POLITECNICO DI MILANO
SCHOOL OF INDUSTRIAL AND INFORMATION ENGINEERING
Master of Science in Materials Engineering and Nanotechnology Department of Chemistry, Materials and Chemical Engineering “G.Natta”
Mechanical characterization of a short glass fiber polyamide blend
reinforced with continuous fiber layers
Supervisors: Roberto Frassine Maria Chiara Marchesi
Rota Giancarlo 817692
Academic Year: 2015-2016
..In una notte senza luna truccò le stelle ad un pilota
e quando l’aeroplano cadde lui disse “E’ colpa di chi muore”..
La cattiva strada, Fabrizio De Andrè
TABLE OF CONTENTS
INDEX OF FIGURES 1
INDEX OF TABLES 2
INDEX OF GRAPHS 2
ABSTRACT 3
1.– INTRODUCTION 5
2. – BACKGROUND ON POLYAMIDES 8
2.1 – POLYAMIDES FROM STRAIGHT-CHAIN ALIPHATIC MONOMERS 9
2.2 – POLYAMIDES FROM LINK-SEPARATED AROMATIC MONOMERS 11
3 - MATERIALS AND METHODS 13
3.1 – METHODS OF PROCESS 13
3.1.1 – CORE PROCESSING 13
3.1.2 – SANDWICH PROCESSING 14
3.2 – DIFFERENTIAL SCANNING CALORIMETRY 15
3.3 – OPTICAL MISCROSCOPY 16
3.4 – MECHANICAL TESTS 17
3.4.1 – THREE POINT BENDING TEST 17
3.4.2 – INTERLAMINAR SHEAR STRENGTH TEST 18
3.4.3 – IMPACT TEST 19
3.4.4 – CLASSICAL LAMINATION THEORY 21
4. - RESULTS 28
4.1 – DSC ANALYSIS 28
4.2 – MICROSTRUCTURE OBSERVATIONS 30
4.3 – FLEXURAL TESTS 32
4.3.1 – THREE POINT BENDING 32
4.3.1.1 – THREE POINT BENDING: EFFECT OF TEMPERATURE 34
4.3.2 – INTERLAMINAR SHEAR STRENGTH TEST 38
4.3.2.1 – INTERLAMINAR SHEAR STRENGTH TEST: EFFECT OF TEMPERATURE 41
4.3.3 – CLASSICAL LAMINATION THEORY 43
4.4 – FALLING WEIGHT IMPACT TEST 48
4.4.1 – FALLING WEIGHT IMPACT TEST: EFFECT OF TEMPERATURE 49
5. – CONCLUDING REMARKS 53
6. – REFERENCES 54
1
INDEX OF FIGURES
Figure 1 - Thermoplastic engine oil pan fitted on ISF 3.8 Cummins Engine 6
Figure 2 - Structural parts of a vehicle 6
Figure 3 - Hybrid seat back-rest: injection molded short fiber-reinforced structure (right) and overmolding
of continuous fiber-reinforced layer (left) 7
Figure 4 - General structure of linear aliphatic polyamides: a) polyamide-x b) polyamide-x,y 8
Figure 5 - Hydrolitic ring opening of ε-caprolactam, followed by polycondensation between two end groups
to give polyamide 6 9
Figure 6 - Polycondensation scheme of polyamide 6,6 9
Figure 7 - Scheme of an hydrogen bonded sheet of PA6 with anti-parallel (a) and parallel (b) orientation of
amide groups 10
Figure 8 - Structures of the α-phase of PA6 and of PA6,6. The left side shows the view of the hydrogen-
bonding planes, and the right side shows the view along the chain axis 11
Figure 9 - Meta-xylylenediamineadipamide (MXD6) and para-xylylenediamineadipamide (PXD6) 12
Figure 10 - Molecular conformation of MXD6 12
Figure 11 - Schematic diagram of a screw injection molding machine 14
Figure 12 - Stacking sequences: a) unidirectional, b) cross-ply 0/90, c) cross-ply 90/0. The arrows indicate
the direction of injection 15
Figure 13 - Mold parts used for the production of the sandwiches 15
Figure 14 - Positions of the sections along the injection molded plate. The red arrow indicates the direction
of injection 16
Figure 15 - Three Point Bending test configuration 17
Figure 16 - Interlaminar Shear Strength test configuration 19
Figure 17 - Left: Lamina in plane state of stress. Right: variation of strains through layer and laminate
thickness typical in-plane stress (a) and corresponding stresses (b) 22
Figure 18 - Basic deformations of the layer: (a) in-plane tension and compression; (b) in-plane shear; (c)
bending; (d) twisting 24
Figure 19 - Preferential fibers' orientations in the injection molded core 24
Figure 20 - Sign convention 26
Figure 21 - Equilibrium of interlaminar stresses in a laminated beam 27
Figure 22 - Schematic rapresentation of the fiber orientation in the core of an injection molded square
plate 30
Figure 23 - Thin section micrographs 31
Figure 24 - Crack development during ILSS testing of UD specimen 39
Figure 25 - UD specimen after ILSS test at T=23°C 40
Figure 26 - CP0/90 (left) and CP90/0 (right) specimens after ILSS test at T=23°C 40
Figure 27 - 00-specimen after ILSS test at T=23°C 40
Figure 28 - Core specimen after falling weight impact test at T=23°C: impacted side (left) and bottom side
(right) 48
Figure 29 - Cross-ply specimen after falling weight impact test at T=23°C: impacted side (left) and bottom
side (right) 49
Figure 30 - Cross-ply specimen after falling weight impact test at T=-20°C: impacted side (left) and bottom
side (right) 52
Figure 31 - Cross-ply specimen after falling weight impact test at T=80°C: impacted side (left) and bottom
side (right) 52
2
INDEX OF TABLES
Table 1 - Melt temperatures of aliphatic polyamides 10
Table 2 - Dimensions of Three Point Bending tests specimens 18
Table 3 - Dimensions of Interlaminar Shear Stress tests specimens 19
Table 4 - Data from DSC analysis 29
Table 5 - Fiber length distribution 31
Table 6 - Mean thickness of the core layers 31
INDEX OF GRAPHS
Graph 1 - DSC analysis of Radistrong 28
Graph 2 - Three Point Bending test of Radistrong core at T=23°C: 33
Graph 3 - Three Point Bending test at T=23°C: 33
Graph 4 - Three Point Bending test at T=23°C: 34
Graph 5 - Effect of temperature on longitudinal core specimens 35
Graph 6 - Effect of temperature on transversal core specimens 36
Graph 7 - Effect of temperature on unidirectional specimens tested with Three Point Bending 36
Graph 8 - Effect of temperature on cross-ply 0/90 specimens with Three Point Bending 37
Graph 9 - Effect of temperature on cross-ply 90/0 specimens with Three Point Bending 37
Graph 10 - Interlaminar Shear Strength test on unidirectional and cross-ply specimens: comparison at
T=23°C 39
Graph 11 - ILSS test: comparison between UD and 00-specimens 40
Graph 12 - Effect of temperature on ILSS unidirectional specimens 41
Graph 13 - Effect of temperature on ILSS cross-ply 0/90 specimens 42
Graph 14 - Effect of temperature on ILSS cross-ply 90/0 specimens 42
Graph 15 - Variation of flexural modulus for core-L specimens (ν12=0.35) 43
Graph 16 - Variation of flexural modulus for core-T specimens (ν12=0.35) 44
Graph 17 - Variation of flexural modulus for UD specimens (ν12=0.3) 44
Graph 18 - Variation of flexural modulus for CP0/90 specimens (ν12=0.3) 45
Graph 19 - Variation of flexural modulus for CP90/0 specimens (ν12=0.3) 45
Graph 20 - Stress distribution across the thickness of core-L samples 46
Graph 21 - Stress distribution across the thickness of core-T samples 46
Graph 22 - Stress distribution across the thickness of UD-samples 46
Graph 23 - Stress distribution across the thickness of CP0/90 samples 47
Graph 24 - Stress distribution across the thickness of CP90/0 samples 47
Graph 25 - Impact energy versus time at T=23°C 49
Graph 26 - Falling Weight Impact Test on core-specimens: effect of temperature 50
Graph 27 - Falling Weight Impact Test on CP0/90-specimens: effect of temperature 50
Graph 28 - Falling Weight Impact Test on CP90/0-specimens: effect of temperature 51
3
ABSTRACT
The mechanical properties of an injection molded polyamide blend reinforced with short glass fibers were
investigated. Flexural tests as three point bending and interlaminar shear stress were utilized together with
falling weight impact test. The same tests were also performed on sandwich structures in which the
discontinuous fiber reinforced blend is embedded between continuous glass fiber reinforced polyamide
tapes. Then comparisons were made between bare core, single-layers and cross-ply sandwiches. Moreover
the change in mechanical response with temperature was explored.
Microscopy observations of thin section of the core material reveal the creation of a three-layer
configuration, inherently present in injection molded samples, due to a particular distribution of the short
fibers. This distribution is seen to influence the mechanical behavior of the core material and it is more
evident with high temperatures. On the other hand, the application of continuous fiber-plies demonstrates
to be responsible for an increase in stiffness. Contrarily to expectations, cross-ply sandwiches have lower
flexural rigidity with respect to single-layer unidirectional sandwiches, and it was hypothesized that in the
latter case a better adhesion of the ply to the core was obtained. Interlaminar shear strength results seem
to provide a confirmation of our interpretation, together with the analytical calculations made following
Classical Lamination Theory results. Considering the mechanical behavior of the sandwiches at different
temperatures, it has been seen that temperatures as low as -20°C have a little effect on their stiffness,
while rising at 80°C polyamide matrixes become more ductile, being above glass transition temperature,
resulting in lower stiffness and higher strains.
