Similar results were also obtained for CO.

26

Figure 12: Influence of the pre-turbo catalyst on HC conversion efficiency; 200 cpsi; length: 30 mm

3.2

The hybrid catalyst

The idea is to influence the light-off behaviour of the catalyst by different substrate thermal masses. Using metal foils as the basis for a substrate, it is easy to vary the foil thickness (and thus the thermal mass) for optimisation of the catalyst. Since the exhaust temperature for Diesel engines always hovers around the light-off temperature (figure 9), it is advantageous to have a low heat capacity in the front part of the catalyst (foil thickness 0.03 mm) to allow quick heating when the gas temperature rises. By using a thicker foil (0.08 mm) a heat accumulator is fitted in the rear part, so that even when the gas temperature falls below the light-off temperature this part of the catalyst remains catalytically active because of the energy stored in it (figure 13).

Figure 13: Hybrid catalyst

Measurement results – A modern 3 litre 6-cylinder in-line engine with common-rail injection and inter-cooled turbo charging (turbine with variable geometry) on a test bench was used

27 for laboratory testing. During all experiments the engine setting was unchanged. Two quickresponse analysing systems were installed for modal analysis of the gaseous concentrations in the untreated exhaust and behind the catalytic converter. The exhaust system was original equipment from a medium-sized passenger car. The sequence of engine speed and torque was set according to measured values during EU III test runs on a chassis dynamometer. The hybrid catalysts were in the close-coupled position of the real car instead of the original ceramic cats. The insulation of the exhaust system was also taken from the car. The exhaust mass flow was calculated by measuring the fuel consumption and the mass flow of the combustion air on a modal basis. Two systems with identical coating were compared: the original ceramic closecoupled catalyst (oval shape, diameters 90 mm × 185 mm, length 114 mm, volume 1492 cm 3 , 62 cells/cm 2 , wall thickness 165 m) and the hybrid catalyst (diameter 127 mm, 62 cells/cm2 , 1. brick length 31 mm with 30 m foil thickness, 2. brick length 70 mm with 80 m foil thickness, volume of the hybrid catalyst 1393 cm3 ) plus the original underfloor catalysts (volume 762 cm3 each, 62 cells/cm 2 , foil thickness 40 m). Figure 14 shows the HC conversion rates for the hybrid and the mass-produced catalyst system in the European test cycle.

Figure 14: EU III test results of a hybrid catalyst compared to the standard production system

The hybrid concept shows clear advantages over the standard system, with HC results reduced by 50 percent from 0.16 g/km to 0.08 g/km. CO emissions were even reduced by 65 %. 3.3

Flow through particulate trap

„Wall flow“ particle traps consisting of ceramic substrates with alternately sealed channels have been available for a number of years and have already been installed in mass produced vehicles. Such a filter can achieve an efficiency of more than 95 % over the total range of particle sizes. In addition to chemical interactions with additives and special coatings, the safe regeneration, i.e. combustion of soot in all sorts of vehicle operation, still cause problems. With excessive amounts of deposited soot the exhaust gas back-pressure increases and during the soot incineration process, temperature peaks develop in the filter leading to mechanical dam-

28 age. Therefore, silicon carbide is used in modern applications in view of its high temperature stability. In order to circumvent the disadvantage of discontinuous regeneration, continuously regenerating filter systems CRT were developed. In such a system particles are incinerated at temperatures above 200 °C via oxidation by NO2. The NO2 is often generated in oxidation catalysts upstream of the particle trap. The exhaust gas temperatures of a Diesel engine, especially in low-load operation, are so low, however, that only an insufficient amount of NO is oxidised to NO2. The oxidation behaviour can be improved by means of pre-turbo charger (3.1) and hybrid catalyst (3.2). In addition, such a trap would have to be installed close to the engine to guarantee the highest possible exhaust gas temperature. Upstream of the filter an oxidation catalyst which oxidises CO and HC and subsequently NO into NO2 has to be installed. It is well known that soot is deposited at the gas inlet front face of the catalyst in the Diesel exhaust gas pipe which partially reacts with NO2. It was the task of the flow through trap development to ensure that the filtering efficiency of a catalytic substrate was increased and that deposited particles would not be blown out again by a sudden increase of the mass flow and the resulting aerodynamic forces. The development is based on a mixing catalyst support originally used for the distribution of Urea in SCR systems (Figure 15). The vanes of the mixing section pass part of the exhaust gas of each cell into a neighbouring channel. The construction of a mixing honeycomb substrate is used but the perforated flat foils are replaced by sintermetal foils so that gases can pass from one channel into the neighbouring channel. If a porous flat foil of wire mesh, fibre material or stretch material is used, some of the particles which are in the exhaust gas can be trapped by passing the porous foil.

Figure 15: Flow through particulate trap

29 The following four constant load points were driven in order to see the trapping efficiency. The results are shown in figure 16. A trapping efficiency between 12 % and 31 % can be seen. Higher efficiencies should be possible by increasing the size of the trap. The results are shown in figure 16.

Figure 16: Particle emissions in different constant load points with and without flow through trap

Table 1: Torque and engine speed of tested constant load points Engine Speed [rpm] 1200 Torque [Nm] 150

4

2 000 195

3 000 194

4 000 155

Mechanical and thermal loads on catalysts

Metal substrates are used today for a great variety of applications. Figure 17 shows the resulting mechanical and thermal loads. Using FEM calculations it is possible to test the construction of the catalyst substrate with respect to application-dependent loads. Figure 18 shows the calculated and measured selfresonance of a metal substrate. It can be seen that both the vibration characteristics and the frequency match well. Values calculated in this way can be compared to stimuli from the vehicle.

30

Figure 17: Thermal / mechanical load distribution of various applications

Figure 18: FF Analysis of metallic substrates; modelling and results of vibration analysis

5

Conclusions

Driven by stricter emissions limits, the effectiveness of catalyst systems must be increased continually. Metal catalyst substrates offer a variety of solutions for all combustion engine applications: · Metal substrate technologies are available for both spark ignition and Diesel engines · All emissions (HC, CO, NOx, PM) can be significantly reduced · New developments like the cone-shaped catalyst and the open particulate traps help to meet future emissions legislation · New, high cell density, ultra thin foil substrates further increase catalyst efficiency · Stress and durability of metal substrates can be precalculated using detailed FEM models before component or engine tests

Materials Issues Relevant to the Development of Future Metal Foil Automotive Catalytic Converters J. R. Nicholls* and W. J. Quadakkers *Cranfield University, Cranfield, Bedford, UK Forschungszentrum Juelich, Juelich, Germany

1

Abstract

Metal supported automotive catalytic converter bodies are based on ferritic steels, with 5–8wt%Al, 17–22wt%Cr plus small additions of reactive elements. To improve the catalyst performance there is a continued drive towards higher operating temperatures, thinner components and alternative geometries offering large surface area to volume ratios. Maintaining acceptable component lives is mandatory, even when thinner support geometries and higher operating temperatures are envisaged. This has led to a number of materials development strategies, including alternative substrate geometries, modifications to the alloy composition, both through the addition of multiple reactive elements and through the close control of trace element additions, and the development of surface treatments to increase the available aluminium reservoir. Each of these strategies will be reviewed. Service life predictions, not only relies on suitable materials, but also on the existence of adequate models and simulation techniques, supported by reliable data. In the present paper, the latest thoughts on modelling the oxidative failure of the FeCrAlRE based materials will be presented.

2

Introduction

Following the Kyoto agreement, there has been a concerted and intensified push to lower the worldwide emissions from industrialised power plants [1]. Motor vehicles and the “mobile society” are significant contributors; within the industrialised countries the norm is 400–600 cars per 1000 population, USA is the highest at 800 cars per 1000 population, whilst the less developed countries are much lower (50–100 cars per 1000 population) but aspire towards the industrialised position [2]. Following this scenario, it is likely that there will be 1,500 million vehicles in use worldwide in 2010, thus low emission vehicles are perceived as a major factor in achieving the Kyoto emissions targets. Lowering of hydrocarbon emission from motor vehicles can be achieved in one of two ways, firstly, making the engines more fuel efficient and secondly, cleaning up the hydrocarbon emissions from the exhaust gas. It can be seen therefore that catalytic converters must play a significant role and will be an indispensable part of future '‘clean engine' design. As a result of this worldwide emissions awareness a demand for new, innovative, efficient catalytic converter technologies has arisen. Ceramic and metal substrate technologies have

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

32 been developed, together with catalytic coatings to provide efficient conversion [3]. Metallic substrates offer improved efficiencies as cell densities can be increased and they produce a lower pressure drop per unit length of converter in the exhaust gas system. This leads to an increased volume converter and therefore further reduction in emissions. Figure 1 illustrates this reduction in total hydrocarbon emissions with both cell density and catalytic converter volume increase. The future direction is obvious, to higher capacity converters with increased cell densities (up to or in excess of 1200 cpsi), whilst maintaining or reducing the exhaust gas pressure drop. This has significant implications on the materials used to design the catalytic converter, particularly the support substrate. Furthermore, the metallic substrate converter has a lower heat capacity and therefore can more rapidly rise to the operating temperature from a cold start. A significant fraction of emissions is produced in this ‘cold start’ phase, particularly for vehicles undertaking short duration journeys. This paper addresses material issues relevant to the development of future metal foil automotive catalytic converters.

Figure 1: Converter emissions efficiency depends on the converter surface area and cell density [reproduced from reference 2 – hydrocarbon emissions measured in bag 1 of an FTP test after a ‘cold start’ operating condition].

33

3

Metal Supported Automotive Catalysts

Metal supported automotive catalytic converter bodies are, in the main, based on ferritic steels with 5–8wt% aluminium, 17–22wt% chromium, plus small additions of reactive elements added to improve the oxidation resistance of the alloy and aid oxide adhesion. These alloys are protected by the formation of a slow growing surface oxide, usually alumina and it is critical to component life that this naturally formed alumina protective layer is maintained. Breakdown of this thermally grown alumina would lead to breakaway oxidation conditions and rapid component failure [4]. Thus the early formation and maintenance of this protective oxide is key to the life of metallic automotive catalytic converter bodies. Table 1 lists a number of commercial alloys that may be used to manufacture the metallic catalyst support. All have 5wt%–6wt% aluminium. Increasing the aluminium content above 6wt% would be an advantage; as discussed later this would increase the available aluminium reservoir and hence component life. However, such increases lead to the formation of intermetallic phases within the alloy, embrittling the metallic foil, and therefore limiting the minimum thickness of foil that can be produced. Foil thickness is a critical parameter in achieving high cell densities; 200cpsi supports are made from 70mm foil, 400 cpsi from 50mm foil, while 600 cpsi supports would have to be made from 30mm foil. Thus to improve component efficiencies there is a drive to use thinner foils and it is likely that components will be manufactured from 20mm material in the foreseeable future, provided the available aluminium reservoir (see Section 3) can be maintained in these thinner section components to provide the desired component life. Figure 2 illustrates the advanced design of foils in an electrically heated converter (EHC) produced by Emitech GmbH [5] including the corregated nature of the catalytic support design (b) and fine dimension brazed joints (c). Table 1: Materials for metal foil automotive catalytic converters Alloy Aluchrom Aluchrom YHf FeCrAlloy FeCrAlloy JA13 Kanthal AF Kanthal APM Nisshin steel Ugine Savoie 12178 Ugine Savoie 12179

Fe Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal.

Cr 20.6 20.3 20.3 16.3 21.1 21.1 19.9 19.9 20.0

Al 5.4 5.6 5.4 5.0 5.2 5.9 5.0 5.0 5.0

Ti La .010 .010 .084 .010 .094 .026 0.120 0.009 0.014

Zr 0.170 0.054 0.080 0.050 0.058 0.110

Y .006 .046 .045 .320 .034

Ce

Hf .031

0.019 0.030

Future designs of advanced catalytic converter will have to address “cold start” behaviour. Some 70% of the total emissions from a vehicle, constitute the cold start phases of engine operation [2]. This is recognised as a major problem for the larger engine and for stop/start operations. The increased cell density reduces the thermal mass and this helps to reduce the problem by ensuring the catalytic converter comes to operating temperature earlier. Sophisticated designs, like electrically heater converter (see Figure 2) [5] help to overcome these problems, but at a cost. Other approaches involve the repositioning of the converter, from under the floor to close coupled to the exhaust manifold. This ensures the converter more rapidly

34 heats to its operating temperature, but also imposes high thermal loads, with a commensurate increase in converter metal temperature. Increased metal support temperatures will further increase converter efficiency, but at the expense of converter component life. Oxidation of the FeCrAl-RE support material depends exponentially on temperature. Thus increased temperatures must lead to a more rapid growth of the alumina scale on these alloys, consuming aluminium from thin sectioned components within the converter body and ultimately leading to chemical failure. The increased thermal loads will also influence the level of mechanical stress and thermal fatigue that the converter body will see.

a) b) c) Figure 2: Advanced catalyst design [5]; a) EHC design of metal foil substrate; b) catalytic support manufactured from flat and corrugated foils; c) brazed joints of foil components

The current temperatures seen by the converter body, in the under floor arrangement, depends on the driving conditions. Figure 3 shows the thermal transient measured in a bench test engine under high load conditions [5]. It can be seen that the matrix core rapidly reaches 1000 °C, whereas the outer support tube only reaches a temperature of 700 °C, and that this peak temperature lags behind that for the matrix core. During cool down the outer support tube cools more slowly (it has a greater thermal mass), thus stop/start condition results in a cyclic thermal stress on the converter, as well as a severe temperature cycle. High speed driving achieves peak temperatures similar to the bench endurance test, depicted in Figure 3. Temperatures of 910 °C have been recorded (100% load : 114Nm) [6], with exhaust gas velocities of 25ms–1. The exhaust gas was predominantly nitrogen, water vapour and CO2 (14.6 vol.%), with a small amount of oxygen (0.43 vol.%), CO (1.05 vol.%) hydrogen (0.35 vol.%), SO2 (6.4 vpm), NOx (2670 vpm) and hydrocarbons (380 vpm). Medium speed cruising and urban driving are less stressful, with temperatures of 505 °C and 280 °C being recorded respectively. In both these latter cases, the NOx level is significantly reduced and the hydrocarbon emissions level increased as one moves from high speed running to urban driving conditions [6]. The other exhaust compositions change little, other than the oxygen content of the exhaust increase from 0.43 to 1.17 vol%. For close coupled converter applications higher temperatures are to be expected. In general, temperatures are expected to be 100 °C hotter, thus, future designs will be based on a continuous service temperatures of 1050 °C, with short peak excursions to 1100 °C [7]. In addition, the converter must stand cyclic stressing due to exhaust gas pulsation, mechanical vibrations from the engine and road surface, thermal shock due to exhaust temperature fluctuations and splash water on the converter casing during normal motoring.

35

Figure 3: Thermal cycle in a metal supported catalyst during bench endurance testing [5]

4

Materials Issues Relevant to the Development of Future Metal Foil Automotive Catalytic Converters

As has been shown from the foregoing sections, to improve catalytic converter performance there is a continued drive towards higher operating temperatures within the converter matrix, coupled with thinner foil components, providing high cell densities, such that high surface area to volume ratios are achieved. Maintaining acceptable component lives is mandatory, even when thinner support geometries and higher operating temperatures are envisaged. The aim is to provide components that outlast the warranted life of the vehicle. Within the USA, this translates to 10 years and/or 100,000 miles for future environmentally friendly vehicles [7]; i.e. 1,100h under high speed conditions (910 °C for an underfloor converter, or 1000 °C for a close coupled converter). It is such high speed operation that will be life limiting from an oxidation/corrosion perspective, since the rate of corrosion increases exponentially with temperature. Thus, any high temperature (high speed) service will be particularly damaging. 4.1

Modelling the Cyclic Oxidation Lives of FeCrAl-RE Foil Materials

Recognition that the lives of thin foil, ferritic steels may be oxidation/corrosion dominated – known as chemical failure – has lead to considerable research effort into the understanding of the oxidation processes and the development of life prediction models to predict such behaviour. Research papers aimed at understanding the oxidation process for this class of alloy are numerous in the literature, as evident by recent research proceedings [3, 8–17], which have dominant sections on the growth of alumina scales on FeCrAl-RE steels. Publications on life models are more limited [4, 17–19] with the recent Dechema Workshop on the “Prediction of Cyclic Oxidation Behaviour” an excellent treatise on current thinking [17].

36 The modelling work of the LEAFA project “Life Extension of Alumina Forming Alloys” [4, 17] provides an excellent forum from which to discuss materials issues relevant to the future development of metal foil, automotive catalytic converters. The model extends earlier modelling work by Quadakkers and co-workers [18, 19] and based on a generic understanding of the growth, aging and failure of an alumina scale formed on high temperature, alumina forming, ferritic alloys [20]. Its development was supported by detailed mechanistic understanding (see references 12, 15–17) and an extensive, ‘full life’ test data set extending out to 20,000 h+ for some materials [21]. It assumes, for the case of FeCrAlbased alloys that the initial weight increase is associated with the formation of an alumina scale. Thus, scale growth leads to a loss of aluminium from the FeCrAl-RE foil. After some long oxidation time (tB/O) the aluminium content within the alloy will decrease below some critical level (CB), such that an alumina scale can no longer form. At this point, less protective oxides and ultimately base metal oxides will form – breakaway oxidation – which marks the corrosion life limit of the component. The model takes the form [4]:

b.

kt B / O Dx

( n -1)

+ (1 - b).( k .t B / O )1/ n = 0.89.

r m (Co - CB ) V . . r ox (1 - CB ) A

(1)

where Hm and Hox are the alloy and oxide densities respectively, V/A is the volume to surface area ratio and allows changes in sample dimensions to be taken into account (this can be a local parameter if the component size is greater than the diffusion path length necessary to ensure a uniform drop in the aluminium content), Co is the aluminium mass fraction in the fabricated alloy and CB is the aluminium mass fraction in the alloy at onset of breakaway. tB/O is the time to breakaway oxidation, k is the oxidation rate constant, n takes values from 2 (parabolic) to 3 (sub-parabolic) and , x is the critical oxide thickness for the onset of spallation. Finally, > is a probabilistic term that incorporates the fractional surface area that may spall during thermal cycling and any variance in the rate of oxidation (k) and its effect on the time to breakaway (tB/O). As spallation is rarely observed for foil components the term > takes values equal to, or close to, zero and component life is dominated by the second term on the left hand side (oxide growth) in equation (1). The first term on the left hand side of the equation accounts for oxide spallation, while the right hand side of the equation is a measure of the available aluminium reservoir, locally in the component (if V/A is the local volume/surface area ratio. This behaviour is illustrated in Figure 4, which provides model predictions (assuming parabolic kinetics) compared with breakaway oxidation lifetimes measured over a range of foil/alloy thicknesses from 70mm to 1.8mm (for rectangular geometry test pieces the V/A ratio is approximately half the sample thickness). The open symbols represent samples that have not yet gone into breakaway, while the filled symbols represent samples that show some evidence of breakaway corrosion, usually at one of the corners. Superimposed on this figure are the median predictions assuming a parabolic rate law and various failure criteria. The line marked ‘InCF’, with > = 0 and CB = 0wt%, corresponds to ‘Intrinsic Chemical Failure’. This is the most likely failure mode for foil material of low strength and results in the longest component lives at any foil thickness (V/A ratio). The lines marked ‘MICF’, with > = 0.1 or 1.0 and CB = 1.7 wt%, correspond to ‘Mechanical Induced Chemical Failure’; under these conditions either local spallation or tensile cracking of the alumina scale occurs, usually in areas of constraint, and the repair, cracking and spallation process rapidly depletes the available aluminium reservoir leading to early failure. The term ‘>’ is a measure of the extent of spallation, > = 1

37 being one limiting case whereby the alumina totally spalls at each shut down cycle. In the LEAFA experimental programme this was never observed to happen for the foil samples, but was approached for strong, thick alloy materials on rapid cooling. >=0.1 1.0E+04

MICF >=1.0

Aluchrom Yhf - 1300C Aluchrom Yhf - 1300C (Not.) Kanthal AF - 1300C

2

kp*tB [ (mm) ]

1.0E+03

Kanthal AF - 1300C (Not.) Kanthal APM - 1300C Kanthal APM - 1300C (Not.) PM2000 - 1300C PM2000 - 1300C (Not.)

1.0E+02

1.0E+01 0.01

C0=5.5 wt.% CB=1.7 wt.% (MICF) CB=0.0 wt.% (InCF)

InCF >=0.0 0.10

1.00

10.00

Volume/Surface Area [mm]

Figure 4: Model of breakaway oxidation, based on parabolic rate law oxidation : the lines super-imposed on the figure correspond to median prediction.

It can be seen from Figure 4 that the parabolic based model predictions prove conservative (predict shorter lives) for all foil materials tested, whether weak (Aluchrom YHf, Kanthal AF) or strong (PM2000). This is because a-alumina growth is generally observed to follow subparabolic kinetics [18,19] with an exponent ‘n’ between 2.3 and 3.0; such behaviour would lower the rate of aluminium consumption and therefore lead to extended component lives. It can further be seen from Figure 4, that > (the propensity to spall) is much less than 0.1 for all foil materials in current usage. This means that the dominant foil failure mechanism is ‘Intrinsic Chemical Failure’ and thus CB (the aluminium content at onset of breakaway) can be expected to reduce to essentially zero. 4.2

The Influence of Component Geometry on Catalytic Converter Body Life

It is evident from the foregoing modelling work that specimen thickness is a significant factor in determining the life of the converter matrix. To be exact it is the local volume/surface area ratio that is critical in determining the onset of chemical failure, that is why most rectangular test samples fail at corners. This aspect has particularly been addressed by Strehl et al [22] where it is shown that material thickness, sample shape and local constraints resulting from component geometry may have significant effects on component life. It is also shown that the local oxygen pressure adjacent to the components surface can be reduced by unfavourable geometries, such as crevices, and this can trigger early breakaway corrosion. Probably the most significant contribution of this study is the recognition that is the ‘local’ volume/surface area ratio that controls component life. A simple analysis for plate material, which is obvious once demonstrated, shows that the volume/surface area ratio at a corner will be one third of that for a free surface, while for an edge the volume/surface area ratio is re-

38 duced by a factor of two. For foil samples, these geometry factors are even more critical; when the foil is thin the aluminium concentration within the foil is in equilibrium with two free surfaces (hence the assumption that for an infinite foil of thickness ‘d’ the volume/surface area ratio is d/2), however, at a corner the local volume/surface area ratio approaches d/4, half of that for the bulk of the sheet, while along edges the value approached d/3. Thus the local aluminium reservoir is significantly reduced at edges and corners because of the change in the local volume/surface area ratio in this region and this accounts for the onset of breakaway oxidation usually being noticed at corners first. This has major implications in manufacturing components and emphasises the need for good design particularly at corners, edges and fixings if premature breakaway is to be avoided. However, the local aluminium reservoir is not the sole controlling parameter; for example two small closely spaced holes could be expected to give rise to a geometry where the local volume/surface area ratio results in premature breakaway, however, this does not occur because additional aluminium can diffuse to this region from the bulk of the component. Thus it is the balance between the local aluminium reservoir, its rate of consumption through oxidation and the rate of supply by diffusion from the bulk of the component that determines whether breakaway will occur or not. Geometry also modifies the constraint that the oxidising surface sees and this can alter the oxidation rate. Thin unconstrained foils are free to expand as a result of growth and thermal stress (length increases of 10% have been measured in oxidation measurements of thin foil [21]). Constrained foils, and thick section components, can generate sufficient stress at surface imperfections that scale cracking, can occur locally [22,23] thus oxidation rates in areas of high constraint may differ from those of unconstrained surfaces. This effect of constraint is illustrated in Figure 5, which plots the mass gain per unit area for various geometry components, including a model catalytic converter body, manufactured from a 58mm Aluchrom YHf foil oxidised at 1000 °C for 700h. The lowest mass gain was observed for the rectangular specimen which was free to deform. Both the ring sample and model catalytic converter body showed an increased mass gain, both of similar magnitude. The additional mass gain was associated with cracks observed in the oxide scale on both the ring samples and catalytic converter body. This cracking has the apparent effect of increasing the rate constant by some 60%, during these discontinuous oxidation studies. In practice, this will significantly reduce component life as component failure switches from intrinsic chemical failure (InCF) to mechanically induced chemical failure (MICF) – see Table 2, and Figures 4 and 9. Table 2 provides predictions of the medium component lives for FeCrAl foil material with 5wt% aluminium, based on the LEAFA model, assuming parabolic kinetics, for the cases of intrinsic chemical failure and mechanically induced chemical failure. In calculating the values for mechanical induced chemical failure it is assumed that the catalytic body is constrained (kp is increased by a factor of 1.6) and the alloy concentration at which alumina can no longer reform (CB) is 1.7wt%.

39

Figure 5: Influence of component geometry on the oxidation of Aluchrom YHf: 58mm foil at 1100 °C [22].

One can see from Table 2 that reducing the foil thickness from 50mm to 30mm reduces the life of the component by a factor of 2.77. However, this reduction is not as great as that imposed by the constraint of the cylindrical geometry, which has the effect of reducing the life by a factor of 3.39 over the unconstrained foil. Table 2: Median predicted component lives for FeCrAl-RE foils as a function of foil thickness and degree of constraint

Component life measured as the kp.tB product [mm2] Foil thickness Unconstrained foil Catalytic converter body 20mm 0.90 0.27 30mm 2.03 0.60 50mm 5.63 1.66 70mm 11.03 3.26 4.3 Changes in Operating Temperature, Brought About by Catalytic Converter Positioning

Current underfloor catalytic converter bodies see peak temperatures of 910 °C, while those for close coupled configurations are expected to see 1000 °C. Assuming that a-alumina is the stable oxide formed on the FeCrAl-RE catalytic converter bodies (it is at more elevated temperatures > 1050 °C), then this temperature increase will raise the oxidation parabolic rate constant by a factor of 5.2, from 2.3 × 10–14 g2 cm–4 s–1 at 910 °C to 1.2 × 10–13 g2 cm–4 s—1 at 1000 °C as can be seen from Figure 6 [24]. In terms of numbers suitable to insert into the above mentioned life model the associated scaling rates are 2.16 × 10–2 mm2 h–1 at 910 °C and 1.13 × 10– 1 mm–2h–1 at 1000 °C. Thus assuming a foil thickness of 50 mm, under unconstrained conditions, then the hot exposure time would be reduced from 260 hours to some 50 hours before chemical

40 failure occurs under peak load conditions. Constraint or thinner foil sections would reduce these lives still further. One conclusion, therefore, is that catalytic converters must spend a considerable period of normal operation way below these peak load temperatures, for under such peak load conditions the foil components would not be able to sustain the warranted lives of the vehicle. To achieve the desired warranted life the mean operating temperature of a current, underfloor, catalytic body would have to 780 °C, based on the above FeCrAl-RE life model. These modelling assumptions are based on the rate of growth of a-alumina scales, however, at operating temperatures in the range 780 °C–910 °C it is more likely that transition aluminas will form. This aspect has been extensively studied. One elegant study by Molins et al [25] on a Ugine Savoie alloy examined the oxidation kinetics over the temperature range 850 °C– 1100 °C in flowing synthetic air. The samples were foils 45mm thick. The results obtained are reproduced in Figure 7, as relative mass gain against time (the norm is taken to be 168h at 950 °C).

Figure 6: Arrhenius plot of the parabolic rate constant as determined from TGA studies, for various FeCrAl-RE alloys, over the temperature range 750-1350 °C [24].

41

Figure 7: Oxidation kinetics for a Fe20Cr5.2Al0.01Ce alloy over the temperature range 850 °C–1100 °C [25]

This research shows that two domains exists, at low temperatures (T<950 °C) faster kinetics are observed, with mass gains greater than that observed at 950 °C. The oxidation rate again increases as the temperature is raised above 950 °C. A further important observation is that at 900 °C a clear two stage process can be seen, commencing with a transient stage before parabolic oxidation is established. Elegant SEM and TEM microscopy shows that at low temperatures, below 950 °C, g-alumina platelets form initially with the rate of g-alumina growth significantly higher than that for a-alumina, later this transforms to a-alumina and a slow parabolic growth is established. At 900 °C, only partial transformation is observed after 60h exposure, although the kinetic curves show a low oxidation rate at this point due to the formation of a-alumina as an interfacial barrier between the FeCrAl alloy and the outer g-alumina microstructure. Above 950 °C, a-alumina predominates. Thus the operation of catalytic converter bodies at temperatures below 950 °C, may result in reduced lives as a result of g-alumina formation on the FeCrAl foil material. g-alumina formation may increase the rate of aluminium consumption by a factor of four, when compared to aalumina growth at a similar temperature [25]. One solution is to preoxidise the catalytic converter body, prior to service at temperatures above 1000 °C. Work by Gobel et al [24] has shown that this has the effect of suppressing transition/base metal oxide formation, even when exposed at 800 °C. A preoxidation in vacuum at 10–7 mbar for as little as 90s seem to be sufficient to ensure a-alumina growth, even at temperatures as low as 800 °C. These observations raise the question of environmental effects and how changes within the combustion environment modify the oxide morphology and growth rate. Preliminary work in this area has commenced [26,27], particularly with regard to the influence of water vapour and its role in stabilising transition alumina’s and this is an area requiring further study.

42 4.4

Modifications to the Alloy Composition

Research into the effect of alloy composition on FeCrAl-RE performance has focussed on three areas and these are discussed below:

4.4.1 Increasing the Aluminium Reservoir by Increasing the Aluminium Content of the Alloy Most commercial FeCrAl-based alloys contain around 5wt%Al and 20wt%Cr, see Table 1. Increasing the aluminium content would obviously provide an extension to the oxidation limited life through an increase in the available aluminium reservoir, and this would be of particular benefit in thin section components, such as foils where the available aluminium reservoir is necessarily limited. To the contrary, it must be considered that high aluminium contents will degrade mechanical properties of the alloys, particularly ductility, and this will make fabrication of foil components more difficult. Studies have shown that increasing the aluminium content increases the lifetime. Experiments by Quadakkers and co-workers [18] have clearly demonstrated this for the ODS alloy, ODM 751, where increasing the aluminium content from 4.5wt% to 6.5wt%, increase the time to breakaway from 4,300 h to 7,200 h at 1200 °C for a 1mm wall thickness; i.e. an increase in life of some 67%. However, work by Klower and Kolb-Telieps [29], in a study of the effects of aluminium and reactive element additions on the oxidation behaviour of thin FeCrAl-RE foils, has shown that increasing the aluminium content does not necessarily increase the time to breakaway oxidation. This is because, as the aluminium content of the FeCrAl-RE foils was increased from 5wt% to 13wt%, the parabolic rate constant was found to increase. This behaviour is depicted in Figure 8.

Figure 8: Effect of aluminium concentration on the parabolic rate constant of FeCrAl-RE foils at 1100 and 1200 °C [29]

For additions between 5wt%–8wt% aluminium, an increase in aluminium content increase the lifetimes to breakaway, although the increase is less than would be expected based on an aluminium reservoir calculation because of the associated increase in the parabolic rate constant with aluminium content. Above 8wt% any benefit is marginal, and additions upto 13wt% may prove detrimental. Calculations have shown that the time to breakaway at 1100 °C is re-

43 duced from 1102 hours to 849 hours as the aluminium content is increased from 8 to 13wt% [30]. These studies have shown that increasing the aluminium content within the foil may be of benefit to the oxidation lifetime of the alloy, but that there is an upper limit of 8wt%. Above this the improvement is marginal and at even higher levels may be detrimental. As increasing the aluminium content is detrimental to the alloys mechanical properties, particularly ductility, research should be undertaken on methods to enrich the aluminium content in prefabricated components. Chemical vapour deposition processes show promise [31,32], but from the above analysis the aluminium content of the foil must be limited to below 8wt%.

4.4.2 The Addition of Multiple Active Elements From Table 1 it can be seen that many, and sometime multiple, active elements have been considered to improve the oxidation behaviour and scale adhesion of alumina forming ferritic steels. Generally the level of addition is less than 0.3wt% metallic addition with the more favoured additions including yttrium, lanthanum, cerium, zirconium, hafnium and possibly titanium at low levels [33]. The property improvement which can be achieved by the addition of reactive elements (RE’s) depends on the type of reactive element used, its amount, distribution and form within the alloy. Numerous papers exists in the literature in this area as evident from the many review papers in reference [8]. Most studies confirm that the dominant mechanisms are tramp element scavenging, such as removing sulphur impurities [34,35] or modifying the transport properties within the oxide scale, due to some modification of the oxide grain boundaries [36]. Both of these mechanisms have been observed, and prevail, for the alumina forming ferritic steels. The questions are therefore, which active element? How much? And in what form? As evident from available commercial alloys (Table 1), research and development favours Y, Ce, La, Zr and Hf and discussion will be limited in the main to these elements. The role of titanium is an interesting one, as should it be considerable as a minor alloying addition, a tramp element (impurity) or an active element? Recent work has shown that small additions of titanium may be beneficial and it may indeed be acting as an active element [33]. In FeCrAl-RE foils reactive elements are commonly added as metallic additions, with additions as low as 0.05wt% providing optimum oxidation behaviour [30]. In fact, the lowest mass gain has been measured for yttrium additions of 0.02wt% [29], however, this alloy showed evidence of spalling after 1000 h at 1100 °C and therefore the slightly higher value is considered optimum. If the yttrium content is increased further to 0.07 or 0.08wt% then internal oxidation is observed at 1100 °C, which is though to be detrimental to the behaviour of foil materials [37]. Hafnium, like yttrium has a beneficial effect on both scale adherence and oxide growth rate [29]. Similarly zirconium may be considered beneficial, when added singly, but in combination with yttrium may be detrimental to thin foil material by increasing the depth of internal oxidation [29,38]. Thus to a Fe-20Cr-5Al alloy containing 0.04wt% Y, the addition of 0.04wt% Zr may be beneficial, however, increasing either the yttrium content or zirconium content above this level can cause extensive internal oxidation, penetrating up to 120mm from each free surface [29]. However, a 0.04wt% addition of hafnium in addition to yttrium is beneficial at temperatures in the range 1100–1350 °C and is the basis of the commercial Aluchrom YHf alloy. Equally 0.10wt% Zr, without yttrium is also beneficial, as demonstrated by the excellent high temperature oxidation behaviour of Kanthal APM. This complexed interplay suggests that

44 there are optimum levels of yttrium, hafnium and zirconium and that these levels may vary with the aluminium content of the alloy. The choice of the optimum combination of active elements within thin foil FeCrAl-RE materials is a clear area for future research as the appropriate selection should reduce oxidation rates and improve alloy adherence without adversely affecting the foil mechanical properties through internal oxidation. Care has to be taken to limit “overdoping” because of the increased oxidation rates and internal oxidation damage that it produces. The role of titanium, a beneficial element or detrimental impurity, is still an open debate. Recent compelling papers by Quadakkers and coworkers [33,39,40] suggest that for small additions it behaves like hafnium and zirconium, but that in combination with other active elements it is easy to overdope the alloy [37,38].

4.4.3 The Interaction of Reactive Elements with Alloy Impurities Minor additions of titanium, zirconium and hafnium have been shown to be beneficial in limiting internal oxidation attack. [29, 37–40]. This is associated with their ability to tie up carbon and nitrogen impurities as tiny carbo-nitrides [39,40]. Other tramp elements, such as phosphorus, vanadium calcium etc and minor additions such as silicon must also be considered. Calcium and vanadium are thought to be benign, not showing a clear effect at concentrations of 5 and 80ppm respectively. Phosphorus equally is benign at 1100 °C and 1200 °C, however at 1300 °C it can drastically decrease the lifetime to chemical failure, associated with a substantial increase in oxide growth rate [40,41]. This detrimental behaviour is associated with phosphorus segregation at oxide grain boundaries and the metal oxide interface [41]. Clearly, these interactions between the reactive elements and minor alloy impurities is an important area of future study. Results strongly indicate that optimum oxidation resistance can be achieved through a combination of multiple active elements. The choice of the right levels, however, can only be made when their interaction with alloy impurities is fully quantified. Thus in future alumina forming ferritic steels, not only the major alloy additions need controlling,but also the combination of active elements and the type, amounts and distributions of tramp elements and other manufacturing alloy impurities.

5

Concluding Remarks

In this paper, pertinent materials issues in the development of future, metal foil, automotive catalytic converters have been addressed by reference to the effect they have on component lifetime. The predicted life of alumina forming ferritic alloys, under oxidising conditions, was modelled based on an aluminium reservoir approach. The reserve of aluminium depends on the original alloy aluminium content (Co), the aluminium level at which breakaway oxidation occurs (CB) and the local volume/surface area ratio (V/A). Thus the aluminium reservoir is given by 0.89

r m (Co - CB ) V . . . r ox (1 - CB ) A

Consumption of aluminium results from the oxidation/corrosion processes, through,

45 · Early scale formation and the growth of a transient oxide. At low temperatures, this may be a transition alumina, for example g-alumina, while at high temperatures a-alumina rapidly forms. For FeCrAl-RE foils 950 °C has been shown to be a critical temperature, below this temperature g-alumina is stable for appreciable times (in excess of 60 h at 900 °C). Since these transition aluminas grow at a faster rate than a-alumina, this can have a significant effect on catalyst life. · Formation of a stable, protective oxide. Ideally this is a-alumina, whether d, q, g, or aalumina form depends on the temperature, environment (i.e. water vapour content) and alloy composition. Spalling is not normally expected to occur for weak FeCrAl-RE foil materials and is usually only observed when the section thickness increases. If spallation occurs the alumina scale fails by mechanically induced chemical failure (MICF) and this can substantially reduce component life, not least because under MICF breakaway triggers when the aluminium level falls below some critical level, circa 1.7wt%Al. For foil materials, stress relaxation processes usually ensure that internal stresses are not sufficient to cause failure and the foil will ultimately fail by intrinsic chemical failure (InCF) where CB approaches zero. However, geometric factors, such as sharp corners, holes, welds and changes in shape can provide stress raisers and may lead to local spallation. · Ultimately, consumption of scale forming elements due to oxidation, leads to non protective conditions. Less stable oxides are formed, coupled with possible internal oxidation and nitridation. This stage is critically dependent on the alloy composition, the type of active elements present and the composition of the gas phase. This internal oxidation rapidly depletes any remaining scale forming elements. · Finally, Breakaway oxidation ensues. Aluchrom Yhf - 1100C

1.0E+06

1.0E+05

2

kp*tB [ (mm) ]

1.0E+04

MICF

Aluchrom Yhf - 1200C

>=0.01 5% 0.1% >=0.1 5% 0.1% >=1.0 5% 0.1%

Aluchrom Yhf - 1300C

1.0E+03

Aluchrom Yhf - 1300C Kanthal AF - 1300C Kanthal AF - 1300C Kanthal AF - 1200C Kanthal APM - 1300C Kanthal APM - 1300C Kanthal APM - 1200C

1.0E+02

PM2000 - 1300C PM2000 - 1300C

1.0E+01 InCF > = 0.0

1.0E+00 0.001

C0=5.5 wt.% CB=1.7 wt.% (MICF) CB=0.0 wt.% (InCF)

0.010

0.100

Volume/Surface Area [mm]

1.000

10.000

PM2000 - 1200C PM2000 - 1200C (Improve) Aluchrom Yhf 1200C (Iso) Kanthal APM 1200C (Iso) PM2000 1200C (Iso)

Figure 9: Stochastic life prediction model for the chemical failure of FeCrAl-RE, alumina forming alloys [4]

46 Thus, chemical failure reflects a balance between the available aluminium reservoir and its rate of consumption due to a complex interplay of oxidation parameters. Key among these is the oxide rate constant, which may follow parabolic or sub-parabolic kinetics, and for thicker sectioned components, or highly constrained geometries, the critical oxide thickness to spall. It has been proposed that the (k × tB) product is a temperature independent parameter, that defines the lifetimes of an alumina forming ferritic steels, whether a foil or sheet materials [4]. Figure 9 presents a plot of (kp × tB) against V/A ratio for a range of alumina forming ferritic steels over the temperature range 1050 °C–1400 °C, confirm the hypothesis that the (kp × tB) product may be used to provide a temperature independent estimate of component life. In Figure 9 alloy lives between a few tens of hours and 20,000 h are plotted, superimposed on the plot are statistical corrosion models, that define the risk of failure [4], two levels of risk are plotted: a 5% chance of failure and a 0.1% chance of failure assuming oxidation follows parabolic kinetics. All foil samples were observed to fail by intrinsic chemical failure (InCF), when unconstrained. A life model, based on intrinsic chemical failure (> = 0.0; CB = 0.0wt%) and parabolic kinetics, provides a conservative estimate of foil component life for a range of FeCrAl-RE materials over the temperature range 1100 °C–1400 °C.

6

References

[1] World Energy Outlook, 2000, International Energy Agency. [2] W. Maus ‘Mobility, Prosperity and Environment Protection – the Catalytic Converter is Indispensable’, in “Metal-Supported Automotive Catalytic Converters” (ed. H. Bode) p3– 13, Werkstoff-Informationsgesellschaft, Frankfurt, Germany (1997). [3] H. Bode (ed) “Metal-Supported Automotive Catalytic Converters” (ed. H. Bode) p3–13, Werkstoff-Informationsgesellschaft, Frankfurt, Germany (1997). [4] J. R. Nicholls, R. Newton, M. J. Bennett, H. E. Evans, H. Al-Badairy, G. Tatlock, D. Naumenko, W. J. Quadakkers, G. Strehl and G. Borchardt, ‘Development of a Life Prediction Model for the Chemical Failure of FeCrAlRE Alloys in Oxidising Environments, “Life Modelling of High Temperature Corrosion Processes’, (eds M. Schutze, W. JH. Quadakkers and J. R. Nicholls) EFC, Publication 28, IoM Communications 2001. [5] H. Bodes ‘Development Status of Materials for Metal Supported Automotive Catalysts’, in “Metal-Supported Automotive Catalytic Converters” (ed. H. Bode) p3–13, WerkstoffInformationsgesellschaft, Frankfurt, Germany (1997). [6] B. H. Engler “Katalysatoren fur den Umweltschutz”, Chem-Ing.Tech. 63, 298–312 (1991); cited in reference 5. [7] T. Nagel and W. Maus “Development of More Exacting Test Conditions for close Coupled Converter Applications”, in Metal-Supported Automotive Catalytic Converters” (ed. H. Bode) p107–126, Werkstoff-Informationsgesellschaft, Frankfurt, Germany (1997). [8] E. Lang (ed) “The Role of Active Elements in the Oxidation Behaviour of High Temperature Materials and Alloys” Elsevier Applied Science, London (1989). [9] M. J. Bennett and G. W. Lorimer (eds), “Microscopy of Oxidation”, Institute of Metals, London (1991). [10] S. B. Newcomb and M. J. Bennett (eds), “Microscopy of Oxidation-2” Institute of Materials, London, 1993.

47 [11] D. Coutsouradis et al (eds) Proc. “Materials for Advanced Power Engineering”, COST 501, Kluwer Academic Publishers, Dordrecht, Netherlands (1994). [12] S. B. Newcomb and J. A. Little (eds) “Microscopy of Oxidation-3”, Institute of Materials, London, 1997. [13] D. A. Shore, R. A. Rapp and P. Y. Hou (eds), Int. Conf. on “Fundamental Aspects of High Temperature Corrosion”, The Electrochemical Soc., USA, (1997). [14] J. Lecomte-Beckers et al (eds), Proc. “Materials for Advanced Power Engineering 1998”, Forschungszentrum, Julich, Germany (1998). [15] M. Schutze and W. J. Quadakkers (eds) “Cyclic Oxidation of High Temperature Materials”, EFC Publication 27, IoM Communications, London, (1999). [16] G. Tatlock and S. Newcomb (eds), Special Issue of Material at High Temperatures on “Microscopy of Oxidation-4” Vol. 17(1) (2000). [17] M. Schutze, W.J. Quadakkers and J. R. Nicholls (eds) “Life-time Modelling of High Temperature Corrosion Processes”, EFC Publication 28, IoM Communications, London (2001). [18] W. J. Quadakkers and K. Bongartz, Werst. U. Korros. 24, 232, (1994). [19] W. J. Quadakkers, K. Bongartz and F. Schutbert in Proc. “Materials for Advanced Power Engineering”, COST 501, (eds D. Coutsouradis et al) part II, p1533, Kluwer Academic Publishers (1994). [20] J. R. Nicholls and M. J. Bennett, “Cyclic Oxidation-guidelines for Test Standardisation aimed at the Assessment of Service Behaviour”, European Federation of Corrosion Publications Vol. 27, pp437–470 (1999). [21] R. Newton, M. J. Bennett, J. P. Wilber, J. R Nicholls, D. Naumenko, W. J. Quadakkers, H. Al-Badiary, G. Tatlock, G. Strehl, G. Borchardt, A. Kolb-Telieps, B. Jonsson, A. Westerlund, V. Guttmann, M. Maier and P. Beaven, “The Oxidation Lifetime of Commercial FeCrAlRE Alloys” in ‘Life Modelling of High Temperature Corrosion Processes’, (eds. M. Schutze, W. J. Quadakkers and J. R. Nicholls) EFC Publication 28, IoM Communications 2001. [22] G. Strehl, V. Guttmann, D. Naumenko, A. Kolb-Telieps, G. Borchardt, J. Klower, P. Beaven, J. R.Nicholls, “The Influences of Sample Geometry on the Oxidation and Chemical Failure of FeCrAl(RE) Alloys”, Life Modelling of High Temperature Corrosion Processes’, (eds M. Schutze, W. J. Quadakkers and J. R. Nicholls), EFC Publication 28, IoM Communications, 2001. [23] H. Al-Badairy, G. J. Tatlock and J. LeCoze, “An Auger Study of Thermally Spalled Oxides on Fe-20Cr-5Al Based Alloys” in “Microscopy of Oxidation-3” (eds S. Newcomb and J. A. Little) p105, Institute of Materials, London (1997). [24] M. Gobel, J. Schimmelpfennig, A. Glazkow, G. Borchardt, “Growth of a-alumina Scales on Fe-Cr-Al Alloys” in Metal-Supported Automotive Catalytic Converters” (ed. H. Bode) p191, Werkstoff-Informationsgesellschaft, Frankfurt, Germany (1997). [25] R. Molins, A. Germidis and E. Andrieu “Oxidation of Thin FeCrAl Strips: Kinetic and Microstructural Studies in “Microscopy of Oxidation-3”, (ed. S. B. Newcomb and J. A. Little, p3, Institute of Materials, London (1997). [26] A. Kolb-Telieps, U. Miller, H. Al-Badairy, G. Borchardt, G. Tatlock, D. Naumenko, W. J. Quadakkers, G. Strehl, R. Newton, J. R. Nicholls and V. Guttmann, “The Role of Bioxidant Corrodants on the Life Time Behaviour of FeCrAlRE Alloys” ‘Life Modelling of

48 High Temperature Corrosion Processes’, (eds M. Schutze, W. J. Quadakkers and J. R. Nicholls), EFC Publication 28, IoM Communications, 2001. [27] G. J. Tatlock and H. Al-Badairy “The Oxidation of Thin Foils of FeCrAl-RE Alloys in Moist Air”, accepted for publication in Materials at High Temperatures (2001). [28] M. J. Bennett, R. Perkins, J. B. Price and F. Starr in Proc. “Materials for Advanced Power Engineering”, COST 501 (eds D. Coutsouradis et al), p1553, Kluwer Academic Publishers (1994). [29] J. Klower and A. Kolb-Telieps, “Effect of Aluminium and Reactive Elements on the Oxidation Behaviour of Thin FeCrAl Foils in “Metal-Supported Automotive Catalytic Converters” (eds H. Bode), p33, Werkstoff-Informationsgesellschaft, Frankfurt, Germany, (1997). [30] J. Klower, Materials and Corrosion, 49, 758 (1998). [31] A. B. Smith, A. Kempster and J. Smith, “Characterisation of Aluminide Coatings formed on Nickel Based Superalloys by Vapour Aluminising” in “High Temperature Surface Engineering” (eds J. R. Nicholls and D. S. Rickerby) p13, IoM Communications, London (2000). [32] L. Vandenbulcke, G. Leprince and B. Nciri, “Low Pressure Gas-Phase Pack Cementation Coating of Complex-Shaped Alloy Surfaces”, Materials Science and Engineering, A121, 379 ff, (1989). [33] W. J. Quadakkers, T. Malkow, H. Nickel and a. Czyrska-Filemonowics in Proc. 2nd Int. Conf on “Heat Resisting Materials”, Gatlinburg USA, p91, ASM International, Ohio, USA (1999). [34] J. G. Smeggil, “Some Contents on the Role of Y in Protective Oxide Scale Adherence”, Materials Science and Eng. 87, 261 (1987). [35] J. L. Smialek, “Sulphur Impurities and the Microstructure of Alumina Scales” in “Microscopy of Oxidation-3” (eds S. B. Newcomb and J. A. Little), p127, Institute of Materials, London, (1997). [36] B. Pint, Oxid. Met. 45, 1 (1996). [37] J. Klower and J. G. Li, Materials and Corrosion 47, 545 (1996). [38] P. A. Van Manen, E. W. A. Young, D. Schlakoord, C. J. an der Wekken and J. H. W. de Wit, Surface and Interface Analysis 12, 391 (1988). [39] W. J. Quadakkers and L. Singheiser, “Practical Aspects of the Reactive Element Effect”, in “High Temperature Corrosion”, Les Embiez, France, (May 2000). [40] D. Naumenko, W. J. Quadakkers, P. Beaven, H. Al-Badairy, G. Tatlock, R. Newton, J. R. Nicholls, G. Strehl, G. Borchardt, J. Le Coze, B. Jonsson, A. Westerlund, “Critical Role of Minor Elemental Constituents on the Lifetime Oxidation Behaviour of FeCrAl-RE Alloys”, “Life Modelling of High Temperature Corrosion Processes”, (Eds. M. Schutze, W. J. Quadakkers and J. R. Nicholls) EFC Publication 28, IoM Communications 2001. [41] H. Al-Badairy, G. Tatlock, H. E. Evans, G. Strehl, G. Borchardt, R. Newton, M. J. Bennett, J. R. Nicholls, D. Naumenko and W. J. Quadakkers, “Mechanistic Understanding of Chemical Failure for FeCrAl-RE Alloys in Oxidising Environments” in “Lifetime Modelling of High Temperature Corrosion Processes”, (eds M. Schutze, W. J. Quadakkers and J. R. Nicholls) EFC publication 28, IoM Communications, London (2001).

