Oxidation of Intermetallics

H. J. Grabke, M. Schutze

Oxidation of Intermetallics

8WILEY-VCH

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H. J. Grabke, M. Schutze

Oxidation of Intermetallics

@ WILEY-VCH

Weinheim . Berlin . New York . Chichester Brisbane Singapore - Toronto

Prof. Dr. H. J. Grabke Max-Planck-Institut fur Eisenforschung Max-Planck-StraBe 1 40237 Dusseldorf Germany

Dr. M. Schiitze Karl-Winnacker-Institut der DECHEMA e.V. Theodor-Heuss-Allee 25 60486 Frankfurt/Main Germany

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

i

Editorial Director: Dr. Jorn Ritterbusch Production Manager: Hans-Jochen Schmitt Cover picture: Metallographic section of the oxide scale and the subsurface zone of a twophase TiAl alloy after oxidation in air at 1350 "C for 1 h (interference layer metallography) Every effort has been made to trace the owners of copyrighted material; however, in some cases this has proved impossible. We take this opportunity to offer our apologies to any copyright holders whose rights we may have unwittingly infringed. 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: Oxidation of intemetallics / H. J. Grabke ; M. Schiitze. - Weinheim ; Berlin ; New York ; Chichester ; Brisbane ; Singapore ; Toronto : Wiley-VCH, 1997 ISBN 3-527-29509-7 0 WILEY-VCH Verlag GmbH, D-69469 Weinheim (Federal Republic of Germany), 1998 Printed on acid-free and chlorine-free paper All rights reserved (including those of translation into 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 a machine language without written permission from the publishers. Registered names, trademarks, etc. used in this book, wen when not specifically marked as such, are not to be considered unprotected by law. Composition: Mitterweger Werksatz, D-68723 Plankstadt Printing: Strauss Offsetdruck, D-69509 Morlenbach Bookbinding: Groabuchbinderei J . Schaffer, D-67269 Griinstadt Printed in the Federal Republic of Germany

Preface

About 10 years ago, many research activities were started in Europe, USA and Japan on the mechanical properties and oxidation resistance of intermetallic alloys at high temperatures, striving for new structural materials for application in engines and turbines. Especially aluminides and silicides appeared to be promising, which show high melting point, low density and high strength, and possibly good oxidation resistance. Vast efforts were made in research and development to attain improved high temperature strength and low temperature ductility, much less research and testing was performed on the resistance of intermetallics against oxidation and high temperature corrosion. Oxidation and corrosion resistance, however, are as important for actual application of materials as mechanical properties. Aluminides and silicides were expected to form protective alumina or silica layers and especially for the compounds with high A1 or Si content no great problems were foreseen. But in fact many phenomena were observed, which are deleterious for corrosion resistance: void formation beneath the scales, cracking and spalling, growth of unprotective oxides such as TiO, and Nb,O, on Ti-and Nb-aluminides, effects of nitride formation, even internal and accelerated oxidation, intergranular oxidation and ‘pesting’, i. e. disintegration of materials ... so a wide field opened for the study of kinetics and mechanisms of oxidation and corrosion attack. But also means of improving the oxidation resistance were found and from the relatively susceptible intermetallic compounds more resistant multi-component and multi-phase alloys were developed. After now, the possibilities and limitations of intermetallic compounds have been widely explored, and the research activities begin to decrease, it was time to try to get an overview on the present knowledge. The call for contributions brought 27 papers from Germany, Great Britain, Italy, and the Netherlands. Additionally, three keynote speakers were invited, G. Suuthoff to give an overview on the present state of intermetallics development and G. H. Meier and S. Tunigicchi to give keynote lectures on the research activities in the USA and Japan, on oxidation of intermetallics. In the workshop (held in Frankfurt, Germany, January 18-19,1996 and attended by 65 participants) the presentations and vivid discussions gave a rather complete view, especially of the oxidation and corrosion of the Ti-, Ni- and Fe-aluminides. Now this and the two next volumes of ‘Materials and Corrosion’ present a part of the papers in full length and another part as extended abstracts with references on where to find the complete work. The organizers of the workshop would like to extend sincere thanks to authors for their great effort in contributing to a fine compilation of work on the ‘Oxidation of Intermetallics’.

April 1996,H. J. Grubke

List of Contributors

P. Andrews Rolls Royce plc Derby United Kingdom

W. Auer Institute of Materials Science Chair for Corrosion and Surface Technology University of Erlangen-Nurnberg Martensstr. 7 91058 Erlangen Germany

M. J. Benett Coatings and Interfaces Section AEA Technology Harwell Laboratory Didcot Oxfordshire OX11 9AH United Kingdom M. Bobeth Max-Planck-Gesellschaft Research Group on Mechanics of Heterogeneous Solids Hallwachsstr. 3 01069 Dresden G e r rna n y

L. B. Bradley Corrosion and Protection Center University of Manchester Institute of Science and Technology Manchester M60 1QD United Kingdom

M. W. Brumm Max-Planck-Institut fur Eisenforschung GmbH Postfach 14 04 44 40074 Dusseldorf Germany S. J. Bull Coatings and Interfaces Section AEA Technology Harwell Laboratory Didcot Oxfordshire OX1 1 9AH United Kingdom

K. T. Chuah Corrosion and Protection Center University of Manchester Institute of Science and Technology Manchester M60 1QD United Kingdom F. Dettenwanger Max-Planck-Institut fur Metallforschung Institut fur Werkstoffwissenschaft Seestr. 92 70174 Stuttgart Germany M. Eckert Research Center Julich Institute of Materials in Energy Systems P.O. Box 1913 52425 Jiilich Germany

VlII

List of Contribicfors

B. Eltester Max-Planck-Institut fur Eisenforschung GmbH Postfach 14 04 44 40074 Dusseldorf Germany

A. H. H. Janssen Netherlands Energy Research Foundation ECN PO Box 1 1755 ZG Petten The Netherlands

A. Gil University of Mining and Metallurgy 30059 Krakow Poland

J. Klower KruppVDM CimbH Plettenberger Str. 2 58791 Werdohl Germany

H. J. Grabke Max-Planck-Institut fur Eisenforschung GmbH Postfach 14 04 44 40074 Diisseldorf Germany M. J. Graham National Research Council Institute for Microstructural Sciences Ottawa, K1A OR6 Canada V. A. C. Haanappel Institute for Advanced Materials Joint Research Center of the European Commission 21020 Ispra (Va) Italy

K. Hilpert Research Center Julich Institute of Materials in Energy Systems PO. Box 1913 52425 Julich Germany M. Hollatz Max-Planck-Gesellschaft Research Group on Mechanics of Heterogeneous Solids Hallwachsstr. 3 01069 Dresden Germany

R. Klumpes Delft University of Technology Lab. for Materials Science Div. of Corrosion Technology and Electrochemistry Rotterdamseweg 137 2628 AL Delft The Netherlands V. Kolarik Fraunhofer-Institut fur Chemische Technologie 76327 Pfinztal Germany R. Krajak Max-Planck-Institut fur Eisenforschung GmbH Postfach 14 04 44 40074 Diisseldorf Germany R. Kremer Institute of Materials Science Chair for Corrosion and Surface Technology University of Erlangen-Niirnberg Martensstr. 7 91058 Erlangen Germany

List of Contributors

C. Lang Karl-Winnacker-Institut der DECHEMA e.V. Theodor-Heuss- Allee 25 60486 Frankfurt/M. Germany J. Leggett Cranfield University Cranfield Bed f ord United Kingdom C. H. M. Maree University of Utrecht Department of Atomic and Interface Physics 3508 TA Utrecht The Netherlands

IX

W. Pompe Max-Planck-Gesellschaft Research Group on Mechanics of Heterogeneous Solids Hallwachsstr. 3 01069 Dresden Germany W. J. Quadakkers Forschungszentrum Julich GmbH IWE 1 52425 Jiilich Germany

J. Rakowski Department of Materials Science and Engineering University of Pittsburgh Pittsburgh, PA 15261 USA

G. H. Meier Department of Materials Science and Engineering University of Pittsburgh Pittsburgh, PA 15261 USA

I. Rommerskirchen Krupp VDM GmbH Plettenberger Str. 2 58791 Werdohl Germany

J. R. Nicholls Cranfield University Cranfield Bedford United Kingdom

M. Ruhle Max-Ylanck-Institut fur Metallforschung Institut fur Werkstoffwissenschaft Seestr. 92 70174 Stuttgart Germany

M. Palm

Max-Planck-Institut fur Eisenforschung GmbH Postfach 14 04 44 40074 Dusseldorf Germany

G. Sauthoff Max-Planck-Institut fur Eisenforschung GmbH Postfach 14 04 44 40074 Dusseldorf Germany

B. A. Pint Oak Ridge National Laboratory PO. Box 2008 Oak Ridge,TN 37831-6156 USA

P. Schaaf Universitat Gottingen 37073 Gottingen Germany

X

List of Conlributors

Institute for Advanced Materials Joint Research Center of the European Commission 21020 Ispra (Va) ltaly

S. StrauB Max-Planck-Institut fur Eisenforschung GmbH Postfach 14 04 44 40074 Dusseldorf Germany

B. Schramm

M. F. Stroosnijder

H. J. Schmutzler

Institute of Materials Science Chair for Corrosion and Surface Technology University of Erlangen-Nurnberg Martcnsstr. 7 91058 Erlangen Germany E. Schramm Delft University of Technology Lab. for Materials Science Div. of Corrosion Technology and Electrochemistry Rotterdamseweg 137 2628 A L Delft ‘The Netherlands E. Schumann Max-Planck-Institut fur Metallforschung Institut fur Werkstoffwissenschaft Seestr. 92 70174 Stuttgart Germany

M. Schiitze

Karl-Winnacker-Institut der DECHEMA e.V. Theodor-Heuss- Allee 25 60486 FrankfurtIM. Germany

F. H. Stott Corrosion and Protection Center University of Manchester Institute of Science and Technology Manchester M60 1QD United Kingdom

Institute for Advanced Materials Joint Research Center of the European Commission 21020 Ispra (Va) Italy

J. D. SunderkGtter Institute for Advanced Materials Joint Research Center of the European Commission 21020 lspra (Va) Italy S.Taniguchi Department of Materials Science and Processing Faculty of Engineering Osaka University’ 2-1 Yamadaoka Suita, Osaka 565 Japan P. ETortorelli Oak Ridge National Laboratory PO. Box 2008 Oak Ridge,TN 37831-6156 USA

J. P. 7’.Vossen Netherlands Energy Research Foundation ECN PO Box 1 1755 ZG Petten The Netherlands

B. Wagemann Max-Planck-lnstitut fur Eisenforschung GmbH Postfach 14 04 44 40074 Dusseldorf Germany

E. Wallura Forschungszentrum Julich GmbH IWE 1 52425 Julich Germany J. H. W. de Wit Netherlands Energy Research Foundation ECN PO Box 1 1755 ZG Petten The Netherlands J. H. W. de Witt Delft University of Technology Lab. for Materials Science Div. of Corrosion Iechnology and Electrochemistry Rotterdamseweg 137 2628 AL Delft The Netherlands

I. G. Wright Oak Ridge National Laboratory P.O. Box 2008 Oak Ridge.TN 37831-6156 USA J. C. Yang Max-Planck-Institut fur Mctallforschung Institut fur Werkstoffwissenschaft Seestr. 92 70174 Stuttgart Germany

N. Zheng Forschungszentrum Julich GmbH IWE I 52425 Julich Germany

Contents

Preface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

V

List of Contributors . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

VII

Part I Introduction 1

2 3

State of Intermetallics Development .......................... G. Sauthoff Research on Oxidation and Embrittlement of Intermetallic Compounds in the U.S. ...................................... G, H. Meier Oxidation of Intermetallics - Japanese Activity . . . . . . . . . . . . . . . . S. Taniguchi

3

15 59

Part I1 Ni-Aluminides 4 5 6

7

8 9

10

The Oxidation of NiAl and FeAl .............................. 79 H. J. Grabke, M. W Brumm, B. Wagemann Sulfidation Behaviour of Nickel Aluminides . . . . . . . . . . . . . . . . . . . 85 B. Schramm, W Auer The Influence of Chromium on the Oxidation 99 of P-NiAI at 1000 "C ........................................ R. Klumpes, C. H. M. Marie, E, Schramm, J. H. W de Wit Oxidation of P-NIAI, Undoped and Doped with Ce,Y, Hf . . . . . . . 109 I. Rommerskirchen, I/: Kolarik TEM Investigations on the Oxidation of NiAl . . . . . . . . . . . . . . . . . . 121 E. Schumann, J. C. Yang, M. J. Graham, M. Ruhle Failure of Alumina Scales on NiAl Under Graded Scale Loading .............................................. 135 M. Hollatz, M. Bobeth, W Pompe The Corrosion Behaviour of NiAl in Molten Carbonate 161 at 650 "C ................................................... J. I? T Vossen,A. H. H. Janssen, J. H. W de Wit

XIV

Contents

Part 111 Fe-Aluminides 11 12

13 14

15

Oxidation of P-FeAl and Fe-AI Alloys ......................... 175 I. Rommerskirchen, B. Eltester, H. J. Grabke The Oxidation Behaviour of ODS Iron Aluminides . . . . . . . . . . . . . 1S3 B. A. Pint, F! F: Tortorelli and I. G. Wright High Temperature Corrosion Behaviour of Iron Aluminides and Iron-Aluminium-Chromium Alloys ....................... 203 J. Klower 221 Oxidation-Sulphidation of Iron Aluminides .................... E H. Stott, K. 7: Chuah, L. B. Bradley Metal Dusting of Fe,AI and (Fe, Ni),AI ........................ 233 S. Strauo., R. Krajak, M . Palm, H. J. Grabke

Part IV Ti-Aluminides 16

17 18

19 20

21

22

23

Determination of Thermodynamic Activities in the Alloys of the Ti-AI System and Prediction of the Oxidation Behaviour of the Alloys ................................................ M , Eckert, K. Hilpert The Initial Stages in the Oxidation of TiAl ..................... C. Lang, M. Schiitze Development and Microstructure of the Al-Depleted Layer of Oxidized TiAl ............................................. E Dettenwanger, E. Schumann, J. Rakowski, G. H. Meier, M. Riihle Beneficial and Detrimental Effects of Nitrogen on the Oxidation Behaviour of TiAl-Based Intermetallics ....................... M! J. Quadakkers, I? SchaaJ N Zheng,A. Gil, E. Wallura Influence of Moisture on the Oxidation of y-XAl . . . . . . . . . . . . . . . R. Kremer, M!Auer Ion Implantation as a Tool to Study the Oxidation Behaviour on TiAl-Based Intermetallics ................................. M. E Stroosnijder, H. J. Schmutzler,V A. C, Haanappel, J. D. Sunderkotter Protection of Titanium Aluminides by FeCrAlY Coatings . . . . . . . M. J. Bennett, S. J. Bull Hot Salt Corrosion of Titanium Aluminides .................... J. R. Nicholls, J. Leggett, F! Andrews

Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

239 245 265 275 289 299 313 329 345

Part I Introduction

Oxidation of Intermetallics by H. J. Grabke, M. Schiitze 0 WILEY-VCH Verlag GmbH, 1998

1 State of Intermetallics Development G. Sauthoff

1.1 Introduction

It is known since long that intermetallic phases excel by special physical and mechanical properties, the latter being high strength at high temperatures and low deformability at lower temperatures [l-31. Since the fifties it has been tried to develop new high-temperature materials on the basis of intermetallic phases without resulting in applications in the past because of processing problems [4,5]. Only more recent developments, which were initiated by the “ductilization” of Ni,AI by the addition of boron [6], have progressed successfully with centres of gravity in USA, Japan and Germany and are at the brink of commercialisation [7-lo]. The developments have been paralleled by extended fundamental research which, however, has not yet led to a clear physical understanding of the observed phenomena [2, 11-13]. In the following the state of development and the prospects of the various candidate phases for high-temperature applications are overviewed.

1.2 Titanium Aluminides 1.2.1 Ti,AI In the USA the a2alloys and super-a, alloys have been developed by alloying Ti,Al with large amounts of Nb which show sufficient ductility at room temperature and aim at applications in aircraft industry [14-161. The materials development is practically complete, production and processing technologies have been developed, and alloys are commercially available. However, these alloys are not yet applied because of insufficient oxidation resistance, higher demands for strength and high economic barriers [17-191. Corresponding developments have not been done in Germany because of the limited potential of these alloys, but possibilities for alloy optimization by microstructure control have been studied successfully 1201.

4

G. Southoff

1.2.2 TiAl Compared with Ti,AI, the Al-richer so-called y titanium aluminide TiAl is regarded as more promising because of lower density, higher strength at higher temperatures and higher oxidation resistance [18]. Thus the world-wide efforts to develop a lightweight alloy for application as high-temperature structural alloys are concentrated presently on ‘IiAI. The intensive developments in USA have led to the y titanium aluminide alloys which contain 48 at.% Al and about 2 at.% Nb, Cr and/or other transition metals and which are two-phase with Ti,AI as second phase [21-241. The mechanical behaviour as characterized by strength, ductility and toughness can be optimized by careful control of processing, i. e. microstructure which is lamellar, duplex or mixed. These developments aim at aerospace applications, whereas a TiAl turbocharger rotor has been developed in Japan for application in car engines [25].Further Japanese developments are subject of a national materials program [26]. The enormous potential of TiAl has been recognized early in Germany, Austria and Switzerland which gave rise to various ongoing developments [27-361. Development aims are applications as valves in car engines, turbocharger rotors, gasturbine blades and lightweight sheet material.

1.3 Nickel Aluminides 1.3.1 Ni,AI Ni,AI is known since long as the strcngthening y-phase in superalloys and has been described repeatedly in review papers [37-391. Characteristic is its anomalous temperature dependence of flow stress, i. e. increasing flow stress with rising temperature up to temperatures of about 700°C and normal softening at higher temperatures. Polycrystalline Ni,Al is practically not deformable at room temperature, whereas singlecrystalline Ni,AI is ductile. Polycrystalline Ni,AI with reduced Al content can be “ductilized” by microalloying with B. Ni,Al suffers from environmental embrittlement at room temperature, which is claimed to be caused by moisture in the atmosphere, and from high-tcmperature embrittlement at about 700°C which is reduced by Cr addition. On the basis of “ductilized” Ni,Al the so-called Advanced Aluminides were developed at the Oak Ridge National Laboratory which are now ready for application [38, 401. These alloys are recommended for applications in water turbines, steam turbines and gas turbines where high strength and high resistance against fatigue, wear, erosion, cavitation and oxidation is important. A problem is the limited hot-forming capability. Ni,Al-base alloys are also promising for applications in combustion engines as is exemplified by a development aiming at turbocharger rotors in large Diesel engines [41]. The potential of Ni,Al-base alloys is regarded as limited since it does not differ sufficiently from that of the Ni-base superalloys [42]. German development work concentrates on production technology and aims at making use of the high corrosion resistance of Ni,Al-base alloys [43].

I Slate of Internwtallics Development

5

1.3.2 NiAl Compared with NiAI, and conventional superalloys, NiAl is regarded as more promising for high-temperature structural applications because of higher melting temperature, lower density and higher thermal conductivity [4446]. Furthermore, NiAl shows an excellent oxidation resistance which excels that of most other intermetallics [47]. Problems are the low deformability at room temperature and the insufficient strength and creep resistance above 1000°C. It is noted that Al-deficient NiAl, which is instable, can transform martensitically, which is used for developing shape-memory alloys [48--501.Intensive developments in USA are aimed at applications as blades in flying gas turbines and have already led to components which have been tested successfully under service conditions 1511.The world-wide development efforts have been paralleled by cxtended fundamental research which has been reviewed repeatedly [52-551. In Germany the potential of NiAl-base alloys has been recognized early which initiated respective developments [56-591. These developments rely on the formation of second phases by alloying with third elements which improves high-temperature strength and creep resistance sufficiently with still tolerable room-temperature brittleness. A special development has been based on eutectic NiAl alloys with a ductile Cr, Mo or Re phase and aims at applications in flying gas turbines and car engines [60, 611. Another development by mechanical alloying makes use of oxide dispersions for increasing the creep resistance and of ductile particles for increasing the toughness [62,63]. A third development results from alloying NiAl with Ta and/or Nb to produce a strengthening Laves phase of type TaNiAl or NbNiAl [57,64-67]. Such alloys show not only an advantageous combination of high creep resistance at 1200°C and acceptable toughness at room temperature, but also an excellent thermoshock resistance and hot-gas corrosion resistance which makes them particularly promising for applications in car engines. A further development makes use of alloying with Fe to produce a second ductile particulate phase, and the resulting alloy with additional Cr shows an advantageous combination of strength, ductility and corrosion resistance against oxidation, carburization and sulphidation with a high potential for applications in coal conversion plants, petrochemical plants and industrial furnaces [68,69].

1.4 Iron Aluminides 1.4.1 Fe,Al The particular magnetic properties of Fe,AI resulted in the development of the magnetic head material Sendust, which is based on Fe,(AI,Si) and is applied world-wide in huge quantities [70).Presently Fe,AI is regarded as promising for structural applications, and respective developments have been started in USA [71-731. Fe,Al-base alloys show higher strength with anomalous temperature dependence in comparison to conventional iron alloys, and they excel by their corrosion resistance in oxidizing and sulphidizing atmospheres. Problems are the little deformability at room temperature and the strength decrease above about 500°C because of lattice transformation [74].

6

G. Sauthoff

The mechanical behaviour is improved by reduction of Al content, alloying with further elements and thermomechanical treatments, which leads to alloys for eventual applications in conventional power plants - in particular in steam turbines - or in coal conversion plants. Exhaust units in cars and electric heating elements are near application. In Germany Fe,AI has found only little interest up to now - apart from fundamental research [75,76] - and the little interest aims at applications in the petrocheniical industry [77].

1.4.2 FeAl FeAl shows the same crystal structure as NiAl, and the behaviour of FeAl is indeed similar to that of NiAI. A ductility of few percent at room temperature could be achieved by reducing the A1 content and appropriate processing [78]. FeAl is regarded as highly promising for high-temperature applications in view of the advantageous deformation and corrosion properties and the comparatively low density, and respective materials developments are on the way in IJSA, including the development of intermetallic matrix composites (IMC) with FeAl matrix [74,79-811.

1.5 Chromides Already in the past the Laves phase TiCr,, which crystallizes with the hexagonal C14 or C36 structure or the cubic C15 structure depending o n temperature and composition, was regarded as promising for high-temperature applications because of high strength and oxidation resistance 182,831. However, its high brittleness at room temperature has precluded any direct application. Present development efforts in IJSA are based on multiphase TiCr,-Ti-Nb alloys with the hard Laves phase in a ductile matrix, which offer an acceptable room-temperature toughness [84,85]. A rather similar Laves phase is NbCr,, which has been selected as candidate phase for high-tcmperature applications because of its high melting temperature, high strength and creep resistance, low density and potential oxidation resistance [86-88]. Again this phase is combined with ductile Cr-rich or Nb-rich phases as matrix or particles to obtain multiphase alloys with acceptable toughness.

1.6 Silicides Silicides represent the transition from intermetallics with predominantly metallic bonding to non-metallic compounds since-silicon is no longer a metal, but a semiconductor. Nevertheless the silicides are often comprised within the intermetallics. Silicides were selected for high-temperature applications already in the past because of their potentially high oxidation resistance at highest temperatures [89-911. Presently

1 Slate of Inrermetallics Developtnent

7

new developments are in progress [92].The general interest is concentrated on MoSi,, which has been applied since long as electric-heating material [93-951. Various other alloy systems have been selected, in particular Nb,Si,-Nb and Cr,Si-Cr, which allow ductile phase toughening [96-991. In Germany various silicides have been selected for materials developments. Ti,Si, excels by its high-temperature capabilities with melting temperature of about 2130°C and its low density of about 4g/cm3 [loo, 1011. However, this phase is brittle at lower temperatures, and its potential can be used only by combining this phase with more ductile phases. It has indeed been shown that Ti$,-Ti,Al alloys on the one hand [102, 1031 and Ti$,-Ti alloys on the other hand [loo, 1041 are promising as high-temperature lightweight materials. The high oxidation resistance of the disilicides MoSi, and TiSiz has initiated additional developments [105-1071. The magnesium-rich silicide Mg,Si has been selected as basis for the development of two-phase Mg,Si-Mg and Mg,Si-A1 lightweight alloys for application as piston material in car engines because of the outstandingly low density of only 1.88g/cm3of this phase [108-110].

1.7 Conclusions and Prospects New structural intermetallic alloys for high-temperature applications are in the centre of the present interest in intermetallics which is still growing. Some few developments, which are based on the classic phases Ni,AI, Ti,AI and TiA1, are on the brink of commercialization, but even these developments are still in an early stage compared with other developments of advanced materials, e. g. the modern engineering ceramics. Much more experimental and theoretical work is necessary for solving the processing problems and for adjusting the property spectra to the specific applications. The much advanced nickel aluminides and titanium aluminides can be used only up to about 1000°C because of limited strength or oxidation resistance or both at higher temperatures as has been stated before [3]. For applications significantly above 1000°C other less-common phases with higher melting temperatures have to be used. Such phases are available, and examples are shown in Fig. 1 [ lll] . In comparison to the nickel aluminides and titanium aluminides, the less-common phases are stronger and more brittle, their crystal structures are complex, their handling is difficult, and thus they are regarded as exotic. However, these exotic phases may fill the gap between metallic high-temperature alloys and ceramics, as is visible in Fig. 1.The brittleness of the less-common phases can be alleviated by combining them with softer phases to form multiphase alloys with adequate microstructures. Even strengthening hard phases may improve the toughness by impeding crack growth. The mechanical behaviour can be optimized by optimizing the microstructure, which requires a careful control of processing. However, it has to be emphasized that one cannot expect to obtain new intermetallic materials with properties similar to existing conventional metallic alloys. The “ductilizations” that have been achieved in few cases - in particular Ni,AI and (Fe,Co,Ni),V - rely on rather specific mechanisms and cannot be expected for other intermetallic phases. Thus intermetallic materials have to be regarded as a materials

8

G. Smithoff

25 E

Y

.s 20 v) v)

a L

t

v,

9 .-a

15

)5

U .-

y-

lo

!!i d 5

v)

0

500

1000 1500 temperature in O C

Fig.1. Specific yield strength (0.2 % proof stress in compression per unit weight density at 10-4s-1strain rate in compression) as a function of temperature for the DO,, phase AI,Nb (112. 1131,the Heusler-type phase Co,TiAl[67], the Laves phasesTiCr, sSi,, andTaFeAl[67,114]. the two-phase alloy NbNiAl-NiAl with 15vol.% NiAl in the Laves phase NbNiAl[67,114], and the hexagonal D8, phaseTi,Si, [lo01in comparison to the superalloy MA 6000 (in tension) [I 151 and the hot-pressed silicon rutride HPSN (upper limit of flexural strength) [116].

class of its own with property spectra which differ significantly from those of other materials and which can be varied within broad limits corresponding to metals on one side and non-metals on the other side. This offers enormous possibilities for manifold developments which are exciting with respect to;ooth practical applications and materials science. Finally it is noted that much interest is concentrated on the development of intermetallic alloys for application as blades in flying gas turbines. This application is most demanding and it is not clear when all the problems with strength, ductility, toughness and corrosion resistance can be solved at economic costs. Less-high-technology applications may be more rewarding presently for the introduction of new intermetallic materials. An example may be the car engine where strong light components with sufficient corrosion resistance are needed. Here brittleness is not the problem since designers have learnt to use ceramic materials for e. g. valves. However, new materials must be compatible with the metallic engine with respect to the physical properties, in

I State of Intermetallics Development

9

particular thermal expansion and thermal conductivity. This requirement corresponds to the characteristics of intermetallic materials which are hard and brittle with mainly metallic atomic bonding, i. e. metallic physical properties. Thus new intermetallic materials are expected to play an important role in the manufacture of car engines and similar applications.

