International Encyclopedia of Composites-Rev290110

8/13/2019 International Encyclopedia of Composites-Rev290110

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It is obvious that the mechanical performance of continuous-fiber- or fabric-reinforced

polymers is superior to that of discontinuous-fiber versions. This disadvantage, which is

attributed to restricted load transfer between the matrix and the fibers, is compensated by

other benefits, i.e. by design freedom, easy processing via injection and extrusion moldings.

Therefore, it is not surprising that the development of discontinuous fiber-reinforced

thermoplastics is well reflected by a steady increase in the aspect ratio (length to diameter,

l/d) of the fibers both in the parent granules and molded parts. The l/d ratio of short fiber-

reinforced thermoplastics (SFRTPs) produced by extrusion melt compounding technique was

20 earlier, nowadays it lies at 50. The next milestone in the development of SFRTPs was

achieved by pultrusion and powder coating techniques, through which granule size fiber

length was set. The related products are termed to long fiber-reinforced thermoplastics

(LFRTPs). In their injection and compression moldable grades the initial aspect ratio of the

discontinuous reinforcement (usually glass fiber) is 1000 and 2500, respectively.

It is doubtless true that the microstructure of injection-molded composites strongly depends

on the processing mode and its conditions. It is also well known that the mechanical

properties of plastics depend on the testing conditions, especially frequency and temperature.

Therefore, these aspects have to be considered when the fracture and failure performance of

discontinuous fiber-reinforced thermoplastics are discussed [10].

The mechanical performance of discontinuous fiber-reinforced thermoplastics is affected by

the followings (cf. Figure 1):

1. Composition and morphology2. Type and amount of the reinforcement

3. Interface (or interphase) between matrix and reinforcement

4. Processing methods and conditions

5. Testing conditions

There is a strong interrelation amongst items 1) to 5). For example both the matrix

morphology and reinforcement structuring may be highly dependent on the processing

methods as in the case of injection molding. On the other hand, the type and amount of the

reinforcement dictate the selection of both suitable processing methods and conditions. Items

1) to 5) list some matrix-, reinforcement-, interface- processing- and testing-related factors

and serve at the same time as an outline for this contribution. The mechanical tests are

grouped into static and dynamic fracture with monotonic increasing load and static and

dynamic (cyclic) fatigue measurements. The test results are interpreted based on fracture

mechanical concepts.

FIGURE 1 Factors influencing the fracture mechanical performance of discontinuous fiber-

reinforced thermoplastic composites

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Development of Microstructure

Changes in the molecular orientation and crystallization behavior in neat and matrix polymers

of (!)"#T$s occur during processing. This is accompanied with fiber structuring (i.e.

orientation and layering) in case of the reinforced grades. %lthough these changes are rather

complex& the resulting microstructures can be explained by the viscoelasticity of the melt and

by the melt flow fields evolved in the mold. The viscoelastic behavior of the melt depends on

several parameters of the polymeric material (molecular weight and its distribution& main

chain flexibility&conformation possibilities of the chain& etc) and the processing conditions

(melt and mold temperatures& plunger speed) that affect the orientation and relaxation of the

polymer. "or the flow field consisting of shear and elongational flows& processing conditions

are not the only important factors' the mold construction (sprue& runner& gate& and cavity

geometry inducing converging and diverging flow during processing) is also relevant.

t is widely accepted that fiber orientation in discontinuous fiberreinforced thermoplastics

can ade*uately be described by the model of Tadmor +11,& which involves the fountain or

volcano effect discussed by #ose +12,. %ccording to this model& the fiber orientation pattern

produced by in-ection molding can be approximated by a threelayer laminate structure. This

is depicted schematically and as it loos in practice in "igure 2. n the surface () layers&

fibers are oriented parallel to the mold filling direction (/"0). This is caused by the shear

flow of the melt along the *uicly solidified layer at the mold wall. n the central (C) layer&

fibers adopt an orientation perpendicular to the /"0 in the plane of the molded pla*ue. This

ind of alignment is due to the elongational flow at the midplane of the cavity. "actorscontributing to this elongational flow are diverging flow at the cavity entrance and the

fountain effect described by #ose +12,. %n additional argument for the transverse fiber

orientation in the C layer was found in the s*ueeze flow of the melt during the pacing stage.

n the literature& examples of a more complex layering of the particulate reinforcement can be

found& as reviewed +11,. 3uite often a random fiber orientation can be produced in the

solidified layer at the mold wall. n the subsurface layer& however& fibers are aligned in the

/"0 is a result of the shear flow evolved in this region. The splitting of the layers in this

way yields a 4fiveply4 laminate structure. The fiber layering can be even more complicated&

since particulate fillers tend to migrate toward the midplane of the molding& where flow

speeds are higher +15,. This change& attributed to normal stress effects& again modifies the

flow profile and thus the layering and orientation of the discontinuous reinforcement.

