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HIGH-TEMPERATURE STRUCTURAL INTERMETALLICSp

M. YAMAGUCHI{, H. INUI and K. ITO

Department of Materials Science and Engineering, Kyoto University, Kyoto 606-8501, Japan

(Received 1 June 1999; accepted 15 July 1999)

Abstract In the last one and a half decades, a great deal of fundamental and developmental research hasbeen made on high-temperature structural intermetallics aiming at the implementation of these intermetal-lics in aerospace, automotive and land-based applications. These intermetallics include aluminides formedwith either titanium, nickel or iron and silicides formed with transition metals. Of these high-temperatureintermetallics, TiAl-based alloys with great potential in both aerospace and automotive applications havebeen attracting particular attention. Recently TiAl turbocharger wheels have nally started being used forturbochargers for commercial passenger cars of a special type. The current status of the research and devel-opment of these high-temperature intermetallics is summarized and a perspective on what directions futureresearch and development of high-temperature intermetallics should take is provided. # 2000 Acta Metal-lurgica Inc. Published by Elsevier Science Ltd. All rights reserved.

Keywords: Optical microscopy; Transmission electron microscopy; Intermetallic compounds; Mechanicalproperties (plastic); Microstructure

1. INTRODUCTION

Vigorous activity has been present in the research

and development of high-temperature structural

intermetallics for the last one and a half decades.

Some alloys based on Ni3Al and iron aluminides

are currently used for some structural applications

such as furnace xtures [1, 2]. There are several po-

tential applications that have been identied forTiAl-based alloys in the aerospace, automotive and

turbine power generation markets. Aircraft engine

manufacturers are pursuing the implementation of

these alloys in aircraft engines. Recent extensive

engine tests of components of TiAl-based alloys

such as low-pressure turbine blades have revealed

that no serious limitations exist to aircraft engine

applications of TiAl-based alloys [3, 4]. The auto-

motive community is pursuing the qualication and

introduction of exhaust valves and turbocharger

turbine wheels of TiAl-based alloys for automotive

engines. Very recently TiAl turbocharger turbine

wheels have started to be used for commercial cars

of a special type [5]. Thus, these high-temperature

structural aluminides are entering the rst phase of

structural applications.

In parallel to these recent advances in the

research and development for structural appli-

cations, considerable progress has been made in the

basic research of high-temperature intermetallic

compounds. First, it should be pointed out that our

understanding of deformation and creep mechan-

isms and property/microstructure relationships in

TiAl-based alloys is substantially deepened. There

has been a decade of good interaction between the

fundamental research and the industry communities

in the eld of TiAl-based alloys and such inter-

action is believed to have played an important role

for progress in the research and development of

structural applications for TiAl-based alloys.

Recent progress in basic research of high-tempera-

ture intermetallics also includes a new interpretation

of the yield strength anomaly of FeAl on the basis

of interaction between dislocations and thermal

vacancies [6], nding the plastic deformability of

MoSi2 single crystals at temperatures as low as1008C [7] and nding a new refractory metal sili-

cide system based on Mo-rich alloys in the MoSi

B ternary system [810].

These achievements in basic research are hoped to

lead to some new development activities on high-tem-

perature intermetallic compounds. For the hope to

be realized, we should rst summarize the current sta-

tus of the research and development of high-tempera-

Acta mater. 48 (2000) 307322

1359-6454/00/$20.00 # 2000 Acta Metallurgica Inc. Published by Elsevier Science Ltd. All rights reserved.

P I I : S 1 3 5 9 -6 4 5 4 (9 9 )0 0 3 0 1 -8

www.elsevier.com/locate/actamat

p

The Millennium Special Issue A Selection of Major

Topics in Materials Science and Engineering: Current

status and future directions, edited by S. Suresh.

{ To whom all correspondence should be addressed.


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ture intermetallic compounds and then clarify which

questions basic research has answered and which

ones we should seek to answer. The purpose of this

paper is to perform these tasks and then to provide a

perspective on the directions for future research on

high-temperature intermetallics such as titanium alu-

minides, nickel and iron aluminides, and transition

metal silicides. The present paper will, to a large

extent, focus on the results published in the last dec-

ade. In the high-temperature intermetallics commu-

nity, some international conferences and symposia

covering somewhat dierent areas have regularly

been held [11]. The reader may consult proceedings

of these international meetings and recently pub-

lished books on intermetallics [1, 12] to obtain a

broad knowledge of the research and development

that has already been made.

2. TITANIUM ALUMINIDES

2.1. TiAl-based alloys

Of the intermetallic compound phases identied

in TiAl alloys, Ti3Al(a2), TiAl(g ), Al2Ti and Al3Ti

phases are stable at room temperature and their

mechanical properties have been investigated using

single-phase specimens. TiAl-based alloys with two-

phase structures consisting of the major g and

minor a2 phases are the most intensively studied

materials among these aluminides and their alloys.

There are two reasons for this. Firstly, their low

density, strength and modulus retention at high

temperatures, some tensile ductility at room tem-

perature, and reasonably good oxidation resistance

are very attractive as a new class of light-weight

high-temperature materials for structural appli-

cations. Secondly, TiAl-based alloys can be pro-

cessed more or less similarly to metals and alloys

through conventional manufacturing processes such

as ingot melting, casting, forging, precision casting

and machining on almost conventional equipment

[3, 4, 13]. In particular, it is essentially important

that TiAl-based alloys are somewhat ductile even at

room temperature and thus they are readily castable

using standard titanium casting processes [13].

Otherwise, the pace of the research and develop-

ment of TiAl-based alloys for structural appli-

cations would have been slackened quickly.

Engineering TiAl-based alloys generally start soli-dication as the b phase and go through the a

single-phase region and aAa g and a gAa2

g reactions, producing a g a2 two-phase struc-

ture. Thus, some dierent two-phase structures can

be obtained by manipulating these phase transform-

ation reactions. Since mechanical properties of

TiAl-based alloys strongly depend on their micro-

structure, their mechanical properties can be tai-

lored to meet the needs for a specic component by

controlling their microstructure. This is of great

merit for TiAl-based alloys and has aroused a great

deal of fundamental research on microstructure/

property relationships in TiAl-based alloys.