1.8 References [l] G. Suuthoffi Intermetallics. In: K. H. Matucha (ed.) Structure and Properties of Nonferrous Alloys. Verlag Chemie, Weinheim (1995) 643-803. [2] G. Sauthoff:1ntermetallics.Verlag Chemie, Weinheim, (1995). [3] G. Suurhoff:Plastic Deformation. In: J.H. Westbrook and R. L. Fleischer (eds.) Intermetallic Compounds: Principles and Practice. John Wiley & Sons, Chichester (1995) 911-934. [4] J. H. Westbrook:Mechanical Properties of Intermetallic Compounds - A Review of the Literature. In: J. H. Westbrook (eds) Mechanical Properties of Intermetallic Compounds. J. Wiley. New York (1960) 1-70. [5] J. H. Westbrook: Silicides, Borides, Aluminides, Intermetallics and Other Unique Refractories. In: High Temperature Technology. McGraw Hill, New York (1960) 113-128. [6] K.Aoki, 0. Zzumi:J. Jap. Inst. Metals 43 (1979) 1190. [7] G. Sauthoff: Neue Strukturwerkstoffe auf der Basis intermetallischer Phasen - Stand und Perspektiven. BMFT-Symposium Materialforschung - Neue Werkstoffe, KFA-PLR, Jiilich (1994) 309-322. [8] H.-J. Engell, A . Von Keitz, G. Sauthoff: Intermetallics - Fundamentals and Prospects. In: W. Bunk (ed.) Advanced Structural and Functional Materials. Springer Verlag, Berlin (1991) 91-132. 191 G. Sauthoff: Internationaler Stand der Werkstoffentwicklungen auf der Basis intermetallischer Phasen. In: F. J. Bremer (ed.) Intermetallische Phasen als Strukturwerkstoffe fur hohe Temperaturen. Forschungszentrum Julich GmbH, Jiilich (1991) 1-13. (lo] G. Suuthoffi Z . Metallkde. 81 (1990) 855. Ill] High-Temperature Ordered Intermetallic Alloys V (MRS Symp. Proc. Vol. 288). Materials Research Society, Pittsburgh, (1993). [12] Structural Intermetallics. TMS, WarrendalelPA, (1993). [13] Ordered Intermetallics - Physical Metallurgy and Mechanical Behaviour. Kluwer Acad. Publ. Dordrecht, (1992). [14] D. Banerjee:Ti,Al and its Alloys. In: J. H. Westbrook and R. L. Fleischer (eds.) Intermetallic Compounds: Principles and Practice, Vol. 2. John Wiley & Sons, Chichester (1995) 91-132. [15]H. A . Lipsitt: Mat. Res. SOC.Symp. Proc.288 (1993) 119. [16] H. A. Lipsitt: Titanium Aluminides - An Overview. In: C. C. Koch, C. T. Liu, and N. S. Stoloff (eds.) High-Temperature Ordered Intermetallic Alloys. MRS, Pittsburgh (1985) 351-364. [17] C. M.Ward-Close,E H. Froes:J. Metals 46 (1994) 28. [18] f? H. Froes, C. Suryanarayana,D. Eliezer: J. MateT. Sci. 27 (1992) 5113. [19] E H. Froes, C. Suryanarayana,D. Eliezer: ISIJ Intl. 31 (1991) 1235. [20] G. Proske, G. Liitjering,J.Albrecht, D. Helm, M . Daeubler: Mater. Sci. Eng. A152 (1992) 310. [21] Y - W Kim: J. Metals 47 (1995) 39. [22] Y - W Kim: J. Metals 46 (1994) 30. [23] S. C. Huang, J. C. Chesnutt: Gamma TiAl and Its Alloys. In: J. H. Westbrook and R. L. Fleischer (eds.) Intermetallic Compounds: Principles and Practice, Vol. 2. John Wiley & Sons, Chichester (1994) 73-90. [24] D. M. Dimiduk, D. B. Miracle, Y - W Kim, M . G. Mendiratta:ISIJ Intl.31 (1991) 1223. [25] Y Nishiyama, Z Miyushita, S. Zsobe, Z Noda: Development of Titanium Aluminide TurboCharger Rotors. In: s.H. Whang, C. T. Liu, D. F! Pope et al. High-Temperature Aluminides and IntermetallicsTMS, Warrendale (1990) 557-584.

10

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[26] M . Yamaguchi, H. Inui: TiAl Compounds for Structural Applications. In: R. Darolia. J. J. Lewandowski, C. T. Liu et al. Structural Intermetallics. TMS, Warrendale/PA (1993) 127-142. [27] H. Clemens: Z . Metallkde. 86 (1995) 814. [28] C. Koeppe, A . Bartels, H. Clemens, I? Schretter, W Glatz:Mater. Sci. Eng.,A201 (1995) 182. (291 R. Wagner, E Appel, B. Dogan, I? J. Ennis, U. Loreni, J. Miillauer, H. I? Nicolai, W Qiradakkers, L. Singheiser, W Stnarsly, W Vaidya, K. Wirrzwallner: Investment Casting of Gamma TiAl Based Alloys: Microstructure and Data Base for GasTurbine Applications. 1n:Y.-W. Kim, R. Wagner, and M. Yamaguchi (eds.) Gamma Titanium Aluminides (ISGTA ‘95).TMS, Warrendale/PA (1995) 387-404. [30] R. Wagner: Intermetallische Gamma-Titan-Aluminide - Von den Grundlagen zur Bauteilanwendung -. BMFT-Symposium Materialforschung Neue Werkstoffe. KFA-PLR, Jiilich (1994) 753-756. (311 M . Nazmy, M. Staubli: Scripta Metall. Mater. 31 (1994) 829. [32] A. Bartels, C. Koeppe, K. Wurzwallner, I? Schretter, H. Clemens: Microstructure and Mechanical Properties of TiAI Alloys after Thermomechanical Processing. In: H. Bildstein and R. Eck (eds.) Plansee Proceedings - Proc. 3‘h Interntl. Plansee Seminar. Vol. 3. Plansee Metall AG, Reutte (1993) 564-577. [33] H. Clemens, I? Schretter, K. Wurzwallner, A . Bariels, C. Koeppe: Forging and Rolling of Ti48A12Cr on Industrial Scale. In: R. Darolia, J. J. Lewandowski, C.T. Liu et al. Structural Intermetallics.TMS, WarrendalelPA (1993) 205-214. [34] (1. Herold-Schmidt, H . Kiihnle, S.Schwantes: Oxidation and Versprodung von TitanaluminidBlechen. In: H. Bildstein and R. Eck (eds.) Plansee Proceedings - Proc. 3Ih Interntl. Plansee Seminar. Vol. 3. Plansee Metall AG, Reutte (1993) 607421. [35] I? A. Beaven, EAppel, B. Dogan, R. Wagner:Fracture and Ductilization of Gamma-Titanium Aluminides. In: C. T. Liu, R. W. Cahn, and G.Sauthoff (eds.) Ordered Intermetallics - Physical Metallurgy and Mechanical Behaviour. Kluwer Acad. Publ. Dordrecht (1992) 413432. [36] G. Frommeyer, W Wunderlich, Z Kremser, Z . C. Liu:Mater. Sci. Eng. A152 (1992) 166. [37] D. L. Anton: Ni,AI in Nickel-Based Superalloys In: J. H. Westbrook and R. L. Fleischer (eds.) Intermetallic Compounds: Principles and Practice, Vol. 2. John Wiley & Sons, Chichester (1995) 3-16. [38] C. Z Liu, D. I? Pope: Ni,AI and its Alloys. In: J. H. Westbrook and R. L. Fleischer (eds.) Intermetallic Compounds: Principles and Practice,Vol. 2. J. Wiley, Chichester (1995) 17-52. [39] A! S. Stofoff:Intl. Metals Rev.29 (1984) 123. [40]C.Z Liu, J. 0.Stiegler, E H. Froes: Ordered Intermetallics. Metals Handbook Vol. 2: Properties and Selection: Non-Ferrous Alloys and Special Purpose Materials. ASM, Materials Park (1990) 913-942. [41] J. W Panen: Nickel Aluminides for Diesel Engines. In: S. H. Whang, C. T. Liu, D. P. Pope et al. High-Temperature Aluminides and Intermetallics. TMS, Warrendale (1990) 493-503. [42],D. M. Dimiduk, D. B. Miracle, C. H. Ward:Mater. Sci.Technol.8 (1992) 367. [43] U.Brill, J. Klower: High Temperature Corrosion of Intermetallic Phases Based on Ni,AI. In: L. A. Johnson, D. I? Pope, and J. 0. Stiegler (eds.) High Temperature Ordered Intermetallic AUoys IV. MRS, Pittsburgh (1991) 963-968. [44] D. B. Miracle, R. Darolia: NiAl and its Alloys. In: J. H. Westbrook and R. L. Fleischer (eds.) Intermetallic Compounds: Principles and Practice, Vol. 2. John Wiley & Sons, Chichester (1995) 53-72. [45] R. Darolia: NiAl for Turbine Airfoil Applications. In: R. Darolia. J. J. Lewandowski, C. 7:Liu et al. Structural Interrnetallics. TMS, Warrendale/PA (1993) 495-504. [46]R. Darolia: J. Metals 43(3) (1991) 44. [47] G. H. Meier, A! Birks, E S. Petiit, R. A . Perkins, H. J. Crabke: Environmental Behavior of Intermetallic Materials. In:R. Darolia, J. J. Lewandowski, C.T. Liu et al. Structural Intermetallics. TMS, Warrendale/PA (1993) 861-877. [48] R. Kainuma, K. Ishida, Z Nishizawa: Metall.Trans. 23A (1992) 1147. [49] R. Kainuma, H. Nakano, K. Oikawa, K. Ishida, 7:Nishizawa: High Temperature Shape Memory Alloys of Ni-AI Base Systems. In: C.T. Liu, M. Wuttig, K. Otsuka et al. Shape-Memory Materials and Phenomena - Fundamental Aspects and Applications MRS, Pittsburgh (1992) 403-408. [50] S. Furukawa, A . Inoue, Z Masumoto: Mater. Sci. Eng. 98 (1988) 515. ~

I Slate of Inrernietallics Development

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[51] R. Daro1ia:Acta Met. Sin. (Engl. Lett.) 8 (1995) 625. [52] 1. Baker: Mater. Sci. Eng. A193 (1995) 1. [53] D.B. Mirac1e:Acta Metall. Mater. 41 (1993) 649. [54] R. D. Noebe, R. R. Bowman, M . k! Nathal: Intl. Mater. Rev.38 (1993) 193. [55] I. Baker, P R. Munroe: Properties of B2 Compounds. In: S. H. Whang, C. T. Liu, D. P. Pope et al. High-Temperature Aluminides and Intermetallics. TMS, Warrendale (1990) 425-452. [56] G. Sauthofi Neue Strukturwerkstoffe auf der Basis intermetallischer Phasen. 2. Symposium Materialforschung des BMFT. KFA-PLR, Jiilich (1991) 877-898. [57] G. Sauthoff: Intermetallische Phasen. Symposium Materialforschung 1988. KFA-PLR, Jiilich (1958) 399-414. [58] M.Rudy, I. Jung, G. Saurhoffi Ferritic Fe-Ni-A1 Alloys for High Temperature Applications. In: J.B. Marriott, M. Merz, J. Nihoul et al. High Temperature Alloys -Their Exploitable Potential. Elsevier Appl. Sci., London (1987) 29-37. [59] M. Rudy, G. Sauthofl: Creep Behaviour of the Ordered Intermetallic (Fe,Ni)Al Phase. In: C. C. Koch, C.T. Liu, and N. S. Stoloff (eds.) High-Temperature Ordered Intermetallic Alloys. MRS, Pittsburgh (1985) 327-333. [60] W Kowalski: Mikrostruktur und mechanische Eigenschaften intermetallischer NiAl-Cr-Legierungen. VDI Verlag, Diisseldorf, (1994). [61] W Kleinekathofer, A. Donne6 H. Meinhardt, M. Hengerer, G. Sauthofif B. Zeumer, G. Frommeyer, H. J. Schiifer: Entwicklung von intermetallischen NiAl-Basislegierungen fur Motorenkomponenten. BMFT-Symposium Materialforschung - Neue Werkstoffe. KFAPLR,Jiilich (1994) 1014-1015. [62] E. Arzt, P Grahle: Mat. R e s SOC.Symp. Proc.364 (1995) 525. [63] E. Arzr, E. Gohring, P Grahle: Mat. Res. SOC.Symp. Proc. 288 (1993) 861. [64]G. Sauthofj W Kleinekathofer: Hochstwarmfeste NiAl-Basis-Legierungen fur Strukturbauteile im Verbrennungsmotor. In: Effizienzsteigerung durch innovative Werkstofftechnik. VDI Verlag, Dusseldorf (1995) 647-654. [65] B. Zeumer, W.Wunnike-Sanders, G. Sauthoffi Mater. Sci. Eng. A 192/193 (1995) 817. [66] G. Saurhoff: High Temperature Deformation and Creep Behaviour of BCC Based Intermetallics. In: 0.Izumi (ed.) Proceedings of the International Symposium on Intermetallic Compounds - Structure and Mechanical Properties - (JIMIS-6). The Japan Institute of Metals, Sendai (1991) 371-378. [67] G. Sauthoffi Mechanical Properties of Intermetallics at High Temperatures. In: S. H. Whang, C. T. Liu, D. P. Pope et al. High-Temperature Aluminides and Intermetallics.TMS, Warrendale (1990) 329-352. [68] J. Klower, G. SauthoB D. Letzig:Alloy 10 A1 - A New Sulphidation and Carburization Resistant Alloy for Fuel Combustion and Conversion. In: Corrosion 96. NACE International, HoustorvTexas (1996) 144/1-144/15. [69] D. Letrig: Zur Entwicklung kaltumformbarer NiAl-Basis-Legierungen. Dr. rer. nat. Dissertation, RWTH Aachen (1995). [70] T. Yamamoto:On the Discovery of the High Permeability Alloy “Sendust” and the Progress of its Industrialization. In: T. Yamamoto (ed.) The Development of Sendust and Other Ferromagnetic Alloys. Committee of Academic Achievements, Chiba (1980) 1-6. [71] S. C. Deevi, I/:K. Sikka: Intermetallics 4 (1996) 357. [72] K K. Sikka: Processing and Applications of Iron Aluminides In: J. H. Schneibel and M. A. Crimp (eds.) Processing, Properties, and Applications of Iron Aluminides. TMS, Warrendale (1994) 3-18. [73] k! K . Sikka, S.Vkwanathan, C. G. McKamey: Development and Commercialization Status of Fe,Al-Based Intermetallic Alloys. In: R. Darolia, J. J. Lewandowski, C. T. Liu et al. Structural Intermetallics. TMS, Warrendale/PA (1993) 483-491. [74] K. Vedula:FeAl and Fe,AI. In: J. H. Westbrook and R. L. Fleischer (eds.) Intermetallic Compounds: Principles and Practice,Vol. 2. John Wiley & Sons, Chichester (1995) 199-210. [75] H. Rosner, E. Nembach: Mater. Sci. Eng. A 196 (1995) L 1. [76] W Schroer, C. Hartig, H. Mecking: Z. Metallkde. 84 (1993) 294. [77] J. Klower: Stand der Entwicklung intermetallischer Phasen auf Basiss der Eisenaluminide Eigenschaften und Herstellung (VDM - internal report Nr. 10/92(1992)).

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[78] 1. Baker, l? Nagpal: A Review of the Flow and Fracture of FeAI. In: R. Darolia. J. J. Lewandowski, C.T. Liu et al. Structural IntermetallicsTMS,WarrendalelPA (1993) 463-473. [79] C. T. Liu. K. S. Kumar: J. Metals 45 (5) (1993) 38. (801K. S. Kutnnr: ISIJ Intl. 31 (1991) 1249. [81] K. Vedula:Strength, Ductility and Toughness of Intermetallic Matrix Composites. In: 0.Izumi (ed.) Proc. Int. Symp. Intermetallic Compounds - Structure and Mechanical Properties (JIMIS-6).The Japan Institute of Metals, Sendai (1991) 901-926. [82] R. D. Grinthnl: WADCTechnical Rep.53-190 (1958) 1. [83] WArbiter:WADCTechnical Rep.53-190 (1953) 1. [84] K. C. Chen, S. M . Alleti, J. D. Livingston: Mat. Res. SOC.Symp. Proc. 364 (1995) 1401. [85] R. L. Fleischer. R. J. Znbala: Metall.Trans.2lA (1990) 2149. [86] D. L. Antoti, D. M . Shah: Mat. Res. Soc. Symp. Proc. 288 (1993) 141. [87] D. L. Anton, D. M . Shah: Ternary Alloying of Refractory Intermetallics. In: L. A. Johnson. D. P. Pope, and J. 0.Stiegler (eds.) High Temperature Ordered Intermetallic Alloys IV.MRS. Pittsburgh (1991) 6348. 1881 M. Takeyamn, C. T Liu: Mater. Sci. Eng. A132 (1991) 61. 1891 E. Firzer: Hochsttemperaturbestandige Werkstoffe durch Silizieren von Wolfrani und Molybdan. In: E Benesovsky (ed.) Pulvermetallurgie (1. Plansee-Seminar. Reutte, Tirol. 1952). Komm.-Verlag Springer, Wien (1953) 244-258. [90] l? Schwarzkopi R. Kieffer: Refractory Hard Metals. Macmillan, New York. (1953). [91] R. Lowrie:Trans.AIME 194 (1952) 1093. [92] K. S. Kumar: Silicides: Science, Technology and Applications. In: J. H.Westbrook and R. L. Fleischer (eds) Intermetallic Compounds: Principles and Practice,Vol. 2. John Wiley & Sons. Chichester (1995) 211-236. [93] J. J. fetrovic: Mater. Sci. Eng. A 193 (1995) 31. [94] D. A. Hardwick, l? L. Martin, S. N. Patankar, J. J. Lewandowski: Processing Mircostructure Property Relationships in Polycrystalline MoSi,. In: R. Darolia. J. J. Lewandowski. C.T. Liu et al. Structural Intermetallics.TMS, WarrendalelPA (1993) 665-674. [95] 11 fetrowic: MRS Bulletin 18(7) (1993) 35. (961 M.R. Jackson, B. E Bewlay, R. G. Rowe, D. W Skelly, H. A. Lipsitt: J. Metals 18(1996) 3Y. (971 I! R. Subramanian, M. G. Mendiratta, D. M . Dimiduk: J. Metals 48 (1996)33. (981 M. Nazmy, C. Noseda, G. Sauthoff; B. Zeumer, D. Anton: Mat. Res. SOC.Symp. Proc. 364 (1995) 1333. [99] S. Mazdiyasni, D. B. Miracle: Survey of Eutectic Systems as Potential Intermetallic Matrix Composites for HighTemperature Application. In: D. L.Anton, P. L. Martin. D. B. Miracle et al. Intermetallic Matrix Composites MRS. Pittsburgh (1990) 155-162. [ 1001 G. Frommeyer, R. Rosenkranz, C. Liidecke: Z . Metallkde. 81 (1990) 307. [ 1011 P A . Beaven, 1 S. Wu, B. Dogan, C. Hartig, J. Seeger, R. Wagner:GKSS-Jahresbericht (1989) 49. [lo21 M. Es-Souni,R. Wagner,I! A. Benven, F-I! Schimansky. R. Gerling: Scripta Metall. Mater. 26 (1992) 1845. [lo31 J. S. Wu, D. Chen, l? A. Beaven, R. Wagner, J. Seeger: Microstructure and Properties of Two Phase Alloys Based on (Ti,Nb),(AI,Si) and (li,Nb)s(Si.Al)3. In: M. Kong and L. Huang (eds.) Structural Materials (Proc. C-MRS International ‘90. Beijing, Vol. 2). Elsevier Sci. Publ. Amsterdam (1991) [lo41 R. Rosenkranz, G. Frommeyer, W Smnrsly: Mater. Sci. Eng. A152 (1992) 288. [lo51 F Jansen, E. Ltcgscheider: Neue Struktur- und Beschichtungswerkstoffe auf MoSi,-Basis fur Hochtemperaturanwendungen. In: Efliienzsteigerung durch innovative Werkstofftechnik. VDI-Verlag, Dusseldorf (1995) 147-1 50. [lo61 R. Rosenkranz, G. Frommeyer: Z . Metallkde. 83 (1992) 685. [lo71 E. Ludetischeider, U. Westermann, J. Wonka, H. Meinhnrdt, H. Neisiris, R. Arnold: Investigations on Molybdenum and Titanium Disilicide as Structural Materials for Highest Temperatures. In: 0. Izumi (ed.) Proc. Int. Symp. Intermetallic Compounds - Structure and Mechanical Properties (JIMIS-6).The Japan Institute of Metals,Sendai (1991) 621425. [lo81 S. Beer, G. Frommeyer, E. Schmid, H. He1big:VDI-Berichte 1080 (1994) 89. [lo91 E. E. Schmid, K. von Oldenburg, G. Frommeyer: Z . Metallkde. Sf (1990) 809. [110] K. von Oldenburg, G. Frommeyer, E. Schmid, W Henning: Microstructures and Mechanical Properties of As Cast and Mechanically Alloyed Mg,Si-A1 Alloys. In: T. Khan and G. Effen-

1 State of Intermetallics Development

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berg (eds.) Advanced Aluminium and Magnesium Alloys. ASM International, Materials Park (1990) 477-484. [ill] G. Sauthoff: Creep Behaviour and Creep Mechanisms in Ordered Intermetallics. In: C.T. Liu, R. W. Cahn, and G. Sauthoff (eds.) Ordered Intermetallics - Physical Metallurgy and Mechanical Behaviour. Kluwer Acad. Publ. Dordrecht (1992) 525-539. [112] C.-I! Reip: Untersuchung des Verformungsverhaltens der DO,-geordneten intermetallischen Phase A1,Nb. Dr. rer. nat. Dissertation, RWTH Aachen (1991). [113] C I ! Reip, G. Sauthoffi Intermetallics 1 (1993) 159. [1141 L. Machon: Untersuchung des Verformungsverhaltens hochschmelzender hexagonaler Laves-Phasen. Dr. rer. nat. thesis. RWTH Aachen, Aachen, (1992). [115] Inco pamphlet (1982). [116] E Porz, G. Grathwohl: KfK-Nachr. I 6 (1984) 94.

Oxidation of Intermetallics by H. J. Grabke, M. Schiitze 0 WILEY-VCH Verlag GmbH, 1998

2 Research on Oxidation and Embrittlement of Intermetallic Compounds in the U.S. G. H. Meier

2.1 Introduction The general aspects of the oxidation of intermetallics are contained in a number of reviews [l-S].lhe purpose of the present paper is to describe current work on this subject being done in the U.S.,for the systems the author feels are significant, and to indicate how the results of this work enhance our understanding. Clearly, previous work in the US., Europe, and Japan has laid the foundations for the current work and, where appropriate, will be described. The fundamentals of intermetallic oxidation will be briefly described followed by discussion of work on the Ni- and Fe aluminides, Ti aluminides, and refractory metal compounds. It will become clear that work is now focussed in two general areas: (i) adding to the fundamental knowledge of oxidation mechanisms and (ii) developing approaches to allow intermetallics to reach the application stage e. g. coating, composite development, and life predictions.

2.2 Background Important features of the selective oxidation process are shown schematically in Figure 1.The slow growth rates of alumina and silica, illustrated in the plot of parabolic rate constants versus temperature at lower right, makes the formation of one of these oxides as a continuous surface layer necessary for long term oxidation protection. This requires that the protective oxide be more stable thermodynamically than the more rapidly growing oxides. The plot of standard free energy of formation as a function of temperature at lower left shows that the Ni-AI system satisfies this condition. Alumina is stable, relative to NiO, even when the activity of aluminum in the alloy is very low. However, when the A1 concentration is low the alumina forms as internal oxide precipitates and is non-protective allowing an external layer of NiO to form (illustrated in the cartoon at top). Therefore, a critical concentration of Al exists above which out-

16

G. H. Meier

Fig. 1. Important aspects of the selective oxidation process: Relative magnitudes of free energies of formation of various oxides, relative magnitudes of the parabolic rate constant for the growth of various oxides, and schematic diagrams showing internal formation of alumina in a dilute NiAl alloy and external alumina on a concentrated alloy such as NiAI.

ward Al diffusion creates a continuous, protective alumina layer. The oxidation morphology also contains transient Ni oxides which form before the alumina becomes continuous. The selective oxidation process is somewhat different for many of the intermetallics of interest. Figure 1 indicates that Ni-, Fe-, and Mo-base systems generally satisfy the thermodynamic condition that alumina or silica is the most stable oxide in the system. However, this is not the case for Nb- and Ti-base systems where the base-metal oxides are of comparable stability to alumina and silica. In these cases the activity of A1 or Si must be maintained at a high level in the alloy to maintain stability of the protective oxide. Furthermore, the potential transient oxides, e. g. TiO,, have much higher growth rates than NiO which makes development of a continuous alumina o r silica layer difficult even when the thermodynamic conditions are clearly favorable. This is particularly the case for the refractory metals Mo, W, Nb, and Ta whose oxides d o not exhibit parabolic kinetics and grow rapidly inward with linear kinetics such that the nuclei of the protective oxide are removed before they can grow together to form a continuous layer. A n additional complication arises because some intermetallic compounds, e. g. MoSi, and NbAI,, have narrow ranges of stoichometry so that a new phase forms beneath the oxide in the zone which is depleted of A1 or Si by formation of the external layer. The newly-formed phase then influences the subsequent stability of the protective oxide as well as its adherence.

2 Research on Oxidation und Embritilement of Intermetallic Compounds h the U S .

17

T

4

3

Fig. 2. Schematic diagrams showing the cracking and spalling of an oxide layer, which often occurs on cooling from the oxidation temperature and comparing the isothermal and cyclic oxidation kinetics.

In order for an alumina or silica layer to remain protective it must remain adherent to the substrate. The spalling of a protective oxide during thermal cycling and the resultant oxidation kinetics are illustrated schematically in Figure 2. The damage to the oxide is generally the result of the thermal stresses arising during cooling because of thermal expansion mismatch between the metal and oxide. Table 1 shows the linear coefficients of thermal expansion for selected materials [&lo]. The data indicate large thermal expansion mismatches between alumina and the aluminides of Ni and Fe whereas the mismatch between alumina and TiAl is small. These data arc also pertinent to the behavior of composites. For example, Sic, a commonly used strengthening phase, has large mismatch with potential matrices such as MoSi, and Ti,AI. The major interest in the intermetallics is for developing systems with higher specific strengths at high temperature. Thus the density is an important property which is the most attractive feature of the titanium aluminides which have densities less than half that of the Ni-base superalloys (Table 1).Additionally important are the ductility and fracture toughness of the internietallics.The latter is seen in Figure 3 to be inferior to that of a typical Ni-base superalloy and of the same magnitude as that of grey cast iron. The mechanical properties of intermetallics are generally further degraded, sometimes catastrophically, by exposure to oxidizing environments.