FlGURE 2:"iber oriention resulting from in-ection molding (a) for 6 wt7 (819. vol 7)

long :" reinforced polypropylene ($$)' (b) schematically. This picture illustrates position and

designation of the compact tension (CT) specimens preferentially used. ;ote that the

designation of the CT specimens considers the loading notching (longitudinal& ! or

transverse& T) directions in respect to the /"0.


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#esults of numerous investigations carried out on in-ectionmolded pla*ues of mm

thicness indicate that (cf. "igure )

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t can be concluded that the microstructural parameters of reinforced in-ectionmolded

composites are fiber layering& fiber orientation (the two latter are commonly termed fiber

structuring)& fiber volume fraction& and effective fiber aspect ratio and its distribution

"Design" of Microstructure

%mong the guidelines for processing of "#T$s and !"#T$s& priority is given to processing

parameters and mold constructions that contribute to preserving the initial aspect ratio& that is&

the fiber length of the reinforcement. %voiding fiber breaage re*uires molding at minimal

frictional heating. En a given reciprocating in-ection molding machine this can be achieved

by slow screw rotation& low in-ection speed& low bac pressure& and high barrel temperature.

$rocessing of !:" reinforced thermoplastics is very similar to that of :" composites. t is

recommended& however& that a 1626FC higher barrel temperature and a special 4low wor4

screw be chosen. This screw is characterized by a long feed section with constant root and

wide& deep flights. This section is followed by a low gradual compression zone without

neading or mixing elements' the screw ends in a constantroot metering section with flat

flights. n addition& certain aspects of mold construction have to be considered (short runners&

large film or fan gates).

ervice conditions for "#T$ composite parts often re*uire a given welldefined fiber

structuring. "or in-ectionmolded items& a new techni*ue called multiple livefeed in-ection

molding was developed. n this method& a pacing head is inserted between the mold and the

head of the in-ectionmolding machine. The melt flow& and thus fiber orientation in thepacing stage& can be modified accordingly by a programmable movement of the pistons of

the pacing head that pressurizes the solidifying melt directly +26,.

Computer aided design (C%0) is a new techni*ue that has been successfully applied to

optimization of mold construction for molded parts. n C%0 design of an in-ectionmolded

part& the first step is to visualize the wea sites& that is& the nit lines (supposing a runner and

gate system). The next step is to change the position andDor type of runner and gate so that

nit lines do not evolve or& if this is impossible& are positioned where low stresses in the part

during service can be predicted. The next phase is modelling the flow in the mold& subdivided

into finite elements& and characterizing the melt flow patterns in these mold segments. "or the

calculation of the flow patterns& rheological parameters& determined experimentally& are used.

The flow modelling is repeated in several steps until optimized mold filling occurs. The aim

during extrusion die design is to get the same material flow in all segments of the die resulting

in smooth surfaced& warpagefree extrudates. "or the flow simulations different software

pacages are available.

Microstructural Caracteri!ation

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%s stated before& the microstructural parameters are fiber layering and orientation& fiber

volume fraction& and fiber aspect ratio and its distribution.

"or the determination of fiber layering by imaging of polished sections or thin slices& light

(reflective or transmission)& scanning electron microscopy (>/)& and contact

microradiography are preferred. "or fiber orientation& microwave& Gray diffraction& sonic&

and thermographic measurements can also be used. n "igure 2a >/ micrographs taen

from polished sections along the thicness/"0 (zx) and yx planes are shown.

The evaluation of fiber alignment and mean fiber orientation in a given plane is very time

consuming& as it involves determining the angle distribution under which fibers are aligned. n

this respect& image analysis offers the new possibility of getting information about not only in

plane but also spatial orientations +21,. "iber orientation can be described either by using

mean orientation factors& such as ermans +22,& Hrenchel +2,& and modified ermans +2,&

or by vectors +25,.

The aspect ratio (since the diameter of the fibers is mainly constant& it can be replaced by fiber

length) distribution curves are generally determined from microphotographs of the fibers

taen after burning away the matrix. n many cases the matrix polymer can also be removed

by solvents. nstead of histograms showing the relative fre*uency of fibers in a given length

interwall& the use of envelope curves& either in differential or in integral form& is preferred.

The abovementioned microstructural parameters are 4integrated4 in a reinforcing

effectiveness term (#). This term previously considered the effects of fiber structuring with

respect to the loading direction and the fiber loading +2?,. This was extended later to includethe aspect ratio and aspect ratio distribution +1&1?&1B,& and generalized in the formA

=i im

in

iequifieffpireldl

dl

d

lVfTR

&

&

&&&&&)D(

)D()( (1)

where Trel,iis the relative thickness of the ith layer normalized to the sample thickness (B, see

Figure 31.4), fpeff,i is the effective orientation in the i thlayer calculated using the function of

planar orientation (fp) vs. fp,effintroduced by Friedrich [26], =f&iis the fiber volume fraction in

the ithlayer, (l/d)equ,iis the equivalent aspect ratio in the i thlayer, (l/d)m,iand (l/d)n,iare the mean

mass- and number average aspect ratios in the ithlayer, respectively [16].