Results from impact tests further showed the strengthening effect of applying skins over the short fiber
reinforced polyamide composite, being the first able to absorb more energy during the impact event and to
avoid a catastrophic failure. As could be expected, high temperatures alter positively the response to
impacts of the core, while at -20°C it becomes more brittle. On the contrary with respect to bending tests,
cross-ply impact testing at 80°C shows only a slight difference compared to room temperature, while low
temperature tests developed in a markedly higher rigidity of the material, thus lower impact resistance.
4
SOMMARIO
Nel corso di questo lavoro di tesi sono state studiate le proprietà meccaniche di un materiale composito
prodotto tramite stampaggio per iniezione, la cui matrice è una miscela di poliammidi e rinforzato con fibre
di vetro corte. Per fare ciò sono stati utilizzati test di flessione, come la flessione a tre punti e l’interlaminar
shear stress test, e di impatto. Le stesse prove sono state effettuate anche su strutture a sandwich in cui il
composito rinforzato a fibre corte era compreso tra film di poliammide, questa volta rinforzata con fibre di
vetro continue. Si sono quindi confrontati il composito a fibre corte, sandwich con un solo film con fibre
continue e sandwich con due film posti ortogonalmente l’uno con l’altro. Ci si è inoltre interessati
dell’effetto dell’aumento e della diminuzione di temperatura sulle proprietà meccaniche studiate.
L’osservazione al microscopio di sezioni sottili ricavate dal composito a fibre corte ha rivelato la presenza di
una struttura a tre strati dettata da una distribuzione delle fibre corte nello spesso re dei provini tipica di
manufatti prodotti tramite stampaggio ad iniezione. Questa struttura ha dimostrato di essere in grado di
influenzare il comportamento meccanico del materiale, specialmente ad alte temperature. L’applicazione
dei film con fibre continue si è rivelata efficace per l’ottenimento di una maggiore rigidezza. Tuttavia,
contrariamente alle aspettative, i sandwich con due film posti ortogonalmente presentano una minore
rigidezza a flessione rispetto a quelli con un solo film, per cui è stato ipotizzato che fosse dovuto
all’ottenimento di una minore adesione nel caso del sandwich con due film a fibre continue. Lo studio dello
sforzo interlaminare tramite i test sperimentali e tramite l’applicazione della Teoria della Laminazione
sembra confermare questa ipotesi. Per quanto riguarda l’effetto della temperatura, si è notato che la
diminuzione ad una temperatura di -20°C non comporta grandi cambiamenti, al contrario di quanto accade
ad alte temperature, specificamente ad 80°C. In questo caso, essendo al di sopra della temperatura di
transizione vetrosa, le matrici poliammidiche diventano più duttili, diminuendo così la rigidezza ed
aumentando la deformazione a rottura.
In linea con le prove a flessione, i risultati delle prove a impatto forniscono un’ulteriore conferma
dell’effetto di rinforzo dei film a fibra continua, poiché in grado di assorbire più energia durante l’evento di
impatto e di evitare la completa penetrazione e una rottura catastrofica. Elevate temperature influenzano
positivamente la risposta all’impatto dei materiali considerati, aumentando l’energia assorbita, mentre a -
20°C si ha un comportamento più fragile. Contrariamente ai test a flessione, in questo caso i sandwich a cui
sono stati applicati due strati mostrano comportamenti simili a temperatura ambiente ed a 80°C, mentre
l’effetto di diminuzione della resistenza all’impatto a basse temperature si dimostra essere più consistente.
(1)
5
1.– INTRODUCTION
A main concern in the automotive industry is the development of highly fuel efficient vehicles with low
emissions. In this field the materials traditionally used are metals, such as aluminum and steel, due to their
established use in most engineering industries and better understanding of application and manufacturing
process. As weight reduction is seen as the most efficient way to achieve fuel efficiency and lower CO2
emissions, lightweight polymers have been used to replace metal parts over the last few decades, from
interior to external body panels (1). An example of a current application is shown in Figure 1 (2). Fiber
reinforced plastic composites offer a range of distinct advantages over conventional materials, such as high
values of specific modulus and specific strength, superior corrosion resistance and improved fatigue
properties (3). Furthermore, thermoplastic matrix composites have enhanced toughness and an unlimited
shelf-life due to their intrinsic recyclability.
As a matter of facts, short-fiber reinforced thermoplastic composites are rapidly being recognized as a valid
substitution of metal parts for under-the-hood applications due to their potential for high-volume
processing combined with lower manufacturing costs. The more diffused fabrication methods are
extrusion, injection molding and compression molding, all already established processes for the production
of polymers. The first is a continuous process, the second an automated one and the latter is a semi-
automated process able to make large and complex parts not easily obtainable with the previous processes
(4). However, high performance levels can only be obtained from a composite part with high fiber
concentrations and if the reinforcing fibers in the final product have a sufficiently high fiber aspect ratio.
Indeed, short-fiber reinforced composites were developed to fill the property gap between continuous-
fiber reinforced and unreinforced polymers and are utilized into lightly loaded secondary structures or in
applications in which temperature causes the polymer alone to not maintain its structural integrity (5).
Nonetheless, structural parts of a vehicle as the chassis (Figure 2) can be manufactured by an hybrid
composite consisting of a short-fiber reinforced thermoplastic core with the application of unidirectional
continuous-fiber reinforced films. In this way the toughness of the discontinuous fiber-reinforced material
is combined with the outstanding strength of the films. An example is the seat back-rest shown in Figure 3.
Polyamides are a widely used class of materials as thermoplastic matrices for composites. They offer an
excellent resistance to mineral oils and a good balance of toughness and strength. Moreover, because of
their high melting point they maintain good mechanical properties up to temperatures in the region 120-
150°C (6). The most abundant polyamides used in the automotive industry are PA6 and PA66 for their
superior heat ageing resistance, good resistance to wearing and good surface finish.
The aim of this thesis work is then to compare the mechanical properties of an injection molded short glass
fiber-reinforced polyamide 66 blend with those of a composite sandwich made applying layers of
continuous glass fiber-reinforced PA6 and to analyze the adhesion between them.
6
Figure 1 - Thermoplastic engine oil pan fitted on ISF 3.8 Cummins Engine
Figure 2 - Structural parts of a vehicle
7
Figure 3 - Hybrid seat back-rest: injection molded short fiber-reinforced structure (right) and overmolding of continuous fiber-reinforced layer (left)
8
2. – BACKGROUND ON POLYAMIDES
Polyamides, also called nylons, were the first materials to be recognized as engineering thermoplastics.
These polymers consists of polyethylene segments (-CH2-)n separated by amide groups (-CONH-) that
provide hydrogen bonding between polymer chains, giving to the polymer a semi-crystalline nature.
Basically polyamides consists of a mixture of carboxylic acid and amine end groups, amide groups and free
or associated water that react by condensation. Chemically, they can be divided into two groups depending
on the type of repeating units involved in the synthesis (Figure 4):
- Those based on the reaction of amino acids or lactams, thus on monomers bearing both the
carboxylic acid and amine groups, yield polyamide-x types,
- Those based on the reaction between diamines and diacids yield polyamide-x,y types.
In this notation, x and y represent the number of carbon atoms between successive nitrogen atoms. Most
of the physical properties of polyamides of these two types are very similar, but synthesis and some
chemical properties are quite different. The most important commercial polyamides are two, one for each
type: polyamide 6 and polyamide 6,6. The first is obtained by hydrolytic ring opening of ε-caprolactam
followed by the condensation that lead to the build up of molecular weight (Figure 5). On the other hand,
polyamide 6,6 is synthesized through melt-condensation of hexamethylenediamine and adipic acid (Figure
6) (7). Apart from these, there exists a wide variety of polyamide products but few of them have find
interest for some specific application, for example nylon 10 and nylon 11 showing piezoelectric properties.
The monomeric content is the most important aspect of a given polyamide, as it largely determines the cost
of the polymer, depending on the availability and the cost of the monomer, the requirements of the
polymerization process and the properties of the final polymer through its internal structure. Thus
polyamides can be categorized depending on the types of monomers and their combinations (8):
- Polymers from exclusively straight-chain aliphatic monomers,
- Polymers incorporating at least one monomer that contains one or more cycloaliphatic rings,
- Polymers incorporating at least one monomer that contains an aromatic ring separated by
methylene groups from the carboxyl and/or amine end groups,
- Polymers incorporating at least one monomer (but not all) that contains an aromatic ring directly
connected to the carboxyl and/or amine end groups,
- Polymers from exclusively monomers that contain aromatic rings directly connected to the end
groups,
- Polymers incorporating at least one monomer that has one or more side substituent.
The interest of this thesis work will be focused on the first and third category.