2.3 Oxidation of Selected Compounds 2.3.1 Nickel and Iron Aluminides The binary phase diagram for the Ni-A1 system is presented in Figure 4. Superposed on this diagram are the oxidation data of Pettit [ll].The crosshatched region indicates

18

G. H. Meier

Table 1. Selected physical property data CTE X106 (“C-’)

Material FeAl Fe,Al NIAl Ni,AI MarM-246 All03 TiAl Ti,AI TiAI, Sic XB, MoSi, Mo Nb

16.5 16.5 15 12.5 16 9 11 9 13 5.5 7.8 8.1 5 7.3

Density (g/cm3)

T, (“C)

-

5.56 6.70 5.86 7.65 8.44

1337 1540 1647 1390 1317

3.91 4.20 3.30

1462 1602 1352

6.30

2030

~

Fracture Toughness of Selected lntermetallics

1

IN738LC

Fig.3. Comparison of the room temperature fracture toughness of selected intermetallic compound and that for a nickel-base supcralloy and cast iron.

that NiAl should form and maintain protective alumina under all conditions while Ni,Al is a “marginal” alumina former at temperatures below 1200°C. The most successful structura! application of any intermetallic has clearly been Ni,AI as the major phase in modern Ni-base superalloys. Research on the oxidation behavior of superalloys is continuing with major emphasis on the adherence of the alumina scales.The influence of reactive elements (Y,Hf) in improving adherence and of impurities (S) in degrading it suggest that further improvement in the oxidation resistance of superalloys is still possible. Figure 5 shows the effect of reducing the sulfur content of two single crystal alloys from about IOppmw to less than 1 ppmw [12].’fie reduction of the sulfur content has clearly improved the cyclic oxidation resistance, particularly for PWA 1484 which contains 0.1 wt% Hf. Such data indicate the distance

2 Research on Oxidation and Enibrifllement of lntertnetnllic Conipoirnds in the U.S.

19

1800

F - l6O0 1400 s!

3 2

1200

g 1000

I-

800

600

Ni 10

20

30

40

Al

60 70

50

(at

80

90

Al

Oh)

Fig.4. Binary phase diagram for the Ni-A1 system showing the compositions which form protective external alumina scales, after Doychuk [5). 10 cu

B

-

w

LOW

Low s 1480

I 0

E

s 1484

0

a

gI-

I

-10

-

-20

Fig.5. Plot of mass change versus time indicating the effect of sulfur content on the cyclic oxidation kinetics of two single-crystal, nickel-base superalloys at 1100°C[12].

the monolithic intermetallics must travel to obtain oxidation resistance sufficient to make them competitive as alternative materials to the superalloys. The areas concerning monolithic intermetallics which have been studied in recent years are (i) the formation of metastable aluminas, and their transformation to stable a-alumina, (ii) the formation of interfacial voids and scale adherence and how these are influenced by reactive elements and sulfur, and (iii) accelerated oxidation at intermediate temperatures. Additionally the applications oriented areas of (iv) coatings, (v) oxidation of composites, and (vi) life predictions have received attention.

2.3.1.1 Formation of a-Alumina from Transient Aluminas This subject has been reviewed extensively by Doychak [5] and will only be treated briefly here, primarily, with regard to recent work.'l%e oxide scales formed on Ni,Al at

20

G. H. Meier

oxygen partial pressures on the order of 1 atm in the temperature range 950-1200”C consist mainly of Ni-containing transient oxides, NiO and NiAI2O4,over a layer of columnar a-alumina. Schumann [13] has recently studied the very early stages of transient oxidation of the (001) faces of Ni,AI single crystals in air at 950°C using crcsssection TEM. After 1min. oxidation simultaneous formation of an external NiO scale and internal oxidation of y’ were observed. The internal oxide particles were identified as y-Al,03, which possesses a cube-on-cube orientation with respect to the Ni matrix. After 6 min. oxidation a continuous y-Al,O, had formed between the internal oxidation zone (IOZ) and the y’ single crystal. Oxidation for 30 min. resulted in a microstructure similar to the 6min. oxidation but the Ni in the two-phase zone was oxidized to NiO. Oxidation for 50 h resulted in a scale consisting of an outer layer of NiO, an intermediate layer of NiAI,O,, and an inner layer of y-Al,O, in which a-A1203grains had nucleated at the alloy/oxide interface.The spinel was presumed to have formed by a solid state reaction between NiO and y-Al,O,. A crystallographic orientation relationship was found between the a and y alumina whereby (OOOl)[lTOO], I1 (11I)[lTO],. i. e. close packed planes and close packed directions of c1 arc parallel to close packed planes and directions in y. Pint and Hobbs [14] studied the oxidation of yttria (2~01%)-dispersed Ni,AI at 1000 and 1200°C and did not report transient alumina formation for oxidation times as short as 1hour. The difference in the results of [13] and [14] may stem from slightly different oxidation temperatures or stabilization of the y-Al,O, on the specific orientation or the single crystal because of surface energy considerations. The transient oxidation of Fe,Al has not been extensively studied, however, the formation of 8-A120, has been invoked to explain anomalies in the oxidation rate at 900°C [ 151. The oxidation of NiAl is somewhat unique in that, at temperatures of 1000°C and above, there are negligible amounts of Ni-containing transient oxides. The transient oxides are all metastable phases of AI,O, (y, 6, and/or 0) [ S ] . The transition of these metastable phases to the stable a-Al,O, results in significant decreases in the scale growth rate and a “ridged” oxide morphology which is distinct from the columnar morphology observed on Ni,AI [5].The transition aluminas have been shown to grow primarily by outward migration of cations while a grows primarily by inward transport of oxygen [5].The effect of oxidation time and temperature on the phases present in the scales has been studied by several authors. Rybicki and Smialek [16] identified t) as the transient oxide on Zr-doped NiAl and found that the transition to a occurred at longer times at lower oxidation temperatures e. g. scales consisted entirely of 8 after 100hours at 800°C while it transformed to a in about 8hours at 1000°C. Only a was observed at 1100°C and 1200°C.Pint and Hobbs [17] observed only a on undoped NiA1 after 160 seconds at 1500°C. Brumm and Grabke [IS] observed two transforniations at 900°C for undoped NiA1.The scale consisted initially of y which transformed to 8 after approximately 10hoursThe transformation of 8 to a occurred at much longer times but accelerated as the oxidation temperature was increased. The transition to a on NiAl has been reported to initiate at the scale/gas interface 1191 which is in contrast to the observations on Ni,AI where it is reported to initiate at the scale/alloy interface [13].The scale transformations have also been found to be sensitive to the presence of third elements in the alloy.Additions of Cr accelerate the transformation to a [18] by a proposed mechanism involving transient Cr,O,, which is isostructural with a , providing nucleation sites. This results in a finer-grained a which grows somewhat faster

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than the a formed on binary NiA1.A study of oxide dispersed NiAl found that the 0 to transformation was slowed by Y, Zr, La, and Hf and accelerated by Ti [20]. The proposed explanation of these results is that the transformation is slowed by large ions which can enter the more open lattices of the transition aluminas.This result is in contrast to the observation that no transition aluminas were observed on Y,O,-dispersed Ni,A1 [ 141. These phenomena have not been well studied for FeAl but limited data suggest a similar set of transformations in the aluminas formed on FeAJ [21]. (Y

2.3.1.2 The Formation of Interfacial Voids and Scale Adherence and how these are Influenced by Reactive Elements and Sulfur The integrity of the alloy/oxide interface and the adherence of the alumina scale to the alloy are critical issues for the application of aluminides or aluminide c0atings.A common feature of the oxidation of the Ni- and Fe-aluminides is the formation of voids at the alloy/oxide interface.'I'his is illustrated, schematically, in Figure 6 for NiAl[3]. The voids result from the diffusion of Ni into the alloy, as A1 is consumed to grow the oxide, which brings vacancies to the interface where they condense. Brumm and Grabke [22] have shown that void formation is negligible on Al-rich NiAl where A1 diffusion to the interface predominates as compared to Ni-rich NiAl where Ni-diffusion away from the interface predominates. The presence of the interface voids has long been known to be a source of poor adherence [23].'I'he interface voids are greatly suppressed by additions of reactive elements [24]. This generally results in improved alumina adherence and better cyclic oxidation resistance. The effect of reactive element additions as oxide dispersions on the cyclic oxidation resistance of NiAl at 1200°C is illus-

Fig. 6. Schematic diagrams demonstrating the fluxes of A1 and Ni caused by oxidation of NiAI, leading to vacancy formation and condensation as voids beneath the oxide scale.

G. H.Meier

22 7 ,

0.1 IZ r

I I

0

-10

,

0

Ref.: B.Pint and L. Hobbs. M R S Symp. Proc.. Vol. 364. 1995 I

I

2

4

I

6

I

I

8 10 12 Cycles (2 hwdcycle)

14

I

16

1

18

1

Fig.7. Cyclic oxidation kinetics of NiAI, containing various oxide dispersions, in air at 1200°C [XI.

Mass Change/Area (rng/cm2) ......... ..

-...a-,.-*.y,--

-

_-0

200

400

600

t (hours)

800

lo00

1200

Fig.8. Effect of sulfur content on the cyclic oxidation behavior of NiAI.The as-received material contained 20ppmw S and the sulfur content of the hydrogen annealed material was calculated to be less than 1ppmw.

trated in Figure 7 [25].Lowering the sulfur content of NiAl also suppresses interface void formation [26] and, as illustrated in Figure 8, can greatly improve scale adherence. The addition of reactive elements has been shown to produce similar improvements in the scale adherence to Ni,AI [14],Fe,Al[15], and FeAl [21].

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A number of studies have introduced the reactive element into the aluminides by ion implantation. The author considers this approach to be of limited value since, as discussed by Pint and IIobOs [27], the high local concentration of reactive element can stabilize the transition aluminas and, at high temperatures, has a short-lived effect. More importantly. from a fundamental standpoint, the implantation process can have a profound effect on the nature of the exposed surface. For example, Schumann [28] has shown that Y implantation into single-crystal NiAl results in a 45 nm thick, finegrained crystalline region which is disordered.

2.3.1.3 Accelerated Oxidation at Intermediate Temperatures A number of intermetallic compounds, which form protective alumina or silica scales at high temperature, undergo accelerated degradation at intermediate temperatures. This subject has been recently reviewed [29].The observations actually involve several different, but related, phenomena which may be subdivided into “accelerated oxidation”, “internal oxidation”, “intergranular oxidation” and “disintegration”. The following definitions will be used throughout this paper.

Protective Oxidation - formation of a continuous alumina or silica surface film with no internal oxidation and minimal penetration of oxygen into the substrate. The overall oxidation rate is determined by transport through this film. Accelereated Oxidation - the alumina or silica is not continuous and significant amounts of the other component(s) of the intermetallic are present in the surface film. The overall oxidation rate is substantially faster than that for the growth of alumina or silica. Internal Oxidation - precipitation of oxides rich in A1 or Si within the intermetallic. Intergranular Oxidation - special case of internal oxidation in which oxides form along grain boundaries within the intermetallic. Pesting - disintegration (fracture) of the intermetallic into smaller particles at the oxidation temperature. The occurrence of protective oxidation precludes the other phenomena from occurring as long as the oxide is not damaged. However, in the absence of protective oxidation, any or all of the other phenomena may occur together. In the case of NiAl this degradation occurs in the temperature range 700 to 1000°C and can take the form of intergranular oxidation at reduced oxygen partial pressures in the range lo-” to 10-’atm. and internal oxidation in the oxygen pressure range to 10 22atm.[30].A particularly complex form of degradation occurs when exposures are carried out in evacuated silica ampoules containing buffer mixtures, such as Cu/Cu,O. This attack, termed “pocks” [30] is illustrated in Figure 9 [31].The oxidation morphology consists of an outer zone of silicides and silicates (Si is the result of SiO vapor transport from the capsule walls) over an internal oxidation zone of alumina in virtually pure Ni. A substantial Al-depleted zone, which has transformed to Ni,Al, is observed beneath the internal oxidation zone. The detailed mechanism for this process is not completely understood but is known to require conditions that prevent the formation of a protective alumina layer (temperatures which favor formation of transition aluminas and possi-

24

G. H. Meier

Porous Al-oxides

Fig.9. Example of pock formation of NiAl oxidized at llOOK in a silica ampoule containing a mixture of Cu and Cu,O powders.

bly contaminants such as S) and a process which rapidly consumes Al (the internal oxidation). Pocks do not form when the buffer mixture is thoroughly dried prior to specimen exposure [30,31], when the NiAl is alloyed with Cr [30], or for aluminide coatings on Ni-base superalloys [31]. These observations are possibily all the result of conditions which favor the rapid establishment of a-alumina rather than transition aluminas. Similar studies have, apparently, not been performed for FeAI. 2.3.1.4 Coatings The aluminizing of Ni-base alloys via the pack-cementation process has been a conimercially viable process for many years [32]. The typical coating morphologies for the

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25

Fig. 10. Schematic diagram showing the as-deposited structurc o f a “low activity” (left) and a “high activity” (right) diffusion aluniinide coating on a nickel-base superalloy. The high activity coating would receive a heat treatment to convert the Ni,AI, to NiAI.

so-called “low activity” and “high activity” processes are illustrated schematically in Figure 10.These coatings oxidize in a similar manner to bulk NiAl except that,superposed on the oxidation process. the coating loses A1 by interdiffusion with the substrate. In recent years the plating of Pt on the substrate prior to aluminizing has been used to produce “platinum-aluminide” coatings [33]. This type of coating, shown schematically in Figure 11, has resulted in improved cyclic oxidation rcsistance and reduced interdiffusion with the substrate. Currently, the improvement in hot corrosion resistance imparted to Ni-AI alloys by Cr additions has resulted in research efforts aimed at codepositing Cr and A1 in aluminide coatings. The simulatenous codeposition of Al and Cr via the pack cemcntation process, using pure elemental powders, is, however, difficult. The large difference in the thermodynamic stabilities of the A1 and Cr halides causes Al-halide species to predominate in thc pack atmosphere [34]. However, by employing binary chrorniumaluminum (Cr-A1) master alloys, the high relative A1 halide vapor pressures can be moderated.This is the result of the fact that chromium-rich master alloys exhibit negative deviations from idcality and the activity of A1 in the master alloy can be reduced by several orders of magnitude. The reduced thermodynamic activity of A1 results in generation of lower vapor pressurcs for the otherwise favored halide species (e. g., AICI, AICI, etc.). Therefore comparable A1 and Cr vapor pressures result. Thus, provided a suitable activator and binary master alloy is chosen, the codeposition of Cr and Al into Ni base materials is possible. Rupp and Bianco [35-371 have used this approach to form coatings containing as much as 13 at% Cr in P-NiAI on Ni-base superalloys and have shown them to have greater hot corrosion resistance at 900°C than aluminide coatings without the Cr modification [37]. Da Costa et al. [38, 391 have achieved Cr concentrations as high as 40 at%, using Cr-rich master alloys and multiple activators (NaCl+NH,CI). Stinner et al. 1401 arc currently studying how the pack

26

G. H. Meier

variables,such as the amount and composition of the source alloy and the activator. a f fect the composition and morphology of Cr-aluminide coatings on Ni-base alloys.

2.3.1.5 Composites The lack of toughness at low temperatures and low creep strength a t high temperatures suffered by many monolithic intermetallics has led to effords to use them in composites to achieve the required toughness and strength. The presence of the second phase not only affects the mechanical properties but also the oxidation behavior, usually, in a detrimental fashion.This has occurred even for NiAI. Doychuk et al. [41] found that incorporation of alumina fibers in a NiAl matrix resulted in oxidation along the fibers, in 1200 and 1300°C exposures, and cracking of the matrix. Perkins [42] found, in air exposures at 800 and llOO°C,that addition of 10~ 01%TiB, to NiAl resulted in an incubation period followed by rapid growth of TiO,. The addition of 20 ~0 1 %TiB, resulted in rapid oxidation from the beginning of the exposure. One strengthening phase which, apparently, does not degrade the oxidation resistance of NiAl is AIN [43].The oxidation of AIN-dispersed NiAl was similar to undoped. monolithic NiAl and, when small amounts of Y,O, were added, showed cyclic oxidation resistance comparable to Zr-doped NiAI.

2.3.1.6 Life Prediction The possibility of using intermetallics as structural materials has increased interest in predicting how long a component will last under a given set of exposure conditions. Nesbitt and coworkers [44-46] have treated this problem for the case of NiA1.The factor which controls life is the loss of A1 in forming alumina on the surface and through spallation during thermal cycling. The procedure involves identifying a failure criterion which can be a critical amount of surface recession, reaching a predetermined minimum Al content in the alloy, or the appearance of spinels in the oxide scale. After selecting the failure criterion the rate of Al loss is needed. This can be obtained by isothermal measurements of the oxide growth rate coupled with measurements of the amount of oxide which spalls on each cycle during cyclic oxidation. However. the measurement of the amount of spalled oxide is difficult, particularly at low temperatures. Therefore, an approximate procedure for estimating the A1 loss, which is good at all but very short times, is to use the linear portion of a cyclic oxidation plot such as that in Figure 2.The time dependence of the A1 content in the specimen can be calculated by solving the diffusion equations in the alloy. A simpler, approximate method may be used for compounds, such as NiAI, which have high interdiffusion coefficients by assuming the A1 content decreases uniformly. Either way the time at which the Al content falls below a certain value can be calculated. If desired the A1 loss can be converted to surface recession using the partial molar volumes for A1 in NiAI. Figure 12 shows the results of predicted life for various intermetallics as a function of specimen thickness using a failure criterion of a critical surface recession of 10% of the original specimen

2 Researcli on Oxidation and Etnbritrlenzenr of Intenneiallic Compounds in rhe U.S.

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1200 "C NiAl-O.1Zr

0.10

0.15 0.20 Thickness (crn)

0.25

_I I

0.1 0

NiAI-0.1 Zr 1400 OC

0.1 5 0.20 Thickness (cm)

I

0.25

Fig. 12. Predicted oxidation lifetime for various intermetallic compounds as a function of specimen thickness using the time to produce a surface recession equivalcnt to 10 % of the original specimen thickness [45].

thickness [45].The longer lives predicted for NiAl-O.lZr, relative to NiAI, is the result of greatly reduced oxide spallation from the Zr-doped alloy.

2.3.2 Titanium Aluminides Alloys in the Ti-A1 system are of interest for high temperature systems such as aircraft engines because they have low density and maintain strength at high temperature. However, their resistance to oxidation and interstitial embrittlernent is a concern. Those alloys which form alumina scalcs have excellent resistance to surface re-

G. H. Meier

28

cession while those which form titania-rich scales oxidize at much higher rates. 'fie modification of the microstructure immediately beneath the oxide layer, resulting from interstitial dissolution and/or selective removal of one or more alloy component. is of particular concern since it can cause a substantial loss of mechanical properties. The discussion in this section covers (i) the thermodynamics of titanium-aluminide oxidation, (ii) oxidation of specific compounds. (iii) embrittlement of titanium aluminides, (iv) the effect of complex environments on the oxidation of titanium aluminides, (v) coatings on titanium aluminides, and (vi) oxidation of titanium aluniinidematrix composites.

2.3.2.1 Thermodynamics A n important aspect of the oxidation of Ti-aluminides, compared to the aluminides of Ni and Fc, is the small difference in standard free energy of formation between alumina and the oxides of titanium, Fig. 1,which is accentuated by a negative deviation from ideal solution behavior in the Ti-A1 system. The aluniinum activity is much smaller than unity in Ti,AI and TiAI. In fact, combining the activties with standard free energy data for the oxides indicates that 'Ti0 is more stable in contact with the alloy than is AI,O, for atom fractions of A1 much less than about 0.5 [47, 481. Thus, AI,O, is unstable in contact with binary a2 and is only marginally more stable than T i 0 in contact with y-TiAl. Figure 13 presents the Ti-AI phasc diagram with the compositions which have been abserved to form continuous alumina scales indicated by crosshatching.

p

1500

v

5 6 !?

I-

loo0

500

Ti

XAl

Al

Fig. 13. Binary phase diagram for thc Ti-A1 system showing regions (crosshatched) where continuous alumina films have been observed to form.

2 Reserirch on Oxidntion trnci Enibritrlmimt of Inrernietcillic Compounds in the U.S.

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2.3.2.2 Oxidation of Specific Compounds Oxidation of 5,AI (a2) The oxidation of Ti,AI alloys would not be expected, in light of the above thermodynamic considerations, to form continuous alumina scalcs. Instead they form mixed rutile-alumina scales [49].The oxidation kinetics of Ti,AI between 600 and 950°C are reported to be essentially those expected for rutile growth [SO, 511. These kinetics result from the development of a complex oxide layer which contains continuous paths of TiO, through which rapid transport occurs. Typical oxide scalcs are shown schematically in Figure 14. The oxidation of Ti,Al is generally more rapid in oxygen than in air, as indicated in the 900°C oxidation data in Figure 14 [ 5 2 , 5 3 ] .This appears to be the result of a layer of TIN, which forms at the scaleialloy interface during air exposures, providing a diffusion barrier. Alloying of Ti,A1 with P-stabilizing elements, particularly Nb, at lcvcls less than about 10at% reduces the oxidation rate [53,54,55] as indicated in Figurc 14. Generally more complete nitride layers form on the Nb-containing alloys which accounts for part of the rate reduction. Wallace et al. [53] found a layer of y-'IiAl beneath a layer of TiN on ?'i-24AI-l1Nb exposed in air at temperatures between 700 and 1000°C. How-

Fig.14. Oxidation rates for a2 and a,+Nb alloys in air and oxygen at 900°C and schematic diagrams of the scales formed under thc various conditions.

30

G. H. Meier

ever, the effect of Nb also occurs in pure oxygen which indicates an additional effect. Niobium has been detected in the scales [52,55] leading to the proposal that the Nb ions produce a doping effect which decreases the concentrations of oxide ion vacancies and/or titanium ion interstitials in the rutile lattice [52]. The scales developed on cx2 alloys containing multiple additions of P-stabilizers have been described by Wallace et al. [56] and Schaeffer [57]. An additional aspect of the oxidation of Ti,AI alloys is dissolution of oxygen into the alloy at the scale/alloy interface. The embrittlement associated with this phenomenon can be more damaging to the mechanical properties than the surface recession caused by scale formation in the temperature range where Ti,A1 will likely be used (< 700°C) [58].This subject will be discussed in a separate section. Oxidation of orthorhombic alloys There are relatively few oxidation data available on alloys in the Ti-Al-Nb system with significant volume fractions of the orthorhombic phase. Howevcr, data are becoming available which indicate that the amount of Nb necessary to stabilize the orthorhvmbic phase results in more rapid oxidation than the a2alloys because of the formation of discrete Nb-oxides in the reaction products [59].The effect of N b on the parabolic rate constant for the oxidation of Ti-25 at% A1 at 800°C [52,54] is presented at right in Figure 15. The rate constant decreases with increasing Nb content up to the 5-10% range but increase with further additions. Additions of ? a have been found to reduce the oxidation rate slightly at a given Nb content [59].The rate at 700°C has been reported to decrease with Nb additions up to 10at% [55]. Additionally, alloys with Nb contents typical of the “orthorhombic alloys” undergo breakaway oxidation following an initial parabolic period as seen at left in Figurc 15.This phenomenon which is more prominent at the higher temperatures has been shown, using acoustic emission mea-

Oxidation Kinetics of Ti-22AI-23Nb at 500 900T in Air

-

Variation of Initial Parabolic Oxidation Rate at

800’C in Air of (Ti -25AI) + Nb

4 d C

-10

’ I

OTi-25AI b)

a)

I

21Nb-2Ta

20 OhNb ( a t . l )

I

40

Time (hours)

Fig. 15. Oxidation kinetics for Ti-22A1-23Nb“orthorhombic alloys” in air at temperatures in the range 500-90O0C(left) and the effect of Nb content on the parabolic rate constant for Ti-25 at% A1 alloys at 800°C (right).

2 Research on Oxidation rrnd Etnhrittlenient of Intermetallic Compounds in the U S .

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surements, to result from oxide cracking as the scale is enriched in Nb-oxides [59]. During oxidation the alloy microstructure is modified beneath the oxide layer causing embritt1ement.This will be discussed in a subsequent section. Oxidation of y-TiAl Choridhitry et al., in a classic paper, [50]studied the oxidation ofTiAl (50at% Al) in 0, and air over the temperature range 800-1200°C. In 0, cast XAl which was abraded through 120 grit S i c formed alumina and exhibited k, values = 10-9g2cm-4hr-’ at 950 “C but polished specimens formed Ti0,-rich scales and exhibited k, values of about 10- to 10-6g2cm-4hr-’. Extruded TiAl formed alumina scales regardlcss of surface preparation. Choudhury et al. explained the effects of extrusion in terms of the absence of macroinhomogeneities which existed in the cast structure. The “surface finish effect”, observed for the cast material was presumed to result from coarse grinding homogenizing the alloy at the surface. A %,A1 layer was reported to form between the oxidc and the alloy for all the exposure. Oxidation behavior in air at 950°C was independent of specimen preparation or fabrication method with titania-forming kinetics (k, about 10-5-10-6g2cm-4hr-’) observed in all cases. The scales were similar to those formed on polished specimens in 0, at 950 “C. Experiments to determine the species responsible for the difference between exposures in 0, and air indicated that CO, CO, and €1,0 impurities were not responsible nor was the difference in poz. It was, therefore, concluded that the increased rate of oxidation in air was a “nitrogen effect” although no N,-containing phases were identified in the scale or substratc. The oxidation of ?>A1also exhibited a “temperature effect” in that kinetics for all alloys at 1100 and 1200”C,regardless of surface preparation or exposure atmosphere, indicated XO, formation. Internal oxidation was also observed at the two higher temperatures. There has recently been a substantial amount of work on the oxidation of TiAl with seemingly contradictory results. Unfortunately, much of the confusion has arisen because the authors are unaware that temperature, surface preparation, and composition of the exposure atmosphere all play a critical role in determining the nature of the rcaction products and the kinetics, as illustrated by Choudhury et al. [50] more than twenty years ago. The “temperature effect” is illustrated for TiAl with a ground (600 grit) surface [60] in Figure 16. For specimens exposed in pure 0, at temperatures below about 800°C a continuous alumina layer, shown schematically at left in Figure 17a, covers the entire specimen surface. At temperatures between 800 and lo00 “C small nodules, containing both alumina and titania, form at various locations, as seen at right in Figure 17a. The area density of the nodules increases as the oxidation temperature is increased but the nodules do not grow with time after their initial formation. At temperatures above 1000°C the nodules grow continuously and develop a continuous mixed oxide layer, similar to that which forms on a2with greatly accelerated kinetics, Fig. 16. Mendiratta and Choudhury [51] reported that varying the Al-content of TiAl(50,53, and 54 at%) did not affect the oxidation behavior. Appalonia et al. [61,62], however found that increasing A1 content, in the range 50 to 56 atyo,increased the temperature at which the accelerated kinetics began. The “surface finish effect” for oxidation temperatures below 1000°C is illustrated in the mass gain vs. time plot in Figure 18 for Ti-50at% A1 oxidized at 900”C.The ground

G. H. Meier

32

A

20,

700

750

800

850 900 Temperature (“C)

950

1000

1 50

Fig. 16. Mass change versus temperature for TiAl oxidized for 58 hours in air, oxygen, and oxygen which is contaminated with a small amount of nitrogen.Air and O?(N?)curves are for 54 at% A1 and 0, curvc is for 52 at% Al.

