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Acta Materialia 53 (2005) 617–628
www.actamat-journals.com
Dynamic densification of metal matrix-coated fibrecomposites: modelling and processing
H.X. Peng a,1, F.P.E. Dunne b,*, P.S. Grant a, B. Cantor a,2
a Oxford Centre for Advanced Materials and Composites (OCAMAC), Department of Materials, Oxford University, Parks Road, OX1 3PH, UKb Department of Engineering Science, Oxford University, Parks Road, Oxford OX1 3PJ, UK
Received 4 March 2004; received in revised form 17 August 2004; accepted 11 October 2004
Available online 11 November 2004
Abstract
The consolidation processing of Ti–6Al–4V matrix-coated fibre (MCF) composite under vacuum hot pressing (VHP) has been
investigated. A new test methodology has been developed for the determination of in situ matrix coating creep properties. In using
the methodology, only a single, simple test is required, together with finite element modelling of the single fibre compression test.
The creep coefficient and stress index have been determined for electron beam evaporated physical vapour deposited Ti–6Al–4V at
900 �C to be 1.23 · 10�5 and 1.3, respectively. Consolidation experiments have been carried out on multi-ply MCF arrays under
vacuum hot pressing. Finite element models have been developed for the dynamic consolidation of both square and hexagonal fibre
packings. The creep constants for the Ti–6Al–4V, determined using the single fibre test, were assigned to the coating in the finite
element models. Excellent agreement between predicted and experimental results was achieved, providing verification of the single
fibre test methodology for the determination of creep constants.
� 2004 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Keywords: Single fibre tests; Matrix-coated fibres; Vacuum hot pressing; Dynamic densification; Titanium matrix composites
1. Introduction
Titanium matrix composites (TMCs) in which the
reinforcement is continuous silicon carbide fibre are
being developed for application in aircraft engine com-
ponents such as compressor bladed rings (blings). TMCs
offer superior properties compared to conventionalmaterials and may allow significant reduction in compo-
nent weight. Among the techniques developed for man-
ufacturing TMCs [1,2], the matrix-coated fibre (MCF)
1359-6454/$30.00 � 2004 Acta Materialia Inc. Published by Elsevier Ltd. A
doi:10.1016/j.actamat.2004.10.015
* Corresponding author. Tel.: +44 1865 273 140; fax: +44 1865 273
905.
E-mail address: [email protected] (F.P.E. Dunne).1 Present address: Department of Aerospace Engineering, Univer-
sity of Bristol, University Walk, Bristol BS8 1TR, UK.2 Present address: Vice-Chancellor�s Office, University of York,
Heslington, York YO10 5DD, UK.
method offers a number of advantages over others, such
as uniform fibre distribution, reduced fibre damage and
better fibre fraction control during consolidation
processing [3–5]. In the MCF method, the fibres are
packed into arrays and consolidated either by hot iso-
static pressing (HIP) or by vacuum hot pressing (VHP)
to provide the finished TMC. Long fibre-reinforced me-tal matrix composites are unlikely to be used in their
own right for component manufacture. Instead they
are likely to be used as local reinforcement in a larger
metal structure. In this context, matrix-coated fibres
can be used to produce locally reinforced TMC compo-
nents and the VHP technique is likely to be more prac-
tical than HIP because it allows the consolidation of the
MCFs directly into a die of desired component material.Thus, the consolidation of MCFs via the VHP technique
is of much interest [6].
ll rights reserved.
618 H.X. Peng et al. / Acta Materialia 53 (2005) 617–628
The consolidation process is important in determin-
ing the properties of the final composite. Pressure and
temperature cycles have to be optimized to ensure the
quality of the final product [1]. In order to optimize
these parameters, theoretical simulations that can pro-
vide an insight into the fabrication process [7–15] andsome experimental studies [6,10] have been carried out.
However, there are two problems with existing work.
First, the material properties used in the models are of-
ten taken from bulk matrix materials such as Ti–6Al–
4V. The properties of bulk material are not the same
as those of the Ti–6Al–4V coating produced during
MCF manufacture by physical vapour deposition
(PVD). Although PVD bulk materials have been pro-duced and investigated to provide improved property
data [16], the properties of the unique microstructure
of the Ti matrix formed during PVD onto fibres are
not well established [17,18]. Second, while many of the
theoretical simulations were carried out by assuming a
square fibre array and only a few studies have consid-
ered the hexagonal array [9,10], the experimental data
were mostly acquired from randomly packed fibres.Akisanya et al. [7] employed steel cylinders with diame-
ters of 12 and 8.2 mm coated with plasticine with thick-
ness of 1.5 and 3.4 mm as an alternative in their
experimental studies to simulate a square array of
MCFs. This approach was adopted because of the tech-
nical complexity and difficulty in achieving a regular fi-
bre arrangement in MCFs. Due to the absence of
regular fibre packing, both matrix deformation and fibrerearrangement occur during the consolidation process.