Figure 4 - General structure of linear aliphatic polyamides: a) polyamide-x b) polyamide-x,y
9
Figure 5 - Hydrolitic ring opening of ε-caprolactam, followed by polycondensation between two end groups to give polyamide 6
Figure 6 - Polycondensation scheme of polyamide 6,6
2.1 – POLYAMIDES FROM STRAIGHT-CHAIN ALIPHATIC MONOMERS
Polymers built from only straight-chain monomers are the original and still dominant group, as they
comprehend nylon 6 and nylon 6,6. The principal structural difference between the various types of nylons
within this group is the length of the aliphatic chain segments separating adjacent amide groups. Due to the
fact that the polar amide groups give rise to high inter-chain attraction in the crystalline zones while the
aliphatic segments impart flexibility in the amorphous zones, these polymer show to be tough above their
glass transition temperature. Furthermore the response to temperature depends on this combination of
properties: in fact polyamides are known to have a melting point above 200°C due to the intermolecular
attraction, but it decreases as the length of the aliphatic chain increases (Table 1). In other words, as the
ratio of amide groups to ethylene groups in the chain increases, the melting point rises. This high
intermolecular attraction is responsible also for the high corrosion resistance of polyamides, as their
structure is highly cohesive. However a big drawback of the presence of the amide groups is their tendency
to attract water, and its absorption can induce dimensional and properties changes as it acts as a plasticizer
(9).
The crystallization of aliphatic polyamides is facilitated by the essentially linear character of the polymer
chain, the moderate molecular weight and the driving force for hydrogen bonding. Despite the fact that a
variety of unit cells are formed by the different polyamides, a regular pattern applies to all their crystalline
structure. Alignment of folded polymer chains in parallel arrays to form the maximum concentration of
hydrogen bonds results in monomolecular thick sheets, held together in the chain direction by the covalent
chain bonds and in the inter-chain direction by hydrogen bonds. The stacking of these sheets results in the
creation of crystallites, which grow in flat branched ribbons that form space-filling spherulites up to 50μm
in diameter. Usually the ribbon growth direction is the fast-growing hydrogen bond direction of the crystal.
These spherulites are held together by tie chains (10). Thus polyamides consists of many closely spaced,
high melting, rigid crystalline units held together by less regularly packed but hydrogen bonded amorphous
regions.
Inside each spherulite, the crystal structures observed for aliphatic polyamides fall in two main categories:
α-phase and γ-phase. In the first case the chains are in all-trans conformation, thus in a fully-extended zig
zag configuration, and the hydrogen bonds occur between anti-parallel chains in the same sheet (Figure 7,
Figure 8). In the second case instead the amide groups are rotated of about 60° around the chain axis and
hydrogen bonding is created between parallel chains among adjacent sheets (11). These phases however
have some differences between each polyamide. For example, in the α-phase of polyamide 6,6 a triclinic
unit cell can be individuated, where the planes are not rectangular but have a slight shear due to a
progressive offset between adjacent chains. These planes are also regularly offset, tilting of 42° relative to
10
the chain axis. The same phase of polyamide 6 instead have a monoclinic unit cell. The γ-phase instead
have a pseudo-hexagonal cell. Due to the fact that the α-phase have shorter CH2 bonds, it results to be the
more thermodynamically favored form and it is the most stable at room temperature. Increasing the
energy available to the system, polyamides are more prone to transform into the second less stable γ-form.
Thus α-phase can be easily produced by slow crystallization from the melt or by crystallization from a
solvent, while γ-phase by rapid crystallization. Noteworthy is to say that nylons show the occurrence of a
third non-equilibrium crystalline phase when heated through the Brill temperature: in the range 120°C –
180°C the α-phase reorganize into a new α’-phase where the triclinic or monoclinic structure is maintained
but with different lattice parameters.
Polyamide Tm (°C)
Nylon-6,6 265 Nylon-6,8 240
Nylon-6,10 225 Nylon-6,12 212
Nylon-6 230 Nylon-7 223
Nylon-11 188 Nylon-12 180
Table 1 - Melt temperatures of aliphatic polyamides
Figure 7 - Scheme of an hydrogen bonded sheet of PA6 with anti-parallel (a) and parallel (b) orientation of amide groups
11
Figure 8 - Structures of the α-phase of PA6 and of PA6,6. The left side shows the view of the hydrogen-bonding planes, and the right side shows the view along the chain axis
2.2 – POLYAMIDES FROM LINK-SEPARATED AROMATIC MONOMERS
Turning now to monomers containing aromatic rings, a class of polyamides based on these molecules has
been defined as araliphatics, in which the ring is separated on both ends by one or more methylene groups
from the amide linkage. If the separation consists of a single methylene unit, then the polymer is defined as
polyarylamide and is a member of the nylon family. As the name indicates they posses characteristics of
both aliphatics and aromatic groups. The importance of the separating unit is given by the fact that, if the
monomer is a diacid, it avoids the conjugation of the amide carbonyl double bond with the double bonds of
the ring. On the other hand the ring, with its bulkiness, have effects on both the crystalline and amorphous
region, changing the glass transition temperature, the melting point and other properties. It is in fact
known that the addition of an aromatic group affects most of all the Tg of a polymer, as it decreases the
chain flexibility and consequently increases the transition temperature. In addition the same effect is due
to the symmetry of the aromatic added groups. Following this reasoning also the melting temperature of a
polymer should increase, but this is not a linear trend as eutectics can occur, leading to a depressed melting
point (12). The aromatic ring can be contained in the diacid, in the diamine or in a xy-monomer. It has been
reported that the addition of a substituted diamine allow to obtain an higher glass transition temperature
as the moiety’s length is lower compared to the same diacid, leading to a better molecular fit and a higher
12
occurence of the ring along the chain. For this reason the araliphatic diamines are the more studied and
one in particular, meta-xylylenediamine (MXD) is the only commercialized member of the family. It is
polymerized with adipic acid to give polyamide m-MXD6. MXD is obtained from meta-xylylene, which is
obtained from coal tar or petroleum as part of a mixture of ortho-, meta- and para-isomers (Figure 9). Even
if PXD6 demonstrates to have higher thermal properties, the difficulties in extracting pure para-
xylylenediamine have led the interest only to MXD (8).
As in the case of aliphatic polyamides, MXD6 crystallize into spherulites composed of stacking of crystalline
sheets held together by hydrogen bonds. It is in fact believed that the unsymmetrical ring do not affect the
stacking of sheets but it interferes only with the polymeric sheet formation. In fact it is inferred that the
molecule is twisted from the planar structure by an angle of about 30° around the chain axis every six
carbon atoms, giving rise to an helical polymeric chain (Figure 10) (13).Furthermore, MXD6 is believed to
have a triclinic unit cell.
This polymer results to have an higher glass transition temperature than polyamide 6 and polyamide 6,6,
owing to the presence of the aromatic ring, but its melting temperature is lower (Tm = 243°C).
Figure 9 - Meta-xylylenediamineadipamide (MXD6) and para-xylylenediamineadipamide (PXD6)
Figure 10 - Molecular conformation of MXD6
13
3 - MATERIALS AND METHODS
This thesis is focused on the mechanical characterization of a composite consisting of a polyamide blend
matrix of polyamide 6,6 and MXD6 reinforced with 50% by weight of short glass fibers. The material was
supplied to us from Radici Group and commercially named as RADISTRONG A X15120 (14). The high fiber
content was a driving factor for the choice of blending polyamide 6,6 with MXD6 as the latter reduces the
viscosity of the polymer, facilitating the processing operations. Moreover it has the ability of increasing the
performances under tensile loads with respect to bare polyamide 6,6. This composite material has overall
high mechanical properties that makes it a valuable option for metal replacement. In addition, a composite
film of polyamide 6 and 60% by weight of unidirectional continuous glass fibers (CETEX TC910) purchased
from TenCate (15) was used as reinforcement to the Radistrong plates in the form of sandwiches. This film
product has an excellent combination of strength and impact resistance together with good resistance to
solvents, moldability and formability. Noteworthy is to say that all the materials were placed inside a
vacuum furnace at 80°C for two days before each test and treatment, in order to remove the presence of
humidity. In fact, polyamide is known to easily absorb water from the atmosphere and consequently all the
results later provided should be considered for materials in the dry-as-molded state.
3.1 – METHODS OF PROCESS
3.1.1 – CORE PROCESSING
Radistrong was supplied to us in the form of injection molded plates of dimensions 180x120mm. Injection
molding is widely used in manufacturing plastics as it allows producing parts even with a complex shape
whit a high dimensional accuracy at an effective cost and relatively short cycle time. This process consists in
feeding a barrel heated at around 280-300°C (Figure 11) with pellets previously obtained by mixing the
fibers together with the base resin in the melt phase. These pellets, which due to the temperature higher
than the melting point of the matrix forms a melt containing the fibers, are continuously mixed and
transferred down the barrel by a rotating screw located in it. When the melt reaches the end of the barrel it
is injected under high pressure through a gate into a closed metal mold heated at 90-120°C and machined
to desired finished part specification. Great importance in this process obviously has the viscosity of the
melt. It is known that the movement of the melt down the barrel and into the mold due to the high
pressures subjects the polymeric molecules to high shear forces. Thermoplastic polymers show an
increasing pseudo-plastic behavior as the shear rates increases and consequently their apparent viscosities
decrease. The addition of fibers has a big effect on the viscosity of the melt, as it is raised. However it
appears to follow the same trend as the neat resin as it decreases for higher strain rates. Interestingly, at
very high strain rates (104-105 s-1) the difference in apparent viscosity between the base thermoplastic resin
and the reinforced one becomes much smaller meaning that they behave most likely at the same way (5).