High Temperature Corrosion of FeCrAlY/Aluchrom YHf in Environments Relevant to Exhaust Gas Systems Angelika Kolb-Telieps1), Gernot Strehl2) , Dmitry Naumenko3), Willem. J. Quadakkers3) , Rachel Newton4) 1)

Krupp VDM GmbH, Kleffstr. 23, 58762 Altena, Germany TU Clausthal, Institut f³r Allgemeine Metallurgie, Robert-Koch-Str. 42, 38678 Clausthal-Zellerfeld, Germany 3) Forschungszentrum Juelich, IWV-2, 52425 J³lich, Germany 4) Cranfield University, SIMS, Cranfield, Bedfordshire, MK43 0AL, UK 2)

1

Abstract

Fe-20Cr-5.5Al-Y/Aluchrom YHf foils were exposed in air, N2 + 5000vppm NO, a simulated fuel-rich exhaust gas, air + SO2 and air + 50vppm HCl at 1200°C. N2 + NO and the exhaust gas act as shielding gases, a behaviour which probably can be explained by the higher oxygen partial pressure in air compared to that of the mixed gas. Air + 0.3% SO2 leads to earlier breakaway, which is supposed to be induced by internal sulphidation. Air + 50vppm HCl seems to result in the formation of volatile species and active oxidation.

2

Introduction

Assessments of different drive concepts with regard to emissions and weight/cost ratios show that the spark-ignition engine will remain the preferred drive concept during the next years [1]. The necessity of further improving the cold-start efficiency, which is highly dependent on the cell density of the catalyst substrate, leads to the objective to provide the largest possible catalytic surface with the lowest possible heat capacity, as can be seen in Figure 1.

Figure 1: Dependence of cold-start factor (catalytic surface/heat capacity) on cell density and foil thickness [1]

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

50 Increasing the cell density also means to reduce the thickness of the foil of the catalyst substrate, for which iron chromium aluminium alloys with additions of reactive elements proved to be an excellent solution. The chemical composition of the foil has been optimised with respect to its oxidation resistance, especially with respect to thickness [2]. However, for a thickness of 50 Ám or less environmental influences have to be considered in more detail than for the thicker foils. Therefore this paper will compare results gained in air to data obtained in atmospheres relevant to exhaust gas systems.

3

Experimental

The tests were performed on Aluchrom YHf, an FeCrAl alloy with additions of reactive elements. The chemical composition is given in Table 1. 0.05mm × 20mm × 10mm coupons were cut from the foil and exposed to multicomponent corrodants at 1200 °C. Table 1: Chemical compositions (mass%) element Cr (wt.-%) Al (Wt.-%) Y (ppm) C (ppm) S (ppm) O (ppm) N (ppm) P (ppm) Zr (ppm) V (ppm) Ti (ppm) Cu (ppm) Ca (ppm) Hf (ppm) Mn (ppm) Si (ppm) Nb (ppm) Mg (ppm) Mo (ppm)

Aluchrom YHf 19.7 5.5 460 2100 1.3 < 10 40 130 540 860 98 110 12 310 1800 2900 < 50 78 100

The atmospheres (in vol.% res. vppm) were the following: 1. N2 + 5000 vppm NO, which simulated the major NOx components in exhaust gas. The only important NOx compound NO was chosen, since engine temperatures reach values from 900 °C to 1300 °C, where the equilibrium partial pressure of NO is most relevant at levels between 500 ppm – 3000 ppm, as can be seen in Figure 2. The cycle time was 20 hours. 2. A simulated fuel-rich exhaust gas (N2 + 12% CO2 + 2% CO + 10% H2O). The oxygen partial pressure at this temperature was calculated to be about 10 Pa–15 Pa. The cycle time was again 20 hours.

51 3. Foils were exposed in air + 0.3 vol.% SO2. Although a lower SO2 concentration, e.g. of 10 vppm, is more relevant in automotive exhaust systems, air + 0.3 % SO2 was chosen as worst case condition. Figure 3 shows the equilibrium partial pressures for both concentrations. From these calculations an influence of the SO2 enrichment at 1200 °C should not be visible for the 10 vppm. The cycle time was 20 hours. 4. 0.05mm and 1 mm thick samples were discontinuously exposed in air + 50 vppm HCl with 100 hours cycles.

Figure 2: Equilibrium composition of air with 5000 ppm NO in the temperature range 0 °C to 1400 °C, calculated with ChemSage [3]

a) Air + 10 vppm SO2

b) Air + 0.3 vol.% SO2

Figure 3: Equilibrium composition of air with (a) 10 vppm resp. (b) 0.3 vol.% SO2 in the temperature range 0 °C to 1400 °C at 1.013 bar, calculated with FactSage

After every cycle the furnace was cooled down to room temperature and the mass changes of the species were measured. No differences between net and gross mass changes were found in

52 N2 + 5000 vppm NO, N2 + 12% CO2 + 2% CO + 10% H2O and air + SO2. In air + 50 vppm HCl the gross mass change could not be determined due to volatile species. The surface coloration was evaluated and a grey colour was associated with (-alumina, a green colour with chromia and red and black with iron oxides. The microstructure has been characterised by optical and scanning electron microscopy. The test results were compared to those gained in parallel tests performed in laboratory air.

4

Results

4.1

N2+5000 vppm NO

As shown in Figure 4, the oxidation rate in the N2-NO mixture is always smaller than in air but shows a similar development. This is also true for the coloration changing from metallic to grey and then to green before breakaway occurs. No evidence of internal nitridation has emerged.

Figure 4: Comparison of the net mass gain of 50 mm foils after cyclic oxidation in N2 with 5000 vppm NO and in laboratory air at 1200 °C [3]

4.2

Fuel-rich exhaust gas (N2 + 12% CO2 + 2% CO + 10% H2O)

Specimen mass gains are shown in figure 5. The mass increase of the foil was slower in the fuel-rich exhaust gas (N2 + 12% CO2 + 2% CO + 10% H2O) than in air.

53

Figure 5: Specimen mass change in artificial exhaust gas at 1200 °C [3]

4.3

Air + 0.3 vol.% SO2

Figure 6 indicates little effect on growth kinetics when adding 0.3 vol.% SO2 to the air environment. But the onset of breakaway at 1200 °C occurred after 150 hours in air in comparison to 120 hours in the sulphur dioxide containing environment. Breakaway occurred with the formation of chromia underlying the alumina scale, and later non-protective iron oxide formation, which resulted in a rapid increase of mass gain.

Figure 6: Comparison between the gross mass gain in air + 0.3% SO2 with that in laboratory air at 1200 °C [3]

4.4

Air + 50 vppm HCl

As can be seen from Figure 7a, the net mass gain at 1200 °C is lower in the HCl containing gas than in air. In contrast to the tests performed in laboratory air, HCl induces spallation in all samples (see Figure 7b). Small red and black spots can be recognised on the surfaces after relatively short exposure times in the HCl containing gas.

54

a)

b)

Figure 7: a) Net mass gain in air and air+50 vppm HCl at 1200 °C [3]; b) cross section of 1 mm thick Aluchrom YHf after 100 hours

5

Discussion

The nitrogen-oxygen-bioxidant and the synthetic exhaust gas retard the breakaway in comparison to exposures in air. In these cases the alumina scale, which could grow due to sufficiently high oxygen levels, protects the alloys against nitridation [4]. So these environments actually can be used as shielding gases for FeCrAlRE alloys. Similar results in exhaust gas have been found for thicker species by Sigler [5] and have been attributed to different oxide morphologies. The oxidation mechanisms are the same in these environments and in air. The formation and almost parabolic growth of alumina is followed by that of chromia and subsequently iron oxides, which then leads to chemical breakaway failure. The experiments show that the growth rates of the aluminium oxide in the two low- pO2 test gases (N2/NOx and exhaust gas) are substantially slower than in air. Two-stage oxidation studies using 18O-Tracer [7] have shown that gas tight alumina scales on FeCrAl alloys containing reactive elements grow by oxygen grain boundary diffusion. For scales exhibiting this growth mechanism, the relation between scale thickness x and oxidation time t is given by:

x2 = k p × t

(1)

The parabolic rate constant kp has been calculated in /8,9/ to be

kp =

4@ DB æ DmO2 ö ç ÷ r è RT ø

(2)

55 with DB the oxygen grain boundary diffusion coefficient along grain boundaries with a width @and a grain diameter of r , DmO2 the chemical potential gradient across the oxide scale, R the gas constant and T the temperature. Previous experiments [7,8] have shown that in an alumina scale with optimum protective properties the grains develop a columnar type structure whereby the grain size tends to increase in growth direction. Thus, the grain size at the scale-metal interface increases with increasing scale thickness, i.e. in the above equation r = r(x). Based on theoretical considerations and experimental observations it was previously shown [7,8] that this increase in grain size can with reasonable accuracy be described by the relation r = r0 × x p

(3)

whereby p is close to one and r0 is the initial grain size, when diffusion along grain boundaries becomes the rate determining process. The scale thickness can easily be transformed into æ Dm ö mass gain ç ÷ using a factor f, because this is the common way to measure oxidation kiè A ø netics. x = 0.53374

mm Dm Dm = f g / m2 A A

(4)

The equations (1) to (4) can now be combined to achieve a comprehensive understanding of the oxidation process.

x2 = k p × t x2 =

4d DB æ DmO2 ç r è RT

Û ö æ DmO2 ÷×t = ç ø è RT

ö 4d DB ÷ ×t × p ø r0 × x

æ 4d DB ö x2+ p = ç ÷ DmO2 t è r0 RT ø

Û

1/ 2 + p

x= f

Dm æ 4d DB ö =ç ÷ A è r0 RT ø

Using the definitions

Dm 1/ n =  kct  . A

Û

 Dm t 

1/ 2 + p

O2

DmO2 4d DB 1 1 = the growth law simplifies to: and kc = f n r0 RT n 2+ p (5)

For the tested foils p equals one. This reveals a cubic time dependence for the scale thickening and thus a scale thickness dependence on the oxygen partial pressure

x : DmO1/2 2 + p = DmO1/23

(6)

The chemical potential gradient across the scale is proportional to the difference between the logarithm of the oxygen activities at the scale/alloy and the scale/oxide interface, respectively.

56 DmO2 = RT  ln aO2 |surface - ln aO2 |int erface 

(7)

The activity of oxygen at the surface is given by the oxygen partial pressure in the surrounding atmosphere. The oxygen activity at the scale/alloy interface has been calculated from thermodynamic data [10] and activity coefficients measured in the LEAFA [3] project to be 2.305 × 10–25 at 1200 °C. Table 2 gives the oxygen partial pressures for the tested atmospheres and the expected retardation in the oxide growth. Table 2: Oxygen partial pressures for the tested atmospheres and the expected retardation in the oxide growth

atmosphere air N2-NO exhaust gas

aO2 0.21 2.5 × 10–3 1 × 10–10

,mO2 / kJmol–1 675.5 621.6 412.8

retardation 1 0.9 0.6

The retardation factor is calculated assuming that breakaway oxidation in air and in the reduced oxygen partial pressure atmosphere occurs for the same mass gain. 1/ n Dm air air 1/ n  kc tB  =  kcatmospheretBatmosphere  Û kcairtBair = kcatmospheretBatmosphere A

retardation =

t Bair t Batmosphere

(8)

atmosphere

=

kcatmosphere DmO2 = kcair DmO2

air

(9)

In Table 3 the calculated data are compared to the measured data. They are the same for the exhaust gas but differ for N2+NO. However, this difference might be explained by the cycle time. The data gained in air were measured with 20 hours cycles whereas those measured in the bioxidant refer to 100 hours cycles. Table 3: Comparison of calculated and measured retardation factors after 100 and 150 hours exposure

atmosphere N2-NO exhaust gas

calculated retardation 0.9 0.6

measured retardation after 100 and 150 hours 0.8 0.6

Results gained for PM 2000, a Plansee FeCrAl alloy, support this idea. Isothermal thermogravimetric tests were performed in Ar + 20% O2 and in Ar + 4% H2 + 7% H2O at 1200 °C. Figure 8 shows the retarded gross mass gain in Ar + 4% H2 + 7% H2O, the atmosphere with the lower oxygen partial pressure. The exponential fits of these data resulted in k values of 0.31 mg2cm4/h for Ar + O2 and 0.19 mg2cm4/h for Ar + H2 + H2O. This method which allows to predict results from retardation factors is not applicable for atmosphere where volatile species or inclusions occur, like in air + HCl or air + SO2. A similar oxidation mechanism is found in air and in air + SO2, but 0.3% SO2 induce earlier breakaway. The SO2 bearing environment affects the oxidation behaviour dramatically by internal sulphidation. The sulphidation can also be explained by the oxygen partial pressures of

57 SO2 and SO2 at 1200 °C, which are shown in Figure 3. Since SO2 contents more realistic in automotive exhaust systems are in the range of 10 vppm, the conclusion is that SO2 has not to be considered for oxidation processes in vehicles with spark-ignition engines.

Figure 8: Gross mass gain of PM 2000 in atmospheres with different oxgen partial pressures gained at 1200 °C

In the above mentioned environments the species fail due to chemical breakaway when most of the aluminium is consumed and other oxides types form. Therefore for the species with 5.5 mass% Al, the foil thickness is critical. Spallation was only noticed on thicker coupons. However, the growth of alumina in the HCl containing atmosphere is not parabolic. Even for the foils the alumina scale spalled. The mechanistic understanding for the behaviour of species in HCl containing atmospheres is based on the observation, that an influence is visible before the first cracking or spalling of the alumina scale occurs. Furthermore very fine spall was found early on and also the scale adherence in some species seems to be weaker. Red and black spots on top of the alumina indicate that iron is involved in the process, but not chromium, because green spots are missing. Together with the oxygen HCl can react at 1200 °C to H2O and Cl2. Similar reactions have already been found at 600 °C [6]. This might promote the formation of volatile chlorides. These compounds deteriorate the alumina scale, which in turn allows access of the gaseous species O2, HCl, H2O and Cl2 to the metal oxide interface. Thermodynamical calculations revealed that the chlorides with the highest fugacity at 1200 °C are aluminium chlorides. These ideas are confirmed by the observation that the processes accelerate as soon as the first cracks appear in the alumina.

6

Conclusions

For wrought Aluchrom YHf N2 + 5000 vppm NO and N2 + 12% CO2 + 2% CO + 10% H2O act as shielding gases at 1200 °C. An explanation for this behaviour is thought to be the higher oxygen partial pressure in air compared to that of the mixed gas.

58 Air + 0.3 vol.% SO2 leads to earlier breakaway due to sulphidation. Air + 50 vppm HCl seems to result in the formation of volatile aluminium chlorides and active oxidation at 1200 °C.

7

Acknowledgements

We are grateful to the European Commission for financial support under the LEAFA project no. BRPR-CT97-0562 and to our partners for the supply of the alloys tested, for the chemical analysis of alloys and for their contribution to scientific input in discussing these results. We also want to thank Dr. K. Hack, Gesellschaft für Technische Thermochemie und -physik, Herzogenrath, for the calculations of figure 3.

8 [1] [2]

References

W. Maus, R. Brück, G. Holy, Int. Congress in Graz, Sept. 2–3, 1999 J. Klüwer, A. Kolb-Telieps, M. Brede, Int. Conf. MACC ’97 in Wuppertal, Oct. 27–28, 1997 [3] BRITE-EURAM project, Contract no. BRPR-CT97-0562 [4] M.J. Bennett, R. Newton, J.R. Nicholls, Eurocorr 2000, available as CD [5] D.R. Sigler, Oxidation of Metals, Vol. 40, No.5/6, 1991, p. 555–583 [6] A. Zahs, M. Spiegel, H.J. Grabke, Materials and Corrosion 50, 1999, p. 561–578 [7] W.J. Quadakkers, H. Holzbrecher, K.G. Briefs, H. Beske, Oxidation of Metals 32, (1/2) (1989), 67–88 [8] K. Bongartz, W.J. Quadakkers, J.P. Pfeifer, J.S. Becker, Surface Science 292 (1993) 196–208 [9] S.N. Basu, J.W. Halloran, Oxidation of Metals 27 (1987) 143 [10] Ihsan Barin, Thermochemical Data of Pure Substances, Weinheim, Basel, Cambridge, New York, VCH, 1989

Improved High Temperature Oxidation Resistance of REM Added Fe-20%Cr-5%Al Alloy by Pre-Annealing Treatment K. Fukuda, K. Takao, T. Hoshi and O. Furukimi Technical Research Laboratories, Kawasaki Steel Corp., Chiba (Japan)

1

Introduction

Fe-Cr-Al alloys exhibit outstanding oxidation resistance at high temperatures because a protective α-Al2O3 scale forms on their surface. As a practical application of this good oxidation resistance, Fe-20mass%Cr-5mass%Al alloy foils have recently been used as a catalytic converter substrate for automobiles, in which the material is exposed to high temperature exhaust gas. (1)-(3) The thickness of these foils is usually limited to 30 µm–50 µm so as not to increase the back pressure in the exhaust system. Since the surface-to-volume ratio of foils is higher than that of thicker alloy sheets, the Al in the foils is consumed as Al2O3 in a shorter time than in sheets. Therefore, reducing the growth rate of Al2O3 scale is important for this application. It is widely known that the addition of reactive elements such as rare earth metals (REM), Ti, Zr, and Hf to these alloys improves their oxidation resistance by preventing spalling of the Al2O3 scale.(4)-(7) Considered these studies, Fe-20Cr-5Al-La-Zr alloy foils have been widely used for catalytic converter substrates. Foils for this application are usually supplied as-cold rolled or after pre-annealing in a reducing atmosphere. However, the effect of pre-annealing conditions on the growth rate of Al2O3 scale formed on Fe-20Cr-5Al with the addition of reactive elements such as rare earth metals (REM), Zr, and Hf alloys had not been clarified. Therefore, in this study, the effect of pre-annealing in a hydrogen atmosphere on improvement of high temperature oxidation resistance in Fe-20mass%Cr-5mass%Al alloy foils with a small content of La-Zr or La-Hf was investigated.

2

Experimental Procedure

2.1

Specimen Preparation

The chemical compositions of the alloys used in this study are shown in Table 1. The basic composition was Fe-20mass%Cr-5mass%Al with a small addition of La, La-Zr, or La-Hf. All the alloys were melted in a vacuum induction furnace and cast as 10kg ingots, and were then hot-rolled to 3mm thick plate. These plates were annealing and cold-rolled to foils with a thickness of 50 mm or 300 mm. Some of the foils were annealed at 1223K and polished with #600 SiC paper, and some of the polished foils were also pre-annealed in hydrogen gas or air at 1223K for 60 seconds. These foil specimens were cut into coupons with a size of 20–30mm, which were degreased in acetone and in alcohol before oxidation. Alloy specimens with a

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

60 thickness of 300 mm were used for observation by scanning electron microscope (SEM) and transmission electron microscope (TEM). Table 1: Chemical compositions of specimens Alloy A B C D E F 2.2

Cr (mass%) 19.8 20.0 20.1 20.1 20.1 20.6

Al (mass%) 5.5 5.7 5.6 5.6 5.7 5.6

La (mass%) 0.024 0.045 0.082 0.098 0.088 0.072

Zr (mass%)

Hf (mass%)

0.029 0.032

Oxidation tests

Oxidation tests were carried out in air at 1373 K or 1423 K using an electric furnace. The mass change was measured by weighing the specimens at certain time intervals after cooling to room temperature. 2.3

Analysis method

The surfaces of the foils after polishing, pre-annealing in hydrogen, and oxidation were investigated by secondary ion mass spectroscopy (SIMS). Mass filtered O2+ primary ions (acceleration voltage: 15 kV) were rastered over areas of 100 × 100, 150 × 150, and 200 × 200 µm on the targets. Specimens of the 300 µm thick sheets of the alloys after continuous oxidation at 1373 K for 172.8 ks were cracked by immersion in liquid nitrogen to observe the cross-section of the scale by SEM. The Al2O3 scale which formed on the specimen foils after continuous oxidation at 1373 K for 172.8 ks was examined with a field emission type TEM equipped with an energy-dispersive X-ray spectrometer (EDX). Specimens of the foils were prepared by mechanical grinding of the alloy substrate and subsequent ion-milling with an Ar gun, aiming at the mid-thickness region of the scale. Specimens of the 300 mm thick sheets of the alloys were prepared by a focused ion beam (FIB) system using Ga-ions to observe the cross-section of the scale.

3

Results and Discussion

3.1

Effect of pre-annealing on oxidation rate

Figure 1 shows the mass change curves of 300 µm thick foil samples of alloy C during oxidation at 1423 K. No spalling of the scale was observed with any of the specimens, and a protective Al2O3 layer formed on all the specimens after the oxidation test. The mass gain results were virtually the same with the as-cold rolled, polished, and pre-annealed in air specimens. The mass gain of the specimen which was pre-annealed in hydrogen gas was lower than that of the other specimens. Figure 2 shows the oxidation gain after oxidation in air at 1423K for 86.4ks with 300 µm thick foil samples of alloys A, B, C, and D with several La contents in the as-cold rolled con-

61 dition and with pre-annealing in hydrogen. The oxidation gain of the as-cold rolled specimens decreased as the La content increased up to about 0.04 wt%, and remained approximately constant 5.0 (g/m2) above about 0.04 wt%. In contrast, the oxidation gain of the samples which were pre-annealed in hydrogen continued to decrease as the La content increased even at values of more than 0.04 wt%. This demonstrates that pre-annealing treatment in hydrogen gas is more effective for decreasing the oxidation rate as the La content increases.

Mass gain (g/m2)

20.0

15.0

as rolled annealed in air polished after annealing annealed in hydrogen

10.0

5.0

0 0

200

400

600

800

Oxidation time (ks) Figure 1: Effect of annealing treatment on oxidation behavior of 300 µm thick sheets of alloy C at 1423K in air

Mass gain (g/m2)

10 as rolled annealed in hydrogen

8 6 4 2 0 0.00

0.02

0.04

0.06

0.08

0.10

La content (mass%) Figure 2: Effect of La content and hydrogen annealing on mass gain after oxidation in air at 1423K for 86.4ks

Figure 3 shows the mass change curves during oxidation at 1373K with 50 µm thick foil samples of alloys C, E, and F in the as-cold rolled condition and with pre-annealing in hydrogen. The oxidation gain results of the as-cold rolled alloys E (La + Zr) and F (La + Hf) were smaller than that of alloy C, which contained only La. The oxidation gain of these La + Zr and La + Hf co-added alloys was also smaller when the samples were given pre-annealing treatment in hydrogen.

62

5

Alloy C as rolled Alloy E as rolled Alloy F as rolled Alloy C annealed in hydrogen Alloy E annealed in hydrogen Alloy F annealed in hydrogen

100 (b) Mass gain (g/m2)

10

(a)

2

Mass gain (g/m2)

15

Alloy C as rolled Alloy E as rolled Alloy F as rolled Alloy C annealed in hydrogen Alloy E annealed in hydrogen Alloy F annealed in hydrogen

10

0.4 0 0

200

400 Time(ks)

600

800

1 10

100 Time (ks)

1000

Figure 3: Effect of hydrogen pre-annealing annealing on oxidation behavior of 50 µm thick foils of alloy C, E, and F at 1373K in air.

The mass change curves of the 50 mm thick alloys during oxidation at 1373 K are expressed in a log-log plot in Figure 3(b) in order to understand the kinetics. The slope of each sample is about 0.4. This means that the oxidation curve basically conforms to the parabolic rate law. 3.2

SIMS analysis

Figure 4 shows the depth profiles of 50 µm thick foils of alloy E in the as-cold rolled condition and with pre-annealing in hydrogen as obtained by SIMS. No apparent peak was detected near the surface of the as-cold rolled foil. However, with the pre-annealed foil, the intensities of Al, La, and Zr secondary ions near the surface of the foil were stronger than that inside the substrate. This indicates that a thin layer of Al2O3 which contained La and Zr had formed on the surface as a result of the pre-annealing treatment. Figure 5 shows the depth profiles after oxidation at 1373 K for 10.8 ks for the 50 µm thick foils of alloy E in the as-cold rolled condition and with pre-annealing in hydrogen as obtained by SIMS. With the as-cold rolled sample, the intensities of Cr and Fe were high at the outer side of the oxide layer. This indicates that the Al2O3 which formed on the as-cold rolled foil contained Fe and Cr oxides. However, with the foils which were pre-annealed in hydrogen, no apparent peaks of Fe or Cr were detected. This means that the Al2O3 which formed on the alloy pre-annealed in hydrogen contained little Fe or Cr.

63 109

(b) annealed in hydrogen

(a) as cold rolled

Secondary ion intensity (counts)

108

27Al

27Al

53Cr

53Cr

107 106 105 104

57Fe

57Fe

103

16O

102

139La

16O 139La

101 100

90Zr 0

500

1000 Time(s)

1500

2000 0

500

1000

1500

90Zr 2000

Time(s)

Figure 4: Depth profiles of 50 µm thick foils of alloy E as obtained by SIMS. (a) As-cold rolled (b) annealed in hydrogen at 1223 K for 60 s in hydrogen; contined litle Fe or CR

109

(a) as cold rolled

27Al

27Al

108 Se co nd ary ion int ens ity (co unt s)

(b) annealed in hydrogen

107 106 105 104

53Cr

53Cr

57Fe

16O

16O 139La

103 102

57Fe 139La

90Zr

90Zr

101 100 0

2000

4000 Time(s)

6000

8000 0

2000

4000 Time(s)

6000

8000

Figure 5: Depth profiles of 50 µm thick foils of alloy E after oxidation at 1373 K for 10.8 ks as obtained by SIMS. (a) As-cold rolled (b) annealed in hydrogen at 1223 K for 60 s before oxidizing

64 3.3

SEM observation

Figure 6 shows cross-sectional SEM images of the Al2O3 scale which formed on the specimens of alloys E and F in the as-cold rolled condition and with pre-annealing in hydrogen after oxidation at 1373 K for 172.8 ks in air. The alloy/scale interface was smooth with no voids at the interface. The Al2O3 scale which formed on the cold rolled foil consisted of two layers. The outer layer was about 0.5 µm in thickness and showed a morphology characterized by equiaxed grains. The inner layer was about 2.0 µm in thickness and had a columnar grain morphology, which is the same morphology as that reported by Golightly et al. (8) In as-cold rolled foils, Al2O3 scale forms in equiaxed grains during the initial oxidation period and then grows by forming columnar grains. On the other hand, the Al2O3 scale which formed on foil that had been pre-annealed in hydrogen consisted of only one layer, which was approximately 1.5um in thickness and had a columnar grain morphology. (a)

(b)

2.0um (c)

2.0um (d)

2.0um

2.0um

Figure 6: Cross-sectional SEM images of Al2O3 scale formed on alloy E and F after oxidation at 1373 K for 86.4 ks in air; (a) and (b) as-cold rolled , (c) and (d) annealed in hydrogen at 1223 K for 60 s before oxidizing

3.4

TEM observation

Figure 7 shows the results of TEM observation of the Al2O3 scale which formed on alloys E and F in the as-cold rolled condition and with pre-annealing in hydrogen after the specimens were oxidized at 1373K for 86.4ks in air. No secondary phases were apparent at the grain boundaries, as shown in Figure 7. As can be seen in this figure, the Al2O3 scale which formed on the as-cold rolled foil consisted of two layers, an outer layer of equiaxed grains and an inner layer of columnar grains. In contrast, the Al2O3 scale which formed on the foil that was preannealed in hydrogen consisted of only one layer and had a columnar grain morphology. Moreover, the size of the columnar grains of Al2O3 which formed on the pre-annealed foil was slightly larger than that of the Al2O3 grains which formed on the cold rolled foil.

65

(a)

(b)

0.5um (c)

0.5um (d)

0.5um

0.5um

Intensity (counts)

Figure 7: Cross-sectional TEM images of Al2O3 scale formed on alloy E and F after oxidation at 1373 K for 86.4 ks in air; (a) and (b) as-cold rolled, (c) and (d) annealed in hydrogen at 1223 K for 60 s before oxidizing

X-ray intensity ratio La-L/Al-K, Zr-L/Al-K

(a)

0.06 0.05

(c)

0.04 0.03

50 (b) 40 Al-K 30 Fe-K 20 O-K Cr-K 10 Zr-L La-L 0 0 2 4 6 8 10 X-ray energy, E (keV)

La/Al Zr/Al

0.02 0.01 0

40 30 20 10 0

10 20 30 40

Distance from grain boundary, d / nm Figure 8: (a) TEM image of grain boundary in parallel section of Al2O3 scale formed on alloy E after oxidation at 1373 K for 86.4 ks in air; (b) EDX spectrum from grain boundary in parallel section of Al2O3 scale formed on alloy E and (c) X-ray intensity ratios of La-L/Al-K and Zr-L/Al-K in EDX spectrum across grain boundary

66 One of the authors has reported previously that the segregation of La, Zr, and Hf at grain boundaries in A l 2 O 3 scale suppresses oxygen diffusion along the Al2O3 grain boundaries, resulting in a decrease in the growth rate of the Al2O3 scale. (9) In the present study, a distinct La-L intensity and Zr-L peak were also detected at each grain boundary by EDX analysis, as shown in Figure8 (b). In Figure 8(c), the X-ray intensity ratios of La-L to Al-K and Zr-L to AlK were taken from the points indicated by the dots in Figure 8(a) and plotted against the distance from one grain boundary. The results revealed that La and Zr had segregated to the grain boundary. Figure 9(a) and 9(b) show TEM images of the grain boundary in a parallel section of the Al2O3 scale formed on alloy F in the as-cold rolled condition and with pre-annealing in hydrogen after the specimens were oxidized at 1373 K for 86.4 ks. In the EDX spectrum from the columnar grain boundary of the Al2O3 scale on the as-cold rolled specimen of alloy F, a distinct Fe-L intensity and Cr-L intensity peak were detected, as shown in Figure 9(c). Moreover, from the equiaxed grain boundary of the Al2O3 scale on alloy F with pre-annealing in hydrogen, LaL intensity and an Hf-L intensity peak were detected, but an Fe-L intensity and Cr-L intensity peak were not detected, as shown in Figure 9(d). These results were in good agreement with the results obtained by SIMS. (b)

(a)

0.5um Intensity (counts)

Intensity (counts)

0.5um 50 (c) Al-K 40 Cr-K 30 O-K Fe-K 20 10 0 0 2 4 6 8 10 X-ray energy, E/keV

50 (d) Al-K 40 30 Cr-K 20 O-K La-L Fe-K Hf-L 10 00 2 4 6 8 10 X-ray energy, E/keV

Figure 9: (a) TEM image of grain boundary in parallel section of Al2O3 scale formed on as-cold rolled alloy F; (b) TEM image of grain boundary in parallel section of Al2O3 scale formed on alloy F annealed in hydrogen; (c) EDX spectrum from grain boundary of columnar grain in parallel section of Al2O3 scale formed on as-cold rolled alloy F; (d) EDX spectrum from grain boundary of equiaxed grain in parallel section of Al2O3 scale formed on alloy F annealed in hydrogen

Figure 10 is a plot of the X-ray intensity ratios of La-L, Zr-L, and Hf-L to Al-K at the Al2O3 grain boundary, against the distance from the alloy/scale interface. These results show that the degree of La, Zr, and Hf segregation at the Al2O3 grain boundaries was higher at the columnar grain boundaries than at the equiaxed grain boundaries. In addition, the intensity ratios of La-

67 L, Zr-L, and Hf-L at the Al2O3 c o l u mn a r grain boundaries were stronger with the preannealed alloys than with the as-cold rolled alloys. Based on the results of SEM observation, Golightly has reported that the Al2O3 scale on FeCr-Al alloys with small contents of rare earth metals forms at the alloy/scale interface mainly by inward diffusion of oxygen through the scale. (7) Similarly, based on experiments using an O18 tracer, Reddy has reported that the Al2O3 scale on an Fe-Cr-Al alloy grew by inward diffusion of oxygen through the grain boundary of the Al2O3. (10) In this study, in Fe-Cr-Al alloys with small contents of La-Zr or La-Hf, the alloy/scale interface was smooth with no voids at the interface, as shown in Figure 5. This means that outward diffusion of Al was suppressed, and the Al2O3 scale grew by inward diffusion of oxygen through the Al2O3 scale. Considering the fact that the oxidation curves of the alloys which were pre-annealed in hydrogen basically conformed to the parabolic rate law, it was inferred that the pre-annealing treatment did not change the mechanism of oxidation, but rather suppressed the inward diffusion of oxygen through the Al2O3 scale. (c) X-ray intensity ratio , La-L/Al-K

as cold rolled annealed in hydrogen

0.01

00.0 0.5 1.0 1.5 2.0 Distance from interface (um) (b) 0.02 as cold rolled annealed in hydrogen 0.01

0 0.0

0.5

1.0

1.5

2.0

Distance from interface (um)

X-ray intensity ratio , Hf-L/Al-K

X-ray intensity ratio , Zr-L/Al-K

X-ray intensity ratio , La-L/Al-K

(a) 0.02

0.02

as cold rolled annealed in hydrogen

0.01

0

0.0 0.5 1.0 1.5 2.0 Distance from interface (um) (d)

0.02

as cold rolled annealed in hydrogen

0.01

0

0.0

0.5

1.0

1.5

2.0

Distance from interface (um)

Figure 10: X-ray intensity ratios, (a) La-L/Al-K, (b) Zr-L/Al-K, (c) La-L/Al-K, (d) Hf-L/Al-K in EDX spectra from grain boundaries in columnar grain and equiaxed grain oxide layers of Al2O3 scale on alloys E and F

As with the cold rolled alloy specimens, Fe and Cr were oxidized together with Al during the initial oxidation period because the partial pressure of oxygen at the surface was high. The large amount of Fe and Cr oxides in the Al2O3 prevented the Al2O3 from growing by forming columnar grains. Instead, Al2O3 grew by forming small equiaxed grains during the initial oxidation period. The growth of these equiaxed grains then reduced the partial pressure of oxygen at the alloy/scale interface, and as a result, only Al was oxidized. After this point, Al2O3 grew by forming columnar grains at the alloy/scale interface, resulting in a two layer structure, and La, Zr, and/or Hf segregated at the grain boundary.

68 However, when pre-annealing in hydrogen was performed, Fe and Cr were not oxidized in this reducing atmosphere. The Al2O3 oxide, including a small amount of La, Zr, and/or Hf, grew by forming large columnar grains beginning in the initial oxidation period. It may be inferred that the low density of grain boundaries in the Al2O3 and the high segregation of La, Zr, and/or Hf at the Al2O3 grain boundaries suppressed oxygen diffusion along these grain boundaries. For this reason, the pre-annealing treatment in hydrogen reduced the growth rate of the Al2O3 scale.

4

Conclusion

The effect of pre-annealing on the oxidation behavior of Fe-20mass%Cr-5mass%Al alloy foils containing a small amount of La-Zr or La-Hf was examined in a cyclic oxidation test at 1373 K in air. 1. The oxidation rate of these alloys was reduced by pre-annealing in hydrogen. 2. In the Al2O3 scale which formed on the pre-annealed alloys, the outer equiaxed grain layer was thinner and the grain size of the inner columnar grain layer was larger than in the scale which formed on as-cold rolled alloys. 3. The segregation of La, Zr, and/or Hf at the columnar grain boundaries of the Al2O3 scale was higher with the pre-annealed alloys than with as-cold rolled alloys. It may be inferred that the low density of grain boundaries in the Al2O3 and high segregation of La, Zr, and/or Hf in the Al2O3 grain boundaries suppressed oxygen diffusion along these grain boundaries, reducing the growth rate of the Al2O3 scale.

5 [1] [2] [3] [4] [5] [6]

References

S. Isobe: Denki-seikoh, 58 (1987), 104. D. R. Sigler: Oxid. Met, 32 (1989), 337. D. R. Sigler: Oxid. Met, 40 (1993), 555. . A. Golightly, F. H. Stott and G. C. Wood: Oxid. Met, 10 (1976), 163. K. Ishii and T. Kawasaki: J. Japan Inst. Metals, 56(1992), 854. H. Hindam and D. P. Whittle: Proc. 3rd JIM Int. Symp. on High Temperature Corrosion of Metals and Alloys, The Japan Institute of Metals, Supplement to Trans. JIM, 24 (1983), 261. [7] F. A. Golightly, F. H. Stott and G. C. Wood: J. Electrochem. Soc., 126 (1979), 1035. [8] T. A. Ramanarayanan, M. Raghavan and R. Petkovic-Luton: J. Electrochem. Soc., 131(1984), 923. [9] K. Fukuda, K. Ishii, M. Kohno and S. Satoh: Proc. Int. Symp. on High-Temperature Corrosion and Protection 2000, (2000) Hokkaido ISIJ, p. 309. [10] K. P. R. Reddy, J. L. Smialek and A. R. Cooper: Oxid. Metals, 17(1982), 429.

Oxidation Induced Length Change of Thin Gauge Fe-Cr-Al Alloys C. Steve Chang, Leigh Chen, and Bijendra Jha Engineered Materials Solutions, Inc., Attleboro, MA 02703 USA

1

Abstract

Alloys of Fe-20Cr-5Al have been used extensively as the material of choice in metallic catalytic converter substrates. The alloy chemistry has been developed through the last decade to provide the oxidation weight gain resistance that was thought to be adequate. Mainly by the addition of rare earth and active elements, the cyclic oxidation behavior has been improved to the point to meet the regulation requirements on the durability of catalytic converters. The major uncertainty on the understanding of oxidation behaviors of thin gauge Fe-Cr-Al foil is the cause of foil length change. It has been suggested that this length change is one of the causes of buckled honeycomb observed in the fractured substrate. Analysis on the stress and strain of thin gauge Fe-Cr-Al foil under the oxidation condition has been published to account for the length changes due to the oxidation. However, the potential metallurgical factors that control the length change has not been rationalized yet. In this paper we will present (1) the oxidation test results on Fe-Cr-Al foils with different gauge and chemistry, (2) the microstructure evolution and (3) oxide scale development during the oxidation test. The oxidation failure mechanisms will be demonstrated. A phenomenological model to incorporate the alloy chemistry, microstructure and oxide scale will be described to account for the oxidation behaviors of these Fe-Cr-Al alloys.

2.

Introduction

Catalytic converters have become the universal solution for automobiles to meet the emission regulations. The catalytic converters with metallic type substrates have seen ever-wider acceptance because of several advantages over conventional ceramic-based converters[1]. One of the advantages of metallic substrate is the thinner substrate wall (30 to 50 microns) which provides lower backpressure and smaller package. However, the desirable thin gauge of alloy foils limit the usable life of the substrates. This limit is due to the fact that the metallic alloys require the formation of stable, protective scales such as aluminum and chromium oxides to slow down the oxidation of the alloy. The scale acts as a barrier to slow the diffusion of oxidizing agents such as oxygen to reach the alloy. However, the continuous oxidation of Al or Cr to form scale is like a reservoir being continuously drained and eventually the protective elements will be exhausted. At this point, oxygen will be able to react with the rest of the alloys to form the non-protective oxides, which causes the alloy to gain weight rapidly. The rapid and catastrophic oxidation weight gain is usually referred to as breakaway oxidation [2].

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

70 The alloy of choice for the metallic converter substrate has been the ferritic stainless steel with a nominal composition of 20wt% Cr, 5wt% Al and the balance of Fe. The addition of 5wt% Al provides the stable scale for the alloy to be used above 1100°C while the 20wt% Cr provides the oxidation and corrosion resistance from the ambient to where the Al oxidation becomes significant. The cyclic oxidation resistance is improved by the addition of rare earth elements such as Y, La and Ce. The addition of active elements such as Zr, Hf provides further improvement on the oxidation resistance [3]. The oxidation resistance is commonly measured by the amount of sample weight gain due to the scale formation. The oxidation weight gain is important, as it is an indication of the effectiveness of protective scale. The oxidation weight gain for the Fe-20Cr-5Al alloy has seen significant improvement and known to be one of the most oxidation resistant materials. Nonetheless, Fe-Cr-Al foils for the catalytic converter applications have to be dimensionally stable to avoid rupture of the substrate during the service [4]. The source of the dimensional instability has been attributed to the stress between scale and ally, which causes the creep deformation of substrate. Heats of Fe-Cr-Al alloys having identical oxidation weight gain behaviors have shown drastically different dimension changes. This paper summarizes the effort to rationalize the oxidation length change mechanism and attempts to draw a guideline to prevent dimensional instability. On the practical aspect of producing Fe-Cr-Al foils, it has been known that the conversion of Fe-20Cr-5Al alloy to thin gauge has been difficult and contributed to the high cost of the materials. A commercially feasible process [5,6] to produce thin gauge Fe-Cr-Al foils has been developed to address the issue. The process starts out with roll bonding the Fe-Cr alloys (AISI 4xx type ferritic stainless steels) to proper amount of Al alloys to form a three layer composites. The composite coil, which was roll bonded with proper attention so it can be cold rolled to an intermediate thickness. The intermediate thickness was selected to allow a homogenization heat treatment to be conducted at the temperature and time, which are commercially acceptable. After the heat treatment, the strip is cold rolled to provide the desirable temper and finish. Obviously, the advantage of alloying the Al to Fe-Cr by the roll bonding process is to circumvent the limits of Al content and the conversion difficulties[7]. In this study, oxidation tests on the oxidation weight gain and dimension stability were conducted on materials taken from roll bonding produced Fe-Cr-Al alloys. The effects of chemistry and heat treatment on the oxidation behaviors, in particular the dimensional stability is rationalized with a phenomenological model developed from examining the length change behaviors of hundreds of samples. This oxidation length change model will attempt to show that from the synergetic effects of physical (e.g. density and thermal expansion) and metallurgical (e.g. scale adhesion) changes, the various length change behaviors can be accounted for.

3

Experimental Procedure

3.1

Materials

The Fe-Cr-Al alloy foils were produced via the roll bonding process. In brief, Fe-Cr alloy (stainless steel) strips were clad with Al strips on both sides by feeding the strips into a fourhigh rolling mill. The cladding process was developed to apply sufficient reduction to form a well adhere three layer (Al/SS/Al) composite. The roll-bonded composites were cold rolled to

71 an intermediate thickness followed by heat treatment to homogenize the Al layers with the FeCr alloy. A finial cold rolling was applied to reduce the alloyed strip to the finish foil thickness. Fe-Cr-Al alloys with Al content range between 5 to 8 wt% are readily produced by this roll bonding process. Rare earth addition in the manner of La+Ce was accomplished by casting the Fe-Cr alloy with misch-metals. Typical alloy chemistry is shown in Table 1 in weight %. Table 1: Nominal composition of Fe-Cr-Al alloys in wt% C Mn P S Si Cr Ni Al N O La Ce 0.02 0.2 0.02 <.002 0.2 21 0.2 6 0.02 0.008 0.01 0.03 3.2

Oxidation Test

The oxidation test was conducted on a honeycomb type sample at the cell density of approximately 400 cpi (cell per square inch). The cold rolled Fe-Cr-Al alloy foils were degreased prior to corrugation and assembling. Honeycomb samples consisted of one flat and one corrugated strip. The two strips were rolled up to form a cylinder along the rolling direction. The cylindrical honeycomb samples had the dimensions of approximately 74 mm in length and 20 mm in diameter. Flat coupons (50 mm by 50 mm) were also used to obtain suitable samples for microstructure and scale characterizations. These coupons were prepared and tested in the same conditions as the honeycomb samples except the dimensional change was not measured. However, oxidation weight gains of coupon and honeycomb type sample were always cross-checked to ensure that the oxidation weight gain followed the same rate. Oxidation tests were conducted on samples in the cold-rolled and vacuum heat-treated conditions. Vacuum heat treatment was performed in a diffusion pump evacuated, cold wall furnace at 1200 °C for 30 minutes. The typical vacuum level at the soak temperature was better than 10–5 torr. The oxidation test was conducted in a semi-cyclic manner. The samples were placed in the furnace at the ambient condition. The furnace temperature was ramp to 1100°C in two hours and held for increasing duration (5, 20 and 25 hours). The duration of hold was increased to 50 hours per cycle afterward until a total of 400 hours was accumulated. After each hold period, furnace was ramp down to ambient in 6 hours. Samples were then removed to measure the weight gain and length changes. The oxidation sample weight was measured in a precision scale with accuracy to 0.00001g. The dimensional change of oxidation sample was measured by a dial indicator on the length of the honeycomb cylinder. Oxidation weight gain results were analyzed by applying the parabolic rate equation (1) [8] between the weight gain (DW) and test time (t) for the alumina forming alloys:

,W = Dt

(1)

Where D is the parabolic rate constant. The rate constant was obtained by a linear curve fit on the square root of time plot. The validity of equation 1 is verified by (1) plot the weight gain against square root of time for a linear dependence and (2) plot log-log of weight gain against time for ½ slope value.

72 3.3

Microstructure and Oxide Scale Characterization

The microstructure of Fe-Cr-Al alloys and the morphology of surface scales were examined by the optical microscope as well as by the energy dispersive x-ray (EDX) equipped scanning electron microscope (SEM). The edge-on type samples for scale morphology examination were obtained by pulling the oxidation-tested samples in a tensile testing machine until fracture.

4

Oxidation Test Results

Six oxidation test samples are shown in Figure 1in the following conditions. From left to right, these samples are: (a) cold rolled, (b) vacuum heat treated, (c) 400 hours tested, normal, (d) 400 hour tested, abnormal length, (e) 400 hour test, Al depleted, no nreakaway, and (f) 400 hour tested, weight gain breakaway. The oxidation weight gain and length change results will be presented in the following sections to show the effect of the sample conditions such as heat treatment and alloy chemistry on the oxidation behavior.

(a) (b) (c) (d) (e) (f) Figure 1: Oxidation samples (a) cold rolled; (b) vacuum heat treated; (c) 400 hr tested, no breakaway; (d) 400 hr tested, length change abnormal; (e) 400 hr tested, Al depleted; (f) 400 hr tested, weight gain breakaway

4.1

Effect of heat treatment

The oxidation weight gains and length changes for the representative cold rolled and vacuum annealed samples with normal length change are shown in Figure 2 (a). The square root of time plat is shown in Figure 2 (b). The samples were 50 micron thick and the Al content of 6 wt%. Length change results were shown together with the oxidation weight gain to delineate the relationship between the two.

73

6

0.0012

5

0.001

4

0.0008

3

0.0006

2

0.0004

1

0.0002

0

0 0

Weight Gain/area (g/cm 2)

a)

100

200 Tim e, hours

-1 400

300

w eight vacuum annealed

w eight cold rolled

length vacuum annealed

length cold rolled

0.0014

6

0.0012

5

0.001

4

0.0008

3

0.0006

2

0.0004

1

0.0002

0

0 0

b)

200

400 600 800 Square Root of Tim e (sec1/2)

Length Change (%)

0.0014

1000

50 micron vacuum annealed

50 micron cold rolled

length vacuum annealed

length cold rolled

Length Change (%)

Weight Gain/area (g/cm 2)

The rapid oxidation weight gain (open square markers) of cold rolled sample quickly depleted the Al in the alloy (at ~150 hours). The oxidation weight gain curve of cold rolled sample in Figure 2 (b) changed its slope at this point.