It seems that the fracture mechanical response of discontinuous fiber-reinforced plastics can

be appropriately related to this reinforcing effectiveness parameter [14].

Fracture Mechanics

Detailed treatment of fracture mechanics is far beyond the scope of this article; it can be

found in the literature [2-9]. Here only a brief overview is given.

A fundamental aspect of fracture mechanics is that the onset of fracture depends not only onthe applied stress but also on the size of intrinsic flaws that act as stress concentrators. The

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presence of such flaws is the reason, for example, that the real tensile strength of solids,

including polymers, is 1/10 of the theoretical value [4]. Such stress concentration sites are

always present in neat and reinforced molded plastics, either as a result of processing

("notches" at the knit lines resulting from compressed air, bare fiber segments as a result of

imperfect wetting by the matrix, voids caused by differences in the thermal expansion

characteristics of the matrix and reinforcement etc.) or caused by use (scratches, damage by

cutting or shaping, impacts etc.). The common effects of stress and flaw size are combined in

linear elastic fracture mechanics (LEFM) which deals only with bodies that obey the

Hookian law, that is, whose deformation is fully elastic in a term called stress intensity

factor (K) or fracture toughness:

2D1)( aYKK

Ic == (2)

where

I 8 applied stress

a 8 crac size

Y 8 geometrical correction factor taing into consideration the finite size of the specimen used

8 tensile opening mode& mode

According to the LEFM theory, fracture occurs when K I> KIc, where KIc, is a critical value of

the stress intensity factor. This material parameter is also termed fracture toughness. From Eq.

(2) it is obvious that KIis a stress-related fracture mechanical criterion. KI, can be determinedfrom monotonic static loading measurements, for example according to ASTM E 399. In the

case of monotonic dynamic loading, KIccan be computed by Eq. (2) as the slope of the plot

Y against a -1/2.

The other LEFM material parameter is an energy-related one that measures the energy

required to extend the crack over a new surface unit. This term is denoted G Ic, and is called

fracture energy, critical strain energy release rate, or specific crack extension force. The onset

of fracture depends again on whether GI is less than or greater than GIc. Kc and Gc are

interrelated by a function whose exact form depends on the stress state of the specimen (planestress or plane strain). For S(L)FRTPs, G Ic is generally determined from high speed impact

tests - using Charpy and Izod test set ups - through the method of Plati and Williams [27] or

its derivatives (ISO 17281).

The main criterion of LEFM, namely fully elastic deformation, is very severe for plastics that

may undergo pronounced plastic deformation (yielding or tearing) during fracture. In this

case, other approaches, also used originally for metals, were pursued for plastics: J-integral,

crack opening displacement (COD), and essential work of fracture (EWF) [6,28]. These

material parameters are included in plastic, elastoplastic, or postyield fracture mechanics

(PYFM). JIcis an energy-related term connected with the onset of stable crack growth. For

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linear elastic bodies GIc= JIc. JIccan be determined for example by ASTM E813 and ISO/CD

28660. In addition, JIccan be deduced from high speed impact tests carried out on sharply

notched Izod or Charpy specimens [28].

The fracture mechanical approach can also be applied for both static and dynamic (cycling)

fatigue of cracked specimens. Static and dynamic fatigue means slow crack growth under

subcritical stresses and stress amplitudes, respectively; that is, the stress intensity factor and

its amplitude lie below KIc. The aim of both measurements is to establish crack extension

characteristics with respect to the stress concentration at the crack tip as a function of either

time (da/dt static fatigue loading) or number of fatigue cycles (da/dN dynamic fatigue).

The latter characterization can be performed by the ASTM E 647 standard, originally

developed for metals.

For the determination of the critical values of fracture mechanical parameters related to the

plane strain condition, the specimens used have to meet different size criteria. These can be

taken from the corresponding standards; alterations to these standards for discontinuous fiber-

reinforced thermoplastics are summarized in Ref. 14. If these criteria are not met, critical

values of KI, are denoted Kc, instead of KIc; the same designation is used also for Gc, and Jc.