Consequently, the fibre arrangement has not been well
controlled to provide valid experimental results for ver-
ifying theoretical predictions.
Existing modelling work has provided valuable in-
sight into the consolidation process, and provides, for
example, stress–strain data during densification, the size
and shape of the yield surface [1,8], and details of thematerial flow [10]. Carmai and Dunne [19,20] have
developed physically based constitutive equations for
the consolidation of MCFs which require just two mate-
rial properties; namely, the creep coefficient and expo-
nent for the titanium matrix coating. The model,
implemented within the FE method, is able to predict
the consolidation behaviour given the process parame-
ters of pressure and temperature, and details of the par-ticular composite selected (fibre size and volume
fraction) [19]. A further simplified model has also been
developed for simulating practical consolidation proc-
esses [21], which also relies on knowledge of the matrix
creep coefficient and exponent.
In the present work, a novel test method is developed
to determine the MCF coating creep properties. This is
achieved by carrying out compressive tests on a singlematrix-coated fibre, under transverse loading. The
extraction of the matrix creep properties relies on finite
element modelling of the single-fibre compression test.
The single-fibre test methodology presented provides a
comparatively simple way to obtain matrix coating
properties by testing a single-coated fibre. Constrained
uniaxial consolidation tests are also carried out in which
square and hexagonal fibre packings are carefully main-tained, using test methods already developed [22]. These
tests are also modelled using FE techniques, using the
creep properties obtained from the single-fibre tests.
The FE predictions are compared with the results of
the experiments, so that the effectiveness of the single fi-
bre test in determining correctly the creep properties can
be established.
The dynamic densification behaviour of MCFs dur-ing uniaxial vacuum hot pressing is then investigated
in which dies made of Ti–6Al–4V alloy have been used
to generate conditions closer to those experienced in
practical consolidation processing. We investigate both
fibre rearrangement and microstructural evolution in
processing, and present results of scientific and practical
interest.
2. Single fibre testing
2.1. Experimental arrangement
Ti–6Al–4V matrix-coated SiC monofilament fibres
with a diameter of �0.245 mm and SiC volume fraction
of �35% were supplied by Rolls-Royce plc, Derby, UK.The single fibre compression tests were conducted using
a vacuum hot press. The experimental arrangement is
shown schematically in Fig. 1. A number of fibres were
positioned parallel to each other between two precisely
machined ceramic discs with diameter of 50 mm and sur-
face roughness of �0.2 lm. The fibres were well spaced
relative to one another to ensure that they did not touch
during the compression testing. The length of each fibrewas no less than 100 times that of its diameter, so that
plane-strain conditions applied. Two identical steel
plates were heated by symmetric induction coils. To as-
sure the highest possible uniformity of compression of
the fibres, a ceramic ball was placed under the hydraulic
ram. The arrangement was centralized along the axis of
the ram. To measure deformation of a few tens of mi-
crons, a specifically designed and manufactured hightemperature extensometer together with its signal condi-
tioner (Model 3448 RCIH-010-015, Epsilon Technology
Co., Jackson, USA) and a data logging unit with 16 bit
resolution were employed. This enabled a deformation
of 0.25 lm to be repeatedly resolved within a 10 mm
gauge length. The extensometer was enclosed in a
water-cooled box with two ceramic rods attached to
the upper and lower ceramic discs where the tempera-ture was monitored using K-type thermocouples. After
pulling the vacuum, the temperature was ramped at a
Fig. 1. Experimental set-up for the single-fibre compression testing.
Upper Movable boundary
Fibre
Matrix
Ux=Uy=0
Lower unit-cell boundary
Lef
tu
nit
-cel
lbo
un
dar
y
Fig. 2. Finite element model developed for simulating single-fibre
compression.
H.X. Peng et al. / Acta Materialia 53 (2005) 617–628 619
heating rate of �10 �C /min to 900 �C and held for
15 min to allow thermal equilibrium to be attained
and the readings from the extensometer to become sta-ble. A constant load was then applied and the nominal
pressure was calculated according to the length and
the number of fibres. The data logger recorded the ver-
tical displacement versus time. After testing, the de-
formed fibres were cross-sectioned, mounted and
polished. Both polished and etched samples were exam-
ined using optical microscopy with DIC/polarized light.
2.2. Modelling the single fibre compression
In the finite element model developed, it was assumed
that the SiC fibre was stiff compared to the titanium al-
loy matrix, especially at elevated temperatures. The elas-
tic moduli of SiC and Ti–6Al–4V at 900 �C were
assumed to be 354 and 28.4 GPa, respectively [10].