-1 1200

Figure 2: Typical oxidation weight gains and length changes for Fe-Cr-Al foils at 1100°C, room air; (a) linear time scale, (b) square root of time scale

Meanwhile, the length change (solid square marker) of cold rolled sample started to show rapid increase at 150 hours as well. However, the breakaway of oxidation weight gain did not start until 300 hours of oxidation time. In contrast, the vacuum annealed sample showed much lower oxidation weight gain. The parabolic rate constants (kp) are 160 and 49.6 (10–14 g2/cm4sec) for the cold rolled and vacuum annealed samples, respectively. The parabolic rate for the cold rolled sample was obtained by linear curve fit on the weight gain between 5 and 150 hours while that for the vacuum annealed sample was obtained between 5 and 400 hours.

74 The depletion of Al in the cold rolled sample was verified by SEM/EDX analysis on the coupon samples being oxidized along with the honeycomb samples. The Al content in cold rolled sample being oxidized for 150 hours was less than 0.5 wt%. The calculated oxidation weight gain of Fe-Cr-Al alloy, assuming that Al was oxidized to form aluminum oxide, for 50 micron gauge and 6% Al composition, would yield weight gain of 0.00095 g/cm2. This value matches very well with the change of weight gain rate (at weight gain of approximately 0.001 g/cm2) as well as the length change breakaway. Based on the experimental results of weight gain and length change, it is concluded that the length change breakaway occurred when the Al is depleted in the cold rolled samples. The Al content and time to reach length change breakaway is always related for the cold rolled samples. The vacuum annealed sample gained approximately 0.0009 g/cm2 of weight at the end of 400-hour test cycle and this is less than the Al depletion point. The vacuum annealing heat treatment reduced the oxidation weight gain rate and thus delayed the onset of length change breakaway. For the vacuum annealed samples, the correlation between the Al content and time to length change breakaway is not as predictable as that for the cold rolled sample. The unpredictable length change breakaway for vacuum annealed samples will be shown in the next section when oxidation behaviors for foils at different gauges are presented. 4.2

Effect of Foil Thickness

The oxidation weight gains for 6wt% Al, Fe-Cr-Al alloys tested at the thickness of 30, 40 and 50 micron are shown in Figure 3 for the vacuum annealed samples. The results are presented as weight gain and length change against square root of time.

Oxidation Behavior and Foil Gage 1100C, High Length Change

8

0.0006

6

0.0004

4

0.0002

2

0 0

200

400

600

800

1000

Length Change, %

Weight Gain/Area (g/cm 2)

0.0008

0 1200

Square Root of Tim e, sec1/2 30 m, DWt

40 m, DWt

50 m, DWt

30 m, DL

40 m, DL

50 m, DL

Figure 3: Oxidation weight gain and length change for materials show abnormal length change behaviors at 30, 40 and 50 mm.

Samples were originally from the same master coil except being rolled to different thickness. The weight gain results are shown in open symbols while length changes are shown in solid

75 symbols. The weight gains of samples at the three gauges clearly followed the same trend, i.e. same oxidation rate. The oxidation weight gain rates were not dependent on the foil thickness. The foil thickness determined the amount of Al that was available for scale formation. The calculated weight gain values of 6% Al alloys, for the 30, 40 and 50 micron foils are: 0.00057, 0.00076 and 0.00095 g/cm2 respectively. The individual weight gain values show that the slope changes for each thickness match very well with these values. The time to deplete the Al can be predicted from the short-term measurement of oxidation rate. The breakaway of oxidation induced length changes for samples in Figure 3 showed some similarity to that of the cold rolled samples, although the length change was on the positive side. As soon as the weight gain reached the critical values at which the Al was depleted, the length changes started to deviate from the “quasi-parabolic”, stable trends. However, this “quasi-parabolic” length increase behavior was not the norm. The oxidation length changes, in the case of vacuum annealed samples, were most likely to stay constant during the oxidation test even though the Al had been depleted. Oxidation test results for samples with the more usual, stable dimensions are given in Figure 4. Oxidation weight gains for the samples with normal length change are approximately the same as given in Figure 3 for the abnormal length change samples. However, the length change behaviors are significantly different from that of the abnormal samples. The oxidation length changes were much lower and what was more significant, the samples in Figure 4 did not show the breakaway behaviors as in Figure 3 when the Al was depleted in the alloy. These results demonstrated that the oxidation weight gain by itself could not be used to predict the length change behavior reliably

Oxidation Behavior and Foil Gage 1100C, Normal Length Change

8

0.0006

6

0.0004

4

0.0002

2

0 0

200

400

600

800

1000

Length Change, %

Weight Gain/Area (g/cm 2)

0.0008

0 1200

Square Root of Tim e, sec 1/2 30 m, DWt

40 m, DWt

50 m, DWt

30 m, DL

40 m, DL

50 m, DL

Figure 4: Oxidation weight gain and length change for materials show normal length change behaviors at 30, 40 and 50 mm.

76 4.3

Microstructure and Oxide Scale

0.001

5

0.0008

4

0.0006

3

0.0004

2

0.0002

1

0

0

0

200

400

600

800

1000

Length Change (%)

Weight Gain/Area (g/cm 2)

The evolution of microstructure and oxide scale in the vacuum heat treated and oxidation tested conditions were investigated for two samples at 50 microns with different oxidation length change behaviors. The oxidation test results for the two samples are shown in Figure 5. The oxidation weight gains were virtually identical for these two samples. However, the oxidation length changes of the two samples were significantly different. The “normal” (triangle markers) sample showed very small length change through the test while the “abnormal” sample showed continuous length increase through the oxidation test.

1200

1/2

Square Root of Tim e, (sec ) Abnormal, DWt

Normal, DWt

Abnormal, DL

Normal, DL

Figure 5: Oxidation weight gain and length change for samples with normal and abnormal length changes

The microstructures of the two samples are shown in Figure 6 from coupons withdrawn after being oxidation tested up to each test cycle. Optical micrographs of samples in the vacuum annealed and oxidation tested (5, 25, 100 and 200 hours) conditions show significant difference in the microstructure development. The vacuum annealed sample with the “normal” length change behavior had the microstructure of approximately two grains through the thickness while the “abnormal” sample showed the classic bamboo structure. The “normal” sample gradually developed a duplex grain structure as the oxidation test proceeded while the grain structure of the “abnormal” sample was quite stable through the oxidation test. The Auger analysis on the surface of vacuum heat-treated samples of “normal” and “abnormal” showed a layer of approximately 200-angstrom thick oxide scale. There is no detectable difference in the composition of scale, which consists of only Al and oxygen. The oxide scales on the two alloys are shown in Figure 7 for oxidation tested to 100 hours. The scale on the sample of “normal” length change behavior shows a mixed morphology of equaxis and columnar oxides. However, the scales on the “abnormal” sample show dual layer morphology. The top layer, which was developed earlier, can be described, as equaxis while the lower layer is clearly columnar.

77

Figure 6: Microstructure for samples show abnormal length changes (left) and normal length change (right) for vacuum heat treated, oxidation tested 5, 25, 100 and 200 hours (from top to bottom)

Figure 7: SEM micrographs of oxide scale on 100 hours oxidation samples show normal (top) and abnormal (bottom) length change behaviors

78

5

Discussion

The oxidation weight gain of Fe-Cr-Al, with rare earth addition, at 1100 °C in air, clearly followed the classic, diffusion controlled, parabolic behavior. The oxidation weight gain rate is thus a useful value to calculate the time to deplete the Al in the alloy. The oxidation tests carried out here has shown that the weight gain breakaway did not start at the depletion of Al. However, the oxidation length change breakaway is closely related to the Al depletion. At least in the cases of cold rolled samples length change breakaway is always associated with the Al depletion. The obvious effect of vacuum heat treatment on the oxidation behavior is the significant reduction of weight gain. The direct consequence is to extend the duration of stable, protective oxide growth. It is not clear if the rapid oxidation weight gain in the cold rolled sample is due to either the doping of Fe/Cr in the scale or thin layer of aluminum oxide formed during the vacuum heat treatment. This thin layer of aluminum oxide might act as a protective layer during the first heating up and enable the formation of pure and highly protective aluminum oxide scale. Another possible effect of prior heat treatment is to avoid the recrystallization of cold rolled structure in the first test cycle, which might disrupt the formation of a stable scale. However, the dimension stability is somewhat more complicate than the weight gain during the oxidation. Although the depletion of Al has been shown to coincide with the length change breakaway of cold rolled samples, the vacuum annealed samples behaved quite unpredictably. An attempt is made to generalize the oxidation length change behaviors for the Fe-Cr-Al foils to provide the base for a mechanism that is useful to analyze the metallurgical variables. Figure 8 depicts the four types of commonly observed length change behaviors. The most common length change is the one of which the foil shows stable length through out the oxidation test (Curve B, diamond). Curve C (check mark) shows that the foil starts out with small length increase then follow a shrinking trend to reach a stable dimension. This type of behavior is most commonly observed in cold rolled samples prior to the depletion of Al. This type of behavior is also observed in foil thinner than 40 micron or with higher Al contents. The third type of length change (Curve A, triangle) is termed as “abnormal” since the sample length increases continuously through the oxidation test. Curve D (square) shows the onset of breakaway length change when the Al was depleted after a period of continuous length increase. Curve D has to be considered for two scenarios. One is attributed to that when foil gauge is so thin that weight gain breakaway occurs prior to the end of oxidation test. The other scenario belongs to the abnormal behavior since the length change breakaway takes place when the Al is depleted (see Figure 3). It is necessary to mention that the length change breakaway will take place eventually when the weight gain breakaway occurs. Within these length change behaviors, it is most critical to understand the “abnormal” behaviors and rationalize against, if available, controllable material variables in order to avoid the undesirable consequence on the catalytic substrate. The sample dimension, during the oxidation test, was under the influence of several stresses of thermal-mechanical nature. Schutze and Przybilla [9] discussed the various origins of stress during the oxidation of Fe-Cr-Al alloy foils. A multitude of stresses can coexist in the alloyscale system while the oxidation proceeds. However, two types of stress have aroused the most attention. The thermal expansion mismatch between the scale and alloy substrate will introduce a thermal stress on the alloy during the cooling, which will be in the tensile state. A growth

79 stress is also presented due to the combined effects of oxide growth as well as structure change might take place in either the scale or alloy itself. The thermal stress values from calculation and experimental measurement have shown to match reasonably well [10]. However, the stress states thus obtained, in the foil, are tensile and this implies that the foil shall be continuously growing as the oxidation proceeds. As being mentioned earlier, the majority of length changes of Fe-Cr-Al foils is near zero or even negative. It is necessary to find a mechanism to provided the compressive stress in the scale-alloy system that can counter the tensile stress from the thermal/growth mechanism. Schutze and Pryzbilla have suggested that there can be a “relaxation” mechanism, which relieves the foil from being stretched by the scale during the cooling. It is worthwhile noting that the relaxation mechanism, which causes the foil to show no significant dimensional change, will not provide a negative length change.

Typical Oxidation Length Change Behaviors

Length Change, %

4 3 2 1 0 -1 0

100

200

300

400

Tim e, hours A: abnormal, quasi-parabolic

B: stable

C: shrinking

D: abnormal, early breakaw ay

Figure 8: Generalized oxidation length change behaviors for (a) abnormal, (b) stable, (c) shrinking and (d) breakaway after Al depletion

Nevertheless, it might be the simple fact that the foil density is changing as the Al is consumed during the oxidation process. The density of Fe-20Cr-6Al alloy is 7.15 g/cm3 while that of Fe-20Cr alloy is 7.5 g/cm3. Since the Al oxidation follows the parabolic rate relation, the volume of alloy decreases follows the density increase. The dimensional change from density increase can amount to approximately a 2% decrease in the linear dimension of foil. Combining the volumetric introduced dimensional change (~-2%) to that of the thermal stress induced creep of approximately 1%. The apparent dimensional change at the complete depletion of Al is about 1%, which has been observed for the majority of the vacuum annealed samples. The aforementioned length change mechanism of which the stretching form the thermal/growth stress is balanced by the shrinking of materials, provides reasonably good length change value that matches the majority of measured values. To support this mechanism, it is worthwhile to examine the length changes of cold rolled sample, which always showed negative length change (i.e. shrinking). If only the scale-induced stress is considered, the rapid

80 weight gain (thickening of scale) in the cold rolled sample should have stretched the sample quickly as well. It is reasonable the oxide scales developed on the cold rolled samples were not as adherent (i.e. protective) as that of the vacuum heat-treated samples since the oxidation weight gains are much faster in cold rolled samples. Then the scales on the cold rolled samples will not be able to exert the tensile stress on the foils to cause the stretching of samples. Thus, the length of cold rolled sample will likely be controlled by the alloy density change due to the Al depletion. The abnormal length changes, in the case of quasi-parabolic trend, can be as high than 6% before the complete consumption of Al in the alloy. This level of length change can not be accounted for by the aforementioned mechanism. Schutze and Pryzbilla have pointed out that structure changes in alloy or scale might introduce additional dimensional changes. The materials showing abnormal length change can usually be identified with several compositional and microstructural abnormalities. These alloys tend to have higher rare earth and/or austenite stabilizer contents and well-developed second phases (e.g. Fe-Cr carbide). These variables might affect the microstructures and scale development during the oxidation process. Another explanation on the quasi-parabolic type, abnormal length change might be that the scale adhesion is so effective that no relaxation is possible between scale and foil. The scale on the abnormal length change sample showed columnar structure, which is usually an indication that the oxide adhere well and is growing without disruption. The microstructure of abnormal length change samples showed no change from vacuum annealing to the end of oxidation test and probably caused no disruption on the oxide growth. The abnormal growth of foil seen through the cyclic test is the accumulated creep of foil after each test cycle. The thermal expansion mismatch from cooling (1100°C to ambient) between the scale of aluminum oxide (CTE = 9 PPM/°C) and foil (CTE = 15 PPM/°C), which are in equilibrium at the test temperature, will be approximately 0.66%. Since each test cycle will introduce this amount of stretching from cooling because there is no relaxation, the 10-cycle (400 hours) oxidation test will yield a 6.6% total length change. The accumulation mechanism is supported by testing sample of 6% total length change isothermally rather than cyclically. The one cycle length change, for example 50 hours at ~0.5% is far less than that obtained ~1.5% after three cycles, for the same total oxidation duration, from the multiple cycle tests [10] . In summary, the “apparent” linear dimension of Fe-Cr-Al foils after each oxidation cycle is controlled by the combined effect of four mechanisms. The first mechanism is the creep of foil due to the thermal/growth induced tensile stress. The second mechanism is the volumetric reduction and thus the shrinking due to the density increase since the Al is being removed from the alloy. The third mechanism is the adhesion of scale to foil. The importance of adhesion is that this will determine the effectiveness of transfer of the thermal/growth stress to the foil. At the extreme case when there is no adhesion, the foil dimension shows shrinking rather than the expected growth. The last mechanism is the structural changes of either foil or scale, as this will provide materials with different thermal expansion characteristics. This mechanism can also affect the adhesion of scale to alloy and complicate the relative effectiveness of mechanism one and two. The structural changes might be the root cause of the early length change breakaway. The Al and Cr additions to Fe are known to stabilize the ferritic structure and the depletion of these elements will cause the austenitic phase, which has different thermal expansion and physical properties to be presented at elevated temperature.

81 Although there are still many unanswered questions on the oxidation induced length change, a general guideline to obtain dimensionally stable foil of Fe-Cr-Al alloy can be drawn. The alloy has to be heat treated to provide a low oxidation weight gain. The alloy shall have as low as possible alloying elements that might cause any phase transformation or second phase precipitation. The rare earth addition has to be at the level that provides sufficient scale adhesion for oxidation resistance. But the optimum amount is limited since the scale adhesion might be improved to the point that no “relaxation” is possible and thus full tensile stress is exerted on the foil to cause abnormal length change. Effects of other elements can be rationalized by considering the effects on the “relaxation” of scale adhesion. For example, free sulfur might be needed as oppose to the conventional, total elimination of sulfur approach for the best oxidation resistance.

5

Conclusion

Oxidation weight gains of Fe-Cr-Al alloys follow the classic parabolic behavior. The oxidation weight gain is significantly reduced by the vacuum heat treatment. The oxidation induced length change behaviors are categorized into four different groups. The most commonly observed behavior is that the sample maintains a stable dimension prior to weight gain breakaway. However, there are also the abnormal cases of which either the length increases continuously or the length change breakaway occurs as soon as the Al is depleted. To rationalize the commonly observed length change, be either stable or negative through the oxidation, mechanisms to counter the tensile stress states between the scale and foil is proposed. The density increase due to the Al consumption is proposed to provide the necessary “shrinking” or volumetric decrease of the foil. The foils showed abnormal length change in the case of continuous length increase are examined. Based on the scale morphology and microstructure evolution, it is proposed that a “perfect” adhesion between scales and foils existed in this type of samples. The perfect adhesion allows the thermal expansion mismatch between scale and foil to exert an approximately 0.6% stretch in each oxidation test cycle. The stretching is accumulated after each test cycle since there is no relaxation between scale and alloy and thus the observed “abnormal” length change. From a theoretic point of view, the ideal foil dimensional change, in the cyclic oxidation condition, can be estimated by the multiplication of thermal expansion mismatch to the number of test cycles. The commercial Fe-Cr-Al alloy, in order to achieve the least length change, has to strike a balance between the perfectly adherent scale for the maximum protection and the somewhat reduced oxidation resistance needed for the relaxation of adhesion.

6

Acknowledgment

Authors appreciate the extensive oxidation test performed by Linda Linehan and Bill Gorman. The permission from the management of EMS to publish the results is greatly appreciated.

82

7 [1]

Reference

S. Pelters, F. W. Kaiser and W. Maus, SAE Paper 89044, Society of Automotive Engineers, 1989 [2] N Birks and G. H. Meier in Introduction to High Temperature Oxidation of Metals, Edward Arnold, London, 1983, 141 [3] K Ishii, S Satoh, M Kobayashi and T Kawasaki in Metal-Supported Automotive Catalytic Converters (Ed.: H. Bode) Werkstoff-Informationsgesellschaft mbH, 1997, 55 [4] K Tanaka and T. Saito, Tetsu to Hagane, 1995, 81, 79–84 [5] C S Chang, A Pandey and B Jha, , SAE Paper 960566, Society of Automotive Engineers, 1996 [6] I M Sukonnik, C S Chang and B Jha, US Patent 5980658, 1999 [7] I. M. Sukonnik, C S Chang and B Jha in Metal-Supported Automotive Catalytic Converters (Ed.: H. Bode) Werkstoff-Informationsgesellschaft mbH, 1997, 93 [8] N Birks and G. H. Meier in Introduction to High Temperature Oxidation of Metals, Edward Arnold, London, 1983, 63 [9] M. Schutze and W. Przybilla in Metal-Supported Automotive Catalytic Converters (Ed.: H. Bode) Werkstoff-Informationsgesellschaft mbH, 1997, 163 [10] S Chang, Internal Study, EMSI, 2001

Improvement in the Oxidation Resistance of Al-deposited Fe-Cr-Al Foil by Pre-oxidation Shigeji Taniguchi1 Atsushi Andoh2 and Toshio Shibata1 1 Department of Materials Science and Processing, Graduate School of Engineering, Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan. 2 Steel and Technology Development Laboratories, Nisshin Steel, 5 Ishizu-nishimachi, Sakai, Osaka 592-8332, Japan.

1

Abstract

The oxidation kinetics of conventional Fe-Cr-Al foil, Al-deposited foil, and Al-deposited and pre-oxidised foil was studied at 1373 K in air. All the foils were 50 µm thick and contained minor additions of rare earth elements. The oxide scales were observed with SEM and TEM combined with EDS, and were examined with X-ray diffractometry. The deposition of Al onto the foil from vapour phase improves the oxidation resistance. Details regarding this matter were reported elsewhere. The combination of the Al deposition and the subsequent preoxidation at 1173 K for 90 ks in air further increases the oxidation resistance, i.e. the smallest parabolic rate constant among the three kinds of specimens and excellent scale adherence. The pre-oxidation enhances the growth of G-Al2O3, which transforms to =-Al2O3 during the subsequent oxidation. However, such =-Al2O3 grains are much larger than those formed on the conventional foils of similar chemical compositions. Small closed voids and small spinel type oxide particles appear in =-Al2O3 grains with the progress of oxidation. The former is explained in terms of the volume decrease accompanying the phase transformation and the latter by small solubility of Fe in =-Al2O3.

2

Introduction

The usage of a metal honeycomb made of Fe-20Cr-5Al (in mass %) alloy foil containing minor amounts of rare earth elements as a catalytic converter for automobile exhaust gas is increasing [1–3], because it has a few advantageous points over conventional ceramic honeycombs. They include smaller weight, greater resistance to vibration and smaller pressure drop for the exhaust gas. As the metal honeycomb is subjected to heating cycles in an oxidizing atmosphere, the oxidation resistance is one of the major concerns. The oxidation-resistant life of Fe-(15, 20)Cr-Al alloy foils [4–6] is essentially controlled by their Al content per unit surface area. Therefore, high Al-content foil is favoured in this respect. However, it is quite difficult and usually uneconomical to increase the Al concentration to a value more than 5 mass %, because such an alloy is brittle and hence induces difficulty during production or lowering of productivity. In addition to the above, the foil thickness is being decreased from the presently used 50 µm to 30 µm or less, in order to decrease the

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

84 pressure drop and to reduce the converter size. The total amount of Al per unit foil area becomes more important for such a thin foil, although mechanical aspects and fabricability are equally important. The deposition of Al from a vapour phase onto foil is one of the measures to circumvent such difficulties. In our previous paper [7,8] we have shown that an increase in surface Al concentration of Fe-Cr-Al alloy foil by such a method leads to much smaller oxidation rates. The Al2O3 scales formed on such specimens transform from C- to G- to =-Al2O3 with the progress of oxidation and these transformations complete in shorter periods as temperature increases. The final scale showed a two-layer structure: an outer layer consisting of equiaxial grains and an inner layer consisting of columnar grains. However, the equiaxial grains are larger and columnar grains are wider than those grown on conventional Fe-20Cr-5Al (in mass %) foils containing minor amounts of rare earth elements. The present paper deals with the influence of pre-oxidation on the oxidation resistance of the Al-deposited foil. Such a research was motivated by an idea that if G-Al2O3 was well developed in the scale, the finally obtained α-Al2O3 grains would become large and thus decrease the oxidation rate, since the scale grows predominantly by the diffusion of oxygen along the oxide grain boundaries. Such idea was confirmed by the experiment.

3.

Experimental Procedures

3.1

Specimens

Two kinds of specimen foils of 50 µm thickness were prepared by the following process. 200 kg ingots were produced by induction melting in a vacuum and they were hot-forged to slabs of 60 mm thickness. After surface defects were removed, they were hot-rolled to 2.5 mm thickness. They were further rolled at room temperature with intermittent annealing to the final thickness of 50 µm. The chemical compositions of the specimens are shown in Table 1. The deposition of Al was performed using a vacuum depositor where Al in a crucible was evaporated by electron-beam heating. The thickness of the deposit was adjusted by controlling the electron beam power and moving rate of the foil. The Al-deposited foil has a smaller Al content than the alloy foil (conventional foil), however, application of Al deposit layer of 1.1 µm thickness resulted in a similar total Al content as the alloy foil. Rectangular foil specimens measuring 60 × 50 × 0.05 in mm were cut out of the rolled sheets. Several of the Al-deposited foils were pre-oxidised. Table 1: Chemical compositions of specimen foils (mass%)

In the following Al-deposited and pre-oxidised foil, Al-deposited foil, and alloy foil are referred to as specimens A, B and C, respectively. They were ultrasonically washed in an acetone bath before oxidation runs except specimen A which received no treatment after the pre-

85 oxidation. The present paper deals mainly with the results of specimen A, since some results of specimens B and C have been shown elsewhere. [8,9] 3.2

Oxidation tests

The pre-oxidation was performed at 1173 K for 90 ks in air using an electric furnace. These conditions were chosen to obtain scales consisting mainly of G-Al2O3, basing on the previous results. [8,9] The oxidation resistance was assessed by cyclic oxidation test with temperature varying between room temperature and 1373 K in static laboratory air. The specimens for metallographic examinations were oxidised similarly in the electric furnace. Each specimen was put in alumina crucibles, which were then placed in the furnace and taken out after specified periods of oxidation. 3.3

Metallographic Examinations

Oxides were identified by X-ray diffractometry (XRD) using Cu K-α-radiation at 40 kV and 150 mA. Specimen surfaces and fractured sections were observed by scanning electron microscopy (SEM) combined with energy dispersive spectroscopy (EDS) unit for analyzing the relevant elements. The acceleration voltage for SEM was 15 kV and beam diameter for EDS was 1 µm. The fractured sections were revealed by bending the specimens in liquid nitrogen. The depth profiles of the relevant elements were obtained by secondary ion mass spectrometry (SIMS). Bright field images and selected area diffraction (SAD) patterns were obtained using transmission electron microscopy (TEM), equipped with an EDS unit. The acceleration voltage for TEM was 300 kV and beam diameter for EDS was 20 to 30 nm. The TEM specimens were prepared in the following manner; the specimen thickness was decreased to about 50 µm by mechanical polishing and then further thinned with Ar ion beam until a very small hole is made.

4

Results

4.1

Oxidation kinetics

The oxidation curves of the three specimens are shown in Figure 1, where the mass gain due to oxidation is plotted against square root of time. From the slopes of the linear parts of the curves parabolic rate constants were obtained and are shown in Table 2. It is clear that specimen A shows the smallest rate constant, indicating the best oxidation resistance among the three specimens. The rate constant was decreased to about a third of that of the alloy foil by the Al deposition and further decreased to about one sixth by the combination of Al deposition and pre-oxidation. The mass gain at the beginning of oxidation for specimen A is due to the pre-oxidation. During and after the oxidation no scale spallation was observed for all the specimens.

86 Table 2: Parabolic rate constants for the oxidation at 1373 K in air

Figure 1: Oxidation kinetic curves for the three kinds of specimens at 1373 K in air

4.2

Phase transformations

Figure 2 summarises phase transformations of Al2O3 scales on specimen A with the progress of oxidation. It is clear that the pre-oxidation resulted in the scale consisting mainly of GAl2O3. The 60 s oxidation resulted in G-Al2O3 and =-Al2O3. Then, peaks for =-Al2O3 become higher with an increase in the oxidation time. Almost all G-Al2O3 transformed to =-Al2O3 by oxidation for 3.6 ks. 360 ks oxidation resulted in spinel type oxides of Fe(Cr, Al) 2O4. 4.3

SEM observations

Detailed SEM observations were performed for the three kinds of specimens oxidised for 3.6, 90 and 360 ks. However, the results for only specimen A are shown here in Figure 3, since the results for the other specimens were already shown elsewhere. [8,9] The surface of specimen A is covered with numerous needle-like crystals, which were shown [9] to grow during Cto G-Al2O3 transformation. The fractured section shows that the scale obtained by the preoxidation is about 2 µm thick with needle-like crystals on it. By 3.6 ks oxidation the scale became thicker and the needle-like crystals became blunt, although the scale kept growing.

87

Figure 2: XRD patterns for the Al-deposited and pre-oxidised specimen, after oxidation at 1373 K for various periods in air

After 90 ks oxidation the scale shows columnar grains near the scale/substrate interface. Contrarily, the columnar structure in the scale on specimen C was well developed. The numerical date for the grain size after 90 ks oxidation are shown in Table 3. After 360 ks oxidation the scale on specimen A shows a two layer structure; the outer layer is about 0.8 to 1.2 µm consisting of equiaxial grains and the inner layer of columnar grains of 1.5 to 1.9 µm height as Figure 3 shows. Table 3: Comparison of =-Al2O3 grain size after oxidation at 1373 K for 90 ks in air

4.4

TEM observation

Figure 4 shows a bright field image of specimen A oxidised at 1373 K for 3.6 ks and diffraction patterns at (a) to (c). The specimen was tilted so that an outer part of the scale makes strong contrast. The diffraction patterns at (a) to (c) completely agree with each other and were identified as =-Al2O3. This indicates that the dark area including (a) to (c) consists of a single crystal of =-Al2O3. This crystal is larger than equiaxial grains of =-Al2O3 observed with specimen C. Many small voids were also found and are indicated by arrows in the figure. These

88 tendencies are the same for several of other =-Al2O3 grains. The scale consists of nearly equiaxial grains.

Figure 3: SEM micrographs for the surfaces and cross sections of the Al-deposited and pre-oxidised specimens, after oxidation at 1373 K for various periods

89

Figure 4: TEM micrograph and diffraction patterns of the Al-deposited and preoxidised specimen, after oxidation at 1373 K for 3.6 ks in air

Figure 5 shows a bright field image of specimen A oxidised at 1373 K for 90 ks and diffraction patterns at (a) to (c). The diffraction patterns are the same to each other again, indicating that the grain containing (a) to (c) is a single crystal of =-Al2O3. In particular, the grain is large enough to occupy almost whole scale thickness. The tilting tests were performed so that nearby grains are confirmed to be very large. The scale developed columnar structure at the bottom half at some areas. This is shown in Fig. 3 already. There are again closed pores the diameters of which are 10 to 160 nm in the α-Al2O3 grains. Figure 6 shows a higher magnification view of an are in Fig. 5, indicating the presence of spinel type oxide particles (dark) and voids (pale gray). There are no such voids at and near the scale/substrate interface which is very sharp as Fig. 7 shows. In general, these closed voids are found in the areas a little apart from the interface.

90

Figure 5: TEM micrograph and diffraction patterns of the Al-deposited and preoxidised specimen, after oxidation at 1373 K for 90 ks

Figure 6: A higher magnification view of a part of Fig. 5, showing spinel type oxide particles (dark particle) and closed voids (pale gray)

91

Figure 7: A higher magnification view of a part of Fig. 5, showing sharp scale/substrate interface and quite few spinel type oxide particles

5

Discussion

5.1

Appearance of metastable alumina

In previous studies [7–9] we have shown that the deposition of Al from vapour phase onto an Fe-Cr-Al foil containing minor amounts of rare earth elements decreases the oxidation rate significantly. This effect was attributed to the =-Al2O3 grains which were larger than those formed on the conventional foils of similar compositions such as specimen C. It is well known that scales consisting mainly of =-Al2O3 grains grow predominantly by the diffusion of oxygen along oxide grain boundaries. The larger grains mean less grain boundary area available for the oxygen diffusion. Thus the oxidation rate was decreased. The deposition of Al was initially intended to increase the oxidation resistant life of the foil by increasing the total Al content. However, additional effect was found, i.e. the oxidation rate itself was decreased. It was also found that the Al deposition stabilised metastable Al2O3, which transformed from C- to G- to =-Al2O3 with the progress of oxidation. The =-Al2O3 grains transformed from G-Al2O3 are much larger than those quickly transformed grains. Therefore, the smallest oxidation rate constant obtained with specimen A is attributable to the large =-Al2O3 grains. 5.2

Closed micropores in α-Al2O3 grains

The formation of closed micropores in =-Al2O3 grains is attributable to the volume decrease accompanying the G- to =-Al2O3 phase transformation. This volume decrease is accommodated by the formation of these voids. A similar volume decrease takes place for the oxidation of NiAl, however in this case many cracks are formed [10,11] in the scale loosing protectiveness for certain periods until these cracks were buried with new oxides. 5.3

Appearance of spinel type oxides

Spinel type oxides appeared when the transformation to =-Al2O3, progressed to some degree. This can be explained in terms of difference in the solubility of Fe between in G-Al2O3 and =-

92 Al2O3. The former has a larger solubility that the latter. This view can also explain the observation that these spinel type oxides were not found near the scale/substrate interface, because the transformation starts at the interface and moves outwards.

6

Conclusions

[1] The pre-oxidation at 1173 K for 90 ks in air improves the oxidation resistance of Aldeposited Fe-Cr-Al foil containing minor amounts of rare earth elements by forming GAl2O3 scales which finally transform to =-Al2O3 scales during the subsequent oxidation at 1373 K. This effect is attributable to =-Al2O3 grains which are larger than those formed on the Al-deposited foil or the conventional alloy foil. [2] These =-Al2O3 grains contain many closed voids. This is explained in terms of the volume decrease accompanying the G to =-Al2O3 transformation. [3] Many small spinel type oxide particles appeared in =-Al2O3 grains. This is though to be due to the difference in solubility of Fe between G -Al2O3 and =-Al2O3.

7 [1] [2] [3] [4] [5]

References

M. Nonnenmann, SAE Tech. Paper Ser. 850131, (1985). P. N. Hawker, C. Jaffray and A. J. J. Wilkins, SAE Tech. Paper Ser. 880317 (1988). T. Takeda and T. Tanaka, SAE Tech. Paper Ser. 910615 (1991). S. Isobe, Denki-Seikoh, 58, 104 (1987). N. Hiramatsu, K. Miyakusu and Y. Uematsu, Nissin Steel Technical Report, No. 63, 145 (1989). [6] K. Ishii and T. Kawasaki, J. Japan Inst. Metals, 56, 845 (1992). [7] A. Andoh, S. Taniguchi and T. Shibata, Oxid. Met., 46, 481 (1996). [8] A. Andoh, S. Taniguchi and T. Shibata, Tetsu-to-Hagane, 83, 205 (1997). [9] A. Andoh, S. Taniguchi and T. Shibata, Tetsu-to-Hagane, 84, 285 (1998). [10] 10. G. C. Rybicki and J. L. Smialek, Oxid. Met., 31, 275 (1989). [11] 11. J. Doychak, C. A. Barrett, and J. L. Smialek, TMS Proc. Corrosion and Particle Erosion at High-temperatures, Ed. by V. Srinivasan and K. Vedula, (1989) 487.

Factors Affecting Oxide Growth Rates and Lifetime of FeCrAl Alloys W.J. Quadakkers (Sp), J. Nicholls*, D. Naumenko, J. Wilber*, L. Singheiser Forschungszentrum Jülich, D-52425, IWV-2, Jülich, Germany * Cranfield University, School of Industrial and Manufacturing Science, Cranfield, Bedford, MK43 0AL, Cranfield, U.K.

1

Introduction

Service lives of FeCrAl-based components at high temperatures are determined by the consumption of the alloying element aluminium, which is necessary to maintain the protective Al2O3-scale. After the alloy Al-content becomes depleted beneath a critical level, a catastrophic breakaway oxidation of Fe and Cr occurs [1]. In thin walled FeCrAl-components, such as e.g. car catalyst carrying foils, the Al-reservoir is depleted due to the scale growth, rather than by scale spalling, which is encountered in the case of components of 1 to 2 mm thickness [2]. Numerous publications issued recently on the oxidation of FeCrAl alloys have made substantial progress in explaining the beneficial effects of minor metallurgical additions of reactive elements (Y, Zr, Hf etc) on the oxide scale growth rate and adhesion [3,4,5]. However, considering the degradation of components fabricated from commercial alloys operating in real conditions is a complex issue, in which numerous factors must be taken into account. In this study a number of experimental results will be presented showing the effect of a typical minor addition of the oxygen active element titanium on the growth rate and adhesion of the protective alumina scale in the commercial and model systems. Furthermore, the oxide growth rate and thus the component life of FeCrAl-components will be shown to depend on the intrinsic mechanical properties of the alloy and/or application relevant component geometries. Finally, the effect of transient oxide formation on scale growth rate will be illustrated.

2

Experimental

The studied commercial and model alloys prevailed as cold-rolled sheets with a thickness varying from 0.05 mm to 2 mm. Some of the materials were delivered as hot extruded bars of 6 mm in diameter. The batch designation, sheet thickness and chemical composition analysed by Atomic Absorption Spectrometry (AAS) for each of the studied alloys are listed in Table 1. Coupons of 20 ´ 10 mm in size or discs of 6 mm diameter were machined from the asdelivered materials, eventually ground to 1200 grit surface finish or polished with 1µm diamond paste, depending on the experimental requirements. All specimens were ultrasonically degreased in a detergent before oxidation. Cyclic lifetime oxidation tests were performed in resistance heated furnaces in laboratory air at temperatures varying between 900 °C and 1200 °C with intermediate weighing of the specimens during cooling intervals. Isothermal

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

94 thermogravimetric (TG) experiments up to 200 h at temperatures between 900 °C and 1300 °C were carried out in synthetic air using a SETARAM® microbalance. The oxidised specimens were investigated using optical metallography, Scanning Electron Microscopy (SEM) with Energy Dispersive X-Ray Analysis (EDX) and Rutherford Backscattering Analysis (RBS). Table 1: Chemical compositions of the studied alloys (mass % or mass ppm) Aluchrom Kanthal AF PM 2000 YHf Supplier VDM Kanthal AB Plansee As received 1,0 mm sheet 1,0 mm 0,5 mm state sheet sheet Batch name FVK FVJ CKS Fe, % Base Base Base Cr, % 19.65 20.83 19.9 Al, % 5.53 5.23 5.3 Si (ppm) 2900 1900 200 Ca (ppm) – – 210 Hf (ppm) 310 3.1 – Mg (ppm) 78 17 – Mn (ppm) 1800 610 2100 P (ppm) 130 140 75 Ti (ppm) 98 940 4900 Y (ppm) 460 340 3700 Zr (ppm) 540 580 40 C (ppm) – – 240 S (ppm) 1.3 1.5 50 N (ppm) – – 143 O (ppm) – – 4611 V (ppm) 860 200 18 Cu (ppm) 110 330 2000 Nb (ppm) <50 <50 3.4 Mo (ppm) 100 58 97 Material

PM 2000 Plansee 1,0 mm sheet CKT Base 19.8 5.2 600 59 – – 1200 37 4900 3600 13 290 30 316 2582 55 580 2.8 29

Model ODS Ti-free Plansee Æ6 mm rod

Model ODS +Ti Plansee Æ6 mm rod

DAG Base 15.9 6.56 120 <10 <50 <10 480 <10 250 3800 110 260 30 218 2470 55 16 <50 –

DAH Base 16.6 6.33 150 <10 <50 <10 610 <10 4000 3800 110 200 70 206 2620 80 13 <50 –

3

Results and Discussion

3.1

Interaction of reactive elements Ti and Y during oxidation of FeCrAl ODS alloys

The commercial FeCrAl alloys normally contain several RE-additions, empirically found to provide better material properties in particular improving the oxidation resistance [6]. For example, the commercial ODS (oxide dispersion strengthened) alloys in addition to 0.5 mass% of yttrium oxide dispersion are usually doped with 0.5 mass% of titanium. In order to study the

95 effects of the second RE-element (Ti) addition, long term oxidation experiments were performed with model FeCrAl ODS alloys with and without Ti (Figure 1a). The poorer alumina scale adherence of the Ti-free model alloy led to a lifetime decrease by approximately a factor of two in comparison to the Ti-doped alloy.

Mass change / mg.cm

-2

4

0

-4

-8

a)

-12 0

3000

6000

9000

Exposure time / h

Mass change / mg.cm-2

1,2 ,m=0.19t0.38; no Ti; Alloy DAG

1,0 0,8

,m=0.25t0.30; 0.4% Ti; Alloy DAH

0,6 0,4 0,2

b)

0,0 0

20

40

60

80

100

Exposure time / h

Figure 1: Specific mass change data for model ODS alloys with (open circles) and without (filled triangles) Ti addition during: (a) lifetime cyclic air oxidation tests of 2mm thick coupons at 1200°C; (b) Isothermal oxidation at 1200°C in synthetic air with scale growth kinetics, calculated by fitting the measured data with power law function

During isothermal exposures for 100 h at 1200 °C the oxygen uptakes by the Ti-free and the Ti-doped alloys were very similar (Figure 1b), meaning that the alumina scale thickness on the two model materials was approximately the same. After mild cooling (10K/min) to room temperature large areas of the alumina scale on the Ti-free alloy were found to spall off, whereas the scale on the Ti-doped alloy remained perfectly adherent to the substrate (Figure 2). The SEM-image in Figure 3 indicates that the scale failure on the Ti-free alloy occurred during cooling and was related to formation of a “wavy” scale morphology rather than to intrinsically deteriorated oxide adhesion. Furthermore, the RBS data in Figure 4 clearly show that the yttrium enrichment on the top of the alumina scale occurs faster for the Ti-containing alloy, than for the Ti-free alloy.

96

a) b) Figure 2: Macrographs of model ODS coupons after 300 h cyclic oxidation at 1200°C in air: (a) Ti-free alloy DAG; (b) 0.4%Ti-containing alloy DAH

Figure 3: Spalled area of oxide scale formed on Ti-free model ODS alloy after 300 h cyclic oxidation (100 h cycles) at 1200 °C in air (SEM image in secondary electrons)

It is well established that additions of a reactive element (RE), such as Y either in metallic form or as an oxide dispersion improve the oxidation resistance of FeCrAl-alloys. Especially the alumina scale adherence seems to benefit from the RE-additions, although a number of other important RE-effects have been reported, including modifications in the scale structure and morphology [3,5]. Many of the mentioned positive RE-effects may be explained by the fact that the RE’s change the oxide growth mechanism by blocking the outward diffusion of aluminium [7]. This assumption has been verified by O18-tracer studies [8], where the alumina scales formed on the RE-doped FeCrAl-alloys were found to grow almost exclusively by inward oxygen transport. In the scale growing by oxygen diffusion the new oxide is formed nearly exclusively at the scale/metal interface. In this case the scale remains flat and adherent since no lateral oxide growth can be expected. In contrast, if the scale grows by mixed trans-

97 port mode of aluminium and oxygen, new oxide is formed within the scale, resulting in lateral growth and formation of convolutions [9]. 5 no Ti 0.4% Ti

Normalised Yield

4 Cr

3

Fe

2

Y

1 0 700

750

800

850

900

Energy channel

Figure 4: RBS-Spectra taken from the surfaces of model ODS alloys with and without Ti addition after 2 h oxidation at 1200 °C in air. High energy edges for Cr, Fe and Y are indicated in plot, showing higher Y-signal intensity for Ti-doped alloy.

Combination of all the above results leads to the conclusion that in the ODS alloy Ti enhances incorporation and mobility of yttrium in the alumina scale, which is a necessary requirement for achieving optimum cyclic oxidation performance. The positive effect is probably related to formation of a mixed Y/Ti-oxide on the alumina grain boundaries [8], however, the exact mechanism is still to be clarified. 3.2

Interaction of titanium with the alloy carbon and nitrogen impurities

It has been known for more than two decades that reactive element additions to FeCrAl alloys prevent the deleterious effect of sulphur on the scale adherence [3,4]. However, the REinteraction with other alloy impurities has hardly been considered. An interesting observation was made by the presenting authors recently when studying the oxidation of two batches of the commercial ODS alloy PM2000. One of the studied batches (CKT) exhibited a clearly higher oxidation rate during cyclic as well as isothermal exposures (Figure 5) than the other batch (CKS). The oxide growth kinetics on FeCrAl-alloys containing RE-additions can frequently be described by a sub-parabolic time dependence of the oxygen uptake, i.e. Dm=k.tn, (Dm - mass change; t-time). The experimentally determined exponent n in the above equation for dense alumina scales has usually a value of around 0.35 [7] and can be correlated with the columnar scale structure formed by the inward oxygen diffusion via grain boundaries. The latter scale growth mechanism apparently prevails during the oxidation of batch CKS (Figure 5-b). In contrast the scale formed on batch CKT exhibits a growth rate exponent of 0.5 indicating another mechanism contributing to the overall oxidation kinetics. Extensive analytical studies were performed on the two batches of which the results are presented in a separate paper [10]. The difference in the oxidation behaviour was attributed to the higher nitrogen content (Table 1) of the batch with the higher oxidation rate. The nitrogen impurity was found to be tied up into titanium (carbo)nitride precipitates in the alloy matrix. During long term high temperature

98 exposure these Ti(C,N) particles become incorporated into the inwardly growing alumina scale. Oxidation of these Ti(C,N) particles within the alumina scale results in microcracking and pore formation as illustrated in Figure 6.

Specimen mass change / mg.cm -2

12 1200°C; Air

10

CKT 0.90 mm thick

8

6

4

2

CKS 0.48 mm thick

a)

0 0

300

600

900

1200

1500

1800

Exposure time / h 5

Specimen mass change / mg.cm

-2

Isothermal Oxidation at 1300°C in synthetic air

CKT 37

4

3

CKS 33

2

1

b) 0 0

50

100 Exposure time / h

150

200

Figure 5: Mass change data of two studied batches (CKS and CKT) of PM2000 obtained during (a) lifetime cyclic air oxidation at 1200 °C and (b) isothermal air oxidation at 1300 °C

This mechanism proposed in reference [10] can explain the observed enhancement in the oxidation kinetics of batch CKT. An important property of the porous scale on batch CKT appeared to be its very good adherence to the alloy substrate. In spite of a high density of defects such as pores and cracks and a thickness exceeding 50 µm, the alumina scale on batch CKT could withstand the severe 20 h cycles oxidation conditions for several thousand hours of accumulated exposure time without being spalled. It seems to be that such a micro-cracked, defective scale can accommodate the large strains imposed by temperature cycling in the scale, e.g. by a cracking and re-healing mechanism as proposed in [11].

99

Figure 6: SEM cross section of oxide scale formed on batch CKT of PM 2000 after 1000 h cyclic oxidation at 1200 °C in air

The same effect of local scale disruption due to incorporation of carbide or carbo-nitride precipitates was observed due oxidation of the commercial wrought alloy Aluchrom YHf. The SEM studies of the formed alumina scales after cyclic air oxidation at 1200 °C support this idea (Figure 7).

Figure 7: Oxide scale formed on 1 mm thick coupon of commercial alloy Aluchrom YHf after 1000 h cyclic air oxidation at 1200 °C: (a) secondary electron image of the scale fracture surface; (b) Backscattered electron image showing transfomation of Zr/Hf-carbide (point A) into oxide (point B)

EDX analyses of the Zr/Hf-rich particles present in the inner part of the scale are apparently carbides formed by reaction between carbon impurity with minor additions of Zr and Hf in the

100 alloy. If the carbides become embedded into the scale, they are gradually transformed into oxides (Figure 7b). This results in void and crack formation in a similar manner as described for batch CKT of the ODS alloy PM2000. It should again be emphasised that this microcracking does not necessarily leads to decreased scale adherence. 3.3

Effect of the intrinsic mechanical properties of the alloy on the lifetime

Considering the stress generation and relaxation in the oxide scales on high temperature materials, the mechanical properties of the substrate metal are of great importance for oxide adherence. This is especially true in the case of the FeCrAl-alloys oxidation since the encountered service temperatures of the material frequently exceed 0.8 of its melting point. Plastic deformation or creep of the substrate alloy may thus contribute to the relaxation of the growth and especially thermally induced stresses [12]. Hence, the scales on thinner and intrinsically weaker substrates (e.g. wrought car catalyst foils) are expected to be more resistant to spalling, than those formed on thicker components (e.g. ODS-alloy tubes of a few millimetre thickness). Knowing the data on mechanical properties of the FeCrAl-materials and the alumina scales allows the definition of conditions (temperature drop, cooling rate etc.) at which the scales are prone to spalling. The different spallation modes can be derived from spallation maps as described in reference [12]. Unfortunately, as far as it is known to the authors, no comprehensive data base exists at the present time with respect to the properties of the numerous commercial alloys at high temperatures as well as for the alumina scales in their typical dimensional and structural features. Therefore, in many cases modelling of scale spalling is difficult. Consequently, to date prediction of the Al-depletion and the time to breakaway taking into account the scale spalling behaviour for a given component in particular service conditions is a complicated task. Comparison of experimental data with lifetime prediction for components of simple geometries was presented for several types of FeCrAl-alloys in reference [13]. The two following cases were considered in the lifetime modelling: (1) Al-depletion as a result of the oxide growth only and (2) Al-depletion as a result of the oxide growth superimposed with spallation of the scale after reaching a critical thickness. Recently an extended lifetime prediction model was presented [14], which incorporated a more detailed statistical description of the scale spalling process. Important parameters incorporated into the lifetime predicting equations are the critical scale thickness (or mass change) for spalling initiation (,m*) and a critical Alconcentration (CB) for occurrence of breakaway. Both of the latter parameters can be affected by the intrinsic alloy mechanical properties [14,15]. Figure 8 demonstrates the lifetime predictions compared to the experimentally determined lives for two commercial alloys having different mechanical properties. To account for the plastic stress relaxation by the alloy substrate, a much lower experimentally determined value of CB and a factor of two higher value of ,m* were taken for the wrought alloy Kanthal AF, than those for the stronger ODS alloy PM2000. As can be seen in Figure 8, the predicted lifetimes are in a reasonable agreement with the experimental data.