(600 grit) surface results in very slow kinetics in 0, as the result of continuous alumina formation as shown schematically in Figure 17a.The polished (1 pm diamond) results in greatly accelerated kinetics as the result of the formation of a continuous layer of a mixed oxide, Figure 17b. Part of the beneficial effect of grinding may be in homogenizing the surface, as proposed originally by Choudhury et al. [50].Indeed, the distribution of the minor a, phase, generally present in TiAI, has been shown to influence the composition of the external scale [63,64]. However, an additional important effect is the recrystallization of the ground surface as illustrated in the transmission electron micrographs in Figure 18 [64]. The deformed layer, produced by grinding, recrystallizes during the initial oxidation exposure to produce y grains with a grain size on the order of 1 pm from a starting y grain size on the order of 100 pm. No recrystallized layer forms on the polished surface. It has been suggested [64] that enhanced Al diffusion in the fine-grained layers promotes the fo5mation of continuous alumina in the same manner as that reported by Giggins a$# Pettit [65] for the development of chromia films on Ni-Cr alloys. The “nitrogen effect” is also illustrated in Figure 16 and the mass gain vs. time plot in Figure 18.The kinetics in air are essentially independent of surface perparation and are much faster than those for a ground surface exposed in pure 02.Thescales formed in air for ground and polished surfaces are shown schcmatically in Figures 17c and 17d, respectively. Similar observations have been made by a number of investigators [50,66,67]. In a study to clarify the “nitrogen effect” [60] protective alumina scales were observed to form on TiAl (52 at% Al), exposed in 0,,up to 1000°C. However, the same exposures conducted in air resulted in the formation of Ti0,-rich scales which grew at rates orders of magnitude faster than pure alumina scales and even trace amounts of N, influenced the oxidation morphology, Fig. 16.The rate of oxidation increased continually as increasing amounts of N, were added to pure 0, at 900”C.The addition of

2 Research oti Oxidation and Etiibrittlemetit of Intermetallic Compounds in the US.

33

Mixed oxide nodule

1

Recrystallized ylayrr

1

a) Oxygen exposure, 600 grit surface preparation

I

Recivstallized Ylaver

c) Air exposure, 600 grit surface preparation

b) Oxygen exposure, 1 micron surface preparation

I

d) Air exposure, 1 micron surface preparation

Fig. 17.Schematic diagrams of the scales formed on y-TiAl in the temperature range 800-900°C in air or oxygen and with two different surface finishes.

10 % N, to 0, resulted in the formation of nodules of intermixed TiO, and AI,O, interspersed with thin areas of protective AI,O, which would cover the entire surface in the absence of N,. The area density of these nodules increased as the concentration of N, increased, until the surface was completely covered with the mixed oxides when the gas contained 90 % N,. One effect of N, appears to involve the nucleation and initial growth of the scale since preoxidation in a nitrogen-free gas develops an alumina scale which remains protective during subsequent exposures in air [60] even under cyclic conditions [68]. The influence of N, on the initial scale development has been investigated using Auger spectra collected while sputtering through scales formed for short times in 0, and air [60].The scales on specimens which were heated to 900°C and immediately cooled were 2000-3000p\ thick with AES profiles which were consistet with intermixed transient oxides of A1 and Ti. However, specimens exposed in air were found to have a N,rich layer at the scale/alloy interface.This layer was postulated to contain TiN but was

34

G. H. Meier

Time (h)

600 grit

1 Pm

Fig.18. Oxidation rates for y-TiAl at 900°C in air and oxygen with two surface finishes and crosssection TEM micrographs of specimens with two different surface finishes exposed in oxygen for 1 hour showing the recrystallized layer which formes on the ground specimen and the cubic “Xphase” which forms in the Al-depleted zone on both specimens.

not unequivocally identified. The scales on specimens which were held in O2at 900°C for 15 minutes were quite different from those in air.The AES profiles indicated that a continuous layer of aluminum oxide was forming below the mixed transient oxide in 0, while in air the scale consisted only of intermixed Ti- and Al-oxides and was almost ten times as thick as that in 0,. It was suggested that a major effect of the N, was in forming a nitride layer which prevented the alumina from developing continuity. Nitrogen has also been observed at the scale/alloy interface in SNMS profiles through scales formed in air for 1.5h on Ti-5OAl at 800°C [69]. The proposal that the “nitrogen effect” involved formation of nitrides which prevented the alumina from bqoming continuous has been verified by elemental spectroscopic imaging (ESI) in a transmission electron microscope 1701. An ESI map for the elements Ti, Al, 0, and N from a cross-section of the scale/alloy interface of Ti50 at% A1 oxidized in air for one hour at 900°C is presented in Figure 19.These maps indicate that the alumina is broken up by islands of TiN. The micrographs in Figure 18 reveal an intermcdiate layer between the oxide and the y phase.?his zone is depleted in A1 and was originally thought to be a2(501.However, an

2 Research on Oxidation and Embrirtlernent of Intermetallic Compounds in the U.S.

35

Fig. 19. Maps formed using electron spectroscopic imaging in a transmission electron microscope from the cross-section of a specimen of y-TiAl oxidized for 1 hour in air at 900°C showing intermixed regions of TiN and alumina.

important observation was made by Dowling and Donlon in 1992 [71] when they identified a cubic phase with a lattice parameter of 0.69nm and a’WAl ratio of approximately 24. These observations were confirmed [70] and the cubic phase was found to have 432 point group symmetry [72]. Extensive TEM investigation of this phase using convergent beam electron diffraction (CBED) [73] showed it to belong to one of two space groups, P432 or P4,32, and EDS analysis indicated an approximate composition of 57at% Ti33at%AI-lOat %O [74]. Increased exposure times and/or higher oxidation temperatures result in the Al-depleted zone becoming two-phase as the result of nucleation of a2 at the interface between the cubic phase and the parent y 1741.This observation confirms earlier suggestions of a two-phase depletion layer based on WDS measurements [66], Auger electron spectroscopy [75], and X-ray diffraction [76]. The influence of the depleted zone on the oxidation behavior of y alloys is not yet clear, however, it is crucial to the mechanical properties, as will be discussed in a subsequent section.

G. H. Meiet

36

The effects of alloying elements on the oxidation behavior of y-TiAl have been tabulated by Doychuk [5] with regard to their positive or negative effects and the subject has been recently addressed by Rahmel et al. [77]. The early work by Choirdhirry et al. [50] indicated that Nb and W additions promoted the formation of continuous alumina on y at 950°C independent of surface preparation or exposure environment (air or 0,)while Hf had a minimal effect and Y and Ga were detrimental. Unpublished activity data are quoted as indicating that Nb and W additions increase the a,,/a, ratio and promote alumina stability. Unfortunately, comparison of much of the more recent work is difficult because experiments have been performed at different temperatures and in different atmospheres and various surface finishes (which are sometimes not clearly specified) have been used. However, the following generalizations can be made. Additions of Nb [SO, 78-81], W [SO, 78,791, Ta [Sl], and Si [79,82] are generally beneficial. Additions of V and Mo promote alumina formation at temperatures above 900°C [78,79] but actually accelerate scale growth at the lower temperatures (= 800°C) where y alloys are of practical interest. Additions of Mn are generally detrimental [83]. Somewhat surprisingly, small additions of P and CI are reported to improve oxidation resistance [79]. Chromium is an alloying element of fundamental and practical significance. Small additions of Cr to y ( 5 5 at%) accelerates the growth of the mixed oxide by doping the TiO, to increase the concentration of oxide ion vacancies and/or titanium ion interstitials [84]. However, at higher concentrations the Cr promotes the formation of a continuous external layer of alumina [82].These effects are illustrated for 900°C oxidation in air in Figure 20 [31]. The micrographs in Figure 21 show the different scale morphologies produced by the different Cr additions. The binary alloy formed a mixed oxide layer approximately 10 p,m thick while the layer on the 4 Cr alloy is more than twice as thick. The scale on the 12Cr alloy is continuous a-alumina and is less than 1p,m thick (Note this micrograph is a taper section which makes the layer appear thicker.) The oxide map in Figure 22 [82] indicates the combinations of Cr and Al contents that result in the formation of external alumina in air. The alloying levels required increase as the oxidation temperature is decreased. The mechanism whereby 12 -,?

rj . z 57

v

108-

d

{ G!$

r2

2

4-

M ."

$ d

Ti-48A1 2-

0

Ti-49AI- 12Cr

0

I

I

I

I

I

I

Time (Hours)

I

I

I

Fig.20. Effect of Cr content on the oxidation rate ofTi-48 at% Al at 900°C in air.

2 Research on Oxidation and Embrilllrrnerii of Intermetallic Compounds in the (J.S.

Fig.21. Cross-section SEM micrographs of the thrce alloys from Figure 20. (a) X-48AI. b) Ti48A1-4Cr,and c) Ti-48A1-12Cr)Note that the micrograph of the 12Cr alloy is from a taper section which magnifies the scale thickness by about a factor of 10.

37

38

G. H. Meier

TiCrAl

Ti

-

1000°C Cyclic

8"

I

0

A1

1

03

0.8

0.7

0.6

05

0.4

03

0.2

0.1

0

Cr

-

500 1000 1500 2000 2500 3000 3 Time (Hours)

IC

Fig.22. Oxide map showing the regions of Cr and A1 contents which form external alumina scales in air at various temperatures (left) and cyclic oxidation kinetics for Ti-Cr-AI alloys over a range of compositions for exposures in air at 1OOO"C(right).

Cr promotes alumina formation is still unclear. The compositions which form alumina all fall outside of the single-phase y field [31,85,86]. See, for example, Figure 21c. Figure 23 from the work of Brudy et al. [85]indicates the stable phases in the Ti-Cr-A1 system at 1000°C.The alloys along the boundary for alumina formation, except those with very high Cr contents, consist of two phases, y and a ternary Laves phase TiCrAl. One likely effect the Cr has on forming A1 is by increasing the activity ratio a,&. Brudy et al. [87] observed that binary alloys with A1 contents between 49 and 53 at% reacted with alumina in diffusion couples at 1000°C. However, a Ti-42A1-27Cr alloy did not react under the same conditions. An additional effect is that the "nitrogen effect" apparently does not operate in the presence of sufficient amounts of Cr even though nitrides form beneath the oxide during the early stages of scale formation [31]. The cyclic oxidation plot in Figure 22 indicates an additional feature of the TiCrAl alloys, the excellent adherence of the alumina layers formed on them. The alloys for which the data are plotted include a wide range of compositions, as indicated by the 1073K Alumina Boundary

A1

*O Cr

Fig.23. Section of the Ti-Cr-A1 phase diagram at 1000°C indicating the equilibrium phases for the alloy compositions which form continuous alumina scales (after Bra& et a].).

2 Research on Oxidation arid Embrittlement of Intermetallic Compounds in the U.S.

39

stars on the oxide map. The spalling resistance of the alumina on these alloys is consistent with the good thermal expansion match between y-TiAl and a-alumina (Table 1). Oxidation of TiAI, The oxidation of TiAl, results in the exclusive formation of external alumina [62] and the parabolic rate constants are essentially the same as those observed for NiAl [88]. However,TiAI, is a line compound and, as such, is difficult to prepare by melting and casting as a single phase material. It has been found [88] that excess Al results in rapid transient oxidation prior to the attainment of the slow steady-state rate.

2.3.2.3 Embrittlement of Titanium Aluminides The effects on mechanical properties of exposure to oxidizing atmosphere constitutes what is currently the most important area of study regarding the use of intermetallics as high temperature structural materials. The creep resistance and ductility of a2 at high temperature are much lower in air than in vacuum [58] and room temperature fracture toughness is severely degraded by exposure to oxygen at elevated temperatures [58,89-911. Ward et al. [89] found that embrittlement of Ti-25Al-lONb-3Mo-lV during tensile testing in air at 550 and 650°C was strain-rate sensitive and proposed a dynamic process whereby oxygen embrittled the alloy in the region of a propagating crack. Rakowski et al. [52] found that room temperature ductility of Ti-21Al-llNb was severely degraded by exposures to oxidizing atmospheres at temperatures from 600 to 900°C and that the embrittlement was sensitive to the composition of the exposure environment,Table 2. Oxygen caused severe embrittlement whereas high-purity nitroTable2. Embrittlement of’I’i-21Al-llNb Exposure Conditions

Elongation (YO) 24 h - 900°C 20 h - 900°C Purged System Evacuate and Backfill

Argon < 1ppm H,O < 10 ppm H,O

3.60

Nitrogen < 1 ppm H,O < 32 ppm H,O

1.55

Oxygen < 3 ppm H,O < 50 ppm H,O

3.06

Air < 3 pprn H,O < 50 ppm H,O

6.87

Air - 46 YO H,O Oxygen - 37 YoH,O Hz - 39 Yo HZO

1.30 1.91 1.11

> 20.3

17.0

3.60

2.70

40

G. H. Mrier Cracks

-

6 -6-

&*Mixed

TI, Nb. Al Oxide

2

10

0

500 c T c 800 C in Air

200

400

600 800 1000

Time (hours)

Fig.24. Schematic diagram of the scale and interstitial affected zone (IAZ) which forms on “orthorhombic alloys” in the temperature range 500-800°C (left) and the time dependence of the IAZ thickness (right).

gen or argon did not. However, the presence of water vapor in the nitrogen or argon caused embrittlement and analysis of the fracture behavior suggested a synergistic effect between hydrogen and oxygen. The detailed mechanisms of these phenomena have not been studied and the effect of alloying additions have not been extensively investigated. The orthorhombic alloys have been reported to be less susceptible to environmental embrittlement than the a2alloys [92], however, more recent work [59,93] has indicated that this is not generally the case.The microstructure of the alloy is substantially modified by the oxidation process, as indicated schematically in Figure 24. An “interstitial affected zone” (IAZ) in which the volume fraction of (x2 is increased, forms in the alloy below the scale. A very small angular phase also develops in this

Variation of RT Ductility After 100 hr. Preoxidation in Air

0

400

(degrees

800

Fig. 25. Room temperature ductility (measured in 3-point bending) of “orthorhombic alloys” after 100 hours exposure in air at various temperatures,

2 Resetrrch on Oxidation nnd Ernbrirtlenient of lnternietallic Compounds in the U.S.

41

zone but it has not yet been identified. The hardness of the (IAZ) has been determined to decrease from a high value just beneath the scale to values typical of the base alloy at its deepest penetration. The hardness profiles have been used to measure the growth of the IAZ, Figure 24. The growth is parabolic as long as the overall oxidation kinetics are parabolic, Figure 15, but establishes a constant thickness when breakaway occurs. The formation of the IAZ greatly decreases the resultant room temperature ductility as indicated in Figure 25. Exposures at temperatures as low as 500°C reduce the plastic strain at fracture to near zero. 'I'he fracture toughncss [94] and ductility [71,95-961 of y alloys are also reduced by exposures in air at elevated temperatures as a result of the layers which form beneath the oxide, as described above. Exposures at temperatures as low as 315°C and for times as short as six minutes at 650°C werc found to embrittle Ti-48A1-2Cr-2Nb [95], Figure 26. An additional factor in the embrittlement is the effect of temperature and atmosphere during the mechanical test. Shrouding the test in argon, Figure 26, or raising the test temperature to 150"C,Figure 27, negate most of the embrittling effects of the prior exposure at elevated temperature [95]. Clearly there is a need for a detailed understanding of the mechanisms for influence of environmental exposures on the mechanical properties of the titanium aluminides.

3

2.5

+

+

A 650°C

X 650T (Argon Test)

unexposed

2

0

X

0

+

'3

1 6

a?

g 1.5

0

.-

X

I

M m

0"

Ei

As-Machined (Argon Test)

0 315°C

1

0

A

A

A

0.5

0

A

8

A Ref.: C.Austin and T.Kelly, Structural Intermetallics.TMS. 1993.p.143. I

0.01

I

0.1

0

0

41

A

A

1

I

I 10 Exposure Time (h)

I

100

I

loo0

Fig.26. Effects of air exposure and test atmosphere on the tensile ductility of y-Ti-48AI-2Cr-2Nb at room temperature (after Austin and Kelly).

G. H. Mrier

A A

0 0 Ref.: C. Austin and T. Kelly, Structural Intermetallics. TMS. 1993. p.143. 0

50

100 Test Temperature ("C)

Fig. 27. Effects of air exposure and test temperature on the tensile ductility of y-Ti-48AI-2Cr2Nb (after Austin and Kelly).

2.3.2.4 Effect of Complex Environments on the Oxidation of Titanium Aluminides There has been relatively little work published on the reaction of titanium aluminides in atmospheres other than air or oxygen. Niu et al. [96] studied the reaction of Ti25A1-11Nb in a simulated combustion atmosphere (N,+1%0,+ O.5%SO2)with and without surface deposits of Na,SO,+ NaCl at temperatures between 600 and 800T. Exposures in the absence of surface deposits resulted in reaction rates similar to those described above for simple oxidation.The rates in the presence of the deposits at 600 and 700°C were initially rapid and then slowed markedly after 25 to 50 hours exposure. The rate at 800°C remained rapid with the kinetics being essentially linear. The major difference in the corrosion morphology at 800°C was the presence of copious amounts of sulfides below the oxide scales. The authors postulate a mechanism of attack involving a combination of sulfidation-oxidation and scale-fluxing. Schaeffer et al. [97] have compared the behavior of the y alloy Ti-48A1-2Cr-2Nb with the nickel-base superalloy RenC 80 in a high velocity oxidation test in a burner rig burning Jet A fuel and a hot corrosion test in which sea salt was injected into the burner rig. The cross-sections of the two alloys after the high velocity oxidation test are compared in Figure 28.The scale on the y alloy is thinner and more compact than that on the superalloy. Figure 29 presents macroscopic photographs of the two alloys following the hot corrosion tests which indicate the hot corrosion resistance of y is equivalent to, or better than, that of RenC 80.The y alloy formed a compact oxide that

2 Research otz Oxidatiotz and Enibrittleinerit of bilernietallic Compounds in the U S . a)

Ti-48AI-2Cr-2Nb

Rent 80

50p.m

Fig.28. Cross-sections of Ti-48AI-2Cr-2Nb and Rene 80 after 200 hours of high velocity oxidation testing at 871 "C (after Schaeffer et al.).

43

44

G. H. Meirr

Gamma (Ti-48A1-2Cr-2Nb) Ren6 80 Fig.29. Macroscopic photographs ofTi-48A1-2Cr-2Nh and 1icnC 80 aftcr 5 0 hours o f h o t coi-rosion testing (after Schrreffer ct al.).

spalled in the region of maximum tcrnperature but did not form liquid corrosion products. RenC 80 underwent oxidation and sulfidation t o ii substantial depth with the formation of liquid corrosion products. It is anticipated that studies, such as those above. will become more numerous as titanium aluminides approach the application stage.

2.3.2.5 Coatings on Titanium Aluminides The rapid oxidation kinetics of the titanium aluminides and. particularly, the embrittlement has created a substantial interest in the possibility of forming a protective coating on these alloys. Coatings of TiAI, have been formed by pack cementation on cx2 [99,100] and y [101].These coatings form continuous alumina scales but TiAI; i\ an extremely brittle compound and tends to crack, particularly for thicker coatings [ 1001. It is also likely that the presence of a layer of TiAl, on the surface of a2or y will be as embrittling as a high temperature oxidation exposure. The ability of TiCrAl alloys to form protective alumina scales raises the possibility of applying them as protective coatings [102]. Coatings have been successfully applied to y substrates by sputtering, low pressure plasma spraying, high velocity oxygen fuel spraying and slurry fusion [102]. Figure 30 shows thuxyclic oxidation kinetics for a sputtered coating and a cross-section of the coating after the oxidation exposure. The oxidation kinetics were similar to those observed for bulk’TiCrA1 alloys and there appears to be niinimal degradation of the coating by either oxidation or interdiffusion with the substrate. Cockerum and Rupp have evaluated the kinetics of silicide coatings on ’Ti [ 1031 and have used a halide-activated pack-cementation method to form boron- and germanium-doped silicide coatings on orthorhombic alloy substrates [104].The coatings greatly decreased the cyclic oxidation kinetics and microhardness measurements did not indicate diffusion of oxygen into the substrate.

2 Research on Oxidatiori m i d Embriitlement of Intermetallic Compounds in the US. h

y

45

0.9 ,

0.2 > , .

0

I

Sputtered Coating on TiAl (28.7Ti-44.5AI-26.8Cr) Cyclic Oxidation at 900°C in Air

500

I

,

.

1

1

1

,

loo0

r

I

1500

,

1

l

.

2Ooo

I

I

,

21

x)

Time (Hours)

Fig.30. Cyclic oxidation kinetics for a sputteredTiCrA1 coating on a y-substrate at 900°C in air and a cross-section of the coating after exposure.

The important aspect of coating any of the titanium aluminides is that thc coating prevent interstitial embrittlemcnt and that thc coating does not function as an embrittlement layer. Additional studies are necessary in this area.

2.3.2.6 Oxidation of Titanium Aluminide-matrix Composites The use of fibres, such as Sic, is being investigated as a way to strengthen a2 and orthorhombic alloys. These can be fabricated by pressing the fibers between thin foils of the titanium aluminide. The presence of these fibers can degrade the oxidation resistance and penetration of oxygen along the fibers can degrade the interfacc and embrittle the surrounding matrix. Additionally, the substantial thermal expansion mismatch between S i c and the matrix (Table I ) can result in matrix cracking during cooling. These phenomena are shown schematically in Figure 31.

Fig.31. Schematic diagram of the oxidation morphology and cracking observed in “orthorhombic

alloy/SiC composites (left) and the oxidation rates of the composites over the temperature range 500-900°C (right).

G. H. Meier

46

2.3.3 Refractory Metal Compounds Refractory metal silicides have been used for many years as furnace heating elements and protective coatings on refractory metals. In recent years there has been interest in using refractory metal intermetallics as structural materials and in composites which have the potential of exceeding the temperature capability of the nickel-base superalloys. However, the oxidation resistance of most of these compounds is poor despite their high contents of Si or A1.The following section describes: (i) the oxidation of specific silicides with emphasis on MoSi2,(ii) silicide coatings. (iii) silicide-based composites, and (iv) oxidation of compounds in the Nb-AI system.

2.3.3.1 Oxidation of Specific Silicides Oxidation of MoSi, Molybdenum disilicide is an intermetallic compound which has been extensively used for high temperature applications, particularly furnace heating elements. The oxides of Mo (MOO,, MOO,) are much less stable than SiO, (Fig. 1) so that silica should be the stable oxide for any but the most dilute Mo-Si alloys [105]. In fact, the nature of the external scale formed is a strong function of temperature. Figure 32 shows the rates of oxidation of MoSi, over the temperature range 500 to 1400°C.The cross-hatched region represents a large amount of data in the range 600-1400°C where the mass changes are small. However, at 500°C the rates are much faster.The oxidation mechanisms can actually be broken into three temperature regimes. 7.8 6.8

5.8

%

h

4.8

\

Z’ a8

3.8

W

2.8

3 m

2

b0 .*

s” Q

1.8

0.8

-0.2

-1 2 0

20

40

60

80

Time, h

100

120

140

160

1

‘O

Fig.32. Oxidation data for MoSi, as a function of temperature.

2 Resenrch ori Oxirlntiori and Emhrittlenient of Interriiernllic Compounds iri the U.S.

47

1. Regime I(lOOO” 600°C was explained as the result of plastic deformation of the matrix near the flaw accommodating the stresses. The cracks were found to be mostly transcrystalline. Westbrook and Wood [113] proposed that the catastrophic nature of the “pest” mechanism was the result of preferential intergranular diffusion of a gaseous element (most likely oxygen or nitrogen), coupled with a temperature dependent hardening reaction. Fitzer et al. [llO, 114,1141 and Schlichting [116] describe pesting of MoSi, as intercrystalline attack whereby each individual grain is enveloped by reaction product.They note that most oxidation occurs in pores or internally along pore canals, and failure occurs as the result of a wedging effect from oxide growth in the defect. Fitzer and Schlichting give no evidence that attack is intergranular and may only assume so, as pores tend to form predominantly along grain boundaries. This is inconsistent with Berkowitz-Mat-

50

G. H. Meier

tuck et al. who found fracture from pesting to be predominantly transcrystalline. Rccent work [lo51 has shown that, while accelerated oxidation is generic to all f o r m s of MoSi,, grain boundaries alone do not result in pesting since dense HIPetl MoSi, did not pest even though it was polycrystalline. Only cast material, which contained prcexisting microcracks, was observed to undergo pesting. It was concluded that pesting was the result of the occurrence of accelerated oxidation within the microcracks. There is a large change in volume going from Mo to MOO; (= 340%), along w i t h the volume change of forming SiO, from Si (= 180 YO).These processes enhance the widening of the pre-existing cracks leading to pesting (i. e. turning to powder).These phenomena are illustrated schematically in Figure 33. 'This mechanism is supported by the observation of oxidation-induced growth of cracks formed in HIPed MoSi, by a microhardness indenter which ultimately resulted in pesting of a material which did not undergo pesting in simple oxidation exposures [117]. McKnrncy et al. [ 1 181 have also concluded that pesting occurs as the result of oxidation in preexisting cracks and pores. There are few studies of alloying effects on the oxidation behavior of MoSi, in the published literature. Recently Yunagihara et al. [ 1191 have reported on alloys in which 15% of the Si was replaced by A1 or 10 % of the Si was replaced by Ta.Ti, Y or Zr.l'he MoSi, had the C l l b crystal structure, the Al-containing alloy the C40 structure. and the others were two-phase Cllb+C40. The alloys were oxidized in air over the temperature range 1435 to 1685"C.Theoxidation kinetics for MoSi, and the Al-;ra-, and --containing alloys were parabolic with MoSi, and MoSi,+-Ta having the slowest rates and MoSi,+Al having a rate constant more than a factor of ten larger. The MoSi,+Ti had rates similar to MoSi, at temperatures below the eutectic temperature in the'I'i0,SiO, system (1550'C) and rates similar to MoSiz+Al above this temperature. The alloys containing Y and Zr exhibited non-parabolic kinetics and oxidation rates which were much greater than for the other alloys. The scales formed on MoSi, and MoSi,+Ta were silica, that on the Al-containing alloy was alumina or a liquid solution. and the scales on the other alloys were complex mixtures of silica and other oxides.

Oxidation of other silicides Accelerated oxidation phenomena have not been studied in detail for refractory metal silicides other than MoSi,. It is generally observed that WSi, has oxidation resistance roughly comparable to that of MoSi,, while TaSi, and NbSi, have considerably poorer resistance, even at very high temperatures [105,120].The Me,Si,-type silicides are only observed to form protective silica films at extremely high temperatures [ 1211. In all cases the rapid oxidation results from the rapid formation of a refractory metal oxide which prevents the formation of a continuous silica layer. Virtually all of the refractory metal silicides have been observad to undergo pesting.