2.1.1. Microstructure. When g/a2 two-phase alloyswith nearly stoichiometric or Ti-rich compositions

Fig. 1. A polycrystalline lamellar structure of a TiAl-basedalloy with a nearly equiatomic composition.

Fig. 2. Crystal structures of the (a) L10 and (b) D019types.

308 YAMAGUCHI et al.: STRUCTURAL INTERMETALLICS


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are prepared by usual melting-and-casting pro-

cesses, a polycrystalline lamellar structure is formed

(Fig. 1). When these two-phase alloys with such a

lamellar structure are heated or hot-worked at tem-

peratures higher than 11508C in the g a region,

the lamellar structure is destroyed and a duplex

structure consisting of equiaxed grains with the g

single-phase and the lamellar structure is formed.

Microstructures of these two types exhibit very

dierent mechanical properties. In general, ne andhomogeneous duplex structures result in good duct-

ility. The lamellar microstructures are poor in duct-

ility; however, they are generally superior to the

duplex structures in other mechanical properties

such as fracture toughness, fatigue resistance and

high-temperature creep strength. Currently cast

TiAl-based alloys are going well ahead of wrought

TiAl-based alloys in the development and im-

plementation status. In addition, lamellar micro-

structures are quite common and persistent after

thermal treatment. Thus, there has recently been

much eort invested in studying lamellar micro-

structures.

The g and a2 lamellae in the lamellar microstruc-

tures are stacked such that a {111}g plane is parallel

to 0001a2 and the closely packed directions on

{111}g are parallel to those on 0001a2 X However,

the "1 10 direction and the other two 10 "1 and

0 "1 1 directions on (111) in the g phase are not

equivalent to each other because of the tetragonal

L10 structure of the g phase [Fig. 2(a)] while direc-

tions of h11 "2 0i on the basal plane in the a phase

(h.c.p.) and a2 phase (hexagonal D019) are all equiv-

alent [Fig. 2(b)]. Thus, when the g phase precipitates

from the a parent phase, the L10 structure can be

formed in six orientation variants corresponding to

the six possible orientations of the "1 10 direction

along a reference h11 "2 0i direction of the a phase

and thus of the a2 phase [14]. When one g plate

impinges on another g plate, one g plate can be

rotated by y, which can be 608 n n 05), and/

or translated by f, which can be 0, 1a2h10 "1 {,

1a6h11 "2 and 1a6h1 "2 1, with respect to the other g

plate (Fig. 3). The three non-zero f vectors corre-

spond to fault vectors for APB, SISF and CSF in

the g phase [15, 16]. When f 0, g/g lamellar

boundaries resulting from such impingement of g

plates are of the true-twin type y 1808), the 1208-

rotational type y 1208, 2408 or the pseudo-twin

type y 608, 3008). Both positions and species of

atoms are mirror images across the true-twinning

plane while only atom positions are mirror images

across the pseudo-twinning plane assuming that the

c/a axial ratio of the g phase is unity. g/g lamellar

boundaries with an APB-type shift were observed in

TiAl alloys [1720]. However, no observations of

lamellar boundaries of the 1208-rotational and

pseudo-twin types with f T 0 have been reported.

Thus, the large majority of g/g lamellar boundaries

are believed to be one of the three types with f 0X

Domains of dierent variant types can coexist

within each g lamella [14, 21, 22]. Boundaries

between such domains are simply termed domainboundaries [14]. Such domain boundaries as well as

g/g lamellar boundaries are all g/g intervariant

boundaries, although domain boundary planes do

not show any preference for a specic crystallo-

graphic plane [22]. Atomic planes parallel to the

lamellar boundaries in domains coexisting in a g

lamella stack in the same sequence, either abcabc F F F

or cbacba F F F [22, 23], and thus two neighboring

variants in a g lamella are always of the 1208-ro-

tational type with or without f T 0X The reason for

this has yet to be claried. However, this is believed

Fig. 3. Translation and rotational operations to create pla-nar faults.

Fig. 4. Schematic illustration of the lamellar structure ofTiAl-based alloys. 1M3M and 1T3T are matrix and twin

variants.

{ {hkl) and huvw ] are often used to dierentiate the rst

two indices from the non-equivalent third one of the tetra-

gonal structure.

YAMAGUCHI et al.: STRUCTURAL INTERMETALLICS 309


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to be closely associated with the growth mechanism

of g lamellae in the parent a phase probably invol-

ving the motion of 1a3h10 "1 0i type Shockley partial

dislocations on alternate basal planes of the parent

phase, similar to the case of h.c.p.-to-f.c.c. struc-

tural change [2426]. Of the three dierent types of

lamellar boundaries, those of the true-twin type

with the lowest energy are most frequently observed

[27]. The growth process of g lamellae must involve

a mechanism to maximize the occurrence of true-

twin type g/g lamellar boundaries [27, 28]. Thus, the

lamellar structure of two-phase TiAl-based alloys

can be schematically described as in Fig. 4 [29].

Energies (for a review see Ref. [30]) and chemistry

[3133] of lamellar boundaries have been studied.

However, it has yet to be understood how they

aect the microstructural variables and deformation

behavior of the lamellar microstructures in TiAl-

based alloys.

2.1.2. Microstructural features and mechanical properties of lamellar microstructures. Mechanical

properties of the lamellar microstructures in TiAl-

based alloys depend on the lamellar orientation

with respect to the loading axis and lamellar micro-

structural variables such as grain size, thickness and

spacing of g and a2 lamellae and g domain size.