The SiC fibre was therefore assumed to be rigid. Thematrix coating was considered to be incompressible
and to deform by power-law creep. Uniaxial power-
law creep is given by
_e ¼ Arn; ð1Þwhere _e is creep rate, r the stress, A the creep coefficient
and n the stress exponent. This can be generalized intomultiaxial form as
_eij ¼3
2Arn�1
e rij; ð2Þ
re denotes the effective stress defined by
r2e ¼
3
2r0ijr
0ij; ð3Þ
where r0ij ¼ rij � dijrm is the deviatoric stress, rm the
mean stress and dij is the Kronecker-Delta symbol.
The constitutive power-law creep equation (2) was
implemented into ABAQUS finite element software
using a UMAT subroutine.
The finite element model developed for the single fi-bre compression using ABAQUS is shown schematically
in Fig. 2. Due to symmetry, one quarter of the fibre was
modelled. Since the fibre was assumed to be rigid, the fi-
nite element mesh was only developed for the matrix
coating which consisted of 2D, four-noded plane strain
elements. All nodes lying on the fibre-matrix boundary
were fixed in both the x and y directions and those on
the lower unit-cell boundary were allowed to move inthe x-direction only while all nodes lying on the left
unit-cell boundary were allowed to move in the y-direc-
tion only. An upper movable boundary was used for the
application of the distributed load simulating the pres-
sure state during the compression process and was al-
lowed to move in the direction of the y-axis. Sticking
friction was assumed between the matrix and the mova-
620 H.X. Peng et al. / Acta Materialia 53 (2005) 617–628
ble boundary. Examination of tested specimens and
comparison with finite element modelling results showed
this to be a good assumption at 900 �C. The pressure
applied on the movable boundary was in the range of
20–40 MPa. The displacement of the upper movable
boundary simulated the extensometer-measured dis-placement in the experiments. By choosing values for
the creep constants A and n, finite element simulations
of the test could be carried out. By carrying out para-
metric finite element studies over ranges of values for
A and n, and by comparing the resulting predicted dis-
placements with experimental measurements, the opti-
mum values of A and n to give the best fit to the
experimental data was obtained. For a particular nomi-nal pressure of 30 MPa, the calculated and experimen-
tally determined displacements are shown in Fig. 3(a)
for the optimum creep constants. In addition to the
extensometer-determined displacements, the displace-
ment was also measured by using an image analyser
Fig. 3. (a) Best fit finite element displacement for the single MCF
compression test at 900 �C and 30 MPa and (b) micrograph of the
deformed MCF with superimposed finite element mesh.
and micrographs obtained under dark mode in which
the coating edges were highlighted. Micrographs were
obtained by carrying out repeat tests to different levels
of displacement. These results are also shown in Fig.
3(a) by the symbols. The predicted and experimentally
determined single matrix-coated fibre deformations atthe end of the test are shown in Fig. 3(b), and good
agreement was obtained.
It is argued that the values of A and n so obtained are
the true creep constants for the coating and that they are
unique. This is because the creep deformation occuring
in the coating is highly inhomogeneous, with heteroge-
neity in strain, stress, and importantly, strain rate. In
some ways, this can be considered similar to carryingout many homogeneous, uniaxial tests on the coating
over a wide range of stress level. This is the conventional
way to identify creep constants. The optimum values of
the creep constants A and n for 900 �C were found to be
1.23 · 10�5 and 1.3, respectively. In the subsequent sec-
tions, we examine in detail whether these are correct. If
it can be demonstrated that they are, then it is argued
that we have developed a simple testing methodologyfor the determination of the creep properties of the fibre
matrix coating.
3. Investigating the validity of the test methodology
In order to test the correctness, or otherwise, of the
creep properties obtained, experimental tests were car-ried out on both square- and hexagonal-coated fibre
packing arrays under conditions of constrained uniaxial
consolidation. Both cases were also modelled, and the
matrix coating assigned the creep constants determined
above. The correct prediction of the consolidation
behaviour in the two cases provided evidence that the
creep properties are correct.
The fibre square and hexagonal arrays were ob-tained by firstly winding fibre onto an aluminium
drum to form a ply with 24 fibres. A small amount
of acrylic binder was used to hold the fibres together.
The binder was subsequently burned off at �320 �C.The fibre-ply were cut into the desired length (no less
than 100 times that of the fibre diameter) and posi-
tioned into a hot press die made of nickel-alloy as
shown in Fig. 4. The advantage of this die assemblywas that it allowed a fine control of the fibre arrange-
ment by varying the width of ram �D� according to the
diameter and number of MCFs in the fibre-ply, since
the fibre sideways motion was well constrained at both
sides. More details can be found elsewhere [22]. The
extensometer was attached on the upper and lower
ram �D�. Due to the limitation on the gauge length
and full-scale travel of the extensometer, the maximumnumber of fibre layers used in these experiments was
Fig. 5. Finite element model developed for square fibre array.