2.3.3.2 Silicide Coatings The refractory metal silicides have been used for many years to protect refractory metals from oxidation in very high temperature, but short duration applications [121]. These coatings have been highly successful but their use in applications which require long term stability have been limited by problems with accelerated oxidation and

2 Research on Oxidation mid Embrirrlement of Intermetallic Compounds in the U.S.

51

pesting, evaporation of SiO at low oxygen partial pressures, interdiffusion with the substrate, and cracking because of thermal expansion mismatch between the coating and substrate.These fractors have been reviewed in detail by Packer [122] and Kircher and Courtright [123].Rccent work by Rnpp and coworkers [124,125] has been directed at improving the resistance of MoSi,-based coatings for Nb-base alloys, for which there is a good thermal expansion match, and Mo, for which there is a relativcly poor thermal expansion match (Table 1).The coatings on Nb, formed by pack cementation, consisted of W additions to the MoSi, to strengthen it and Ge additions to increase the thermal expansion of the protective SiOz layer to improve thc cyclic oxidation resistance and to lower the viscosity to reduce accelerated oxidation at low tempcratures. The coatings were reported to provide cyclic oxidation resistance on Nb for 200 hours at 1370°C [124].The coatings on Mo have made use of Ge doping and also the abovedescribed Na-doping [109], by means of a NaF activator in the coating pack, to successfully limit accelerated oxidation and pesting at low temperatures 11251. Data with regard to coating cracking during thermal cycling from high temperatures, which would be expected to be severe because of the poor match between the thermal expansion coefficients of Mo and MoSi,, have not yet been reported.

2.3.3.3 Silicide-Based Composites The compounds, such as MoSi,, are intrinsically brittle at low temperatures and weak at elevated tempcratures. Composites have been developed to toughen and strengthen the matrix. However, the presence of the second phase also influences the oxidation behavior, usually in a detrimental manner. Figure 36 [126] shows thc effects of 30 vol% of various reinforcing phases on the oxidation of MoSi, in air at 1200°C. The presence of TiB, and HfB, result in an increase in scale growth rates because of the incorporation of TiO, and HfO, into the scales (see schematic in Figure 37). The

Tim (hr)

Fig. 36. Effect of various second phases on the oxidation rate of MoSi, at 1200°C.

52

G. H. Meirr

Fig.37. Schematic diagram of the oxidation morphology developed on a MoSi,-HfB, compo~itt.at 1200°C.

presence of Sic, which is also a silica-former, has relatively little effect on the oxidation rates at this temperature [l26,127].The reinforcing phases have little effect on the low temperature oxidation kinetics under isothermal conditions, i. e. the composites undergo accelerated oxidation at 500°C similar to that shown in Figure 33. However, under thermal cycling conditions, the thermal expansion mismatch between the MoSi, matrix and reinforcing phases, such as S i c (Table l ) , causes matrix cracking. Cracks resulting from discontinuous S i c [127] and from S i c fibers (1281 have been shown to cause complete pesting of MoSi,-Sic composites during oxidation at 500°C.

2.3.3.4 Oxidation of Compounds in the Nb-AI System The oxidation behavior of compounds in this system was first extensively studied by Svedberg [129] who found that the only binary compound which formed a continuous alumina scale was NbA1,. The lower compounds formed scales consisting of NbAIO, and Nb,O, and oxidized nearly as fast as Nb. However. NbAI, is a "line compound" and the A1 depletion caused by the formation of the alumina scale results i n the immediate formation of the lower aluminide Nb,AI below the scale. Thus, after rupture of the initial alumina layer, rapidly growing NbAIO, and Nb,O, form until the Nb,AI layer is consumed and alumina can again be formed o n theNbA1,. Repetition of this process results in a layered scale and nearly linear oxidation kinetics [2,130-1321. Excess Al prevents the formation of the layered scale but degrades the mechanical properties and the long term oxidation resistance because of Al evaporation and alumina growth in the grain boundaries.

2 Rcseardi on Oudiitiori and Ernbrrrrl~meiitof Itirermrrullrc Compoitnds in the U S .

53

The oxidation of NbA1, at intermediate temperatures is a striking example of pesting. It has been observed [130] that NbAI, is susceptile between 550°C and 950°C with a maximum between 650°C and 850°C. Gruhke and coworkers have intensively studied NbAI, of stoichiometric composition [ 132-1341 applying thermogravimetry and Auger electron spectroscopy (AES). The maximum rates were observed at 750°C and reduced pressures between and 0.1 bar 0,.At 750°C and pressures between 0.1 and lO-"bar 0, (He-0, mixtures) a stepwise disintegration was observed. At 2xlO-'O to 7 ~ 1 0 - ' ~ ) b O2 a r (in ampoules with C U - C U ,or ~ Ni-NiO) the disintegration was abrupt and after 24 h the specimens were converted completely to powder. At bar 0, (Nb0,-Nb,O,, Cr-Cr203)no attack or disintegration occurred. However, after oxidation at 750°C under all the conditions listed, AES investigations of the specimens fractured in a IJHV system showed oxygen penetration into the grain boundaries. At bar 0, there was no AI,O, formation in the grain boundaries, but at the higher pressures after the preceding grain boundary saturation with oxygen, oxide formation was observed progressing from the surface to the interior. The very rapid disintegration was ascribed to the wedging effect of the inwardly growing A1,0,. To explain the very fast ingress of oxygen into the NbAI, at temperatures about 750"C, corresponding to a diffusivity of about 10-('cm2/sec,a mechanism was put forward in which formation of fissures is assumed at the grain boundaries, where oxygen can enter by gaseous or surface diffusion [135].The pesting of NbAl, was explained as follows: (i) selective oxidation of aluminum with AI,O, scale formation results in Aldepletion of the NbAI, phase, preferentially along grain boundaries, (ii) the Al-depletion leads to a phase transformation to Nb,AI at the grain boundaries, NbAI, transformation causes fissure formation at the grain boundaries and cracking of the outer scale. (iii) oxygen from the atmosphere penetrates into the fissures and A&O, is formed on the surfaces of the scparated grains within the material. (iv) dislocations and low-angle boundaries can also act as short-circuit paths for aluminum diffusion to the grain surfaces, then Al-depletion and Nb,A1 formation also opens cracks into the grains. Doychak and co-workers have also performed extensive studies on accelerated oxidation of binary NbAI, and material alloyed with Cr and Y [136-138].The high temperature oxidation is improved by the alloying additions [137]. The Cr additions resulted in a layer of AlNbCr forming beneath the scale.This compound maintained the stability of the alumina scale better than the Nb,AI which forms on the binary compound. The Y additions reduced convolutions in the scale. The alloyed NbAI, also undergoes pesting [136, 1381 but the presence of Cr apparently prevents grain boundary oxidation. A careful TEM study of the initial stages of accelerated oxidation, which leads to pesting, revealed an external scale covering an internal oxidation zone consisting of A120, and Nb.'The rupture of the external scale and rapid oxidation of the Nb was the cause of the accelerated oxidation. The authors suggest that a similar phenomenon may be occuring in the grain boundary region of binary NbAI,. A number of studies have evaluated the effects of alloying elements on the oxidation behaviour of Nb-A1 alloys [132,136-1411. In some cases the alloys formed alumina scales but none of the alloys were resistant over a broad enough range of exposure conditions to be considered for extensive application. This is particularly the case when the compitions are restricted to those which provide even marginal mechanical

54

G. H. Meier

properties. The author believes that, despite the existence of ongoing programs, attempts to protect Nb-base compounds by the selective oxidation of A1 or Si are fruitless.

2.4 Concluding Remarks The above discussion indicates that there are relatively few intermetallic compounds that form protective alumina or silica scales over wide ranges of exposure conditions. 'The use of high-volume fraction y' (Ni,AI) Ni-base superalloys and aluminide coatings on superalloys remain the most successful technological applications of intermetallic compounds in structural materials. The titanium aluminides are approaching the application stage. However, exposure to oxidizing environments degrades the mechanical properties of essentially all the titanium aluminides. The author believes that developing an understanding of the mechanisms associated with this phenomenon and finding solutions are the most important problems currently in intermetallics oxidation. The possibilities of other intermetallics reaching widespread application as structural materials seems unlikely because of poor fracture toughness and/or poor oxidation resistance. Additionally, it is clear that addition of second phases for strengthening or toughening generally degrades the oxidation resistance of those compounds which are oxidation resistant. The compounds Ni,Al, NiAI, Fe,AI, and FeAl have sufficient oxidation resistance for high temperature applications if their mechanical property shortcomings can be resolved. The only refractory metal compound with adequate oxidation resistance is MoSi,.

2.5 Acknowledgements The author gratefully acknowledges the contributions of D. Berztiss, R. Cerchiara, 1. Rakowski, C. Sarioglu, and C. Stinner for generating many of the results discussed in this paper and for preparation of the illustrations. The author wishes to thank colleagues M.P Brady, J. Doychak, J. A. Nesbitt, R. A. Perkins, B. A. Pint, R. A . Rapp, J. C. Schaefer, and J. L. Smialek who provided copies of their papers, in some cases prior to publication. Finally, the author acknowledges the continued, valuable collaboration with N. Birks and E S.Pettit at the University of Pittsburg.

2.6 References [l] E. A . Aitken in: Intermetallic Compounds, J. H. Westbrook el., p. 491, Wiley (1967). (21 G. H. Meier in: Oxidation of High Temperature Intermetallics,T. Grohstein and J. Doychak

eds., p. 1,TMS (1989). (31 G. H. Meier, E S. Pettit: Mater. Sci. and Eng. A153 (1 992) 548.

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[4] G. f1. Meier, N. Birks, E S. Pettit, R. A. Perkins, H. J. Grabke: in Structural Intermetallics. R. Darolia J. J. Lewandowski. C. T. Liu. P. L. Martin. D. B. Miracle. M. V. Nathal eds., p. 861,TMS (1993). [5] J. Doychak in: Intermetallic Compounds. J. H. Westbrook. R. L. Fleischer eds., p. 977, Wiley (1994). (61 A. K. Vasudevan, J.J. Petrovic: Mater. Sci. and Eng.AZ55 (1992) 1. 171 R. Gibala, A. K. Ghosh, D. C. Van Aken, D. J. Srolovitz, A. Basu, H. Chang, D. P Mason, W. Yang: Mater. Sci. and Eng. A155 (1992) 47. [8] J. Cook, A. Khan, E. Lee. R. Mahapatra: Mater. Sci. and Elng. A255 (1992) 183. 191 R. 7:DeHoff:Thermodynamics in Materials Science, McGraw-Hill, (1993) p. 492. [lo] C. Dior,P Choquet, R. Mevrel in: Proc. Internat. Conf. on Residual Stresses (ICRS-2) G. Beck, S. Denis. A Simon eds., Elsevier, London, (1988) p. 273. [ l l ] ES.Petrit:Trans. Mct.Soc.AIME239 (1967) 1296. 1121 G. H. Meier, E S. Pettit, J. L. Smialek: Materials and Corrosion. 46 (1995) 232. 1131 E. Schumann, M. Riih1e:Acta Metall. Mater. 42 (1994) 1481. 1141 B. A. Pint, L. W. Hobbs: “The Oxidation Behavior of Y,O,-Dispersed Ni,AI”. in: Oxide films on Metals and Alloys, B. R. McDougall et al. eds.. Electrochem. SOC.Proc., Vol. 92-22,1992, p. 92. [15] B. A. Pint, K. B. Alexander, F? 1: Tortorelli: “The Effect of Various Oxide Dispersions on the Oxidation Resistance of Fe,AI”, Mat. Res. SOC.Symp. Proc.,Vol. 364.1995, p. 1315. [16] G. C. Rybicki,J. L. Smialek: Oxid. Metals31 (1989) 275. [17] B. A. Pint, L. W Hobbs: Oxid. Metals 41 (1994) 203. [18] M. M. W Brumm, H.J. Grabke: Corr. Sci.33 (1992) 1677. [19] J. Doychak, M. Riih1e:Oxid. Metals32 (1989) 431. [20] B. A. Pint, M. Treska, L. W Hobbs: submitted to Oxid. Metals. [21] L. Smialek, J. Doychuk, D. J. Gaydosh: Oxid. Metals 34 (1990) 259. [22] M. W Brumm, 11.J. Grabke: Corr. Sci.34 (1993) 547. [23] J. L. Smialek: Me1.Trans.A 9A (1978) 309. [24] E. Schumann, J. C. Yang, M. J. Graham, M. Riihle: Materials and Corrosion 46 (1995) 218. [25] B. A. Pint, L. W Hobbs: “The Cyclic Oxidation Behavior of Oxide-Dispersed P-NiAI”, Mat. Res. SOC.Symp. Proc., Vol. 364,1995, p. 987. 1261 C. Stinner, F;S. Pettit, G. H. Meier: unpublished research, Univ. of Pittsburgh, 1994. [27] B. A. Pint, L. W Hobbs: J. Electrochem. SOC.141 (1994) 2443. [28] E. Schumann: Oxid. Metals 43 (1995) 157. [29] H.J. Grabke, G. H. Meier: Oxid. Metals 44 (1995) 147. 1301 M. W Brumm, H. J. Grabke, B. Wagemann: Corr. Sci. 36 (1994) 37. 1311 D. A. Berztiss: Ph. D Thesis, Univ. of Pittsburgh, 1996. [32] G. W Coward, L. W Cannon: “Pack Cementation Coatings for Superalloys: A review of History,Theory, and Practice”, Paper 87-GT-50,Gas Turbine Conf., ASME, 1987. [33] J. S. Smith, D. H. Boone: “Platinum Modified Alurninides-Present Status”, Paper 90-G‘r-319, Gas Turbine and Aeroengine Congress, ASME, 1990. 1341 S. C. Kung, R. A. Rapp: Oxid. Metals 32 (1989) 89. [35] R. Bianco, R. A. Rapp: J. Electrochem. SOC.140 (1993) 1181. 136) K. Bianco, R. A. Rapp: in High Temperature Materials Chemistry-V, W. B. Johnson and R. A. Rapp Eds.,The Electrochem. SOC.1990 p. 211. 1371 R. Bianco, R. A. Rapp, J. L. Smialek: J. Electrochem. SOC.140 (1993) 1191. [38] W Da Costa, B.Gleeson, D. J. Young: J. Electrochem. SOC.141 (1994) 1464. [39] W Da Costa, B. Gleeson, D. J. Young: J. Electrochem. SOC.241 (1994) 2690. [40] C. Stinner, E S. Pettit, G. H. Meier: unpublished research, IJniv. of Pittsburgh, 1995. [41] J. Doychak, J. A. Nesbitt, R. D. Noebe, R. R. Bowman: Oxid. Metals 38 (1992) 45. [42] R. A. Perkins: unpublished research, Lockheed Palo Alto Res. Lab., 1993. 1431 J. Whittenberger in: Structural Intermetallics, R. Darolia, J. J. Lewandowski, C.T. Liu, P. L. Martin, 11.B. Miracle, M. V. Nathal eds.,p. 819,TMS (1993). [44] J. A. Nesbitt, C. A. Barrett in: Structural Intermetallics, R. Darolia, J. J. Lewandowski, C.T. Liu, P.L. Martin, D. B. Miracle, M. V. Nathal eds.. 1.’ 601,TMS (1993). [4S] J.A. Nesbitt, C. E. Lowel1:“Prediction of the High‘remperature Oxidative Life of Intermetallics”, Mat. Res. SOC.Symp. Proc.,Vol. 288,1993, p. 107.

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(46) J. I,. Smialek,J. A.Nesbitt, W J. Brindley, M . P Brariy, J. D ~ ~ h nR.kM. . L)ickcwm.U R. l l ~ t l l : “Service Limitations for Oxidation Resistant Intermetallic Compounds”. Mat. Rcs. Soc. Symp. Proc.,Vol. 364,1995, p. 1273. [47] A. Rahmel, PJ.Spencer: Oxid. Metals35 (1991) 53. [48] K. L. Luthra: Oxid. Metals 36 (1991) 475. [49] J. M . Rakowski: Senior Thesis, Univ. of Pittsburgh, 1992. [50] N S. Choudhurv, H. C. Graham, J. W Hinze: “Oxidation Behavior o f?’itanitmi Aluminides”. in Properties of High Temperature Alloys. Z . A. Foroulis and F.S. Pettit etls., The E1ectr.ochcm SOC..1976,p. 668. [Sl] M. G. Mendiratta, N. S. Choudhurv: “Properties and Microstructure o f High-Tempcraturc Materials”. AFMLTR-78-112, Contract No. F33615-7.5-C-l0OS,(Systems Research Lnlioratories, Inc., 0hio.August 1978). [52] J. Rakowski. G. H. Meier. R. A . Perkins:The Oxidation and Embrittlement of CY: (Ti;AI)l i t a nium Aluminides, in Microscopy of Oxidation 2. Inst. of Materials. 1993.p. 376. (531 7:A. Wallace,R. K. Clark, K. E. Wiedemann:Oxid. Metals 42 (1994) 4.51. 1541 M. Khobnib, E W. Vahldiek: “High Temperature Oxidation Behaviour o f Ti;AI Alloys”, Second International SAMPE Metals Conference. F. H. Froes and R. A. Cull etls.. C’OVIIIX Ca, 1988,pp. 262-270, [55] C. If. Koo, J. W Evans, K. I.:Song, 7:I I . Yu:Oxid. Metals -12 (1Y94) 529. [56] 7:A. Wallace,R. K. Clark. K. E. Wiedetnann, S. K. Sankaran: Oxid. Metals 37 (1992) 1 I 1. [57] J. C. Schaeffer: Scripta Met.28 (1993) 791. [58] S. J. Balsone: “The Effect of Elevated Temperature Exposure on the Tensile a n d C’rrep Properties of Ti-24AI-1 INb” In Oxidation of High Temperature Interinetallics.‘T.Grobstcin and J. Doychak eds.,?’MS, 1989,p. 219. [59] R. Cerchiara, J. Rakowski, E S . Pettit, nnd G. H. Meier: Unpublished Research, I!nivrrsity oS Pittsburgh, 1993. [60] G. H. Meier, E S.Pettit, S. l f u :“Oxidation Behavior of Titanium Aluminides” J. de Physiclue IV, Colloque C9.1993, p. 395. 1611 D. S. Appolonia: “Thc Oxidation of Gamma-Titanium Aluminides”, Bachelor of Science Thesis,University of Pittsburgh, Pittsburgh, PA (1988). [62] G. H. Meier, D. S Appalonia, R. A. Perkins, K. 7: Chiang in: Oxidation o f Iligh-Tempcraturc Interrnetallics,T. Grobstein and J. Doychak. eds.. (TMS, 1989) p. 185. (631 A. Gil, H. Hoven, E. Watlura, and W J . Quadakkers: Corr. Sci.34 (1993) 615. [64]J. M. Rakowski,E Dettenwanger, E. Schitmann, G. ?I.Meier: E S.Pcttit. ill. Riililc: ”The Effect of Surface Preparation on the Oxidation Behavior of Gamma TiAl Base Intermetallic Alloys’’, submitted to Scripta Mater. [65] C. S. Giggins, E S. Pettit:Trans. Met. SOC.AIME 245 (1969) 2509. [66] S. Becker,A. Rahmel, M. Schorr, M. Schiitze: Oxid. Mctals 38 (1992) 425. [67] N. Zheng, W J. Quadakkers,A. Gill, H. Nickel: Oxid. Metals 14 (1995) 417. [68] E. Kobnyashi, M. Yoshihara, R. Tanaka: H$h Tcmp.Tech. 8 (1990) 179. [69] CL Figge, A. Elschner, N Zheng, H. Schitsfk( W J . Quadakkers: Fresenius J. Anal. Cheni. 3.46 3 (1993) 75. [70] J. M. Rakowski, E S. Pettit, G. H. Meier. E Dettenwnnger, E. Schumnnn. ill. Riihle: Scripla Met. et Mater.3.7 (1995) 997. [71] W E. Dowling Jr., W.7:Donlon: Scripta Met. et Mater. 27 (1992) 1663. [72] R. Field, J. Schaeffer, C. Austin, 7:Kelly: unpublished, GE Aircraft Engines. 1993. [73] I! Cheng, f?Dettenwanger, J. Mayer, E. Schumnnn, M . Riihle: Scripta Met. et Matcr. 3-1 ( I 9%) 707. [74] E Dettenwanger, E. Schumann, J. Rakowski, G. H. Meier, M. Riihle: submitted to Materials and Corrosion. [75] K. W Beye, R. Gronsky:Acta Met. et Mater. 42 ( I 994) 1373. [76] N. Zheng, W Fischer. H. Griibmeier. I/:Shemet, W J. Quadakkers: Scripta Met. et Matcr. .?.? (1995) 47. [77] A. R a h m , W J. Quadakkers, M. Schiitze: Materials and Corrosion 46 (199s) 271. [78] R. A&.??kins. K. 7:Chiang, G. H. Meier, R. A. Miller: “Formation o f Alumina on Niohium a n d Titanium Alloys”, in: Oxidation of IIigh-Temperature Intermetallics, 71 Grohstein and J. Doychak, eds.,The Min., Met., and Materials SOC.,1989 p. 1.57.

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[79] Y Shida, H. Anntlu: Mater.Trans. JIM 35 (1994) 623. [SO] H . Nickel, N. Zheng,A. Elschner, W J. Quodnkkers:Mikrochim. Acta. 119 (1995) 23. [81] T A . Wnlluce, R. K. Clark, S. N. Sarikarnn, K. E. Wiederrrinnri:In Environmental Effects on Advanced Materials, R. H. Jones and R. E. Ricker eds.,TMS 1991,p. 79. [82] R. A. Perkins. G. H. Meier: “Oxidation Resistant Aluminides for Metal Matrix Composites”, in Advanced Materials Confcrence 11,F. W. Smith ed.. Advanced Materials Institute, 1989.p. 92. [S3] Y Shirlo, H . Anudn: Corr. Sci. 35 (1993) 945. [84] C. Stinner:M. S.Thesis, University of Pittsburgh, 1992. [85] M . P Brndy, J. L. Sminfek, F: Terepkn:Scripta Met. et Mater. 32 (1995) 1659. [86] M. R Brndy, J. L. Sniinfek,D. L. llumphrey: Mat. Res. SOC.Symp. Proc..Vol. 364.1995, p. 1309. [87] M . P Hrod-y, J. L. Sniiolek, D. L. Ilunzphrey: “Mechanism of Alumina Formation in li-Cr-A1 Alloys“, Extended Abstract. Fall Meeting.The Electrochem. SOC.,1995. [88] J. L. Sniialek, D. L. llumphrey: Scripta Met. et Mater.26 (1992) 1763. [89] C. H . Word,J.C. Williarns,A. W Thompson: Scripta Met.28 (1993) 1017. [90] C. H. Wurk 1nternat.Mater. Rev.38 (1993) 79. [91] Y Saifohand K . Mino: Mater.Trans. JIM34 (1993) 393. [92] P R. Smith, J. A. Groves, C. G. Rhodes: “Preliminary Mechanical Property Assessment of a SiC/OrthorhombicTitanium Aluminide Composite’‘ in Structural Intermetallics, R. Ilarolia, J. J. Lewandowski. 10GPa when formula (7) is applied. Presumably, the stresses relieve during the relatively slow bending deformation by Coble creep of the oxide or by a slight scale wrinkling [27].This suggestion was confirmed by RT stress measurements on the side face of bending bars using OFS. The measured stress on the compressive side of the bending bar was considerably smaller than that expected from formula (7). For high temperature applications of materials, spallation of oxide scales has to be prevented. A high oxide-metal adhesion is, of course, most important in order to avoid a large-scale delamination of the protective oxide. However, also for high oxide-metal adhesion, the scale can fail by buckling under compression when large voids are present. After scale buckling, spalling occurs by the propagation of cracks through the scale. Thus, a high fracture toughness of the oxide is also important in order to suppress compressive scale failure. The fracture toughness is probably easier to evaluate by tensile tests as discussed above.

9.8 Experimental Difficulties and Theoretical Problems The investigation of defect patterns in oxide scales formed under a graded scale loading seems to be promising since critical load parameters could be analysed ex situ. However, there are a series of experimental difficulties and open theoretical problems in realising this approach. Within the proposed approach, the fracture toughness is determined from the position of the crack tip in a nonuniform stress field. Because of the small crack opening width near the crack tip, the tip position can be detected with high precision probably only by electron microscopy. A difficulty arises in measuring the distance of the tip

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from the specimen edge. In the present work. optical microscopy was applied t o nieasure this distance where features of the scale were used to identify the tip position at lower magnification. When cracks are produced at elevated temperatures and analysed at KT. there is an uncertainty in the detection of the original crack tip position since the cracks can partly close during cooling of the specimen. In particular. this could happen when the failure strain of the oxide is smaller than the strain due to the thermal expansion mismatch. In-situ observation at high temperature avoids this problem. A further difficulty is to determine the stress field in the scale with sufficient accuracy. This concerns especially the determination of the tensile strength or fracture toughness from RT experiments on scales which exhibit high residual stress (> 1 GPa). Since the total strain in the scale results as the small difference of two large terms, cS- kRl(cf. formula (2), E ~ 0< ) ,both terms have to be known with high accuracy. This problem is not specific to the present approach. It applies also to other tests of tensile failure.The problem does not arise for small residual strains, i. e. for materials with low thermal expansion mismatch, and for tests at elevated temperatures. The best way to determine the stresses in the scale would be a direct measurement. However, X-ray methods have usually a limited spatial resolution which makes it difficult to measure nonuniform stress fields. 'The application of OFS for scales consisting of a-A1203provides a sufficient spatial resolution and permits, i n principle, to examine stress variations in the scale. However, only the trace of the stress tensor can be nieasured for an untextured polycrystalline scale. Thus, anisotropic stress states have to be analysed in combination with a mechanical modelling of the scale loading in order to deduce the stress components from the trace of the stress tensor. The intent in the present work was to determine the scale loading by measuring the isotropic residual strains in the scale at RT and by calculating the additional strains due to a controlled substrate deformation. However, the calculation of the substrate deformation in the case of single crystalline NiAl is very difficult because of its strongly anisotropic plastic behaviour. To simplify matter, isotropic material properties were supposed in the above analysis. For the case of spherical indentation, it was found that isotropic models are insufficient to describe the observations. Thus, in the case of NiAI, indentation can serve only as a semi-quantitative method for comparative studies. In the case of the bending deformation, the deformed specimens gave no evidence for strong deviations of their shape from that expected for isotropic properties. Nevertheless, a more detailed analysis of the scale deformation due to specimen bending is necessary. The above estimates of fracture-mechanical parameters by assuming isotropic material properties represent only rough approximations. The fracture analysis given above concerned the propagation of a single crack. The observed crack patterns suggest however the propagation of a front of parallel cracks. An accurate analysis has to take into account the mutual unloading of parallel cracks. Crack interaction effects have been studied in [9,41] for the case of an elastic substrate, and in [lo, 111 the ductility of the substrate has been taken into account.'The crack cl&sity can increase with substrate deformation due to the formation of new cracks between earlier cracks. However, due to the ductility of the substrate, a maximum density will be reached as outlined in [lo].At high tensile deformation, also scale delamination can occur by the deflection of trough-scale cracks into the oxide-metal

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interface [15]. The onset of delamination can be used to characterise the scale adherence. Related problems have been investigated in [9, lo].

