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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.
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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.
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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.
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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
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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].
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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
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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-
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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.
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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-
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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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