However, the lamellar orientation has far more in-

uence than lamellar microstructural variables. A

new approach for studying mechanical properties of

the lamellar microstructure of TiAl-based alloys

was introduced in 1990 by producing crystals where

the entire crystal consists of only a single lamellar

grain [34]. Since numerous thin twin related lamel-

lae are contained in the major constituent g phase,

these crystals were named polysynthetically twinned

(PST) crystals [34] from analogy with the phenom-

enon ``polysynthetic twinning'', which is often

observed in mineral crystals [35]. Since then, lamel-

lar microstructural features and fundamental prop-

erties of the lamellar microstructure such as

microstructural characterization, deformation, frac-

ture toughness and macroscopic ow behavior have

all been extensively studied by making the best use

of the fact that the PST crystal is in a sense a single

crystal of the fully lamellar polycrystalline alloy.

2.1.2.1. PST crystals. What the studies using PST

crystals have typically claried is the eect of lamel-

lar orientation on the mechanical properties and the

anisotropic macroscopic ow behavior of the lamel-lar microstructural form. The tensile properties of

PST crystals depend strongly on f but not signi-

cantly on c (Fig. 5). PST crystals exhibit the high-

est strength at f 908, however, tensile ductility at

f 908 is almost zero. A good balance of strength

and ductility is obtained at f 08, where strength

is not as high as that for f 908, but tensile duct-

ility is as large as 510% at room temperature.

When f is in the range of 308608, yield stress is

much lower and elongation is much higher than

f 08 and 908 [36]. This trend remains unchanged

Fig. 5. Loading axis orientation for polysyntheticallytwinned (PST) crystals. Macroscopically PST crystals pos-sess hexagonal symmetry with respect to the direction per-pendicular to the lamellar boundaries because each g

lamella consists of domains of six orientation variants.Thus, a loading axis orientation given by a set of f and cis equivalent to that given by a set off and 2c an inte-

gral multiple of 608.

Fig. 6. Macroscopic deformation of PST crystals.



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almost up to 10008C [37]. The f dependence of the

yield stress and ductility of PST crystals results

from the fact that shear occurs parallel to the lamel-

lar boundaries (deformation in soft mode) for f 308608 but it occurs mostly on {111} planes inter-

secting the lamellar boundaries (deformation in

hard mode) when f is close to 08 and 908. The

large dierence in yield stress between orientations

for deformation in soft and hard modes can be

mostly interpreted in terms of the HallPetch mech-

anism and the Schmid factors on the operative slip

and/or twinning systems [29, 38, 39]. The mean free

path of the dislocations in the soft mode corre-

sponds to the average size of domains in g lamellae

which is about two orders of magnitude greater

than the average thickness ofg lamellae correspond-

ing to the mean free path of the hard mode dislo-cations.

Ordinary slip on f111gh110, superlattice slip on

f111gh101 and twinning on f111gh11 "2 can be oper-

ative in the g phase in g/a2 two-phase alloys and the

dierence in the critical resolved shear stress

between these systems is not signicant [37, 40]. In

general, a combination of ordinary slip, superlattice

slip and/or twinning systems operates in domains of

each orientation variant. The combination of oper-

ative systems and the amount of shear produced by

each slip or twinning system has been found to be

determined so that deformation incompatibility at

the lamellar and domain boundaries is minimized

[29, 41]. Macroscopic plastic deformation of PST

crystals is generally given as shown in Fig. 6 on this

basis. The results of the gure are in good agree-

ment with experimental observations [40, 41]. When

f 08, where deformation is typically anisotropic,

ey is much smaller than ex and ez. In particular,

when f 08 and c 08 in compression, ey is

exactly zero. Recently such plain strain deformation

was fully conrmed by precise strain gage measure-

ments of the three axial strains [42].

Micromechanical models in which the PST plas-

ticity is implemented have been proposed [4345].

Such models are attractive since they may be used

to predict the plastic response and texture develop-

ment in polycrystalline lamellar structures. If tensile

tests of microspecimens consisting of a single g

domain are carried out using the recently developed

microsample tensile machine [46], the eects of

lamellar and domain boundaries on the PST plas-

ticity would be more clearly understood.The anisotropic PST plasticity suggests that large

deformation incompatibility may arise between the

two neighboring grains in TiAl-based alloys with

polycrystalline lamellar structures and may exert a

strong inuence on their deformation. In order to

avoid such diculty and to design polycrystalline

lamellar alloys with the optimized mechanical prop-

erties, an approach with a potentially great payo

has been proposed. It is to use directional solidica-

tion techniques to produce a columnar grain ma-

terial with the lamellar orientation aligned parallel

to the growth direction (Fig. 7) [4750]. A recent

deformation study on bi-PST crystals consisting of

component PST crystals with f 08 rotated about

the loading axis indicates that deformation incom-

patibility arising at the grain boundary does not sig-

nicantly aect the deformation behavior of these

bi-PST crystals since component crystals deform in

hard mode and thus the ow stress of bi-PST crys-

tals is already high enough to activate additional

deformation in the vicinity of the grain boundary to

compensate for the incompatibility [51]. Thus, the

good combination of tensile properties displayed by

each columnar grain may be directly reected in the

tensile properties of directionally solidied (DS)

ingots such as shown in Fig. 7.

Fracture toughness of PST crystals is also sensi-tive to the relative orientation of the notch and

lamellar boundaries. It is high when the notch

orientation is of the crack-arrester or the crack-divi-

der type and it is very low for the crack-delamina-

tion orientation [5255]. The fatigue [56, 57] and

creep strength [58, 59] of PST crystals depend on f

similarly to their yield strength. Thus, the best com-

bination of the high-temperature strength, room-

temperature ductility and fracture toughness is

expected to be achieved through growing composite

microstructures, such as that of Fig. 7, by direc-

Fig. 7. Schematic presentation of directionally solidiedTiAl ingot.



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tional solidication. Recent creep tests of direction-

ally solidied [60, 61] and strongly textured TiAl-

based alloys [62] provide evidence that this can

indeed be the case.

