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Transcript of Impacto Baja Velocidad
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S1359-835X( N)OOW4-7
Composites Purr A 27A (1996) 112331131
Copyright 0 1996 Elsevier Science Limited
Printed in Great Britain. All rights reserved
1359-835X/96/ 15.00
ELSEVIER
Review
Review of low-velocity impact properties
of composite materials
M. 0. W. Richardson” and M. J. Wisheart
Institute of Polymer Technology and Materials Engineering, Loughborough University of
Technology, Loughborough, Leicestershire LEll 3TU, UK
Received 28 October 1994; revised 19 April 1996)
This paper is a review of low-velocity impact responses of composite materials. First the term ‘low-velocity
impact’ is defined and major impact-induced dama ge mo des are described from onset of dama ge throu gh to
final failure. Then, the effects of the composite ’s constituents on impact properties are discussed and post-
impact perform ance is assess ed in terms of residual stre ngth. Copyright 0 1996 Elsevier Science Limited
(Keywords:
ow-velocity impact; damage modes; constituent properties; specimen geometry; post-impact residual strength;
review)
INTRODUCTION
This review ha s been carried out as part of a research
program me investigating the application of pultruded
composite systemsle3
to the construction of freight con-
tainers. This is a realistic goal, ow ing to opportunities
which have arisen in the container industry&I6 and tech-
nological advanc es in the putrusion process’7-‘9. Con-
tainers have traditionally been metal constructions, and
extensive rese arch h as been performed on the impa ct
response of meta ls over a wide range of velocities. Impac t
damage in metals is easily detected as damage starts at
the impacted surface; however, damage in composites
often begin s on the non-imp acted surface or in the form
of an internal delam ination.
Impact damage is generally not considered to be a
threat in metal struc tures becau se, ow ing to the ductile
nature of the material, large amo unts of energy may be
absorbed . At yield stress the mate rial may flow for very
large strains (up to 20% ) at constant yield before work
hardening. In contrast, composites can fail in a wide
variety of mode s and contain barely visible impa ct
damage (BVID) which nevertheless severely reduces the
structural integrity of the comp onent. Mo st comp osites
are brittle and so can only absorb energy in elastic
deformation and through dam age mechanisms, a nd not
via plastic deformation. The term damage res is tance
refers to the amount of impact damage which is induced
*
To
whom correspondence should be addressed
in a composite system. Clearly, the vast majority of
impacts on a composite plate will be in
the
transverse
direction
but due to the lack of through-thick ness rein-
forcement, transverse dam age resistance is particularly
poor. Interlamina r stresses-she ar and tension-are
often the stresses that cause first failure due to the
correspondingly low interlaminar strengths. A s a result,
design failure strains of 0.5% are used to guard against
impac t failure, resulting in a failure to take advanta ge of
the excellent in-plane strength and stiffness properties of
composites.
DEFINITION OF LOW-VELOCITY IMPACT
Generally, impa cts are categorized into either low or
high velocity (and sometim es hyper velocity), but there is
not a clear transition between c ategories and authors
disagree on their definition.
Sjijblom et aL2’ and Shivakumar and co-workers2’
define low-velocity impa ct as events whic h can be treated
as quasi-static , the upper limit of which can vary from
one to tens of ms-’ depen ding on the target stiffness,
material properties and the impacto r’s ma ss and stiff-
ness. High-velocity impac t response is dom inated by
stress wave propagation through the material, in which
the structure does not have time to respond, leading to
very localized damag e. Boun dary condition effects can
be ignored because the impac t event is over before the
stress waves have reached the edge of the structure. In
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Re v i e w o f l o w v e l o c i t y im p a c t p r o p e r t i e s : M . 0 . W . R i c h a r d s o n a n d M . J . W i s h e a r t
low-velocity impa ct, the dynam ic structural response
of the target is of utmost importance as the contact
duration is long enough for the entire structure to
respond to the impact and in consequence more energy is
absorbed elastically.
Cantwell and Morton2’
conveniently classified low
velocity as up to 10 m s-i, by considering the test tech-
niques which are generally employed in simulating the
impact event (instrumented falling weight impact testing
(IFWIT), Charpy, Izod, etc.) whilst, in contrast, Abrate2 3
in his review of impact on laminated composites stated
that low-velocity impa cts occur for impac t spe eds of less
than lOOms_‘.