101 100000

without spalling

Time to breakaway / hours

with spalling

10000

"Weak" substrates: 2 CB=0.1%; ,m*=4mg/cm

1000

"Strong" substrates: 2 CB=1.21%; ,m*=2mg/cm

100

PM 2000 Kanthal AF

10 0,01

0,10

1,00

10,00

Specimen thickness / mm

Figure 8: Oxidation diagram for two commercial alloys at 1200 °C, showing effect of component wall thickness and substrate creep strength on time to breakaway. Predicted data (dashed and solid lines for “strong“ and “weak“ substrates respectively) are compared with experimental results (symbols).

3.4

Effect of mechanical constraints on the lifetime of FeCrAl-components

The oxidation behaviour of real components (e.g. a car catalyst substrate) can be different from that of free hanging laboratory specimens due to geometrically imposed mechanical constraints. Figure 9a shows that the scale formed at 1100°C on a free hanging commercial foil exhibited slower oxidation kinetics than mechanically constrained ring and model catalyst specimens manufactured from sheet of the same alloy batch. The macrophotograph of the model catalyst after the testing (Figure 9b) clearly shows that on part of the specimen chromia has formed and this chromia formation is accompanied by macroscopically visible scale cracking.

Mass change / mg.cm

-2

1,5

1

C

A B

0,5

0 0

200

400

600

800

1000

1200

1400

Exposure time / h

a) b) Figure 9: (a) Mass change data for Aluchrom YHf foil-based components during 100 h cycles oxidation at 1100 °C in air: A-model car catalyst; B-ring specimen; C-free hanging foil coupon. Tests were interrupted upon visual observation of green colouring due to formation of chromium oxide (indicated by arrows); (b) Macroscopic pictures of model car catalyst prior to exposure (left) and after 670 h oxidation at 1100 °C in air (right). Compare mass change data for specimen A in (a).

102 As it was claimed in the previous chapter, the oxidation induced stresses can be relaxed either by the plastic deformation of the substrate or by scale cracking. For free hanging, wrought foils the first relaxation mechanism will prevail. The reason for the accelerated oxide growth on the ring and model catalyst specimens can be related to the inability of the constrained substrate to relax the oxidation induced stresses by plastic deformation. No macroscopically visible cracks were seen on the specimen surfaces till chromia started to form. For the free hanging specimen this was observed at the same mass gain as for the constrained model components. Hence, in case of the constrained specimens, relaxation of the oxide stress most probably occurs by a scale microcracking mechanism described elsewhere [11]. Consequently, molecular oxygen transport through the microcracks results in accelerated scaling kinetics and consequently a reduced lifetime. 3.4

Lifetime of thin-walled FeCrAl components at temperatures below 1000 °C

In oxidation of FeCrAl-alloys, especially when used as thin foils for car catalyst bodies, the service temperature is a crucial parameter determining the lifetime since it directly affects the scale growth rate. Considering the Arrenius temperature dependence of the scale growth rate [16] a dramatic life-time extension for the thin foil FeCrAl-specimens is expected when the temperature is decreased from 1200 °C to 900 °C [17]. However, extrapolating the scaling kinetics obtained at high oxidation temperatures (1000 °C–1200 °C) to lower temperatures (800-900°C) can only be made if no change in the scale growth mechanism occurs within the temperature range considered. Isothermal testing of two commercial wrought foil materials at 900 °C revealed that one of them (batch HKF) exhibited initially higher oxidation rate than the other one (batch HKL, Figure 10). The mass change data in Figure 10 is in good agreement with the scale morphologies of HKF and HKL after long term oxidation presented in Figure 11. The thick outward growing scale on HKF contrasts with that on HKL, where a generally flat scale featured only a few areas with outward growing nodules. 0,8

Mass change / mg.cm

-2

,mB for 50 µm thick foil 0,6 HKF 0,4

,mB for 20 µm thick foil

0,2

HKL

0 0

10

20

30

40

50

60

70

Exposure time / h

Figure 10: Mass change data for two 50 µm thick FeCrAl foils during isothermal oxidation at 900 °C in air. The values for ,mB indicate the mass changes which will result in total Al-depletion and thus occurrence of breakaway oxidation.

103 The reason for this rapid initial scale growth at temperatures of around 900 °C is the formation of metastable modifications of alumina such as q. This results in an oxide growth rate, which is much higher than that of a-alumina [18]. Formation of the metastable alumina is apparently responsible for the enhanced oxidation kinetics of material HKF. Why alloy HKF is susceptible to the effect and HKL not has not been fully investigated and is subject of an ongoing study. First results indicate that it might be related to differences in formation of transient oxides from the other alloying elements. Of practical importance is the Al-depletion caused by the high oxidation rate of batch HKF. In Figure 10 the mass changes corresponding to the complete Al-consumption, i.e. breakaway oxidation are shown. Although, for 50 µm thick HKF-foils such a high initial oxidation rate might have no significant effect on the lifetime, for 20 µm thick foils it could lead to breakaway oxidation after just 40 h of exposure at 900 °C.

Figure 11: Scale cross sections of FeCrAl alloys HKF (a) and HKL (b) after 2000 h oxidation at 900°C in air

4

Conclusions

A number of factors must be taken into account when considering the lifetime oxidation performance of commercial FeCrAl-alloys. For achieving the best oxidation resistance with respect to oxide growth and adhesion it seems to be necessary that the minor alloy composition includes a combination of various reactive elements. For example, in yttria containing ODS alloys titanium addition appears to be of vital importance for maintaining optimum scale adhesion during cyclic oxidation. At the same time, the impurity elements, such as carbon and nitrogen should be kept at minimum levels in order to prevent enhanced oxidation due to incorporation of matrix carbo-nitride precipitates into the alumina scale. During cyclic oxidation the mechanical properties of the substrate material may significantly affect the scale spallation resistance and consequently the lifetime. It is not only the scale

104 thickness for spall initiation, which seems to be influenced by the mechanical properties of the metallic substrate but also the critical Al-concentration for the occurrence of breakaway (CB), i.e. CB decreases with decreasing alloy or component strength. For real FeCrAl-components the lifetime assessment based on the results of short term oxidation testing of laboratory specimens at high temperatures must be made with great care. If the FeCrAl-alloy is used as construction material in a constrained component, higher overall oxidation rates may occur. This results in shorter lifetimes than those predicted from the data obtained for free-hanging specimens in laboratory tests. At lower temperatures (» 900 °C) a further enhancement of Al-depletion and a shorter lifetime than expected from extrapolation of high temperature (1000 °C–1200°C) data can occur if the alloy is prone to formation of metastable aluminas.

5

References

[1] [2]

W.J. Quadakkers and M.J. Bennett, Mat. Sci. Tech. 1994, 10, 126–131. D. Naumenko, L. Singheiser and W.J. Quadakkers, in Proc. Int. Conf. on Cyclic Oxidation of High Temperature Materials (Ed.: M. Schütze and W.J. Quadakkers), Feb. 1999, Frankfurt am Main, Germany, 1999, 287–306. D.R. Sigler, Oxidation of Metals. 1989, 32, Nos 5/6, 337–355. H.J. Grabke, D. Wiemer and H. Viefhaus, Jour Appl. Surf. Sci. 1991, 47, 243–250. B.A. Pint, Oxidation of metals, 1996, 45, Nos 1/2, 1–31. J. Klöwer, A. Kolb-Telieps, M. Brede, in Proc. Int. Conf. On Metal Supported Automotive Catalytic Converters (Ed.: H.Bode), Oct. 1997, Wuppertal, Germany, WerkstoffInformationsgesselschaft mbH, Frankfurt, 1997, 33–46. W.J. Quadakkers, Werkstoffe und Korrosion 1990, 41, 659–668. W.J. Quadakkers, A. Elschner, H. Holzbrecher, K. Schmidt, W. Speier and H. Nickel, Michrochimica Acta, 1992, 107, 197–206. H.J. Grabke, M. Siegers and V.K. Tolpygo, Zeitschrift für Naturforschung, 1995, 50a, 217–227. W.J. Quadakkers, D. Naumenko, L. Singheiser, H.J. Penkalla, A.K. Tyagi and A. Czyrska-Filemonowicz, Materials and Corrosion, 2000, 51, 350–357. M. Schütze, Protective oxide scales and their breakdown, John Wiley & Sons Ltd, 1997. H.E. Evans, Int. Mat. Reviews, 1995, 40, No. 1, 1–40. W.J. Quadakkers, K. Bongartz, Werkstoffe und Korrosion, 1994, 45, 232–241. J.R. Nicholls, R. Newton, M.J. Bennett, H.E. Evans, H. Al-Badairy, G. Tatlock, D. Naumenko, W.J. Quadakkers, G. Strehl and G. Borchardt, Development of a life prediction model for the chemical failure of FeCrAl alloys in oxidising environments, in Proc. Int. Workshop on Lifetime modelling of high temperature corrosion process, 22–23 February 2001, Frankfurt am Main, Germany, in press J.P. Wilber, M.J. Bennett and J.R. Nicholls, in Proc. Int. Conf. on Cyclic Oxidation of High Temperature Materials (Ed.: M. Schütze and W.J. Quadakkers), Feb. 1999, Frankfurt am Main, Germany, 1999, 133–147. P.Kofstad, High Temperature Corrosion, Elsevier, 1988

[3] [4] [5] [6]

[7] [8] [9] [10] [11] [12] [13] [14]

[15]

[16]

105 [17] W.J. Quadakkers, in Proc. Int. Conf. On Metal Supported Automotive Catalytic Converters (Ed.: H.Bode), Oct. 1997, Wuppertal, Germany, Werkstoff-Informationsgesselschaft mbH, Frankfurt, 1997, 149–159. [18] B.A. Pint, J.R. Martin and L.W. Hobbs, The oxidation mechanism of q-Al2O3 scales, Solid State Ionics, 78, p.99–107 (1995)

On Deviations from Parabolic Growth Kinetics in High Temperature Oxidation G. Borchardt, G. Strehl Institut für Metallurgie, TU Clausthal, 38678 Clausthal-Zellerfeld, Germany

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

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108

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113

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116

Effect of Reactive Elements and of Increased Aluminum Contents on the Oxide Scale Formation on Fe-Cr-Al Alloys Vladislav Kolarik1, Angelika Kolb-Telieps2, Heike Hattendorf2, Maria del Mar Juez-Lorenzo1, Harald Fietzek1, Ralf Hojda2 1

Fraunhofer-Institut für Chemische Technologie, Pfinztal, Germany ² Krupp VDM GmbH, Werdohl, Germany

1

Introduction

Foils with a thickness of 50 µm of iron-chromium-aluminum alloys containing about 20 mass % chromium and 5.5 mass % aluminum are currently used as substrate in metal-supported automotive catalytic converters [1]. These alloys achieve a high oxidation resistance forming at high temperatures a thin, protective alumina scale [2]. Responding to the societal demand to reduce hazardous vehicle emissions an increase of the catalytic converters efficiency is of essential importance [3]. Thinner foil dimensions down to 30 µm are therefore required and hence the danger of an accelerated Al consumption and thus a shorter life time due to breakaway oxidation is more likely [4]. To reach the required life times using foils with reduced thicknesses, either the aluminum concentration has to be increased or the parabolic rate constant has to be decreased. New foil production techniques have been developed for achieving higher aluminum contents [5] and the effect of reactive elements on the oxidation behavior of Fe-Cr-Al alloys has been extensively investigated [2], [6]. In the present research work in situ – studies by high temperature X-ray diffraction have been performed in order to investigate the oxidation behavior at 950 °C and 1100 °C of two Fe-Cr-Al alloys with 5.5 mass % Al containing different reactive element additions and of a prototype alloy with 7 mass % Al as well as of an Al – coated material. The method allows an in situ identification of the oxides and their modifications, and it monitors the formation of each phase as a function of time[7].

2

Experimental

2.1

High temperature X-ray diffraction

The experimental set-up for the high temperature X-ray diffraction consists of an X-ray diffractometer and a high temperature device with a programmable temperature controller. Isothermal measurements and freely selectable temperature programs can be performed between room temperature and 1600 °C under oxidizing conditions. Series of X-ray diffraction patterns with defined time intervals or temperature steps are recorded in situ yielding the structural changes in the sample as a function of time or temperature. Phase changes, formation of new products and dilatation as well as contraction of the lattice are detected continuously during the experiment.

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

118 2.2

Kinetic evaluation of the in situ – measurements

For a kinetic evaluation of the oxide formation from a series of in situ X-ray diffraction patterns the intensities of the oxide peaks are determined as a function of time by a summing method. This procedure calculates the peak intensities summing the counts of each channel in the range of the peak. The background is subtracted to eliminate the influence of changing background intensity. The resulting intensity curves iz(t) for each oxide and its modification show their formation as a function of time taking into account the absorption of the X-ray beam in the growing oxide layer [7].

3

Investigated Alloys and Experimental Conditions

The in situ – studies by high temperature X-ray diffraction were applied to investigate the oxidation behavior of Aluchrom Y B and Aluchrom Y Hf with 5.5 mass % Al and of the prototype alloy E2 with 7 mass % Al as well as of the Al – coated Fe-Cr-Al alloy Aluchrom P (Table 1). Isothermal experiments of 100 h were performed in air at 950 °C and 1100 °C recording an X-ray diffraction pattern every 2 h. A scintillation counter with a secondary monochromator and fixed angle geometry with an incidence angle . = 10° was used with a measuring time of 1 h per pattern. After the in situ – experiments the sample surface was analyzed by scanning electron microscopy and the oxide scale was studied in cross-section by electron microprobe analysis. Table 1: Composition of the investigated alloys in mass % Alloy

Fe

Cr

Al

Si

Mn

Y

Zr

Ti

Aluchrom Y B Aluchrom Y Hf E2 Aluchrom P

Bal. Bal. Bal. Bal.

22.6 20.4 19.9 20.2

5.5 5.6 6.9 7.75

0.14 0.19 0.33 0.33

0.25 0.19 0.29 0.19

0.11 0.06 0.02 0.05

0.04 0.02 – 0.04

0.08 0.01 – 0.01

4

Other elements Hf Hf Hf

Results

The oxides identified on the investigated alloys at 950 °C and 1100 °C as well as the oxide scale thicknesses are summarized in Table 2. The alloys Aluchrom Y B and Aluchrom Y Hf with 5.5 mass % Al form =-Al2O3 and small amounts of FeAl2O4 at 950 °C. On Aluchrom Y Hf the additional formation of G-Al2O3 is found during the first 40 h. On further oxidation the G-Al2O3 transforms into the stable =-Al2O3. During the formation period of G-Al2O3 an accelerated growth of =-Al2O3 is observed, which flats down after the phase transformation at 40 h (Figure 1). With this mechanism Aluchrom Y Hf forms in the first 50 h a thicker oxide scale than Aluchrom Y B. With longer oxidation times however, the =-Al2O3 – intensity curve iz(t) of Aluchrom Y B still increases approaching to intensity values of Aluchrom Y Hf. The oxide scales on Aluchrom Y B after 74 h and on Aluchrom Y Hf after 100 h show a good adherence

119 to the substrate and a comparable thickness (Table 2). On Aluchrom Y Hf the side view of the plate-like structure figures of the G-Al2O3 are visible. (Figure 2). 2500

950°C Aluchrom Y Hf

2000

Intensity [cps]

1500

Aluchrom Y

=-Al2O3

1000

500

Aluchrom Y Hf

FeAl2O4

Aluchrom Y 0

G-Al2O3 Aluchrom Y Hf

-500 0

20

40

60

80

100

120

Time t [h]

Figure 1: Intensity curves iz(t) for the diffraction peaks of the oxides formed on Aluchrom Y and Aluchrom Y Hf at 950 °C in air

a) Aluchrom Y B

b) Aluchrom Y Hf Figure 2: Back scattered electron micrographs showing the oxide scales after oxidation at 950 °C on (a) Aluchrom Y after 74 h and (b) Aluchrom Y Hf after 100 h (2000:1)

At 1100 °C on both Aluchrom Y B and Aluchrom Y Hf =-Al2O3 and FeAl2O4 are formed and on Aluchrom Y B small amounts of FeCr2O4 are observed additionally (Figure 3). The intensity curves iz(t) of =-Al2O3 show a faster growth on Aluchrom Y B than on Aluchrom Y Hf, which even slightly increases at 70 h. The resulting oxide scale on Aluchrom Y B is with

120 12 µm notably thicker than the 7 µm thick scale on Aluchrom Y Hf. Both scales have a homogeneous thickness and a good adherence (Figure 4). Table 2: Oxides formed on the studied alloys at 950 and 1100 °C in the first 100 hours and the resulting oxide scale thicknesses Alloy

Al – content

950 °C Oxides

Aluchrom Y 5.5

=-Al2O3 FeAl2O4

Aluchrom Y 5.6 Hf

=-Al2O3 G-Al2O3 (40 h) FeAl2O4 =-Al2O3 G-Al2O3 (60 h) =-Al2O3

E2

6.9

Aluchrom P

7.75

Thickness 1.2*

1.7

1100 °C Oxides

Thickness 12

=-Al2O3 FeAl2O4 FeCr2O4 =-Al2O3 FeAl2O4

7

1.2

=-Al2O3

7

1.7

=-Al2O3

3.8**

*) after 74 h **) after 40 h 800

1100°C

Aluchrom Y

700

600

=-Al2O3

500

Aluchrom YHf

400

300

200

Aluchrom YHf FeAl2O4 Aluchrom Y FeCr2O4 Aluchrom Y

100

0 0

10

20

30

40

50

60

70

80

90

100

110

120

Time t [h]

Figure 3: Intensity curves iz(t) for the diffraction peaks of the oxides formed on Aluchrom Y and Aluchrom Y Hf at 1100 °C in air

121

a)

b)

Figure 4: Back scattered electron micrographs showing the oxide scales after 100 h oxidation at 1100 °C on (a) Aluchrom Y and (b) Aluchrom Y Hf (2000:1)

The prototype alloy E2 with 6.9% Al forms at 1100 °C the protective =-Al2O3 only. The growth rate is sub-parabolic and a continuously increasing intensity curve iz(t) is observed (Figure 5). The resulting oxide scale is well adherent and shows a homogeneous thickness of 7 µm after 100 h (Figure 6a). At 950 °C however, E2 forms the metastable G-Al2O3 in the initial state. The amounts are quite considerable and a complete transformation into =-Al2O3 occurs at 65 h. Simultaneously with the G-Al2O3 formation and its transformation into =-Al2O3 the latter shows an accelerated quasi-linear growth during the first 35 h. After the transformation the growth rate of =-Al2O3 decreases notably (Fig. 5). The oxide scale shows in the crosssection wide areas of spallation and on its top the figures of the plate-like ex-G-Al2O3 structure in quite high density (Figure 6b). The scale thickness measured between the ex-G-Al2O3 plates is 1.2 µm. The Al – coated alloy Aluchrom P forms at both 950 °C and 1100 °C the protective =-Al2O3 only. No other oxide phases are observed. The intensity curves iz(t) show an accelerated alumina formation during the first 5 h and a more moderate and continuous growth on further oxidation (Figure 7). The oxide scales show a good adherence to the substrate and a homogeneous thickness at both temperatures (Figure 8).

122 600

=-Al2O3

Intensity iz(t) [cps]

500

400

E2, 1100°C 300

200

=-Al2O3

E2, 950°C

100

3-Al2O3 0 0

10

20

30

40

50

60

70

80

90

100

110

120

Time t [h] Figure 5: Intensity curves iz(t) for the diffraction peaks of Al2O3 formed on the prototype alloy E2 at 950 °C and 1100 °C in air

a)

b) Figure 6: Back scattered electron micrographs showing the oxide scales on the prototype alloy E2 after 100 h oxidation (a) at 950 °C and (b) at 1100 °C (2000:1)

123

Aluchrom P 57962/BU022, 100 h

5000

=-Al2O3 1100°C

4000

3000

950° 2000

1000

0 0

10

20

30

40

50

60

70

80

90

100

110

120

Time t [h]

Figure 7: Intensity curves iz(t) for the diffraction peaks of =-Al2O3 formed on Aluchrom P at 950 °C and 1100 °C in air

a)

b) Figure 8: Back scattered electron micrographs showing the oxide scales on Aluchrom P (a) after 100 h at 950 °C and (b) 40 h at 1100 °C (2000:1)

5

Discussion

For optimizing the oxidation resistance of Fe-Cr-Al – foils for automotive catalytic converter supports both the aluminum content as well as the reactive element additions are parameters influencing the oxide scale formation. The two alloys with 5.5% Al, Aluchrom Y B and Aluchrom Y Hf, show the influence of the reactive elements on the oxidation. The Aluchrom Y B,

124 which contains no Hf and higher amounts of Y, Zr and Ti, tends to higher =-Al2O3 growth rates at 1100 °C resulting in thicker oxide scales and forms furthermore FeCr2O4 (Figures 3 and 4). The addition of Hf obviously reduces the =-Al2O3 growth rate at 1100 °C. At 950 °C however, it seems to favor the formation of the metastable G-Al2O3, although in very low amounts (Figure 1). The favoring of the G-Al2O3 formation by Hf – additions at 950 °C was also previously observed on >-NiAl [8]. Raising the Al – content to the range of 7% and taking out Zr and Ti suppresses the formation of other oxides besides alumina at both temperatures. After 100 h at 1100 °C the scale thickness is the same as that on Aluchrom Y Hf and it does not contain FeAl2O4 (Figure 6). At 950 °C however, an enhanced G-Al2O3 formation is observed in the first 60 h (Figure 5). After 100 h, when no G-Al2O3 is detected anymore by X-ray diffraction, the plate-like figures typical for G-Al2O3 are still observed (Figure 6b). This means that during the phase transformation the crystallographic structure changes, but not the morphology. The compositional modification performed on the prototype alloy E2 improves the protective oxide scale formation at 1100 °C, but it leads to an undesirable G-Al2O3 formation at 950 °C. The Al – coated FeCrAl – alloy Aluchrom P with a total Al – content of 7.75 % achieves the formation of a pure =-Al2O3 scale at both 950 °C and 1100 °C (Figure 7). The reactive element additions are practically those of Aluchrom Y Hf, but increasing the Al – content by the coating technique suppresses not only the formation of mixed oxides, but also that of the metastable G-Al2O3. Obviously, the final Al – content must be high enough for not forming metastable phases in the initial state like reported in [9]. Special attention should be paid to the formation of G-Al2O3 and its subsequent transformation into =-Al2O3. It is remarkable, that in the cases of an initial G-Al2O3 formation with a subsequent fast transformation into =-Al2O3 oxide scale spallation is found in wide areas. A transformation within 30 h as observed for the alloy E2 at 950 °C is already considered as fast (Figure 5). Smaller amounts of G-Al2O3 and/or a slow phase transformation are not or at least less detrimental for the scale adherence as observed for Aluchrom Y Hf (Figures 1 and 2). For reducing the foil thicknesses to 30 µm or even 20 µm the suppression of the G-Al2O3 formation is therefore essential for two reasons. First, the higher growth rate of G-Al2O3 depletes the Al in the foil faster than an =-Al2O3 formation only and shortens thus the life time. Secondly and may be more important, G-Al2O3 in an amount above a critical level transforming quickly into =-Al2O3 produces stresses in the oxide scale due to the volume difference of the two alumina modifications, which often leads to spallation. A foil with a faster depleted Al – content by an initial G-Al2O3 formation will loose the capacity of rehealing the alumina scale earlier with reduced foil thickness.

6

References

[1]

H. Bode in Metal-Supported Automotive Catalytic Converters (Ed. H. Bode), WerkstoffInformationsgesellschaft, Frankfurt am Main, Germany, 1997, pp. 17–31 J. Klöwer and A. Kolb-Telieps in Metal-Supported Automotive Catalytic Converters (Ed. H. Bode), Werkstoff-Informationsgesellschaft, Frankfurt am Main, Germany, 1997, pp. 33–46

[2]

125 [3] [4]

[5]

[6]

[7] [8] [9]

W. Maus in Metal-Supported Automotive Catalytic Converters (Ed. H. Bode), Werkstoff-Informationsgesellschaft, Frankfurt am Main, Germany, 1997, pp. 3–13 E. Wahlström and S Olsson in Metal-Supported Automotive Catalytic Converters (Ed. H. Bode), Werkstoff-Informationsgesellschaft, Frankfurt am Main, Germany, 1997, pp. 69–77 A. Kolb-Telieps, J. Klöwer, R. Hojda and U. Heubner in Metal-Supported Automotive Catalytic Converters (Ed. H. Bode), Werkstoff-Informationsgesellschaft, Frankfurt am Main, Germany, 1997, pp. 99–104 H.J. Grabke, M.S. Siegers and V.K. Tolpygo in Metal-Supported Automotive Catalytic Converters (Ed. H. Bode), Werkstoff-Informationsgesellschaft, Frankfurt am Main, Germany, 1997, pp. 139–148 V. Kolarik, W. Engel and N. Eisenreich, Mater. Sci. Forum. 133–136 (1993) 563 I. Rommerskirchen and V. Kolarik, Materials and Corrosion 47 (1996) 625–630 A. Andoh, S. Taniguchi and T. Shibata, High Temperature Corrosion and Protection 2000, 2000 Science Reviews, 2000, pp. 297–303

High Temperature Strength of Metal Foil Materials M. Cedergren, K. Göransson R&D, AB Sandvik Steel, Sandviken, Sweden

1

Abstract

The current trend for metallic catalytic converter steel foils is towards thinner thicknesses due to harder regulations. This leads to several problems; one is the demand on better oxidation properties because of less available aluminium per surface area unit. Another problem is that a material with better high temperature fatigue strength will possibly be needed. In the paper a number of methods for improving the high temperature strength of FeCrAl foils are suggested. The methods include strengthening by particles and solid solution hardening.

2

Introduction

Thin foils of ferritic Fe-Cr-Al alloys are today used as carrier materials for catalytic converters in the purification of exhaust gases from internal combustion engines. The usefulness of such an alloy lies in the formation of a thin, adherent aluminium oxide film on the surface. The surface oxide film protects the metal from rapid oxidation at the temperature at which the catalytic converter is used. In order to form such a protective oxide a minimum of 4.5% aluminium is necessary in the alloy. The protective properties of this aluminium oxide are known to be improved, especially with respect to thermal cycling, if the alloy contains small amounts of one or more of the so called reactive elements (RE), such as Mg, Ca, Zr, Hf or rare earth elements (Sc, Y or one of the lanthanide elements). Such alloys can be conventionally produced by melting, refining, casting, billet rolling or forging followed by hot and cold rolling to produce thin strips with a final thickness of 10 µm– 100 µm. The alloys can also be produced by using a pre-rolled strip with a lower aluminium content than the desired catalytic converter carrier material and deposit on the surface of this material a layer of an aluminium rich alloy. The deposition can be made in many different ways, e.g. by dipping the strip in a molten Al alloy [1], or by roll-bonding (cladding) an Al alloy on top of a ferritic steel [2]. Sandvik has also described how a PVD process can be used to deposit the Al rich alloy [3]. In all these examples, the thickness of the strip with deposit may be the final thickness, or the strip may be rolled down to a smaller thickness after the deposition has been performed. The deposited composite of ferritic alloy and aluminium alloy may be heat treated to produce a homogeneous alloy, or an alloy with an increasing aluminium concentration towards the surface.

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

127 The mechanical properties of ferritic Fe-Cr-Al alloys are known to be poor at high temperature. Several ways of improving these properties are known, such as the production of fine dispersions of oxide or nitride phases by powder metallurgical processes. These processes involve expensive operations during production and are hence not suitable for manufacturing of alloys that are to be used in large quantities. Another way of increasing the high temperature strength is by precipitating minute particles of nickel aluminides. This has been described for a number of Fe-Ni-Cr-Al alloys with a basically ferritic structure or a mixed austenitic/ferritic structure [4], [5]. However, such alloys are only known by the present authors in hot rolled or cast condition. If they are to be useful in catalytic converters, the alloys must be able to be formed by cold rolling, down to a final thickness of less than 100 µm. Fe-Ni-Cr-Al alloys with a basically austenitic structure are known to be possible to produce in thin sections. The nickel content of such alloys must be above 30%, which makes the raw material cost high. Also, the oxidation properties of such alloys are in general less good than those of ferritic alloys. It is known [6] that addition of approximately 10% nickel to a Fe-Cr-Al alloy furthermore improves the resistance to thermal shock, a phenomenon that is known to cause reduced life in catalytic converters. The purpose of the present work was thus to investigate high temperature properties and possible production routes for a ferritic Fe-Ni-Cr-Al alloy.

3

Experimental Procedure

The effects of several alloy modifications have been evaluated in terms of oxidation resistance, processability and high temperature mechanical properties. Some of the alloys are presented in Table 1. The alloys were produced by induction melting. The cast ingots were rolled to billets which subsequently were hot-rolled down to a thickness of 3 mm. Cold-rolling trials were performed. High temperature tensile tests were performed on all alloys between 600–1000 °C according to the standard SS-EN 10002-5. The Young’s modulus was measured directly using strain gauges mounted on the specimen. Oxidation properties of the alloys were evaluated at 1100 °C and 1200 °C in normal atmosphere. The samples were removed from the furnace at pre-set intervals and weighed in order to monitor the weight gain. Table 1: Chemical composition Element C Cr Ni Mo + W Al REM Zr + Hf

FeCrAl <0,02 20 <0,5 <0,02 5,5 0,04

5Ni2Mo <0,02 20 5 2 5,0 0,025

12,5Ni <0,02 20 12,5 <0,02 5,9 0,04

5Ni <0,02 20 5 <0,02 6,2 0,08 0,08

128

4

Results

The large scale microstructure of the alloys is identical to that of a normal FeCrAl. However, SEM and TEM analyses show that alloys of these compositions contain minute nickel aluminide particles with the CsCl type structure. The particles form evenly spaced within the ferrite grains. The hardness of the material after hot rolling is high: in the range 400–520 HV1. By annealing, the hardness in one case could be brought down from 490 to 320 HV1. The high hardness of the material has caused process disturbances, resulting in the large scale cold rolling trials having to be postponed. No such results can thus be presented at the present time. Figs. 1–4 show the measured high temperature mechanical properties of the alloys. The Young’s modulus of the experimental alloys is generally higher than that of the standard FeCrAl. One interesting effect is the measured increase of Young’s modulus in the two 5% Ni alloys above 900 °C. 160

Elongation (%)

140

FeCrAl 5Ni2Mo

120

12,5Ni

100

5Ni

80 60 40 20 0 550

600

650

700

750 800 850 Temperature (ºC)

900

950

1000

1050

Figure 1: Elongation at fracture

In the temperature range below 750 °C–800 °C, the experimental alloys have greater mechanical strength than the FeCrAl. However at higher temperatures the difference between the alloys is within the experimental uncertainty of the equipment used, with one exception. The yield strength of the 12,5Ni alloy is significantly higher at 900 and 1000 °C than that of the other alloys. The experimental alloys show consistently less elongation at fracture (cf. Fig. 1). This effect is highest in the highest Ni alloy and in the Mo alloyed alloy. As regards oxidation resistance this is shown in figs. 5–6. At 1100 °C, the 5Ni2Mo alloy shows a lower weight gain than the FeCrAl. At 1200 °C, the three experimental alloys behave individually different: the 5Ni alloy shows a normal oxidation behaviour with a 30% lower total weight gain than the FeCrAl. The 5Ni2Mo alloy starts out with a low weight gain but starts spalling after a short initiation period. The rate of spallation is comparatively low and

129 only starts to accelerate after 350 h. The rate of spallation is significant from the beginning for the 12,5Ni alloy, although it appears to decline after some period of operation. 140

5Ni2Mo 5Ni 12,5Ni FeCrAl

E-modulus (GPa)

120 100 80 60 40 20 0 550

600

650

700

750 800 850 Temperature (ºC)

900

950

1000

1050

Figure 2: Young’s modulus

700 FeCrAl

600

5Ni2Mo 12,5Ni

500 Rm (MPa)

5Ni

400 300 200 100 0 550

600

Figure 3: Tensile strength

650

700

750 800 850 Temperature (ºC)

900

950

1000

1050

130 500 FeCrAl 5Ni2Mo

400

Rp0,2 (MPa)

12,5Ni 5Ni

300

200

100

0 550

600

650

700

750 800 850 Temperature (ºC)

900

950

1000

1050

Figure 4: Yield strength 20

18

16

Mass change [g/m2]

14

12

5Ni2Mo FeCrAl

10

8

6

4

2

0 0

50

100

150

200

250

300

350

400

Time [h]

Figure 5: Oxidation test 1100 °C

5

Discussion

The high hardness of the material is partially due to the presence of Ni aluminides. A calculated phase diagram section for the system Fe-Ni-20Cr-5Al is shown in fig. 7. The phase diagram was calculated with Thermocalc [7]. It shows that NiAl is likely to be stable even at very

131 low Ni contents in the alloy. The dissolution temperature of NiAl is approximately 900 °C for a 5% Ni alloy and 1050 °C for a 12,5% Ni alloy. No austenite is expected to form below a total Ni content of 14%. The lattice parameter mismatch between NiAl and ferrite in equilibrium is expected to be small, and precipitation of NiAl appears to occur coherently. The presence of NiAl in the 12,5Ni alloy in the hot tensile tests above 900°C explains the improved yield strength. 50,00

40,00

30,00

Mass change [g/m2]

20,00

10,00

0,00 0

50

100

150

200

250

300

-10,00

-20,00

-30,00

-40,00

-50,00

Time [h]

Figure 6: Oxidation test 1200 °C

a

a+g

a+g+NiAl

g

g+NiAl g+g’

a+NiAl

a+g+NiAl+s

g+g’+NiA l

Figure 7: Section through the Fe-Ni-Cr-Al phase diagram at 20%Cr, 5% Al

350

400

5Ni2Mo FeCrAl 12,5Ni 5Ni

132 The unexpected temperature dependence of the Young’s modulus for between 900 and 1000 °C for two of the alloys can not be explained by the authors, however it may be connected with the dissolution of NiAl. The actual numbers for the Young’s modulus are however still much higher than those for the standard FeCrAl. It must be noted that measurement of the Young’s modulus is less accurate at high temperatures than at room temperature. The mechanical strength is improved below 800°C. At higher temperatures, the effect is less clear. The strengthening effect of Mo appears to be small above 600 °C with respect to the yield strength. In order to evaluate the usefulness of these alloys in practical applications, high temperature fatigue tests as well as creep tests are probably necessary. The inital tests performed in the present paper, however, indicate that these alloys are promising candidate materials for catalytic converter bodies in mechanically challenging applications. The oxidation properties of the experimental alloys are unexpectedly good; in several cases superior to that of standard FeCrAl. In other cases, spalling is found, although the rate of spallation is not too serious for possible use of the material in other applications than catalytic converters.

6

Conclusions

By adding Ni (2,5%–15%) and Mo+W (<3%) it is possible to improve the high temperature strength compared to FeCrAl catalytic converter steel, without deteriorating the oxidation resistance. The alloy may also be useful in other high temperature applications, e.g. in heat treating furnaces. The material is extremely difficult to manufacture by conventional production due to brittleness, thus a prefered manufacturing is by coating a low-Al alloy with Al in one of the final steps in the production. The coating may be applied by e.g. dipping, cladding or a PVDprocess. The alloys of the present work are basically ferritic Fe-Ni-Cr-Al alloys strengthened by the presence of minute particles of nickel aluminides and if necessary further strengthened by the presence of substitutionally dissolved elements such as Mo or W. Owing to a high Al content and the presence of reactive elements, the resistance to oxidation at high temperatures is good. Thus, this is a suitable alloy for use as a carrier material in metallic catalytic converters, especially such that are exposed to a combination of high temperature and mechanical load.

7

Acknowledgements

This paper is published by permission of AB Sandvik Steel. The authors wish to express their gratitude to Dr. T. Thorvaldsson for placing the excellent R&D facilities at our disposal and also for encouragement. We also thank the staff at AB Sandvik Steel that have been involved in the project, especially Mr. Lennart Törngren, Dr. T. Helander, Mr. Håkan Myrberg and Mr. Stefan Larsson.

133

8

References

[1]

A. Kolb-Telieps, J. Klöwer, R Hojda, U. Heubner in Metal-Supported Automotive Catalytic Converters (Ed.: H. Bode), Werkstoff-Informationsgesellschaft, Frankfurt, Germany, 1997, p 99. I. M. Sukonnik, S. Chang, B. Jha in Metal-Supported Automotive Catalytic Converters (Ed.: H. Bode), Werkstoff-Informationsgesellschaft, Frankfurt, Germany, 1997, p 93 and references therein. Andersson-Drugge, I.-M., WO 98/08986, 1996. J. H. Davidson, Electricité et Progrès. Une réalité dans l’industrie des métaux, 1978. Hamada T., Yamada S., Tsuji E., Mizukoshi T., US Patent 5,089,223, 1990. Bartlett C. A., Holt W. H., US Patent 4,780,276, 1987. B. Sundman, B. Jansson, J-O Andersson, Calphad 9, 1985, p 853.

[2]

[3] [4] [5] [6] [7]

Lifetime Predictions of Uncoated Metal-Supported Catalysts via Modeling and Simulation, based on Reliable Material Data Hans Bode University of Wuppertal (D)

Christian Guist BMW AG, Munich (D)

1

Introduction and Purpose of Research Work

The complex interrelations which determine the mechanical stability of metal catalyst supports cause high levels of time and cost investments in establishing assurances with regard to lifetime predictions via experimentation. For this reason, the emphasis is shifting with regard to assurance in the direction of simulation models. In order to be able to make precise statements using simulations, exact descriptions of the relationships concerned and known facts with regard to parameters, especially material parameters, are of major importance. The method used to establish these relationships and parameters will be explained in the following. The method has been validated using a reference system.

2

Construction of a Model for Lifetime

The approach here is to bring the task and technical system onto a virtual level and to construct a model. This model should be able to offer: · Precise concept design · Better evaluation of experimental results · Efficient transfer of knowledge and experience to subsequent projects · Specific error search in the case of damage etc. The created and proposed model is shown in Figure 1 and is be explained as follows: The load profile of the component is given in the left part of the diagram, considerations with regard to mechanical strength of the component (stability of component) are given on the far right-hand side. A comparison between load and stability leads to the lifetime prediction based on the virtual level in the much needed very early stages of product development. In the construction design of the model, thermal load, flow-related load and vibration-related load are considered. Taking a closer look at thermal load, hot exhaust gas flows through the catalytic converter during vehicle operation, imposing thermal loading of the metal support. Also the catalytic reaction leads to increasing thermal loads. Speed of hot exhaust gas, pressure of gas and temperature will result in flow-related load. Unfortunately, thermal- and flowrelated loads are not uniform, leading to additional loads. Acceleration and spectra of vibration will result in vibration-related load and here again, these parameters are not uniform.

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

135

Figure 1: Model

Influencing parameters with regard to load are type of engine and vehicle, as well as driving conditions and the geometry of the catalytic support. Parameters with regard to the mechanical strength of the component are: alloy composition, brazing – if applied, the way in which the component is shaped and eventually also welding. The parameters of the vehicle and driving conditions, considered in the research work related to lifetime prediction are: untreated exhaust gas, speed, pressure, temperature, acceleration and frequency spectrum. For a selected reference system, values for stresses and strains of loads have to be calculated or measured and these values have to be compared with permissible stresses and strains of the component. Material data will become very important and here, not only yield and tensile strength, but also dynamic strength and deformation characteristics have to be taken into consideration. Special measurement techniques, such as holography for deformation characterization (strains), have to be applied. Additional details can be found in [1].

3

Reference System

The reference system (Figure 2) consists of a BMW 6-cylinder engine (left side) and an EMITEC closed coupled catalyst-system (center). Each support consists of an outer shell, often called the mantle and an inner matrix (example right side). The matrix is composed of flat and corrugated foils. In certain areas these are brazed together at their contact points. Details of this kind of support are given by [2.The matrix is also brazed to the mantle in various possible ways. Therefore, properties of foils in the “asdelivered” state and in the “brazed” state are different or in other words, the actual supports have different properties to the thin foils in the “as-delivered” state. The geometrical data of

136 the supports is also provided. The ferritic foil material used in the materials tests was Aluchrom ISE (see [3]) in the “as-delivered” and in the “brazed-simulated” condition. Components composed of these foils were always temperature brazed with a Nickel-based material.

Figure 2: System of reference

In order to be able to quantitatively determine the input values for stresses and strains, these values have to be derived for the reference engine. The relationships explained in the following are valid with regard to new systems. The following groupings can be derived: · Vibration-related load is vibration due to inertial forces and flow loads, leading to stress values, and both values can be grouped together. · High frequency radial load is natural vibration (self vibration), leading to a value for deformation characteristics, i.e. strain. · Thermal load, leading again to a value for stress.

4

Determination of Axial Load and Comparison with Mechanical Strength Stability

Flow-related load (Figure 3) has been determined by numerical simulation (CFD). Values of flow-related loads were found to be very small (in the order of less than 1 N/mm2 in the reference system). Details of that calculation are given in [1].

137

Figure 3: Load caused by flow and vibration

The vibration-related load on the substrate in axial direction is caused by excitation by the engine, the road surface and gas pulsation. Excitation by the engine is the most extreme of these, in terms of its acceleration and frequency spectrum. It has therefore been considered as the only significant source of vibration. The effects of this vibrational excitation have to be examined from two different points of view: Inertial forces and natural vibrations (self vibration), see Chapter 5. If the catalytic converter is viewed as a rigid physical system, acceleration leads to the exertion of inertial forces mainly in an axial direction. These forces generate stress at the connecting points between the mantle and the matrix (Figure 3). The tension (Jload) is calculated from the inertial force and the load bearing section. The load bearing section is the sum of the foil sections that are brazed to the mantle. To compare load with stability, Jload can be used directly as can comparative stress (Iv). According to the shear stress hypotheses, the comparative stress is two times larger than the inertial load. These values have to be compared with the strength of the foil material in the “as-brazed” condition. Two possibilities exist for the determination of strength: · Determination of allowable stresses due to dynamic testing and / or · taking the yield stress as reference, divided by factor 3 for dynamic loadings [4]. Figure 4 shows the results of dynamic testing of foils in a heat-treated condition (left part) and the yield strength (static testing) as a function of temperature (right part). The heat treatment parameters correspond to those applied during brazing of supports. Parameters in the dynamic material tests are given in the figure on the left-hand side. The test was not conducted up to sample breaking point but to an allowed deformation of 2% due to

138 the requirement of cell stability and adherence requirements of ceramic coatings placed on the supports prior to starting service life. This was the limiting factor and the diagram is therefore a modified Smith-Diagram. The reduction of yield strength due to the brazing process is significant, but the differences in allowable max. I-amplitudes are small between “brazed” simulated samples and samples in an “as-delivered” state, i.e. the brazing parameters do not significantly influence the dynamic properties but rather the static properties. The comparison of high frequency axial load in the reference system with the mechanical strength stability of the material is also shown in Figure 4. Values of JLoad in total (flow-related and inertial-related) are about 4 N/mm2 for the first support and they are equivalent to a comparative stress sV of 8 N/mm2. These load values from the reference system are much lower than the maximum permissible values in the modified Smith diagram for the whole temperature range. This is confirmed by using or Jpermissible as a function of temperature.

Figure 4: Comparison of load and stability

Concerning Rp0.2/u as a function of temperature, there is an intersection of allowable shear stress with load stress above 900 °C. According to experiences in service, the critical point is the area close to the mantle. However, temperature of 900 °C will not be reached in this area, so flow forces and inertial forces are not significant and will not limit the lifetime of the reference products.

139

5

Determination of High Frequency Radial Load and Comparison with Mechanical Strength Stability

Because high frequency vibrations occur along the radial axis as well, this means that natural frequencies now have to be considered. By considering the catalytic converter as a complex spring-mass system instead of a rigid body (discussed previously), the result will be that a number of different natural frequencies occur along with corresponding deformation effects. The following approach has been chosen for determining these deformation effects: Only one of four modes has been selected because it is considered to be most critical [1]. At the beginning of the analyses of a given engine, the frequency spectrum initiated in the catalytic converter (support) and the energy required to produce excitation of the different frequencies has to be determined. The natural frequencies of the catalytic converter and the corresponding deformation effects are determined in parallel. The deformation in form of extension- or compression- related strains, which occurs during vehicle operation, is then determined because the stress as a result of this deformation is difficult to calculate. Therefore, strains are used as comparative values instead of determining stresses. Deformation can be determined in a planar manner by means of holography (Figure 5, righthand side) as a function of frequency. The deformation-related strain values determined in this way are in the order of less than 1 mm/mm. Details are given in [1]. Figure 5, left-hand side shows determined permissible deformation values (here elongation). Brazed samples have been tested as a function of different phase relationships.

Figure 5: Comparison of load and stability

140 Phase relationship describes how the sinus waves are arranged. Wave crest over wave crest means 0 degrees, wave crest over wave trough means 180 degrees. Tests have been carried out at room temperature and 800 °C. These tests contained samples composed of 2 layers or 5 layers. However, 5 layers were only used at 800 °C just in order to show the tendency in mechanical behavior. Forces and calculated stresses were applied in tests in order to reach 107 cycles. The maximum strain value was then measured. It can also be seen from Figure 5-lefthand side, that the phase has a very strong influence on mechanical strength stability. Lowest stability is given at 90 degrees and higher. The maximum permissible strain for reaching 107 cycles is about 6 mm/mm at room temperature and 5 mm/mm at 800 °C. When the number of layers is increased to 5, the permissible strain to reach 107 cycles at room temperature is lower, here 4 mm/mm. Further tests at different temperatures and with other numbers of layers should be carried out. By comparing the degree of extension-related strain expected during vehicle operation with the permissible degree of strain at room temperature for the two supports used in the reference system it has been established that natural frequencies will not lead to plastic deformations. The investigations have been confirmed experimentally.

6

Determination of Thermal Load and Comparison with Mechanical Strength Stability

The connections between the catalytic converter and the engine and its position in the vehicle determine the distribution of flow and the boundary conditions of heat transfer. A gradient of temperature between the outer shell and inner parts of the matrix has to be considered as well as transient temperature distributions due to the transient nature of vehicle operation. In addition, the maximum temperature has to be taken into account for thermal loading. The nonuniform thermal load will have an influence on a variety of properties, such as physical and mechanical properties of the support, elongation/extension of the support and oxidation/corrosion of the support. The resulting stresses will not be uniform. In view of the catalytic converter’s complex geometry, the load was not modeled by means of FEM programs. A separate mathematical model has been developed just for calculating thermal stresses in order to describe the characteristic status. It contains the principal contributory effects. Figure 6 shows how stresses resulting from thermal load were determined. A spiral wound support was used for the model. Sinus-shaped cells were taken as quadratic cells. The difference between outer radius Ra and inner radius Ri is the length of such a quadratic cell. If a temperature is applied, Ra and Ri will change, so A can be calculated (Figure 6, left-hand side). When such a segment is heated up, the outer shell (mantle) will impose pressure tensions on the cells. When cooled down, tensile tensions occur, starting at the point of pressure deformations which have possibly been caused previously. By simply considering the warm-up scenario, the cell is now viewed as a cylindrical container and the comparative stress can be calculated according to the given rule [5] for sV in Figure 6. The relationship between e and pa has been established in [6] with regard to matrix structures. The temperature dependence has to be taken into consideration. The relative change in the elasticity modulus is seen in Figure 6-center. In order to calculate the comparative stress it has to be multiplied by the temperature factor for the E-module as stated here. The relevant

141 property for the stability of the component is the yield strength of foils in the “as-brazed” condition. Their dependence on temperature has already been stated in Figure 4.

Figure 6: Determination of load (stress)

The comparison of load and stability is made as a function of cross section by referring to the radii in the previously given equation for the comparative stress and leads to the following prediction of deformation occurrence for an extreme thermo-shock loaded system, not the reference system. It emerged, as shown in Figure 7, that permanent deformation is to be expected in the outer zone, since Rp0.2 is less than the actual comparative stress. Plastic deformation occurs. The result was confirmed in experiments. Deformation caused by thermal effects is the most critical type of load on catalytic converters. An example for a damaged matrix is given in [2]. The objective of development work is to reduce this deformation to a minimum so that no changes in the catalytic converter’s characteristics occur. The mathematical model now provides a tool by means of which this can be done. Figure 8 shows the relevant parameters at the critical point close to the casing for catalyst 1 (outer zone) of the reference system. The related values for yield strength (tension and pressure) at the critical point can be seen. Comparative stress is also shown and, with regard to the reference system, is much lower than the permissible stress. No damage is predicted. This result has been verified by actual tests with the reference system.

