\input texp.tex  \parskip=3mm
 \def\picture#1#2#3#4#5{{\dimen1=#2mm \divide
 \dimen1 by 1000 \multiply\dimen1 by #4 \dimen2=#3mm
 \divide \dimen2 by 1000
 \multiply\dimen2 by #4 \midinsert{$$\vbox to \dimen2 {\vfil\hbox to
 \dimen1 {\special {picture #1 scaled #4}\hfil}}$$}#5
 \endinsert }}
 \def\b{$\beta$}
 \def\bp{$\beta'$}
 \def\gp{$\gamma'$}
 \def\g{$\gamma$}
 \def\um{$\mu$m}
 \def\degg{$^\circ$C}
 \def\eb{E$_{\rm B}$}
 \def\fb{F$_{\rm B}$}
 
 \centerline{Optimisation of a Multiphase Intermetallic}
 \centerline{Metal--Metal Composite Material}
 \bigskip \singlespace
 \centerline{J. D.  Robson\footnote{$^{\dag}$}{\ninerm Now of the
Manchester Materials Science Centre, Grosvenor Street, Manchester, M1 7HS}
N. Duvauchelle, A. Lugan, J. Street and H. K. D. H.  Bhadeshia}
 \medskip
 
 \parindent=0pt
 
 \centerline{University of Cambridge}
 \centerline{Department of Materials Science and Metallurgy}
 \centerline{Pembroke Street, Cambridge CB2 3QZ, U.K.}
 \bigskip
 \vfill\eject
 
 \doublespace
 \sec{ABSTRACT}
 
Metal--Metal composites (MeMeC's) manufactured by coextruding 
\b[Ni(Al,Ti)], \bp[Ni$_2$AlTi] and \gp[Ni$_3$(Al,Ti) show promise as
high--strength heat--resistant materials. The best combination of strength
and plasticity are obtained by using a high fraction (0.8) of \gp\ as the
matrix phase combined with equal fraction of \b\ and \bp\ reinforcement.
Further improvements in properties may be obtained by increasing the
fracture strength of the matrix phase following the methods used in
monolithic \gp\ alloys. Boron additions are shown to lead to significant
improvements in tensile fracture strength. The highest levels of
plasticity were obtained in a MeMeC material which had a matrix
composition designed to allow the more ductile \g\ phase to precipitate in
the \gp, as well as containing boron to increase grain boundary cohesion.
Fractography suggests that tensile failure occurs in this material in a
much more ductile way, with evidence of local plastic deformation and a
largely transgranular failure mode rather than the intergranular fracture
observed in the other MeMeC materials.
 \vfill\eject
 
 \sec{INTRODUCTION}
 
 \x Single phase  intermetallic materials, while often possessing
desirable properties such as low density and good oxidation resistance,
are often limited in their application either by poor low temperature
ductility or inadequate elevated temperature strength. Many efforts have
therefore focused on the  development of multiphase intermetallic
materials, with the complementary phases used  either to enhance strength
or ductility. Such materials may be produced by solid state precipitation
and directional solidification [1-4]. Another possibility which allows
greater control over the final microstructure than obtained with
solid--state precipitation is to mechanically combine the individual
constituent phases in the required proportions. This approach was
used by Hsuing and Bhadeshia [5] and Robson and Bhadeshia [6] to produce
three phase \b[Ni(Al,Ti)]--\bp[Ni$_2$AlTi]--\gp[Ni$_3$(Al,Ti)] materials
which they referred to as Metal--Metal Composites (MeMeCs). In these
materials the strong but brittle \b\ and \bp\ phases act as reinforcement
in a matrix of the more ductile \gp\ phase. 

In both studies, the \g, \b\ and \bp\ were manufactured
independently in powder form.  The
powders were then coextruded to give the required intimate dispersion of
the three phases in the final material. The  chemical compositions of the \g, \b\
and \bp\ powders had been chosen to correspond to the corners of a
tie--triangle defining the three phase equilibrium at 900\degg. This
ensured the high--temperature stability of the microstructure following
consolidation. It was discovered that this method could be used
successfully  to produce a thermally stable multiphase intermetallic
material which exhibited good compressive strength levels, coupled with
significant plasticity before failure [5]. By refining the size of the
powders and optimising the extrusion conditions further improvements in
compressive failure strain were achieved [6]. Nevertheless, failure in
compression still occurred after relatively low plastic deformation
($\simeq$ 11 \% at room temperature). Studies of failed specimens showed
that the grain boundaries in the \gp\ phase were a source of weakness.
Furthermore, the tensile properties of the MeMeC materials are yet to be
characterised. The aim of the present work was to modify the compositions
of the starting phases in an attempt to improve ductility, as well as
determine the tensile properties of the MeMeC material.