Fig. 6. (a) Micrograph showing a square fibre array early in the
consolidation process; (b) the deformed finite element mesh superim-
posed on the micrograph showing excellent agreement; and (c) the
resulting fully dense rectangular fibre array obtained under contrained
uniaxial compression.
Fig. 4. Die assembly for MCF consolidation experiments.
H.X. Peng et al. / Acta Materialia 53 (2005) 617–628 621
10. To make a square array, the number of fibres in
each layer was the same; 24 in this case. For a hexag-
onal array, the number of fibres was alternated layer
by layer with a difference of one fibre in each layer,
that is, with fibre numbers of 24/23/24/23� � �24/23.The fibre arrangement had a significant effect on the
measured dynamic densification behaviour, in particu-
lar, the displacement versus time curve. In addition,the two arrangements led to quite different initial rela-
tive densities; hexagonal packing provided the higher
initial relative density.
The vacuum chamber was purged with Ar followed
by evacuation. The die assembly was then heated with
a ramping rate of �10 �C/min. After the temperature
reached 900 �C and the extensometer reading became
stable, a constant load was applied for a specified time.At the end of the test, the induction power was cut off
and the die assembly was subjected to natural cooling
in vacuum. Cooling to 300 �C took �2 h. A data logger
recorded displacement, load and temperature. The
cross-sectioned sample was polished and examined using
optical microscopy with DIC/polarized light to reveal
the microstructure of the composite, with particular
attention to fibre arrangement.
3.1. Consolidation of MCFs with square array
configuration
The finite element model developed to simulate uni-
axial constrained consolidation of multiple fibres ar-
ranged with square packing is shown schematically in
0
100
200
300
400
0 500 1000 1500 2000
Time / s
Dis
pla
cem
ent
/m
0
400
600
800
1000
Lo
ad /
kN, T
emp
erat
ure
/ oC
FEM-disp.
Temp.
E XP-disp.
Load 5
Fig. 7. Predicted and experimental displacement versus time curves
obtained for the consolidation of a square fibre array under
constrained uniaxial compression at 900 �C and 30 MPa.
622 H.X. Peng et al. / Acta Materialia 53 (2005) 617–628
Fig. 5. For reasons of symmetry, only one quarter of the
coated fibre was modelled. The right non-movable
boundary was also a line of symmetry to produce a
repeating cell in which no fibre displacement in the hor-
izontal direction was allowed, simulating the con-
strained uniaxial compression. Pressure was applied to
the upper movable boundary, as in the experiments.
Sticking friction was assumed between the matrix andthe upper movable boundary because of the rapid effects
of diffusion bonding when the titanium alloy coating of
one fibre comes into contact with another. The friction
condition between the matrix and the right non-movable
boundary had very little effect on the predicted results
and hence no friction was assumed. The creep constants
assigned to the fibre coating were those determined
above from the single fibre test.Fig. 6(a) shows a micrograph of the square packed fi-
bre array early in the consolidation process. Note the
good control of fibre packing. Fig. 6(b) shows a closing
void in more detail together with the excellent agreement
with the superimposed finite element model prediction
Fig. 8. Finite element model develop
at the corresponding time in the consolidation process.
On completion of consolidation, the voids were com-
pletely removed as seen in Fig. 6(c). Due to the con-
strained uniaxial compression, the centre to centre
distance between fibres in the vertical direction was less
than that in the horizontal direction, with the latterremaining almost unchanged. This results in the near-
perfect rectangular fibre array at the end of the process,
shown in Fig. 6(c). The experimental and predicted
overall vertical displacements versus time are shown in
Fig. 7, and again there was excellent agreement, which
was reproduced for similar experiments carried out at
20 and 40 MPa. However, the sample consolidated at
20 MPa for 50 min was not fully dense at the end ofthe test indicating that this pressure was too low and,
consistent with Schuler�s work [10], suggested a thresh-
old value of �30 MPa for densification in reasonable
times.
3.2. Consolidation of MCFs with hexagonal array
configuration
The finite element model developed for hexagonal fi-
bre packing is shown schematically in Fig. 8. The repre-
sentative unit cell did not apply for the fibres near to the
vacuum hot press die. This would be likely to lead to dif-
ferences in predicted and experimental results if a small
number of fibres were used in the experiments. However,
because each layer in our experiments contained 23 or
24 fibres, this was unlikely to be relevant. Sticking fric-tion was assumed between the contacting fibre coatings.
The upper movable boundary was a rigid surface that
was allowed to move only in the vertical direction. A
force was applied to the rigid surface to generate the re-
quired applied pressure. The creep constants assigned to
the fibre coating were, again, those determined above
ed for a hexagonal fibre array.
Fig. 9. Micrograph showing a consolidated hexagonal fibre array with
finite element model superimposed, showing good agreement.