Fracture and Related Failure

In both static and dynamic fracture measurements, breakdown is caused by monotonically

increasing load. The only difference between them is related to the strain rate or frequency

range; however, the threshold value is rather arbitrary. Measurements carried out below across-head speed v of 1 m/min are referred to as static, whereas impact measurements with a

striker speed above 1 m/s are referred to as dynamic fracture tests.

tatic !oading

EFFECT OF MICROSTRUCTURE. Fiber loading. Fiber reinforcement may affect fracture

toughness in different ways. It can be improved, worsened, or held at a constant level by fiber

incorporation, depending on the matrix of the composite [14]. The run of K c, as a function of

Vfcan hardly be predicted, because of competitive micromechanisms that either increase or

decrease Kc. Nevertheless, discontinuous fiber reinforcement is always the right tool to

increase the fracture toughness of low molecular weight polymers prone for brittle fracture.

The effects of reinforcement-matrix bond quality and of matrix toughening are worth

mentioning here. Improving the coupling between fiber and matrix is not necessarily

beneficial for Kc. Strong bonding may hinder the deformability of the composite so that Kc

tends to decrease [29]. This is in accordance with the Hahn-Rosenfield equation [30]. This

equation explicitly shows that Kcdoes not depend solely on strength but also depends on

ductility parameters:

( ) 2D1

LEK BBc = ()where

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E 8 > modulus

b& Jb 8 tensile strength and elongation at brea& respectively

L8 proportionality constant including strain hardening& in length dimension

It is well known that toughening of the matrix results in increased ductility; however, this is at

the cost of stiffness and strength. Thus matrix toughening may also be connected with

deterioration in Kc.

The course of Jc as a function of V fdepends on the corresponding Kc - Vf and E - Vf

functions:

( )E

KGJ

C

cc

2

== ()

provided the plane stress condition satisfies the LEFM theory. If the Kcincrement due to the

square function overcompensates for the increment in E modulus, Jcincreases; when it does

not, the opposite tendency becomes evident. It should be noted here that J lcvalues can be

found scarcely for discontinuous fiber-reinforced thermoplastics in the literature [14,28].

In spite of the very complex fracture mechanical response to fiber loading, the following

conclusions can be drawn:

1. An increase in Kcfrom fiber loading is more probable the higher the E modulus

and the lower the ductility of the unfilled matrix (brittle, low molecular weight,

degraded polymers, especially polycondenzates).

2. For ductile materials a relative increase in both Kcand Jccan be achieved by using

fibers of higher aspect ratio (e.g. LGF instead of SGF).

3. The effects of matrix toughness and fiber-matrix bonding are hardly predictable.For relative improvements in Kcand Jc, the strength and ductility characteristics of

the matrix have to be balanced by the reinforcement.

Fiber structuring. The layering and orientation of the fibers in injection-molded items were

already shown in connection with Figures 2 and 3. On the fracture surface of the specimens,

fibers lying parallel or longitudinal to the crack plane (L fibers) can clearly be distinguished

from those oriented perpendicular or transverse to it (T fibers) (Figure 4).

FIGURE A "racture surface at the razor notch of 6 wt 7 (819. vol 7) :" reinforced

in-ectionmolded $$. (n this T! type CT specimen& ! fibers can be found on the surface&

whereas T fibers in the central layer& as indicated' cf. "ig. 2. #azor blade notch is mared by

arrow.)

It is doubtless true that the anisotropic structuring of the fibers yields different fracture

mechanical values when specimens with various notch directions (T and L; see Fig. 2b) are

tested [14,17,26,29]. The load bearing capacity of T fibers aligned in the load direction is

considerably higher than that of the L fibers, which have practically no reinforcing effect.

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Therefore, the fracture mechanical response depends on the relative thickness of the layers

containing T and L fibers, respectively.

The degree of fiber orientation in these layers is also important. T fibers completely aligned in

the load direction guarantee the best stress transfer and thus the greatest reinforcement. Fiber

misalignment along the load direction necessarily reduces the overall reinforcing effect.

Friedrich introduced an effective fiber orientation term that takes this fact into account cf.

Figure 5 [26].

FIGURE #: #elationships between the effective (fp&eff) and planar fiber orientations (fp)

considering the actual mechanical loading direction

Many investigations carried out on SGF and LGF composites (e.g., Refs. 14,16,17,18) have

indicated that the anisotropy in the mechanical response of the LGF-reinforced systems is not

very pronounced, in spite of the fact that the three-ply laminate structure caused by the

injection molding still exists. This observation suggests the important role of the aspect ratio.

Fiber aspect ratio. The influence of filler shape on fracture toughness at a given filler loading

strongly depends on the matrix characteristics [29]. However, with increasing fiber aspect

ratio Kcalways increases, at least above a given threshold l/d. This is connected with an

increase in the loadability of the discontinuous-fiber-reinforced composites, since their

strength increases with increasing aspect ratio [31]. The l/d ratio can be increased either by

using longer fibers of the same diameter or by using smaller diameter fibers of the same

length. Reinforcing with small diameter fibers is beneficial only in a given aspect ratio range.