\sec{METHOD}
 
\ssec{Modified Compositions}
 
 Previous work demonstrated that MeMeC material with a high  proportion of
\gp\ phase (80\% \gp , 10\% \b , 10\% \bp ) shows the best compressive
strain to failure without significant loss of strength when compared with
MeMeC variants with higher levels of \b\ and \bp\ [5,6]. In compression,
MeMeC has been shown to fail by the propagation of cracks which initiate in
the brittle \b\ and \bp\ phases or at the interphase boundaries. Although
the \gp\ matrix offers more resistance to crack propagation than the \b\
or \bp, cracks nevertheless occur along grain boundaries in this phase.
These link up and lead to general failure. 

There probably is little scope for increasing greatly the crack resistance
of the \b\ and \bp\ phases [7,8]. However, the ductility may be increased
by enhancing the resistance of the matrix \gp\ to transgranular crack
propagation. In the present study, modifications were made to the matrix
phase composition to achieve this.
 
 It is well established that boron additions to polycrystalline \gp\ phase
can enhance its tensile ductility [9]. Boron segregates to the
grain boundaries where it is thought to enhance the cohesion of
atoms, thereby inhibiting intergranular cracking. It is also well known
that this effect works  only for \gp\ alloys which are on the nickel--rich
side of the stoichiometric Ni$_3$(Al,Ti) composition; i.e.
the \gp\ should contain more than $75\,$at.\% Ni for the boron to be
effective. One aim of this investigation was to determine whether the
boron effect can be exploited successfully for MeMeC, and what
effect this would have on strength and ductility.
 
 The compositions of the powders used in the original work were chosen to
 coincide with the three corners of the \b, \bp, \gp\ tie--triangle on the
isothermal section at 900\degg, of the ternary phase diagram (\fagg) [10].
The position of this tie--triangle does not change much with temperature
and thus there should be little tendency for dissolution or
re--precipitation of any of the three phases up to this temperature.
The three original powder compositions are shown in \tablaa. It can be seen
that the nickel concentration in the original \gp\ powder is below the
critical level required for boron additions to be effective. New \gp\
powder was therefore manufactured with a nickel level of $76\,$at\%, in
excess of the stoichiometric concentration (\tablee). Both boron
doped (0.1 wt\% B) and boron--free variants of this powder were produced.
 
Another possible method for ductilising the matrix phase is by designing
its composition so that it contains a small fraction of the disordered,
more ductile $\gamma$ phase as well as \gp. To investigate this, two more
powders were manufactured with compositions which lie in the two phase
$\gamma$ + \gp\ region on the phase diagram (\figg). One of these powders
was doped with boron to investigate the combination of both effects. 
 
 \ssec{MeMeC Fabrication}
 
  The powders were manufactured following the same methodology as in
previous investigations [6]. The \b\ and \bp\ powders were prepared by
mechanically crushing cast ingots. All the other powders were prepared by
argon gas atomisation. The mean particle sizes of the new \gp\ and \g\ +
\gp\ powders were similar to those of the original \gp\ powder
(diameter$\simeq$ 70\um ). X--ray diffraction studies of each powder were
performed using a diffractometer  with Cu  K$_\alpha$ (doublet  $\lambda$
= 1.5406 and 1.54439 \AA) radiation to determine which phases were
present. The powders were then blended together in the proportions used in
previous work (10\% \b, 10\% \bp, 80\% \gp\ or \g + \gp ) and mixed using
a small mechanical mill. 

Extrusion  was performed at INCO Alloys U.K. using the following
procedure: the blended powder was sealed in an Incoloy 800 alloy
can, preheated to 1170\degg\ for 2~h and then extruded using a 16:1
reduction ratio, a ram speed of 25$\,$mm$\,$s$^{-1}$ and two layers of
C--glass as lubricant. \tablaa\ summarises the nomenclature used to
identify varieties of samples manufactured in this and previous studies.
  