9.9 Summary The intent of the present paper was to focus the attention on the analysis of crack patterns in scales caused by a graded loading. In this way, the scale can be tested simultaneously under different loads which should permit to derive critical loads or failure strains by ex-situ observations provided the scale loading is known. In the present work, the graded loading was established by four-point bending (observing the side face of bending bars) and spherical indentation of oxidised specimens. Spherical indentation is easily to perform. However, a determination of the resulting stress state in the scale seems to be too difficult.Thus, this test can be used mainly for comparative studies. The bend test is presumably more appropriate for determining fracture-mechanical properties. Other specimen deformation can be more suitable for special purposes. An ex-situ analysis of crack patterns formed under a nonuniform scale loading can provide independent information on the oxide fracture toughness and tensile strength. Whether the crack pattern characterises the toughness or strength depends on the choice of the direction of the stress gradient relative to the direction of the tensile stress. In this way, the fracture toughness could be determined, in principle, without knowledge of the size of initial flaws in the scale. A difficulty of the proposed approach is the determination of the graded scale loading. For scales consisting of a-A1203,optical fluorescence spectroscopy permits to measure stresses with high lateral resolution. A disadvantage is that only the trace of the stress tensor is obtained. Thus, an interpretation of spectroscopic data has to be complemented by a mechanical modelling of the scale loading. In the present analysis, large-scale spallations of alumina scales on NiAl were always connected with the presence of interfacial voids. In the absence of large voids and for small strain rate, it was found that the scale was well adherent under compressive substrate deformation even for large specimen deformations (cf. Fig. 7a).This suggests the presence of an effective stress relief mechanism by a slight scale wrinkling or by oxide Coble creep. In summary, the establishment of a graded scale loading in combination with spatially resolved stress or strain measurements could be an effective method for an ex-situ investigation of failure mechanisms and fracture-mechanical parameters.

9.10 Acknowledgements This work was supported by the Deutsche Forschungsgemeinschaft. The authors would like to thank Th. Hutzler, W Loser and G. Vuerst for providing NiAl single crystals and performing initial bend tests, and J. Edelmunn for his assistance in the

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SEM analysk’lle authors are very grateful to D. R. Clarke and D. M . Lipkin for the support in measuring stresses by OFS, and to 11. Balke and V: Marx for helpful discussions concerning the FEM modelling.

9.1 1 References [l] A. G Evans, R. M. Cannon: Mater. Sci. Forum 43 (1989) 243. [2] J. Robertson, M. I. Manning: Mat. Sci.Technol. 6 (1990) 81. [3] M. Schiifze:Die Korrosionsschutzwirkung oxidischer Deckschichten unter thermisch-chemisch-mechanischer Werkstoffbeanspruchung, Gebriider Borntraeger Berlin. Stuttgart (1991). [4] M. M. Nagl, W i? Evans: J. Mater. Sci. 28 (1993) 6247. [5] H. E. Evans: Matcr. Sci. Rev. 40 (1995) 1. [6] H. 6 Evans, R. C. Lobb: Corr. Sci.24 (1984) 209. [7] A. G. Evans, J. W.Hutchinson: Int. J. Solids Structures20 (1984) 455. [8] H. E. Evans, G. Z? Mitchell, R. C. Lobb, D. R. J. Owen:Proc. R. SOC.L0nd.A 440 (1993) 1. (91 J. W.Hutchinson, Z. Suo: Mixed mode cracking in layered materials, in: Advances in Applied Mechanics (eds. J. W. Hutchinson and T. Y. Wu) vol. 29, p. 64, Academic Press, New York (1991). 1 [lo] M. S Hu, A. G. Evans: Acta metall. 37 (1989) 917. [l 11 F Delannuy, fl Warren: Acta metall. mater. 39 (1991) 1061. [12] S.R. J. Saunders, H. E. Evans, J. A. Stringer (eds.): Mechanical properties of protective oxide scales,Workshop,Teddington, UK, June 1994,in: Mater. High Temp. 12 (1994) 83. [13] S. R. J. Saunders, M. M. Nagl, M. Schiitze:Mater. High Temper. 12 (1994) 103. [14] M. M. Nagl, S.R. J. Saunders, W I: Evans, D. J. Hall Corr. Sci.35 (1993) 965. [15] M. M. Nugl, W I: Evans, D. J. Hall, S.R. J. Saunders: Oxid. Met. 42 (1994) 431. [16] Z? Huncock, J. R. Nicholls, K. Mahmood Corr. Sci. 35 (1993) 979. 1171 I? I!Hou, J. Sfringer:Oxid. Met. 38 (1992) 323. [18] M. D.Drory, J. W Hutchinson: Science 263 (1994) 1753. [19] D. R. Sigler: Oxid. Met. 40 (1993) 555. [20] D. M. Lipkin, D. R. Clarke: Oxid Met., submitted [21] Q. Ma, D. R. Clurke:J. Am. Cer. SOC.76 (1993) 1433 [22] D. M. Lipkin, D. R. Clarke, M. Hollatz, M. Bobeth, W Pompe: Corr. Sci.. submitted. (231 D. B. Mirac1e:Acta metall. mater. 41 (1993) 649. [24] J. K. Doychak, i? E. Mitchell,J. L. Smialek: Mat. Res. SOC.Symp. Proc. Vol. 39 (1985) 475. [25] M. W Brumm, H. J. Grabke: Corr. Sci. 33 (1992) 1677. [26] E; Schumann: Oxid. Met. 43 (1995) 157. [27] Z. SUO:J. Mech. Phys. Solids 43 (1995) 829. [28] R. D. Noebe, R. R. Bowman,M. K Nathal: Inter. Mater. Rev.38 (1993) 193. [29] Thermophysical Properties of Matter, The TPRC Data Series, Ed. Y.S. Touloukian, Vol. 13, IFIIPlenum, New York (1977). [30] M. Schiitze: Oxid. MFt. 44 (1995) 29. [31] J. He, D. R. Clarke: J.Am. Cer. SOC.78 (1995) 1347. [32] M. W Brumm, H. J. e a b k e : Corr. Sci.35 (1993) 547. [33] M. ZIollatz, M. Bobefh,W Pompe, K Marx: Acta metall. mater., submitted. [34] K. L. Johnson: Contact Mechanics, Cambridge University Press, Cambridge (1985). [35] R. Hill, E R. S., B. Storakers, A. B. Zdunek: Proc. R. SOC.Lond. A 423 (1989) 301. [36] S.Biwa, B. Storakers: J. Mech. Phys. Solids 43 (1995) 1303. [37] R. D. Noebe, R. Gibala: Scr. metall. 20 (1986) 1635. [38] Z? Lipetzky,A. Wanner, B. Schietinger, E. Arzt: in: Numerical Predictions of Deformation Processes and the Behaviour of Real Materials, Proc. 15th Riso Int. Symp. on Materials Science, eds. S. I. Andersen, J. B. Bilde-Sorensen. T. Lorentzen, 0.B. Pedersen and N. J. Sorensen. Roskilde, Denmark 1994,p. 399.

9 Failure ofAlunzina Scules on NiAl Under Graded Scale Loading [39] l? Hancock, J. R. Nicholls: Mater. Sci.Technol.4 (1988) 398. [40] M . S. HLL: Mater. Res. Soc. Symp. Proc.Vol.130 (19x9) 213. [41J M. D. Thotiless. E. Olsson. A. Gupta:Acta rnetall. rnater. 40 (1992) 1287. [42] J. R. Nicholls, C. Mendes, l? Hancock: Mater. HighTemp. 12 (1994) 85. [43] M. Boberh, E. Bischoff; E. Schurnann, M. Rocksrroh, M. Riihle: Corr. Sci. 37 (1995) 657.

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Although originally AISI 347 stainless steel seemed a reasonable choice, it was soon shown that it is not suited for long term operation over several years. The best corrosion protection would be provided by materials that can passivate. AISI 316 or AISI 310 stainless steel would in principle provide such a relatively stable and protective oxide scale in the presence of molten carbonate at least under cathodic conditions. IJnder fuel gas conditions AISI 316 is unsuitable since no protective layer is present [1,21. A systematic evaluation of various possible materials classes showed that the choice had to be made from nickel alloys as Inconels, Incoloys and Hastelloys, or austenitic stainless steels such as the already mentioned AISI 316L or AISI 310s. At the same time it became clear that no single material would be able to perform well under both reducing and oxidising conditions. An extra problem was created by the conflicting demand for corrosion resistance - as provided by a ceramic layer under anodic conditions, for example - and the necessary good electrical conduction.l'he conclusion from experimental work in many different laboratories was that different coatings would be necessary for different parts of the separator plates [2]. A nickel or copper coating would be appropriate for the active part at the anode, while aluminium would be a good choice for the wet seal area, where a LiAIO, oxide layer, which would be formed, could provide the necessary high ohmic resistance. Aluminides or aluniinised diffusion coatings have shown to be effective. Ni and Al both being interesting materials for the MCFC, we have investigated pure Ni and various Ni alloys including the intermetallic NiAl for their corrosion resistance under fuel cell conditions [l,3-61. The objective of this paper is to compare the corrosion behaviour of NiAl in molten carbonate with that of metallic nickel.

10.2 Experimental The experiments were performed in a so-called pot-cell which is schematically given in Fig. 1. The inlet gas composition is the MCFC anode inlet gas composition: 64'26 hydrogen, 16% carbon dioxide and 20% water vapour; the carbonate used is an eutectic mixture of 62 Yo Li,CO, and 38 % K,CO,. The cell temperature was kept constant at 923 K. The outlet gas was led through a water lock to prevent air inlet. The reference electrode consisted of an alumina tube filled with carbonate and with a small hole (app. 0.1 mm diameter) in the bottom. The function of the hole is to realise contact between the melt in the reference electrode compartment and in the working electrode compartment. The reference gas was a mixture of 14 % oxygen and 30 YOcarbon dioxide, balanced with nitrogen. The nickel used was of 99.98 % purity as purchased from Goodfellow Metals Ltd., Cambridge, England. The specimens had dimensions of 5 x5x 1mm3 and were ground to grit 1000.The counter electrode consisted of 99.99 YOpure gold foil that was bent in a cylindrical shape. The nickel electrode and the counter electrode were spot-welded to a 99.99 % pure gold wire of 0.5 mm diameter.

10 The Corrosion Behaviour of NiAl in Molten Carbonate at 650°C

163

Fig. 1. Schematic set-up of a pot-cell typc set-up. 1.Gas inlet reference electrode; 2. Gas inlet; 3. Gas outlet; 4. Glass flange glued to the jackct; 5. Outer alumina jacket; 6. Catalyst in gas inlet; 7. Alumina crucible; 8. Reference electrode gas outlet;9. Silicon rubber stops; 10. Glass lid; 11.Flag electrode. CE: Counter electrode. WE: Working electrode.TC: Thermocouple. RE: Reference electrode

The nickel-aluminium was supplied as a small bar, made upon special order by Philips Research Laboratories, Eindhoven, The Netherlands, by mixing the proper amounts of nickel and aluminium and subsequent arc-melting. The bars were cut into slices. The dimensions of the specimens used for the electrochcmical measurements were 6 x 6 ~ mm3.The 1 surface of the specimcns was not pre-treated because this alloy was considered to be too brittle to apply abrasive techniques without damage to the specimens. Specimens of the NiAl intermetallic were laser welded to a nickel-wire because gold and aluminium can form an eutecticum with a liquidus temperature below 873 K [7]. The polarisation curves werc recorded immediately after immersion with an initial potential of - 1700mV (Ni) or - 1100mV (NiAI) and a final potential of + 100mV. Potential steps of 5 mV werc applied, the current was measured after 50 s. To investigate the steady-state corrosion layers, quenching experiments were performed. Shortly after immersion, the working electrode potential was fixed at a chosen value for 4 hours. After this period the specimens were taken out of the melt while still being under potential control and cooled down to room temperaturc in air as fast as possible.The carbonate adhering to the spccimen prevented contact and interaction

164

J. l? 7: Vossen.A . H.H. Janssen and J. H. W de Wit

between the specimen and air. Sometimes longer polarisation times of 16 or 24 hours were applied to enable the formation of more oxide. The carbonate was removed by rinsing in water (pure nickel) or in a mixture of acetic acid and acetic acid anhydride (NiAl). This mixture was chosen to prevent the dissolution of the aluminium oxide scale in water. Quenching experiments were performed at potentials of 1 I00 niV 01' more anodic, since these potentials are of interest for MCFC operation. The impedance measurements were recorded in the potentiostatic mode w i t h ;I Zahner IM5d-system and IMSd-software. The amplitude of the applied potential for the impedance measurements was 10 mV (peak-peak).For our system this value gave a good signalhoke ratio and was sufficiently small to avoid clcviation from linearity [4].111e impedance was recorded in the range of 50mHz to 10 kHz.The measurements were started immediately after applying the dc-potential, from the lowest frequency, going up to higher frequencies and finally back to the lowest frequency. The impedance spectrum was measured 5 times at each frequency. The reproducibility was good. i. e. in a Nyquist plot, the results of duplo measurements were within 5 '%deviation of the average. For more experimental details we refer to earlier publications (3-61. The surface and the cross section of the quenched specimens were investigated with SEM/EDX, the corrosion products were investigated with ex-situ X-ray powder diffraction (Guinier-de Wolff method, Du-K,,.,). ~

10.3 Results and Discussion 10.3.1 Quenching Experiments Metallic nickel is stable at - 1100mV and -900 rnV. Therefore no oxide scale is formed on the surface at these potentials. Analysis of NiAl quenched after 4 hours polarisation at - 1lOOmV or -900mV shows that the surface is covered with small crystals with an octahedral and cubic morphology. as can be seen in Fig. 2a. SEMEDX-analysis showed that the corrosion product contains aluminium and oxygen (lithium can not be detected by EDX). The cross section of this specimen (Fig. 2b) shows that the corrosion product is present as a thin oxide layer. The only phase that could be detected by X-ray diffraction is a-LiAlO,. After longer oxidation times (up to 24 hours) the dimensions of the a-LiA102-crystals increased slightly,but t M i c k n e s s of the oxide scale did not increase significantly. a-LiAIO, could alobserved on aluminised stainless steel exposed under MCFC fuel and oxidant conditions for 5400 hours [8]. Approximated phase diagrams [9] and data of other groups [ l O , l l ] indicate that y-LiAlO, is probably the stable phase in molten carbonate of the chosen composition at 650"C.Thc a- and P-LiAIO, phases may also exist, but are not expected to be stable on the long term. Therefore a-LiA102is expected to be formed as an intermediate for the y-LiAlO,, which is formed slowly. aLiAIO, may be formed preferentially a consequence of epitaxial growth, while nucleation problems for the phase transition from a-LiAlO, to y-LiAlO, may occur. After quenching Ni specimens from potentials of - 775 mV and more anodic, oxide crystals can be observed also on the surface of these specimens as can be seen in

10 The Corrosion Behavioiir of NiAl irr Molten Carbonate at 650°C

165

Fig. 2. Surface (a) and cross section of NiAl (b) after 4 hours polarisation at .1100mV in molten carbonate

Fig. 3. SEM/EDX analysis confirms the presence of oxygen on the surface. X-ray diffraction analysis shows that the corrosion product is NiO. Whcn NiAl is quenched after 4 hours polarisation at potentials of -700mV or more anodic, the surface morphology and the cross section are similar to that of the specimcn quenched at -1100mV (Fig. 2), even though the potential is sufficiently anodic for the nickel oxidation to take place.This indicates that a protcctive aluminium oxide layer has been formed before inserting the specimen into the mclt or in the period just after inserting the specimen into the mclt and before applying the potential. The protective properties of the lithium aluminate scale seem very good.

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J. f? 7:Vossen.A .H. H. Janssen and J. H . W de Wit

Fig. 3. Surface of nickel after 4 hours polarisation at -750mV in molten carbonate.

10.3.2 Polarisation Measurements 4

The potential limits for performing electrochemical experiments are limited by the stability range of the carbonate mixture. The anodic limit is given by reaction 10 while the cathodic limit is given by the decomposition of carbonate to carbon not given here. Rut already at less cathodic potentials we are dealing with the reduction of water to hydrogen, according to reaction 4. The quasi-stationary polarisation curves of nickel and NiAl are given in Figs. 4 and 5. The polarisation curves of these materials are similar.Therefore we will discuss the plot for NiAl by close comparison with Ni. The melt reactions and the oxidation reactions of nickel that correspond with the potential regions in the polarisation curves are:

H,O + CO, + 2e- + H, f C02(-1600 < E < -1070mV) H, t CO,,- -+ H,O + CO, (-1070 < E < -800mV) Ni + NC;, + 2eaccompanied by

+ 2e-

Ni + C0:- -+ NiO + CO, + 2e (-800 < E < -700mV)

(4)

(5)

(6) (7)

Ni + Ni;:, + 2e(8) (dissolution through passive layer; -700 < E < -525mV)

10 The Corrosiori Bekavioiir of NiAl in Molten Carbonate at 650°C

Nf?,:, + NiCi0 + e(-525 < E < 0 mV) accompanied by reaction (6) and (7). CO:- + CO, (E 2 50 mV)

167

(9)

+ Vz02 + 2e-

Since these reactions have been discussed in previous publications, these reactions will only be discussed here briefly [3-61.

2.00 h

N

E

aE 0

0.00

v

.-

-2.00

-4.00I

I

-1.50

-1.oo

0.00

-0.50

E (V vs. 14% 02,30% C02, bal. N2)

Fig. 4. Quasi-stationary polarisation curves of nickel in molten carbonate. Initial potential - 1700 mV, final potential I- 100 mV, 5 mV/step, 50 s/step.

3.00 -

E aE v

.-

Ni-50 Al

2.00 -

1.00 -

0.00

1

-1.oo

-0.50

0.00

E (V vs. 14% 02,30% C02, bal. N2)

Fig. 5. Quasi-stationary polarisation curves of NiAl in molten carbonate. Recorded immediately after immersion. Initial potential: OCP, final potential + 100mV, 5 mV/step, 25 s/step.

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J. I? 7:Vossen.A.H. H. Jnnssen and J. H. W. de Wit

For a discussion of the stability of Ni the potential region of -800mV to 0 m V is most relevant. From -800mV to -525 mV Ni is oxidised and a NiO layer is formed, while from -525 to 0mV some of the Ni'+ is oxidised to Nij' accompanied by a further growth of the oxide layer due to increased ionic and electronic conductivity. Impedance measurements led to the following reaction mechanisms for the formation of the NiO scale and the partial oxidation of Ni'' to Ni3+in the NiO scale respectively [6]. The reaction mechanisms are given schematically in Figs. 6 and Fig. 7. NiO formation

accompanied by the further oxidation of nickel: Ni + C0:-

+ NiO + CO, + 2e

(22)

The NiO layer is not very protective due to the restricted potential region where it is stable. It can be easily reduced cathodically, which is accompanied with detachment of the oxide layer [4]. Because we already know from the quenching experiments that an oxide layer is formed on NiAl at -1100mV, where Ni is not yet oxidised, the anodic current at this potential is ascribed to the oxidation of aluminium. For Ni the current &nearly zero at this potential, due to near-equilibrium of the reactions 4 and 5. PossIMe reactions for the oxidation of aluminium are: 2AI

+ 302- + 2A120, + 6e-

c0;- -+ 0,- + co, Al + Lit + 2CO:- + LiAIO, + 2C0, + 3e-

(23a) (23b) (24)

co:

I

"p' Ni'-O2-Y-l NiZ*

NiO

Fig. 6. Schematic representation of the nickel oxidation and passivation mechanism in molten carbonate.

o

e

e.

Fig. 7. Schematic representation of the reaction mechanism for the transformation of bivalent nickel ions to trivalent nickel within the oxide scale.

The formation of AI,O, can only take place as an intermediate or as a thin interphase because it is not stable in contact with molten carbonate. As a matter of fact the electrolyte tile matrix (LiAIO,) is made in situ by reaction of the carbonate phase with A1,0,. Direct formation of LiA10, according to equation 24 seems possible too. Thermodynamic calculations [12] show that aluminium oxide/lithium aluminate is stable at much more cathodic potentials than -1100mV.Therefore it seems likely that the oxidation reactions proceed at all potentials in the carbonate stability range. Thus some oxide will probably be formed before the polarisation mcasurements are started when the electrode is dipped into the molten carbonate. Thc constant current from -1000mV to -800mV reflects some further but limited growth of the oxide layer, which reaches a thickness of 1 km after 24 hours (see Fig. 2).The crystallite size of the oxide also increases with exposition time especially during the first few hours. After about 4 hours the morphology hardly changes anymore. In this potential region up to -800mV no clear hydrogen oxidation (on Ni from -1070 to -800mV) either dissolved in the metal or dissolved in melt proceeds on the LiAIO, layer. The current is much lower than on Ni and no peak current is observed, which is typical for the hydrogen peak (see also Fig. 4) on Ni due to the limited amount available in the metal [4].This reflects the early protection of the base alloy by LiAlO,, so that no hydrogen is formed. The formation, onrNiO takes place between -800mV and -700mV, as is clearly visible in Fig. 4, while further oxidation to Ni3+takes place around -500mV. Onc would expect that these reactions would not be visible in the polarisation diagram of NiAl, if the oxide layer would be impervious for Ni ions. Judging from the relatively high current density around -1OOOmV, where only LiAIO, has been formed, which is a poor conductor (see Fig. 5 ) it seems justified to conclude that the layer is not blocking ion

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.I.E! 7:Vossen,A. H. N. Jarissen and J. H. W de Wit

transport completely, possibly due to grain boundary diffusion or transport through macroscopic defects. Nevertheless, the fact that no NiO phase can be found by X-ray diffraction on top of the LiAIO, after further anodic polarisation leads to the conclusion that the anodic currents in Fig. 5 above -800mV also reflect the oxidation of the thin Ni wire to which the NiAl electrode was attached and not only the oxidation of Ni ions transported through defects in the oxide layer. Long term behaviour of NiAl in molten carbonate is probably much better than suggested by the short term electrochemical measurements. This is also suggested by the good behaviour of Kanthal-A1 (with only 5.8% of aluminium) at long term i n mersion, which was ascribed also to the formation of an aluminium oxide layer, in that case between the outer Li(Fe,Cr)O layer and the base alloy. Small amounts of Al were found to be beneficial also by Uchida [13]. Polarisation plots of preoxidised Kanthal did not show any anodic peak currents, due to this protective layer. For more details we refer to reference [l].Due to the problems with the unavoidable Ni wire no such experiments could be done on NiAI.

10.4 Conclusions The oxidation product in molten carbonate on Ni is NiO and on NiAl is a-LiAIO,.The latter layer with a thickness of about 1 pm after a few hours exposure either at the OCP or under potential control in a wide range of potentials, provides good protection against further attack, due to its low conductivity. The protection is much better than provided by NiO, which is quite insufficient due to the limited potential stability region. Long term protection by LiAIO, is probably much better than suggested by the current density around -1000mV, due to the further blocking of the fast transport paths as has been shown before for other alloys with a very much lower Al content.

10.5 Acknowledgements Financial support by the Netherlands Agency for Energy and the Environment (NOVEM) and the Netherlands Ministry of Economic Affairs is gratefully acknowledged.

10.6 References 111J.I?T Vossen, L. Plomp, J. H.W de Wit, G. Rietveld J. Electrochem. SOC.142 (1995) 3327.

[2] J.H.W de Wit:in: “Case Studies in Manufacturing with Advancerd materials”, Vol 11. Eds. A. Demaid and J.H.W. de Wit, North Holland, Amsterdam, 1995,pp 137-177. (31 J.E!T Vossen: Ph. D. Thesis: “Corrosion of Separator Plate Constituents in Molten Carbonate”, 1994,Delft University of Technology,The Netherlands. [4] J.PT Vossen, L. Plomp, J. H. W de Wit J. Electrochem. SOC.141 (1994) 3040.

I 0 The Corrosion Behaviour of NiAl in Molten Carbonate

at

650°C

171

[ S ] PC.H. Anzenf:ECN-report number: 93-R-93-027.1993. [6] J.PT Vossen, P C H . Amenr, J.H. M! rfe Wit:“Mechanisms for the Oxidation and Passive Behaviour of Nickel in Molten Carbonate‘., accepted by J. Electrochem. SOC..March 1996. [7] T B . Mussalski (Editor in Chief): “Binary Alloy Phase diagrams”,Vol. I., Am. SOC.for Metals. October. p. 90,142,1986. [S] G. Rielveld: ECN, priv. comm., March 1995. 191 J.R. Selman, H.C. Maru: in: “Advances in Molten Salt Chemistry”, G. Mamantov, ed., Plenum Press, New York, p. 159.1981. 101 K.K. Ghosh: Bull. Electrochem. 7(1991) 512. 111 G.H. Kucerrz: Argonne Natl. Lab., CONF-811014-4. 121 D.A. Shores, P Sitigh: Proc. Electrochem. SOC.84-13 (1984) 271. 131 I. Uchida, 7:Nishinu: 1992 Fuel Cell Seminar, Nov. 29-Dec.2 Tucson,AZ, p 550. Courtesy Associates., Washington. DC. 1992.

Part 111 Fe-Aluminides

Oxidation of Intermetallics by H. J. Grabke, M. Schiitze 0 WILEY-VCH Verlag GmbH, 1998

11 Oxidation of P-FeAl and Fe-A1 Alloys I. Rommerskirchen, B. Eltester and H. J. Grabke

11.1 Introduction

Whereas the oxidation behaviour of P-Ni AI has been investigated under various conditions [l-51, the oxidation bchaviour of the related P-FeAI phase is entirely unknown. Therefore, as a comparison, the oxidation of P-FeAI was studied by means of thermogravimetrical experimcnts in the temperature range from 800°C to 1100°C. Similar phenomena as for NiAl have been observed concerning formation and conversion of A1,0, modifications and occurrence of porosity beneath the scale. At temperatures higher than 650°C the P-FeAI phase is connected to the wFe(A1) phase field without phase boundary. To elucidate the effect of decreasing order with decreasing aluminium content on the oxidation behaviour, samples from the whole range of this phase with 5,10,15,20,25,30,40 and 48 at.% A1 were investigatcd in exposure experiments. The isothermal oxidation behaviour was studied in situ by thermogravimetrical experiments in hclium with 1 vol.% 0, in the temperature range between 800°C and 1100°C (Fig. 1). The p-FeAI samples investigated had the compositions Fe,,AI,, and Fe,Al,. To analyse the transition from external to internal oxidation, exposure experiments in H,/H,O mixtures at very low oxygen partial pressures have been carried out using samples with varying aluminium content between 5 and 48 at.%.The oxygen partial pressurc was controlled with an oxalic acid water saturator to a value below the thermodynamic stability range of iron oxide (Fig. 2). In the thermogravimetrical cxperimcnts the sample is hanging from a microbalance in the temperature constant zone of a vertical furnace (Fig. 1).The microbalance records the mass gain due to oxidation and gives this information to the connected x-twriter. The temperature of the furnace is controlled by means of a Pt-Rhl8 thermocouple. The oxidizing gas is a mixture of commercial helium and oxygen which are dried by P,Os columns and then controlled by a capillary flowmeter to a total flow velocity of 2 ml/s. To analyze the transition from external to internal oxidation dependency on aluminium content, it is important to establish a small oxygen partial pressure which allows only aluminium oxide growth and not iron oxide growth. Therefore, commercial hydrogen gas was cleaned from rest oxygen by a Cu catalyst and then dried by a P,O, co-

176

I. Rornmerskirctieii, B. Eltester und H. J. Grabke 1

exhaustpipe

7

mnba Row

Fig. 1. Flow apparatus for thermogravimetrical investigations 1.P,O, column, 2. capillary flowmeter, 3. three step cock, 4. thermocouple, 5. sample holder with sample, 6. resistance furnace, 7. electronic microbalance, 8. x-t recorder

Fig. 2. Experimental equipment for exposure experiments at very low oxygen partial pressure 1.active Cu catalyst, 2. P,O, column, 3. oxalic acid water saturator, 4. capillary flowmeter, 5. fur. nace

lumn (Fig.2). A defined concentration H 2 0 is added to the H, gas while flowing through an oxalic acid water saturator containing a mixture of 90 % C,H20, . 2 H,O and 10% C,H,O,. Corresponding to the dissociation equilibrium of H,O at the given furnace temperature, a very low oxygen partial pressure in the range lo-" to 10 bar is attained in the reaction atmosphere (Table 2).