Liu and M alvern24 and Joshi an d S un25 suggest that
the type of impac t ca n be classed ac cording to the
damage incurred, especially if damage is the prime con-
cern. High velocity is thus characterized by penetration-
induced fibre breakag e, and low velocity by delam ination
and matrix cracking.
Davies and Robinson26*27
define a low-velocity impac t
as being one in which the through-thickness stress wave
plays no significant part in the stress distribution and
suggest a simple model to give the transition to high
velocity. A cylindrical zone under the impac tor is con-
sidered to undergo a uniform strain as the stress wave
propagates through the plate, giving the compressive
strain as26
E, =
impact velocity
speed of sound in the material
For failure strains between 0.5 and l%, this gives the
transition to stress wave dom inated events at lo-
20 m s-* for epoxy comp osites.
MODE S OF FAILURE IN LOW-VELOCITY
IMPACT
The heterogeneous and anisotropic nature of fibre-
reinforced plastic (FRP) lamin ates gives rise to four major
mode s of failure (although many others could b e cited):
1) ma t r i x mode -cracking occurs parallel to the fibres
due to tension, comp ression or shear;
2)
de l am i n a t i o n
mode-produ ced by interlamina r stresses;
3)
j i b r e
mode -in-tension fibre breakage and in-
compression fibre buckling; and
Ii _11
mmx
J--k
._
(a) uansvcrse view
impacted O/90/0 composite plate
4) pene t r a t i on - t he impactor completely perforates the
impacted surface.
It is very important to identify the mode of failure
because this will yield information not only about the
impact event, but also regarding the structure’s residual
strength. Interaction between failure mode s is also very
important in understanding damage mode initiation an d
propagation24.
M a t r i x damage
The majority of the impa ct test work reported in the
literature has involved low-energy testing, in the range of
1 to 5 J approximately (i.e. that which ca uses only mini-
mal damage). It is this work that has revealed informa-
tion about m atrix cracking and delamination initiation.
Ma trix dam age is the first type of failure induced by
transverse low-velocity impac t, and usually takes the
form of matrix cracking but also debonding between
fibre and matrix. Matrix cracks occur due to property
mismatching between the fibre and matrix, an d are usu-
ally oriented in planes p arallel to the fibre direction in
unidirectional layers. Joshi and Sun2’ reported a typical
crack and delamination pattern show n in
Figure I .
The matrix cracks in the upper layers Figure la ) and
the middle layer Figure lb ) start under the edges of the
impactor. These shear cracks29 are formed by the very
high transverse shear stress through the material, and are
inclined at approximately 45”. The transverse shear
stresses are related to the contact force and contact area.
The crack on the bottom layer of
F igu re la
is termed a
bend ing crack because it is induced by high tensile bend-
ing stresses and is characteristically vertical. The bending
stress is closely related to the flexural deforma tion of the
laminate.30 Lee and Sun31 reached the same conclusions
in their analyses. Cantwell and Morton32 em phasized
that the type of matrix cracking which oc curs is depend-
ent on the global structure of the impac ted specim ens.
For long thin specimens bending cracks in the lower
layers occur due to excessive transverse deflection a nd
subsequent m embrane effects predominate, whereas short
thick spe cimen s are stiffer and so higher p eak contact
forces induce transverse shear cracks under the impactor
in the upper plies.
Liu and M alvern24 presented a detailed view of matrix
cracking which agreed with the above, w hilst Wu and
delamination
(b) longitudinal view
Figure Initial damage
in an
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Review of low-velocity impact properties: M. 0. W. Richardson and M. J. Wisheart
Springer3” reported detailed loca tions of matrix crack ing
for graphite/epoxy plates of various sta cking sequen ces.
Chang and co-workers29,34-38 have performed much
research in this area, and postulated that the bending
crack in the 90” layer is caused by a comb ination of (Tag,
gIl and g13 (Figure 2) for line-loading impact damage.
Their analysis also concluded that 033 was very small
relative to gll and ~7~~hroughout the impact event, and
that there is a critical energy below which no dama ge
occurs.