This is due to the fact that the deleterious effect of stress concentration at the fiber ends

should be compensated for by toughness-enhancing effects, such as increasing interface,

improved stress interaction between fibers, and enhanced crack path [32]. A decrease in the

critical fiber length for smaller diameter fibers is expected to lead to an increase in the clack

path length, which promotes fiber pullout and reduces fiber fracture [14].

A definite answer on the effect of fiber aspect ratio distribution cannot be given. There are,

however, several indications [26,32] that use of reinforcing fibers with different aspect ratios

as a result of varying diameters can be beneficial for fracture mechanical characteristics.

EFFECTS OF TESTING CONDITIONS. Temperature. The fracture toughness, measured at

low cross-head speed, decreases as a function of temperature T for both matrix and its fiber-

reinforced versions [16,29]. A steeper decrease in the related plot can always be found in the

vicinity of the glass transition temperature (Tg) of the matrix. Here the enhanced molecular

mobility of the matrix induces a change in deformation mode from ductile to viscous. So, T g

is always the upper threshold for the applicability of the LEFM theory. K cvalues calculated

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according to Eq. (1) for temperatures above Tg no longer have meaning for fracture

toughness; they can be treated only as trend data of a mechanical property thus defined.

Cross-head speed. At v = 1000 mm/min, the trend of Kcwith T is basically different from that

at v = 1 mm/min. Kcvalues start at a relatively low level in the subambient temperature range

before they increase at the Tg [29, 33, 34]. This is attributed to a clear change in the stress state

of the samples (plane strain to plane stress) and related failure manner (brittle to viscous as a

result of adiabatic heating at the crack tip).

FAILURE BEHAVIOR. It is obvious that big differences due to testing conditions are related

to substantial changes in microscopic breakdown events. The microscopic failure mechanisms

occurring in S(L)FRTPs are shown schematically in Figure 6. They can be grouped into

matrix-related (crazing and shear yielding) and fiber-related (fiber fracture, pullout, and

debonding) events. For longer-fiber-reinforced injection-molded composites, the latter can be

extended by fiber bridging (unbroken fibers connect the crack sides) and by cleavage and slip

of fibers within bundles or rovings during debonding and pullout. It should be noted here that

the relative orientation of the fibers (L and T) strongly influences the relative probability of

the individual fiber-related energy absorption mechanisms [14,29,35].

FIGURE $: "ailure mechanisms for discontinuousfiberreinforced thermoplastic at a

microscopic level

Based on fractographic results, the characteristic failure modes of the composites can be

summarized in failure maps. In such maps the dominant matrix- and fiber-related breakdown

processes are indicated as a function of the testing conditions T and v [18,35,36]. Failure

maps not only give guidance for selecting composites for particular circumstances but also

suggest methods of increasing the toughness. In the patent 1iterature one can find abundant

examples and ideas for such improvement.

Dynamic Loading

EFFECTS OF MICROSTRUCTURE. Fiber loading. The plot of dynamic fracture toughness

as a function of Vf can be as complex as that for the static case. Generally the dynamic

fracture energy (Gd) decreases with increasing fiber loading. This effect may or may not be

compensated for by the increasing E modulus with respect to the resulting K d. Since this

effect of the E modulus is very closely matched to the static one, the plot of K das a function

of Vf depends mostly on the Gd - Vf function of Eq. (4). The effect of V f on Kd is

demonstrated in Figure 7 for SGF- and LGF-reinforced composites. It can be concluded that

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increasing the aspect ratio of the reinforcement yields an improvement in Kd. The Gd - Vf

function, on the other hand, should depend on the matrix toughness and fiber-matrix bond

quality.

FIGURE %:Change in the static (Hc) and dynamic (Hd) fracture toughness as a functionof fiber loading (=f) at ambient temperature for :" and !:"reinforced in-ectionmolded

$$ (Hc determined on CT specimens at v 8 1 mmDmin& whereas Hd deduced from lzod

measurements)

At this point attention has to be paid to the internal flaws induced by specimen machining

(sawing, cutting, polishing, and the like). The presence of such flaws causes trouble during

curve fitting and a very big scatter in Gd, especially for long fiber-reinforced systems.

However, if failure of the specimens occurs not at the notch introduced but near to it means

that cutting introduced a flaw (either by debonding or by intrabundle fiber cleavage) that

acted as a stress concentrator site. In some cases this threshold or latent notch value was

found to be = 0.5-0.6 mm [37].

Fiber aspect ratio. This effect is surprisingly pronounced if one compares K cand Kdfor SGF-

and LGF-filled composites (see Fig. 7). This suggests some differences in the load transfer

during static and dynamic measurements that should be clarified.