I I Oxidation of P-FeAI and Fe-AI Alloys

177

Table 1. Composition of the investigated alloys and iron aluminides Itd. chemical analysis A1 wt.O/,

Fe wt.%

2.47 5.12 7.86 10.9 14.0 16.6 23.7 30.3

98.1 95.4 92.4 89.1 86.0 83.0 76.2 69.2

Ni

C

S

0

wt.%

wt.%

wt.%

ppm

0.0018

49 81 81 81 100 92 127 34

18 at% Al to H,S-H,-H,O mixed gases, but more resistant than conventional iron-chromium-aluminium alloys. Others have found that the corrosion rates of iron-chromium-aluminium alloys in sulphur-rich gases decrease with increasing alloy aluminium concentration [ 121. However, is has also been reported that additions of only 2 % Cr to iron aluminides containing 26 to 28 at% Al can be detrimental to sulphidation resistance [13], although tests in that research involved higher temperatures and gases of higher H,S concentrations than the present work. One aim of this research has been to investigate the effects of yttrium on the effectiveness of the A1,0, scales in giving protection, based on the well-known beneficial effccts of this element in improving the mechanical integrity and spallation resistance of such scales during high-temperature oxidation. However, in thc present tests, thc thickness of the A1,0, scales were only of the order of 200 to 300 nm and scale spallation was not a factor for any of the iron aluminides. Nonethelcss, the two alloys containing the most yttrium, FA 57 and FA 58, developed extensive intergranular precipitates in the substrate. These precipitates were rich in both zirconium and yttrium. The causes of such prccipitation are unclear at present, but its occurrence is much more damaging than the small extents of sulphide formation in the tests. Although M A 956 and Fecralloy are essentially Al,O,-forming alloys at high temperatures, it is interesting that the scales developed during preoxidation and during exposure in environment 1 at 700°C were not entirely AI,O,; some areas were covered by a more granular Cr,O, scale. Presumably this is because the rates of growth of the initial AI,O, nodules are very low at the relatively low temperature and low oxygen potentials of the tests while the rate of aluminium diffusion to the surface is also low at this temperature, enabling a Cr,O,-rich scale rather than an Al,O,-rich scale to become established in local regions on the surface. O n exposurc to environment 2, it is also difficult for an Al,O,-rich scale to develop and faster growing sulphides are able to become established rapidly, particularly on M A 956. The relatively poor performance of 310 SS and Alloy 800 in the mixed gases is consistent with the widely-reported susceptibility of Cr,O,-forming alloy t o sulphidation in such environments. The importance of a high temperaturc for establishment of a protective oxide scale is also illustrated by the results in environment 3 at 500°C. Mere, an effective A1,0, scale was not able to be established fully on either FA 56 (containing 8.1 YOA1 and 5.0 YoCr) or M A 956 (containing 4.5 YOAl and 20.0 YOCr) and the weight gains were relatively large, consistent with formation of significant amounts of sulphide. More effective scales were established on the other iron aluminides, containing greater concentrations of aluminium, and the weight gains were relatively low.

230

E H. Stott, K . 7:Chuah and L.B. Bradley

Generally, this research has shown that iron aluminides can have good resistance to mixed gases of high sulphur and low oxygen potential. However, the results also indicate that, for alloys containing 5 YOCr, aluminium concentrations above 10 wt% give better protection than those between 8 and 10 wt%.As alloys containing less than 8 or 9 wt% A1 have much higher room temperature ductilities than those containing over 10 wt% Al,it is apparent that further development of such alloys could involve a trade off between mechanical properties and sulphidation resistance, with emphasis on the effects of other minor alloying elements, such as the yttrium used in the work, on overall performance.

14.5 Conclusions 1. Iron aluminides containing 5 wt% Cr, 0.2 wt% Zr and 8 to 16 wt% A1 show better resistance to a mixed-gas, HJ1.5 %H,S/4.2 %H,O environment at 700°C than conventional A1,0,- and Cr,O,-forming alloys, with the extent of sulphide formation for the aluminides increasing with decreasing aluminium concentration in the substrate. 2. Under these conditions, the presence of 0.2 to 0.3 wt% Y in the iron aluminides has little influence on the effectiveness of the thin A1,0, scale in giving protection; however, extensive intergranular precipitates, rich in yttrium and zirconium, are formed in alloys containing 0.3 wt% Y. 3. It is more difficult to establish a protective A1,0, scale at lower temperatures than at higher temperatures, particularly on alloys of relatively low aluminium concentrations; thus, an iron aluminide containing 8 wt% A1 formed more extensive sulphides at 500°C than at 700"C, although a protective oxide scale prevented significant sulphidation of an alloy containing 12 wt% Al at either temperature.

14.6 Acknowledgements The experimental work was undertaken in MSc dissertation programmes (by KTC and LMB), sponsored by Shell Research, Arnhem. The authors are also grateful to Shell Research, particularly ?: Wolfert,for providing the iron aluminides and for useful discussions.

14.7 References [l] C. G. McKarney, J. H. DeVan, I! E Tortatelli, V K . Sikka: J. Mater. Res.6'(1991) 1779. [2] C. M . Packer, R. A. Perkins: Roc. Conf. on High-Temperature Alloys in Aggressive Environments, 813,The Metals Society, London (1980). [3] E H. Stott, E M. E Chong, C. A. Stirling: Materials Science Forum 43 (1989) 327. [4] S. W Green, E H. Stotf:Corros Sci. 33 (1992) 345.

14 Oxidation-Sulphidationof Iron Aluminides

[S] [6] [7] [8] [9) [lo] [Ill 1121 [13]

S. W.Green, E H. Stott: Oxidation Metals 36 (1991) 239.

231

M. H. La Branche,A. Garratt-Reed, G.J. Yurek:J. Electrochem. SOC.130 (1983) 2405. C. R. Wang, %. B. Zhao, S. K . Xia, W Q. Zhang: Oxidation Metals 32 (1989) 24. E H. Stott, E M . E Chong, C. A . Stirling:High Temperature Corrosion in Energy Systems, ed. M. Rothman, 253, Met. SOC.AIME, New York (1985). P A. Mari, J. M. Chaix,J. I? Larpin: Oxidation Metals 17(1982) 315. J. M.DeVan, P E Tortorelli:Corros. Sci. 35 (1993) 1065. b! K. Sikka: Proc. 6th Annual Fossil Fuel Energy Conf. ORNI.IFMP-92Il. 195, Oak Ridge Nat Lab,Oak Ridge.TN (1992). S. Mrowec, M . Wedrychowska: Oxidation Metals 13 (1979) 481. J. If. DeVan: Oxidation of High-Temperature Intermetallics, cd. T. Grobstein and J. Doychack, 107.Minerals, Metals and Materials SOC.(1989).

Oxidation of Intermetallics by H. J. Grabke, M. Schiitze 0 WILEY-VCH Verlag GmbH, 1998

15 Metal Dusting of Fe,AI and (Fe, Ni),AI S. Straufl, R. Krajak, M. Palm and H.J. Grabke

15.1 Introduction

Sevcral new materials have been dcveloped on the basis of the aluminides Fe,AI, Ni,AI (Fe, Ni),AI, alloyed with Cr and doped with small additions of Zr or Hf [14]. These materials proved to be rather resistant against high temperature corrosion in oxidizing, sulfidizing and carburizing environmcnts. The carburization resistance was tested in CH,-H, atmospheres at high temperatures and carbon activities a, < 1 [4]. However, at intermediate temperatures 400-900°C and carbon activities a, > 1 there is another dangerous corrosion phenomenon ‘metal dusting’ which leads to disintegration of Fe-, Ni- or Co-base materials to a dust of fine metal particles and carbon (graphite) [5-141. Low alloy steels show general metal wastage and high-alloy materials show pitting and holc formation. It was hoped that the M,Al based materials would be resistant against metal dusting and might be used as coatings or construction materials under ‘metal dusting’ conditions. Thus, their metal dusting resistance was testcd under standard conditions.

15.2 Experimental The M,Al based materials were exposed in flowing CO-H,-H,O mixtures at 650°C for different periods of time. Corresponding atmospheres are used in many processes attained from methane conversion for use in synthesis of ammonia, hydrocarbons, alcohols etc. or in the direct reduction of iron ores. The materials tested were (concentration in wt%): 1. 2. 3. 4. 5.

Fe,AI Fe,Al-2.2% Cr Fe-16 %AI-4.8%Cr-0.15 %Zr Ni,Al Ni-26 %Fe-10 %A1-8 %Cr-0.1 %Hf

The mass of thc samples was determined before and after the exposure, and after removal of the corrosion product, the ‘coke’ which had grown on the specimens, locally from pits or on the whole surface.

S. StrutiJ3,R. Krujuk, M . Pulni and H.J. Grabke

234

15.3 Results All M,AI aluminides containing Fe were attacked by metal dusting, most severe attack was observed on Fe,AI (Fig. 1).Vast protrusions of coke were growing from pits on Fe,AI, from the material (2) with 2.2% Cr only a few protrusions grew from the unpolished edges, and on the material ( 3 ) only a thin coke layer was observed without pitting. In contrast, the pure Ni,AI showed only a thin carbon layer on the unpolished edges. Some data on the mass changes by coke growth (removable), metal wastage resp. oxidation and carburization (remaining weight change) are compiled in Fig. 2.

Fig. 1. Metal dusting attack on Fe,AI, vast outgrowth of coke after 7 d exposure at 650°C and a, = 10,the arrow indicates the sample from which the coke has grown.

1000

time: Id, 2d

Mo

g

10

2 .-C

%

1 0.1 0.01

Y) VI

E 0.001 -0.01 -0.1

-1

Fig. 2. Mass changes of four alloys in exposures at 600°C and a, = 41.5for one day and a second day. black columns: removable coke and hatched columns: remaining mass change by metal wastage (loss), carburization and oxidation (gain).

I5 Meral Ditsrirrg of Fe.,Al and (Fe. Ni).,Al

235

The metal dusting of Fe and low alloy steel proceeds via intermediate formation of the instable carbide Fe,C [6,7,8] which then decomposes according to Fe,C -+ 3 Fe + C. For Fe,Al an intermediate formation of the K-phase Fe,AIC [lS] might occur, which also would be instable if formed at a, > 1.Therefore. specimens of Fe,AI were investigated after metal dusting for 48h at 650°C: and a, = 1.5, using metallography, scanning electron microscopy with EDAX and X-ray structure analysis to search for the K-phase. However, no indication of its presence was found but clear evidence of the presence of Fe,C. This instable carbide was observed as a thin layer on thc Fe,AI, obviously growing from the inside and decomposing to the outside, as in the case of metal dusting of iron and steels [6,7,8].The A1 obviously is oxidized under formation of A1,0, but this oxidation does not lead to a protective scale. Protection appears to be improved by alloying with Cr, whereby the metal dusting attack is decreased. Ni,AI is not attacked by metal dusting, but materials containing Fe show pitting after prolonged exposures (see Fig. 3).

Fig. 3. Pitting by metal dusting of a Ni-base aluniinide, after 330 h at 650°C. and a, = 20.

236

S.StrauJ3, R. Krajuk, M. Palm and H.J. Grabke

15.4 Conclusions Fe,AI based materials are susceptible to metal dusting. In CO-H,-H,O atmospheres such materials disintegrated to a ‘dust’ of iron, carbon and alumina. By alloying with Cr the attack is decreased somewhat.The disintegration proceeds via intermediate formation of the instablc carbide Fe,C, as in the case of iron and low alloy steels Ni,AI is not susceptible to metal dusting, however, aluminides (Ni,Fe),AI containing Fe are subject to metal dusting. After long-tcrm exposures such materials show pitting such as high alloy steels [7].

15.5 References [l] J. H. DeVan, C. A. Hippsley: in: Oxidation o f High Temperature Internietallics. Ed.T. Grohstein. J. Doychak,The Minerals, Metals and Matcrials SOC.1989,p. 31 and p. 107. [2] J. H. DeVan: in: Heat-Resistant Materials, Proceedings of the First Int. Conference, Fontana. Wisconsin, Sept. 1991. Eds. K. Natesan, D.J.Tillack ASM 1991, p. 235. [3] J. H. DeVan, P E Tortorelli: Materials at HighTemperature 11 (1993) 30. [4] J. KMwer:Workshop ‘Oxidation of Intermctallics’, in press [5] J. C. Nava Paz, H. J. Grabke: Oxid. Metals 39 (1993) 437. 161 H . J. Grabke, R. Krajak, J. C. Nava Paz: Corrosion Sci.35 (1993) 1141. [7] H. J. Grabke, R. Krajak, E. M. Miiller-Lorenz:Werkst. Korros. 44 (1993) 89. [ 8 ] H. J. GraDke, C. B. Bracho-Troconis, E. M. Miiller-Lorenz:Werkst. Korros. 45 (1994) 215. [9] H. J. Grabke, R. Krajak: Harterei-Techn. Mitteilungen 49 (1994) 150. [lo] H. J. Grabke, E. M.Miiller-Lorenz, 1;. Pippel, S. StrauJ3: Conf. Papers U K Corrosion & EUROCORR ‘94,Vol.3, Chameleon Prcss Ltd. London 1994,p. 361. [ l l ] H. J. Grabke: Corrosion NACE 51 (1995) 711 [12] E. Pippel, H.J. Grabke, S. StrauJI, J. Woltersdorf: steel res. 66 (1995) 217 [13] H. J. Grabke: Solid State Phenomena 41 (1995) 3. [14] H. J. Grabke, E. M. Muller-Lorenz: steel res. 66 (1995) 254. I151 M. Palm, G. Inden: Intermetallics3 (1995) 443.

Part IV Ti-Aluminides

Oxidation of Intermetallics by H. J. Grabke, M. Schiitze 0 WILEY-VCH Verlag GmbH, 1998

16 Determination of Thermodynamic Activities in the Alloys of the Ti-A1 System and Prediction of the Oxidation Behaviour of the Alloys M. Eckert and K. Hilpert

Titanium aluminide alloys based on Ti,Al and TiAl are of interest as construction material for high temperature components particularly in aerospace industry. Good mechanical properties can be attained with alloys consisting of y-TiAl with 3 to 15 vol% a2-’fi3A1.The disadvantages are the low ductility and the inadequate oxidation resistance at service temperatures of 700-900°C [l].A fundamental understanding of the oxidation behaviour is necessary in order to improve the corrosion resistance. The formation of the oxides on the alloy surface depends on the temperature, the oxygen partial pressure of the corrosive atmosphere, and the thermodynamic activities of Ti and A1 in the alloys. The measurement of the thermodynamic activities of Ti and Al in the alloys of the Ti-AI system is the aim of this study in order to render possible an improved understanding of the formation of the various oxides. The measurements of the thermodynamic activities of Ti and Al in Ti/AI alloys carried out so far are insufficient.They were essentially conducted by Hoch and Usell [2] using Knudsen effusion mass spectrometry and by Samokhval et al. [3] using e. m. f. measurements. Hoch and Usell have measured the thermodynamic activities of Ti and Al at 1780K in the @-Tiphase for the composition range 0 < xA, 5 0.4. Samokhval et al. measured the A1 activities at 960 K in the composition range 0.05 5 xA, 5 0.5. Recently,Jacobsen et al. [4] reported A1 pressures for three samples of the phase compositions %,A1 + TiAI, Ti-0.52A1, and TiAl + TiAI, obtained by Knudsen effusion mass spectrometry. The thermodynamic studies by Hoch and Usell as well as Samokhval in addition to many other investigations [5-131 are the basis for the recent calculation of the phase diagram of the Ti-A1 system by Katfneret al. [14]. The present study was carried out to determine for the first time the thermodynamic activities of Ti and Al over the complete composition range 0 < xA, < 1 in the temperature range from 1200 to 1600K. A detailed presentation of the results obtained will be reported in Ref. [15].The study reported here is part of our systematic investigations on the influence of ternary addiditives to the Ti/AI alloys. Vaporization studies by Knudsen effusion mass spectrometry are used for the thermodynamic investigations. The results obtained by this method for many alloy systems are summarized in a review article by Hilpert [ 161. ‘The measurements reported in this work complement

M . Eckert and ti.Hilpert

240

our investigations on the alloys of the Ni-A1 system [ 17,181 including the intermetallic phase Ni,AI [17,19]. Nineteen Ti/AI alloy samples with different compositions covering the complete composition range of the Ti-AI system were prepared by inductive melting in a purified argon atmosphere using the levitation technique (Fig. 1). The oxygen content of the alloys was additionally determined after the vaporization measurements, because of the known high oxygen solubility of Ti. The 0, concentrations in ppm given in parentheses were obtained in the alloys with Al concentrations of 0.05 (759 i 98). 0.10 (1119 -C 54), 0.15 (2766 2 31), 0.20 (2340 -C 162), 0.25 (708 % 59), 0.30 (861 75). 0.35 (17610 2 1550),0.40 (571 % 11),0.45 (957 % 168),0.47 (461 % 48),0.50 (673 ? 28),0.55 (165 ? 2), 0.65 (65 2 15), 0.69 (32 2 15), 0.725 (31 -C l l ) , 0.80 (39670 -? 26680), and 0.90 (329 % 183). The measurements were carried out with the substantially modified Knudsen cell mass spectrometer system described in Ref. [20].The Knudsen cell mashined from Mo was completely lined with yttria.The geometry of the Knudson cell with lining was the same as described and depicted in Ref. [17]. Ti and A1 partial pressures were determined in the temperature range of the mcasurements (cf. Fig. 1) over the different samples. Thermodynamic activities of Ti and Al resulted according to the relation a, = p, / pv (i = Ti,AI) by comparing the pressures over the alloy samples, p,, with those of pure Ti(s) and AI(I), py.The results obtained at the mean temperature of 1473K are given in Fig. 2. The thermodynamic activity of Ti was, additionally, evaluated from the measurement results by the ion intensity ratio integration method. The results obtained by this method are also shown in Fig. 2. The two different evaluation methods agree in general very well. Knowledge of the thermodynamic activities of Ti and A1 in the alloys is essential for predictions on the stability of oxide scales. Such predictions were carried out so far by -f

T

-

I

-

I

-

t 2000

1800 1600 Y

;1400 1200 1000 800 0 Ti

02

0.4

06 xAl

08

1 Al

Fig. 1. Phase diagram by Kattner et al. [14] with the composition of the prepared samples and the temperature ranges of the vaporization measurements

16 Determinniiori of ThermodyaniicAciiviiies in the Alloys of the Ti-A1System

241

1

0.9 0.8

0.7

c ‘

< .,”

G

0.6

05 0.4

0.3 0.2 0.1 U I

XAl

L

*

Fig.2.Thermodynamic activities of X and A1 at 1473 K in solidTi/Al alloys of differcnt compositions obtained by the ion intensity ratio method ( 0 )and by the equation a, = pi / pi0 (+,0)

Rahmel and Spencer [21] as well as Lurhra [22]. Only estimated thermodynamic activities were, however, used by these authors. The computations by Rahmel and Spencer as well as Lurhru are repeated in the following by employing the thermodynamic activities determined in this work. The oxygen partial pressures of the reactions 3 2A1 + 2 O,(g) w A1,0, (s) and (1) 1 Ti + 2O,(g) w TiO(s) (2)

were computed.The computation was carried out at 1373K in order to render possible a comparison with the results obtained by Rahmel and Spencer [21]. The result of this computation is shown in Fig. 3. It is obvious, that the oxygen partial pressures belonging to Eqs. (1) and (2) intersect at an A1 concentration of about xAl= 0.54. Rahmel and Spencer [21] determined the Al concentration for this intersection as xAl = 0.61, while Lirthra [22] obtained xA, = 0.55 for this intersection at 1073 K. This means, that the results of this work lead to the prediction of a stable A1,0, scale for lower A1 concentrations than by Rahmel and Spencer [21]. Moreover, the data obtained by us show that this A1 concentration decreases with decreasing temperature. Dissolution of oxygen in the alloy was neglected in our computations and in the predictions by Rahmel and Spencer. Work is, therefore, in progress to elucidate the influence of oxygen dissolution o n the thermodynamic activities ofTi and Al.

242

M. Eckerf and K. Hilpert

Fig.3. Variation of the metal/oxide equilibrium pressures in the Ti-A1-0 system at 1373 K

Acknowledgements The authors thank the Deutsche Forschungsgemeinschaft for financial support within the program ,,High Temperature Corrosion“.

16.1 References [I] A. Rahmel, WJ. Quadakkers, M. Schiitze: Materials and Corrosion 46 (1995) 271. 121 M. Hoch, R.J. Usell:Met.Tfans. 2 (1971) 2627. [3] K V Samokhval, RA. Poleshchuk,A.A. Vecher: Russ. J. of Phys. Chem. 45 (1971) 1174. [4] N S . Jucobsen, M.I? Bra$y, G.M. Mekrota: 188th Meeting of The Electrochemical Society. Chicago, Oct. 1995. Symp. on “High Temp. Materials/Corrosion” - High Temp. Mat. Chemistry 7. exten&&bstract. (51 K. Shibata, T Sato, d Dhira: J. Cryst. Growth 27 (1978) 1329. [6] E. W Collings: MetaKTrans. IOA (1979) 463. [7] A. Abdel-Humid, C,H.Alliberf, F: Durand Z. Metallkd. 75 (1984) 455. IS] R.D. Shul1,A.J. Me-Alisfer,R.C. Reno:Titanium Sci.Tech.3 (1985) 1459. [9] R.M. Waterstrat: NISlIR 88-3856, United States Department of Commerce, (1988). [lo] K. Kalfenbach, S. Gama, D.G. Pinatti, K . Schu1ze:Z. Metallkd.BO(1989) 511. [ l l ] C. McCullorcgh,J.J. Vulenciu, C.G. Levi, R. Mehrabian: Acta Metall. 37 (1989) 1321. [12] J.C. Schuster, H. Ipser: Z. Metallkd.81 (1990) 389. [13] J.C. Mishurda, Perepezko: in: “High-TemperatureAluminides and Intermetallics”, edited by S.H. Whang et al. (TMS, Warrendale, Pennsylvania, 1990). [14] U R .Katfner, J.C. Lin, Y A . Chang: Met.Trans.23A (1992) 2081.

16 Delerniinntion of Tlzerniori),namicActivilies in the Alloys of the Ti-A1 System

243

115) M. Eckerr. L. Bericze, D. Knth. H. Nickel, K . Hilperl: Ber. Ihnsenges. Phys. Chem., in press. [ 161 K . Ifilpert:Structure and Bonding 73 (1990);Springer.Heidelberg. [17] K. Hilput, D. Kobertz, V: Veriugopnl, M. Miller, H. Gerads9 EJ. Bremer, H. Nickel: Z. Naturforsch.42a (1987) 1327. [18] K . Hilpert, M. Miller, H. Gerads, H. Nickel Ber. Bunsenges. Phys. Chem. 94 (1990) 40. [19] M. Albers, M. Sai Baba, D. Kath, M. Miller, K . Hilpert: Ber. Bunsenges. Phys. Chem. 96 (1992) 1663. [20] K. Hilpert: Habilitationsschrift,Technische Hochschule Darmstadt 1981; Report from the KFA Jiilich. Jul-1744,p. 28. [21] A. Ruhnrel, PJ. Spencer: Oxid. Met.35 (1991) 53. [22] K.Z.. Lirflira:Oxid. Met.36 (1991) 475.

Oxidation of Intermetallics by H. J. Grabke, M. Schiitze 0 WILEY-VCH Verlag GmbH, 1998

17 The Initial Stages in the Oxidation of TiAl C. Lang and M. Schiitze

17.1 Introduction The high temperature oxidation behaviour of titanium aluminides has been investigated by several authors in recent years [l-51. It is well known from these studies that in most cases the titanium aluminides based on y-TiAl do not form an protective A1,0, oxide layer and therefore exhibit an insufficient oxidation behaviour at temperatures above 800°C. Instead of a dense Al,O, layer the oxide scale usually exists of a layered structure TiO2//AI20,//TiO2 + A120, mixture//Al depletion zone from the outside to the inside.The A1,03 in the second layer from the outside provides protection as a barrier layer during a certain time of the oxidation process, but is not stable for longer oxidation times [4]. Often a better oxidation resistance of y-TiAI based titanium aluminides is observed if the oxidation tests are carried out in pure oxygen instead of air, since the absence of nitrogen in the oxidizing environment has a beneficial effect on oxidation kinetics in most cases [S-71. Also thc addition of ternary elements can improve t h e oxidation behaviour of TiAl [8]. One of the most effective elements to decrease the oxidation rate is niobium, as several investigations have shown, e. g. [l, 9-11]. In spite of this extensive number of results on TiAl oxidation the mechanisms are not yet understood in detail. In particular a lack of information exists on the processes taking place in the early stages of oxidation. Because it is to be expected that transient oxidation at the beginning influences the oxidation behaviour in the later stages it is the aim of this paper to elucidate the mechanisms of the initial stages of TiAl oxidation with respect to long-term behaviour.The influence of nitrogen and the effect of niobium are considered in particular.

17.2 Experimental Procedures The experiments were carried out with 'K36A1 (mass-%) and Ti35A15Nb (mass-%). The materials were produced by vacuum melting and casting, leading to a microstructure consisting of a y-TiAl matrix with embedded two phase colonies of lamellar.a,Ti,A1 and y-TiAl.The overall content of a2-Ti3AIwas very small. In Table 1 the chemical composition of the investigated materials is given.

246

C. Lung and M.Schiirze

Table 1. Chemical composition of thc materials investigated

(in mass-%)

Material

Ti

Ti36A1 Ti3SA15Nb

62.30 37.25 59.25 35.65 4.32

-.

A1

Nb ~

C

H

N

0

0.009 0.003 0.011 0.31 0.039 0.004 0.008 0.34

Specimens with the dimensions of 10 X 1.5 X 1mm were cut from cast rods and wcre ground with S i c paper to a 4000 grit surface finish. Beforc the oxidation tests the samples were cleaned ultrasonically in ethanol. Two kinds of oxidation experiments were carried out: isothermal exposure tests at tcmperatures of 800 to 1000°C:for 0.5 h up to 4 h and continuous mass change measurements at 900°C, both in laboratory air. To ensure rapid heating up the exposure tests were started by introducing the samples directly into the hot furnace. The same test conditions were used for the continuous thcrmogravimetric measurements by raising the hot furnacc around the specimen chamber. Because of the expected small dimensions of the oxidation products and of the oxide scale after short oxidation times it was necessary to use electron microscopic methods to characterize the samples after oxidation. Besides scanning electron microscopy (SEM) transmission electron microscopy (TEM) was mainly used to describe the morphology of the oxide scale and to identify the oxidation products by energy dispersive X-ray analysis (EDX) as well as electron diffraction. A detailed description of the preparation of TEM cross-sections and of the experimental procedure is given in [ 121.