An important drawback of such kind of process is the fiber breakage that occurs due to the high shear rates
generated by the action of the screw in the barrel, but also to the presence of narrow openings and the
previous pelletizing step. This factor inherently limits the length of the fibers and as a consequence the
mechanical properties of the final composite. Moreover, they also depend on the orientation distribution
of the fibers in the injection molded part. In fact the fiber orientation is governed by the melt flow inside
the mold and it greatly varies across the thickness. As it is nearly impossible to predict it a priori, numerical
simulations could be done in order to optimize the processing parameters and maximize the mechanical
properties of these short fiber composites.
14
Figure 11 - Schematic diagram of a screw injection molding machine
3.1.2 – SANDWICH PROCESSING
Three types of sandwiches were fabricated:
- Unidirectional sandwiches, consisting of the Radistrong core embedded between two Cetex skins
with the continuous fibers directed along the flow direction (0°/core/0°, Figure 12a);
- Cross-ply sandwiches, consisting of the Radistrong core and two continuous fibers skins over each
face, with two different stacking sequences:
- In one case the outer skins have the fibers oriented as the mold flow direction of the core, while
the fibers of the inner ones are placed transversally (0°/90°/core/90°/0°, Figure 12b);
- In the other case the outer skins are placed with the fibers perpendicular to the mold flow direction
and those of the inner skins parallel to it (90°/0°/core/0°/90°, Figure 12c).
Noteworthy is that all the configurations are symmetric across the thickness with respect to the midplane
cutting the core in half. In accordance with a previous work (16), these layered composites are made with
the use of an aluminum mold consisting of two halves: as it can be seen in Figure 13, one half has a
rectangular hole measuring 175 x 115 mm and eight screws which guides the male half in it. Different steps
were rigorously followed in the manufacturing of the sandwiches. First of all, the mold was cleaned with
water and air sprayed to remove any residual of previous operations and dust particles, then it was sprayed
with PTFE in order to avoid the sticking of the melted polymeric matrix inside the mold. Additionally, films
of polyimide were deposed on the planar faces and on the edges of the mold so that the possibility that the
polyamide would attach to the mold is further decreased. Subsequently the Cetex skins and the Radistrong
core were deposited inside the mold, which then was closed and placed inside an oven. The male part
weights about 5 kg and additional weights were added, thus resulting in an applied pressure of 5.4 kPa,
which was sufficient to ensure a proper consolidation of the composite. The oven was then heated up until
the mold reaches 253°C, temperature that is maintained for 15 min before a natural cooling. This
temperature was chosen as it is higher than the melting temperature of the polyamide 6 film but slightly
lower than that of the polyamide 6,6 core: in this way it was possible to achieve a good adhesion melting
the outer skins, that were pressed against the softened core.
15
Figure 12 - Stacking sequences: a) unidirectional, b) cross-ply 0/90, c) cross-ply 90/0. The arrows indicate the direction of injection
Figure 13 - Mold parts used for the production of the sandwiches
3.2 – DIFFERENTIAL SCANNING CALORIMETRY
Differential scanning calorimetry is a thermal analysis technique that is able to investigate thermal
transitions in materials with respect to temperature through the measurement of heat flow to and from a
specimen within a temperature-controlled environment. Specifically a Mettler Toledo DSC 822 was utilized.
It is a power compensation DSC, consisting of a furnace containing a reference crucible and a second
crucible in which the sample to be analyzed is placed. The data are acquired by measuring the heat flow
needed to maintain the temperature of the sample equal to that of the reference. The temperature is
detected through area thermocouples and the furnace is uniformly heated or cooled by nitrogen as purge
gas in a range between -20°C and 300°C. The device has thus two electrical control systems: one able to
maintain the programmed specimen temperature and one able to equalize the temperatures of the two
crucibles. The latter thus acts as recorder of the heat flow curve against temperature. This method is able
to detect a variety of parameters, including the glass transition temperature, melting point, crystallization,
polymerization oxidation and degradation (17).
In this thesis differential scanning calorimetry was performed on the short fiber reinforced polyamide 6,6
core, mainly to detect its melting temperature so that the sandwich processing could be set up. A specimen
of 11,2 mg was obtained and it was subjected to an heating cycle from room temperature up to 300°C, then
cooled down to -20°C and subsequently heated again to 300°C with a rate of 10°C/min. From the data
gathered the glass transition temperatures and the melting temperatures were individuated. Moreover,
the enthalpies of fusion could be calculated as an integral of the area under the melting peaks.
16
3.3 – OPTICAL MISCROSCOPY
The fiber orientation distribution of the short fibers throughout the plates was analyzed. Samples were cut
out from the as-received plates and subsequently incorporated into a transparent resin and lapped with
abrasive papers of progressively lower roughness until a thickness of approximately 100μm was reached.
These thin sections were then analyzed with the use of the optical microscope in transmission. The samples
were taken from different locations across the plate in order to investigate the differences that exists in the
orientation and dimensions of the short fibers. Six samples were extracted parallel to the injection flow
direction along the edges of the plate while three were perpendicular and at the center of the plate(Figure
14). Next, measurements were made with the use of a software of image analysis, ImageJ.
Figure 14 - Positions of the sections along the injection molded plate. The red arrow indicates the direction of injection
17
3.4 – MECHANICAL TESTS
In this section the tests performed with the aim of assessing the mechanical properties of the considered
materials are illustrated. Flexural tests were conducted using the dynamometer Instron 1185 R 5800 which
electronically records the flexural modulus and the displacement of the specimen. Impact tests were
performed with a drop weight tower Ceast Fractovis. The machine gives us a force-time diagram,
constructed by the use of a cell monitoring the load and a photo cell measuring time. The striker has an
hemispherical tip and a clamping system prevents it to rebound on the specimen after the first strike.
3.4.1 – THREE POINT BENDING TEST
The three point bending test was chosen as the method to evaluate the strength and modulus of our
materials. Flexural tests are widely used owing to the relative simplicity of test method, instrumentation
and sample preparation required. As can be seen in the (Figure 15), the three point bending test is
performed placing a flat rectangular sample over two supporting pins and loading it at the center: with this
configuration a tensile stress arises at the face in contact with the supporting pins, as the beam is deflecting
downward, while a compressive stress can be found on the upper face. Thus the specimen is subjected to a
combination of non uniform stresses, as we have maximum tension at the lower surface, maximum
compression at the upper surface and maximum shear at the mid-plane . Elsewhere the beam is in a
combined stress state of normal and shear stress and the samples can fail due to the single contribution or
the combination of the described stresses. To ensure that the primary failure comes from the compressive
or tensile stress the shear one must be minimized, and this can be done by carefully controlling the
geometry of the specimen, in particular the span to depth ratio L/h, that is the ratio of the distance
between the supporting pins and the thickness of the sample. Knowing that the shear stress is independent
on the specimen length, as it is uniform along the length of the beam, while the bending moment is directly
proportional to it as it linearly increases from zero at the supports to the maximum at the centre, large L/h
ratios (higher than 16) have to be chosen so that specimens are unlikely to fail in shear (18). In this thesis
three point bending tests are performed according to the standard ISO 14125 (19). Table 2 reports the
dimensions of the specimens tested. Their calculation was made considering the bare Radistrong (core) and
the sandwiches (one 0° skin per face, UD; two 0°/90° skins per face, CP) as class I and class III materials
respectively: in the two cases the span to thickness ratio are L/h = 16 and 20. The test speed was set to 1
mm/min for all the tests. In order to note if the layers of fibers inside the core material have some effect on
the mechanical properties, specimens were cut along and transversally to the mold flow direction.
Figure 15 – Three Point Bending test configuration
18
As a first approximation, we assume the material properties to be uniform throughout the thickness of the
specimens. However this hypothesis is not appropriate as the materials considered are characterized by an
high degree of anisotropy, both the short fiber reinforced core and especially the continuous fiber
reinforced tapes. For this reason, a model able to consider the directionality dependency will be applied
later on.
Homogeneous beam theory can be used to calculate the flexural stress and strain:
�� =3��
2�ℎ� ; �� =
6�ℎ
��
in which:
�� is the flexural stress (MPa),
F is the recorded load (N),
L is the span (mm),
b is the width of the specimen (mm),
h is the thickness of the specimen (mm),
�� is the flexural strain in the outer surface of the specimen,
s is the recorded beam mid-point deflection (mm).
From this equation the flexural strength can be obtained considering the maximum load until fracture
sustained by the sample. The flexural modulus is instead evaluated as the slope of the curve fitting the
flexural stress-flexural strain curve at the beginning of the test.
h (mm) l (mm) L (mm) b (mm)
Core 3.4 70 54 10 UD 3.8 132 80 15 CP 4.2 132 86 15
Table 2 - Dimensions of Three Point Bending tests specimens
3.4.2 – INTERLAMINAR SHEAR STRENGTH TEST
ILSS test is a flexure test performed again with a three points loading configuration (Figure 16) but in an
opposite situation with respect to three point bending test. In fact in this case the highest probability that
specimens fail under the action of shear stress is wanted and thus again the span to thickness ratio is
considered. As L/h is reduced, the importance of shear stresses inside the sample increases at the expenses
of normal stresses. As said before, according to classical beam theory the shear stress distribution in short
beams loaded in a three point flexure configuration has a parabolic shape through the specimen thickness.