0
30
60
90
120
150
180
0 200 400 600 800 1000 1200 1400Time / s
Dis
pla
cem
ent
/ µm
0
200
40
600
800
1000
Lo
ad /
kN, T
emp
erat
ure
/ oC
Tem
Exp.-
FEM-
Load 5
Fig. 10. Predicted and experimental displacement versus time curves
for a hexagonal fibre array under constrained uniaxial compression at
900 �C and 30 MPa.
H.X. Peng et al. / Acta Materialia 53 (2005) 617–628 623
from the single fibre test. The results of a constrained
uniaxial compression test carried out at 900 �C and at
30 MPa on a hexagonally packed fibre array are shown
in Fig. 9, at the end of the experiment, during which thehexagonal packing was carefully maintained. Superim-
posed on the micrograph is the finite element model pre-
diction for the corresponding time, showing good
agreement. The predicted and experimental overall ver-
tical displacement versus time curves are shown in
Fig. 10, and again, good agreement between prediction
and measurement is achieved. Comparing Figs. 10 and
7 shows that the time required to achieve fully densematerial was shorter for the hexagonal packing (�1200
s compared to �1500 s for the square array), and of
course that the displacement required to achieve this
was smaller (approximately half of that required for
the square array).
4. Effect of fibre rearrangement and microstructural
evolution in consolidation processing
4.1. Fibre rearrangement
It has been demonstrated that the fibre arrangementcan be maintained in either the square or hexagonal
packing formats which allows the dynamic densification
behaviour of MCF arrays to be extracted. This is mainly
attributed to the well-constrained uniaxial compression,
i.e. sideways motion of the fibres was prohibited by fine
control of the die assembly. The negligible deformation
of the nickel alloy dies in the temperature and pressure
ranges used in the present study also helped to retainthe fibre array.
In a practical consolidation process, however, fibre
rearrangement may inevitably occur due partially to
the die deformation at high temperature and hence less
constraint imposed on the fibre sideways motion [22].
In addition, the MCFs diffusion bond themselves to the
die in which they are being consolidated. To demonstrate
this, the consolidation of the MCFs into a titanium alloy(Ti–6Al–4V) die has been carried out. The die assembly
was heated up with a ramping rate of about 10 �C/min
to 900 �C. After compression of the MCF arrays for a gi-
ven time, the die assembly was allowed to cool naturally
for 2 h to 300 �C in vacuum. In order to obtain insight
into the deformation process, several experiments were
interrupted at different times and therefore extents of
consolidation. These samples were sectioned, mountedand polished. Precautions were taken to ensure the sur-
faces perpendicular to the fibre axes were polished to
the highest standards. Both polished and etched samples
were examined using optical microscopy with DIC/po-
larized light. Because the tight fit of the fibre-die assem-
bly at room temperature, the initial fibre arrangement in
the die was likely to be constrained to square rather than
hexagonal array packing, and this was confirmed byexamining the MCF array at the initial stage of the con-
solidation process as shown in Fig. 11(a). As consolida-
tion proceeded, it was reasonable to assume that
rearrangement and deformation co-existed as the fibre
array tended to the closest packed state – a hexagonal ar-
ray. The cross-section of a near-fully consolidated TMC
is shown in Fig. 11(b) comprising a fibre array with much
strengthened hexagonal character. Complete diffusionbonding between the matrix coating and the titanium al-
loy die is shown in Fig. 11(c).
The dynamic densification behaviour for a practical
fibre arrangement of mixed square and hexagonal
characters is shown in Fig. 12 together with those
for carefully controlled square and hexagonal fibre ar-
rays under identical processing conditions of 900 �Cand 30 MPa. The ram displacement was normalizedby the measured total displacement for each particular
fibre array. To densify a square array required a larger
Fig. 11. The cross-section of the fibre array: (a) at the initial stage of the consolidation process showing a near square fibre array; (b) the near-fully
consolidated composite for a near hexagonal fibre array; and (c) sound bonding with upper die at the initial stage.
0
0.2
0.4
0.6
0.8
1
0 400 800 1200 1600
Time / s
No
mo
rlis
ed D
isp
lace
men
t
Fig. 12. The normalized displacement versus time curves for a
practical process (heavy line), square fibre array (square symbol) and
hexagonal fibre array (diamond symbol) under identical processing
conditions.
624 H.X. Peng et al. / Acta Materialia 53 (2005) 617–628
displacement (typically, twice) than for a hexagonal ar-
ray because of the different initial packing density.
From Fig. 12, for the practical case where fibre rear-
rangement occurs, dynamic densification behaviour
falls between those for hexagonal and square array
packings. The vertical displacement (Dh) caused by
the fibre rearrangement during the transition from asquare to hexagonal array in uniaxial constrained
compression is [22]:
Dh ¼ ðN � 1Þð1� coshÞd; ð4Þ
where N is the number of fibre-layers, d is the fibre diam-
eter and h is the angle representing a particular stage of
the rearrangement, where 0 6 h 6 30�. The fibres form asquare array with h = 0 and a hexagonal array at
h = 30�. Fibre rearrangement is expected to occur
throughout the consolidation process and the value of
h at a particular stage is hard to quantify.