EFFECTS OF TESTING CONDITIONS. Gdand Kdvalues derived from Izod and Charpymeasurements are very closely matched. This is due to an analogous stress state during

impact, which is carried out at practically the same deformation rate. The plot of Kdand Gdas

a function of T can be calculated from Eq. (4) provided that the E(T) and G d(T) or Kd(T)

functions are known at the given frequency. Kdgenerally increases with decreasing T as a

result of the increase in E modulus, whereas Gdremains practically constant. The plots of both

fracture mechanical parameters depend on the temperature range investigated; maxima and

minima in the plots can also be found. Their appearance is attributed to primary and

secondary relaxation transitions of the matrix and its components [4,5,38].

FAlLURE BEHAVIOR. It has to be emphasized that dynamic failure mechanisms are the

same as those shown and discussed with respect to static loading. Although failure mapping

has not been performed for dynamic measurements, the following findings are expected:

1. The frequency embrittlement of the matrix promotes brittle matrix cracking, the

onset of which depends on the frequency-dependent Tg. Among the matrix-related

failure mechanisms, crazing is more common than shear yielding.

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2. Among the fiber-related failure events, fiber pullout and fracture tend to dominate.

Their relative proportions depend not only on the testing conditions but also on the

fiber-matrix bonding.

Fatigue and Related Failure

S(L)FRTP parts are widely used in fields in which constant and cyclic subcritical loading

occur. The response to long-term constant loading (static fatigue) is often called either stress

corrosion (SCC) or environmental stress corrosion cracking (ESC). In this case, crack

propagation is induced in different environments at the condition K 0< KIc, where K0denotes

the initial stress intensity factor. The result of this kind of measurement is either the stress

corrosion threshold (KI,SCC) or the crack growth rate (da/dt vs. KI), or both. KI,SCCmeans a

threshold KI, value below which no crack growth takes place in a given environment.

In dynamic or cyclic fatigue, the crack growth rate per cycle is established as a function of thestress intensity factor amplitude ( K). A threshold value K th, which is connected with the

onset of fatigue growth, can be read from the fatigue crack propagation (FCP) curve. The

characteristics of static and dynamic fatigue with respect to the measurements and results to

be discussed are summarized in Figure 8.

FIGURE 8:The fatigue behavior of injection-molded composites.

Static Fatigue

EFFECT OF MICROSTRUCTURE. When Kc, from static fracture, increases with V f, one can

expect a similar trend in the resistance to SCC for the given composite. Incorporation of a

rubbery impact modifier in the matrix results in further improvement in SCC resistance of the

related composite. Fractographic analysis supports the conclusion that this is due to better

wetting of the fibers and thus better protection of them against acidic attack. On the other

hand, when Kcdecreases with Vf, an acceleration in SCC growth can be predicted in relation

to the corresponding matrix. The da/dt-KI curves, or at least their segments in double

logarithmic representation, can be approximated by straight lines. This indicates the validity

of the Paris-Erdogan relationship (often called the Paris power law [3,4,39]):

m

IK

dt

da= (5)

The crack growth kinetics depend on both the microstructure and the environment [29,40,41].

Note that the final breakdown of the specimens (usually CT type) does not necessarily occur

at the Kcderived from static loading measurement.

In SCC tests, KIat the specimen breakdown may be higher than Kcwhen a material with

rather ductile behavior is investigated. This is a consequence of the very low frequency of this

kind of measurement, which generates a process or damage zone that is fundamentally

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different from that found in static fracture. Alterations can be observed in both the size of the

zone and the related failure mechanisms therein. It can be concluded that final breakdown of a

specimen with ductile features (either due to the material or due to the stress state) takes place

near the static Kmaxvalue. This value can be determined by Eq. (1) using the maximum load.

Such response can be observed for composites either in air or in surrounding media that

plastify their matrix. Final failure of the specimen may also occur at K I< Kcas a result of

aggressive attack by a given environment. If the initial stress intensity factor K 0is very small

but higher than KI,SCC, the surrounding medium can penetrate deeply into the specimen,

causing failure events (e.g. multiple fiber breakage, surface degradation) that decrease the

SCC resistance (diffusion-assisted SCC). Therefore, it is highly reasonable to always indicate

K0. For composites with a ductile matrix or one that is "ductilized" during the measurement,

the crack tip at its advance can hardly be resolved. Instead of a sharp crack, a well-evolved

damage zone can be observed. The propagation of this zone as a whole can be treated and

adequately described by the crack layer theory (e.g. [42]).

Changes in the time to failure curves due to microstructural and environmental effects are

very similar to those discussed above with respect to SCC growth [40,41]. However, final

failure may occur at a Kcwhich is either independent or dependent of K0. The latter case

suggests that failure depends additionally on K0, that is, on the corrosion loading history of

the specimen. This corrosion loading history is rather complex, since it involves the

immersion time and all changes in both the structure and the stress state that are caused by the

diffusion and penetration of the environment. This observation means that Eq. (5) no longerholds, since da/dt depends in addition on K0; therefore, talking about a material parameter

according to the fracture mechanical concept is very questionable.