The stress is maximum at the mid-plane and zero at the upper and lower surfaces. However this test can
only provide an apparent shear stress. In fact a very complex stress state is induced in the short beam as
contact points have an high influence: finite element analysis have demonstrated that the through
thickness shear stress distribution is skewed near the load and reaction points and that the maximum
interlaminar shear stress is positioned between the mid-plane and the upper zone, close to the loading
zone (18). However, as three point bending, this test is widespread thanks to its ease of specimen
preparation and testing equipments needed.
19
In this thesis interlaminar shear stress tests were performed following the directives of the standard ISO
14130 (20). Table 3 summarizes the specimens’ dimensions. Noteworthy is to say that in this case the span
to thickness ratio is L/h = 5. It must be specified that only sandwiches were tested as this method was used
to identify the goodness of the attachment of the unidirectional films on the core layers. As defined by the
normative the speed of testing was set to 1 mm/min. from the recorded load versus displacement values
the apparent interlaminar shear stress is calculated as:
� =3�
4�ℎ
in which:
F is the load recorded (N),
b is the width of the specimen (mm),
h is the thickness of the specimen (mm).
Figure 16 - Interlaminar Shear Strength test configuration
h (mm) l (mm) L (mm) b (mm)
UD 3.8 38 19 19 CP 4.2 42 21 21
Table 3- Dimensions of Interlaminar Shear Stress tests specimens
3.4.3 – IMPACT TEST
Generally, impact tests can be categorized into low or high velocity. Low velocity impact events can be
treated as quasi static and concern speeds between one to tens m/s. The structure has the time to respond
to the impact as the contact duration is long enough and in consequence more energy is absorbed
elastically. On the other hand, high velocity impact response is dominated by stress wave propagation
through the material, in which the structure does not have time to respond leading to very localized
damage (21). These kind of tests offer a measure of the actual energy required to break the specimen from
the recorded force vs time diagram.
In this thesis a drop weight impact test was chosen, which enables to measure the impact behavior of a
structure for out of plane impact loads. It is a relatively simple impact strength test which is carried out
20
using a machine in which a weight carrier is able to follow rails down in a tower and with an impacting tip
to ensure that we have a punctual impact location. The height and the mass of the striker can be altered
and the minimum value of the product of falling height and mass which causes fracture is the impact
strength of the material. As reported in reference (22) to carry out the test a trial run is undertaken first.
The weight and height of the striker are set so that the energy measured is equal to the expected impact
strength. If the result is a broken specimen, a second one is tested with a lower impact energy and so on
until a specimen does not break. Then the remaining specimens will be tested with the minimum impact
energy required to break the specimens founded. On the contrary, if after the trial run the sample does not
break, the impact energy is raised.
The falling weight test was used to evaluate impact properties of the Radistrong core and of the cross-ply
sandwiches. The standards taken as reference are ISO 6603-1 (23)and ISO 6603-2 (24). As defined in these
standards, the specimens were squares of approximately 60x60mm for both the materials tested. A
constant falling mass procedure was used, keeping the mass of the striker fixed at 4.158kg and varying its
height in order to find the minimum impact energy needed to break the specimens.
As already said Ceast Fractovis records the puncture force as a function of time. In order to evaluate the
impact energy, the deflection of the test specimen has to be calculated from the force-time trace using the
following equation:
�(�) = ��� −1
��� �� �(�)���
��
�
�
�
�
�� +1
2���
where
�� is the impact velocity (m/s),
t is the time after impact at which the deflection is to be calculated (s),
F(t) is the force measured at any time after the impact (N),
l(t) is the deflection, more precisely the relative displacement between the striker and the
specimen support (m),
�� is the falling mass of the energy carrier (kg),
� is the acceleration due to gravity (m/s2).
Once the force and deflection are known for identical times during impact, the energy can be calculated by
integrating the area under the force-deflection curve up to a specific time t1:
�� = � �(�)��
��
�
in which
F(l) is the force measured at any deflection l (N),
l is the deflection (m),
E is the impact energy (J).
Usually brittle composites as our materials absorb energy by damage mechanisms such as fiber debonding,
fiber pull out, fiber breakage and matrix cracking.
21
3.4.4 – CLASSICAL LAMINATION THEORY
It is well known that composite materials, especially laminates, are anisotropic in nature. Even short fibers
reinforced plastics, that have an high tendency in being isotropic if they have a random distribution of
fibers, show the occurrence of the already mentioned three-layer configuration if processed through
injection molding and so they can be considered as laminated materials. On top of these consideration it
becomes straightforward that the methods utilized by the standards are not applicable to our materials as
they refer to homogeneous and isotropic ones. In order to overcome this issue classical lamination theory
has been used as a tool to calculate the stress distribution throughout the specimens tested.
The theory (25) is based on the fact that laminated composites are stacks of fiber reinforced plies and
therefore at first the mechanical properties of the basic building blocks have to be defined. In formulating
the constitutive equations of a unidirectional lamina it is assumed that every lamina is a continuum and
behaves as a linear elastic material. From the first statement, with classical lamination theory we can
investigate the macro-mechanical behavior of a lamina without a micro-mechanical approach able to
consider events such as fiber breakage or fiber-matrix debonding. At this level the properties of a
composite are derived from a weighted average of those of the constituent materials, consequently a
lamina is assumed as homogeneous despite its intrinsic heterogeneous characteristics. On the other hand
the second assumption implies that the generalized Hooke’s law is valid and for an anisotropic material it
states:
�� = �����
where the tensors �� and �� contain the stress and strain components and the ��� are the material’s
stiffness coefficients, all referred to an orthogonal Cartesian coordinate system. By means of the
assumption that the material is hyperelastic the stiffness matrix is symmetric (��� = ���) and thus for the
most general elastic material 21 independent coefficients must be known to define its stress state.
However composite plies reinforced with unidirectional fibers are orthotropic, as the coordinate system
identify three axes of symmetry (Figure 22a) . Moreover most laminates are typically thin and experience a
plane state of stress: the transverse stress components are very small in comparison to those in the other
directions and can be neglected (��� = 0). As a result the constitutive equation can be reduced to:
�
��
��
��
�
�
= ����
���
0
���
���
0
00
���
�
�
�
��
��
��
�
�
in which ���� are the components of the reduced stiffness matrix referred to the k-th layer and are related
to the engineering constants as follows:
���� =
���
1 − ���� ���
�
���� =
���� ��
�
1 − ���� ���
�
���� =
���
1 − ���� ���
�
���� = ���
�
Figure 17- Left: Lamina in plane state of stress. Right: variation of strains through layer and laminate thickness typical instress (a) and corresponding stresses (b)
It can be noted that the reduced stiffnesses inv
of the laminae in direction �� and �
coefficients ���� and ���
� . The stiffness coeffici
single one, with one axis, usually ��
with different orientations to obtain the desired mechanical properties. Thus the fa
equation can be transformed in such a way they are all relative to a common coordinate system:
where:
Ǭ�� = ������
Ǭ�� = ������
Ǭ�� = (��� +
Ǭ�� = (��� − ���
Ǭ�� = (��� − ���
Ǭ�� = (��� + ���
in which � is the angle between the fibers in each lamina
the principle of virtual works it can be develope
and forces acting on the laminate to its strains. Although strains are continuous through the thickness,
stresses are not due to the changes in material coefficients (
calculated integrating the stresses through the laminate thickness with a lamina
Left: Lamina in plane state of stress. Right: variation of strains through layer and laminate thickness typical in
It can be noted that the reduced stiffnesses involve 4 independent material’s constants: the Young’s moduli
�� , respectively ��� and ��
�, the shear modulus ���
. The stiffness coefficients of each lamina refer to the coordinate system of the
, aligned to the fiber orientation. Usually laminae are stacked together
with different orientations to obtain the desired mechanical properties. Thus the factors of the constitutive
equation can be transformed in such a way they are all relative to a common coordinate system:
�
���
���
���
�
�
= �
Ǭ��
Ǭ��
Ǭ��
Ǭ��
Ǭ��
Ǭ��
Ǭ��
Ǭ��
Ǭ��
�
�
�
���
���
���
�
�
����� + �������� + 2(���+2���)����������
����� + �������� + 2(���+2���)����������
+ ���−4���)���������� + ���(����� + �����
��−2���)��������� + (��� − ���+2���)�������
��−2���)��������� + (��� − ���+2���)���������
���−2���−2���)���������� + ���(����� + ���
is the angle between the fibers in each lamina �� and the principal laminate direction
the principle of virtual works it can be developed a constitutive law relating the resultants of the moments
and forces acting on the laminate to its strains. Although strains are continuous through the thickness,
stresses are not due to the changes in material coefficients (Figure 17b) and hence the resultants can be
calculated integrating the stresses through the laminate thickness with a lamina-wise integration.
22
Left: Lamina in plane state of stress. Right: variation of strains through layer and laminate thickness typical in-plane
olve 4 independent material’s constants: the Young’s moduli
��� and the Poisson’s
ents of each lamina refer to the coordinate system of the
, aligned to the fiber orientation. Usually laminae are stacked together
ctors of the constitutive
equation can be transformed in such a way they are all relative to a common coordinate system:
�)
���������
�����
�����)
and the principal laminate direction �. From
d a constitutive law relating the resultants of the moments
and forces acting on the laminate to its strains. Although strains are continuous through the thickness,
) and hence the resultants can be
wise integration.