H.X. Peng et al. / Acta Materialia 53 (2005) 617–628 625
4.2. Microstructural variation during consolidation of
regular fibre arrays
Fig. 13 shows the microstructure of a fully consoli-
dated composite with �square� fibre array and with com-
plete diffusion bonding. Because of the largerdeformation along the vertical direction, a near equi-
axed matrix structure was formed in the region between
each pair of vertical fibres in Fig. 13(a), while a colum-
nar structure was retained between each pair of horizon-
tal fibres in Fig. 13(b). After experiencing the thermal
history described above, the matrix structure consisted
of a coarse a-plate and b-phase columnar structure. As
discussed later, the break up of the lamellae in the regionbetween each pair of vertical fibres was likely to be due
to the larger deformation compared with that between
each pair of horizontal fibres. The microstructure of a
hexagonal array deformed at 900 �C and 30 MPa for
only �5 min is given in Fig. 14(a). The regular fibre ar-
ray was near perfectly retained during consolidation.
Microstructural coarsening probably occurred during
the slow cooling process from the consolidation temper-
Fig. 13. The matrix microstructure with a square fibre array showing
(a) near-equiaxed structure between each pair of vertical fibres and (b)
columnar structure between each pair of horizontal fibres (the pressing
direction is vertical).
Fig. 14. Micrographs of the hexagonal fibre array: (a) at the start of
consolidation showing coarsening of the lamellar structure; (b) just-
consolidated composite showing incomplete diffusion bonding; and (c)
further-consolidated material showing improved diffusion bonding.
ature. Fig. 14(a) shows that the contact area betweeneach pair of �vertical� fibres was larger than that between
the horizontal fibres. The microstructure of a �just� fullyconsolidated sample (as indicated by the horizontal re-
gion of the displacement versus time curve on the data
logger screen) is shown in Fig. 14(b). All voids were
eliminated, and the relative density of the composite
was unity. Again, the fibre array was well maintained.
626 H.X. Peng et al. / Acta Materialia 53 (2005) 617–628
A particular feature of the microstructure shown in Fig.
14(b) was that the boundaries between adjacent fibre
coatings were still visible. This indicates that diffusion
bonding was not yet complete. Additional time at pres-
sure and temperature was necessary to ensure full bond-
ing. The microstructure of a sample compressed for anextra 8 min is shown in Fig. 14(c) where better bonding
was achieved. These findings may be of significant
importance for developing the manufacturing process
for a real component. Because of the reduced overall
matrix deformation required to consolidate a hexagonal
array, the break up of lamellae was less pronounced
than that observed in a square array as shown in Fig.
13(a). Nonetheless, some fine equiaxed a-phase was pre-sent at the triple points where the material free surfaces
impinged during the final stage of consolidation. Slight
a-phase refinement was also noted along the centre line
where the matrix free surfaces impinged.
Complete diffusion bonding in the square array oc-
curred at a comparatively early stage, in Fig. 13, because
of the high local contact stress between each pair of ver-
tical fibres. For a hexagonal fibre geometry, the localstress at the contacting area is expected to be lower than
that for square fibre arrays under equivalent external
load because of the larger number of vertical load-carry-
ing contacts per MCF.
4.3. Mechanisms responsible for the changes of the matrix
microstructure
The microstructural variation during the heat treat-
ment and hot working of Ti–6Al–4V with a lamellar
transformed microstructure has been a subject of a num-
ber of investigations [23–29]. The b transus temperature
above which the two phase a + b mixture transforms to
b is generally assumed to lie within the range of 995 �C[23] to 1015 �C [27], considerably above the consolida-
tion temperature of 900 �C used in the present study.Therefore, deformation of the Ti–6Al–4V matrix in the
MCFs takes place in the a + b two phase field. Research
on the thermomechanical processing of Ti–6Al–4V [28]
concluded that there was no dynamic (or metadynamic)
recrystallization within this temperature range. It has
been established reasonably well that the break up of
the a lamellae structure is caused by the globularization
or spheroidization, either static [29] or dynamic [30–32],of the colony structure. In view of the MCF consolida-
tion process, it is possible that both dynamic and static
globularization occurs during the pressing and cooling
processes. According to Weiss et al. [33], there are two
possible mechanisms for the breakup of the a lamellae
into low-aspect-ratio segments during hot deformation.
One is that both low and high angle boundaries across
the a-plates are formed, with misorientation angles var-ying from a few degrees up to about 30�. The b-phasecan then penetrate into the a-plates along these sub-
boundaries. If the a lamellae width is small, i.e. less than
two times the penetration distance of the b cusps, the
separation of a lamellae occurs readily. The penetration
rate depends on the diffusion of the alloying elements.