142

Figure 7: Comparison of load and stability (cross-section, extreme thermo shock loaded system)

Figure 8: Comparison of load and stability (critical point, reference system)

143

8

Conclusions

As has been shown, it is indeed possible in spite of the extremely complex interrelationships, to make well-founded prognoses with regard to the mechanical stability of metallic catalyst supports. The model illustrated forms the basis for this. However, additional specific material values will be required. Values taken from literature are not suitable for accurate lifetime predictions. Using this work as a basis, the influence of further parameters can be investigated, for example, the influence of lesser foil thickness and higher cell densities as well as the influence of coating composition and coating thickness. Also the changes in foil properties with time is of importance. Due to the findings established as part of the investigations described in the above, the use of metals as catalytic supports does allow calculations about lifetime predictions and this is a major advantage when modifying or developing new supports in the shorter times now required.

9

References

[1] [2] [3]

C. Guist, Dissertation D468, Bergische Universität-GH Wuppertal, 1998 H. Bode in Conference Proceedings MACC 1997 (Ed. H. Bode), Wiley VCH, 1997 J. Klöwer, H. Bode, M. Brede, R. Brueck, L. Wieres, Werkstoffwoche, München, (Materials Week, Munich) Section 2, 1998 K. Giek, und R., Technische Formelsammlung, (Collection of Technical Formulae) Giek Verlag, 1989 AD-Merblatt B6, Zylindrische Mäntel unter äußerem Druck, (Cylindrical casings under external pressure) 1977 J. Pollack, Diplomarbeit (Thesis), TU Dresden, 1996

[4] [5] [6]

Elastic-Plastic Thermal Stress Analysis for Metal Substrates for Catalytic Converters Shogo Konya Nippon Steel Corporation, Futtsu, Japan

Atsushi Kikuchi Nippon Steel Technoresearch Corporation, Futtsu, Japan

1

Introduction

Emissions regulations for automotive vehicles have been more stringent for these years. Especially, reducing hydrocarbon emissions during cold start has been important so catalytic converters have been used at the position closer to engines where the exhaust gas temperature is high in order to improve light-off property. However, as catalytic converters are exposed to higher temperature, more severe thermal load is applied to catalytic converters. Therefore, higher durability of catalyst substrates has been required. Core / Jacket joint

Metal honeycomb-core Brazed

The Strengthened outer layers

Non-brazed

Jacket

The displacement Metal support A Displacement before 150 cycles

The pushing-out Metal support B No displacement after 900 cycles

Metal support C No displacement after 900 cycles

Figure 1: Three types of joint structure of metal substrates (by Takada et al.) [2]

Particularly in metal substrates, mechanical durability depends upon their joint structures [1–2]. It is considered that thermal stress simulation, using such as finite element method

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

145 (FEM), is effective for mechanical durability evaluation and optimal design of metal substrates. However, because the honeycomb structure of a metal substrate consists of a lot of layers of flat and corrugated stainless steel sheets, modeling the structure with complete fidelity causes excessive increase of the number of elements in FEM. Therefore, stress and strain to be generated in honeycomb structures are often simulated by replacing flat and corrugated sheets with equivalent solid elements [3]. However, since it is difficult to calculate stresses or strains at specific points using the equivalent elements, the authors considered that elastic-plastic FEM analysis with shell elements which directly represent flat and corrugated sheets is necessary to investigate deformation behavior caused by thermal loads within the honeycomb structure in detail. Takada et al. reported the durability of several joint structures of metal substrates against heat cycles [2]. The metal substrate (C) shown in Figure 1, which had an asymmetrical joint structure with “Strengthened Outer Layer”, had good mechanical durability, while the durability of the substrate (A) having no Strengthened Outer Layer was not enough. In this study, effects of joint structures of the metal substrates on mechanical durability are discussed using the results of elastic-plastic FEM analysis.

2

Model

MARC ver. K7 (Trademark of MSC. Software Corporation) was used as the codes for thermal elastic-plastic analysis in this study. (a) Top view

(b) Squint view

Figure 2: Fan shaped periodical symmetric model

Metal substrates in which flat and corrugated sheets of Fe-20mass%Cr-5mass%Al stainless steel – generally used for metal substrates – wound spirally in outer jackets of AISI430 steel were analyzed in this study. The honeycomb structure was assumed as a sector part of a fan

146 shaped periodical symmetric model as shown in Figure 2 considering computing speed and efficiency. In the periodical model, the honeycomb structure was expressed as a structure that consists of 30 flat sheets and 29 corrugated sheets arranged alternatively. As shown in Figure 2(a), wave peaks of the corrugated sheets were arranged in order in the G direction along the edge of right side, while cell shapes did not hold their original shapes along the left side edge. The outer jacket, flat sheets, and corrugated sheets were represented by bilinear thick-shell elements with four nodes, which were intersected into three layers in the thickness direction. The element length in the qdirection was small where the large amount of bending of the sheets would occur, and large in the z direction where the influence of bending would be comparatively small as shown in Figure 3.

Substrate (A) and (B)

Substrate (C)

Figure 3: FEM elements for thermal stress analysis

In this model, all joints between the outer jacket and the honeycomb core or between flat and corrugated sheets were assumed to be joined by brazing and mechanical properties in the brazed areas where Ni-based materials are often applied were assumed to be the same as those of Fe-Cr-Al alloy. While the brazed area consists of the flat and corrugated sheets and brazing material in between in reality, it was regarded as one sheet of Fe-Cr-Al alloy with double thickness, 60 mm in this model. The nodes of the corrugated sheets were tied with the nodes of the double thickness sheet under complete constraint. Brazed width was set around 370 mm. The nodes of the elements for the outer jacket at the area brazed with the honeycomb core were also tied with the confronting nodes in the outermost flat sheet elements. For the elements where the flat and corrugated sheets or the outer jacket and the honeycomb core were not brazed, contacts of those elements were judged using nonlinear springs connecting the nodes of the flat sheet elements with those of apexes of the corrugated sheets, or with those of the outer jacket. Spring constants were provided only in the r direction and zero in the G and z directions. Three types of substrates shown in Figure 1 were analyzed in this study. As joint structures, all flat and corrugated sheets were joined 20 mm in depth from the gas inlet side surface for the substrate (C), further those sheets in the gas outlet side were also joined for the substrate (A)

147 and (B). The outermost three layers (three corrugated sheets and four flat sheets) were joined in the whole length in the z direction in the substrate (B) and (C) as the Strengthened Outer Layer cited by Takada et al. [2]. The honeycomb core and the outer jacket were joined with the length of 25 mm at the gas outlet portion in the substrate (C) and at the central portion in the substrate (A) and (B). The dimensions of the substrates are summarized in Table 1. Table 1: Dimensions of metal substrates Dimension 80 mm 100 mm 30 mm 1.5 mm 1.25 mm 2.5 mm 62 cells/cm2 (400 cells/inch2)

Diameter Height Thickness of the flat and corrugated sheets Thickness of the outer jacket Wave height of the corrugated sheets Wave pitch of the corrugated sheets Cell density

As mechanical properties, thermal expansion coefficients, Young’s moduli, and yield points dependent on temperature were applied from the measured data. Von Mises yield criteria and work hardening coefficient dependent on Young’s modulus were applied for plastic analysis. The temperature distributions in the metal substrates were given from the measured data obtained from a heat cycle test using an engine bench as shown in Figure 4, which was applied as periodically symmetric one. 1000

Temperatute T / oC

480sec(Tmax)

center

900 800 700

480sec(Tmax)

84sec(,Tmax) 570sec

5th 1st

84sec(,Tmax)

jacket

600

570sec

500 400 1300sec

300

1300sec

200 100 0 0

250

500

750

1000

1250

0

10

Time t / sec

a)

20

30

Distance from the center d / mm

b)

40

0

10

20

30

40

Distance from the center d / mm

c)

Figure 4: Temperature distribution applied for simulation, (a): temperatures as a function of time at the gas inlet side. The ordinal numbers mean the order of corrugated sheets from the outer side. (b) and (c): temperature distribution in the r direction, (b): gas outlet side, (c): gas inlet side.

In the heating stage, the temperatures in the peripheral region were relatively low and those near the central portion were high. Especially, steep temperature gradient appeared from the outermost layer to the fifth layer. The temperatures at the gas inlet portion were higher than those of the outlet side in the peripheral region. The maximum temperature difference between the outer jacket and the central portion (DTmax) appeared at around 84 seconds from the begin-

148 ning of a heating stage. Temperature reached maximum at 480 seconds (Tmax) and temperature differences both in the r and z directions were small then. In the cooling stage, the honeycomb core was cooled by low temperature idling gas for 90 seconds. Therefore, the temperatures became lower and the temperature gradient in the r direction was inversed at the gas inlet side. After the engine stop, the temperature became higher in the central portion and lower in the peripheral portion again because the substrate was cooled in an ambient atmosphere. It takes 1300 seconds for 1 cycle in the heat cycle test.

3

Results of Analysis

3.1

Effect of asymmetric joint structure

Figure 5 shows the displacement in the z direction of the flat sheets in the substrate (B) and (C) at both surfaces of gas inlet and outlet side at the time the maximum temperature difference was generated within the metal substrate in the heating stage (DTmax). The displacements at the gas inlet side edge of the outer jacket are zero because the nodes at the inlet side edge were constrained in the z direction as a boundary condition. Due to the temperature distribution in the r direction and smaller thermal expansion coefficient of the outer jacket (AISI430) than that of the honeycomb core (Fe-20mass%Cr-5mass%Al), expansion of the honeycomb core is larger than that of the outer jacket and the displacement values of the gas inlet side are positive and those of the outlet side are negative.

Support C

Support B 0.6 0.4

0.2

0.2

-0.8 -1.0

28

24

20

16

12

8

4

Jacket

-1.2

# of flat sheet (from outer side)

Outlet side

-1.2 28

Outlet side

24

-1.0

-0.6

20

-0.8

-0.4

16

-0.6

-0.2

12

-0.4

0.0

8

-0.2

4

0.0

Inlet side

Jacket

Inlet side

0.4 Displacement d / mm

Displacement d / mm

0.6

# of flat sheet (from outer side)

Figure 5: Displacement of each flat sheet in the z direction

Sheet displacement differences among the flat sheets are larger at the gas inlet side in the substrate (B) in which the flat and corrugated sheets are brazed in both the gas inlet and outlet region, while displacement differences are larger at the gas outlet side in the substrate (C). Furthur, the displacements in the gas outlet side up to the 4th sheet, where the flat and corru-

149 gated sheets are joined as the Strengthened Outer Layer, are considerably different from those inside the 4th sheet in the substrate (C). The equivalent plastic strain at the gas inlet portion in the 4th flat sheet in the substrate (B) reached nearly 5% and plastic strain amplitude was large, which indicated that thermal fatigue occurred there. On the other hand, plastic strain in the substrate (C) was no more than about 0.3% and the amplitude was small. The substrate (C) was designed in order to prevent the “pushing out“ observed in the substrate (B) by permitting displacement of the sheets in the z direction in the gas outlet side (2), and the simulation results in this study are consistent with those phenomena. 3.2

Effect of Strengthened Outer Layer

The substrate (C) had neither macroscopic displacement nor pushing out (2), but cracks were generated in the sheets after a heat cycle test performed by the authors. The cracks state is schematically drawn in Figure 6. Cracks penetrated the outermost flat sheet at the edge of the core / jacket joint and propagated in the z direction within the adjacent corrugated sheet and stopped on their way. Outer jacket Outermost flat sheet

Outermost corrugated sheet

Crack

Core / jacket joint

Figure 6: Schematic diagram of cracks in the substrates

Figure 7(a) shows equivalent plastic strains (ep) at the points slightly apart from the core / jacket joint in the outermost flat sheet for the substrate (A), (B), and (C), respectively. Figure 7(b) shows Ap at the edge of the core / jacket joint in the adjacent corrugated sheets. Those points correspond to areas cracks were generated. The Ap values became lower than those in Figure 7(b), as the location was apart from the cracked area within the same corrugated sheet. Although the substrate (A) had the shortest life among those three substrates, the values of Ap obtained from the simulation were the lowest. In the substrate (A), the honeycomb core and the outer jacket were joined via only the outermost flat sheet. Therefore, if cracks penetrate the sheet thickness, e.g. when external forces are applied to the substrate, the connection between the honeycomb core and the outer jacket is lost and macroscopic displacement of the honeycomb core occurs in spite of its low thermal stress and strain.

150

Equivalent Plastic strain

Ap

0.06

(a) Outermost flat sheet

0.05 0.04 0.03

B

0.02

C

0.01

p

0.06

(b) Outermost corrugated sheet

AEqu

ival 0.05 ent 0.04 Plas tic 0.03 stra in 0.02

B

C

0.01

A

0

0 0

500

1000 1500 2000

2500

0

500

Time t / sec

1000 1500 2000 2500 Time t / sec

Figure 7: Equivalent plastic strain in the outer most flat and corrugated sheet as a function of time

On the other hand, the substrate (B) and (C) with the Strengthened Outer Layer seem to be disadvantageous for mechanical durability because the values of plastic strain are high and cyclic plastic deformation occurs. Initial cracks are indeed easily to be generated in those substrates. However, the Strengthened Outer Layer is able to hold the honeycomb core even if the crack penetrates the outermost flat sheet, unlike the substrate (A). Figure 8 and 9 show normal stresses in the G and z direction in the outermost flat and corrugated sheets, respectively, in the substrate C at the points same as those shown in Figure 7. In Figure 8, the stresses on the outer surface of the outermost flat sheet are shown, where the values on the inner surface were almost same. Figure 9(a) and 9(b) show the stresses on the outer surface and inner surface respectively. In contrast, the signs of positive and negative are completely opposite between the outer and inner surface especially in IG, which indicates cyclic bending occurred in the corrugated sheet. 400

IGor Iz / MPa

Iz 200

IG 0 -200 -400 0

500 1000 1500 2000 2500 Time t / sec

Figure 8: IG and Iz in the outermost flat sheet as a function of time

Considering the direction of crack propagation, although stress state would be different after crack initiation, Iz is dominant in the outermost flat sheet. It can be estimated that crack generated in the outermost flat sheet propagates in the r or G direction. In the outermost corrugated sheet, IG is dominant due to bending, which leads cracks in the z direction.

151 In the substrate (C), although micro cracks are inevitably generated in the sheets constituting the honeycomb core, the location of initial cracks and their propagation direction are controlled. Therefore, fatal destruction is avoided and high mechanical durability against heat cycles is achieved. (a) Inner surface

400

(b) Outer surface

IG or Iz / MPa

200

IG

0

Iz

Iz

-200

IG -400 0

500

1000 1500 2000 2500 Time t / sec

0

500 1000 1500 2000 2500 Time t / sec

Figure 9: IG and Iz in the outermost corrugated sheet as a function of time. (a) inner surface, (b) outer surface

4

Conclusion

Thermal stresses and strains to be generated in metal substrates for catalytic converters were simulated by elastic-plastic analysis. The flat and corrugated sheets constituting a honeycomb structure were directly modeled by bilinear thick-shell elements. The model could show the pushing out of the honeycomb core where both gas inlet and outlet side are joined and cracks propagation state of the substrate with the Strengthened Outer Layer. Substrates having the Strengthened Outer Layer with asymmetric joint structure in which sheets are joined only in the gas inlet side have two effects for achieving high durability, which are permitting expansion and contraction of the sheets in the z direction at the gas outlet region and controlling the location of cracks initiation and the direction of cracks propagation.

5

Acknowledgement

The authors would like to thank Dr. Yasuo Takahashi, associate professor at Osaka University, for the useful comments on this study.

6

References

[1]

Bode, H., Metal-Supported Automotive Catalytic Converters (Ed.: H. Bode), WerkstoffInformationsgesellschaft mbH, 1997, 17–31 T. Takada et al., SAE 910615, 1991 K. P. Reddy et al., SAE 940782, 1994

[2] [3]

A New Type of Metallic Substrate Reijo Lylykangas, Heikki Tuomola Kemira Metalkat Oy

1

Abstract

To meet ever tightening exhaust emission limits, improved engines and exhaust gas after treatment systems must be developed. The principal trend is to install the catalytic converter as near as practically possible to the engine to hasten light-off performance of the catalyst. A second trend is to reduce the thermal mass of the converter by reducing wall thickness of the substrate, and thirdly, increasing the cell density to improve mass and heat transfer from the bulk gas to the catalyst surface. At the same time as above, demands for improved mechanical durability of the converter are required. These improvements have not only to withstand higher temperatures but also higher accelerations caused by engine vibration and exhaust gas pulsation. A new type of metallic substrate is introduced to meet more stringent requirements. The essential feature of the substrate is that flow channel is not straight and is specifically designed to mix gas flow instead of the traditional laminar flow pattern. This factor improves heat and mass transfer. The new substrate has excellent mechanical durability and superior performance durability compared to that of conventional metallic and ceramic substrates. Theoretical calculations related to mechanical durability together with practical emission and mechanical durability tests are presented.

2

Introduction

The new stringent emission standards have hastened technical development of catalytic converters. Both wash coat formulations and substrates are in the fast development process. Cell densities of ceramic converters are increasing from 400 cpsi to 900 cpsi and wall thickness is reducing from 0,153 mm (6 mils) to 0,052 mm (2 mils). Metallic converters are increasing cell densities up to 2000 cpsi and reducing foil thickness to 0,02 mm. Targets have been to increase geometric surface area and to minimise thermal mass of a converter. To speed up light off time converters are installed as near the engine as possible. In this location converters have to face very demanding conditions. Temperature is higher and acceleration created by engine vibration and exhaust gas pulsation greater than in the under floor position. In a close-coupled position temperature can arise over 1100 °C and acceleration can reach 100g. The weakest links in the ceramic converter systems are the assembling mat and thermal shock. The ceramic blanket does not stand higher temperature than 800 °C [1]. Thermal shock can brake the substrate in case temperature gradient in axial or in radial direction is too high.

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

153 In brazed metallic converters the most critical failure risks are: · The joint between the shell and the substrate is the most critical area. The matrix is heating and cooling faster than the much thicker shell. This causes different thermal expansion rate and high tensile stress to the brazed joints. · The melting point of the brazing material is about 1160 °C. It is near the operating temperature. · Failures of metallic converters start often so that the straight foils start to break. Uneven temperature and flow distribution and high vibration create high tensile stress inside the matrix. A corrugated foil can bend easier than a straight foil can stretch. This can cause the straight foils to break. [2,3]. Light off time of a converter has the most essential influence to the CVS- emission results. Thermal mass of a converter and heat transfer from the exhaust bulk gas to the catalyst surface has vital influence to light off time. After light off period mass transfer dominates reaction rate.[4,5] The flow channels in the catalytic converters are small and straight. In these channels a laminar gas flow develops in the first few millimeters of the substrate. When the flow is laminar mass and heat transfer is not effective. By using discontinuous flow channels can be improved mass and heat transfer in a converter. The main target of this project was to develop mechanically sufficient substrate for the most demanding close-coupled car applications. Simultaneously we were trying to reduce thermal mass of the substrate and improve mass and heat transfer between bulk gas and the catalyst surface.

3

A New Metallic Substrate

Metallic substrates have some essential advantages compared to ceramic substrates in designing the structure of the converter. The channel shape can be discontinuous or channels can be in connection with each other. This gives opportunity to improve mass and heat transfer between bulk gas and the catalyst surface. Also it is possible to mix exhaust gas inside the converter and improve gas flow distribution when gas is passing through the substrate. If a converter locates immediately after a manifold the gas flow is not mixed properly when it comes in the converter. Fast reactions in the converter can cause local overheating. Kemira Metalkat has developed a new kind of metallic substrate called EcoXcell. The structure is based on the classical principle of a static mixer. Corrugated foils are stacked one above the other so that culmination lines are crossing each other. Figure 1 shows the structure. Foils are fixed together in the cross over points by using the new welding process. The matrix made this way is fixed to the shell by using laser welding through the grooves made in the shell and the matrix. The welding are made relative near to each other to avoid failures caused by different expansion of the shell and the matrix during heating and cooling phases. The latest laser welding technology is used to achieve maximal mechanical durability. The shape of the substrate, shell density and all dimensions can be easily varied. In this converter there is no straight foil. Gross angle of the culmination lines can be also changed.

154 A converter with two substrates can be installed so that stacked foil layers have 90° angle between each other. Then converter is mixing gas flow efficiently way and expanding out gas flow in all directions, shown in the figure 2.

Figure 1: New welded structure

Figure 2: New welded two substrate structure

4

Theoretical Calculation

Laboratory test offers durability evaluation for catalytic converter substrate, which could provide pass/failure information under certain vibration environment. Comparing to laboratory test, FEM is another tool, which could be employed for improving durability performance while reducing the product development cost and time. FEM not only evaluates the durability performance, but also explores many possible design options and promises to yield in optimum design with significant cost and time saving. The FEM analysis of a catalytic converter includes calculating on anisotropic material model, thermal stress, and natural frequencies. In

155 this paper an anisotropic non-traditional superelement model is present and all calculations are developed based on MSC.visualNastran for Windows.

5

Load Types

Typical load types on the catalytic converter are: · high temperature · exhaust gas pressure fluctuations · mechanical vibrations from engine and road excitation · high temperature gradient between center of matrix to mantle · thermal shocks due to exhaust gas temperature fluctuations and high gas flow · external temperature shocks due to splash water 5.1

Substrate Non Traditionel Super Element (NTSE) Model

New metallic catalytic converter substrate geometry is too complicated for any exact analytical and numerical solutions method. Typically over 2 million thin shell elements and 12 million degrees of freedom would be need for a discretized model. Homogenization is a procedure in which a real geometry is replace by an anisotropic homogenous material. Individual solid element stiffness on finite element method are based on equation 1 [6 - 10]. 1 1 1

[ K ] = ò ò ò [ B ] [ E ][ B ] J d xd hd z T

(1)

-1 - 1 - 1

[K] = Individual element stiffening matrix [B] = Individual element kinematics matrix [E] = Individual element 6x6 material matrix | J | = Jacobian matrix determinant The properties of the NTSE model are defined in such a way that its global behaviour will be identical to the original. In the case of the catalytic converter body the small piece of substrate in relation to whole substrate dimensions, allows an anisotropy continuum to be defined to replace the original thin shell structure. In homogenisation a NTSE is used to calculate the constitutive equations for the material. This is accomplished by loading the NTSE by a unit force to only one direction at time, and calculating the strain in all other directions. Each strain component will yield a column in the strain matrix. Finally, an inverted matrix on strain matrix is taken. For the converter core the NTSE is a volume containing small piece of substrate. Figure 3. Unit force method based on equation 2–4.

s=

F Þ s = E ×e A

Þ

F = E ×e A

t=

Q Þ t = G×g Þ A

Q =G×g A

and and

E= G=

F A×e

Q A× g

(2) (3)

156 V1 1 C1

Y

Z

X

Figure 3: Catalytic converter structure NTSE model

é ¶u ê ¶x ê ê ¶v ê ¶y ê ê ¶w ê ¶z e=ê ê 0 ê ê ê 0 ê ê ê êêë 0 5.2

¶u ¶x ¶v ¶y ¶w ¶z

¶u ¶x ¶v ¶y ¶w ¶z

0

0

0

0

0

0

0

0

¶u ¶v + ¶y ¶x

0

0

0

0

¶v ¶w + ¶z ¶y

0

0

0

0

ù ú ú 0 ú ú ú ú 0 ú ú 0 ú ú ú 0 úú ú ¶u ¶w ú + ¶z ¶x úúû 0

(4)

Develop the Material Matrix

A three-dimensional anisotropic material is calculated from small piece of substrate. Foil thickness is 0.05 mm. To constrain the model one of the center grid point is fixed in all six component directions. Six self-equilibrating load cases are used to represent each of the six stress components. Each direction was applied as a separate subcase. Self-equilibrating normal loads consist of applying equation and opposite forces to opposite faces of the cube (i.e., Force, Fx, Fy, and Fz). For the shear load, self-equilibrating loads consist of applying four forces around the cube in order to plane the cube in a state of pure shear. Each load case represents a stress, and the resulting strains are the strains due to stress. Figure 4.

157 V1 1 C1

Y

Z

X

Figure 4: Real model transferred to anisotropy solid model

This anisotropic matrix is possible reduce to ortotropic continuum. Equation 5 represents NTSE model strain matrix. 0 0 0 ù -0.00019 -0.00048 é 0.0030 ê -0.00019 0.00334 -0.000398 0 0 0 úú ê 0 0 0 ú 1 ê -0.00048 -0.000398 0.001108 =ê ú 0 0 0 0.0227 0 0 ú E ê ê 0 0 0 0 0.0065 0 ú ê ú 0 0 0 0 0 0.0029 úû êë

(5)

Equation 5 presents a strain matrix that, when inverted, is the material matrix [E]. Note that since the strain matrix is symmetric, the material matrix [E] is also symmetric. The result is the material matrix [E] for our anisotropic material shown in equation 6. Now anisotropy stiffening matrix includes only ortotropic terms 0 0 ù é360 42 174 0 ê 42 317 133 0 0 0 úú ê ê174 133 1027 0 0 0 ú E=ê ú 0 0 44 0 0 ú ê 0 ê 0 0 0 0 153 0 ú ê ú 0 0 0 0 340 úû êë 0

(6)

Foil material modulus of elasticity temperature effect is possible to solve and connect with the FE– model. Equations 7 represent substrate-stiffening matrix at 900 °C.

158

E9000 C

0 0 0 ù é 135 15.8 65.3 ê15.8 119 49.9 0 0 0 úú ê ê65.3 49.9 385.1 0 0 0 ú =ê ú 0 0 16.5 0 0 ú ê 0 ê 0 0 0 0 57.4 0 ú ê ú 0 0 0 0 128úû êë 0

(7)

Of course, it is possible to include the elastic-plastic material model in this model. This full material model is not yet ready. 5.3

Verify the NTSE Model

The model must be verified before taking it into commercial use. The way is to find correlation between standard material test, and the NTSE model. Test results are shown in table 1. Table 1: Test result correlation

Ex Ey Ez

NTSE Model 360 MPa 317 MPa 1027 MPa

Mechanical Test 354 MPa 300 MPa 1015 MPa

Results show a very good correlation between NTSE and material tests. This project is not yet finished.

6

Thermal Stress Analysis

In general, close-coupled converters for today’s applications have to be designed to withstand service temperature of approx. 1050 °C and short time peak temperature of 1100 °C. Measurements have proven that the catalyst core temperature is highest at center parts at the front and decreases towards the end. The casing shell is always colder than the core because it is not a partner in exothermic catalysis reaction, and because outside air is cooling. Typically the center part of a thermally insulated catalyst is about 200 °C warmer than casing shell. In stress calculation the situation is more complicated. All material properties are non-linearly dependent on temperature. Thermal load induced by a constant 200 °C-temperature difference between the substrate core and its casing shell are shown in figure 5.

159

Y Z

X

Figure 5: Thermal loads. Gradient between substrate core to shell is 200 oC

Thermal load is induced by constant 200 °C temperature difference between the catalyst body and its casing shell. Maximum Von Mises stress of 3.3 MPa at substrate and casing shell interface area is obtained. Casing shell maximum Von Mises stress is 13 MPa. Substrate and casing shell Von Mises stress is shown in figure 6. 12995740. 12183943. 11372146. 10560349. 9748552. 8936755. 8124958. 7313161. 6501364. 5689567. 4877770. 4065973. 3254176. 2442379. Y Z

1630583. X

818786. 6989.

Figure 6: Homogenised thermal stress distribution for 200 °C temperature difference. B 90 × 74.5 circular metal substrate. One-quarter substrates are present.

160

7

Frequency Analysis

Natural frequencies for different configurations of the catalytic converter were calculated in the 3D- models. In the 3D- case in order to model a whole catalytic converter body. Substrate must be modelled with solid anisotropic NTSE elements. Casing shell was modelled with thin shell elements. Inlet and outlet flange must be modelled with solid isoparametric elements. Catalytic converter is connected on exhaust manifold rigidly. This rigid connection is modelled with fixed boundary condition on inlet flange bolthole. Natural frequencies are directly proportional to the square root of the modulus of elasticity, which decreases with increasing temperature. The results at 100 °C and 900 °C are given in Table 2, and one mode shape in figure 7. Table 2: Natural frequency on two different temperature. First 3 mode is present.

Temperature o 100 C o 900 C

Mode 1. 1253 Hz 1020 Hz

Mode 2. 1318 Hz 1128 Hz

Y Z

X

Figure 7: Natural mode of a welded metallic catalyst substrate at 1253 Hz. Horizontal displacement

8

Experimental Test Procedure

8.1

Hot Shake Test (KHST-2010)

The hot shake test simulates mechanical vibrations in an exhaust, and respectively converter system, at constant temperature. In most cases the direction of excitation is uniaxial, and the test components are mounted either in horizontal or vertical direction relative to the vibration device or under a defined angle relative to this. Electrodynamic shaker is used to generate the vibration. A 2.0l 16V-gasoline engine used as the heat course. This new cc-catalyst hot shake test includes three different test periods. First period is a horizontal test. Second and third is a

161 vertical test. Table 3, shows test parameters of cc-catalyst hot shakes tests (KHST-2010). Each test period test time is 30 hours and total test time is 90 hours. This periodical test simulates harmonic, and pure random white noise vibration which engine generates in real life. A schematic test set up is shown in figure 8.

Shaker Horizontal Vibration Shaker Vertical Vibration

Y Z

X

Figure 8: Schematic hot shake test set up for CC-applications.

Table 3: Proposal for CC-catalyst hot shake test conditions

Hot vibration

Hot shaker test Period 1 engine 950 °C

Heat source Inlet gas temperature Midbed temperature Exit gas temperature Gas mass flow 300 kg/h Flow distribution Wave fom random Frequency 50 Hz–2000 Hz Acceleration level 14.4 g’s RMS (60 g’s Peak) Direction of simulation vertical Duration time (Hours) 30 hour Total duration time is 90 hrs 8.2

Hot shake test Period 2 and 3 engine 950 °C

300 kg/h sine (peak) 90 Hz 40 g’s horizontal /vertical 30 hour each position

Engine Test Bench by Thermal Cycling (KTC-2010)

The thermal shock test is intended to simulate the cyclic thermal loads, which occur in actual operation. The primary focus here is on the loads resulting from the difference in temperature between the substrate structure and the converter housing, or canning. Converter is connected directly to exhaust manifold. High upper temperature, and temperature transient, and random vibration simulates real life loads. Test cycle is shown in figure 9, and engine speed is shown in figure 10. Table 4 shows the test parameters of cc-catalyst engine test bench (KTC-2010).

162 Table 4: Proposal for CC-catalyst thermal cycling tests.

Thermal cycling Heat source Upper temperature Lower temperature Delta temperature Heating time Cooling time Heat up gradient (dT/Dt) Cool down gradient (dT/dt) Mass flow Cycle time Duration hrs

KTC-2010-01 engine + air injection 1070 °C Midbed 350 °C 720 °C 50 s 50 s +7000 °C/min -3500 °C/min 380 kg/min 100 s 100 h

KTC-2010-02 engine without air injection 1000 °C midbed 600 °C 500 °C 50 s 50 s +3500 °C/min -500 °C/min 380 kg/min 100 s 100 h

TEMPERATURE PROFILE 1100 1000 900

inlet substrate outlet

C°

800 700 600 500 400 300 200 0

100

200

300

400

500 sec.

Figure 9: Gas, and substrate temperature profile of test cycle KTC-2010-01 50 sec. WOT, 5500-4800 rpm, air injection 6000 5000

rpm

4000 3000 engine speed

2000 1000

7 sec. Idle-5500 rpm

0 0

100

50 sec. Idle, air injection

200

300

400

500

sec.

Figure 10: Engine speed profile of test cycle KTC-2010-01

163

9

Mechanical Durability Test

9.1

KTC-2010

The mechanical durability of the catalyst were tested in close couple configuration, figure 11, by fixing the catalyst directly to the engine exhaust manifold. The test cycle is the thermal cycle (KTC-2010).

Figure 11: Schematic test set up for KTC-2010-test.

The foil material in all test parts is the same. Each part has nearly the same dimensions. Test includes brazed structure, ceramic structure and new welded metallic structure. All tested part were without thermal insulation. The more accurate information about this test is shown in figure 12. KTC-2010-01 120

100

100

80

60 50

40

20 5 0 Commercial Brazed Structure

Ceramic Structure

Welded Structure

Figure 12: Test results; the endurance time (h) in close-coupled test by two different constructions

164 The test was stopped after first failure found on catalyst structure. Total test time was 100 hours. Typically failure on brazed structure is straight foils broken, see on figure 13.

Start of failure mechanism

Figure 13: Principal failure mechanics of brazed structure

9.2

KHST-2010

Second mechanical durability test was the hot shake test (KHST-2010). This test includes three different load cases. Each period test time was 30 hours. New metallic catalytic converter was again in a very good condition after this test. We did not find any failure on the substrate nor the casing shell. As an example, figure 14. Shows gas inlet front face of a welded metallic converter 80 mm × 74.5 mm, 500 cpsi, tested over 90 hours according to the above proposed KHST-2010 test.

Figure 14: Gas inlet side of a close-coupled metallic converter after 90 h engine test bench running. Durability test parameter is KHST-2010.

165

10

Emissions, Back Pressure and Flow Distribution

Emission tests were carried out using as a test vehicle Fiat 600, may 2000. The engine displacement was 1,1 liter. Test converters were made as similar as possible. Total PGM- loading and wash coat formulation and cell densities were equal. Because of different structure the geometric surface area and volume of the converters differ slightly from each other. Physical parameters are shown in the table 5. Table 5: Physical parameter

Parameter

Standard metallic converter

EcoXcell converter

Volume, dm³

1,05

1,11

Cell density, cpsi

500

500(*

Total geometric surface area, m²

3,9

3,0

Total PGM loading g/cat

1,24

1,24

Pt:Pd:Rh

0:7:1

0:7:1

Aging of the converters was conducted by using in house aging cycle 20 h RAH. This thermally very severe cycle is described in figures 15 and 16.

Figure 15: Schematic test condition

166

Figure 16: Test cycle

MVEG-B emission test results are shown in the figure 17.

Figure 17: Test results

Results show that all emissions with the EcoXcell converter are better than with the standard converter despite of 23 % lower geometric surface area than in the conventional converter. Emission test results of EcoXcell converter indicate improved mass and heat transfer from the bulk gas to the catalyst surface. Back pressure measurements were conducted by engine test bench. Back pressure of the standard converter was slightly lower than in the EcoXcell converter where the corrugation angle was 20°. Results can be seen in the Figure 18.

167 500

800

450

700

400 600

500

300 250

400

200

300

Temperature [°C]

Pressure [mbar]

350

EcoXcell20 ° Conventional T_average

150 200 100 100

50 0 2500

0 3000

3500

4000

4500

5000

5500

RPM

Figure 18: Back pressure measurements

It was found that back pressure is quite sensitive to the angle in which corrugations are crossing each other. The smaller angle the lower pressure drop in the converter. In the Figure 19 are two similarly stacked EcoXcell converters with 10° and 20° angles is compared.

Figure 19: Back pressure measurements

168 By doubling the angle back pressure was more than doubled. So it is an important parameter when a converter design is optimised. Gas flow distribution was measured with air in room temperature. With the same converters also gas flow distribution was measured. See the Figure 20. 20

Velocity [m/s]

15

EcoXcell Standard met. catalyst

10

5

0 0

10

20

30

40

50

60

70

80

90

100

110

120

Distance [mm]

Figure 20: Flow distribution

Flow distribution with EcoXcell converter was more even than in the standard converter. Vmax/Vaverage was with EcoXcell converter 1.32 and with the standard converter the ratio was 1.70.

11

Summary and Conclusions

The objective of this study was to develop a more durable metallic substrate than any other of the existing substrates. There are today many close-coupled applications where all commercially available substrates have difficulties with mechanical durability. Driving by WOT under lambda control can increase temperature in the converter until 1100 °C. Starting point for this work was to analyse the weakest points of the existing products and to find improvements for those. Our main focus was laying on the metallic substrates but also a ceramic close-coupled converter was tested as a reference. Three weakest points of metallic brazed substrates were found: 1. Fixing between the shell and the matrix. During fast acceleration temperature in the matrix rises faster than temperature of the shell. During deceleration the same happens opposite way. Different thermal expansion causes high tensile stress between the shell and the matrix. 2. Thermal shock and high vibration in radial direction causes higher tensile stress to the straight foil than to the corrugated foil so that the straight foil breaks normally first.

169 3. Melting point of the brazing material is about 1160 °C. It is near the operation temperature. In the new structure improvements for the above mentioned weak points are as follows: 1. Fixing between the matrix and the shell was made by welding. Ring weldings are locating near to each other to minimise tensile stress caused by different expansion of the shell and matrix. 2. By eliminating the straight foil completely the structure was made flexible in radial directions. In axial direction the new structure is firm. 3. All the joints are made by welding. Melting point of the welding spots is about 1500°C. Which is giving more margin to the operation temperature. Mechanical durability test results were showing clearly superior behaviour of the new EcoXcell substrate compared to any of the commercial products. Duration time in the most demanding engine bench tests were more than ten times longer than with than the most common metallic converters and more than double compared to the tested close coupled ceramic converters. According the conducted tests the main objective of the project could be achieved. More tests will be needed. All the other characteristics came as consequence of the new structure. Emission test results and gas flow distribution were improved due to the mixer structure, which enhanced the mass as well as heat transfer in the matrix. There is plenty of room for further optimisation and improvements. Next step will be to study influence of different grossing angles to the emissions, back pressure and mass and heat transfer. Also tests with thinner foil materials will be included to the test programme in the near future.

12

References

[1] [2] [3] [4]

www.unifrax.com, unifrax Product Information Sheet Määttänen M. and Lylykangas R., Mechanical Strength of a Metallic Catalytic Converter Made of Pre-coated Foil Määttänen M. and Avikainen T., Metallic Catalytic Converter Cross Axis Strength Considerations Luoma M., Lappi P. and Lylykangas R., Evaluation of High Cell Density Z-Flow Catalyst Luoma M., Härkönen M., Lylykangas R. and Sohlo J., Optimisation of the Metallic Three-Way Catalyst Behaviour T.R. Chandrapatala. A. D. Belegundu. Introduction To Finite Elements In Engineering. Present Hall, Engelwood Cliffs, New Jersey 07632 Dr. W. Elspass. Design of high precision sandwich structure using analytical and finite element models. MSC World User’s Conference Los Angeles, California March 26 – 30, 1990. Paul R. Woodmansee Master Student, Howard D. Gans, Ph.D Assistant Professor of Aerospace Engineering Air Force Institute of Technology. Finite Element Analysis of Porosity on Material Properties Using MSC/Nastran.

[5] [6] [7] [8]

[9]

170 [10] Craig S. Collier, P.E and Kevin A. Spoth Lockheed Engineering and Sciences Co. Thermomechanical Finite Element Analysis of Stiffened Unsymmetric Composite Panels With Two Dimenssional Models. NASA Langley research Center Hampton, VA. [11] MSC/Nastran The Nastran Theoretical Manual Level 15.5. The Mac Neal-Schwendler Corporation

III Ceramics

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

Development Status of Ceramic Supported Catalyst Claus-Dieter Vogt, Etsuji Ohara NGK Europe GmbH

Mikio Makino NGK Insulators Ltd

1

Abstract

The ceramic honeycomb substrate has been an integral part of automotive emission technology for the past 20 years. From the first ceramic monolithic substrates for catalytic converters to the new more advanced designs and shapes used in various applications and industries, the ceramic substrate has provided years of durability and performance. This paper will discuss the evolution of the ceramic honeycomb substrate from its origin through the latest advanced applications.

2

Introduction

As concerns about increasing atmospheric pollution continue to rise around the world, special environmental interest and political groups continue to push for more demanding pollution regulations for industry to reduce further damage. The internal combustion engine used in automotive transportation has been cited as one of the major contributors to atmospheric pollution. By the year 2010 over 1 billion vehicles are projected to be on the road worldwide (1) Technology continues to strive toward the maximum reduction of the main toxic and environmentally damaging components of exhaust emissions. An example of this technology is the three way catalytic converter which can convert more than 90 % of automobile exhaust pollutants into harmless gases. Since the beginning of emission regulations, the focus has been primarily directed at reducing hydrocarbons, carbon monoxide, and nitrogen oxides. Hydrocarbons are organic compounds of hydrogen and carbon which when exposed to sunlight and nitrogen oxides, react to form oxidants which can irritate mucous membranes, and in some forms are considered carcinogenic. Carbon Monoxide is a colorless, tasteless, and odorless gas produced by the incomplete burning of carbon, which if inhaled with a volumetric concentration of 0.3 % would result in death in 30 minutes. More than 90 % of the carbon monoxide emitted in cities comes from motor vehicles. Nitrogen monoxide is a colorless, tasteless gas produced by the internal combustion engine, but when exposed to air becomes NO2. Pure NOx is a poisonous reddish-brown gas that can irritate mucous membranes. NO and NO2 are sometimes referred together as nitrogen oxides (NOx) and are responsible for photochemical smog and acid rain (2). These represent the majority of toxic and environmentally detrimental gases in automotive engine exhaust.

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

174

3

History of Automotive Emission Control

The automotive industry has been an ongoing matter of interest concerning air pollution from the internal combustion engine since the early 1950’s, when the photochemical reactions that cause smog first became an issue in the Los Angeles area. Pollution attributed to the automobile slowly became evident through the 50’s and became a nationwide issue in the late 1960’s with the establishment of emission regulations by the state of California. This led to emission standard amendments to the Clean Air Act of 1965 passed by the Congress of the United States. In 1970, Congress passed the so-called Muskie amendments which required a 90 % reduction of automotive exhaust emissions between 1970 and the 1975–1976 model years. Initially, exhaust emissions were reduced by engine calibrations such as leaner air/fuel ratio and spark retardation, but many auto manufacturers believed that the continued compromising of engine calibrations would lead to unacceptable fuel consumption and driveability. The solution was an aftertreatment application called the catalytic converter, which reduced the harmful effects of automotive exhaust emissions without sacrificing engine performance. The first catalytic converters, called oxidation catalysts, only controlled HC and CO, because NOx regulations were not strict and could be met by using exhaust gas recirculation, which led to less nitrogen oxide formation from the engine. These catalyst systems were supported by one-eighth inch diameter, thermally stable, transitional alumina coated pellets. Thus, these early converters were named pellet-type converters. They were used for the first few years until durability, available space under the car, and pressure drop issues became more important. During the early 70’s another type of catalyst support, the ceramic monolith (honeycomb) substrate, was developed for the catalytic converter. Figure 1 shows various honeycomb substrates and a cut-away of a typical catalytic converter. The honeycomb substrate gets its name from the array of parallel channels or cells that create the look of a honeycomb. This new design possessed compatibility with catalysts and coatings, flexibility in shape and size, high geometric surface area, durability, low flow restriction, space efficiency, and fast warm-up.

Figure 1: Ceramic honeycomb substrate and automotive catalytic converter

In 1977 Congress amended the Clean Air Act and set revised standards to achieve a 90 % reduction in hydrocarbons by 1980, and 90 % carbon monoxide and 75 % nitrogen oxide reduction by 1981. To meet the newer regulations, the requirements for the catalyst support

175 became more demanding, requiring higher thermal durability, better thermal shock resistance, better corrosion resistance, lower pressure drop, faster warm-up, and higher mechanical strength. This hastened the move toward the ceramic honeycomb substrate, which provided these characteristics, and by the mid 1980’s most automotive manufacturers had changed from the pellet-type support to the substrate support. The original ceramic honeycomb substrates were primarily used with oxidation catalyst systems which required additional air (O2) for catalytic reaction. This air came from engine calibration adjustments (lean fuel ratio) and/or added pump configurations to inject air into the system. The oxidation systems, which by themselves could only control HC and CO, and relied on exhaust gas recirculation (EGR) to reduce NOx, were used until the three-way catalyst system emerged in the early 1980’s. The majority of the early substrates in the oxidation systems were designed to have 300 cells per square inch with a wall thickness of 12 mil (0.30 mm), known as a 12 mil/300 cpsi configuration. As emission standards tightened, the 6 mil/400 cpsi substrate appeared, which presented a larger geometric surface area, quicker warm-up, and lower pressure drop. The three-way catalyst system controls all three main components using a loop fuel metering control system with an oxygen sensor and a catalytic converter. Most of these systems use the ceramic honeycomb substrate and rely on the loop fuel metering system to control the air/fuel ratio near the stoichiometric point, or the point where there is no excess fuel or air in the mixture. Near this stoichiometric point of air/fuel mixture, the following chemical reactions can optimally reduce the three main components of exhaust gas HC, CC, and NOx

2CO + O 2 ® CO 2 2C2 H 6 + 7O 2 ® CO 2 + 6H 2O 2NO + 2CO ® N 2 + 2CO 2 Therefore a small window exists where the catalyst can reduce all three of these emission components as long as the air/fuel ratio is within this specified area, as shown in Figure 2.

Figure 2: Conversion efficiencies of a typical three-way catalyst

176

4

Substrate Requirements

Three-way catalyst technology is the preferred method of catalyst reaction, but tightening emission regulations and increasing performance needs continue to make the requirements for the catalyst support more stringent. The requirements and capabilities of the converter usually dictate the location in the exhaust flow, as shown in Figure 3. The thermal shock resistance, thermal durability, mechanical integrity, warm-up, and low pressure drop, of the catalyst support continue to be improved, but at more critical levels and for a longer duration.

Figure 3: Converter location in passenger car

The catalytic support must be able to withstand the high and varying degrees of exhaust gas temperatures associated with the automotive engine. Catalytic converter support temperatures associated with normal stop and go driving can range from atmospheric temperature to 900 °C extreme exhaust temperatures, which usually occur due to engine ignition problems such as misfiring, can cause temperatures to exceed 1,000 °C . The ceramic honeycomb provides high temperature and thermal shock resistance due to a high melting point and low coefficient of thermal expansion (3,4). The catalyst substrate requires high mechanical or structural strength due to the severe vibration that the substrate must endure by the tight mounting in the converter. The mechanical integrity of the substrate is directly proportional to the integrity of the cell structure of the honeycomb, and the shape of the honeycomb. The canning procedure inflicts a large amount of compressive pressure on the substrate during the clamping of the shell around the substrate. The heat mass of the substrate is the main factor affecting light-off performance. The greater the heat mass, the longer it takes to transfer the needed heat across the substrate, and to reach full operating temperature. Light-off performance is important in the overall conversion efficiency of the catalyst. The converter in the exhaust system produces pressure drop that has an immediate and counteractive effect on engine performance and fuel economy. This pressure change is related to the frictional loss across the honeycomb matrix, created by the cell wall thickness and cell pitch. Reducing the pressure drop is a function of increasing the open frontal area and creating larger cells, which the thinner wall configurations, such as the 6 mil/400 cpsi substrate, can accomplish.

177

5

Production Process of Ceramic Honeycomb

Since the ceramic honeycomb substrate’s inception in 1975, it has been made using Cordierite (2MgO + 2ALO3 + 5SiO2), which is mainly kaolin, talc, and alumina. This composition provides low thermal expansion, high temperature stability, good porosity, and excellent oxidation resistance. Table 1 shows the typical properties of cordierite honeycomb. Table 1: Typical material properties of cordierite honeycomb substrate

Item

Properties Cordierite 2MgO-2Al2O3-5SiO2

Crystal structure

Thermal properties

Physical properties

Thermal expansion (x10– 6/°C) (40 °C–800°C) Specific heat (cal/g°C) Softening temperature (°C) Melting point (°C) Total pore volume (cm3/g) Porosity (%) Mean pore diameter (µm)

Mechanical properties

Compressive strength (kg/cm2)

Thermal shock resistance

Electric furnace-room atmosphere (°C Difference)

< 1.0 0.2 1410 1455 0.2 35 4 A-axis B-axis C-axis

> 85 > 11 >1

> 650

The production process of ceramic honeycomb consists of four phases, the preparation of raw materials forming, firing, and inspection and testing of the finished product. As for raw material, the mineral impurities and particle size of raw materials greatly affect the cordierite characteristic’s such as water absorption, coefficient of thermal expansion and thermal durability of the substrate. The raw material is mixed with organic binders and water, and kneaded to produce a uniform clay with desired hardness and temperature. These factor is important for forming process. As for forming, the extrusion method performs two important tasks: providing the formation of the honeycomb in many sizes and shapes, and lowering the coefficient of thermal expansion of the ceramic material. These tasks are important in meeting the increasing requirements of catalyst supports for thermal shock resistance, and mechanical strength and durability. Figure 4 illustrates the structure of a ceramic honeycomb extrusion die, The mechanism to form square cells begins as material flows into one of the back holes of the extrusion die (a), through the die (b), and out of the corresponding slit junction (c). As the material continues to flow out of the slit junction , it connects with the material flowing from other slit junctions to form a connecting point of the honeycomb wall (d). There are possibility of square, hexagonal, triangular, or rectangular on the cell structure. The square cell configuration has been widely selected and used in this field due to the balance

178 between mechanical strength and pressure drop. As the special application, the hexagonal cell also applied currently.