2.1.2.2. Materials with polycrystalline lamellar

microstructures. Hot extrusion of powder- andingot-metallurgy alloys above the a-transus tem-

perature produces unique rened grain/ultrane

lamellar microstructures with a lamellar thickness

of typically 100200 nm [6365]. These hot-extruded

alloys exhibit an unprecedented high strength of

800 MPa together with 3 5% tensile elongation and

30 MPa m1/2 fracture toughness at room tempera-

ture. Where does such high strength come from?

Lamellar thickness can be further rened to less

than 10 nm by quenching from the a phase eld

and aging between 400 and 8008C, although only

microhardness has been measured for materials

with such a small lamellar width of the order of a

nanometer [66]. The eects of microstructural fea-

tures on the mechanical properties of TiAl-based

alloys with lamellar microstructures are being clari-

ed since materials with lamellar microstructural

parameters controlled over a wide range are becom-

ing available. Rening the grain size of both

wrought-processed and cast alloys can be eectively

made by adding boron [67, 68]. More recently, it is

becoming possible to control the grain size and

lamellar microstructural variables through heat

treatment or hot working above the a transus and

subsequent controlled cooling (Supertransus

Processing) [69, 70]. The a grain size, i.e. the lamel-

lar grain size, is determined during heat treatmentor hot working above the a transus and the lamellar

characteristics are determined by how the specimen

is cooled. Lamellar thickness and spacing can vary

widely. However, they are reasonably approximated

by log-normal distribution regardless of the lamellar

microstructure scale [27, 39, 71].

The uniaxial yield stress (sy) of TiAl-based alloys

with polycrystalline lamellar microstructures is

given as the sum of the intrinsic strength of the g

phase (s0), the lamellar hardening and the grain-

size hardening where the last two terms are given as

a HallPetch type function of lamellar spacing (l )

and grain size (d), respectively [Fig. 8 ] [64, 71]

sy s0 kdad1a2 klal

1a2X 1

Dimiduk et al. [71] have evaluated the last two

HallPetch terms on the basis of dislocation pileup

models, deformation experiments on PST crystals in

soft and hard modes and the experimentally deter-

mined relation, d al2 (where a is a constant), and

have drawn a very informative conclusion that the

key to high strength in fully lamellar TiAl-based

alloys lies in rening the lamellar size rather than

the grain size. Thus, substantial room-temperature

yield strength improvements of b 800 MPa in TiAl-

based alloys with rened grain/ultrane lamellar

structures [64, 65] are attributed to lamellar rene-

ment. The correlation occurring between yield stress

and lamellar spacing at a constant grain size of

about 75 mm yields a kl value in the range of 0.1

0.2 MPa m1/2

[71] which is in agreement with the

HallPetch slope obtained for PST crystals [29].There are shear mists between g lamellae of dier-

ent variants and both shear and biaxial mists

between the g and a2 lamellae and thus coherency

stresses arise within the lamellae and they increase

both in absolute magnitude and relative to yield

stress as the lamellar spacing decreases [7274].

Thus, such coherency stresses should be included in

the model interpreting the lamellar spacing depen-

dence of yield stress of alloys with very ne lamellar

structures.

Both crack initiation and crack propagation

toughness increase with increasing lamellar spacing

in a manner similar to the HallPetch relation since

a small lamellar spacing hinders translamellar

microcracking and thus linkage of the main crack

with interlamellar microcracks [75, 76]. This shear

ligament model predicts that as grain size increases

the size of ligaments formed by mismatched crack

planes increases and thus the crack propagation

toughness increases [68, 75, 76]. However, since

lamellar spacing increases with increasing grain size,

the predicted high propagation toughness is not

often observed and the propagation toughness

reaches a maximum and then gradually decreases

with increasing grain size [68, 75, 76]. In contrast to

the propagation toughness, tensile elongation

increases with decreasing grain size in a wide tem-perature range since tensile fracture is controlled by

the propagation of microcracks with a length com-

parable to the grain size [64, 68, 75, 76].

The long-crack growth threshold (DKth) for fati-

gue was suggested to be proportional to the crack

initiation toughness for a given material [68] and

the 107-cycle fatigue strength is roughly pro-

portional to tensile strength [7779]. Thus the useful

fatigue life region given by the Kitagawa diagram

[80] is expected to be widened by rening both

lamellar spacing and grain size. Ligaments formed

Fig. 8. Grain size (d) and lamellar width (l ) for polycrys-talline lamellar materials.



7/16

in the wake of short or long fatigue cracks are no

longer benecial to increasing the crack propa-

gation toughening since the ligaments are destroyed

by fatigue [81]. Similar crack advance mechanisms

occur at room temperature and at least as high as

6008C [68, 77]. Thus, the high cycle fatigue strength

does not depend on temperature up to 6008C [78,

79].

TiAl-based alloys are candidate materials for gas

turbine applications. Blades in the last stage of

industrial gas turbines will be exposed to a stress of

about 150 MPa at 7008C [82]. Since the creep strain

is required to be smaller than 1% after service for

104 h under these conditions, the creep rate should

be as small as 1010/s [82]. For another application

requiring a shorter period of service, creep strain

for 1000 h is required to be smaller than 0.5% at

200 MPa and 7508C [83], and in addition at least

1.5% room-temperature elongation is required [83].

Aligning the lamellar orientation along the growth

direction using DS techniques is an approach of

great potential to achieve a combination of highcreep strength and high room-temperature tensile

ductility. Directionally solidied (DS) ingots of a

Ti46Al1.5Mo0.2C (at.%) alloy with an aligned

lamellar microstructure in the PST form show a

creep strain as small as 0.15% for 700 h and a

steady state creep rate of 4X2 1010as under a

creep condition of 210 MPa/7508C (the results of

creep tests of these DS ingots were partly published

in Ref. [84] and the rest will be published else-

where). Their room-temperature tensile ductility is

more than 3%. Some polycrystalline alloys with

ne-grained ne lamellar structures exhibit a creep

rate of the order of 1010/s under a creep condition

of 7608C/70 MPa [65, 85] and creep rates of the

order of 109/s under creep conditions such as

7608C/100300 MPa [65, 85] and 7008C/300 MPa

8008C/180 MPa [83]. However, no polycrystalline

lamellar materials comparable with these DS ingots

in a combination of creep strength and room-tem-

perature tensile ductility are currently available.