23
�
���
���
���
� = � � �
���
���
���
� ��
����
��
�
���
; �
���
���
���
� = � � �
���
���
���
� ���
����
��
�
���
Expanding the results we obtain:
�
���
���
���
� = �
���
���
���
���
���
���
���
���
���
� �
���
���
���
� + ����
���
���
���
���
���
���
���
���
� �
���
���
���
�
�
���
���
���
� = ����
���
���
���
���
���
���
���
���
� �
���
���
���
� + ����
���
���
���
���
���
���
���
���
� �
���
���
���
�
where ���, ��� and ��� are called extensional, bending and bending-extensional coupling stiffnesses,
respectively. They are coefficients of 3x3 symmetric matrices and are defined as:
��� = � Ǭ���
�
���
(���� − ��)
��� =1
2� Ǭ��
�
�
���
������ − ��
��
��� =1
3� Ǭ��
�
�
���
(����� − ��
�)
On the other hand [�] and [�] are the vectors of the membrane and bending strains (Figure 18). It can be
seen that the stacking sequence and the material properties of each lamina are embedded in the
coefficients of the stiffness matrices, from which they rule the stress-strain relationship of the laminate.
In order to investigate our specific case, some considerations have to be made. First of all, the Radistrong
core, tested by itself or as part of the sandwiches, is not a continuous fiber reinforced laminate as the
materials considered in classical lamination theory, but it is made of a matrix containing chopped, short
fibers. However, as previously reported, three different layers appear though its thickness due to the flow
of the melt during injection molding. Inside each layer, the fibers have a preferential orientation: in the
outer layers they tend to align parallel to the flow direction while in the central layer transversally to it. This
configuration can be approximately view as a laminate composed of four layers symmetrically distributed
with respect to the midline: each of these layers is considered as a continuous fiber-reinforced ply so that
the core can be studied using the theory introduced (Figure 19). This introduces an important simplification
as due to midplane symmetry there is no coupling between extension and bending so ��� = 0. If we now
place the principal coordinate system of the composite making the flow direction coincide with the x axis,
we note that the core can be viewed as a balanced laminated composite constituted by four plies in which
fibers are directed at 0° and 90° with respect to the flow direction. Placing the continuous fiber reinforced
PA6 film over both the faces of the laminate we obtain our sandwiches: thus in all our configurations we
deal only with plies containing fibers oriented only at 0° or 90°. As a consequence our laminates are all
orthotropic so the components of the transformed stiffness matrix along the transverse direction Ǭ�� and
Ǭ�� are zero:
24
�Ǭ��
Ǭ��
0
Ǭ��
Ǭ��
0
00
Ǭ��
�
�
Figure 18 - Basic deformations of the layer: (a) in-plane tension and compression; (b) in-plane shear; (c) bending; (d) twisting
Figure 19 - Preferential fibers' orientations in the injection molded core
The last consideration concerns the specimens of the test methods. Considering that their width b is small
with respect to the length l and that the loading condition makes displacements function of the x-
coordinate only, they can be treated as beams. In order to analyze the stress distribution in our specimens,
for sake of simplicity a pure bending condition has been considered. Due to the fact that for symmetrically
laminated beams the equations for bending deflection and those for stretching displacements are
uncoupled, we can consider just the firsts. Moreover, if the in plane forces are assumed to be zero, the in
plane displacements are zero and the problem is reduced to the solution for bending deflection and
stresses. From the classical laminated plate theory, the constitutive equations, in the absence of in-plane
forces, are given by:
�
���
���
���
� = − ����
���
0
���
���
0
00
���
� �
���
���
���
�
where:
�
���
���
���
� =
⎣⎢⎢⎢⎢⎢⎡
����
���
����
���
2����
����⎦⎥⎥⎥⎥⎥⎤
25
Inverting the equation and considering that ��� = ��� = 0, we find:
����
���= −���
∗ ��� ; ����
���= −���
∗ ���
in which ���∗ are the elements of the inverse ��� matrix. From the relations above it can be noted that the
deflection �� cannot be independent of the coordinate y due to Poisson effect (���∗ ). Again, these effects
can be neglected only for length-to-width ratio high enough. Assuming that this is our case the transverse
deflection will be treated as a function of x and time, �� = ��(�, �). Introducing the flexural modulus ����
we can recall the Euler-Bernoulli equation from classical beam theory (Figure 20):
���� =
12
ℎ����∗ =
�
������∗ ; ��� =
�ℎ�
12
����
���= −
����
�������
where � and ℎ are the width and the total thickness of the beam, while ��� is the moment of inertia. Our
tests can be depicted as a simply supported beam with a center point load �� where the deflection is
symmetric about the center point of the beam and the bending moment is:
��� =���
2
The deflection of the beam can be calculated by direct integration of the differential equation above.
Imposing the boundary conditions in order to define the integration constants:
��(0) = 0 ; ���
���
�
2� = 0
we obtain:
��(�) = −���
������� �
��
12−
���
16�
where � is the span between the supporting pins. The maximum deflection occurs at midspan and it is:
�� ��
2� = ���� =
�����
48�������
The in-plane stresses in the k-th layer can be computed by the transformed constitutive equation, arriving
to the results:
���� = �����Ǭ��
� ���∗ + Ǭ��
� ���∗ �
���� = �����Ǭ��
� ���∗ + Ǭ��
� ���∗ �
���� = �����Ǭ��
� ���∗ �
In this theory we computed the constitutive equations considering the interlaminar stresses ��� and ���
null. However in reality these stresses exist and may be responsible for failures in composite laminates due
26
to the low shear and transverse normal strength. Considering the equilibrium equations for tridimensional
elasticity:
0 =����
��+
����
��+
����
��
0 =����
��+
����
��+
����
��
0 =����
��+
����
��+
����
��
and remembering that all variables are independent from y, the interlaminar stresses are calculated
integrating these equations with respect to z for each layer. Focusing on the stresses between our plies:
���� (�, �) = −���Ǭ��
� ���∗ + Ǭ��
� ���∗ � �
�� − ���
2� + ��
where �� is the integration constant, again evaluated using boundary conditions:
���� (�, ��) = 0 , � = 1
and interface continuity conditions between each ply (Figure 21):
���� (�, ����) = ���
���(�, ����)
Using the composite beam theory is then possible to analyze the stress distribution across the thickness of
our laminates, identifying the layers able to withstand more stress.
Figure 20 - Sign convention
27
Figure 21 - Equilibrium of interlaminar stresses in a laminated beam
28
4. - RESULTS
4.1 – DSC ANALYSIS
Graph 1 shows the result of the Differential Scanning Calorimetry test performed on Radistrong. Table 4
reports the enthalpies of fusion and crystallization for the three thermal cycles.
The first heating ramp from -20°C to 300°C shows a broad endotherm between 20°C and 100°C and a single
melting peak at 261°C, preceded by a small exothermic peak. While the source of the first endotherm,
which is superimposed to the glass transition temperature (in the range50-80°C for polyamides), is of
doubtful interpretation, the exothermic peak at 235°C can be attributed to a cold crystallization process.
This phenomenon consists in the re-organization of bad formed crystalline structures, quenched in during
the processing, thanks to the increased mobility of the polymeric chains at high temperatures. However,
due to the fact that the first heating cycle is seen to be highly dependent on the previous processing
operations, the only information considered was the melting point in order to set up the temperature at
which sandwiches would be molded. This data, together with the known melting temperature of the
Graph 1 - DSC analysis of Radistrong
29
Tg (°C) Tm (°C) Tc (°C) Enthalpy (J/g)
First heating cycle 77 261 - 10.4
Cooling cycle - - 225 7.7
Second heating cycle 60 247, 257 - 9.4
Table 4 - Data from DSC analysis
reinforcing PA6 film, brought us to the result that a temperature around 250°C was a good choice. In fact at
this temperature the films melt while the core just soften, allowing us to achieve a good attachment and at
the same time, as the core do not melt, avoid the fibers to move. As a matter of facts, the stability of the
fibers was more difficult to obtain in the cross ply sandwiches because the presence of two layers of
unidirectional film resulted in a less stable substrate for the outer one.
During cooling a single crystallization peak appears denoting that the distribution of crystallites sizes is
monomodal (26). The lower enthalpy of crystallization with respect to those of fusion can be related to the
fact that some portions of the polymeric blend do not crystallize back due to the presence of an high
fraction of fibers that may restrain the crystal growth (27). In fact, it is known that fibers can act as
nucleating sites, but as the temperature is lowered and the density of the crystalline nuclei increases, they
begin to impinge one another preventing a thorough crystal growth. This is in accordance with the narrow
distribution of crystallites sizes cited before.