Figure 4: Ceramic honeycomb extrusion dies

The extrusion process also provides the necessary anisotropic arrangement of the kaolinite crystals for lower thermal expansion, as shown in Figure 5. These crystals are flat and the shearing force of the extrusion through the narrow channels of the die orients them. During the firing process the cordierite crystals are produced so that the negative thermal expansion c-axis is oriented vertically to the c-axis of the original kaolinite and in parallel with the honeycomb wall. This production of cordierite with favorable orientation (anisotropy) of the crystals is what allows the substrate to improve upon the coefficient of thermal expansion characteristics of the combined raw materials (3,4).

Figure 5: Orientation of cordierite crystal

179 The firing process affects important characteristics of the substrate such as coefficient of thermal expansion, water absorption, shrinkage percentage of the product, and thermal stability. The temperature is increased to approximately 1,400 °C and remains at this temperature to assure complete chemical reaction.

6

Advanced Ceramic Honeycomb

Automobile manufacturers and emission system suppliers continue to provide new technology for low temperature capability and high temperature durability. The targets of higher engine power (reduced pressure drop), higher fuel economy, and durability continue to be expected from the automotive industry as a whole. Various new technologies are presently available for these requirements and will be briefly discussed. Thin Wall Substrate – Bulk density and cell structure have a direct effect on the warm-up capability of the substrate. By reducing the bulk density and increasing geometric surface area, the substrate not only warms up faster, but also reduces pressure drop. Existing technology has produced thin wall honeycombs with these attributes. These substrates have 4 mil wall thickness and either 400 cpsi or 600 opal cell density for improved catalyst performance. The thin wall substrates 4 mil / 400 cpsi and 4 mil / 600 cpsi reduces bulk density 30 % and 15 % respectively compared to the 6 mil / 400 cpsi standard substrate. Geometric surface area increases by 5 % (4 mil / 400 cpsi), and 25 % (4 mil / 600 cpsi). The 4 mil / 400 cpsi reduces pressure drop 15 %, but the 4 mil / 600 cpsi increases pressure drop 26 % as shown in Figure 6.

Figure 6: Pressure drop and geometric surface area

Effect of Substrate Geometric Surface Area – Figure 7 shows two tested catalytic converter layouts. The catalyst have the identical diameter, and length, and catalyst coating consists of Pd, Pt and Rh with total precious metal loading of 150 g/ft3. The catalysts were aged at 850 °C for 50 hours with 60 seconds cruising and 5 seconds fuel cut cycle at engine dynamometer.

180 The catalysts were tested with two type of vehicles, (TLEV and ULEV) with FTP cycle. The catalysts were installed at close coupled position, 400 mm downstream from engine, for TLEV vehicle, and at under-floor position,1100 mm downstream from engine, for ULEV vehicle. The vehicles are applied two typical approach to reduce cold emission, which either secondary air injection or lean start control at cold start. Close- Coupled (C.C.) – only

‘96 TLEV 2.2Liter L- 4 Gasoline E/G

Secondary Air Injection at Cold Start

for TLEV Vehicle

O2 Sensor

400mm

Pd/Pt/Rh Catalyst 1.0 liter 106Dia.x114Lmm

‘98 ULEV 2.3Liter L- 4 Gasoline E/G

Under- Floor (U.F.) – only

O2 Sensor

Original lean A/F Control at Cold Start

for ULEV Vehicle

Pd/Pt/Rh Catalyst 1.0 liter 106Dia.x114Lmm

1100mm

Figure 7: Tested catalytic converter layouts

Total NMHC and total NOx emissions are shown in Figs. 8–11 in relation to geometric surface area. Figures 8–11 demonstrate the effectiveness of thin wall and high cell density substrate design. These figures generally indicate GSA (Geometric Surface Area) dependency of total NMHC and total NOx emissions. The increased GSA effectively reduced total NMHC and total NOx emissions.

TOTAL NMHC-Emissions ,g/mile

0.07 Test Vehicle : '96TLEV 2.2 liter L- 4 Gasoline Cold-Start Control: Secondary Air Injection Test Cycle : FT P- 75

5/300 (Prototype)

0.06

6/400

0.05

4/600 4/400

Converter: 1.0 liter C.C.-only Substrate : 106 Dia. x 114 L mm

0.04

0.03 20

Catalyst

: Pd/Pt/Rh, 150g/ft 3

Aging

: max. 850oC x 50hrs

25

O 2 Sensor 400mm

30

35

40 2

GSA of C.C.Substrate, cm /cm Figure 8: Total NMHC emissions from TLEV vehicle

3

181 These figures also demonstrate the effectiveness of close-coupled installation of catalyst. TLEV vehicle with close-coupled installation achieved similar emissions to those of ULEV vehicle with under-floor installation.

TOTAL NOx-Emissions ,g/mile

0.35 Converter: 1.0 liter C.C.-only Substrate : 106 Dia. x 114 L mm

0.3

Catalyst

O2 Sensor

: Pd/Pt/Rh, 150g/ft3 o

Aging

: max. 850 C x 50hrs

0.25

400mm

5/300(Prototype)

0.2

6/400

4/400

4/600

Test Vehicle : '96TLEV 2.2 liter L- 4 Gasoline Cold-Start Control: Secondary Air Injection Test Cycle : FTP- 75

0.15 0.1 20

25

30

35 2

GSA of C.C.Substrate, cm /cm

40 3

TOTAL NMHC-Emissions ,g/mile

Figure 9: Total NOx emissions from TLEV vehicle

0.1 0.09

Test Vehicle : '98ULEV 2.3 liter L- 4 Gasoline Cold-Start Control: Original lean Start Test Cycle : FT P- 75

0.08 5/300 (Prototype)

0.07

4/400

0.06

6/400

0.05 0.04 0.03 0.02

4/600 Converter: 1.0 liter U.F.-only Substrate : 106 Dia. x 114 L mm Catalyst : Pd/Pt/Rh, 150g/ft3 Aging

20

o

: max. 850 C x 50hrs

25

O 2 Sensor

1100mm

30

35 2

GSA of U.F.Substrate, cm /cm

40 3

Figure 10: Total NMHC emissions from ULEV vehicle

Effect of Substrate Bulk Density – The catalysts were tested with additional secondary air injection during cold start. Catalytic performance was evaluated under FTP cycle and the first 140 sec of Bag-1 in FTP cycle in order to compare the light-off performance in each system emission level The catalytic coating consisted of Pd only with total precious metal loading of 120 g/ft3 for close-coupled converter, and of 40 g/ft3 for under-floor converter. The catalysts were aged at

182 750 °C for 100 hours with 60 seconds cruising and 5 seconds fuel cut cycle. The emission performance test was conducted with configuration of a close-coupled converter and an underfloor converter. The close-coupled catalyst converter was installed at 400 mm downstream from engine, and the under-floor catalyst converter was installed at 600 mm downstream from close-coupled converter.

TOTAL NOx-Emissions ,g/mile

0.4 Test Vehicle : '98ULEV 2.3 liter L- 4 Gasoline Cold-Start Control: Original lean Start Test Cycle : FTP- 75

0.35

5/300 (Prototype)

0.3 0.25

6/400

4/600

4/400

0.2

Converter: 1.0 liter U.F.-only Substrate : 106 Dia. x 114 L mm

0.15

O2 Sensor

3

Catalyst

: Pd/Pt/Rh, 150g/ft

Aging

: max. 850 C x 50hrs

o

1100mm

0.1 20

25

30

35

40

GSA of U.F.Substrate, cm2/cm3 Figure 11: Total NOx emissions from ULEV vehicle

Total HC is shown in Figs. 12 in relation to Bulk Density. While the reduced BD, namely low heat mass, is significantly reducing bag-1A emissions, namely cold-start HC emissions. Figures 13 also shows the relation between the BD of the close-coupled substrate and total HC emissions. Since HC cold emission dominates the total HC emission, the lower BD effectively reduces the bag-1A and the total HC emissions. 0.08

Bag-1A & Total HC-Emissions (g/mile)

Total 6/400

0.06

4/400 4/600 0.04

Bag-1A Catalyst System: Pd-only Cold-Start Control: Secondary Air

0.02

C.C.Substrate: Volume 0.69dm3 U.F.Substrate: Volume 1.70dm3 6mil/400cpsi 0 0.2

0.3

0.4

3 Bulk Density of C.C.Substrate (g/cm )

Figure 12: BD vs Bag1A & total HC emissions

0.5

183 Ultra Thin Wall (UTW) substrate – To meet the most stringent regulations such as SULEV, UTW substrate (3 mil / 600 cpsi, 2 mil / 900 cpsi, and 2 mil / 1200 cpsi) also are used in the United States. (27,28,29,30) As shown in Fig.13, 2/900 has a potential of 30 % precious metal saving compare to 3/600 by introduction of higher GSA.

Figure 13: Demission performance of ultra thin wall substrate

7

Conclusion

The honeycomb structure can provide maximum contact surface between a gas and a solid with minimum pressure drop. Cordierite honeycomb established its unique position because it can withstand severe operating conditions including high temperature, and can be produced at a low cost. Cordierite honeycomb has been contributing to air pollution control worldwide as a ceramic substrate and filter element for automotive and industrial emission sources. Thin wall and high cell density substrates, 4 mil / 400 cpsi, 4 mil / 600 cpsi, shows significant catalytic performance improvement. In this study, it is demonstrated the bulk density and geometric surface area is the most significant factors of the improve the emissions of HC and NOx.

184

8

References

[1]

J. Mackenzie, M. Walsh; “ Driving Forces: Motor Vehicle Trends and their Implications for Global Warming,” Energy Strategies and Transportation Planning, World Resources Institute, December 1990. Bosch GmbH: “Automotive Handbook,” Vol. 3,1993. S. Mochida; “Ceramic Honeycombs in the Spotlight,” Chemical Engineering MOL, 1984. I. Lachman, R. Bagley, and R. Lewis; “Thermal Expansion of Extruded Cordierite Ceramics,” Ceramic Bulletin, Vol. 60, No.2,1981. J. Cook, C. Fucinari, and C. Rahnke; “ Application of Performance and Reliability Concepts to the Design of Ceramic Regenerators,” SAE No. 770334, Feb. 1977. S. Gulati; “Impact of Washcoat Formulation on Properties and Performance of Cordierite Ceramic Converters,” SAE No. 912370, Oct. 1991. S. Matsuura, A. Harai, K. Arimura, and H. Shinjoh; “ Development of Three-Way Catalyst Using Only Pd as Activator,” SAE No. 950257, Feb. 1995. R.J. Brisley, GR. Chandler, H.R. Jones, P.J. Anderson, and P.J. Shady; “The Use of Palladium in Advanced Catalysts,” SAE No. 950259, Feb. 1995. A. Punke, U. Dahle, S.J. Tauster, H.N. Rabinowitz, and T. Yamada; “ Trimetallic ThreeWay Converters.” SAE No. 950255, Feb. 1995. E AE/STF-3; “Motor Vehicle Emission Regulations and Fuel Specifications –1991 update,” Report No. 3/91, Brussels, May 1991. M. Machida, T. Yamada, and M. Makino; “Study of Ceramic Catalyst Optimization for Emission Purification Efficiency,” SAE No. 940784, Feb 1994. H. Mizuno, F. Abe, S. Hashimoto, and T. Kondo; “A Structurally Durable EHC for the Exhaust Manifold,” SAE No. 940468, Feb. 1994. L. Socha, D. Thompson, and P. Weber; “Optimization of Extruded Electrically Heated Catalysts,” SAE No. 940468, Feb. 1994. K. Kollman, J. Abthoff, and W. Zahn; “Concepts for Ultra Low Emission Vehicles”, SAE No. 940469, Feb. 1994. P. Burk, J. Hochmuth, D. Anderson, S. Sung, S. Tauster, C. Tolentino, J. Rogalo, G. Miles, M. Niejako, A. Punke, and U. Dahle; “Cold Start Hydrocarbon Emissions Control”, “SAE No. 950410, Feb. 1995. M. Hattori, G. Uesugi, and M. Machide; “Optimization of Catalytic Converter Location Achieved with a Curved Catalytic Honeycomb Substrate,” SAE No. 940743, Feb. 1994. H. Hwang and J. Mooney: “Catalytic Converter Potential for 2-Stroke Motorcycle Engine Emission Control,” International Seminar on Catalytic Converters: Fresh Steps, Bangladore, India, Feb.1995. S. Gulati, J. Ten Eyck, and A. Lebold; “Durable Packaging Design for Cordierite Ceramic Catalysts for Motorcycle Application,” SAE No. 930161; March 1993. J. Johnson, S. Bagley, L. Gratz, and D. Leddy;” A Review of Diesel Particulate Control Technology and Emissions Effects,” SAE No. 940233, Feb. 1994.

[2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15]

[16] [17] [18] [19] [20]

185 [21] Y. lchikawa, S. Yamada, and T. Yamada; “Development of Wall-Flow Type Diesel Particulate Filter System with Efficient Reverse Pulse Air Regeneration.” SAE No. 950735, Feb. 1995. [22] M. Walsh; “Global Trends in Diesel Particulate Control – A 1995 Update,” SAE No. 950149, Feb. 1995. [23] B. Engler, E. Lox, K. Ostgathe, W, Cartellieri, and P. Zelenka; “Diesel Oxidation CataIysts with Low Sulfate Formation for HD-Diesel Engine Application.” SAE No. 932499, Sep. 1993 [24] P. Zelenka, G. Hohenberg, and U. Graf; “Diesel Oxidation Catalyst Application Strategies with Special Emphasis on Odour Reduction, ” SAE No. 942066, Sep. 1994. [25] R. Heck, J. Chen, and B. Speronello; “Commercial Operating Experience with High Temperature SCR NOx Catalyst,” Air and Waste Management Association, Annual Meeting and Exhibition, June 1993. [26] K. Umehara, T. Hijikata, and F. Katsube, “Catalytic Performance Improvement by High Cell Density / Thin Wall Ceramic Substrate,” ISATA Paper 96EN044, 1996 [27] K. Umehara, T. Yamada, T. Hijikata, Y. Ichikawa, and F. Katsube, “Advanced Ceramic Substrate: Catalytic Performance Improvement by High Geometric Surface Area and Low Heat Capacity,” SAE Paper 971029, 1997 [28] S. Kikuchi, S. Hatcho, T. Okayama, and S. Inose, K. Ikeshima; “High Cell Density and Thin Wall Substrate for Higher Conversion Ratio Catalyst, ” SAE No.1999-010268,1999. [29] J. Schmidt, A. Waltner, G. Loose, A. Hirschmann and A. Wirth, W. Mueller, J. A. A. van den Tillaart, L. Mussmann, D. Lindner and J. Gieshoff, K. Umehara and M. Makino, K. P. Biehn, A. Kunz; “The Impact of High Cell Density Ceramic Substrates and Washcoat Properties on the Catalytic Activity of Three Way Catalysts,” SAE No. 1999-010272, 1999. [30] H. Kitagawa, T. Mide, K. Okamatsu, and Y. Yasui; “L4-Engine Development for a Super Ultra Low Emissions Vehicle,” SAE No. 2000-01-0887, 2000. [31] K. Nishizawa, S. Momoshima, M. Koga and H. Tsuchida; “Development of New Technologies Targeting Zero Emissions for Gasoline Engines,” SAE No. 2000-01-0890, 2000

Evaluation of In-Service Properties and Life Time of Automotive Catalyst Support Materials Dr. Uwe Tröger, Matthias John Lang Zeuna Stärker GmbH & Co. KG, Augsburg

1

Introduction

The process of automobile catalytic converter manufacturing comprises production of a monolithic substrate for the ceramic coating, the coating itself, and a housing suitable for insertion of the catalytic converter into the exhaust system of the automobile's engine. One important step in this process is the so-called canning of the coated substrate within it's housing. As in most cases the substrate is made from ceramic materials, it cannot sustain high stresses neither during the manufacturing process nor during subsequent service conditions. The housing is made form heat resistant steel sheet. One way to convey stresses between the substrate and the housing is to insert a supporting material.

2

General Properties of Support Materials

As can be seen from the extreme conditions of operation of a catalytic converter - temperatures up to 950 °C, and a thermal expansion that is an order of magnitude larger for the housing than for the substrate – the supporting material plays a key role for proper and long-term service. In detail, the supporting material has to fulfill the following requirements: · fixation of substrate within housing · compensation of CTE mismatch between substrate and housing · compensation of geometrical tolerances of the substrate · resistance against corrosion due to acidic gas components · heat resistance · resistance against erosion due to high gas velocities and gas pulsation · low gas leakage Additionally, low thermal conductivity and capacity are advantageous properties. When the first cars where equipped with catalytic converters, wire mesh was chosen as support material. As wire mesh does not conform with a number of the before-mentioned criteria, e.g. it has a high gas leakage, support materials made from mineral components came into use. An example of a mineral-base support material is the intumescent mat. In this material, artificial mineral fibers and natural mineral Vermiculite grains are compounded by means of organic binders and pressed into sheet form. The binder assures the integrity of the support mat during manufacturing and is burned out during service, while the fibers serve to embed the Vermiculite grains during operation. The Vermiculite, owing to it's crystallographic structure, expands upon heating, increasing thereby it's volume by up to 1000 %. From this volume in-

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

187 crease stems the excellent ability of intumescent support mat materials to provide a required minimum of stress conveyance in manufacturing and a maximum of fixation force in service. However, as catalytic converters are positioned closer to the engine in order to achieve quicker heat-up to operating temperature, intumescent support materials approach their limits regarding thermal stability and erosion resistance. An alternative are non-intumescent support mats. In these the Vermiculite is replaced by a higher density and modified morphology of the fibers. In the following properties of intumescent and non-intumescent mat materials are compared with respect to their stress relaxation and mechanical aging behavior, and their erosion and corrosion resistance.

3

In-Service and Life-Time Related Properties of Support Materials

Mineral-base support materials consist of the following microstructural features: 1. fibers 2. Vermiculite grains (intumescent mat only) 3. binder islands 4. bonded fiber junctions 5. open fiber junctions 6. fiber bends 7. droplets 8. shot During the canning process, the substrate is enveloped in a sheet of supporting mat and inserted or wrapped into the housing, the supporting mat being thereby compressed. If the resulting pressure between mat and housing is plotted against time (see figure 1), it can be seen that initially an exponential increase takes place until a peak pressure is reached. Afterwards the pressure decays towards a relaxed pressure. The peak pressure and the relaxed pressure are closely related to the closing force used during canning, the gap between substrate and housing, and the support mat material's properties, e.g. density, microstructure, and composition. In the initial part of the pressure-time curve the support material's fibers are deflected according to their moments of inertia and the free loop lengths between junctions. As closing force is further increased and the material is being compressed, new open fiber junctions form and the loop lengths are shortened. The material becomes therefore stiffer and pressure increases exponentially. Upon maximum closing force the peak pressure is reached, and relaxation sets in. Relaxation is due to viscous fiber deformation and fiber rearrangement against frictional forces. With increasing manufacturing speed higher peak pressures are obtained. The resulting reduced relaxed pressures are ascribed to fiber breakage and junction rearrangement. The described behavior has been observed with both intumescent and non-intumescent support mats, although intumescent materials show a less pronounced relaxation behavior as the Vermiculite grains do not participate in viscous deformation or geometrical rearrangement (see fig. 1).

188

600mm/min

pressure

10mm/min

0

200

1mm/min

400

0,25mm/min

600

800

1000

1200

time [s] Figure 1: Pressure between supporting material and housing during the canning process for different closing rates (intumescent mat)

Mechanical aging of the support material is a result of cyclic expansion and compression during engine start-up caused by the differences in thermal expansion between substrate and housing. The measurement of pressure versus time under cycled external compression reveals that in the decompression cycle a relaxation behavior equivalent to that in the compression cycle occurs (see fig. 2).

0

500

1000

1500

2000

2500

time [s] Figure 2: Pressure between supporting material and housing under cyclic loading (intumescent mat)

189 When comparing relaxation under static conditions with that under cycled conditions it is found that in the latter case lower relaxed pressures are attained. This is assumed to arise from a rearrangement of fibers and open fiber junctions into geometrically favorable sites under the action of cycled pressure, as is the case with hysteretic dislocation movement. Also, fiber breakage may contribute significantly. The reduced relaxed pressures under cycled conditions requires compensation by correspondingly modified manufacturing or material parameters, i.e. increased closing force, reduced gap, or improved support mat properties. Under the pulsating pressure of the exhaust gases erosion of exposed edges of the support material may lead to extensive material removal and the formation of gas bypasses or substrate dislocation. In a test simulating gas pressure pulsation the difference in erosion resistance between intumescent and non-intumescent materials is evident (see fig. 3). The erosion, measured as mass loss per time, is about one order of magnitude higher for intumescent material than for non-intumescent material. Main contributors to erosion are the Vermiculite grains which are not effectively bonded to the surrounding fibers, but move relatively freely and can in some cases be observed to act as miniature grinding stones on neighboring fibers. In order to analyze short and long-term corrosion behavior of different support materials acid resistance and base resistance tests were defined (see table 1). Under short term acid attack intumescent and non-intumescent materials similarly show almost no weight loss. However, in the long term test significantly lower acid resistance was observed with intumescent material. It remains to be evaluated which components of the support materials examined are especially susceptible to corrosive attack. 100

erosion (g/h)

intumescent mat 10

intumescent mat with rigidizer 1

non intumescent mat 0,1 0,25

0,5

0,75

mount density (g/m3) Figure 3: Erosion of differerent support materials versus mount density. The mount density is the density of the support mat after canning.

190 Table 1: Mass loss in percent during immersion in acids and bases

intumescent material non-intumescent material

4

acid resistance (18,5 % HCl) 10 min 10 days 99,81 70,92 99,95 99,72

base resistance (20 % NaOH) 1 day 79,45 94,90

Summary

Support materials play a key role in the manufacturing and operation of catalytic converters. In the present work, an attempt was made to correlate macroscopic mechanical and chemical properties to the microstructure of support materials. It was shown that under service conditions, which means cyclic compression and decompression, microstructural processes lead to a reduced relaxed pressure. Manufacturing parameters have to be adapted correspondingly. Significant differences between intumescent and non-intumescent support materials with respect to erosion and corrosion resistance were observed. A more detailed examination of the microstructure-property relationship, with a focus on the relaxation phenomena and high temperature aging, is in progress.

5

References

[1]

R. J. Locker, C. B. Sawyer, G. Eisenstock, SAE 2001-01-0223

Loads, Design and Durability Evaluation of Mount Systems for Ceramic Monoliths G. Wirth J. Eberspächer GmbH & Co., Esslingen (D)

1

Environment

Stringent emission legislation leads to more sophisticated emission control systems in passenger cars. One solution is the moving of the catalytic systems closer to the engine, since the catalytic coatings withstand higher temperatures, and the use of catalytic coated substrates with higher cell densities and lower wall thicknesses. As a result the heat loss of the exhaust gases before the catalyst is lower and the catalytic contact surface is larger. The system shows faster light off of the catalytic reaction at reasonable backpressures under full load. For the development of the exhaust system this means higher temperatures at the catalytic converter and higher accelerations on the whole system, which is mounted directly at the engine, without decoupling bellows between engine and converter. In underfloor converters most of the vibrations are perpendicular to the gas-flow-direction, but the space restrictions in the engine compartment often leads to a converter mounting where the gas flow in the converter is almost parallel to the piston movement. This results in an addition of the forces from vibration and the pressure loss over the monolith in the mount system. The minimum air cooling on the outer surfaces of the catalytic converter in the restricted space in the engine compartment together with the higher exhaust gas temperature leads to higher surface temperatures at the can. Ultra Thin Wall Monoliths, with high cell densities to create the large surfaces at low backpressure, have in contrast to the raisen loads lower strength and worse strength conformity than the usual substrates, as it is very hard to produce this fine ceramics with porous material with wall thicknesses of down to 0,05 mm and cells with a minimum width of 0,68 mm. Keeping dimensional tolerances at the substrates the change from underfloor to closed coupled catalysts led to the following numbers in the mounting environment: Table 1: Dimensional tolerances Max. exhaust gas temperature Max. acceleration in flow-axis Max. can temperature Min. isostatic strength (monolith) Diameter tolerance (2 × contour)

Underfloor catalyst 850 °C < 15 g < 400 °C > 1.1 Mpa +/– 1.6 mm

Closed-coupled catalyst 1050 °C ~ 30 g up to 600 °C > 0.7 Mpa +/– 1.6 mm

This data was generated with todays systems, which are mainly built as modular converters with a ceramic substrat in a case covered by expanding paper or ceramic fiber mat and a metal can. An air gap insulated cone is welded on both ends of this middle part to make the connec-

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

192 tion to inlet and outlet tubes with lower diameters. Depending on the space and shape of monolith the cross sectional area of the monolith is always much larger than the tube diameter to maintain a low backpressure and a low space velocity for the exhaust gases to get good catalytic effects. For european driving conditions the high loads and temperatures are common for long time use in cars and systems have to be tested for the integrity of the system for more than 100 000 miles (more than 160 000 km).

2

Requirements for the Canning of Ceramic Monoliths

The change in the environment of the converter lead to the search for new mount materials for the ceramic monoliths and for a new design of the mount system, as it was soon found out, that the standard expanding paper mounting could not work under all of these conditions. The design had to fullfil all the following requirements for the canning of the bricks: · Maintain sufficient holding forces under all loads between ceramic monolith and can · Compensation of all tolerances between monolith, can and mat and all influences from the production process to prevent monolith crushing in mass production and under load · High temperature stability and aging resistance in hot, acidic and humid atmosphere for the cars live · Stability against erosion from the exhaust gas pulsation, even under can deflection and humid atmosphere (enlarged gaps) · Insulation properties to reduce can temperature for minimum gap changes through thermal expansion and for low heat transfer into the engine compartment · Gas sealing properties to prevent a bypass without catalytic treatment Besides these technical requirements there is always the demand for a reasonable price of the system and as a general target a minimum space for the mount system around the monoliths and a reliable mass production process. In practice the mounting is done by wrapping an elastic ceramic fiber mat or an expanding paper around the substrat and compressing this material with a metal sheet case on to the monolith. The holding force is given by the friction between the monolith, can and mat, the contact surface and the compression wich is given by the mat (Figure 1). It has to be higher than the actual load from the pressure drop over the monolith and the inertia force of the monolith under vibration of the system. In service the whole system heats up from inside to outside. As the substrates are specially made with a very low coefficient of thermal expansion (good thermal shock durability), the system mainly changes dimension at the outer shell, even if this is at lower temperature than the monolith. This leads to an enlargement of the mounting gap of 7 % to 20 % under high load and temperature. A further gap enlargement of up to 20 % is from the can deflection under the mat pressure on limited areas of the system at flat parts of the can (Oval or Racetrack cross section).

193

Figure 1: Monolith wrapping

To limit the heat expansion of the can it is common to use ferritic steels for the can with a relatively low coefficient of thermal expansion. The effects on the gap are also reduced, if enough cooling air comes to the surface of the part, the absolute gap height is enlarged (more space and mount material required) or a material with better insulation properties is used (Figure 2).

Figure 2: Gap enlargement

The actual loads on the system are measured at the engine in the car or on a test bench. Temperatures are taken with thermocouples and infrared scanners as Figure 3 shows:

Figure 3: Infrared thermography of a closed coupled catalyst with ceramic monoliths under full load

194 The mechanical loads in the system are taken by water cooled accelerometers. Piezo–accelerometers are mounted on different measuring points of the whole exhaust system in three directions (multiaxial) and the system is stimulated by the engine at all revs. and loads. The maximum acceleration in the flow direction at the mount system is calculated out of the directional values. To withstand the mechanical loads under thermal changes in the system the mat has to hold a high pressure on the monoliths surface. A limitation for the maximum value is given by the strength of the monolith. Ultra-Thin-Wall Substrates show remarkaby less strength and larger spread of strength in serial production than common products. As an example the following strength data shows almost comparable average strength for new products in relation to todays products (6/400) with 6 mil wall thickness and 400 cells per square inch (acceptable for 4/600, 4/400 and 3/600).

Figure 4: Isostatic strength of thin wall substrates, relative to standard substrates 6/400

A point of concern is the lower strength of a significant number of parts as the lower wall thicknesses lead to a broader strength distribution in the production process. Therefore the usable maximum pressure on the mat is limited to lower values to avoid monolith crushes in mass production. With lower wall thicknesses the actual shapes of the parts show bigger deviations than standard substrates. The contour tolerance, describing a mixture of dimensional and shape deviation is now merely used for shape deviations, which are hard to overcome with monolith measurements and calibrated canning systems. As a result higher mechanical and thermal loads on todays and future catalytic converter systems in close coupled position allow lower maximum mounting pressures on the mat systems. The mounting has to show more constant pressures under all loads and tolerances of the parts. This can be done by larger gaps and thicker mount mats, lowering the relative influences on the mat or with better mount materials.

195

3

Mounting Systems

To fix the substrates the mat system has to exert higher forces than the load. This is done by: · The radial pressure in the system [p] · The contact area between mounting mat and monolith/can [U] · The coefficient of friction between the parts [m] F=p×U×m Pressure and friction change their values in service with time, temperature, humidity and corrosion. Practical values are: p = 10 ... 1000 kPa, m = 0.2 ... 0.6 The main factor to influence the holding force in canning is the radial pressure on the mounting mat (the gap). These forces are applied primarily in the canning of the monoliths. There are mainly three methods to do the mounting (Figure 5).

Figure 5: Half shell canning and tourniquet technique (wrapping)

Figure 6: Tube stuffing

All these methods compress the mat to a predetermined height around the monolith.They are all in serial production with their special pro’s and con’s. While the half shell canning has fixed dimensions of the shells coming out of the tools, the tourniquet technique and the tube

196 stuffing (Figure 6) have some possibilities to match the dimension of the case to the actual dimension of the monolith in the production process. Depending on the facilities it is possible to calibrate the tourniquet can to different dimensions or to calibrate the tube for stuffing in different sizes, matching to the actual sizes. But the main part of tolerance compensation, especially if the faults in shape are larger than the size-deviations, has to be done by the mounting mat.

4

Mount Materials-Tests and Results

Today ceramic fiber mats with vermiculite as expanding papers or alumina fibermats are the common mount material to fix the ceramic monolith in the metal can. These mats fill the gap between the parts and hold the ceramic part to prevent it from moving and hitting the metallic can. The composition of these mats is seen under the microscope (Figure 7).

Figure 7: Mat composition, left: expanding paper with vermiculite, right: alumina fibermat

The expanding paper (XP) consists of vitreous alumina-silica-fibers and grainy vermiculite, a mineral that expands under high temperatures. The alumina-silica-fibers are made by blowing a molten ceramic beam into fibers and cooling it with the driving air. This leads to a broad range of fiber diameters and length and to impurities in the form of glassy beads and drops in the material. For years XP was the common material for underfloor catalysts. The fibers show sufficient elasticity to hold the monolith under low loads. With increasing temperatures at higher loads the vermiculite expands and braces the parts with high pressure against each other, which is explained in the following thermo mechanical analysis (Figure 8). As the temperature activity of the vermiculite diminished over the years, some systems showed erosion in the mount mat especially at underfloor catalysts with large gaps as a result of can deflection and low temperatures. This problem was overcome with expanding papers with less vermiculite and higher content of washed fibers. These high fiber content expanding papers showed much better resiliency after canning and worked well, even at exhaust gas tem-

197 peratures below 400 °C. The elastic effect is shown in the following compression diagram. In the second cycle the high fiber content XP shows much better resilience (Figure 9).

Figure 8: Thermomechanical analysis of XP

Figure 9: Compression of expanding papers in first and second cycle

But for new systems in close coupled position to the engine the expanding papers still showed problems at high gas temperatures above 900 °C. The whole mat loses tension as the crystal water leaves the vermiculite and the alumina-silica fibers sinter together. Furthermore the expanding papers still are very progressive in the compression graph against the gap. With common tolerances at the gap the highest pressures at narrow gaps are higher than the guaranteed isostatic strength of the Ultra-Thin-Wall monoliths. Tests on different mount systems showed, that crystalline alumina fibermats for high temperature insulations worked well under these conditions and are now under improvement. This

198 material is good for long time use at more than 1000 °C without becoming weak or brittle. These mats have lower pressure progression at narrow gaps compared to XP’s (Figure 10).

Figure 10: Pressure progression in mats

In the graph the upper curves show the pressure against the gap height in the first compression. In the car the gap opens with every drive cycle mainly because of thermal expansion at the can. With mechanical cycling in a compression tester these gap changes are simulated for 1000 cycles from an estimated gap to a maximum value as it can be seen in the actual mount and load situations. The results are the „aged“ graphs at the bottom. Without thermal expansion from the vermiculite the preheated expanding paper shows almost no resilient pressure at the large gap (high load-hot can) while the alumina fiber mat still shows reasonable values. The improvement in this system goes to lower mass of the expansive fibers (narrow gaps with worse insulation effects) and to oriented fibers in the mats to use the strength of all fibers – leading to the use of less material with the same elasticity of the mat. Also there are developments with other materials to have similar temperature stability with blown fibers or to have good elasticity at lower temperatures (diesel applications).

5

Practical Tests at the Whole System

As the mat tests only show an estimated behaviour in the laboratory, practical tests were developed and ran at the whole system. The next step is the hot shake test of the whole system in the flow direction to give the monoliths inertia load to the mount system at practical temperatures for the monolith and can. This is done by a servohydraulic actuator, while the system sees hot gas from an oil burner. With gas temperature, gas flow and an axial acceleration at a given frequency, all calibrated to the actual load in the car, the system is aged. With increased levels of load over long times the limit to failure is proved. If the system does not fail in the test, a push out test shows the left load bearing capability and is normalized as shear stress at slipping

199 of the monolith. As an example the following graph will show the maximum shear stress in the mount system for different expanding papers, gaps and temperatures (Figure 11).

Figure 11: Shear stress at XP after different preconditioning

While these tests reflect the maximum mechanical stabiltiy of the system quite well, some other effects in the car need to be simulated on the test bench with real engines. As an example remarkable erosion has been detected in systems with large gap enlargement at new expanding papers in car durability runs after very short tests (15 000 km’s). With a standard thermal shock test for 250 hours one can’t see any of these effects on the test bench. Optimised tests, related to the data of the test runs in the cars, resulted to a simulation of internal cooling and condensate. With such thermal shock tests similar effects over 50 to 80 hours at the test bench can be observed as at the car (Figure 12).

Figure 12: Car bench test, left: 15000 km car test, right: 80 h test bench

200 These tests lead to a good separation between the standard expanding papers and the different expanding papers with high fiber content. Compared to very good old expanding papers we foundacceptable materials for standard applications in underfloor condition (Figure 13).

Figure 13: Durability runs to failure at the engine test bench with different expanding papers

Alumina fibermats did not show any erosion after 500 hours in these tests, but the high material costs did not lead to the use in this underfloor position. To test close coupled systems, thermal shock tests are mainly used. As an example the engine runs for 10 Minutes with almost maximum load and speed (adjusted to the desired exhaust gas temperature) and 10 minutes at idle. Under these conditions with exhaust gas inlet temperatures up to 1050 °C in the system expanding papers failed. Sintered inlet edges and erosion in the paper led to significant damage after 20 to 50 hours. With alumina fiber mats we found no changes in the mat after durability runs under these conditions of more than 250 hours. Checking the holding forces left (shear stresses in the mount system) with push out tests after the durability runs showed remarkably high values at room temperature. Even with worst case samples with large gaps and estimating the thermal expansion of the system, the values showed good security against practical loads.

6

Conclusion

The loads on catalytic converters and how to measure them was explained and the methods to test mount materials for practical use and the general design of the system was shown. Changes in the properties of expanding papers and limits of their use in future systems require

201 new mount materials. Due to the high price of the usable alumina-fibermats they can only be recommended for critical applications. So the special properties of the mounting materials lead to the following fields of use: Expanding paper with high fiber content for underfloor catalysts Alumina-fibermat for systems with gas temperatures above 900 °C and or in combination with Ultra-Thin-Wall substrates For diesel applications XP with high fiber content is used, only in special applications with high accelerations, low temperatures or weak monoliths alumina-fibermat is recommended.

High Performance Packaging Materials M. Vermoehlen, D. Merry 3M Deutschland GmbH, Neuss,

Steffen Schmid Corning GmbH, Wiesbaden

1

Abstract

Catalytic converters have been undergoing major changes in recent years in order to meet tighter emission regulations. Catalytic converters are being moved closer to the engine, sometimes even attached directly to the exhaust manifold, in order to achieve quicker light off to reduce emissions during cold start. At the same time, there has been a trend to thinner wall monoliths often accompanied by higher cell densities that also improve emission performance. An important component of a catalytic converter is the mounting system. It has traditionally been a heat-expanding mat that acts to hold the ceramic monolith securely in place despite the significant difference in thermal expansion of the ceramic monolith and metal housing. Additionally, the mat provides a seal to prevent exhaust gas from bypassing the monolith and thermal insulation to keep to keep the housing cooler and to reduce thermal gradients in the monolith. Recently, mounting mats have been changing to reflect the more stringent requirements of close-coupled converters and/or thinner wall monoliths. This requires mats with higher temperature capability and mats with reduced holding force for the thinner wall monoliths. This paper will examine recent developments in this area of high performance mounting materials for the new requirements and compare the properties of these new mats with traditional mats.

2

Introduction

Mounting materials for catalytic converters have always had to adapt to changing requirements. When catalytic converters were first introduced in the US in the middle of the 1970’s, the predominant mounting system was wire mesh. As emission laws tightened, exhaust gas bypass through the wire mesh could no longer be tolerated and gas seals had to be added to the wire mesh to prevent gas bypass. This lead to the introduction of the first intumescent mounting mats which were invented by 3M. These mats consisted of ceramic fibers, an organic binder and a heat expandable material, vermiculite, which generates pressure as it is heated. The latter is needed to compensate for the thermal expansion difference between monolith.and.shell. As catalytic converters were introduced in Europe in the middle 1980’s, the thickness of the intumescent mats increased somewhat to compensate for the higher temperatures of the high speed German driving, but the general composition of the mats remained essentially the same.

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

203 In the middle 1990’s, certain OEM’s, particularly Japanese OEM’s began to use thin wall 400/4 (i.e. 400 cell per square inch/6 mil wall thickness) and 600/4 monoliths. These monoliths generally were considered to have lower compression strength due to their thinner walls. 3M introduced the first lower vermiculite mats to mount these somewhat weaker monoliths. Wall thickness of monoliths has continued to decrease and now non-intumescent ceramic mats are also available as a possible solution for the ultra thin-wall monoliths (600/3, 900/2, 1200/2). In addition, the temperatures in today’s close coupled and manifold mounted converters can sometimes exceed the limits of standard intumescent mats. For this reason mats with higher temperature capability have been introduced that can handle the most severe temperatures.encountered.to.date. Tests have been developed that can measure and compare various properties of mounting mats and help select the proper mat for a given application. Test descriptions and results for three new mats will be given in the next section and compared with standard intumescent mat.

3

Description of mats investigated

This paper will examine and compare the properties of four different mat types and assess their suitability as high temperature, thinwall and close coupled catalytic converter mounting materials. The four mat types are standard intumescent mat, reduced vermiculite intumescent mat, non-intumescent mat and hybrid mat, see description below: · Standard intumescent mat This type of mat consists of three basic raw materials (Table 1). Table 1: Basic raw materials of mats Ceramic fibers (alumina silicate) Unexpanded vermiculite

Organic binder

provide high temperature resiliency and contain the vermiculite provides the intumescent property i.e. creates pressure when the mat is heated and the mounting gap is opening due to the different expansion coefficient of cordierite and steel shell. provides strength to the product during manufacture and assembly

These mats are widely used in millions of converters all over the world to mount ceramic monoliths into converters. The most commonly used basis weight is 4070 g/m² mat which is appropriate for a 4 mm mounting gap. By increasing the basis weight, for example to 6200 g/m² (6 mm mounting gap), this mat can be suitable for a properly designed close coupled converter system. · Intumescent mat with reduced vermiculite level This mat consists of the same basic raw materials as standard mat but employs a lower vermiculite level and cleaner fibers. This result in a mat offering lower peak pressures at temperature and enhanced erosion resistance. · Non-intumescent mat Non-intumescent mats contain no vermiculite and use a special high alumina, polycrystalline fiber. These fibers maintain their resiliency at temperatures over 1000 °C. These fibers

204 contain essentially no shot (nonfiberous particles), resulting in excellent resiliency. Mats made with these fibers offer low, uniform holding pressure over the converter temperature range and enhanced erosion resistance. · Hybrid mat consisting of two layers The hybrid mat is a coformed mat consisting of two distinct layers - non-intumescent layer, which is positioned against the monolith - intumescent layer, which is facing the shell. This results in a mat capable of withstanding very high temperature and offering lower compression pressures as well as reduced peak operating pressures.

4

Property Comparison

Each of the four mats were evaluated for canning pressure simulation (compression test), operating pressure simulation (real condition fixture test), shear stress (resistive thermal push test) and high temperature resistance (engine dyno test). 4.1

Canning Pressure Simulation

The compression characteristics of intumescent and non-intumescent mats are important because they determine the amount of pressure exerted on the monolith during canning operations. If the compressive force is too high, monolith breakage can occur. For each mat there is specific mount density (i.e. density of the mat after canning) recommended to assure proper mat performance. Compression force is measured at target gap i.e. the gap that corresponds to the recommended mount density. To account for tolerances of monolith, can and mat, the compression force is also measured at gaps larger and smaller than target gap. F @ Speed 25.4 mm/inch

Uncompressed Thickness

Figure 1: Compression test

Compressed Thickness

205 The compression force of the mat is tested using equipment shown in figure 1. A sample diameter of 25.4 mm is used and the closing speed of the crosshead is 25.4 mm/min. Each mat is compressed to its recommended target gap and also compressed to gaps that are plus and minus 1 mm of target to simulate the tolerance stack up of the assembly. A target gap of 6.0 mm is used for all mats except the non-intumescent mat. Because of the lower thermal conductivity and higher temperature resistance of the non-intumescent mat, it can generally be used in a smaller gap than intumescent mats. Therefore, a target gap of 4.0 mm was used for the nonintumescent mat. Both the instantaneous or peak pressure value and the relaxed compression value are recorded. The relaxed pressure value is the pressure recorded 15 seconds after closure. The results of this experiment can be seen in figure 2 and 3. Figure 2 shows the peak compression values and Figure 3 shows the relaxed compression values for all four mats. The highest peak pressure recorded is for the standard intumescent mat, followed by the intumescent mat with reduced vermiculite level, and the Hybrid mat. The non-intumescent mat has the lowest peak compression pressure. Actual canning results are dependent on canning technique, know-how and closing speed, but in general the following can be concluded from this compression test: · Non-intumescent mats are suitable for all wall thicknesses. · Hybrid mats are suitable for all wall thicknesses. · Low vermiculite mats are suitable for thin-wall monoliths and possibly ultra thin-wall monoliths. · Standard intumescent mats are suitable for normal wall thickness monoliths. Peak Compression Pressures (Closing Rate 30.5 cm/min) 2500

Standard Intumescent Mat 2000

Pressure [kPa]

Reduced Vermiculite Mat 1500

Hybrid Mat

1000

Non Intumescent Mat 500

0 Target Gap +1 mm

Target Gap

Figure 2: Peak compression force at target gap and +/– 1 mm

Target Gap -1 mm

206 Relaxed Compression Pressures (15sec) (Closing Rate 30.5 cm/min) 700 Reduced Vermiculite Mat 600 Standard Intumescent Mat

Pressure [kPa]

500

Hybrid Mat

400

Non Intumescent Mat 300

200

100

0 Target Gap +1 mm

Target Gap

Target Gap -1 mm

Figure 3: Relaxed compression force at target gap +/– 1 mm

Looking now at the relaxed compression values seen in figure 3, it is clear these values are considerably lower for all gaps than the peak values. This is due to the viscoelastic nature of the mat. It is interesting to note that the relaxed compression values of the low vermiculite intumescent mat are higher than the relaxed compression values for the standard intumescent mat. This is the result of the lower vermiculite mat employing cleaner fibers, which gives the mat greater resiliency and results in less relaxation. 4.2

Operating Pressure Simulation

A real condition fixture test (RCFT) is used (see figure 4) to measure mat pressure under simulated catalytic converter conditions. The fixture consists of two heated plates and thermocouples for temperature control. The upper plate is used to simulate the hot side or monolithmat interface temperature and the lower plate is used to simulate the mat-can interface temperature. A measuring device measures the position of the plates or test gap at any given time. The heated plates, thermocouples, and gap measuring device are connected to a computer. A thermal cycle is chosen based on the actual temperatures of monolith and can that occur in a catalytic converter during operation. From these temperatures the gap change during operation is calculated. The temperature model and gap change are programmed into the computer. The computer controls the temperatures of both plates and the gap desired for each temperature condition within the cycle precisely and repeatably to these programmed conditions. The pressure exerted by the mat is recorded during the cycle. The RCFT is very useful in converter and mounting system development to understand the mat pressure generating characteristics as they relate to the converter design parameters and operating conditions.

207 Load Cell Water-cooled Heatshield

Gap Measurement Device

Heated Top Plate Heated Bottom Plate

Thermocouples Insulation Test sample

Figure 4: Real condition fixture test equipment

Operating Pressure Simulation @ Target Mount Density 1400

First Cycle Second Cycle -----------1200

Standard Intumescent Mat

1000

Pressure (kPa)

15 min. soak 800

Reduced Vermiculite Matt

600

Hybrid Mat 400

200

Non Intumescent Mat

50 50

100 60

150 95

200 130

250 155

300 180

350 215

400 250

450 275

500 300

550 329

600 358

650 387

700 416

750 445

800 474

850 502

900 530

900 530

900 480

900 430

850 325

800 220

750 185

700 150

650 125

550 85

600 100

500 70

450 60

400 50

350 45

300 40

250 38

200 35

150 30

25 25

100 25

0

skin monolith

Temperature (C)

Figure 5: RCFT measurement on four different mats

RCFT pressure curves for the four mats tested are shown in figure 5. Both first and second thermal cycles are shown for each product. There is not much change during subsequent cycles, so only the first two cycles are shown.

208 For all products there is a considerable difference in operating pressure of the first cycle as compared to the second cycle. This is because the organic binder portion of the mat burns out during the first cycle causing the pressure to decrease as it is heated. For the intumescent products, the pressure starts to increase once the expansion temperature of the vermiculite is reached. The increase in operating pressure of the hybrid mat starts later than either the standard or low vermiculite mat because the ceramic fiber layer acts as insulation. Since the ceramic fiber layer is against the hot side of the RCFT, it prevents the vermiculite from heating up as quickly. All the intumescent mats build their maximum pressure during the first cycle. Subsequent cycles show significantly lower peak pressure, but always have their maximum pressure at high temperature and lowest pressure at room temperature. The standard intumescent mat has the highest operating pressure followed by the low vermiculite mat and hybrid mat, respectively. The non-intumescent mat shows the lowest pressure at temperature. The second cycles of the intumescent mats are quite different from the first cycles. There is no pressure decrease because the binder has already been burned out. Also the pressure starts to build very quickly since the vermiculite has already been activated. Peak pressure still reaches a maximum at high temperature and minimum at room temperature. The same order of pressure is maintained i.e. the standard intumescent mat exerts the highest pressure, followed by the low vermiculite mat and hybrid mat, respectively. The operation pressure of the non-intumescent mat differs from the intumescent mats in several ways. First, the pressure does not vary as much over the entire temperature range. Second, the maximum holding force is at room temperature, and the minimum at high temperature, which is the opposite of intumescent mats. This is because it contains no intumescent material, so the pressure decreases as the gap opens up during heating. Finally, the non-intumescent mat has the lowest peak operating pressure of any of the mats making it very compatible with ultrathinwall monoliths. The hybrid mat also shows a peak operating pressure that is compatible with ultra-thinwall monoliths, but has the advantage of increased pressure at high temperature where normally the maximum vibration level occurs. 4.3

Temperature Resistance in Resistive Exposure Testing (RTE)

The temperature characteristics of the intumescent and non-intumescent materials are tested in a heated push test. The axial push test is conducted while the converter is heated. A useful way to do this employs a resistive thermal exposure (RTE) technique developed by Corning, Inc1,2. Resistive wires are inserted into the outermost cells of the ceramic monolith close to the outer skin as shown in figure 6. The ends of the wires are spot welded together to produce a continuous circuit resistance heater. The heater is attached to a temperature controller and the substrate can be heated to the desired temperature. In this way, the substrate can be aged at a predetermined temperature while having a temperature gradient across the mat. This temperature gradient is normally monitored with two thermocouples, one located on the monolith skin and the other at the can surface. The intumescent, non-intumescent (not 3M) and hybrid mats were aged using the RTE test. To increase the severity of this test and to simulate a close coupled situation the RTE converter was aged inside a furnace (see figure 6, first step).