The results of creep tests made on fully lamellar

polycrystalline specimens with some dierent lamel-

lar spacings [8688] indicate that the minimum

creep rate for specimens with smaller lamellar spa-

cings is signicantly smaller than that for specimens

with larger lamellar spacings, however, the dier-

ence decreases with increasing temperature anddecreasing stress [86, 87]. The larger grain size may

be benecial to creep strength [86]. In addition, the

thermal stability of lamellar microstructural features

exerts signicant inuences on the creep of lamellar

materials. Lamellar renement by mechanical twin-

ning or similar processes that transform the crystal

orientation by a moving interface is a typical

example of microstructural changes during creep

[89, 90]. Dissolution of a2 lamellae also occurs [90

94] and leads to a growing supersaturation of inter-

stitial impurities in the g phase and the eventual

decoration of sub-boundaries and dislocations with

precipitates [82]. An initially rapid creep rate up to

0.10.2% strains which are often observed for fully

lamellar materials is of major concern. This rapid

primary creep is believed to be caused by these

microstructural changes and interface-related defor-

mation mechanisms such as the multiple generation

of dislocation loops from lamellar boundaries [82]

and the motion of dislocations in the lamellar

boundaries [95]. The rapid primary creep strain

induced by interface-related deformation can be

reduced by pre-straining, although this remedy may

not be applied to cast components. The primary

creep strain for DS ingots is much smaller than that

for conventional polycrystalline lamellar materials.

Aligning the lamellar orientation along the loading

axis is very eective in reducing the undesirable pri-

mary creep strain.

In view of what has been known about the

microstructure dependence of mechanical properties

of polycrystalline lamellar TiAl-based alloys, con-

trolling grain size appropriately and making lamel-lar spacing as ne as possible are required to

achieve the best combination of mechanical proper-

ties. Otherwise, aligning the lamellar orientation by

DS processing is required.

2.1.3. Alloy development. To date, various TiAl-

based alloys have been developed [3, 67, 78, 82, 83,

9698]. Adding transition metals with high melting

temperatures is generally benecial to increasing the

high-temperature strength of TiAl-based alloys.

Non-metallic elements such as C, N and Si are also

eective in increasing their strength because of both

solid solution and precipitation hardening (for a

review, see Ref. [82]). Changes in Al content signi-

cantly inuence their strength through the Al-

induced changes in microstructure [3, 99]. The eect

of Al variations is often larger than the eect of

transition elements [3, 100]. Thus, strengthening

wise, there are many possibilities. However, a good

combination of mechanical properties of TiAl-based

alloys can be achieved only when their microstruc-

ture is properly controlled. Aligning the lamellar

orientation by DS processing is an extreme example

of such microstructure control. However, funda-

mental information on solidication pathways and

high-temperature solid-state phase transformations

in important ternary systems is still lacking,although there had been some progress in studies

on phase transformations and their kinetics in the

binary and some ternary systems [69, 70, 101103].

From the point of view of microstructure control,

boron is an important alloying element since adding

boron or a combination of boron and a transition

metal is eective in grain rening and stabilizing

the lamellar structures [65, 68, 96].

In contrast to the mechanical properties, oxi-

dation resistance is independent of microstructure

for a given composition [104]. For binary alloys, a



8/16

reduction in Al content from 50 to 48 at.% results

in increasing the oxidation rate at 9008C by a factor

of four [104]. Of the commonly used alloying el-

ements, Nb is the most important element to pro-vide TiAl-based alloys with good oxidation

resistance. TiAl-based alloys containing Nb, in par-

ticular those containing Nb as high as 10 at.%,

whose microstructure can be still in the fully lamel-

lar form, suppress rutile growth and form a thick

continuous outer Al2O3 scale [98, 105, 106]. Figure

9 shows a turbocharger wheel that has recently

started being used for commercial cars of a special

type [5, 107]. These wheels are made from a TiAl

alloy containing Nb by the counter gravity low

pressure casting process [108] and are said to be im-

plemented in turbochargers without any surface

coating. Since the castability of TiAl-based alloys

with a high Nb content is not generally good, there

must have been considerable progress in casting

technologies for TiAl-based alloys, although pub-

lished information in this area is very limited.

It has been shown that substantial variations in

microstructure and porosity distribution result from

varying cooling rates during casting and such vari-

ations lead to a substantial variability in tensile

ductility for a given composition [99]. It is said that

high ductility is not necessarily required for good

performance of TiAl cast components such as tur-

bocharger wheels. However, ductility is needed for

resistance to damage during manufacturing oper-

ations. Poor ductility leads to high production cost.

If demand for both mechanical properties and oxi-

dation resistance cannot be met by alloying and/or

microstructure control, a cost eective surface coat-

ing method should be developed for TiAl-based

alloys. Recently, a new WO3-uidized bed andWO3-shot blast processes have been developed to

produce a thick continuous Al2O3 layer [109, 110].

2.2. Al2Ti, Al3Ti and L12 variations of Al3Ti

Some independent investigations on the phase

eld and crystal structure of Al2Ti have been made,

but there are still some discrepancies between the

results of these studies (for reviews, see Refs [111,

112]). Recent experimental work [113] shows that

Fig. 9. A TiAl turbocharger wheel (from MitsubishiHeavy Industies Ltd, 1998).

Fig. 10. Composition ranges where protective alumina scales are formed at 800, 1000 and 1200 8C in air[120] and projected on the isothermal section of the AlTiCr system at 8008C; taken from Ref. [119].



9/16

Al2Ti crystallizes into a structure of the tetragonal

HfGa2 type containing 24 atoms per unit cell and it

is a very brittle compound as the complexity of its

crystal structure suggests [113].