 \ssec{Microscopy and Mechanical Testing}
 
Following extrusion, specimens were polished and etched with Kalling's
solution (2$\,$g~CuCl$_2$, 40$\,$ml HCl,  50$\,$ml ethanol) for optical
examination. Compression test samples were prepared by cutting rectangular
blocks 2.5$\,$mm  $\times$  2.5$\,$mm $\times$ 6.0$\,$mm from the extruded
bar in either a longitudinal or  transverse orientation  using a SiC
slitting wheel. Prior to testing the surface of the samples was lightly 
ground using 1200 grit  SiC paper. Compression testing was performed using
a 50$\,$kN servohydraulic testing machine. Samples  for elevated
temperature testing were allowed thermal expansion under a small constant 
load and were  held for 30$\,$minutes at the testing temperature to ensure
thermal equilibrium had  been achieved within the furnace. Tests were
carried out at temperatures up to 800\degg, with each test repeated at
least three  times. The strain rate for all tests was about
2.3$\times$10$^{-4}\,$s$^{-1}$.

Standard cylindrical threaded Rolls--Royce RLH10186 tensile specimens
were
manufactured by P. S. Marsden, U. K. using cylindrical grinding. The screw
threads at the ends of the specimen were also manufactured by grinding as the
hard and relatively brittle nature of the MeMeC material prohibited the use of
more common methods. The specimen axis was in all cases parallel to the
extrusion direction. Tensile testing was performed at room temperature
using a 100$\,$kN servohydraulic testing machine. Three tests were
performed for each MeMeC variant. Following testing, scanning electron
microscopy (SEM) was used to examine the fracture surfaces of the failed
samples.

 \sec{RESULTS}
 
 \ssec{Powder Compositions}
 
 X--ray  diffraction patterns of the \gp\ and \g + \gp\ powders confirmed
the absence of unexpected phases. The original \gp\ powder, used in
previous work, had contained measurable quantities of \bp\ and the
hexagonal close packed $\eta$ phase (Ni$_3$Ti). These phases were not
detected in the new powders. 

The presence of \g\ in the nominally \g\ + \gp\ powder can be inferred from
the intensity ratio of selected reflections. The face--centred cubic (fcc)
\g\ phase will only contribute to the diffracted intensity when $h,~k$
and $l$ are either all odd or all even. The intensity of the \gp\
superlattice reflections ($h,~k,~l$ mixed odd and even) relative to that of
the fcc lattice reflections will thus be reduced by the presence of the
\g\ phase. The ratios of intensities of the (100) and (110) superlattice
reflections to that of the (200) reflection are shown in \tablaa. The (200)
reflection was taken as the reference as it does not overlap with any \b\
or \bp\ reflections. The intensity ratio of the superlattice reflections
is significantly lower in the nominally two phase \g + \gp powder than the
\gp\ powder, suggesting that \g\ is indeed present.

 \ssec{As Extruded MeMeC}
 
Optical micrographs of MeMeC~E and F are shown in \fagg. A fibrous,
intimate dispersion of the \b, \bp\ and matrix phases is observed as
expected. It is not possible to differentiate between the \g\ and \gp\ in
the matrix of F using optical microscopy. The fraction of the matrix
phases (\g + \gp ) in F and \fb\ is also greater than in E and
\eb. This suggests that some of the \b\ and \bp\ have dissolved at the
expense of \g\ and \gp\ as the fractions of the phases move towards their
equilibrium values during the preheat and extrusion. The boron--doped
materials, \eb\ and \fb, show no significant
differences optically when compared with the boron free materials.
 
 \ssec{Compression Tests}
 
 \x  A typical plot of the room temperature compressive stress {\it vs}
strain for the new materials and MeMeC~D [6] tested in the transverse
direction is shown in \fagg. \eb\ is slightly stronger than
E but shows a similar strain to failure. Both E and \eb\ are
not significantly different from D, either in terms of yield strength
or strain to failure. In contrast, both F and \fb\ show a
significant increase in strain to failure at room temperature with a
maximum failure strain of 33\% recorded for \fb.