Delamination
A delamination is a crack which runs in the resin-rich
area (approximately 0.0007 mm in graphite/epoxy lami-
nates39 ) between plies of different fibre orientation and
not between lamina in the same ply group33,40>41.
Liu and Malvern24
compiled detailed connections
between delaminations and the areas over which matrix
cracks w ere found, for various lay-ups. Liu4* explained
that delamination was a result of the bending stiffness
mism atch between adjacen t layers, i.e. the different fibre
orientations between the layers. In his experime ntal work
he found that delamina tion areas were generally oblong-
shaped with their major axis being coincident with the
fibre orientation of the layer below the interface. For
O/90 laminates the shape became that of a distinct
peanu t. These results have been widely reported else-
where in the literature28,33,34,4346.He also stated that it is
the bending-induced stresses which are the major cause
of delamination, as both experiment and analysis
revealed that along the fibre direction the plate tends to
bend conca ve, whilst the bend is convex in the transverse
direction. Liu defined a bending mism atch coefficient
between the two adjacent laminates which includes
bending stiffness terms and predicts the peanut shape
reported for O/90 lamin ates. The greater the mism atch
(O/90 is the worst-case fibre orientation), the greater the
delam ination area will be. This is also affected by
material properties, stacking sequence and laminate
thickness4’.
Dorey48-50 has worked widely in this field and pro-
vides a simple expression for the elastic strain energy, E,
absorbed at the point of delam ination failure, which
suggests that this damage mode is more likely to occur
for short spans and thick laminates with low interlami-
nar shear strength.
21-*WL3
Energy = 9Et
f
where t = thickness, 7 = interlaminar shear strength
(ILSS), u’ = w idth, L = unsupported length and Ef =
flexural modulus.
Delamination initiation an d interaction with matri.x
cracking.
Delamination caused by transverse impact
only occurs after a threshold energy has been reached
and it has been observed that delamination only occurs
in the presence of a matrix crack35. Much detailed
work h as been performed to verify this and explain the
stress states which could cause this interaction.
Takeda et aL51 revealed for the first time the asso-
ciation between matrix cracking and delamination, and
showed that delaminations do not always run precisely in
the interface region, but can run slightly either side . Joshi
and Sun’* studied the delamination-matrix crack inter-
action for O/90/0 lamin ates subject to transverse point
impact. They concluded that when the inclined shear
crack in the upper layer
Figure la)
reaches the interface
it is halted (by the change in orientation of the fibres) and
so propagates between the layers as a delamination. This
delam ination is generally constrained by the middle
transverse crack
Figure lb).
The vertical bending crack
Figure lb)
is thought to initiate the lower interface
delamination, the growth of which is not constrained.
Matrix cracks which lead to delamination are known as
critical matrix cracks34.
Chang and colleagues29.34’35.52 erformed a series of
line-loading, low-velocity im pact tests and reported a
typical damage pattern for a O/90/0 composite as shown
in Figure 3. Change and Choi34 simulated these matrix
cracks in their three-dime nsional finite elemen t analysis
to study the stress in the vicinity of the cracks. They
concluded that delamination was initiated as a mode I
fracture process due to very high out-of-plane norma l
stresses caused by the presence of the matrix cracks and
Figure 2 Diagram of the stress components contributing to a bending
matrix crack in a transverse layer
matr i xcrack
\
clamped boundary
Figure 3
Typical matrix crack and delamination pattern from line-
load impact on a O/90/0 composite
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Revi ew of ow -vel oci ty impact propert i es: M. 0. W . Ri chardson and M . J. W ishearl
high interlamin ar shear stresses along the interface. In his
review on delamination, Gargs3 proposed that matrix
crack initiated delamination was due to the development
of the interlaminar normal and shear stresses at the
interfaces.
Liu and co-workers52 created an analytical model to
study the interaction of damage mechanisms due to line-
load impact, utilizing a fracture mechanics approach.
They showed that both bending cracks and shear cracks
could initiate delam ination, but that delamination
induced by shear cracks is unstable and that bending
crack induced de laminations grow in a stable manner
and proportional to the applied load.
Finn and Springer47’54
described in detail the stresses
which they believed cause impact-induced delamination.