Being the first two steps of the DSC analysis performed to erase any previous modifications underwent in
our material, the second heating cycle can give more information on its effective micro-structural
characteristics. The glass transition temperature is shown at around 70°C and is now more evident,
supporting the idea that the endotherm recorded during the previous heating step was influenced by the
processing of the material. After the Tg a broad and slightly marked endotherm appears until approximately
204°C are reached, after which two melting peaks are clearly visible at 247°C and 257°C. Literature reveals
to be in disagreement on the interpretation of DSC thermographs of polyamide 6,6. On one hand, the two
endothermic peaks are associated to the melting of two different phases of the crystalline structure of
polyamide 6,6: the lower melting temperature corresponds to the pseudo-hexagonal γ phase while the
higher one to the more stable triclinic α phase (27), (28). Contrarily, other researchers believes that the
double melting peak is an evidence of a melt-re-crystallization process before which a Brill transition,
associated to the broad endotherm, occurs (29), (30), (31). The Brill transition is believed to be the
transformation during heating from the more stable α phase to the less stable form γ. Supported from XRD
analysis performed simultaneously with DSC, which demonstrates that two reflecting peaks appearing at
low temperatures and corresponding to the spacings of a triclinic structure converge into one at the Brill
temperature, and from FTIR studies showing that the absorption band corresponding to the α phase
weakens at 100°C and disappears before melting, they concluded that the two melting peaks shows the
behavior of the γ phase only. For these reasons it is believed that upon heating above the Brill transition we
have first of all the melting of regular γ crystallites, while ill-formed ones undergo a re-crystallization and
subsequently melts at a slightly higher temperature. However, this re-crystallization phenomenon should
be accompanied by an exothermic peak, not evident here.
Recent studies (32) have demonstrated, focusing on the behavior of hydrogen bonds along the polyamide
chains with the use of H-NMR, that in reality the Brill transition is not a transformation between a triclinic
and a pseudo-hexagonal crystal structure, as no 60° flip-flop occurs. On the contrary, the α phase
30
transforms into another triclinic phase with different lattice parameters, conserving the hydrogen bonded
sheets. As a consequence, the two melting peaks correspond to the two crystalline phases, γ and the new
α’.
MXD6 does not have an high influence on the thermal behavior of the polyamide blend considered and the
possible reasons are two. First of all, it could be that its amount in the blend composition is so low that its
behavior is overwritten by that of polyamide 6,6. In fact, while the glass transition temperature lies more or
less in the same range as that of PA6,6, the melting peak could be hidden by the tail of the Brill transition.
On the other hand it has been demonstrated that, depending on the blending method, the MXD6 aromatic
moieties, if present in a low amount, do not co-crystallize with polyamide 6,6 but they are mostly
incorporated into the amorphous phase, explaining the fact that no other peaks other than those of
polyamide 6,6 are recorded.
4.2 – MICROSTRUCTURE OBSERVATIONS
Figure 23 shows extracts taken from the micrographs of the thin sections cut, while Table 5 lists the
average lengths of the fibers in the different slices. It can be noted that the length of the fibers has an high
standard deviation, as their length is determined by the processing. However no differences in the fibers’
diameters can be seen throughout the plate, assessing to an average value of 10μm. Regarding the
orientation of the fibers, the three-layer configuration described in literature is easily identified in sections
A, B and C, with two outer slabs made of fibers mainly oriented along the injection flow and a core slab of
fibers in the direction perpendicular to the injection flow. The thicknesses of the layers are reported in
Table 6: their mean values are 913μm and 1251μm concerning the outer layers and the core one,
respectively. It can be thus noted that the inner layer corresponds to a slightly lower value than half the
thickness of the plate. On the other hand, the inspection of the thin sections cut out parallel to the injection
flow along the edges of the plate reveals that there is no formation of the three layers. This is not
unexpected as in the outermost regions of the plate the melt coming out from the injection molding
machine encounter the walls of the mold at a lower temperature and solidifies rather rapidly, constraining
the glass fibers as can be seen in Figure 22 . Thus it can be inferred from the micrographs that the skin/core
proportion and fiber alignment are maintained uniform throughout the plate, being disturbed only by the
edge.
Figure 22 - Schematic rapresentation of the fiber orientation in the core of an injection molded square plate
31
Average fiber length (μm) ±Er (μm)
Section 1 301 9.4 Section 2 326 55 Section 3 329 28 Section 4 288 27 Section 5 321 33 Section 6 261 4.4
Section A 264 63 Section B 219 28 Section C 270 62
Table 5- Fiber length distribution
Figure 23 - Thin section micrographs: a) Section 1, b) Section 2, c) Section 3, d) Section 4, e) Section 5, f) Section 6,
g) Section A, h) Section B, i) Section C
Layer 1(μm) ±Er (μm) Layer 2(μm) ±Er (μm) Layer 3(μm) ±Er (μm)
Section A 813 39 1378 56 770 38 Section B 1042 15 1155 35 1122 88 Section C 982 15 1219 13 748 79
Table 6 - Mean thickness of the core layers (Layer 1 = upper layer, Layer 2 = central layer, Layer 3 = bottom layer)
32
4.3 – FLEXURAL TESTS
4.3.1 – THREE POINT BENDING
Concerning the Radistrong core, differences are recorded between samples cut along the flow direction (L)
and transversally (T). It can be readily viewed that L-specimens have an higher rigidity, meaning that the
three-layer distribution of the fibers across the thickness of the injection molded core has an high influence
on the mechanical response of the material. This behavior can be explained considering the orientation of
the short fibers with respect to the loading plane. It is known that, if a fiber reinforced single-layer beam is
subjected to a deflection, the highest resistance is met if the fibers are placed longitudinally along its length
while a minimum can be found if the direction of the fibers form an angle of 90° with the loading plane.
This is due to the fact that while in the first case the load is distributed along the fibers, in the latter the
most load-bearing component is the matrix and the fiber-matrix interface, which are the weakest part of
the composite. If we now consider a laminated beam, the stacking sequence has a great importance as
higher is the distance of the 0°-plies from the neutral axis, higher is the flexural rigidity. In this way we can
exactly depict our situation: in the L-specimens, the layers in which the preferential orientation of the short
fibers is 0°, thus along the flow direction, are the outer one. On the contrary, in the T-specimens we see the
opposite situation, as the fibers oriented at 0° lies in the inner layers and the outer ones show a 90°
preferential orientation. As shown in Graph 2, this results in a maximum strength of about 345MPa
±11.9MPa for the L-samples and 177MPa ±3.5MPa for T-samples, corresponding to average maximum
strains of 3% and 5%, respectively. It can be thus seen that there is a decrease of 95% in strength and an
increase of 60% in strain from L- to T-specimens. Following the same pattern, the flexural modulus of L-
specimens is 14.4GPa ±0.6GPa and that of T-specimens is 6.7GPa ±0.08GPa. From the graph it can be also
inferred that the samples failed in a brittle manner, independently on fiber direction.
Having determined that core-L specimens have the best performance in terms of stiffness among the bare
Radistrong, we compare these results with those of the unidirectional sandwich (UD)(Graph 3). Following
the same reasoning as before, it is straightforward that the addition of two 0°-laminae on the surfaces of
the core plates increases the stiffness of our specimens. In fact the average maximum stress for UD samples
is 528MPa ±17.2MPa, that means a rising of about 53%, while the average flexural modulus is 20.5GPa
±0.3GPa. Such an high increase is mainly due to the fact that the skins of the sandwich are reinforced with
continuous fibers rather than chopped one.
The average maximum strain instead is 3% and it is the same as that of core-L. This means that the skins
have no effect on the maximum deflection of our specimens, but it is ruled by the core: as the maximum
strain is reached, the core fails and, being no more able to keep the unidirectional skins together, the
sandwich fails. At last we analyze the behavior of the cross ply sandwiches(Graph 4). Contrarily to what can
be thought, the addition of a second skin on each face of the specimens, even if with the continuous fibers
perpendicular to the flow direction, do not lead to beneficial effects. The average flexural strength is
460MPa ±21.2MPa in the case of CP0/90 and 440MPa ±39MPa for CP90/0 specimens, while the flexural
moduli are 18.8GPa ±0.6GPa for CP0/90 and 17.4GPa ±0.1GPa for CP90/0. Concerning the average
maximum strains, CP0/90 samples show a value of 2.8% while CP90/0 samples a value 3%. The higher
strength displayed by CP0/90 with respect to CP90/0 is explained by the fact that in the first case the 0°
plies are more distant from the axis of inertia. However the presence of the 90° skin decreases the strength
as can be seen if compared to the UD samples. Also in this case it can be seen that the maximum deflection
is determined by that of the core and that a brittle failure occurs in both configurations.