The variation with temperature of the compressive yield strength and the
compressive strain to failure in the transverse direction are shown in
\fagg. In all cases, the elastic strain was subtracted to give only the
plastic strain to failure. D and \eb\ show similar behaviour across
the whole temperature range, although \eb\ appears to retain its yield
strength better at higher temperatures. E shows a lower yield
strength but slightly greater strain to failure at temperatures up to
650\degg. Above this temperature the yield strength of E starts to
approach that of \eb, exceeding that of D. In all three materials,
the strain to failure and yield strength remain approximately constant up
to 500\degg\ above which temperature the strength falls and the strain to
failure rises. 

As already noted, F and \fb\ show significantly higher room
temperature strain to failure. This is reduced at 300\degg\ in both
variants, although still exceeds that of the other MeMeC materials. At
500\degg\ the strain to failure increases once more and is comparable with
the room temperature value. Above 500\degg\ compressive failure did not
occur in F and \fb\ \ie specimens could be deformed until almost
completely flat. The yield strength of F and \fb\ are both slightly
less than that of the other variants at room temperature, although still
in excess of 1000~MPa. As the temperature is increased the variation in
the yield strength of F is similar to the other MeMeC materials, in
that it remains approximately constant and then begins to fall
at the highest temperatures. In contrast, \fb\ shows a peak in its
compressive yield strength between 500\degg\ and 700\degg. This peak is
the result of the behaviour of the \gp\ phase which is well know to exhibit
a rise and then fall in strength with increasing temperature. The peak in
strength of \gp\ is expected at $\simeq 700$\degg\ [11]. It is notable that
F, which has a similar fraction of \gp\  does not show this
behaviour. This may be because plastic yielding of the matrix is a more
dominant deformation mechanism in \fb\ than in F, where brittle
failure is more prevalent (see the section on fractography).
 
 \ssec{Tensile Tests}
 
The average failure strengths of the MeMeC materials tested in tension are
shown in \fagg (a).  In all cases, tensile failure occurred by fracture
before yielding. Nevertheless, the fracture strength was greatly increased
by the presence of boron and \g\ in the matrix. For example, the average
fracture strength of \fb\ was 833~MPa compared with a value of
435~MPa for E. It can be seen that, in tension, boron has a more
significant effect on the fracture strength of both E and F
than it has on the strength in compression. This is not surprising since
boron is expected to act by enhancing the grain boundary cohesion which
will have a greater influence on the tensile properties. It was quite
remarkable that the fracture surfaces of failed specimens of \fb\
were always inclined to the tensile axis, presenting a shear mode of
failure as opposed to those of E, \eb\ and F which
were normal to the tensile axis (\fagg (b)). This demonstrates a
more ductile failure mode for \fb compared with the other variants.
Whereas \fb samples always failed within the gauge length
to give valid tensile tests, all other samples exhibited brittle
characteristics with the first fracture event causing a further scission
along the gauge length as a result of the shock of failure. The latter
tensile specimens were therefore always in three pieces following
testing. The results for \fb are therefore extremely encouraging
and demonstrate the alloy design concepts used to enhance ductility.

 \ssec{Fractography}
 
\fagg\ shows two scanning electron microscope (SEM) fractographs. \figg (a)
shows the fracture surface after tensile failure of \eb. Similar
fracture surfaces were also observed for E and F. This type of
fracture surface is characteristic of brittle failure. The surface is
planar, individual grains are easily recognised and the failure has
occurred by intergranular crack propagation. In contrast \figg (b)\ shows
a far more irregular fracture surface, found to be characteristic of
\fb. Examples of both transgranular and intergranular fracture can be
observed. The increased tensile fracture strength is undoubtedly
attributable to this change in fracture mode due to the strengthening of
grain boundaries by boron and because of the presence of the \g--phase in
the microstructure. 
  