All the modes w hich could be induced by impact-
bending, twisting and transverse shear-were con-
sidered, as were the restraints on the affected ply due
to layers above an d below. They concluded that if the
cracked ply group is above the interface, then (if the
upper interface of the ply group is unrestrained) cl2 and
23 contribute to delam ination; if the cracked ply group
is below the interface, ~7~~ nd g1 2 contribute to delami-
nation as long as the ply group lower interface is
unrestrained.
Mo st fracture m echanics analyses of the initiation and
growth of delam ination are difficult to apply becau se
they assum e a n initial flaw or crack size5’; however
Davies and Robinson27,
in a highly simplified isotropic
axisymm etric analysis for the threshold force for growth
of an internal circular d elamin ation in the mid-plan e,
show (surprisingly) that mode II strain energy release
rate is independ ent of delam ination radius. Therefore an
initial flaw size is not required and the threshold force is
given by:
p2 = gr2Eh3Guc
c
9(1 - z/2)
where
PC
threshold load, GIIc = critical strain energy
release rate, Y = Po isson’s ratio, h = plate thickness and
E = modulus.
The predictions from this equation for delam ination
initiation agreed well with their experime ntal data on
quasi-isotropic laminates.
Delami nation growt h.
Choi and Chang45 reported
that delamination growth was governed by interlaminar
longitudina l shear stress (c13) and transverse in-plane
stress (022) in the layer below the delam inated interface
and by the interlaminar transverse shear stress (~7~4 n
the layer above the interface.
Several investigators have introduced artificial delam i-
nations by including a thin foil between plies in the
manufacturing stage to assess delamination growth
from a known initial size56. Doxsee et a1.57 alculated
the energy abso rbed per unit area of delam ination
growth and found that this was constant (595 Jmp2).
Jih and Sun3’ concluded that the interlaminar fracture
1126
toughness was independent of delamination size and
that delamination area could be predicted from peak
impact force generated. Wu and Shyu4 also found that
there was a linear relationship between the peak force
and delamination area and, by extrapolating from the
results, they found a threshold force value for the onset
of delamination.
In their num erical simula tion of impac t-induced
delamination growth, Razi and Kobaysh?* concluded
that mode II was the dominant failure mode for propa-
gation, a view also put forward by Guild et aZ .43.
Fibre ail ure
This damage mode generally occurs much later in the
fracture process than matrix cracking and delamination,
and, as research has concentrated on the low-energy
mode s of dam age, there is less information on this area.
Fibre failure oc curs under the impactor due to locally
high stresses and indentation effects (mainly governed by
shear forces) and on the non-imp acted face due to high
bending stresses. Fibre failure is a precursor to cata-
strophic penetration mode . A simple equation for the
energy req uired for fibre failure due to back su rface
flexure is given by Dorey49 a s:
a2wtL
Energy = 18E
f
where 0 = flexural strength, Ef = flexural modulus,
w = width, L = unsupported length and t = specimen
thickness.
Penetration
Penetration is a macrosco pic mode of failure and
occurs when the fibre failure reach es a critical extent,
enabling the impac tor to completely penetrate the
materia159. Research into penetration impact has mainly
concentrated on the ballistic range; however, some low-
velocity impact work has been performed. Cantwell and
Morton32 showed that the impact energy penetration
threshold rises rapidly w ith ‘specimen thickn ess for
carbon fibre-reinforced plastic (CFR P). They also ana-
lysed the penetration process to calculate the energy
absorbed by shear-out (i.e. removal of shear plug),
delam ination and elastic flexure. This simplified analysis
predicted shear-out as the major form of energy absorp-
tion (50-60% depending on plate thickness).
El-Hab ak6’ tested a variety of glass fibre-reinforced
plastic (GFRP) composites at penetration loads and
concluded that the glass fibre treatment played a key
role in determin ing the perforation load. Whilst the
matrix had little effect, polyester wa s preferable to epoxy.
Dorey49 provided a very simplified analytical mode l of
penetration to give the energy absorbed as:
Energy = qtd
where y = fracture energy, d = diamete r of impacto r,
and t = plate thickness.