17.3 Results 17.3.1 Thermogravimetric Measurements In Fig. 1 the weight gain curves of Ti36A1 and Ti35A15Nb obscrved in the continuous thermogravimetric measurements at 900°C in air are presented. As it is known very well in the literature (see beforc) the oxidation resistance of niobium containing titanium aluminides is much better compared to that of niobium free TiA1. This is also found in the early stages of o3idation. After 25 h of oxidation the relative mass gain of Ti36A1 is distinctly higher thah that of Ti35AlSNb. The oxidation behaviour of both materials is characterized by a strong increase in weight during the very beginning of oxidation.This stage with a very high oxidation rate lasts up to about 1 to 2 h and is referred to as stage I. After stage I the oxidation rate decreases strongly and an almost linear oxidation behaviour is observed. This second stage of oxidation ends after about 10h for Ti36Al and after about 20 h for Ti35A15Nb when the oxidation kinetics start to obey a parabolic rate law [12].The different oxidation rates in the early stages of oxidation can be clearly observed in a parabolic plot of the mass gain which is shown in Fig. 2. Due to the absence of a linear dependence of the square of mass gain on oxidation time at the beginning it is evident that TiAl oxidation does not follow parabolic kinetics in the early stages.

I7 The Initial Stages in the Oxidation of TiAl

2,0r

9OO0C, air

t

Ti36A1

r

u,u

0

247

I I 10

5

20

15

25

tinh

Fig. 1. Mass gain vs. oxidation time of Ti36AI and Ti35A15Nb at 900°C in air.

2.0

1 m

1.5-

6 . t

N

0,

1.0 -

E

.-c N

a

E

0.5

0.0

0

1

2

3

4

5

6

7

8

9

101112 131415

t in h

Fig.2. Parabolic plot of the mass gain of Ti36AI and Ti35A15Nb at 900°C in air.

17.3.2 Characterization of the Oxide Scale and the Subsurface Layer 17.3.2.1 Ti36A1, 9OO0C, 0.5 h In Fig. 3 a TEM micrograph and a schematic illustration of the complete oxide scale and of the subsurface layer of Ti36AI after oxidation at 900°C in air for 0.5 h is presented. The oxide scale consists of an oxide mixture of AI,O, and TiO, in the outer part of the scale. Beneath this oxide mixture a porous partial layer rich in TiO, is found. At

248

C. Lang and M . Schiitze

Fig.3.TEM micrograph and schematic illustration of the oxide scale and of the subsurface layer of Ti36A1 after oxidation at 900°C in air for 0.5 h.

the metal/oxide interface AI,O, and Ti-nitrides are formed, which penetrate into the A1 depleted metal subsurface 1ayer.This is also shown at a higher magnification in Fig. 4. Because of the small dimensions of the oxides and nitrides at the metalioxide i n k t face an unambiguous identification of the phases is not free of problems. Several electron diffraction patterns indicate, however, the existance of a metastable AI,O, or an aluminium oxynitride similar in structure to Al,,O,,N [13]. Furthermore TIN is found at the metal/oxide interface.The analysis of the metal subsurface layer by electron diffraction (Selected Area Diffraction (SAD)) and energy dispersive X-ray analysis (EDX) revealed that the Ti-rich phase of this zone, which contains approximately 62 at.-% Ti and 38 at.-% Al, is neither a,-Ti,Al nor y-XAl. Although there is no additional phase known from the binary phase diagram of the system TiAl[14] several authors reported of a new phase rich in titanium [4, 151. The composition of this new phase which was identified as a cubic structure (161 was notified to be similar in composition to Ti,Al, [17], Ti,AI [18, 191 or Ti,A1,0, [20] and the lattice parameter was given in the range of 0.66 to 0.69nm. The indices of this new cubic phase (NCP) fit very well to different SAD patterns of the same grain of the Ti-rich subsurface layer. An example of a SAD pattern of the new cubic phase is given in Fig. 4.

17.3.2.2 Ti36AI, 9OO0C,4 h Figure 5 shows the backscattered electron (BSE) image and element distribution maps of the oxide scale on Ti36A1 after a 4 h exposure at 900°C in air.'The structure of the oxide scale is already similar to that observed after long-term exposure, cf. e. g. [4]. Beneath an outer TiO, layer a partial layer enriched in A1,0, had formed which can be taken as a preceeding step to the formation of the Al,O,-barrier observed later on [4]. The inner part of the scale consists of a mixture of A1,0, and TiO,. An extensive

I7 The Initid Stciges in the Oxidation of TiAl

249

Fig.4. a)'TEM micrograph of the metal/oxide interface and metal subsurface layer of Tl36A1 after oxidation at 900°C in air for 0.5 h. b) SAD pattern of the new cubic phase.

depletion zone in the metal subsurface is not observed. This is confirmed by TEM analysis of the metal subsurface zone, the melal/oxide interface and of the inner oxide scale. As can be seen in Fig. 6 beneath a porous, Ti0,-rich partial layer Tknitrides and aluminium oxide are observed at the metauoxide interface. Only small grains of a Tirich metal phase similar in composition to Ti,AI can be detected by EDX [12].A more detailed analysis of the aluminium depleted layer by electron diffraction revealed that the new cubic phase yields a satisfactory indexing of the SAD patterns. The nitrides which form at the metal/oxide interface are found to be Ti,AIN and TiN, whereas the aluminium oxide at this part of the oxide scale (contrary to the outer part of the scale) is not a-Al,O, [12]. A series of SAD patterns of the same aluminium oxide grain at various sample orientations shows that AI,,O,,N is formed (see Fig. 6) at the mctal/oxide interface.

250

C. Lung and M . Schutze

Fig.5. BSE image (a) and element distribution (b: oxygen, c: aluminium. d: titanium) of the oxide scale and mctal subsurface layer ofTi36A1 after oxidation at 900°C in air for 4h.

17.3.2.3 Ti36A1, 8OO0C,4 h After 4 h oxidation of Ti36Al at 800°C in air an outer oxide scale consisting of a mixture of X O , and A1,0, is also found as in the cases describcd before. Beneath this outer mixture a porous titania rich partial layer is observed (see schematic illustration in Fig. 7). At the metal/oxide interface aluminium oxide and Ti-rich phases are present similar to the sample oxidized at 900°C.From Fig. 8 which shows the metal/oxide intcrface it is obvious that due to the small grain size and the irrcgular appearance it is difficult to identify the phases developed at this site. In some regions where a high Ti content is found Ti,AIN is present according to the SAD patterns analysed. For SAD patterns, which cannot be assigned to Ti,AlN, other Ti-rich phases like Ti,AI or the new cubic phase may occur, but these could not be ascertained. The aluminium oxide at the metal/oxide is found to be A127039Nas at 900°C.

I7 The Initid Stages in the Oxidation of TiAl

25 1

Fig.6. a) TUM micrograph of the rnetal/oxide interface ofTi36A1 after oxidation at 900°C in air for 4 h. b) SAD patterns of AI,,O,,&

In some parts of the sample a certain featurc of thc metal/oxide interface is striking. As it is shown in Fig. 8 a necdle shaped grain is pcnetrating through the metal/oxide interface layer into the base metal. The analysis of this phase shows that Ti,O, is formed in this area of thc metal/oxide interfacc.

252

C. Lung and M. Schiitze

TiAl

Pig.7. Schematic illustration of the oxide scale and of the subsurface layer of'fi36AI after oxidation at 800°C in air for 4 h.

17.3.2.4 Ti36A1, 1000°C, 4 h Due to the higher temperature a much thicker oxide scale is formed on'I'i36Al during a 4 h oxidation at 1000°C compared to 800°C and 900°C. In Fig. 9 a BSE imagc of the cross-section of the oxide scale and of the base metal is presented. While the layered structure of the oxide scale is in principle the same as after oxidation at 900°C for 4 h the metal subsurface zone shows a significant difference. Contrary to an oxidation temperature of 900°C the bright layer beneath the oxide scale indicates the cxistance of an aluminium depletion zone. This is confirmed by the clement distribution images (see Fig. 9). In addition in the BSE imagc the depletion layer appears not to be continuous, but interrupted by a second phase. A more detailed analyses by TEM revealed the presence of single a,-Ti,AI grains which occur at equal distances at the interface (see Fig. 10) [12].They are surrounded by titanium nitrides and aluminium oxide with the latter also growing along the grain boundaries into the metal subsurface layer.'I'he phase identification of the nitrides yields the presence of TiN, whereas the SAD patterns obtained for the aluminium oxide indicate the formation of Al,,O,,N. Beneath the Ti,AI grains a layer with an aluminium content between X3AI and TiAl is obscrved.The analysis of this layer by electron diffraction yields the same result as after 0.5 h of oxidation at 900"C.The suggested new cubic phase fits very well to the diffraction patterns.The res'yks are summarized in a schematic illustration in Fig. 11.

17.3.2.5 Ti35A15Nb, 9OO0C,4 h In Fig. 12 a TEM micrograph and a schematic illustration of the oxide scale and o f thc metal subsurface zone of Ti35A15Nb after 4 h of oxidation at 900°C is shown. Beneath an outer oxide mixture of AI,O, and XO, a layer of coarse-grained XO, is observed.

I7 The Initial Striges in the Oxidation of 7iAl

253

Fig.8.TEM bright field image of the metal/oxide interface of Ti36A1 after oxidation at 800°C in

air for 4 h.

Towards the inner oxide scale a porous layer containing 30, and titanium nitrides (TiN, Ti2AIN) follows. At the metal/oxidc interface Al2,0,N is predominantly formed as results from clcctron diffraction indicatc. Contrary to the observations with Ti36A1

254

C. Lung rind M . Schiirze

Fig.9. BSE image (a) and element distribution (b: oxygen, c: aluminium, d: titanium) of the oxide scale and metal subsurface layer of'fi36A1 after oxidation at 1000°C in air for 4h.

n o depletion in aluminium can be detected in the metal subsurface zone. Even an enrichment in aluminium is locally ascertained. In addition internal oxidation occurs leading to the formaticm'hf internal aluminium oxide (see Fig. 12). Neithcr in the oxide scale nor at the metalhxide interface niobium oxides are found. Except from the base metal a signifkant niobium content only exists in the coarse-grained 'IiO? layer where the niobium is dissolved in the'I'iO,.

I7 The Initial Srages in the Oxidation of TiAl

255

Fig. 10.'IEM bright field image of the metal/oxide interface of Ti36AI after oxidation at 1000°C in air for 4h.

H

1w

TiAl

/

TuAl

b03'

NCP

AION

Fig. 11. Schematic illustration of the oxide scale and of the subsurface layer of'I'i36AI after oxidation at 1000°C in air for 4 h.

256

C. Long a d M. Schiitze

Pig. 12. TEM micrograph (a) and schematic illustration (b) o f the oxide scale and metal subsurfacr layer ofTi3SA15Nb alter oxidation at 900°C in air lor 4 h.

17 The IniiialStapes in the Oxidation of 1iAI

257

17.4 Discussion 17.4.1 Oxidation Progress 17.4.1.1 Stage I As can be seen from the weight gain curves the oxidation process can be divided into several stages. This is also shown in the schematic illustration in Fig. 13. Stage I of the oxidation process represents the initial stage with a very high oxidation rate. At the very beginning TiO, and A1,0, form due to the high oxygen partial pressure at the metal surface which is higher than the equilibrium oxygen partial pressure of the oxides [12,21]. During further oxidation aluminium is preferentially oxidized because of the higher thermodynamic stability of A1,0, compared to TiOiTiO, in equilibrium with the y-?'iAl at the metal/oxide interface [21].This leads to the formation of an aluminium depletion layer in the metal subsurface zone [4,22,23]. From investigations of the first minutes of oxidation [12] and from the results presented in this paper it can be concluded that at the metal/oxide interface no a-Al,O, is formed but either a metastable AI,O, or an aluminium oxynitride similar in structure to AI,,O,,N. ?'he AI,,O,,N has a crystal structure which is comparable to a metastable, intermediate stage of the Mg- or Ni-spinel [24,25]. It is conceivable that Al,,O,,N may be stabilized on the one hand by the presence of nitrogen in combination with the low oxygen par-

A

1

j

I1

III

- t

I:

n:

111:

Iv:

Gas

TiN,AlON NCP TiAl

Pig. 13. Schematic illustration of oxidation mechanisms at the mctal/oxidc intcrfacc during stage 11.

258

C. Larig nnd M. Scliiiize

tial pressure at the metal/oxidc interface and on the other hand by the dissolution of titanium. In contradiction to the binary phase diagram of TiAl the depletion layer in the metal subsurface zone after a 0.5 h oxidation at 900°C consists of a phase with a composition between y-TiAl und a2-Ti3AI.The analysis of this phase by electron diffraction reveals that this phase is equal to the recently reported new cubic phase (NCI’) which is similar in compositions toTi,Al,,?’i,AI or TiSAl3O2[ 17,18,19,20]. By the depletion of the metal subsurface zone in aluminium due to rapid oxidation during stage I the activity of titanium a.fiincreases. The oxygen partial pressure at the metalbxide interface is sufficiently high to form aluminium oxide but too low for titanium oxide formation. In some way nitrogen can penetrate the scale and the nitrogen partial pressure allows the observed formation of titanium nitrides at the metal/oxide interface when a certain value of the titanium activity is reached. ‘The high oxidation rate is caused by the fast formation of nitrides (which later become oxidized to TiO,) in addition to the slowly growing discrete aluminium oxide particles without significant barrier effect. ‘The aluminium depleted layer is consumed by the rapid formation of further nitrides as the absence of a depletion layer after 4 h oxidation at 900°C indicates. Dettenwanger et al. [26] who also report of titanium nitride formation at the mctal/oxide interface did not observe a depletion zone after 1 h oxidation, either. Because of the thin and non-protective oxide scale at this point of time a sufficient supply of nitrogen to the metalhxide interface is presumed to promote this process of nitride formation. Figge et al. [22] also found the formation of nitrides after 1.5 h of oxidation by SNMS investigations.

17.4.1.2 Transition from Stage I to Stage I1 Because of the observed microstructures of the metal/ oxide interface and of the metal subsurface zone after 0.5 h and 4 h of oxidation it is most likely that the transition of stage I to stage I1 is characterized by the consumption of the aluminium depletion layer due to nitride formation and the enrichment of Al,O,-particles in an outer layer of the scale which represents a preceeding step in the formation of a temporary A1,0,barrier. The consequence of the consumption of the aluminium depletion layer is a decrease in titanium activity a, at the metal subsurface zone. ‘llerefore further titanium nitride should be suppressed. However, due to the formation of A127039Nat the metal/oxide interface a local depletion of the metal subsurface layer in aluminium occurs again. This again leads to the formation of titanium nitrides. However, the ratedetermining step for nitride formation is the formation of AI,,O,,N, because Al2,0,,N formation is the basic requirement for aluminium depletion and subsequent titanium nitride formation. In contrast to stage I where an extensive aluminium depletion layer already exists when nitride formation starts the nitride formation during stage I1 is triggered by the preceeding A1,,0,9N formation.

17 The Initial Stuges in the Oxidation oJTiAl

259

17.4.1.3 Stage 11 During stage I1 linear oxidation kinetics are observcd. It is assumed that this linear oxidation behaviour is caused by the repeated cycle of aluminium oxide formation, subsequent local depletion of thc metal subsurface zone in Al, consumption of the Tirich metal phase by nitride formation since nitrogen is still present at the interface. 'fie nitrogen which is needed for this process is set free by the oxidation of the titanium nitrides to T i 0 2 at increasing oxygen partial pressure when the oxide scale grows inwards. 'I'he suggested mechanisms which hamper the renewed formation of an aluminium depletion zone and which lead to the formation of the mixed inner oxide scale are schematically illustrated in Fig. 14 [12]. It shows the situation at the mctal/oxide interface at three successive points of time. At the beginning of stage I1 (t = 1,) A12,0,,N (AION) and titanium nitrides are present in the oxide close to the metal/oxide interface within the metal subsurface zone beneath the interface, a Ti-rich

metastable Ah03 /

A1ox

metal / oxide interface

'l'i-richmetal

'lil

(NW

j

1 11

0

TiN

metastable A1201 I AlON

'Ti-rich metal

(NCP)

I

dissolution of metastable A1203 / AION and outward diffusion of A1 oxidation ofTiN to T i 0 2

I

t =t2

interface

EN rii

t-

Ti01

pore

t3

......... .........

POX

AION

%-rich metal

(NCV

TiAl

b 1: ......

Ti&

Ti02

metastable Ah01 I /\ION

Ti-rich metal

(NCP)

mctastable A l 2 a I AlON TiAl

......

Ti-rich metal

(NW

TiAl

Fig. 14. Schcrnatic illustration of thc stagcs ofTiAl oxidation by means of the weight gain curve and the corresponding microstructures of the oxide scale.

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metal phase designated as new cubic phasc (NCP) being in contact with AlON (due to aluminium depletion by AlON formation) and y-TiAl being in contact with TiN (due to aluminium enrichment by TiN formation). As thc oxide scale grows inwards (t = tr) at the sites of high ’I’i activity a,,, titanium nitrides form. whereas aluminium oxide is formed at sites of sufficient aluminium activity aA,(TiAl).’I’he nitrides of the previous interface layer are oxidized to ’Ii02 when they become incorporated into thc oxide scale where a higher oxygen partial pressure exists. Furthermore it is assumed that the AION-phase becomes unstable in the vicinity of high’li activity and of nitride forming conditions and dissolves in the titania of the adjacent oxide scale. ‘I’hc dissolved aluminium diffuses outwards and reprecipitates in the form of a-Al,O, in the outer oxide scale at a higher oxygen partial pressure pO,.‘Ihc PO, dependend dissolution of AI,O, in TiO, is known from investigations of the long-term oxidation behaviour 141 and may explain the observed formation of a titania rich and porous partial layer adjacent to the metal/oxide. The transition of stage I1 to stage 111 (parabolic growth), which was not investigated in this study, occurs when in the outer oxide scale the outward diffusing dissolved aluminium forms a dense AI,O, barrier. The breakaway which takes place after 300-500h of oxidation (stage IV) was found to be caused by the dissolution o f this outer A1,03 barrier 141.

17.4.2 Temperature Dependence The observed differences in scale thickness and microstructure between the oxide scales and subsurface zones at the various oxidation temperatures seem to be mainly attributed to the different diffusion rates at the respective temperatures. Since the oxidation products formed do not show any differences in the temperature range of 800°C to 1000°C it is concluded that no significant effect of the thermodynamic stability on the composition and structure of the oxidation products occurs. From the calculations of Rahmel and Spencer [21] it is known, however, that the activity of A1 and Ti in the system Ti-AI varies depending on the tempcrature. Thus it has to be taken into account that the temperature may have an influence on the expansion of the phasc fields of some importar$ phases in the system ’li-Al-N-0. Nevertheless it is evidently the temperature which-’mainly influences the kinetics because the structure of the mctaYoxide interface, the formation of titanium nitrides, AI,,O,,N and an aluminium depleted metal phase is on principle always very similar. In this way the effect of different temperatures can, to a certain degree, be interpreted as that of a shift in the different stages of the oxidation process. ’bus, at 800°C an earlier state should be expected for the same oxidation time compared to oxidation at 900°C. However, an oxide scale which exhibits the same characteristics is found at 800°C indicating that at the investigated point of time stage I1 of oxidation has already been reached. One feature only found at 800°C is the formation of needle-shaped Ti,O, grains at the metWoxide interface (see Fig. 8). Because of the crack-like appearance one might speculate that by the growth of the oxide scale tensile stresses arisc in the metal subsurface layer which lead to fracture in this area due to embrittlcment of the metal subsurface layer by dissolved oxygen 118, 19, 27, 281.

I7 The Initial S~agesin rhe Oxidation of TiAl

261

Thus a conncction to rcgions of higher oxygen partial pressure is obtained and instead of nitrides titanium oxide is formcd at the mctal/oxide intcrfacc. As a consequence local temporary regions of high oxidation ratc arc formcd within the scale which may lead to nodules observed on thc surface of the oxidc scale. At highcr temperatures such cracks are not found possibly because of an increased plasticity of the aluminium depleted subsurface zonc or because of a rapid healing rate duc to enhanccd diffusion ~71. In contrast to the situation at 800°C the stage of oxidation invcstigatcd at 1000°C corresponds to a latcr s t a g cornpared to 900"C.'fie main difference conccrning the other samples investigated (i. c. at 800°C and 900°C) can be seen in the occurance of a two-phasc depletion zonc consisting of a?-Ti;Al and the new cubic phase. Evidcntly after stage I1 of oxidation an aluminium depletion zone is formed again as it is known from studies of thc long-term oxidation bchaviour, cf. e. g. [4]. It is also known that this zone is a two-phase dcpletion laycr aftcr longer oxidation times [4,20J.Thus, it is concluded that because of the higher temperature this stage is already rcachcd. The progressive formation of AlON at the metal/oxide interface leads to further depletion of the new cubic phasc in aluminium. As a consequcncc the formation of local a2-X3A1 grains at the mctal oxide interface is observed.The titanium activity increases at these sites and in the further oxidation process the formation of Ti oxidcs or nitridcs is favoured again thcrc. In this way the two-phasc subsurface laycr with diffcrent activities of titanium and aluminium leads to an inner oxidc scale mixcd of'TiO, and Al,O,.'Ihe differcnces found in the form that at 1000°C only AlON is found in contact with NCP while at 900°C it is AlON and TiN can be cxplained by thc effect of the higher scale thickness at 1000°C at thc same oxidation timc.Thc partial pressures at the interface are influenced by the thickness and the quality of the growing oxide scale. Assuming that the nitrogen partial pressure at the interface is defined by thc transport proccsses through the scale it is expected that with increasing oxide thickness the nitrogen partial pressure pN, at the interface decreases. By this means the nitrogen partial pressure pN, may drop below the cquilibrium partial prcssurc of nitride formation and thercforc a higher titanium activity may bc necded to form titanium nitrides.This can be an explanation why at a later stage of oxidation A1,0, is thc favoured phasc in contact with the aluminium depleted ncw cubic phase compared to the nitrides in the carlier stages. Quadakkers et al. 129,221 and Meier et al. 161 observed the disappearancc of nitrides aftcr somc timc which formed in the early stagcs.

17.4.3 Effect of Niobium Addition The striking features after oxidation of Ti35A15Nb for 4 h at 900°C are the slight enrichmcnt of aluminium in the mctal subsurface zone instcad of aluminium depletion, the preferred formation of AlON in wide parts of the metal/oxidc interface and the development of a rathcr densc, coarse-grained partial layer consisting of titania in the oxide scale. Several rcasons for the beneficial effect of niobium addition on the oxidation behaviour are discusscd in the litcraturc [ S . 10,111. Bcsidc the influcnce of niobium on thc a.JaA, ratio and expansion of the y-TiAI phasc field thc effect of doping of titania by niobium is often discusscd. By doping of titania with niobium the concentra-

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tion of oxygen vacancies and titanium intcrstitials should be reduced and thus the transport of oxygen inwards and titanium outwards may be slowed down [30]. If a doping mechanism is assumed it is also cxpectcd that in the same scnsc thc solubility of A1,0, in X 0 2 is decreascd, since the aluminium is dissolved as an interstitial [4]. Thus thc dissolution of AlON in titania and thc outward diffusion of aluminium is reduced. This may stabilize the AlON at the metalloxide intcrfacc leading to a dense AlON barrier there and, thus, to slow oxidation kinetics.

17.5 Conclusions From the prcsent investigations it is concluded that the dccisivc processcs of TiAl oxidation take place at the mctal/oxidc interface. First of all no a-A120, is formed at thc mctal/oxide interface, but a metastable AI,O, or aluminium oxynitride similar in structure to Al,,O,,N which dissolves in the surrounding titania whcn during oxide scalc growth the titanium activity incrcases around the AlON particles and titanium nitride formation takes placc. Thc titanium nitridcs occur at the metal/oxide interface when critical values of the pO,/pN, ratio and of the titanium activity duc to aluminium depletion of the metal subsurface zone are rcachcd. Thc aluminium deplcted metal subsurface zone consists of a ncw cubic phase in thc initial stages of oxidation and in addition of a,-Ti,Al aftcr further aluminium depletion of the metal subsurface layer. Thc intcrrnediate formation of titanium nitrides which are oxidizcd to TiO, during further oxidation prevents the formation of a continuous A1,0, layer, resulting in cnhanccd oxidation. 'lhe alternating processes of A1,0, formation, subsequent aluminium depletion and titanium nitride/oxide formation lead to an inncr oxidc scalc mixed of AI,O, and TiO,. In the niobium containing alloy which shows a bettcr oxidation rcsistance the doping of titania with niobium may reduce the dissolution of AION. By this means a thin layer AlON is formed at the interface lcading to a reduced oxidation rate.'Thus it is assumed that the oxidation behaviour of titanium aluminidcs could bc improvcd by stabilizing the aluminium oxidc at the mctal/oxidc interface cither by prevcntion of aluminium depletion of the metal subsurface zone or by reduction of AI,O, dissolution in TiO,.

17.6 Acknowledgements This work has been financially supported by the Dcutsche Forschungsgcmcinschaft (DFG), which is gratefully acknowledgcd.

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17.7 References N.S. C‘houdhury,I I . C Graham, J. W.Hinze: in “Proc. of the Symp. on Properties of IIigh Temperature Alloys”, Iilectrochemical Society, Princeton, N.J. (1977) 668. R.A.I’erkins, K.7: Chiang, G.II. Meier: Script. Met.21 (1987) 150.5. G. Welsch.A.I. Kahveci: in “Oxidation of high-temperature intermctallics”, T. Grobstein, J. Doychak (cds.), the Minerals. Metals and Materials Society (1989) 207. S. Becker,A. Rahmel, M. Schorr, M. Schiitze: Oxidation of Metals 38 (1992) 425. A. Rahmel, WJ. Qiiadakkers. M. Schiitze: Materials and Corrosion 46 (1995) 271. G. I I . Meiw I,.’S. Pettii, S. Hid: J. Physique 1V 3 (1993) 395. N. Zheng, W.J. Qundukker.s,A. G I , I I . Nickel Oxidation o f Metals 44 (1995) 477. K Shidn, H.Anadn: Oxidation of Metals 45 (1996) 197. K. Maki, M. Shioda, M. Sayashi, 7: Shimizu,S. Isobe: Matcr. Sci. and Eng. A153 (1992) 591. II. Nickel, N. Zheng, A. Elschner, WJ. Quadakkers: Mikrochimica Acta 119 (1995) 23. WJ. Qundakkers, A . Llschrier, N. Zheng, H. Schuster, H. Nickel in “Microscopy o f Oxidation 2”; S.B. Ncwcomb, M.J. Bennett (eds.).’Ihe Institute of Materials London (1993) 488. C. I,ang, M. Schiitze:Oxidation o f Metals 46 (1996) 255. Joint Committee on Powder Diffraction Standards, File No. 2633. G. t’etzow, G. Effenberg: “Ternary Alloys vol. 7”,VCH Weinheim (1993) 386. R.W.Beye, I