Al3Ti has a relatively simple structure of the

tetragonal D022 type, which is derived from the

cubic L12 structure by introducing an antiphase

boundary on every (001) plane. Al3Ti is also very

brittle but otherwise attractive as a high-tempera-

ture structural material because of its low density

(3.4 Mg/m3), relatively high melting temperature b

13508C and good oxidation resistance. Thus, a

great deal of both theoretical and experimental

work on alloying Al3Ti with ternary elements has

been carried out to change the crystal structure to

the higher symmetry L12 structure in the hope that

the increased number of independent slip systems in

the cubic structure improves tensile ductility at low

temperatures (for reviews, see Ref. [114]). However,

no substantial improvement in tensile ductility of

Al3Ti has been achieved. Thus, bulk L12 trialumi-

nides do not seem to be feasible for structural appli-cations.

However, coatings based on Cr-modied L12 tria-

luminides are still of much interest. The Cr- and

Mn-modied L12 trialuminides exhibit a small but

denite plastic strain at room temperature [115,

116]. A tensile strain up to 0.16% was recorded for

a Cr-modied trialuminide with a composition of

Al8Cr25Ti (at.%) in room-temperature bending

tests with strain gages attached to the tension side

of the specimens [116]. The amount of linear ther-

mal expansion of the Cr-modied trialuminide from

room temperature to 8008C is about 1.5% and the

coecient of thermal expansion of the trialuminide

is about 30% higher than that of Ti alloys. Thus, if

the trialuminide is used as an oxidation resistant

coating for Ti alloys, cooling from 8008C to room

temperature gives rise to about 0.35% tensile mis-

match strain in the coating layer. This indicates

that the current level of room-temperature ductility

of L12 trialuminides is about half the tensile duct-

ility needed for using them as coatings for Ti alloys.

It may be possible to ll the gap. In fact, a Cr-

modied L12 trialuminide containing more Cr [Al

14Cr25Ti (at.%)] was reported to exhibit more

than 0.7% plastic strain [117], which was measured

in the same way as that in Ref. [116].

The phase equilibria of the Al-rich portion of theAlTiCr system have been investigated by three

dierent research groups [117119], although there

are still some disagreements between the results of

these studies. Figure 10 shows the protective

alumina-forming composition range in the AlTi

Cr system at 800, 1000 and 12008C [120] projected

on the schematic isothermal section of the system at

8008C taken from Ref. [119]. Interest in developing

engineering coatings based on the trialuminide

phase with compositions forming protective

alumina scales would be reawakened when Ti alloys

are used at higher temperatures and TiAl-based

alloys become widely used as high-temperature

structure components [120, 121]. Also there is a

possibility of developing oxidation-resistant coating

alloys based on the (Al, Cr)2Ti Laves phase, which

has excellent oxidation resistance, paired with more

ductile second phases such as the g phase [122, 123].

More basic work on diusion, plastic deformation

and the phase stability in the AlTiCr system is

required for the optimum development of such

coatings, which may be called TiCrAl for Ti and

TiAl-based alloys like NiCrAl for Ni-based superal-

loys.

3. NICKEL AND IRON ALUMINIDES

3.1. Nickel aluminides

Ni3Al is the intermetallic compound that has

been most intensively studied from both fundamen-

tal and practical points of view. In particular, a

great deal of eort has been made to clarify theow stress anomaly observed at intermediate tem-

peratures. These eorts are well documented and

reviewed [1, 124126]. A variety of Ni3Al-based

alloys have been developed for cast and wrought

applications [1] and some of them are used for

high-temperature furnace applications. Boron-

doped Ni3Al can be extensively cold rolled and

sheet materials that can be cold formed are avail-

able.

The microstructure and texture evolution during

cold rolling has been investigated using both single

and polycrystalline specimens of Ni3Al [127131].

The rolling texture of polycrystalline specimens is

of the copper-type with strongly scattered com-

ponents, which remains very weak up to high

degrees of rolling deformation. After large rolling

reduction, the brass-type texture exhibits the stron-

gest intensity [127130]. The microstructure of

extensively rolled polycrystalline specimens is ex-

tremely inhomogeneous; no dislocation cell struc-

tures are formed. With increasing rolling reduction,

microbands and then shear bands develop. An

increase in the brass-type texture with increasing

rolling deformation was reported from the pro-

nounced microband and shear-band formation

[127]. In comparison to disordered alloys, cross slip

is normally suppressed in Ni3Al because of dislo-cations dissociated into superlattice partial dislo-

cations on {111}. However, local disordering occurs

on heavily activated slip planes [129, 132, 133], dis-

location mobility is greatly improved and massive

cross slip may occur. This promotes shear-band for-

mation. Because of such inhomogeneity of defor-

mation, the rolling texture of polycrystalline Ni3Al

is weaker than those of pure metals and alloys.

Recrystallization in extensively cold-rolled poly-

crystalline specimens starts at shear bands at low

annealing temperatures TH5008C), while the



10/16

matrix needs a much higher annealing temperature

Tb 7508C to be completely recrystallized [128,

130, 131, 134]. Inhomogeneous deformation struc-

ture and low grain boundary mobility result in a

locally very ne but inhomogeneous recrystallized

structure [128]. The recrystallization textures are

very weak. This has yet to be claried.

NiAl is less dense than current Ni-based superal-

loys and it has a high melting temperature, excel-

lent oxidation resistance and high thermal

conductivity. However, the structural use of NiAl

is hindered by low fracture toughness at low tem-

peratures and low strength and low creep resist-

ance at high temperatures. Slip in NiAl occurs

primarily on h001if011g and h001if100g systems

and thus only three independent slip systems are

available. It has been reported that the brittle pro-

blem of NiAl may not be overcome without modi-

fying the slip systems [135]. However, this has not

been possible to date. Developmental eorts have

been made to improve its creep strength by pre-

cipitates and solution alloying. Tensile strengthand stress-rupture properties that compete with

current Ni-based superalloys have been achieved

through precipitation of an ordered L21 Heusler

phase in NiAl single crystals [136]. However, the

precipitation of the Heusler phase makes such NiAl

single crystal alloys more brittle than binary NiAl.