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Review of low-velocity impact properties: M. 0. W. Richardson and M. J. Wisheart
Damage m odes in randomly orientedfibre laminates
Most of the work reported above w as on laminates
consisting of unidirectional plies with varying fibre
orientation. In layers in which the fibres are unidirec-
tional it is quite straightforward to predict the orienta-
tion of matrix cracking. W hen the fibres are oriented
randomly, then crack patterns are less easy to establish .
Sheet moulding compound (SMC ) panels and continu-
ous filament mats (CF M) used in pultrusions are common
exam ples of randomly oriented short and long fibre
layers, respectively.
Clearly, a different approac h to defining damage
mode s is required for these composites. In their research
on SMC panels, Liu and Malvern24 found that matrix
cracks on the impacted surface were short and formed a
series of rings away from the point of contact, and
deduced that these were caused by the tensile strain wave
moving out from the centre of impac t. Both Chaturve di
and Sierakowski6’ and Khetan and Chang62 performed
work on glass/polyester SM C panels with air-gun equip-
men t (i.e. high velocity). W hilst the latter au thors
suggested that damag e could be quantified by a ‘damage
area’, the former au thors concluded from tensile residual
strength tests that more information was required on
failure mo des to be able to predict stiffness and strength
degradation. Liu
et a1.63,
in their work on the repair-
ability of SM C comp osites for the automotive industry,
defined three types of impac t-induced dam age: (1)
indentation (crushing of matrix under the impactor),
(2) bending fracture, and (3) perforation (i.e. dam age
resulting from penetration and associate d fracture).
INFLUENCE OF CONSTITUENTS ON THE
IMPACT RESPONSE OF COMPOSITE
MATERIALS
A fibre-reinforced compo site consists of two major con-
stituents (fibre and matrix) an d the interphase region,
which is the area of bond between fibre and m atrix.
The properties of each of these constituents affect the
threshold energies or stresses required to initiate the
different failure modes induc ed by impac t.
Fibres
Fibres are the main load-bearing constituent, provid-
ing the comp osite with the majority of its strength and
stiffness. The most comm on fibres are glass, carbon and
Kevlar. Carbo n is widely used in the aircraft industry
and in many structural applications as it has the highest
strength and stiffness values; how ever, it also is the most
brittle, with a strain to failure of 0.5 to 2.4%. Glass fibres
have a lower strength and stiffness but have a higher
strain to failure (~3.2%~ ) and are less expensive than
carbon fibres. The mec hanical properties of Kevlar lie
between those of carbon and glas?. Carbon’s design
ultimate allowab le strain is only 0.4% currently, whilst
improvements in damage tolerance performance would
allow a 50% improvement on this@‘.Thus a great deal of
the fibre’s superior performance characteristics cannot
be taken advantage of owing to its weakness with respect
to impact.
For resistance to low-velocity impa ct, the ability
to store energy elastically in the fibres is the fundam ental
parameter22.
This corresponds to the area under the
stress-strain curve, which is dictated by the fibre
modulus and failure strain. E-glass can therefore
absorb approxima tely three times the elastic energy of
carbon. Hybrid comp osites are often formed by adding
glass or Kevlar55>65
o carbon comp osites to improve
impact resistance, but moduli mismatching between
fibres increases the complexity of the design of hybrids.
Strain rate sensitivity of glass jibres.
There is con-
flicting information in the literature regarding the strain
rate sensitivity of glass fibres. In general, carbon fibres
are thought of as not being strain rate dependent20’48’67%68
and glass fibres as having a modulus and stiffness which
increase with strain rate48,68s70.Howeve r, in their review
in 1983, Sierakowski and Chaturvedi” concluded that
there was not enough information available to assess
the role of rate sensitivity of comp osite systems. That is
still the case today.
In their impact tests from 1 to 5.5ms-‘, Caprino
et a1.72
reported no strain rate effects for glass cloth/
polyester. However, over a wider strain rate range,
Sim s73 reported increasing flexural strength s for a glass
mat/polyester lamin ate over a range of 10e6 to 10-l m s-’
displacem ent rate for Charpy testing of com posites.