209 Heating Wires @ 980°C

Furnace @ 550°C

Load Cell

Heating Wires temperature off

Furnace @ 550°C

Mat Mount Sample 1st step 100h aging

2nd step push test at 550°C

Figure 6: RTE test equipment

Three samples of each mat type were heated/aged to 980 °C as shown in the first step of figure 6. After this aging the mats were pushed at 550 °C (second step, figure 6). Three additional samples were pushed at the same temperature (at 550 °C) but without prior heat aging. The results can be seen in figure 7. 800

unaged 700

Residual Shear Strength @ 550°C [kPa]

600

500

400

unaged aged @ 980°C

300

aged @ 980°C 200

unaged

100

aged @ 980°C

0

Intumescent Mat

Figure 7: Results of RTE testing

Non Intumescent Mat

Hybrid Mat

210 Before the shear strength of the mat is reached slippage occurs. The residual strength therefore refers to the force needed to cause slippage divided by the area of the mat. In the RTE test the intumescent mat loses approximately 73 % of its shear strength after aging (from 725 kPa unaged to 197 kPa aged). Intumescent mats have an upper temperature limit for the mat – substrate interface of approximately 950 °C and an average mat temperature limit over the entire mat thickness of 700 °C. When temperature exceeds this limit the mat in contact with the monolith sees the high temperature first. The fibers in this area close to the substrate can become rigid and prevent the vermiculite in the outer mat layers from transmitting pressure to the monolith. At the same time some of the vermiculite is losing its effectiveness because of the high temperature. This combined effect causes the holding pressure to decrease. This effect is the same for all intumescent mats independent of their vermiculite level and explains the large reduction in shear strength. The residual shear strength of the non-intumescent mat decreased approximately 23 % (from 85 kPa unaged to 65 kPa.aged). Because of the different type of fibers used there is no rigidizing effect as described above. However the pressure level of this mat as reflected in shear strength is much lower when compared to vermiculite containing intumescent and hybrid mats. The best results are obtained with the hybrid system showing a minimal decrease of approximately 11% (from 312 kPa unaged to 276 kPa aged) and overall higher values than for non-intumescent mat. This mat construction combines the best features of both mats (i.e. the intumescent pressure building property of vermiculite containing mat with the high temperature resistance of the non-intumescent fiber mat.) The above test results of the hybrid mat encouraged us to test this construction in high temperature engine tests. 4.4

High Temperature Resistance in Engine Tests

Final testing included only the hybrid mat. Engine dynamometer testing was conducted at mat-monolith interface temperatures up to 1100 °C. The conditions were as follows: Test conditions: · 7.5 Liter V8 engine · 3000 rpm/299 Nm torque · 1 converter per exhaust bank (close coupled) · Secondary air added at exhaust manifold · 1050 °C inlet gas temperature · 1100 °C Monolith-Mat interface temperature · 100 hours test duration · · · ·

Converter design Monolith Ø 85.5 mm, 80 mm long 350/5,5 cpsi endcones with and without insulation hybrid mat

211

Figure 8: Shows the substrate wrapped with the hybrid mat

At high temperature the whole converter system is under severe stress. It is therefore important that all components are designed in the most robust way. For this reason a round substrate with insulated double wall end cones was selected. Figure 9 and 10 show the dramatic difference in converters with and without insulated double wall endcones, respectively.

Figure 9: Showing a converter system with insulated endcones

212

Figure 10: Converter design without insulated inlet and outlet cones

The converter with the insulated endcones successfully completed the 100 hour test duration. No failure and no erosion observed after disassembly. Based on the good test results Aston Martin decided to use a hybrid mat system for their new 450 horsepower 6.0 litre V12 Vanquish model. The catalyst system consists of two manifold mounted 400/4 and one close coupled 400/6 substrate. The actual converter system is shown in figure 11.

Figure 11: Aston Martin converter system

213 The Aston Martin will be launched autumn 2001. The whole car was developed together with Ford RVT (Research & Vehicle Technology). According to Autotechnology3 Aston Martin ran a test program with 50 prototypes covering more than a million miles.

5

Conclusions

The four tests that were used to determine the suitability of mats for high temperature and thin wall mounting work quite well and can be used to help select the proper mounting mat. This is especially important in areas where there is not as much long-term field experience such as in high temperature, ultra thin-wall converters. Of course, the actual design of the catalytic converter can also have a very strong positive or negative influence on the performance of a mat4 (e.g. insulated double wall end cones can have a very positive influence, whereas insulated heat shields can have very negative influence). For any given converter design selection of the right mat is essential. Based on results of the four mats tested, the following conclusions can be made. · For standard applications where temperature requirements are within recommended guidelines standard mats are an economically and technically sound solution. Many of today’s applications are using this mat in close-coupled situations. It is important that the overall design is robust e.g. uses insulated cones etc. A recent SAE paper5 describes the durability of a thin-wall, converter system in a close coupled application in a high speed road test · Low vermiculite mats are better suited for thin wall monolith mounting than standard intumescent mats but have the same temperature limitations. · Non-intumescent mats are well suited for ultra-thin-wall monoliths and also for converters exceeding temperature guidelines for standard intumescent mats. Their disadvantage is reduced pressure performance caused by increasing gap at temperature. Maximum holding force is realized at low temperature. · Hybrid two layer constructions can be used for all described applications. They offer high temperature capability, lower peak pressures suitable for ultra thin-wall applications and increased holding pressure with temperature.

6

References

[1]

K.P Reddy, J.D. Helfinistine, and S.T. Gulati, Corning Inc “New test for characterizing the durability of a ceramic catalytic converter package”, SAE Paper 960559 R.J. Locker and C.B. Sawyer, and M.J. Schad, Corning Inc. “Quantification of ceramic preconverter hot vibration durability”, SAE Paper 960563 Auto Technology Volume No 1 August 2001 P.D. Stroom, R.P. Merry, 3M Company and S. T. Gulati, Corning Inc. “System approach to packaging design for automotive catalytic converters”, SAE Paper 900500 J. Kallenbach, and P. Floerchinger, Corning GmbH Wiesbaden “Durability of a thinwall converter system in close coupled application“, SAE 982897E, SAE Brazil 98

[2] [3] [4] [5]

IV Catalysts

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

Three-Way Catalyst Deactivation Associated With Oil-Derived Poisons Joseph Kubsh Engelhard Corporation, Environmental Technologies Group Iselin, New Jersey, U.S.A.

1

Introduction

During much of the 1990s the development efforts associated with three-way catalysts for automotive applications focused on improving the thermal durability of these catalysts (1–3). This stemmed largely from the general trend in the industry of moving these catalysts closer to the engine in order to accelerate their activation during the early stages of vehicle operation. These thermal durability development efforts resulted in the commercialization of new generations of three-way catalysts capable of long-term operation in exhaust environments at temperatures in excess of 1000 °C. The success of developing catalysts with high temperature durability has more recently focused more attention on the deactivation mechanisms for threeway catalysts associated with the deposition of poisons present in the exhaust environment. More specifically, deactivation of three-way catalysts by poisons derived from the consumption of engine lubricating oils such as phosphorus and zinc has been revisited (4–6). These poisoning mechanisms have been found to be of importance in advanced emission systems targeting ultra-low emission performance such as the recently adopted California Low Emission Vehicle II Program (LEV 2) requirements for ULEV2 (ultra-low emission vehicle 2) and SULEV (super ultra-low emission vehicle) classifications. In these near-zero tailpipe emission applications catalyst light-off often dominates a vehicle’s emission performance. Poison deposition on the inlet section of a close-coupled converter could be sufficient to depress the catalytic light-off characteristics of the three-way catalyst outside of these very low emission performance limits. The aim of this work was to revisit oil-derived poisoning issues on three-way catalysts. In particular an investigation was made into the development of an engine-based, catalyst aging protocol that could be used to more closely impart the impacts of oil-derived poisons found under real world driving conditions.

2

Experimental Details

Previous studies aimed at understanding the impacts of oil-derived poisons on three-way catalysts have used a variety of accelerated engine-based aging protocols. In most cases these aging protocols have attempted to accelerate catalyst deactivation by using high levels of oil additives that contain catalyst poisoning agents such as phosphorus or heavy metals such as zinc. An example of one such lubricant additive is zinc dialkyldithiophosphate (ZDDP). An additive such as ZDDP can be incorporated into an aging protocol in several manners to deac-

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

218 tivate a catalytic converter. These options include: 1) doping gasoline with the oil additive of interest, 2) doping the engine oil with higher levels of the additive, or 3) injection of doped oil containing the additive into the exhaust manifold upstream of the catalytic converter. In all cases the intent is to expose the catalyst to high levels of poison precursors that deposit on the catalyst surface and deactivate catalyst performance. In this study, the option of oil injection was selected for investigation. Initial investigations indicated that the effectiveness of this oil injection procedure in deactivating the catalyst was a strong function of the exhaust temperature present in the catalytic converter, as well as the time of catalyst exposure under potential catalyst poisoning situations. For example, using a 75 h catalyst aging protocol in which the bed temperature of the catalyst never reached temperatures below 850 °C, catalyst performance remained similar with and without ZDDP-doped oil injection into the engine exhaust manifold upstream of the converter. In contrast, dropping the exhaust temperature to 450 °C and exposing the catalyst to an additional 24 h of aging with oil injection resulted in a significant degree of catalyst deactivation as measured by the hydrocarbon light-off characteristics of the catalyst. The amount of ZDDP present in the oil injected in the manifold was also expected to have a strong impact on the level of catalyst deactivation observed in such an accelerated engine aging protocol. As a result, a design of experiments approach based on Box-Benkhen methods was used to investigate the impacts of these three key parameters (aging temperature, aging time, level of ZDDP present in the oil used for injection into the exhaust manifold) on three-way catalyst performance. 2.1

Design of Experiment Aging Variable Definition

A 75 hour engine aging protocol was used as the basis for this investigation into catalyst poisoning. The aging protocol contained two primary modes of operation: a high temperature, exothermic mode with a catalyst inlet temperature of 850 °C (maximum catalyst bed temperature of 1000 °C), and a low temperature, mode with a catalyst inlet temperature of either 600 °C or 400 °C. At this low temperature mode the aging engine is run at an idle condition. When oil-containing ZDDP was injected into the exhaust manifold of the engine upstream of the catalyst, this injection occurred during both primary modes of engine operation. The oil injection rate was 0.95 liter of oil over 24 hours of engine operation. For the injected oil, the ZDDP level was doped to be 1.5 wt.%. In some aging experiments no manifold oil injection was used, in which case the catalyst exposure to potential oil-derived poisons stemmed only from the oil consumption rate of the engine itself. This oil consumption rate was also approximately 0.95 liter of oil over 24 hours of engine operation. The concentration of ZDDP in the oil used for normal engine lubrication was 0.1 wt.%. Table 1 provides details on the levels selected for each of the three key design variables: low temperature aging temperature, time at the low temperature aging mode, and wt.% ZDDP in the oil exposed to the catalyst. In aging runs where the low temperature mode was set to be 50 % of the total aging time, the engine was cycled between high temperature and low temperature modes every 20 minutes. For runs in which the low temperature mode was only 20 % of the total aging time, a 20 minute low temperature mode was followed by 80 minutes of time at the high temperature mode.

219 Table 1: Aging cycle parameters for catalyst poisoning studies Aging run 1 2 3 4 5 6 7 8

Low temperature Injection (°C) 600 600 600 600 400 400 400 400

Time at low temperature mode (%) 50 50 20 20 50 50 20 20

ZDDP in oil (wt.%) 0.1 1.5 0.1 1.5 0.1 1.5 0.1 1.5

The catalyst used in all experimental aging runs was identical. A Pd-only, three-way catalyst was used with a total Pd content of 1.90 g Pd/liter of monolith. This catalyst was coated on a ceramic honeycomb monolith with 62 cells/cm2 (0.165 mm wall thickness). The total volume of catalyst present in each experimental converter was 0.69 liters.

3

Aging Results

After aging the converters with the aging cycle options discussed above, catalyst light-off characteristics were used as a measure of catalyst activity. Light-off characteristics of the aged converters were measured on an engine using a standard procedure with a stoichiometric exhaust gas composition. This light-off test procedure used a temperature ramp rate of approximately 20 K/min with a catalyst space velocity of approximately 80,000/hr (at STP). The light-off curves (catalyst conversion vs. inlet exhaust temperature) generated in this fashion were then used to determine the appropriate T50 light-off temperatures (inlet exhaust temperature at which 50% conversion efficiency is observed) for each converter with respect to both hydrocarbons and NOx emissions. These light-off characteristics became the observed outputs for the design of experiment analysis. Analysis of the data was completed using Design-Expert (version 5) software. This software package generated the typical contour plots that stem from the statistical analysis of a Box-Benkhen design. From these contour plots the relationships between the design variables and catalyst performance (T50 light-off temperatures) can be observed. Representative interaction results that stem from the statistical analysis are shown in Figure 1 for catalyst hydrocarbon and and NOx performance. These results are interpolated from the data set for an aging cycle that uses 35 % of the total time at the low temperature mode. With respect to both hydrocarbons (HC) and NOx, catalyst light-off performance degrades when oil is injected into the engine exhaust manifold upstream of the converter (i.e., exposure to higher ZDDP concentrations results in more severe catalyst deactivation). The degree of performance degradation is also more severe when the low temperature aging mode is lowered from 600 °C to 400 °C.

Temperature for 50% Efficiency (C)

220 500 450 400

HC

350

NOx

300 250 600 C, no oil injection

600 C, oil injection

400 C, no oil injection

400 C, oil injection

Aging condition (35% duration for low temp. mode) Figure 1: Impact of aging conditions on catalyst light-off for hydrocarbons (HC) and NOx

4

Aged Catalyst Characterization

Catalysts aged with the various two-mode aging cycles described previously were also characterized for the levels of P and Zn accumulated on the catalyst during the aging protocol. These characterizations included chemical analyzes for total P and Zn levels on the catalyst and electron microprobe investigations to determine the distribution of P and Zn within the catalytic coating. For chemical analyses of P and Zn, aged catalysts were cut in half so information regarding the axial distribution of these poisons could be determined. In general these chemical analyses showed the expected higher distributions of poisons at the inlet side of the converter with lower poison concentrations at the outlet end of the converter. Converters aged with oil injection into the exhaust manifold, and its higher ZDDP levels, showed more than an order of magnitude higher P and Zn levels on the catalyst compared to converters aged without oil injection. For example, with oil injection P levels on the front half of the aged converters were in the 4-5 wt.% range, compared to 0.1-0.25 wt.% for converters aged without oil injection. Poison concentrations in the rear half of the aged catalysts were in general no more than 50% of poison concentrations found in the inlet half of the catalyst. P and Zn levels measured on the converters aged with oil injection were slightly higher for the aging runs completed with the longer duration low temperature mode at 400 °C compared to the 600 °C case. Aging completed with oil injection and the shorter duration low temperature mode showed similar levels of P and Zn for both the 400 °C and 600 °C cases. Based on the throughputs of P and Zn associated with the oil consumption of each aging run and the chemical analyses done on the aged catalysts, the capture rate of P and Zn for each aged catalysts was calculated. Figure 2 summarizes these capture efficiencies for aging runs done with oil injection. With oil injection into the exhaust manifold and the high levels of ZDDP used in the injected oil, capture efficiencies for P and Zn in these aging runs were always greater than 20 % and 10 %, respectively. Consistent with the lower levels of P and Zn measured on the catalysts aged without oil injection, capture rates for P and Zn on aging runs completed without oil injection were significantly lower as shown in Figure 3. P capture efficiencies without oil injection were always less than 9 % and less than 3 % for Zn.

Capture Efficiency (%)

221

40 30 20 10 0

P Zn 600 C, 50% low T 600 C, 20% low T 400 C, 50% low T 400 C, 20% low T Aging condition (all with oil injection)

Capture Efficiency (%)

Figure 2: Capture efficiency of P and Zn by the catalyst during aging with oil injection

10 8 6 4 2 0

P Zn

600 C, 50% low 600 C, 20% low400 C, 50% low 400 C, 20% low T T T T Aging condition (all without oil injection)

Figure 3: Capture efficiency of P and Zn by the catalyst during aging without oil injection

Microprobe analyses of catalysts aged with oil injection in the exhaust manifold showed significant penetration of P and Zn into the washcoat materials coated on the ceramic substrate. The amount of P and Zn, as well as the degree of penetration, was most severe at the inlet end of the catalyst. Consistent with the elemental analyses, the poison levels observed by microprobe decreased in moving toward the rear, outlet end of the catalyst.

5

Discussion and Summary

The results presented here indicate that oil-derived catalyst poisons such as P and heavy metals like Zn are most significantly deposited on three-way catalysts during low temperature engine operating modes. The design of accelerated catalyst aging protocols require modes with low temperature operation with high levels of potential poisons in order to impart the kinds of catalyst deactivation associated with poison accumulation on the catalyst surface. Combining low temperature modes that facilitate poison deposition on the catalyst surface with higher temperature modes appears to drive the poisons deeper into catalyst structure. The aging cycle developed in this study makes use of oil injection doped with high levels of ZDDP upstream of the converter in combination with both low temperature and high temperature operating modes. Adjustable aging parameters such as the level of ZDDP present in the injected oil, the relative duration of the low temperature mode, and the temperature associated with the low temperature mode can all be used to try and create poison profiles and catalyst performance observed after real world vehicle operation. The penetration of P and Zn into the

222 washcoat of three-way catalysts observed by microprobe on the catalysts aged here, has been observed on real vehicle fleets that are known to have large relative consumption levels of engine oil and considerable operating time under low speed, low temperature exhaust conditions. Significantly higher catalyst light-off temperatures for regulated emissions such as hydrocarbons and NOx characterize the catalyst deactivation observed with this oil injection aging protocol and under real world conditions that favor high P and Zn levels on three-way catalysts. As discussed by Darr et al. (6), even modest amounts of poison deposition on a close-coupled catalyst’s inlet face that lead to modest deactivation of a catalyst’s light-off characteristics could be the difference between meeting a SULEV emission requirement or failing to meet these near-zero emission levels. In order to minimize these negative impacts on catalyst performance stemming from oil-derived poisons, it is critical that engine lubricant consumption be maintained at low levels throughout the catalyst regulated lifetime, or alternative lubricant additive packages that minimize the levels of potential poisons such as P and Zn be developed for emission critical applications.

6

References

[1] [2] [3] [4] [5] [6]

Z. Hu and R. Heck, SAE Paper No. 950254 (1995). P. Andersen and J. Rieck, SAE Paper No. 970739 (1997). P. Andersen and T. Ballinger, SAE Paper No. 1999-01-0308 (1999). J. Thoss, J. Rieck, and C. Bennett, SAE Paper No. 970737 (1997). D. Ball, A. Mohammed, and W. Schmidt, SAE Paper No. 972846 (1997). S. Darr, R. Choksi, C. Hubbard, M. Johnson, and R. McCabe, SAE Paper No. 2000-011881 (2000).

223

Catalytic Reduction of NOx in Oxygen-Rich Gas Streams, Deactivation of NOx Storage-Reduction Catalysts by Sulfur Ch. Sedlmaira,b, K. Sehanb, A. Jentysa and J. A. Lerchera a b

Technische Universität München, Institut für Technische Chemie II, Garching, Germany University of Twente, Faculty of Chemical Technology, Enschede, The Netherlands

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

224

225

226

227

228

Catalytic Reduction of NOx in Oxygen-Rich Gas Streams: Progress and Challenges in Catalyst Development Wolfgang Grünert Lehrstuhl Technische Chemie, Ruhr-Universität Bochum

1

Introduction

The forthcoming introduction of tighter emission regulations for car exhaust has created much research effort aiming at new NOx abatement technologies for Diesel and lean-burn engines. This task presents the challenge to reduce NO and NO2 to nitrogen in the presence of excess oxygen, which tends to compete for the reductant employed. Such “selective catalytic reduction (SCR) is well known from flue-gas treatment in power plants where the reductant is ammonia. Since the storage of ammonia in cars is impractical, different solutions have been sought, most of them involving heterogeneous catalysts. The catalytic approaches may be divided into storage-reduction and SCR processes, the latter being differentiated according to the reductant used – hydrocarbons, urea (as an ammonia source), or soot. In the NOx storage-reduction (NSR) approach [1], NOx is stored on a catalyst that contains a storage medium (mostly barium compounds) together with a noble-metal component. When the storage capacity becomes exhausted, the stored nitrate is catalytically reduced during an excursion into rich regime (fuel injection). With real exhaust, the sulfur present in the fuel blocks the storage component rapidly and has to be removed by a high-temperature treatment. It is not yet clear if the forthcoming reductions of fuel sulfur content will be sufficient to pave the way for commercialization of NSR. Selective catalytic reduction of NOx by hydrocarbons (HC) was discovered during research about other NOx abatement approaches – SCR with urea [2] and NO decomposition [3]. The latter has not gained practical interest since then. SCR with hydrocarbons would be the most practical approach utilizing the fuel as an on-board reductant. However, the selectivity of hydrocarbons toward NO (as opposed to O2) is much less than that of ammonia, hence the technique involves a severe fuel penalty. Moreover, SCR with hydrocarbons is much more demanding toward the catalyst than SCR with ammonia (urea), hence progress in catalyst development was unexpectedly slow. On the other side, SCR with urea involves the problem of urea distribution while the catalytic processes (including a urea decomposition stage) appear to be well controlled. There is, however, concern about N2O that may be formed under certain reaction conditions. In addition, the catalysts contain vanadium, which should not be spread into the environment. SCR with soot describes the attempt to use NOx for the combustion of soot held in traps with walls that contain catalytically active materials. This should be differentiated from the Continuously Regenerating Trap (CRT), where soot is burned non-catalytically by NO2 previously produced from the NO present on a noble-metal catalyst. In the CRT, NO2 is reduced only to NO while N2 is the desired product of SCR with soot. Among these techniques, SCR with urea is already commercialized [4] while NSR and SCR with hydrocarbons have real chances for a breakthrough to practical application [5]. While

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

230 NSR is subject of another paper in this book, the present contribution summarizes recent progress in catalyst development for SCR with hydrocarbons and ammonia.

2

Selective Catalytic Reduction with Hydrocarbons

2.1

Basic Problems

SCR with hydrocarbons was originally found with ZSM-5 catalysts highly loaded with copper, which were considered most promising for quite a long time. Meanwhile, many catalyts have been discovered [6–8], most of them consisting of redox components introduced into zeolite matrices. Analogous catalysts based on refractory-oxide supports, e.g. Al2O3, were long considered insufficiently active. A typical feature of most SCR catalysts is a relatively narrow temperature window, in which NO is reduced to N2. Figure 1 shows this for several catalysts. It reflects mostly the situation found in studies with small molecules (propane, propene, isobutane), which are preferred in academic studies. The use of higher hydrocarbons sometimes leads to a significant broadening of this temperature window.

100

a

NO conversion, %

80

c 60

d 40

b

20

0

500

600

700

800

Figure 1: Temperature dependence of NO conversion found with typical SCR catalysts. a) Cu-ZSM-5, reductant propene, 32 000 h–1, from [9], b) Fe-ZSM-5, reductant isobutane, 30 000 h–1 [10], c) Fe-La-ZSM-5, reductant isobutane, 42 000 h–1, from [11], d) Pt/Al2O3, reductant propene, ca. 100 000 h–1, from [12]

231 The SCR reaction appears to be initiated by an oxidative attack of adsorbed NO2 or nitrate intermediate on the HC reductant, which subsequently reduces the N species. This constitutes a high reductant excess (according to stoichiometry, one CH4 molecule reduces 4 NO molecules), and the non-used parts of the reductant molecule are apparently prone to unselective attack by the oxygen present. This leads to a rather modest hydrocarbon utilization, which generally decreases with increasing reaction temperature. It is in the order of 8 %–10 % for good catalysts and small HC molecules at the temperature of peak NO conversion and tends to decrease with increasing chain length of the reductant. Only for methane, which is however difficult to activate, HC utilization may be in the order of 40 % at peak NO conversion. Although Cu-ZSM-5 is considered particularly bad in this respect (which depends, however, on the reductant considered), the catalysts found subsequently offer only gradual improvements of hydrocarbon utilization. Selectivity to N2 is another important issue in SCR with hydrocarbons because N2O may be formed as an undesired side product. Indeed, N2O selectivities between 70 % and 20 % [12] are the major drawback of supported Pt catalysts, which are promising in many other respects (high activity at low temperatures – though with narrow temperature window, poisoning resistance, durability). This disadvantage has not been overcome by now and rules out Pt-based SCR catalysts for practical application. Formation of N2O (even of NO2, at total HC consumption [13]) may be also a problem with oxide catalysts while N2 is almost exclusively formed over redox zeolites. The major obstacles to the practical application of the HC-SCR approach have been, however, the poisoning of many catalysts by H2O and SO2 and the durability problem. While Cu-ZSM-5 is only moderately poisoned by water and SO2 [3], it proved unstable in real exhaust, at least at temperatures above 500 °C [14]. Major goals of catalyst development have been, therefore, to avoid N2O formation with Pt-based systems, to stabilize redox zeolites without compromizing activity properties, and to increase the activities of the more rugged oxide-based systems. It appears that the best success has been achieved so far with the latter catalyst type. 2.2

Examples of new catalyst developments

The effort to counteract poisoning and to stabilize redox zeolite catalysts for the use in real media has lead to several interesting results. Thus, it was found that In-ZSM-5, which belongs to the few systems able to activate methane for the SCR reaction, can be effectively promoted by noble metals [15] and by CeO2 [16]. Examples of the latter case are given in Figure 2, where the pronounced effect of physical CeO2 admixture is demonstrated. In dry medium, NO conversions of >70 % are achieved even at 100 000 h–1 with the methane reductant. Unfortunately, at lower temperatures, which are more relevant to application, Ce-In-ZSM-5 suffers strongly from water poisoning. Figure 2 shows also activity data for a catalyst prepared via a different route (RSSIE, reductive solid-state ion exchange, cf. [16]), where an increased In content resulted in improved durability at a moderate loss of activity. Meanwhile it was found that the low-temperature activity of this system is improved when other hydrocarbon reductants are used. Notably, no loss of activity was found in SCR with higher hydrocarbons after 162 h treatment of the catalyst at 773 K in moist air [17].

232

100

Mixture CeOx/In-ZSM-5

X(NO), %

-1

80

30 000 h -1 100 000 h

60

Ce-In-ZSM-5 (RSSIE) -1 24 000 h 7 % H 2O

40

20

In-ZSM-5 30 000 h

0

600

700

800 T, K

-1

900

Figure 2: NO conversions in the SCR with methane over Ce-promoted In-ZSM-5. In-ZSM-5 with and without CeO2 physically admixed; Ce-In-ZSM-5 catalysts prepared by RSSIE of In and precipiation of Ce [16]. 1000 ppm NO, 1000 ppm methane, 2 or 10 % O2, water content and space velocity indicated in the panel.

Cobalt-exchanged ZSM-5 was the first catalyst found to activate methane for the SCR reaction [18]. It is also useful with other reductants [19], but its stability in feed containing water and SO2 was insufficient although these poisons alone did not deactivate the catalyst [20]. Remarkable progress was made by introducing Co ions into zeolite Beta [21, 22]. This catalyst, which exhibits remarkable stability (cf. Figure 3) is being commercialized for applications with stationary sources. 100

Co-Beta

X(NO, CH4)

80 60

CH4

CH4

40

NO

Co-ZSM-5

20

NO 0

0

1000

2000 3000 Time-on-stream, h

4000

Figure 3: Stability test with Co-ZSM-5 and Co-Beta. Feed: 150 ppm NO, 500 ppm C3H8, 1000 ppm CH4, 10 % O2, 9 % H2O, 0.3 ppm SO2, 500 ppm CO, 250 ppm H2; T = 673 K, 15 000 h–1. From [21].

233 Recently, much attention has been paid to “overexchanged” Fe-ZSM-5, in which iron is introduced into ZSM-5 to a Fe/Al atomic ratio of 1. This interest was generated by work of Feng and Hall [23] describing excellent activity and stability for overexchanged Fe-ZSM-5 obtained by a preparation, which has never been reproduced despite considerable effort. Later, Chen and Sachtler [24] proposed a preparation based on chemical vapor deposition of FeCl3, which produces catalysts that are less active but still have high resistance to poisoning and encouraging stability properties. Figure 1 gives some activity data for this type of catalyst. The best results were obtained with La-promoted materials [11]. Figure 4 illustrates the stability of the Fe-La-ZSM-5 catalyst at the temperature of peak NO conversion [11]. The slow activity decay is due to coke deposits, hence, catalyst regeneration is straightforward.

Conversion, Yield, %

100

NO 80

isobutane 60

CO

40

CO2

20 0

10 % O2/He 2 h, at 773 K 0

20

40 60 80 100 Time-on-stream, h

120

Figure 4: Stability test with Fe-La-ZSM-5. 2000 ppm NO, 2000 ppm isobutane, 3 % O2, 10 % H2O, T = 623 K, 42 000 h–1. From ref. [11].

Fe-ZSM-5 remains a challenge for academic research since recent, still unpublished results of several groups suggest that the original claims by Feng and Hall [23] were real, but that different ways have to be found to obtain the corresponding catalyst structure. The present versions of overexchanged Fe-ZSM-5 are considerably less active with other hydrocarbons (propane, propene). Severe durability tests have not been reported by now. However, there is considerable potential for improvement even with Fe-ZSM-5 prepared by CVD of FeCl3. Alumina-based SCR catalysts were developed by several groups, typically with Ag or Co promoters. While their stability was attractive, their activity was insufficient. Significant progress has been made recently by application of a new mesoporous alumina developed in the Institute of Applied Chemistry Berlin-Adlershof. Figure 5 shows activity data of a material containing 0.6 % Ag and 0.4 % Co on this alumina [26]. High conversions are obtained in a broad temperature window with the reductant n-decane. The N2 selectivity, which is sometimes a problem with alumina-based SCR catalysts, was >90 % at all temperatures. A completely different, interesting concept was proposed by Iwamoto [26]. Based on the fact that NO2 is an intermediate in many catalytic systems, the NO oxidation (mostly over Pt) and the remaining steps are performed in separate catalyst layers at different temperatures, with intermediate reductant addition. Significant improvements of activity and HC utilization were obtained.

234

4

Selective Catalytic Reduction with Ammonia (Urea)

Although there are many potential catalysts for the SCR with NH3, the use of promoted V2O5/ TiO2 catalysts, in particular V2O5/WO3/TiO2 is predominant in practical applications. These catalysts are very active and robust and have, therefore, been employed also in NOx converters for diesel engines. Similar to SCR with hydrocarbons, the reductant ammonia becomes oxidized at high temperatures. Over SCR catalysts based on the V2O5/TiO2 system, this oxidation is not selective for the harmless product N2, but NO and N2O are also formed. The former limits the NO conversion, while the latter will be emitted.

100

N2 yield, %

80

60

60

40

40

20

20

0 500

Decane conversion, %

100

80

0

600

T, K

700

800

Figure 5: Catalytic behavior of a Co-Ag/Al2O3 catalyst. 1000 ppm NO, 550 ppm decane, 6 % O2, 10 % H2O, 12 % CO2, at ca. 100 000 h–1 [25]

Recently, highly active catalysts for the SCR of ammonia have been developed on the basis of redox zeolites. Among the new catalysts are Ce-MOR [27], Cu-zeolites (Y, MOR, ZSM-5, [28]), and overexchanged Fe-ZSM-5 [29, 30]. Figure 6 shows the example of Fe-ZSM-5 pre-prepared by CVD of FeCl3. With a 1:1 NO/NH3 feed, NO conversions of 80 % are obtained at 300 000 h–1 over a wide temperature range, and the reaction is even promoted by moisture in the feed. With slight NH3 excess, the conversions are even higher, and the excess NH3 is oxidized to N2 over the whole temperature range [29]. In a 50 h durability test at 848 K (16 h each in dry feed, moist feed (2.5 % H2O) and SO2-containing feed (+300 ppm SO2)), no evidence of deactivation was found. An interesting stabilization effect has been observed with Cu-ZSM-5 catalysts, the activity of which is similar to that of overexchanged Fe-ZSM-5 [31, 32]. Steaming at 873 K (17 h, 3 % H2O/Ar) led to a significant loss of NO conversion at all reaction temperatures studied. However, this effect was absent when the zeolite crystals were supported on a Ni surface (Raney Ni), instead, the NO conversion at high temperatures increased after steaming. It was proposed that Ni ions migrate into the zeolite where they become catalytically active and exert a stabilizing influence on the zeolite matrix[32].

X(NO), X(NH3), %

235

100 + 2.5 % H O 2 NH3 NO 80

dry

60 40

CVD

conventional

20 0 500

600

700

800

900

Figure 6: SCR of NO with ammonia over Fe-ZSM-5 prepared by CVD of FeCl3, comparison with conventional preparation. 1000 ppm NO, 1000 ppm NH3, 2 % O2, 300 000 h–1. From [29].

5

Conclusions

Since the early days of SCR with hydrocarbons, new materials have been found that give rise to the expectation of a breakthrough of this approach to commercial use including mobilesource applications. New zeolite-based catalysts for SCR with ammonia also offer an interesting application potential.

6

References

[1] [2] [3] [4] [5]

N. Takahashi, H. Shinjoh, T. Ijima, T. Suzuki, et al., Catal. Today 1996, 27, 63. W. Held, A. König, T. Richter, L. Puppe, SAE-Paper 1990, No. 900 496. M. Iwamoto, H. Yahiro, S. Shundo, Y. Yu-u, N. Mizuno, Shokubai 1990, 32, 430. G. Fränkle, W. Held, W. Hosp, W. Knecht et al., VDI-Ber. Ser. 12, 1997, 365. J. Armor, “Opportunities, Strategies and Innovation for Catalysis in the New Millenium” (Plenary lecture), Europacat-V, Limerick (Ireland), September 2001. A. Fritz, V. Pitchon, Appl. Catal. B, 1997, 13, 1. V. I. Parvulescu, P. Grange, B. Delmon, Catal. Today 1998, 46, 233. Y. Traa, B. Burger, J. Weitkamp, Microporous Mesopor. Mat., 1999, 30, 3. T. Liese, W. Grünert, J. Catal. 1998, 172, 34. F. Heinrich, E. Löffler, W. Grünert, Stud. Surf. Sci. Catal. 2001, 135, 30–P-27. H.-Y. Chen, W. M. H. Sachtler, Catal. Lett. 1998, 50, 125. R. Burch, P. Millingham, Catal. Today 1995, 26, 185. F. C. Meunier, J. P. Breen, V. Zuzaniuk, et al., J. Catal. 1999, 187, 493. D. R. Monroe, C. L. Di Maggio, D. D. Beck et al., SAE-Paper 1993, Sp-968, 195. E. Kikuchi, M. Ogura, N. Aratani, Y. Sugiura, et al., Catal. Today 1996, 27, 35.

[6] [7] [8] [9] [10] [11] [12] [13] [14] [15]

236 [16] F.-W. Schütze, H. Berndt, M. Richter, B. Lücke, C. Schmidt, T. Sowade, W. Grünert, Stud. Surf. Sci. Catal., 2001, 135, 10-O-01. [17] F.-W. Schütze, H. Berndt, to be published. [18] Y. Li, J. N. Armor, Appl. Catal. B 1992, 1, L31. [19] S. Sato, Y. Yu-u, H. Yahiro, N. Mizuno, M. Iwamoto, Appl. Catal. 1991, 70, L1. [20] Y. Li, J. N. Armor, Appl. Catal. B 1993, 3, L257. [21] T. Tabata, M. Kokitsu, H. Ohtsuki, O. Okada, et al., Catal. Today 1996, 27, 91. [22] T. Tabata, H. Ohtsuki, et al., Microporous Mesopor. Mat. 1998, 21, 517. [23] X. Feng, W. K. Hall, Catal. Lett. 1996, 41, 45. [24] H.-Y. Chen, W. M. H. Sachtler, Catal. Today 1998, 42, 73. [25] M. Richter, M. Langpape, S. Kolf, R. Fricke, lecture at Europacat-V, Limerick (Ireland), September 2001. [26] M. Iwamoto, Stud. Surf. Sci. Catal. 2000, 130, 23. [27] E. Ito, R. J. Hultermans, M. H. W. Burgers et al., Appl. Catal. B 1994, 4, 95. [28] S. Kieger, G. Delahay, B. Coq, B. Neven, J. Catal. 1999, 183, 267. [29] A. Z. Ma, W. Grünert, J. Chem. Soc., Chem. Commun. 1999, 71. [30] Q. Sun, Z. X. Gao, H. Y. Chen, W. M. H. Sachtler, J. Catal. 2001, 201, 89. [31] A.-Z. Ma, M. Muhler, W. Grünert. Chem. Eng. Technol. 2000, 23, 3. [32] A.-Z. Ma, M. Muhler, W. Grünert, Appl. Catal. B 2000, 27, 37.

Atomic Structure of Low-Index CeO2 Surfaces Holger Nörenberga, J. H. Hardingb and S. C. Parkerc a

University of Oxford, Department of Materials University College London, Department of Physics and Astronomy c University of Bath, Department of Chemistry b

1

Introduction

Polycrystalline cerium oxide is used in catalytic converters for cars as an additive to the washcoat [1]. The cerium oxide acts as an oxygen buffer for reduction and oxidation reactions because it can easily transform between CeO2 and Ce2O3. Polycrystalline materials contain crystallites, which mainly exhibit low-index surfaces because they are usually the low-energy surfaces. We have chosen single crystals of low index (111) and (001) orientation to study the atomic structure of CeO2 because polycrystalline material is unsuitable for detailed investigation by Scanning Tunnelling Microscopy (STM). The geometric and electronic structure of sufficiently conductive surfaces can be studied by STM with atomic resolution.

2

Experiments and Modelling

Commercial Crystal Labs Inc. supplied single crystals of CeO2, cut and polished to obtain lowindex surfaces. The crystals were placed in the UHV-chamber and then annealed to remove surface contamination [2]. STM experiments were carried out at elevated temperatures with a JEOL JSTM 4500-XT at a base pressure in the low 10–10 mbar range at temperatures up to 400 ºC. Auger-electron spectroscopy and X-ray photoelectron spectroscopy were used to check for the cleanness of the surface. No contamination was detectable. We have simulated the surface structures using the MIDAS and CHAOS [3] programs. The first program considers the crystal as a stack of planes defined by the interface normal. Atoms in the planes close to the interface are relaxed to positions of zero force. Atoms in the outer planes are held rigid, but the stack is allowed to move to permit the interface to dilate. The second program takes the relaxed surface structure from the first program and inserts the point defect. The crystal is then divided into two regions. In the inner region the ions are relaxed to positions of zero force. In the outer region, the response of the crystal to the defect charge is calculated using a dielectric continuum approximation. We use a central-force pair potential model for ceria (including electronic polarization by the shell model). The electrostatic terms are handled by the standard two-dimensional summations. A number of model potentials are available for CeO2; we use that of ref [4]. A more detailed discussion of methods and the model can be found in ref. [5]

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

238

3

Results and Discussion

3.1

CeO2(111)

Figure 1a shows an STM image of the CeO2(111) surface after preparation. The surface is (1×1) terminated and shows no reconstruction. This is in agreement with previous calculations, which found the (111) surface to be lowest in energy [6]. Oxygen is located in the top layer and the top layer is imaged by STM [2]. After annealing to about 1000 °C defects appear on the surface. Figure 1b shows a number of triangular defects. It has already been shown that arrangements of oxygen vacancies in pairs lead to a lower energy than two single oxygen vacancies on the surface. In order to study defect agglomeration in more detail we have calculated the energy of a number of defect arrangements consisting of three oxygen vacancies. Preliminary results suggest that indeed the triangle appears to be the energetically most favourable vacancy arrangement. The cluster consists of the triangle of vacancies together with the polarons (Ce3+ ions) on the layer below to balance the charge. This cluster is stabilised by the attraction of the polarons to the vacancies and also by the polarisation of the lattice arising from the defect dipole.

Figure 1: STM images of the CeO2(111) surface at different stages of annealing under UHV conditions, image size 6 nm × 4 nm, a) showing a (1×1) termination after annealing to 950°C, TSTM = 300°C, VBIAS = –2.5 V, IT = 0.2 nA, b) triangular defects after annealing to 1000 °C, TSTM = RT, VBIAS = –2.5 V, IT = 0.01 nA c) line defects after annealing to 1030 °C, TSTM = 500 °C, VBIAS = –2.5 V, IT = 0.08 nA, d) “black holes” after repeated annealing to 1030 °C, TSTM = 500 °C, VBIAS = –2 V, IT = 0.05 nA

239 Figure 1c shows line defects consisting of oxygen vacancies introduced after further annealing. They are running along the [1¯10], [0¯¯ 11] and [10¯1] directions. At this stage there is not enough space left on the surface to form additional triangular defects. In the vicinity of the triangular defects the stoichiometry is Ce2O3 rather than CeO2. Figure 1d shows the surface structure after repeatedly annealing the same crystal to 1030 °C. At this stage the atomic ordering of the surface is destroyed and a hexagonal pattern of “black holes” has formed. It is noteworthy that the distance between these “black holes” is about 1.6 nm, which is the minimum possible distance between two triangular defects on the (111) surface. The rough surface structure (apparent height of 0.25 nm for the “black holes”) indicates the involvement of deeper layers in the formation of this structure. The temperature above 1000 °C (at which the crystal was annealed) is the temperature range where catalytic converters irreversibly lose their efficiency. 3.2

CeO2(001)

Reconstruction is expected to occur on this surface because an unreconstructed surface would have a macroscopic dipole and therefore infinite surface energy. This is due to the stacking sequence of alternating planes with charge densities of equal magnitude but opposite sign in the [001] direction. Figure 2 shows an STM image of the CeO2(001) surface. The surface reconstructs as Ö2/2(3×2)R45°. The bright, vertical rows in figure 2a are separated by 1.15 nm which is equal to three times the lattice constant in the [110] direction. On these bright lines, atomically resolved features 0.77 nm apart (twice the lattice constant in the [¯ 110] direction) are visible. Because an oxygen terminated CeO2(001) surface is much lower in energy than a Ce-terminated one we conclude that these bright features in fig. 2a are related to oxygen.

[¯110] [110] 1 nm a) b) Figure 2: a) STM image of the CeO2(001) surface, TSTM = 300°C, VBIAS = –2.5 V, IT = 0.2 nA, the inset shows an area which the model in b) refers to; b) model of a Ö2/2(3×2)R45° reconstructed surface with an oxygen coverage of 33 % (top view and side view)

We have used the experimental observation as a starting point for atomistic modelling [5]. We have calculated the surface energy for the CeO2(001) surfaces. We varied the symmetry of the surface and the amount of oxygen in the top layer. The result of the calculations is that indeed a Ö2/2(3×2)R45° reconstruction is lowest in energy. This type (3×2) symmetry can be maintained over a rather wide range of oxygen coverage, namely between 50 % and 25 %.

240 Figure 2b shows a Ö2/2(3×2)R45° reconstructed surface with an oxygen coverage of 33 % which has a 0.7 eV (per oxygen atom removed) lower energy compared to reduction in the bulk. In this structure two thirds of the cerium ions in the second layer have been reduced from Ce4+ to Ce3+ in this structure. Based on the apparent height in the STM images and our calculations we conclude that the bright features shown in fig. 2a are oxygen atoms in the top layer of the surface. 3.3

CeO2(110)

The (110) surface of CeO2 has a higher surface energy than the (111) surface. After annealing to about 940 °C we found a (2×1) surface reconstruction [7]. After further annealing to about 1030 °C the surface forms (111) facets to reduce the surface energy. Ridges (>130 nm long) running along the [1¯¯ 10] direction are formed.

4

Conclusions

We have shown the structure of low-index surfaces of CeO2 with atomic resolution. STM at elevated temperatures and/or very low tunnelling currents are crucial to obtain this resolution. The reduction of CeO2 as it may happen in a catalytic converter was illustrated on the (111) surface of a single crystal of CeO2. Starting from a (1x1) terminated surface, agglomerations of oxygen vacancies in triangular arrangement could be observed after annealing the crystal under extremely oxygen lean conditions (UHV). Upon further annealing more oxygen vacancies are introduced forming lines in different, equivalent crystallographic directions. The CeO2 (001) surface shows a strong tendency for reconstruction. Calculations showed that it is able to maintain the same surface symmetry for oxygen coverage between 0.25 and 0.5.

5

References

[1]

G. Ertl, E. S. J. Lox, B. H. Engler, Environmental catalysis – mobile sources, in: G. Ertl, H. Knözinger, J. Weitkamp (Eds.) Handbook of Heterogeneous Catalysis, Vol. 4, VCH, Weinheim, 1997, p. 1559. H. Nörenberg and G. A. D. Briggs, Phys. Rev. Lett. 1997, 79, 4222–4225, Surf. Sci. 1999, 424, L352–L355. P.W. Tasker, Philos. Mag. A 39, 1979 119–136; D.M. Duffy and P.W. Tasker, A guide to CHAOS; a program for the calculation of point defect energies near interfaces in polar crystals 1983 Harwell Report AERE-R 11059. G.V. Lewis and C.R.A. Catlow, J. Phys. C 1985, 18, 1149–1161 H. Nörenberg and J. H. Harding, Surf. Sci. 2001, 477, 17–24. J. C. Conesa, Surf. Sci. 1995, 339, 337–352. H. Nörenberg and G. A. D. Briggs, Surf. Sci. 1999, 433–435, 127–130.

[2] [3]

[4] [5] [6] [7]

Nanostructured Ceria-Zirconia as an Oxygen Storage Component in Three-Way Catalytic Converters-Thermal Stability Boro Djuričić1 and Stephen Pickering2 1 2

Austrian Research Centers, A-2444 Seibersdorf, Austria Institute for Advanced Materials, 1755 ZG Petten, The Netherlands

1

Abstract

Nanostructured CeO2 and CeO2-ZrO2 solid solutions with Ce/Zr ratios of 0.75/0.25 and 0.50/0.50 were produced from aqueous solutions by precipitation/coprecipitation and hydrothermal crystallization in an autoclave at 180 °C for times up to 10 h. The thermal stability of these systems was studied by X-ray diffraction (XRD), differential thermal analysis (DTA) thermogravimetric analysis (TG) and transmission electron microscopy (TEM). Initial crystallite size was about 3nm for all powder compositions as measured by XRD peak broadening. Powders were calcined in air at 300 °C, 500 °C, 800 °C and 1000 °C for periods of up to 600 minutes. The crystallite sizes in pure Ceria , Ce75/Zr25 and Ce50/Zr50 powders were 137 nm, 23 nm and 37 nm respectively after 600 minutes at 1000 °C indicating a significantly higher thermal stability, i.e. greater resistance to coarsening, in nanostructured ceria-zirconia solid solutions. TEM images of the pure ceria powder showed clear evidence of sintering after calcination at 1000 °C. In contrast, TEM images of the ceria-zirconia solid solution powder calcined under similar conditions showed it to be still loosely agglomerated with significantly smaller crystallite size. This finding suggests that ceria-zirconia solid solution have higher effectiveness in service compared with pure ceria due to higher thermal stability.

2

Introduction

The three-way catalytic converter has probably made a greater contribution to reducing emissions from automobile exhausts than any other technology. Nevertheless, there is a continuing need to improve the performance of catalytic converters, particularly with respect to their ability to withstand high temperatures in order to allow them to be mounted closer to the engine and to extend their useful operating life. Cerium oxide plays an important role in catalytic converters as an oxygen storage medium. To maximise its effectiveness in this role, the cerium oxide should be distributed uniformly throughout the washcoat layer as very fine particles to ensure that a high specific surface area is exposed to the exhaust gas stream. Unfortunately, nanostructured cerium oxide powders suffer from pronounced coarsening on heating to temperatures typically encountered in exhaust systems. This paper demonstrates that crystal growth in ceria can be significantly reduced by means of a suitable additive.