EFFECT OF ENVIRONMENT. Both the da/dt-KI, and the K0-time (t) curves depend strongly

on the environment. They can be grouped by whether their degradative attack relates mostly

to the matrix, to the fibers, or to the fiber-matrix interface. In addition, the SCC response is

highly affected by the corrosion loading history.

FAILURE. First impressions about the failure mode can be got from the surface appearance

of the broken specimens. A nearly planar fracture surface indicates fiber degradation, which

mostly occurs in acidic environments [43,44]. A zig-zag fracture surface path, reflecting the

fiber alignment, on the other hand, illustrates matrix and matrix-fiber interface attacks. The

failure micromechanisms during SCC usually agree with those in static fracture; however,

their relative occurrence varies considerably with the aggressive nature of the environment

and the K0of the test.

Dynamic Fatigue

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Since flaws are always present in S(L)FRTPs, one can conclude that the main constituent of

durability is the propagation of such flaws rather than their initiation and development. In this

case, the results of FCP measurements are relevant and important from the point of view of

construction.

EFFECT OF MICROSTRUCTURE. Fiber loading. It should be emphasized that FCP

resistance due to fiber loading changes along with the Kc-Vf function. Increasing fracture

toughness thus correlates with improved FCP resistance (Fig. 9), while decreasing K cwith Vf

yields FCP acceleration [14,29]. Figure 9 again shows the validity of the Paris power law (Eq.

6) for the stable FCP range:

mK

d!

da)(= (?)

FIGURE &A Changes in "C$ behavior due to microstructural parameters& schematically

The exponential term m of Eq. (6) generally increases with V f(Figure 10). This change is

accompanied by a shift of K th, toward higher K values; that is, fiber incorporation enhances

the threshold limit below which no crack growth takes place. The same trend in K thcan be

observed when results achieved on L- and T-cracked specimens are compared. This is due to

a change in the stress state of the specimen as it approaches the plane strain condition as a

result of increasing fiber loading and load direction aligned fiber structuring. The onset of

unstable crack growth (final fast fracture) occurs near either Kcor Kmax, just as it does in static

fatigue. This upper limit and the stable FCP range itself are strongly affected by the

viscoelasticity of the material under the given testing condition (e.g., crack tip heating effects

[45]).

FIGURE 1'A $aris range in the "C$ curves of $$ and its :" and !:"reinforced grades at

different fiber volume fractions (indicated in vol.7)

Fiber structuring. L-T specimens, because of the higher quantity of T fibers oriented in the

load direction (see Fig. 2b), exhibit higher FCP resistance than T-L-cracked ones. This

finding is again analogous to the results of static loadings. The effect of fiber structuring

becomes more and more pronounced with increasing Vffor SGF than for LGF reinforcement.

This supports the statement made before in connection with monotonic loading, namely, that

incorporation of LGF diminishes the mechanical anisotropy.

Fiber aspect ratio. When the FCP curves of SGF- and LGF-reinforced PPs are compared a

clear improvement in FCP resistance can be seen as a result of the use of fibers of higher

15


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aspect ratio (Figure 9). The relative improvement diminishes with increasing Vf. FCP

enhancement of LGF reinforced systems was attributed to the evolution of a more extended

damage zone and longer debonding and pullout routes [46].

The effects of the above microstructural parameters on the FCP behavior are summarized

schematically in Figure 9. For the microstructural interpretation of the FCP response, a

qualitative [47] and a quantitative model [16,18,29] exist. The latter is based on the

reinforcing effectiveness (R) and microstructural efficiency (M) concepts [16,26]. Figure 11

shows the difference in the FCP rates between SGF- and LGF-reinforced PP composites as a

function of the M values. One can clearly see that with increasing M the crack rate, at a given

K value, decreases markedly [18].

FIGURE 11: "C$ rates of :" and !:"reinforced $$ grades as a function of themicrostructural efficiency (/) at a given stress intensity factor amplitude (H81.5 /$am1D2)

EFFECTS OF TESTING CONDITIONS. Information on the effects of external testing

conditions (waveform, frequency [47], main load, temperature, environment, etc.) on the FCP

response of SFRTPs is very limited. These effects seem to be highly material-dependent;

therefore, general conclusions cannot be drawn.

FAILURE. Analogies between the fracture mechanica1 responses given for fracture and forfatigue suggest that the individual failure mechanisms at the microscopic level are the same

(see Fig. 6). The first step in failure is again debonding at the fiber ends. This is followed by

pullout and further debonding for T- and L-oriented fibers, respectively. The fatigue crack

path in Figure 12 demonstrates this. In this picture, stress concentration at the fiber ends is

clearly perceptible. In the damage zone, in addition, a stress concentration field was

developed and preserved by matrix deformation. The matrix underwent crazing and shear

yielding, initiated by fiber debonding and pullout processes.