Another possible mechanism is that localized shear
and rotation of the a lamellae can occur during hotdeformation, and the misorientation across the shear
band can reach 20�. The b-phase may also penetrate
readily along the shear bands and partially or fully sep-
arate the a lamella into short segments. Other studies,
such as those by Ari-Gur and Semiatin [34] also con-
cluded that, at sub-transus temperatures, the globulari-
zation of a-plates did not occur by recrystallization
but by localized deformation of the plates and penetra-tion by the b-phase. Only moderate deformation was re-
quired to effect globularization whereas heavy
deformation normally yielded a large fraction of globu-
larization [33,34].
4.4. Dynamic densification behaviour of MCFs
Both experimental results and FE predictions indi-cate that a longer consolidation time is required for den-
sifying square fibre arrays than that for a hexagonal
array under the same conditions of temperature and
pressure. The typical dynamic densification curves (rela-
tive density against time) for consolidating both fibre ar-
rays under uniaxial constrained compression are very
similar to the curve shown in Fig. 12. It is suggested
that, for a practical process where fibre rearrangementmay occur, the densification curve falls between these
two curves, for example as indicated by the solid line,
in Fig. 12.
A high densification rate was observed at the begin-
ning of the vacuum hot pressing consolidation. Previous
investigations [10,35,36] showed an �S�-shaped densifica-
tion curve for the consolidation of a square fibre array by
hot isostatic pressing (HIP). These studies suggested thatdensification occurred in three stages as proposed by
Ward-Close and Loader [18]: an initial stage of low den-
sification rate as the consolidation conditions develop in
the HIP cycle; a second stage of rapid densification dri-
ven by high local contact stresses; and a final low densi-
fication rate as the final voids are consolidated. The low
densification rate at the initial stage was attributed to a
HIP encapsulating canister shielding effect, and simi-larly, a HIP canister used in consolidation also played
an important role in the densification of TMCs reported
by Vancheeswaran et al. [9]. In the present work, in the
absence of a canister, the first stage does not exist during
vacuum hot pressing and a rapid densification rate oc-
curred as a direct consequence of the high local contact
stresses at the initial consolidation stage. As consolida-
tion proceeded, for square fibre arrays, cusp-shapedvoids were formed and retained as shown in Fig. 11. This
is an important feature for consolidating MCFs because
H.X. Peng et al. / Acta Materialia 53 (2005) 617–628 627
micromechanical analyses of the creep collapse of cusp-
shaped pores indicated that their collapse rate could be
as much as an order of magnitude faster than that
of their spherical pore counterpart [37]. A recent study
of the effect of pore shape on densification concludes that
the densification rate of the material with circular poreswas much less than that of an identical material with
cusp-shaped pores [9]. Consequently, it is suggested that
the densification process can be made more time efficient
by retaining the initial cusp-shaped nature of pores.
5. Conclusions
The creep constants determined from the single fibre
test were used in the finite element predictions of the
densification of carefully controlled square and hexago-
nal fibre packing arrangements that showed quite differ-
ent densification time and displacement behaviour. In
general, excellent agreement between predicted and
experimentally determined results was achieved, and it
is concluded that the creep constants determined forthe fibre coating are therefore correct. We have therefore
been able to validate the single fibre test methodology
for the determination of fibre coating creep properties.
This offers advantages over alternative approaches in
that the test is simple to do, and that only one test is nec-
essary. In addition, the method enables the physical va-
pour deposited (PVD) coating itself to be tested, while
on the SiC fibre core, rather than an artificially gener-ated �bulk� Ti–6Al–4V material which typically has a dif-
ferent microstructure and physical behaviour to that of
Ti–6Al–4V on the SiC fibre.
The consolidation behaviour of matrix-coated fibre
composite under constrained uniaxial compression (vac-
uum hot pressing) has been investigated experimentally.
In contrast to most existing studies, square and hexago-
nal fibre arrangements have been achieved and main-tained during experiments by careful control of the die
assembly.
This control allowed the true dynamic densification
behaviour to be investigated because the effects of fibre
packing rearrangement were eliminated.
Finite element models based on a simple power law
constitutive equation describing the matrix deformation
have been developed and verified for simulating the con-solidation processing of matrix-coated fibres. Excellent
agreement between predictions and experiments was
achieved for both square and hexagonal fibre packings.
The Ti–6Al–4V coating deposited onto continuous
SiC monofilament fibre by PVD exhibited a basket-
weave microstructure containing very fine lamellar
a + b colonies. The a colonies coarsened rapidly during
the heating and cooling processes required for the con-solidation of MCF arrays into dense titanium matrix
composites. The break up of a lamellae and formation
of equiaxed �grains� occurred in regions of the MCF
coating where the largest deformation occurred because
of the static and dynamic globularization induced by the
localized deformation of a lamellae and the penetration
of b-phase.After the MCF reached full density, additional ther-
mal exposure under pressure was required to promote
complete diffusion bonding between adjacent MCFs,
particularly in the case of hexagonal fibre arrays.