It may still be a long time before structural appli-

cations of NiAl become feasible. The considerable

eorts on the study of the physical and mechanical

properties of NiAl and the development of NiAl-

based alloys have been well documented and

reviewed [137, 138].

3.2. FeAl-based alloys

FeAl-based alloys with 3540 at.% Al have out-

standing oxidation, suldation and corrosion resist-

ance in an aggressive chemical environment because

they form a stable and adherent alumina lm.

However, they show relatively low ductility in air

that is caused by environmental embrittlement [139]

and relatively low high-temperature strength. FeAl-

based alloys, based on an Fe36 at.% Al compo-

sition, have been developed adding alloying el-

ements such as Cr, Nb, Mo, Zr, C and B not only

to optimize a combination of mechanical propertiesbut also to improve weldability (for reviews, see

Refs [140, 141]). Micro-additions of boron (0.01

0.02 at.%) and grain rening increase resistance to

moisture-induced hydrogen embrittlement and add-

ing carbon is eective in suppressing hot-cracking,

which makes FeAl-based alloys unweldable [141].

Plastic ow in FeAl with the B2 structure occurs by

slip along h111i at lower temperatures, while at

high temperatures it occurs by the motion of h100i

dislocations. The simple core structure of the h100i

dislocations operative at high temperatures suggests

that introducing strengthening phases such as car-

bides, nitrides and borides for precipitation harden-

ing is the only way to signicantly increase tensile

and creep strength at temperatures above 6008C

[142]. In fact, the high-temperature strength of

FeAl-based alloys containing Zr and C is enhanced

by ZrC precipitates [141].

Recently, FeAl-based alloy sheets with 40 at.%

Al content have been manufactured through an

innovative combination of roll compaction of FeAl

powders and thermomechanical processing. Roll

compacted green sheets were de-bindered, partially

sintered and then cold rolled to a nal thickness of

0.2 mm through several intermediate annealing at

or above 11008C [143]. Controlled recrystallization

is of critical importance to achieve optimized mech-

anical properties in FeAl-based alloys [144, 145].

Full dense sheets have a ne-grained microstructure

with an average grain size of 20 mm [143]. Tensile

elongation is close to 5% at room temperature and

increases with temperature [143]. These FeAl sheets

have high electrical resistivity. Resistive heating el-ement applications have been developed making the

best use of their high electrical resistivity and good

cold workability. Since fully dense sheets with a

ne-grained microstructure can be subjected to var-

ious metal forming, more applications will be facili-

tated for their excellent oxidation and corrosion

resistance benets.

In addition, there has been signicant progress in

fundamental studies of the mechanical properties of

FeAl. A new vacancy-hardening model has been

introduced for the anomalous yield strength peak

observed at 4006008C [146] (for a review, see Ref.

[147]). The onset temperature and formation

enthalpy of thermal vacancies in intermetallic com-

pounds with b.c.c.-derivative ordered structures are

generally low and thus the concentration of thermal

vacancies in compounds with these structures is

generally high. Since the vacancy concentration CVvaries as exp1aT and the hardening due to

vacancies varies as C1a2V [148], the model predicts at

intermediate temperatures an exponential increase

in strength with increasing temperature. This

increasing strength is terminated by dislocation

creep since more vacancies are produced above the

peak temperature but they are mobile and aid dislo-

cation climb instead of impeding dislocations [146].

This model is consistent with many experimentalobservations, although the orientation and loading-

mode dependencies of yield stress may not be well

interpreted by the model [146, 147]. The Burgers

vector of dislocations playing a major role in plastic

deformation changes from h111i to h100i at a tem-

perature close to the yield stress peak. However,

this Burgers vector change is not directly related to

the yield stress peak [146, 147].

Since the enthalpy of the vacancy formation is

low but the migration enthalpy is high, high

vacancy concentrations are easily retained in FeAl-



11/16

based alloys after annealing at elevated tempera-

tures and cooling slowly [146, 147]. Thus, the

strength at low temperatures strongly depends on

annealing temperature, cooling rate and the amount

of alloying element forming vacancy complexes

[149].

4. TRANSITION METAL SILICIDES

MoSi2, Mo5Si3, Ti5Si3 and C40 silicides formed

with Cr, V, Nb and Ta have been attracting atten-

tion as candidate structural materials to be used inoxidizing environments at temperatures higher than

the upper limit for Ni-based superalloys (for

reviews, see Refs [150153]). The increasing interest

in these transition metal silicides is reected in the

three international conferences [154156]. Of these

silicides, MoSi2 oers a combination of excellent

oxidation resistance, a high melting point (20208C),

relatively low density (6.24 g/cm3), high thermal

conductivity and thermodynamical compatibility

with many ceramic reinforcements [157]. Thus, con-

siderable eort has been devoted to developing

MoSi2-based composites [151, 157]. However, there

are still some key issues for the future development

of these materials. One of the issues is to improvethe low-temperature fracture toughness by a second

phase reinforcement. Fracture toughness values of

10 MPa m1/2, which are said to be typically required

for industrial applications, should be achievable at

competitive cost.

Fig. 11. Isothermal section at 16008C for the Mo-rich side

of the MoSiB system [10].

Fig. 12. Creep strain rates at 13008C plotted as a function of stress for some binary MoSi2 and ternary(Mo, X)Si2 single crystals with [001] (hard) and [0 15 1] (soft) orientations. Creep strains at 1200 8C for

some relevant materials are shown for comparison. (a)(e) refer to Refs [167171], respectively.



12/16

Recently, the oxidation resistance of Mo5Si3,

which is the most refractory compound in the Mo

Si binary system, was found to be improved by

adding less than 2 wt% boron to a level near that

of MoSi2 in the temperature range of 80014508C

[8, 158, 159]. The boron addition responsible for

the formation of the molybdenum borosilicide

improves the oxidation resistance dramatically.