Hayes and Adam s74 constructed a specialized pendu-
lum impactor to study tensile strain rate effects, as
impact speeds increased from 2.7 to 4.9 m
s-’
for glass/
epoxy. They also performed static tensile tests. The
elastic modulus and strength in general increased w ith
impact velocity, but the trend was not consistent through-
out the dynamic range, an d the values at static loading
did not support this trend. In contrast to the belief that
carbon fibres are non-rate depend ent, they reported that
graphite/epoxy’s modulus decreased w ith impact speed
and at dynamic loads the ultimate strength and energy to
ultimate stress were lower than the static values.
Li et af.75 reported an increase in the tensile and
comp ressive strength and stiffness for glass from quasi-
static rates of strain (0.001 to 10~~‘) to high-velocity
impa ct (350 to llOOs_‘). They also noted through-
thicknes s strength increases for glass weave. How ever,
low-velocity impa ct indu ces stra in rates which lie in a
muc h narrower range than the work of Li; therefore, in
general, for low-velocity impac t testing fibre strain rate
effects can be ignored.
Matrix
In an FRP the polymeric matrix (usually a thermoset)
provides several key functions: it transfers the load to the
fibres, protects the fibres from damagin g themselv es an d
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aligns/stab ilizes the fibres. The majority of structural
application s employ epoxy resins as they meet the hot/
wet comp ressive strength requirements. However, epoxy
is brittle an d has poor resistance to crack growth (i.e.
delamination). Attempts to reduce matrix damage and
improve the interlamina r fracture toug hness of thermo-
set resins has involved incorporating plasticizing modi-
fiers, or adding rubbe r or thermop lastic particles to
the resin66. Howeve r, increa sed interlam inar fracture
toughness invariably reduces mechanical properties and
improvements made to the pure matrix are never trans-
ferred fully to the comp osite due to the presence of the
brittle fibres, which prevent growth of plastic zo nes in the
matrixz2. The inclusion of a thin discrete layer of very
tough, high shear strain resin can also be employed to
minimize delamination76.
The use of thermop lastic resins [e.g. poly(ether ether
ketone)] c an give an order of mag nitude increase
in fracture toughn ess over thermoset comp osites. Low
thermal stability and chem ical resistance, poor fibre-
matrix interfacial bonding and creep problem s have his-
torically prevented the use of thermop lastic compo-
sites66. The need for new production techniques still
holds back the use of thermoplastics, but as these
problems are overcome so thermoplastic-based compo-
site systems become more com petitive.
In t e r p h a s e r e g i o n
The interphase region betwee n fibre and matrix is of
vital importanc e. Usually, the surfaces of carbon fibres
are treated with an oxidative process in order to improve
the level of adhesio n between fibre and matrix, whilst
glass fibres are treated with a coupling agent. The inter-
phase region can affect the failure mode which occurs at
a given load; i.e. poor adhesio n results in failure a t low
transverse stress, leaving clean fibres. The bond strength
can be manipulated to improve the toughness by absorb-
ing energy in fibre-matrix debond ; ho wever, this reduces
the mech anical properties.
IMPACT PERFORMANCE OF COMPLEX
GEOMETRY SPECIMENS
The vast majority of impa ct te sting has been performed
on flat coupons in beam or plate format, either clamp ed
or simply supported. By simplifying the geometry, struc-
tural effects are minim ized and more information can be
gleaned about the material behaviour. However, many
composite components have a complex geometry and the
response of stiffened panels is a particularly importan t
area.
Dorey4* reported that the energy to cause BV ID
dropped significantly near the stiffeners, where the struc-
ture was less compliant, and that the stiffeners caused
damage to spread asymm etrically, as would be expected
over an area of non-uniform stiffness. Davie s and co-
authors27946 tated that impact forces will be higher in the
1128
stiffened regions, but that reduced deflections may lead
to sma ller strains and therefore less strain-induced
failure. At the edge of the stiffeners delam inations were
formed, whilst impacts directly over the stiffener caused
debonding between plate and stiffener. The damage
tended to extend down the stiffener (because the induced
forces will follow the stiffest path), which wo uld hav e
disastrous effects for a comp ression-loaded panel. Crater-
ing also occurred due to the very high forces induced in
the stiffened regions. Due to fear of stiffener-panel
debond, many manufacturers are using mechanical join-
ing technique s to avoid this problem, indicating that it is
an area of some concern.