FIGURE 12:"atigue crac profile on the surface of a CTspecimen of 26 wt 7 (8 B. vol 7)

:" reinforced $$A ()behind the crac tip' (B)before the crac tip. (Crac direction from

left to right.)

Among the discrepancies between fracture and fatigue, the size of the damage zone and the

sharpness of the crack have to be mentioned; the zone is smaller and the crack is sharper

during fatigue [33].

Further differences between the failure modes are connected with crack tip heating and

accompanying changes in both matrix- and fiber-related micromechanisms caused by cyclicloading. At the onset of stable FCP, the matrix fails in a semibrittle manner with multiple

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fractures. Therefore, the mean pullout length here is relatively small. In addition, this multiple

matrix fracture yields a stress field that promotes fiber breakage. At the end of stable FCP, the

ductility of the matrix increases considerably; crazed arrays and viscous torn matrix parts can

also be evident. Among the fiber-related mechanisms, pullout dominates-however; with an

increased average length [48,49].

On the FCP curves of many S(L)FRTPs, a stable delayed crack growth region can also be

resolved just before stable acceleration crack growth occurs (see Fig. 13a). This is connected

with an evolution and stabilization of the damage zone. This is the right place to call the

attention to differences in cyclic and static fatigue responses. Under static conditions the

stable crack deceleration occurs at much higher actual stress intensity factor compared to the

cyclic one (see Fig. 13b). This is due to the formation of a more extended equilibrium

damage zone the formation of which was supported by the fact that the apparent frequency of

static fatigue is lower than the cyclic one [18,48]. Under apparent frequency the reciprocal

value of the time causing the specimen fracture is meant.

FIGURE 13"table crac deceleration ranges registered during cyclic (a) and static fatigue

(b) for :" and !:"reinforced $$ grades with different fiber volume contents

The fiber-related micromechanisms strongly depend on the relative fiber angle to the crack

plane [29,35,50]. In addition, crack closure may also appear, especially in T-fiber regions

with higher fatigue crack lengths. This decelerates the FCP rate and is connected with a mixed

mode (modes I and II) stress state. It would be highly reasonable to construct fatigue failure

maps analogous to those for fracture. Because experimental results are lacking, however,

mapping cannot be performed yet.

Design Aspects

It has been shown above that the fracture mechanical characterization of injection-molded

thermoplastics and their composites contributes to a better understanding of their performance

in different loading situations. Although in many cases real plane strain fracture parameters

can hardly be deduced because the mean thickness of injection-molded parts (3-6 mm) does

not reach the required one, their measured values can be used for design and construction. For

this purpose, however, an adequate fracture mechanical characterization method should be

selected.

The proper choice of method when a fracture-resistant part is to be constructed depends on the

behavior of the material. For composites with high stiffness and low ductility, which fail by

brittle fracture, stress-related terms (Kc) are used, whereas for those of high ductility, energy-

related fracture mechanical terms (Gc, Jc, COD and EWF) should be preferred.

In the case of a part that is to withstand fatigue loading, the first question to be answered iswhether failure occurs mainly by crack initiation or by crack propagation. When the cycles to

1@


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failure at different stresses are plotted for a given material, the response curve can be divided

into two regions: crack initiation and crack propagation (Fig. 14). It was shown above that for

SGF and LGF reinforced injection-molded composites, crack propagation is of basic

importance. In this case, therefore, threshold values (KI,SSC, K th) derived from static and

dynamic fatigue measurements can be used for construction purposes.

FIGURE 1: 0esign factors for fatigueresistant composites depending on whether crac

initiation or crac propagation dominates the failure behavior.

(ummar)

This article has dealt with microstructural development& microstructurerelated fracture& and

fatigue performance at static and dynamic loadings and corresponding failure behavior of

short and long discontinuous fiberreinforced in-ectionmolded composites. The response of

these systems to different loading conditions was treated by fracture mechanical concepts

using stress and energyrelated terms. The accompanying failure was analyzed by

fractography and grouped into matrixand fiberrelated events. n addition& the changes related

to different loading situations were demonstrated and discussed. %ttempts were made to

determine general trends in microstructural development& fracture mechanical response& and

connected failure behavior& and to summarize them schematically. %nalogies and

discrepancies between fracture and fatigue were emphasized and discussed. "inally&recommendations were given for design of in-ectionmolded thermoplastic matrix composite

parts under increasing (fracture) and alternating (fatigue) load. The literature cited offers

further detail on this topic.

+ee also /olding& $olymer n-ection KKK,

L. HargerHocsis

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