When a small amount of boron is added to Mo5Si3,

a non-porous, protective scale forms [8, 158, 159].

The reasons for this are reported to be two-fold [8,

158, 159]; rstly, boron modies the ow behavior

of the scale and allows viscous sintering to occur to

close pores that form when MoO3 volatilizes; and

secondly, after the initial rapid mass loss due to vol-

atilization of MoO3, a protective borosilicate layer

forms. Also in the NbSiB system, two ternary

borosilicide phases, Nb14Si3B3 and Nb5Si3B2 exist

[160]. However, since Nb2O5 does not volatilize and

thus a thin protective borosilicate layer is not

formed on Nb5Si3B2, the boron addition responsible

for the formation of Nb5Si3B2 does not improve theoxidation resistance of Nb5Si3 to a level near that

of NbSi2 [161].

The MoSiB phase diagram was rst investi-

gated by Nowotny and his colleagues (their 1873 K

isotherm is available in Handbook of Ternary Alloy

Phase Diagrams [162]) and recently the Mo-rich

portion of the system was investigated in more

detail [10]. The results of the recent investigation

show that a ternary borosilicide phase exists with a

composition range around the stoichiometric

Mo5SiB2 value (the T2 phase in Fig. 11). It has

been suggested that there are some possibilities for

developing dierent microstructural morphologies

in (Mo solid solution T2 phase mixture, including

the application of rapid solidication and the pre-

cipitation of Mo within a T2 matrix for strengthen-

ing and toughening [10]. A recent study on the

microstructure and mechanical properties of Mo

Mo3SiMo5SiB2 silicides prepared by arc-melting

followed by drop casting into Cu chill molds indi-

cates that the precipitation of platelets of Mo solid

solution in the T2 phase strengthens the T2 phase

[163]. Alloys containing about 25 vol.% of Mo

solid solution were reported to exhibit some degree

of oxidation resistance [163]. These recent results

warrant further investigation on silicide alloys

based on the T2 phase.Considerable progress has also been made in our

basic understanding on the mechanical properties of

MoSi2 through deformation experiments of MoSi2single crystals in compression [7, 164, 165]. Five

slip systems, f110h111, f011h100, f010h100,

f023h100 and f013h331 have been identied to be

operative in MoSi2 single crystals depending on

crystal orientation and temperature. Critical

resolved shear stresses for these slip systems have

been measured. Importantly, the work of Ref. [7]

has shown that MoSi2 single crystals with orien-

tations other than [001] exhibit macroscopic com-

pressive deformation as low as 1008C. Only

f013h331 has a non-zero Schmid factor for the

exact [001] orientation. However, the critical

resolved shear stress for the slip system strongly

depends on crystal orientation, in particular near

[001]. It drastically increases as the compression

axis approaches to [001]. Thus, [001]-oriented crys-

tals can hardly be deformed at low temperatures.

More recent deformation experiments of MoSi2single crystals at high temperatures show that the

exact [001] orientation shows not only extremely

high strength at strain rates of the order of 104/s

but also excellent creep resistance [166]. Figure 12

shows creep strain rates at 13008C for some binary

MoSi2 and ternary (Mo, X)Si2 single crystals with

the hard [001] and soft [0 15 1] orientations as a

function of stress. For comparison, creep strain

rates at 12008C for some relevant materials are also

shown in the gure [167171]. The exact [001]

orientation is seen to be comparable to or a little

inferior to the Si3N4SiC composite in creep resist-ance. Currently, some attempts to produce compo-

sites with the [001] MoSi2 matrix are in progress

[166].

5. CONCLUDING REMARKS

Reducing the weight of wheel is essential to

increase turbocharger performance. Weight wise,

TiAl is denitely superior to Ni-based superalloys.

In addition, the exhaust gas temperature of auto-

motive engines, in particular diesel engines, is

higher than the upper limit of Ti alloys and is not

high enough to make the best use of ceramics.

Thus, TiAl-based alloys are the best candidate ma-

terials for turbocharger wheels. In fact, TiAl wheels

have recently started being used for turbochargers

to be tted to passenger cars of a special type. In

view of the fact that tting turbochargers to cars

eectively reduces carbon dioxide emissions, turbo-

charger production is expected to increase and pas-

senger cars not only of a special type but also of

ordinary types would be soon equipped with turbo-

chargers. In order to bring mass production to tur-

bocharger rotors with TiAl wheels, a cost eective

production route must be developed. It would be

useful for the purpose of developing TiAl alloys

with better castability and a cost eective joiningprocess of TiAl-based alloys with other materials

such as steels. Fundamentally research wise, more

work is required to clarify the solidication path-

ways, defect formation mechanisms and solid-state

transformations in the relevant alloy systems.

The recent success in developing TiAl turbochar-

ger wheels for automotive engine applications indi-

cates that in order to nd new industrial

applications for a material, it is essential to identify

applications where the best use of its advantages

over its competitors can be made. Identifying resis-



13/16

tive heating applications for FeAl sheets leads to

making good use of the unique properties of FeAl

sheets. This is another example of how it is import-

ant to nd applications which t with the properties

of the material in order to nd users of the material

in industry. More eorts are needed to nd appli-

cations for high-temperature intermetallic com-

pounds driven not only by structural properties but

also other properties such as electrical conductivity

and oxidation and corrosion resistance.

Ti3Al-based alloys and novel types of intermetal-

lics such as Laves phases and A15 compounds have

not been referred to in this paper. The current sta-

tus of the research on these materials is well docu-

mented in relevant review papers in recent

proceedings of international meetings [11] and

recently published books on intermetallics [1, 12].

Acknowledgements This work was supported by JSPS-RFTF 96R12301 grant and Grant-in-Aids for ScienticResearch from MESSC (No.10355026) and MESSC (No.1045026).

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