Carefully controlled square-only and hexagonal-only
MCF arrays provided the lower and upper bounds,
respectively, for the dynamic densification behaviour
of MCFs under uniaxial, hot pressing. In the practical
case where MCF rearrangement occurred, the resultingdensification curve falls between these two bounds.
Acknowledgements
The authors are grateful to Rolls-Royce and the UK
Engineering and Physical Sciences Research Council for
support.
References
[1] Guo ZX, Derby B. Prog Mater Sci 1995;39:411.
[2] Subramanian PR, Krishnamurthy S, Keller ST, Mendiratta MG.
Mater Sci Eng A 1998;244:1.
[3] Wadley HNG, Davison TS, Kunze JM. Comp Part B
1997;28:233.
[4] Baik KH, Grant PS. Metall Mater Trans A 2001;32(12):3133.
[5] Ward-Close CM, Chandrasekaran L, Robertson JG, Godfrey AP,
Murgatroyde DP. Mater Sci Eng A 1999;263:314.
[6] Kunze JM, Wadley HNG. Mater Sci Eng A 1998;244:138.
[7] Akisanya AR, Zhang Y, Chandler HW, Henderson RJ. Acta
Mater 2001;49:221.
[8] Akisanya AR, Zhang Y, Chandler HW, Henderson RJ. Eur J
Mech A 2001;20:77.
[9] Vancheeswaran R, Kunze JM, Elzey DM, Wadley HNG. Met
Trans A 2000;31:1271.
[10] Schuler S, Derby B, Wood M, Ward-Close C. Acta Mater
2000;48:1247.
[11] Vancheeswaran R, Meryer DG, Wadley HNG. Acta Mater
1997;45:4001.
[12] Vancheeswaran R, Elzey DM, Kunze JM, Wadley HNG. Mater
Sci Eng A 1998;244:49.
[13] Vancheeswaran R, Wadley HNG. Mater Sci Eng A 1998;244:58.
[14] Olevsky E, Dudek HJ, Kaysser WA. Acta Mater 1996;44(2):703.
[15] Olevsky E, Dudek HJ, Kaysser WA. Acta Mater 1996;44(2):715.
[16] Warren J, Hsiung LM, Wadley HNG. Acta Mater 1995;43:2773.
[17] Ward-Close CM, Partridge PG. J Mater Sci 1990;25:4315.
[18] Ward-Close CM, Loader C. In: Fores FH, Storer J, editors.
Recent advances in Ti–metal matrix composites. (PA,
USA): TMS Warrendale; 1995. p. 19.
[19] Carmai J, Dunne FPE. Int J Plasticity 2003;19(3):345.
[20] Carmai J, Dunne FPE. Mat Sci Tech 2004;20:478.
[21] Carmai J, Dunne FPE. Mater Sci Tech 2003;19(7):919.
[22] Peng HX, Dunne FPE, Baik KH, Grant PS. Mater Sci Eng A
2003;346:246.
[23] Semiatin SL, Bieler TR. Acta Mater 2001;49:3565.
[24] Semiatin SL, Bieler TR. Metall Mater Trans A 2001;32:1787.
628 H.X. Peng et al. / Acta Materialia 53 (2005) 617–628
[25] Seshacharyulu T, Dutta B. Scr Mater 2002;46:673.
[26] Seshacharyulu T, Medeiros SC, Frazier WG, Prasad YVRK.
Mater Sci Eng A 2002;325:112–25.
[27] Roush ED, Kobryn PA, Semiatin SL. Scr Mater 2001;45:717.
[28] Ding R, Guo ZX, Wilson A. Mater Sci Eng A 2002;327:
233–245.
[29] Stefansson N, Semiatin SL, Eylon D. Metall Mater Trans A
2002;33:3527.
[30] Semiatin SL, Seetharaman, Weiss I. Mater Sci Eng A
1999;263:257.
[31] Shell EB, Semiatin SL. Metall Mater Trans A 1999;30:3219.
[32] Semiatin SL, Kobryn PA, Roush ED, Furrer DU, Howson TE,
Boyer RR, et al. Metall Mater Trans A 2001;32:1081.
[33] Weiss I, Froes FH, Eylon D, Welsch GE. Metall Trans A
1986;17:1935.
[34] Ari-Gur P, Semiatin SL. Mater Sci Eng A 1998;257:118.
[35] Akisanya AR, Zhang Y, Chandler HW, Henderson RJ. Acta
Mater 2001;49:221.
[36] Akisanya AR, Zhang Y, Chandler HW, Henderson RJ. Eur J
Mech A 2001;20:77.
[37] Liu YM, Wadley HNG, Duva JM. Acta Metall
1994;42(7):2247.