POST-IMPACT RESIDUAL STRENGTH
As stated previously, due to the susceptibility of com-
posite materials to impact damage, dramatic loss in
residual strength and structural integrity results. The
term d ama g e t o l e ra n c e refers to a system’s ability to
perform post-impact. Even BVID can cause strength
reductions of up to 50%. Re sidual strengths in tension,
compression, bending an d fatigue will be reduced to
varying degrees depending on the dominant damage
mode.
Re s i d u a l t e n s i l e s t r e n g t h
Resid ual tensile strength2 3 normally follows a curve as
shown in
F i g u r e 4 .
In region I, no damage occurs as the
impac t energy is below the threshold value for dam age
initiation. Once the threshold has been reached, the
residual tensile strength reduces quickly to a minimum in
region II as the extent of dam age increases. Region III
sees a constant value of residual strength becaus e the
impact velocity has reached a point w here clean
perforation occurs, leaving a neat hole. In this region
the tensile residual strength can be estimate d by con-
sidering the dama ge to be equivalent to a hole the size of
the impactor. The minimum in region II is less than the
constant value in region III because the damage spreads
over a larger area tha n is produced at a higher velocity
when the dam age is more localized (resulting in a cleaner
hole)48. As the fibres carry the majority of tensile load in
the longitudina l direction, fibre dam age is the critical
Residual
Tensile
snength1\_.
II : II
III
Y
Impact Energy
Figure 4 Characteristic residual strength
v rsus
impact energy curve
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Re v i e w o f l o w v e l o c i t y im p a c t p r o p e r t i e s : M . 0 . W . R i c h a r d s o n a n d M . J . W i s h e a r t
damag e mo de. C aprino68 developed a linear elastic frac-
ture mechanics model to predict residual tensile strength
as a function of impa cting k inetic energy which gav e
good correlation to experimental results.
Residual compressive strength
Poor post-impact compressive strength (PICS) is the
greatest weakness of composite laminates in terms of
residual properties. This is mainly due to local instability
resulting from delamination causing larg e reductions in
compressive strength48’n.
As delamination can be pro-
duced by low-energy impacts, large strength reductions
in compression can occur for BVID . D elamination
divides the laminate into sub-laminates which have a
lower bending stiffness than the original laminate and are
less resistant to buckling loads&. Under a compressive
load, a delamination catl cause buckling in one of three
mod es23: global instability/buckling of the lamina te,
local instability (buckling of the thinner sub-laminate),
or a combination of the above. The mod e of failure
generally changes from global, to local, to mixed mode as
the delamination length increases .
lOm s_’ which are ordinarily introduced in the labora-
tory by mechanical test machines such as the IFWIT
technique. The contact period is such that the whole
structure has time to respond to the loading. The modes
of impact damage induced range from matrix cracking
and delam ination through to fibre failure and penetra-
tion. Damage mode interaction must also be understood
when attempting to predict initiation and propagation
of a particular form of dam age. Toughened resins or
thermoplastics can reduce matrix-dominated damage
but the fibres have the most bearing on impact response
and, over the narrow velocity range under consideration,
the strain rate sensitivity of fibres can be ignored. Post-
impact performance is related to the major damage
mode, therefore a com bination of tension an d compres-
sion residual strength testing is required to characterize
the laminate.
PICS testing is often avoided due to the dilhculty in
providing a large enough gauge section to accomm odate
the damage. This necessitates the use of complex anti-
buckling guides which must support the specimen to
prevent global buckling, but at the same time must not
prevent local instability .
Much research has been performed on simple geo-
metry carbon/epoxy cross-ply lamin ates consisting of
plies at various fibre orientation, due to their importance
in the aerospace industry. The low-velocity impac t
response of random fibre/unidirectional lamina te com-
binations (such as are found in pultrusions) and impacts
on complex geometry are less well documented, and
progress is required in these areas if compo site lamin ates
are to be employed in more structu ral application s.
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Residual flexural strength
Less
work has been
done in this area, but it has been
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phite/epoxy exhibited no losses until complete failure
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Jones et a1 79 eported tha t comp ression-comp ression
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10
11
12
13
14
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17
CONCLUSIONS
18
19
‘Low-velocity impac t’ refers to impa cts in the range 1 to
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