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

242 Zirconia is a suitable additive because it can also contribute to oxygen storage through its ability to change stoichiometry. About 3–5 times higher oxygen storage capacity in ceriazirconia system than in ceria have been reported [1] which is attributed mainly to the enhancement of the bulk O2– ion mobility via vacancy diffusion in ceria-zirconia system. The goal was to produce material that could maintain a high specific surface area and good oxygen storage characteristics at high temperature. Compositions of ceria-zirconia mixed oxide with Ce/Zr ratios of 0.75/0.25 and 0.50/0.50 were produced for comparison with pure ceria.

3

Experimental

3.1

Preparation of powders

The powders were produced from aqueous solution by a novel two-stage precipitation process using hydrogen peroxide and ammonium hydroxide solutions. A 0.1 mol/l solution was made by dissolving zirconyl oxynitrate solution in de-ionised water. This solution remained slightly white and cloudy which may be attributed to the tendency of many Zr(IV) species to polymerize [2]. Addition of hydrogen peroxide resulted in a fully transparent solution with a pale yellow-green color. These changes in color and transparency indicate that a more soluble zirconium species was present after addition of hydrogen peroxide. The addition of cerium(III) nitrate solution to hydrogen peroxide is believed to result in the oxidation of cerium ions to cerium(IV) which is more easily hydrated and readily forms complexes to initiate the precipitation process. Subsequent addition of ammonium hydroxide completed the precipitation with the formation of an orange-brown gel accompanied by an intensive evolution of gas. The color of the precipitate changed on washing to orange-brown and finally to yellow-orange. The powder was autoclave treated at 180 °C to improve dispersability by removing hard agglomerates. For ceria this two-stage precipitation process [3] was earlier found to yield smaller crystallites (3nm) than by precipitation with ammonia (15nm). Samples of the powders were calcined in air at 300 °C, 500 °C , 800 °C and 1000 °C for periods of up to 600 minutes. 3.2

Analysis of powders

The ceria and ceria-zirconia powders were analysed by various techniques to determine phase composition, particle size and particle morphology. The precipitated powders were characterized by differential thermal analysis and thermogravimetry (DTA/TG, Netsch STA 409) in a dry air atmosphere using a heating rate of 5 °C/min to a temperature of 1400 °C. Samples of precipitated and calcined powders were prepared for examination by transmission electron microscopy (TEM, Philips EM 400) by ultrasonically dispersing the powder in ethanol and dipping a carbon-coated copper grid into the dispersion. After removing the grid from the suspension, the alcohol was allowed to evaporate. The phase composition was determined by x-ray powder diffraction (XRD, Philips PW 3710). A Voigt fitting routine was used to measure the broadening of the (220) and (311) peaks and the apparent crystallite size was calculated using the Scherrer equation: Dapp = Kl/Bcosq where K = 0.9, B = peak width, and l = 1.5406 Å for CuK= radiation. The line broadening due to the limits of instrument resolution was determined using a LaB6 standard.

243

4

Results

4.1

Differential thermal analysis/ thermal gravimetry

The precipitated pure ceria powder precursor showed a weight loss of 17 %–18% on heating to 500 °C and a well-defined exothermic peak at 260 °C–280 °C which probably corresponds to crystallization of CeO2. These data indicate that the precipitated species was CeO2.2H2O which has a theoretical weight loss of 17.3 %. 0.15

100

95

CeO 2

0.05

90

0.00

85

95

Ce0.75Zr0.25O 2

0.05

90 0.00

TG, % initial weight

0.10

0.10 DTA, mV/mg

100

85 -0.05 0

200

400

600

80 800 1000 1200 1400

-0.05 0

200

400

600

800 1000 1200 1400

Temperature, oC

Temperature, o C

Figure 1: DTA and TG analysis of ceria and ceria-zirconia powders after drying at 85 °C

The DTA and the TG data for the ceria-zirconia powders were broadly similar to that for pure ceria powders. In particular there was no sign of the ‘glow exotherm’ at about 400 °C typical of pure zirconia powders. The data therefore indicate that the two cation species precipitated as a solid solution rather than as separate phases. 4.2

X-ray diffraction

The initial crystallite size was about 3nm for all powder compositions as measured by XRD peak broadening The X-ray spectra for pure cerium oxide samples showed that cubic cerium oxide was the only phase present. The amount of peak broadening was the same for samples that were air-dried at 85 °C, autoclave treated at 180 °C for 4 h, or calcined in air at 300 °C, and corresponded to a calculated size of 3 nm. Calcination at higher temperatures resulted in narrower peaks corresponding to significantly larger crystallite sizes as shown in Table 1. After calcination at 1000 °C the powder was seen to be more coarsely agglomerated. Table 1: Average crystallite size [nm] after calcination in air 300 °C 60 min Ce100 3 Ce75Zr25 3 Ce50Zr50 3

500 °C 60 min 6 4 4

6 min 19 7 6

800 °C 60 min 25 8 8

600 min 40 10 10

6 min 72 18 13

1000 °C 60 min 79 25 14

600 min 137 37 23

244 The X-ray spectra of the ceria-zirconia powders were all generally similar to those for pure ceria, showing that cubic solid solutions of zirconia in ceria were present. The spectra for the Ce75Zr25 samples corresponded most closely to JPDS card 28-0271 for Ce0.75Zr0.25O2. The spectra for the Ce50Zr50 samples, however, corresponded slightly better to JPDS card 38-1439 for Ce0.6Zr0.4O2 (cubic) than to card 38-1439 for Ce0.5Zr0.5O2 (tetragonal). This shift is indicative of a change in lattice parameters of the solid solution. The valence change from Ce4+ with ionic radius 1.18 Å [4] to Ce3+with ionic radius 0.88 Å [5] causes phase separation [6]. The reduction of Ce4+to Ce3+, means that the system ZrO2-CeO2 transforms into the system ZrO2-CeO1.5 (Fig 2) with an entirely different phase composition; in fact the ternary system ZrO2-CeO2-CeO1.5 needs be considered. The presence of many defective tetragonal and cubic structures at temperatures below 1000 °C were reported for 50 mol% CeO2 in zirconia [10] and 65 mol% CeO2 in zirconia [11]. The cubic compound Ce2Zr2O7 with the pyrochlore (P) structure is stable in the composition area from about 44 mol to about 57 mol CeO1.5 [7,8,9]. Ce2Zr3O10 (B-phase) has been reported over a wide range of compositions in the CeO2-ZrO2 system at temperatures below about 800 °C [12,13]. The change in crystal symmetry is strongly effected by oxygen stoichiometry (i.e. by the atmosphere to which the material is exposed). Despite of the slow rate at which phases separate to the equilibrium state, these phase separations might be capable of playing a crucial role in the behaviour of this catalyst component, both with respect to the oxygen storage and catalytic activity. T °C

T°C

3000

3000

liquid

liquid

C2 + A

C2

A

C1

C

2000

C11 + P

T + C1

2000

T T

P

t+p T+P

1000

P + C2

1000

T+C

M M+P M

20

ZrO2

40 6060

Mol% Mol% Mol%

8080

CeO2

20

ZrO2

40

60

Mol%

80

CeO1.5

Figure 2: Phase diagram: ZrO2-CeO2 [7] and ZrO2-CeO1.5 [8,9]

If the reversible phase separation were to occur rapidly then, in principle, recrystallisation should lead to a “grain-refining” effect in which crystallites which have grown large at high temperatures would nucleate small crystallites of different phase composition, thereby maintaining a high specific surface area. This phenomenon could make it possible to design materials with a “self-repairing” oxygen storage capability.

245 The peak broadening of samples treated at up to 300 °C was the same in the ceria-zirconia samples as in the pure ceria samples, corresponding to a crystallite size of 3nm. After calcination at 800 °C and 1000 °C both the Ce75Zr25 and Ce50Zr50 ceria-zirconia samples showed a small but systematic shift in peak positions to smaller angles. The pure ceria samples showed no such shift. The peaks widths of the ceria-zirconia samples remained broader than for the pure ceria samples calcined under the same conditions, indicating significantly less crystallite growth than in pure ceria as shown in Table 1. The crystallite sizes in pure ceria, Ce75/Zr25 and Ce50/Zr50 powders were calculated to be 137 nm, 23 nm and 37 nm respectively after 600 mins at 1000 °C indicating a significant beneficial effect of the zirconia additive. 4.3

Electron microscopy

Figure 3 shows the morphology of weak agglomerates in ceria and ceria-zirconia powders after hydrothermal crystallization.

A

10nm

B B

10nm

Figure 3: Powders after hydrothermal crystallisation; (A) Pure ceria and (B) Ce50/Zr50

Figures 4 and Figure 5 show the morphology of calcined ceria and ceria-zirconia powders. The crystallite sizes visible in TEM images always appeared to be considerably smaller than those calculated from XRD peak broadening. To some extent this is to be expected as peak broadening values are very sensitive to the presence of a few large particles. Peak broadening values relate more to the mass median diameter whereas TEM image appearance relates more closely to the count median diameter so that considerable differences can be expected between the two values for the type of log-normal particle size distributions typically found in ceramic powders. Nevertheless, the TEM images of the pure ceria powder after calcination systematically showed a significantly larger crystallite size than in the Ce75/Zr25 and Ce50/Zr50 powders, thus confirming the same general trend shown by the calculated XRD values. There was clear evidence of sintering in the pure ceria powder after calcination at 1000 °C for 10h. In contrast, TEM images of the Ce75/Zr25 and Ce50/Zr50 powders calcined under similar conditions showed them to be still loosely agglomerated and the powder could be easily dispersed in alcohol to form a stable suspension.

246

A

B

10nm

C

10nm

10nm

Figure 4: Powders after calcination in air at 800 °C for 1 h; (A) Pure ceria, (B) Ce75/Zr25 and (C) Ce50/Zr50

50nm

A

B

50nm

Figure 5: Powders after calcination in air at 1000 °C for 10 h; (A) Pure ceria and (B) Ce50/Zr50

5

Conclusions

Ceria and ceria-zirconia mixed oxide powders were produced by a two stage precipitation process using hydrogen peroxide and ammonia. This precipitation process produced a significantly smaller crystallite size than is obtained by conventional precipitation by ammonium hydroxide alone. The ceria-zirconia compositions showed considerably less crystallite coarsening than pure ceria on calcination in air at temperatures up to 1000 °C. The ceria-zirconia compositions therefore have potential advantages over pure ceria for application as the oxygen storage component of 3-way catalytic converters.

247

6

References

[1] [2] [3] [4] [5] [6] [7] [8]

C. E. Hori et al, Applied Catalysis B: Environmental,1998, 16, 105–17. A.Clearfield, Rev. Pure and Appl. Chem. 1964, Vol 14, 91–108. B.Djuričić and S.Pickering, J. Eur. Ceram. Soc. 1999, Vol 19, 1925–1934. W. M. Goldschmidt, Berichte Deutsche Keram. Ges. 1963, 60, No 5. W. H. Zachariasen, Phys. Rev. 1948, 73, No 9. C. Leach and N. Khan, J. Mater. Sci. 1991, 26, [8], 2026–30. P. Duwez and F. Odell, J. Amer. Ceram. Soc.1950, 33, [9], 274–83. A. I. Leonov et all, Izvestiya Akademii Nauk USSR, Neorganicheskie Materialy, 1966, Vol. 2, No 1, 137–44. A. I. Leonov et al, Ogneupory,1966, No 3, 42 – 48. M. Yashima et al, J. Amer. Ceram Soc.1993, 76, [7], 1745–50. M. Yoshima et al, J. Amer. Ceram. Soc.1993, 76, [11], 2865–68. E. Tani et al, J. Amer. Ceram. Soc. 1993, 66, [7], 506–510. P. Duran et al, J. Mater. Sci. 1990, 25, 5001–06.

[9] [10] [11] [12] [13]

V Recycling

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

Recycling Technology for Metallic Substrates: a Closed Cycle Clemens Hensel Demet Deutsche Edelmetall Recycling AG & Co. KG, Alzenau (Germany)

1

Abstract

Emissions during the useage phase of vehicles are of increasing interest in environmental protection, and consequently, there is considerable interest in exhaust systems. The automotive exhaust system including the catalytic converter is, because of the precious metals in the catalyst, of particular interest for the recycling of automotive parts. The paper will describe the recycling technology of ceramic and metal catalyst substrates. The process will be analyzed in detail with the example of metal supports. As a result the complete life cycle and the recycling efficiency are presented.

2

Introduction

As early as the sixties, the first attempts were made to reduce the harmful substances in car exhaust gases by means of catalytic converters. As a result, converters were first employed in the USA in the seventies and were introduced to Europe in the eighties. While the emphasis at first was on the oxidation of carbon monoxide (CO) and hydrocarbons (HC), nitrogen oxides (NOx) are also reduced in today's three-way converters. The chemical reaction is achieved in a stoichiometric equilibrium with precious metal catalysts, platinum, rhodium and palladium. In the reduction processes, this chemical equilibrium ensures that precisely enough oxygen is released for oxidation. As its name indicates, the catalyst is not a collector (filter) but a converter of exhaust gases. A large reactive surface of the catalyst is necessary to convert the exhaust gases which are led through the exhaust train at high speed. This is ensured by the structure of the substrate and by the surface-enlarging wash coat. The surface of a conventional substrate is approx. 3 m² at 1 liter of volume and is increased to several thousand m2 by the wash coat. Efforts to increase catalyst efficiency tends to bring about further increases in specific geometrical surface area. The first catalysts for motor vehicles were manufactured from coated loose material (pellets). However, the relative movements of the parts led to wear and tear due to friction. Ceramic monoliths therefore displaced them almost completely. Increasingly stricter exhaust gas legislation made it necessary for the efficiency of converters to be continually increased. The metal converters – developed at the end of the eighties which are manufactured from thin brazed metal foil – offer some advantages in this respect. One advantage is the increased specific geometric surface area compared to that of the ceramic substrate. As efficiency is dependent first and foremost on the area of the geometric surface (where the parameters are otherwise

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

252 identical), the metal catalyst can be reduced in volume while remaining equally effective. This means that the quantity of precious metals required is likewise reduced. In addition, the employment of especially thin metal foil reduces the thermal mass of the substrate so that it reaches its working temperature far more quickly. Metal substrates are therefore employed to an increasing degree. The recycling procedure used for ceramic substrates is only partly suitable for metal substrates. For this reason, a new process, described in the following, was developed.

2

Differences in the Structure of Catalysts

2.1

Structure of a ceramic monolith catalyst

The ceramic substrate has a honeycomb structure. The exhaust gas flows through the coated channels and contacts the precious metals. The monolith is surrounded by an insulating element (ceramic fibre mat) whose purpose is not just to insulate the catalyst thermally but also to secure it mechanically. As with a silencer, the steel case is the outer “packaging”.

Casing Ceramic fibre mat Ceramic substrate

Figure 1: Ceramic catalyst incl. ceramic fibre mat and casing [1]

2.2

Structure of a Catalytic Converter with Metal Substrate

The metal substrate is made of flat and corrugated metal foils. These are stacked in layers and are wound into a cylindrical or elliptical body. This body is then pressed into a metal casing and is joined by a high-temperature brazing process. The alternating flat and corrugated layers produce a channeled structure which can then be coated in order to increase its surface area – just like the ceramic substrate. Due to the integrated casing, no further measures are required to

253 secure the substrate. If need be, additional heat insulating sheets can be applied for purposes of thermal shielding. These two types of monoliths differ in form, size, weight and precious metal loadings as a result of different designs for the exploitation of specific advantages of the two techniques.These differences in structure bring about different requirements with regard to the recycling of these components.

Figure 2: Structure of a catalyst with metal substrate

3

Recycling Procedures

3.1

Requirements for Recycling

As the precious metals are applied only to certain parts of the converter (in the wash coat), the first step to be taken in the technical process of recycling is a relatively simple one: mechanically seperate the washcoat and precious metals from the substrate. The next step is then to homogenize the precious metal fraction. This is crucial for determining the value of recycling batches, i.e. statistically representative samples to acquire an analytical sample which is analyzed for its precious metal content. The results of this analysis form the basis for the settlement of accounts with the supplier. As described earlier, the ceramic monolith is only layed into the casing, surrounded by the ceramic fibre mat. Since the mechanical properties of the monolith are extremely different

254 from those of the outer steel mantle, the monolith can be seperated relatively easily. The ceramic material obtained by breaking down the catalyst can be easily ground to a homogeneous powder which is highly suitable for sampling. Initially, the first recycling procedures used for metal substrates had to do without this sort of separation. Instead, the entire catalyst was processed in a melting procedure. Apart from the high costs of melting, this has a further disadvantage in that instead of concentrating the precious metals, the precious metals are diluted throughout the whole melt (including the case). Such a process is less effective; in addition, it is difficult to obtain a representative sample from the homogeneous material. In the procedure described below, these disadvantages can be eliminated. 3.2

Newly Developed Separation Process for the Recycling of Metal Catalysts

3.2.1 General Considerations, Patent Special emphasis was placed in development to find a useful sampling procedure for metal catalysts. Melting a few complete components is costly and not sufficiently representative. An obvious approach was to process all of the components into a “homogeneous mass” – and to do this by mechanical rather than pyrometallurgical means. According to this approach, a sample could then be taken from this “homogeneous mass”. However, the first attempts ended with an alarming result: the crushing produced a fine dust rich in precious metals, and this dust was in danger of being lost unless a costly dust collection process was employed. The wash coat had become partly seperated from the substrate! This counterproductive effect would have meant the end of these attempts had it not been for a further attempt to deliberately turn the negative effect to good use. Thus the idea came about to no longer produce a uniform “homogeneous mass” but to remove the entire wash coat with its precious metal content from the remaining material by means of an explicit separation procedure. This was subsequently demonstrated in testing. This procedure – since patented throughout the world (US Patent 005279464A) – enables not just the wash coat with its precious metals to be recovered, but also the steel fraction. 3.2.2 Stages of the Process This procedure is based on the empirically acquired experience that by observing certain process parameters, the wash coat can be seperated from the substrate by mechanical means. If this produces a mixture of particles differing from each other in form, size and weight, then these can be seperated from each other. The end result is that different but homogeneous fractions of high purity are recovered. a)

Mechanical Reduction

In the first stage, the catalysts are reduced to a defined particle size by means of a shredder. Through the high degree of deformation of the substrate foil, caused by the applied mechanical energy, a large part of the wash coat already becomes detached and can be collected by means of air separation.

255 b)

Magnetic Separation

In most cases (depending on the materials used), the remaining metallic fraction can be magnetically seperated. The casing material – often austenitic – is not processed any further. On the other hand, a residue of wash coat still adheres to the substrate foil. c)

Mechanical Aftertreatment

In a further processing stage, the residue of precious metals on the substrate foil is seperated from the steel parts by means of a mechanical separation unit. A hammer mill is employed to “beat” the remaining particles of the wash coat away from the foil; here too, the foil and ceramic particles can be seperated from each other by means of air separation. Metal Substrate Converters

Shredder Washcoat

Metal

Magnetic Separation

Substrate Foil

Stainless Steel

Mechanical Separation Washcoat

Blending & Sampling

Figure 3: Separation process

256 All fractions of the wash coat are brought together for homogenization and sampling. The substrate foil and casing material can be reutilized in steel smelting. Fractions after Separation After separation, the fractions of the wash coat containing the precious metals, the substrate foil, and the casing material are available individually for further processing as shown in Figure 4.

Mantle

Foil Washcoat Figure 4: Fractions after separation

3.2.2 Sampling All subsequent steps for the recovery of the precious metals now focus exclusively on the seperated wash coat. As already mentioned, sampling is of major importance in the recycling of precious metals so that the content of the material can accurately be determined by means of an analysis. For ceramic materials, multiple-stage homogenization and separation processes are generally employed. The first step here is a synchronous milling and sieving procedure. Steel residue can thus be seperated and larger parts of the wash coat can be reduced to a uniform maximum size. The sample produced is also subjected to a drying process to fix its moisture content. As a ceramic material, the wash coat always contains a slight amount of moisture; this can amount to over 3 % of the weight of the catalysts if these are stored under unfavorable conditions.

257 Washcoat

Blending

SampleSplitting

Moisture Determination

Fine Milling

Screening

Basis for settlement of account

Wash-Coat for Refining zum Refining

Fines

Coarse

SampleSplitting

SampleSplitting

Sample

Sample

Figure 5: Sampling

4

Comparison of Catalyst Types in Recycling

4.1

Advantages and Disadvantages in Summary

The decanning of ceramic catalysts can be effected with far simpler means when compared to the shredding process employed for metal catalysts. The result in both cases is a ceramic material, which means that no costly technical procedure is required for the sampling. The only constituent requiring further processing in the case

258 of metal catalysts is the wash coat with precious metals, with ceramic catalysts the substrate is also involved. This has a substantial effect on costs. The subsequent recovery of the precious metals is therefore less costly because only a far smaller quantity of material needs to be refined. In addition, the substrate fraction of the metal catalyst can be used for steel smelting while the ceramic substrate is reduced to slag. The greater expenditure on pretreatment is rewarded by lower costs in further processing. Neither of the procedures can be credited with advantages of economy. From the ecological point of view, however, it is advantageous when only a small fraction of the material which cannot be returned to the materials cycle - this is the case with metal catalysts. To determine the financial expense involved in the recycling of metal and ceramic catalysts, two different systems employed in a modern 2 liter engine are examined. The two systems are equally effective and have similar geometric surface areas. The characteristics of the metal substrate are such that the metal catalyst system has less volume. Comparison is made by setting the values to 100 % for ceramic catalysts as the basis. The differences then are shown by calculating the ratio, giving the percentage of Metalits. Table 1: Geometrical data of catalysts Size Cell density and foil thickness Catalyst volume Geometric surface area

Metalit™ Ø118×174 mm 400 cpsi / 50 µm 1.9 l 6.1 m²

Ceramic substrate Ø143×152 mm 400 cpsi / 6.5 mil 2.44 l 6.2 m²

With the same volume-specific quantity of wash coat, this likewise means a smaller quantity of precious metals. Table 2: Data of catalytic coating Washcoat mass Precious metal loading Value of precious metals

Metalit™ 150 g/l – 332.5 g 50 g/ft³, Pt/Rh 5:1 – 3.39 g 77.9 %

Ceramic substrate 150 g/l – 427.0 g 50 g/ft³, Pt/Rh 5:1 – 4.357 g 100 %

There are considerable differences in the methods employed in the recycling process, especially in the way the materials are broken down. With the ceramic catalyst, the substrate including the coating is ground into a homogeneous powder. With the metal catalyst, the metallic substances and the wash coat powder have been seperated in the mechanical shredding and separation processes. It is somewhat more expensive to break down the metal catalyst with the separation of the coating powder but it results in a higher concentration of precious metals. Table 3: Downsizing of substrates Costs of shredder process Costs for dismantling and milling

Metalit™ 568 %

Ceramic substrate 100 %

259 This means less powder is subsequently processed in order to recover the precious metals, which leads to substantial cost savings in this stage of the process. Table 4: Precious metal recycling costs Value of precious metal content Metal fees Processing costs powder Value of recycled PM Overall loss

Metalit™ 77.9 % 77.9 % 21.3 %

Ceramic substrate 100 % 100 % 100 %

77.9 % 78.9 %

100 % 100 %

A further advantage over processes formerly employed is the possibility of recovering the materials of the converter seperated into fractions. Table 5: Value of steel recycled Metalit™ Mantle + matrix 292 % Mantle

Ceramic substrate 100 %

As a result of these cost savings, the absolute costs for the recycling of catalyst systems are somewhat less for the metal catalyst and the value of the material recovered in relation to the costs of recycling is comparable for both types of catalyst substrates. Table 6: Overall recycling costs Costs / substrate Recovered value / costs

5

Metalit™ 79.5 % 2.58 DM/DM

Ceramic substrate 100 % 2.58 DM/DM

Summary

It has been demonstrated that compared to the traditional melt process, a more efficient alternative is mechanical pretreatment. · Sampling The original development objective of producing a material capable of being sampled was attained despite the change in research plan. The wash coat powder recovered is ideal for sampling as it is a conventional ceramic material. · Recycling of steel In contrast to the fusion procedure, mechanical separation allows the steel fractions to also be reutilized. This is advantageous not only economically but also – and especially – ecologically. · Costs From today's point of view, the recycling of both designs is to be categorized as equally expensive even if the respective cost-intensive factors are in different stages of the processing.

260 In conclusion, it can be stated that with the mechanical pretreatment of catalytic converters having metal substrates, a further step has been taken that contributes both economically and ecologically towards the improved recycling of automobiles.

VI Miscellaneous

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

Hot-Corrosion of Metal and Ceramic Honeycombs by Alkaline Metals for NOx Adsorption Mikio Yamanaka Nippon Steel Technoresearch, Futtsu, Chiba

Yuichi Okazaki Nippon Steel, Toukai, Aichi, Japan

1

Introduction

Alkali metals and alkaline-earth elements are used for the NOx adsorption catalyst of lean burn engines. The NOx storage efficiency is the higher with the higher basicity of these metals (1)(2). But highly basic metals cause a hot-corrosion especially on ceramic honeycombs. In the previous paper, the authors dealt with the hot-corrosion caused by Li, K and Cs sulfates and it was turned out that a serious hot-corrosion takes place on cordierite honeycombs with g-Al2O3 coating containing K2SO4 above 1173 K (900 °C) (3). These corroded honeycombs lose their mechanical strength. In addition, reactive g-Al2O3 is easily converted to a type in the condition where the hot-corrosion takes place. On the other hand, no damage was observed for 20Cr-5Al metal foils under the same condition, but their weight increase was linear with the oxidation time. In case of real catalyst, alkali metals and alkaline-earth elements are added to g-Al2O3 as carbonate compounds for NOx storage. Thus in this paper, the effect of K2CO3 is presented in comparison with those by alkali metal sulfates, and their mechanism will be discussed focusing on the difference of hot corrosion behavior by the kind of alkaline elements.

2

Experimental Details

Test conditions were nearly the same with those of the previous experiment, except that K2CO3 was used as the additive to g-Al2O3 and that the slurry was applied only on the outer surface of honeycombs, not inner surface of caves this time. 0.2 mol of K2CO3 was added to 100g of g-Al2O3 with distilled water. The slurry of their mixture was applied on the surface of cordierite honeycombs and 20Cr-5Al foils. The size of each specimens was 20 × 20 × 2.8 mm for cordierite honeycombs and 20 × 20 mm × 50 mm for metal foils. After dried, each test piece was heat treated in wet air (DP = 323 K) at 873 K, 1023 K and 1173 K (600 °C, 750 °C and 900 °C) for 360 ks (100 hrs). These samples were weighed and subjected to folding test for metal foils and to 3 points bending test for ceramic honeycombs as shown in figure 1. The mechanical strength of these ceramic samples s was calculated by the equation as follows:

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

264 s = 3PL/2WT2 where P is the maximum load before failure, L the distance between two bars beneath the test piece, W and T the width and thickness of the sample, respectively. Both the ceramic and metal structure were observed by an optical microscopy and were analyzed by EPMA .

Figure 1: Schematic illustration of 3 points bending test for ceramic samples

3

Experimental Results

No nodule or abnormal oxidation was observed for metal foils heat treated even at 1173 K (900 °C). Their weight increases are listed in Table 1 with data given by the previous experiment using alkali metal sulfates. No cracking took place in the folding test of any metal foil after heat treatments, but the surface film of a foil heat treated at 1173 K (900 °C) was not uniform, as shown in figure 2. EPMA mapping of the metal foil is shown in figure 3. A few percent of K (potassium) is detected in the surface film but not in the metal matrix.

Figure 2: Cross section of metal foil folded face to face after heat treatment at 1173K

265 Table 1: Weight increase(×10–2kg/m2) of metal foil by heat treatment for 3.6 × 103 Ks (100 h) Additive to g-Al2O3 Li2SO4 Cs2SO4 K2SO4 K2CO3

873 K ~0 ~0 ~0 ~0

1023 K 0.01 0.02 0.03 0.03

1173 K 0.09 0.15 0.18 0.56

Figure 3: EPMA mapping of metal foil after heat treatment at 1173 K with the coating containing K2CO3

266 On the other hand, ceramic honeycombs heat treated above 1023 K (750 °C) with the coating containing K2CO3 changed their color from yellow to a little reddish, and some of them showed cracking, as shown in figure 4. Mechanical strengths of ceramic honeycombs with no crack were plotted in figure 5 with data given by the previous experiment. Ceramic honeycombs heat treated above 1023 K (750 °C) lost their mechanical strength. The ceramic structure of a honeycomb heat treated at 1173 K is shown in figure 6. A distortion and microcracks can be seen in the honeycomb lattice. Figure 7 shows the EPMA mapping of a ceramic honeycomb after heat treatment at 1173 K (900 °C). K (potassium) was observed to invade from the outer surface of the honeycomb sample to the whole wall thickness

Figure 4: Ceramic honeycombs heat treated at 1023 K with the coating containing K2CO3

Figure 5: Mechanical strength of ceramic honeycombs before and after heat treatment

267

Figure 6: Structure of ceramic honeycomb after heat treatment at 1173 K with the coating containing K2CO3

4

Discussion

In the previous paper, the authors reported that 20Cr-5Al metal foils for metal supports have a hot-corrosion resistance to K sulfate at least up to 1173 K (900 °C), but that cordierite honeycombs are attacked by K2O of K2SO4 and lose its mechanical strength at 1173 K. But in case of real catalyst, alkali metals and alkaline-earth elements are added to g-Al2O3 as carbonate compounds for NOx storage. Thus in this experiment, effects of K2CO3 both on metal foils and on ceramic honeycombs were examined. Even to K2CO3, metal foils exhibited a hot-corrosion resistance up to 1173 K (900 °C) as mentioned above. But the weight increase at 1173 K was three times of that by K2SO4 and five times of that with no coating at the same temperature given by another experiment. The value at 1173 K (900 °C) is about 50 % of the weight increase when the whole Al in the foil of 50 m is consumed for the surface oxide film. But this time, it contains the weight increase by K2O which attacked the Al2O3 film, because a few percent of K was detected in the surface film of the metal foil as shown in figure 3. Its surface film is not uniform and thick compared with one given by the mere oxidation in air or exhaust gas from an engine without coating. Nevertheless, the surface film prevented K from entering into the metal matrix. In case of cordierite honeycombs, K2CO3 attacked the ceramic matrix at the lower temperature of 1023 K (750 °C) than K2SO4, though the coating was applied only on one side of the honeycomb wall in this case. Cordierite 2(MgO)2(Al2O3)5(SiO2) is based on silica network and is presumed to be somewhat acidic compound. K2CO3 is more basic than K2SO4 , thus cordierite is more sensitive to K2CO3 than to K2SO4. In the previous experiment, cordierite honeycombs heat treated with the coating containing Li2SO4 or Cs2SO4 did not lose their mechanical strength, but Cs was detected to invade into the surface layer of ceramic matrix by EPMA. Thus the order of alkali attack activity to cordierite can be regarded as follows: Li2SO4
268 of three metals and oxygen were listed with their ionic radii (5). According to these electronegativities, Cs oxide is most basic and should be most active to cordierite honeycombs. But the Cs ion is large enough to suppress its diffusion into cordierite matrix. Thus Cs attacking is limited to the surface layer of cordierite.

Figure 7: EPMA mapping of ceramic honeycomb after heat treatment at 1173 K with the coating containing K2CO3

269 Table 2: Electroneagativities and ionic radii of three metals and oxygen element Li K Cs O

Electroneagativity 0.97 0.91 0.86 3.50

Ionic radius/A 0.90 (6) 1.52 (6) 1.81 (6) 1.21 (2)

Numbers in parentheses after ionic radii are coordination numbers.

Those phenomena can be explained by another aspect. In table 3, the lowest eutectic temperatures of binary metal oxide systems (6) are shown. In case of metal foils, Al2O3 surface film was attacked by K2O at 1173 K (900 °C), but the melting temperature of the yield is at lowest 1723 K (1450 °C). Thus the surface film kept protective to some extent against hotcorrosion, though the surface film grows faster and thicker than one by the mere oxidation with no coating. On the other hand, the lowest eutectic temperature of K2O-SiO2 system is as low as 1053 K (780 °C), thus the silica network of cordierite was easily attack by K2O and lost its strength. But Li2O-SiO2 system has rather high eutectic temperature, so cordierite honeycombs were hardly attacked by Li2O. Table 3: Lowest eutectic temperatures of binary metal oxide systems Binary oxide system K2O-Al2O3 Li2O-SiO2 K2O-SiO2 Cs2O-SiO2

5

Lowest eutectic temperature 1723 K (1450 °C) 1300 K (1027 °C) 1053 K (780 °C) 1143 K (870 °C)

Conclusion

K2CO3 exhibited a higher activity both on the Al2O3 surface film of metal foils and on cordierite honeycombs. But in case of metal foils, the surface film kept protective to some extent and prevented K from entering into metal matrix. On the other hand, a serious hotcorrosion took place on cordierite honeycomb at the lower temperature of 1023 K (750 °C) than by K2SO4, and the ceramic honeycomb lost its mechanical strength. The order of alkali attack activity to cordierite can be regarded as follows:Li2SO4
270

6

References

[1] [2] [3]

S. Matsumoto, ‘Shokubai’, 1997, 39, 210. T. Kobayashi , T. Yamada and K. Kayano, SAE, 1997, 970745. M. Yamanaka and Y. Okazaki, Metal-Supported Automotive Catalytic Converter (Ed.:H. Bode), Werkstoff-Informationsgesellshaft mbH, 1997, 127. A. L Allred and E. G. Rochow, J. Inorg. Nuclear Chem., 1958, 5, 264. F. A. Cotton and G. Wilkinson, Advanced Inorganic Chemistry, 5th ed., John Willy & Sons, p.1387. Phase Diagrams for Ceramists, American Ceramic Society, 1987, Fig. 1350 and 1979, Fig. 365.

[4] [5] [6]

The Effect of Trace Amounts of Mg in FeCrAl Alloys on the Microstructure of the Protective Alumina Surface Scales Pudji Untoro, Mohammad Dani National Nuclear Energy Agency, Kawasan PUSPIPTEK, Serpong 15314, INDONESIA

Hans Joachim Klaar, Joachim Mayer Gemeinschaftslabor für Elektronenmikroskopie, Rheinisch Wesfälische Technische Hochschule Aachen D52074-Aachen, GERMANY

Dmitry Naumenko, Jui-Chao Kuo and Willem Joe Quadakkers Institut für Werkstoffe und Verfahren der Energietechnik (IWV-2), Forschungszentrum Jülich, D-52425 Jülich, GERMANY

1

Abstract

Minor alloying additions of reactive elements are of crucial importance for the protective properties of alumina scales on FeCrAl alloys used for manufacturing of modern car catalyst carriers. The effects of the reactive elements, such as Y, Zr, Hf and/or lanthanides etc on the high temperature oxidation behaviour of metallic materials have been studied in literature since at least 30 years. However, the commercial FeCrAl alloys often contain a number of additional, intentionally added or unavoidable impurities, which might significantly affect the oxidation behaviour. In the present study a range of analytical methods were applied to study the effect of magnesium impurity (80 ppm mass) in a commercial FeCrAl alloy . The isothermal oxidation kinetics were determined at 1200 °C in air using Thermo Gravimetric Analysis (TGA). The results of the TGA measurements have been correlated with SIMS (Secondary Ion Mass Spectrometry) and microstructural investigations with TEM/EDX (Transmission Electron Microscopy/Energy Dispersive X-ray-analysis) and Electron Energy Loss Spectroscopy (EELS). The scale cross-sections for the TEM investigations were manufactured with the Focused Ion Beam technique (FIB) to reach large area transparent samples which can be analysed from the top of the oxide layer to the alloy matrix. Comparison of the data from different analyses techniques reveals that the alumina scale grows primarily by inward oxygen grain boundary diffusion. In addition to this main oxidation mechanism, outward transport of Mg, in combination with Mn, occurs via the scale grain boundaries. Hence, magnesium becomes enriched in the outer part of the oxide scale forming a layer of a MgAl2O4 spinel. An important effect associated with the spinel formation is the development of porosity in the outer part of the scale.

2

Introduction

FeCrAl based alloys are well known for their superior oxidation resistance up to very high temperatures and they are therefore used as construction materials in a large variety of indus-

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KGaA ISBN: 3-527-30491-6

272 trial applications [1–3]. Since recently, FeCrAl alloys are being used as substrate material for modern car catalytic converters [4]. The long term stability of the converters relies significantly on the thickness of the oxide, its scale growth rate, microstructure, and adherence. Several studies have examined the role of reactive elements (RE) and other alloying additions on the oxidation behaviour of FeCrAl alloys [5-7]. It had been clearly established that RE additions improve the alumina-scale adhesion and the stability of the alumina based scale on FeCrAl alloys [8–13]. However, commercial FeCrAl alloys often contain a number of additional, intentionally added or unavoidable impurities, which might significantly affect the oxidation behaviour. Recent studies have shown that for instance traces of titanium and phosphorous can dramatically affect the oxide scale properties [14, 15]. The aim of the present study was to investigate the effect of magnesium (80 ppm mass), an impurity which is sometimes encountered in metallic high temperature materials, on the oxidation kinetics of FeCrAl alloys at 1200 °C in air. For obtaining detailed information on the mechanism of Mg incorporation in the alumina surface scale a number of conventional test and analysis techniques were combined with TEM studies of oxide scale cross sections prepared by a focussed ion beam (FIB) technique.

3

Experimental Procedure

The composition of the FeCrAl alloy used in this study is shown in Table 1. Specimens 20 × 10 × 2 mm in size were abraded on silicon-carbide papers, then mechanically polished with 1 µm diamond paste and degreased ultrasonically in a detergent prior to oxidation. The isothermal oxidation kinetics were measured with a Setaram thermobalance in synthetic air at 1200 °C for 100 h. After oxidation, the specimens were examined by optical metallography and transmission electron microscopy (TEM) coupled with energy-dispersive X-ray analysis (EDX), x-ray diffraction (XRD) and secondary ion mass spectrometry (SIMS). The methods used to quantify the SIMS depth profiles were published elsewhere [16]. To prevent scale loss during TEM-preparation, specimens were sputtered with a thin gold layer and tungsten before cross sections preparation by a focussed ion beam technique (FIB). Phase analysis in the TEM cross sections was carried out by EDX and electron energy loss spectroscopy (EELS). Table 1: Chemical composition of the studied FeCrAl alloy in mass % or mass-ppm Fe (%) Base

Cr (%) 20,10

Y (ppm) Ti (ppm) 600 35

4

Al (%) 5,68

Ni (ppm) 1550

Mg (ppm) 80

Mn (ppm) Mo (ppm) Si (ppm) 2040 60 900

Zr (ppm) 33

Hf (ppm) 340

S (ppm) 2

N (ppm) 2.1

C (ppm) 23

Results and Discussion

The mass change as function of exposure time during isothermal oxidation at 1200 °C is shown in Figure 1.

273 1,0

mass change / mg.cm

-2

0,8 0,6 0,4 Experimental data

0,2

Fit: ,m=0,18.t0,36

0,0 0

20

40

60

80

100

time / h Figure 1: Isothermal oxidation of FeCrAl alloy at 1200 °C in air

a)

b)

c) Figure 2: Etched cross sections of FeCrAl alloys (a) as received condition; (b) after 100h exposure; (c) after 1000h exposure

274 From the result it can be seen, that the oxidation kinetics obey a sub-parabolic timedependence, frequently observed for alumina scales growing by oxygen grain boundary diffusion [10, 17]. The microstructure of the alloy was characterized by optical microscopy and TEM. In the as-received state, the grains are deformed along the rolling direction (Figure 2-a) whereas a 100 h or 1000 h exposure at 1200 °C results in pronounced recrystallization and grain growth (Figure 2-b and 2-c). SEM/EDX analyses of the oxide scale formed after different exposure times revealed the scale to mainly consist of alumina with small impurities of a Hf/Zr-rich phase (Figure 3). This could be confirmed by the TEM cross-section analysis (Figure 4). In the outer scale a slight enrichment of Mn and especially Mg could be detected. Depth profiling of a specimen oxidized for 100 hours using MCs+-SIMS confirmed the presence of both elements in the outer scale (Figure 5). XRD analysis of the same specimen revealed the scale to consist mainly of a-Al2O3 and a few weak lines were found indicating presence of a spinel type oxide (Figure 6). Cross section TEM observations (Figure 4) were then used to study scale microstructure and to more precisely locate the magnesium in the oxide layer. The scale appeared to consist of columnar grains near the scale alloy interface and equiaxed grains near the oxide surface.

Figure 3: EDX element mapping of cross sectioned scale formed on FeCrAl alloy after 1000 h cyclic oxidation at 1200 °C in air

275 The TEM investigations not only allowed to locate grains of spinel (MgAl2O4) at the top of the oxide layer, but also magnesium at the grain boundaries of the alumina, whereas magnesium could not be identified within the alumina grains. However, the highest magnesium signal was detected near the scale/gas interface and the results of EDX analyses correspond to a MgAl2O4 phase. This is confirmed by EELS data in Figure 7. According to the various analyses, it turned out that the Mg enrichment in the outer scale occurs by diffusion of Mg from the bulk alloy into the oxide scale. An important observation was, that substantial porosity was present in the outer part of the scale (Figure 4). A posible explanation for this phenomenon could be the transformation of transient metastable alumina into stable a-Al2O3, a process which is known to be connected with a strong volume decrease [14]. This mechanism seems, however, not very likely, because at 1200 °C metastable alumina formation is generally not observed [14,18]. Therefore, a better explanation for the formation of the large voids in the outer scale seems, that it is related to the spinel formation, induced by Mn and especially Mg, diffusing from the bulk alloy through the inner a-Al2O3 layer to the scale surface. A reaction of the mentioned impurities with the pre-existing Al2O3 according to Mg(Mn) + Al2O3 + ½ O2 Û (Mg,Mn) Al2O4 is likely to result in a volume change induced by the change in lattice structure (hexagonal to cubic), the differences in atomic volume of alumina and spinel and/or the stresses induced by the phase transformation.

Figure 4: TEM-image from cross section and eDX analysis after 100 h isothermal oxidation at 1200 °C in air

276

Figure 5: MCs+-SIMS profile of oxide scale (outer part) on FeCrAl alloy formed during 100 h isothermal oxidation at 1200 °C in air

Figure 6: XRD spectrum of FeCrAl after 100 h isothermal oxidation at 1200 °C in air

277

Figure 7: EELS data of the outer part of scale of (Mg, Al)-rich oxide

4

Conclusions

· Magnesium is easily incorporated in alumina scales on FeCrAl alloys although it is present in the alloy in very low concentrations (~ 80 ppm). · Magnesium, together with Mn, diffuses through the inner a-Al2O3 layer in direction of the scale/gas interface. Reaction of the magnesium with alumina leads to spinel formation accompanied by initiation of voids in the outer part of the scale. · The amount of magnesium impurity should therefore be carefully considered in FeCrAl alloy manufacturing, because it can substantially affect the properties of the protective aumina layer.

5

Acknowledgements

The authors are very grateful to the ministry of education and research (BMBF) and National Nuclear Energy Agency (BATAN) for financial support under the Project IDN 99/004 and Dr. A. Schertel (FEI Deutschland GmbH) for FIB-Preparation.

6

References

[1]

W.J. Quadakkers, K. Schmidt, H. Grübmeier, E. Wallura, Mat. at High Temperature, 1992, 10. 23–32 H. Nickel, W.J. Quadakkers, Proc. Conf. On Heat Resistant Materials, Fontana, Wisconsin, USA ,1991, p. 87 F. Starr, Mat. at High Temperature 1995, 13, 185–194 J.M. Herbelin, M. Mantel, J.Y. Cogne, in Proc. Int. Conf. On Metal Supported Automotive Catalytic Converters (Ed.: H. Bode), Frankfurt , 1997, 79–92

[2] [3] [4]

278 [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15]

[16]

[17] [18]

W.J. Quadakkers, A. Elschner, H. Holzbrecher, K. Schmidt, W. Speier, H. Nickel, Michrochimica Acta, 1992, 107, 197–206 J.P. Pfeifer, H. Holzbrecher, W.J. Quadakkers, W. Speier, Fresenius Journal of Analytical Chemistry, 1993, 346, 186–191 J. Jedlinski, G. Borchardt, Oxidation of Metals, 1991, 36, 317–337 L.B. Pfeil, UK Patent no. 459848, 1937 D.P. White, J. Stringer, Phil. Trans. Roy. Soc., 1980, A295, 309–329 W.J. Quadakkers, Werkstoffe und Korrosion, 1990, 41, 659–668 R.A. Versaci, D. Clemens, W.J. Quadakkers, R. Hussey, Solid State Ionics, 1993, 59, 235–242 B.A. Pint, A.J. Garrat-Reed, L.W. Hobbs, Mat. at High Temp., 1995, 13, 3–16 W.J. Quadakkers, H. Holzbrecher, K.G. Briefs, H. Beske, The Role of Active Elements in the Oxidation Behaviour of High Temperature Metals & Alloys, 1989, p. 155 D. Naumenko, Dissertation, Technical University, Aachen, FRG, 2001, in press D. Naumenko, W. J. Quadakkers, V. Guttmann, P. Beavan, H. Al-Badairy, G. J. Tatlock, R. Newton, J. R. Nicholls,G. Strehl, G. Borchardt, J. Le Coze, B. Jönsson, A. Westerlund, EFC-Workshop “Life Time Modelling of High Temperature Corrosion Processes”, Frankfurt, D, 22–23 February, 2001, Proceedings in EFC Monograph, in press W.J. Quadakkers, H.Viefhaus, EFC-Workshop "Methods and Testing in High Temperature Corrosion, Frankfurt, FRG, 20–21 Jan. 1994, Proceedings in „EFC Publications“, nr 14, Guidelines for Methods of Testing and Research in High Temperature Corrosion, Edt. H.J. Grabke, The Institute of Materials, London, 1995. p. 189–217 Z. Liu, W. Gao, Y. He, Oxidation of Metals, 2000, 53, 341–350 W.J. Quadakkers, L. Singheiser, in Int. Conf. On High Temperature Corrosion, Les Embiez, France, 2000, proceedings in press

Author Index* A

K

Andoh, A. 83

Kikuchi, A. 144 Klaar, H. 271 Kolarik, V. 117 Kolb-Telieps, A. 49 Kolb-Telieps, A. 117 Konya, S. 144 Kubsh, J. 217 Kuo, J. 271

B Bode, H. 134 Borchardt, G. 106 Brück, R. 19 C Cedergren, M. 126 Chang, C. 69 Chen, L. 69 D Dani, M. 271 del Mar Juez-Lorenzo, M. 117 Djuricic, B. 241 F Fietzek, H. 117 Fukuda, K. 59 Furukimi, O. 59 G Göransson, K. 126 Grünert, W. 229 Guist, C. 134 H Harding, J. 237 Hattendorf, H. 117 Hensel, C. 251 Hojda, R. 117 Hoshi, T. 59 J Jentys, A. 223 Jha, B. 69 *

L Lang, M. 186 Lercher, J. 223 Lylykangas, R. 152 M Makino, M. 173 Mayer, J. 271 Merry, D. 202 N Naumenko, D. 49 Naumenko, D. 93 Naumenko, D. 271 Newton, R. 49 Nicholls, J. , 31 Nicholls, J. 93 Nörenberg, H. 237 O Ohara, E. 173 Okazaki, Y. 263 P

Quadakkers, W. 49 Quadakkers, W. 93 Quadakkers, W. 271 S Schmid, S. 202 Searles, R. 3 Sedlmair, C. 223 Sehan, K. 223 Shibata, T. 83 Singheiser, L. 93 Strehl, G. 49 Strehl, G. 106 T Takao, K. 59 Taniguchi, S. 83 Tröger, U. 186 Tuomola, H. 152 U Untoro, P. 271 V Vermoehlen, M. 202 Vogt, C. 173 W Wilber, J. 93 Wirth, G. 191 Y Yamanaka, M. 263

Parker, S. 237 Pickering, S. 241 Q Quadakkers, W. 31

The page numbers refer to the first page of the article

Material Aspects in Automotive Catalytic Converters, Hans Bode Copyright © 2002 Wiley-VCH Verlag GmbH &Co. KgaA ISBN: 3-527-30491-6