\sec{CONCLUSIONS}
 
\x  A variety of metal--metal composite materials based on the
Ni--Al--Ti system have been manufactured, mechanically tested and
examined. The production route allows the initial composition of either
the matrix (\gp, \g + \gp) or reinforcement (\b, \bp) to be modified
independently. In this study, various matrix compositions were chosen to
give (a) a \gp\ matrix containing $>$25$\,$at\% Ni either with or without
boron additions and (b) a \gp\ + \g\ matrix, again either with or without
boron. The following conclusions may be drawn:

{\parindent 5pt
\zz{1.}{ The mechanical properties of MeMeC~E (which has a Ni rich \gp\
matrix) and \eb (with the same matrix, boron doped) are comparable
to those of MeMeC materials produced in previous studies. In contrast,
F and \fb\ which have a matrix sufficiently high in nickel to
contain \g\ show a dramatic increase in compressive strain to failure
(which is approximately doubled at room temperature) compared with the
other MeMeC variants. This increase in strain to failure is not
accompanied by a similar reduction in yield strength -- in both F
and \fb\ a room temperature compressive yield strength in excess of
1000~MPa is retained.}

\zz{2.}{ E and the boron containing \eb\ show similar compressive
mechanical properties. Boron additions also do not make a significant
difference to the compressive behaviour of F, except at
$\simeq$500\degg\ where the boron containing \fb\ shows a peak in
strength which is not exhibited by F. The main role of boron is to
enhance grain boundary cohesion in the matrix and thus suppress
intergranular fracture. The compressive stress acting on the majority of
the grain boundary area during testing will itself reduce the tendency for
this failure mode, and thus any effect of boron is likely to be masked in
compression.}

\zz{3.}{ All MeMeC materials tested in tension fail by fracture before
general plastic yielding. Nevertheless, the tensile fracture strength is
greatly improved by the presence of boron and \g\ within the matrix, with
\fb\ showing the highest fracture strength (833~MPa). Boron
additions improve the tensile fracture strength of both E and
F by $\simeq$150~MPa, the most likely mechanism for this being an
enhancement of the grain boundary cohesion in the matrix.  }

\zz{4.}{ Fractography of specimens after failure in tension reveals that
the failure mode in \fb\ is markedly different from that of all the
other variants. E, \eb\ and F exhibit planar fracture surfaces
which are perpendicular to the tensile axis and failure occurs by
intergranular fracture. In contrast \fb\ shows a much more irregular
fracture surface, inclined to the tensile axis. Failure occurs by a
combination of transgranular and intergranular fracture. This change in
fracture mode may explain the substantial improvement in tensile fracture
strength exhibited by \fb.}

\sec{ACKNOWLEDGMENTS}

\x The authors are grateful to the Defence Evaluation and Research Agency, U.K. and the Civil Aviation Research and Development programme for funding this work. They would also like to thank Mike Henderson, Neil Jones, John O'Driscol and David Knowles for helpful discussions.

\vfill\eject

\sec{REFERENCES}

 {\parindent=10pt \narrower
\def\ref#1#2#3#4#5#6{\item{#1. } #2: {#3} { #4}
{(#5) #6} }

\ref{1}{Misra A. and Gibala R.}{\it Metall. Trans. A}{\bf 28}{1997}{795}

\ref{2}{Whittenberger J. D., Noebe R. D., Johnson D. R. and Oliver B. F.}{\it Intermetallics}{\bf 5}{1997}{173}

\ref{3}{Misra A., Wu Z. L., Kush M. T. and Gibala R.}{\it Phil. Mag. A}{\bf 78}{1998}{533}

\ref{4}{Johnson D. R., Oliver B. F., Noebe R. D. and Whittenberger J. D.}{\it Intermetallics}{\bf 3}{1995}{493}

\ref{5}{Hsuing L. C. and Bhadeshia H. K. D. H.}{\it Metall. Trans. A}{\bf %%@
26A}{1995}{1895}
\ref{6}{Robson J. D. and Bhadeshia H. K. D. H.}{\it Mat. Sci. Tech.}{\bf 16}{2000}{349}

\ref{7}{R. Darolia}{{\it J. Metals}}{\bf 43}{1991}{44}

\ref{8}{Yamaguchi M. Umakoshi Y. and Yamane T.}{{\it Phil. Mag. A}}{{\bf %%@
50}}{1984}{205}

\ref{9}{Lui C. T., White C. L. and Horton J. A.}{\it Acta. Metall.}{\bf 38}{1990}{561}


\ref{10}{Yang R. Saunders N. Leake J. A. and Cahn R. W.}{\it Acta Metall.}{\bf 
40}{1992}{1553}

\ref{11}{Stoloff N. S.}{\it Int. Mat. Rev.}{\bf 34}{1989}{153}

\ref{12}{Yang R.}{PhD Thesis, University of Cambridge}{}{1992}{75}

\vfill\supereject
 