33
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5 5.0 5.5
0
100
200
300
400
core-T
Fle
xura
l Str
ess (
MP
a)
Flexural Strain (%)
core-L
Graph 2 - Three Point Bending test of Radistrong core at T=23°C: comparison between longitudinal (L) and transversal (T) specimens
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5
0
100
200
300
400
500
600 UD
Fle
xura
l Str
ess (
MP
a)
Flexural Strain (%)
core-L
Graph 3 - Three Point Bending test at T=23°C: comparison between longitudinal core-L and unidirectional sandwich (UD)
34
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5
0
200
400
600
Fle
xura
l str
ess (
MP
a)
Flexural Strain (%)
CP0/90
CP90/0
UD
Graph 4 - Three Point Bending test at T=23°C: comparison between unidirectional sandwich (UD) and cross-ply sandwich (CP0/90, CP90/0)
4.3.1.1 – THREE POINT BENDING: EFFECT OF TEMPERATURE
In order to evaluate the effect of temperature on the flexural properties of the material considered, the
stress-strain curves at -20°C and 80°C were plotted together to the already shown curves obtained at room
temperature. Graph 5 andGraph 6 show the results obtained from the Radistrong core. At high
temperatures lower strength and higher strains are achieved. Specifically, regarding core-L it has been
calculated a decrease in ultimate strength of about 40%, down to 208MPa, and a maximum strain around
5%, thus 60% higher with respect to room temperature. In fact, at 80°C, so over the glass transition
temperature, the polymeric chains are more mobile and allow for an higher deformation. Being the
material more ductile, the curves show a great deviation from the linear elastic behavior, accounting for
plastic deformations before fracture. The same trend is followed by core-T samples: the maximum stress
decreases to 85MPa ±3.5MPa and the maximum strain increases to 15%. Obviously in this case the
differences are amplified as the matrix plays a more important role. It must also be noted that in the core-T
specimens the maximum stress is reached well before failure, meaning that some mechanisms of crack
propagation are occurring and that they are not instantaneous as in the other cases. As the flexural stress,
the bending moduli decreases increasing the temperature showing the lower stiffness of the composite
material. They have values of 9.3GPa ±0.9GPa and 2.8GPa ±0.2GPa for core-L and core-T samples,
respectively. On the contrary, at -20°C degree Radistrong is more rigid and shows only a linear elastic
behavior. Regarding core-L specimens, higher ultimate strength (405MPa ±12MPa) and higher modulus
(14.1GPa ±0.4GPa) are reached at the same average maximum strain of room temperature, while core-T
specimens showed lower maximum deflection (3%) and higher modulus (7.3GPa ±0.1GPa) at approximately
the same ultimate stress level. Again, this discrepancy can be explained by the fact that in core-L specimens
0°-layers have more importance so the deflection is constrained by the fibers placed longitudinally, while in
core-T specimens it can be said that the limit is the maximum stress that the matrix and the fiber-matrix
interface are able to withstand.
35
Considering now UD-samples, it can be seen from Graph 7 that they follow the same trend with respect to
temperature as bare Radistrong. Increasing the temperature to 80°C we encounter a decrease in maximum
flexural stress, which for UD-samples is 343MPa ±5.2MPa, and in modulus (17.7GPa ±1GPa), together with
an increase of the average ultimate strain to 3.7%. On the contrary, the curves obtained at -20°C show an
increase of the maximum flexural stress and strain to about 611MPa ±37MPa and 3.3%, respectively, while
the modulus attains a value of 19.6GPa ±0.4GPa, very close to that of the tests at 23°C. Making some
calculations, it appears that the rise of temperature resulted in a decrease of about 35% of the maximum
stress, while the lowering of it brought to an increase of 15.7%. Again, the effect of higher temperatures is
more important than that of lower temperature. Moreover, comparing the effect of temperature upon UD
and core samples, it can be noted that the decrease in stiffness at 80°C is mitigated by the presence of the
continuous fibers reinforced film. We may also say that even if a brittle failure is experienced by UD-
specimens at 23°C and -20°C, at 80°C, together with large plastic deformations, UD-samples show an
evident step-propagation of the cracks until failure, which not always results in a complete fracture of the
specimens, contrarily to the other cases.
Regarding the cross-ply configuration at 80°C (Graph 8 andGraph 9), the average ultimate strength is
278MPa ±11MPa for CP0/90 specimens and 300MPa ±4.6MPa for CP90/0 specimens, while the maximum
strain is 4.1% and 3.3% for CP0/90 and CP90/0, respectively. The moduli instead attain values of 14.3GPa
±0.11GPa for CP0/90 and of 13.5GPa ±0.17GPa for CP90/0. It can be thus seen that again at higher
temperature a decrease in stiffness occurs. However we observe different behaviors in the two cases: while
CP0/90 displays a lowering of strength of 39% and a rising of strain of 46%, CP90/0 has a strength lower
than that of samples tested at 23°C of about 32% and an higher maximum strain of about 10%. From the
results thus appears that the CP90/0 configuration has a slightly higher stiffness than CP0/90, while the
latter shows to be able to bear markedly higher deformations. The testing of cross-plies at -20°C, instead,
gave results which are very similar to those obtained at 23°C: in the case of CP0/90, the average maximum
stress and strain are 497MPa ±15MPa and 3%, with a flexural modulus of 18.7GPa ±0.7GPa; in the case of
CP90/0, the ultimate strength and strain are 468MPa ±18MPa and 3.3%, with a flexural modulus of 17GPa
±0.7GPa.
0 1 2 3 4 5 6
0
50
100
150
200
250
300
350
400
450
core-L(80°C)
core-L(23°C)
Fle
xura
l Str
ess (
MP
a)
Flexural Strain (%)
core-L(-20°C)
Graph 5 - Effect of temperature on longitudinal core specimens
36
0 1 2 3 4 5 6 7 8 9 10 11 12 13 14
0
50
100
150
200
core-T(80°C)
core-T(-20°C)
Fle
xura
l str
ess
(M
Pa)
Flexural strain (%)
core-T(23°C)
Graph 6 - Effect of temperature on transversal core specimens
0 1 2 3 4
0
100
200
300
400
500
600
700
UD(80°C)
UD(23°C)
Fle
xura
l str
ess
(M
Pa)
Flexural Strain (%)
UD(-20°C)
Graph 7 - Effect of temperature on unidirectional specimens tested with Three Point Bending
37
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5
0
100
200
300
400
500
600
CP0/90(-20°C)
CP0/90(80°C)
CP0/90(23°C)
Fle
xura
l str
ess (
MP
a)
Flexural Strain (%)
Graph 8 - Effect of temperature on cross-ply 0/90 specimens with Three Point Bending
0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0 4.5
0
100
200
300
400
500
600
Fle
xura
l str
ess (
MP
a)
Flexural Strain (%)
CP90/0(-20°C)
CP90/0(23°C)
CP90/0(80°C)
Graph 9 - Effect of temperature on cross-ply 90/0 specimens with Three Point Bending
38
4.3.2 – INTERLAMINAR SHEAR STRENGTH TEST
As can be deduced from Graph 10, ILSS tests at room temperature reveal that UD sandwiches have an
higher adhesion with respect to both CP configurations. In fact in the first case, a load of about 6.15kN
±173N is necessary to the specimens to fail, corresponding to an average apparent interlaminar shear
strength τ of 63.3MPa ±1.3MPa, while it decreases to 5.38kN ±297N (τ = 43.2MPa ±4.3MPa) for CP0/90 and
to 5.88kN ±344N (τ =51.2MPa ±2.6MPa) for CP90/0. It is known that the shear stress attains a maximum at
the center of a beam under deflection, so it could be thought that the failure of the specimens begins from
the core. In reality, in composite laminates the weakest part is the interface between two layers. It has
been observed, recording an ILSS test on a UD sandwich at 23°C with an high velocity camera, that the
failure begins right under the central loading pin. Specifically, looking at Figure 24, we have at first a crack
running at the interface between the core and the UD skin. As the load rises the crack increases its length
along the interface, until when it starts to propagate perpendicularly to the interface across the core. At
this point the failure occurs rapidly and in a brittle manner.
From visual inspection of the samples failed (Figure 25 Figure 26), the superficial skins reinforced with
continuous fibers showed to be severely damaged in the UD specimens, with fibers broke and torn away
from the film. Looking instead at the cross-plies, the 0°-fibers do not appears to break. It can be inferred
that the critical point in these materials is the interface lying between the 0°-skin and the layer above
(counting from the bottom of the sandwich), regardless if is the Radistrong core or the other Cetex skin. In
fact, over the core in the CP0/90 sample shown in Figure 26 a light gray layer appears: it is the 90°-skin,
thus failure occurred at the 0°/90° skins interface, with the 0°-layer detached and the 90°-one attached to
the core. Both layers instead were peeled off from the core in the CP90/0 case, as the crack runs through
the interface between the 0°-ply and the core. These findings are in line with the loads displayed in the
graph before: in UD sandwiches the adhesion between the different layers is stronger. Moreover the 0°-ply
seems to create an higher adhesion with the core rather than with the 90°-ply, as the outermost layer of
the core is composed of short fibers aligned in the same direction, allowing for a stronger bond without the
impinging of fibers that may occur between two orthogonal plies.
As a further proof, the ILSS test was performed on a sandwich made with two 0°-skins. As depicted in Graph
11, a load 28% higher than for the UD is necessary to make the sample fail: 7.9kN ±487N, which
corresponds to a shear strength level of about 68.6MPa ±2.9MPa. (Figure 27)
39
0.0 0.5 1.0 1.5
0
1000
2000
3000
4000
5000
6000
7000
CP0/90
CP90/0
Load (
N)
Displacement (mm)
UD
Graph 10 - Interlaminar Shear Strength test on unidirectional and cross-ply specimens: comparison at T=23°C
Figure 24 - Crack development during ILSS testing of UD specimen
40
Figure 25 - UD specimen after ILSS test at T=23°C
Figure 26 - CP0/90 (left) and CP90/0 (right) specimens after ILSS test at T=23°C
Figure 27 – 00-specimen after ILSS test at T=23°C
0.0 0.5 1.0 1.5
0
2000
4000
6000
8000
00
UD
Load (
N)
Displacement (mm)
Graph 11 - ILSS test: comparison between UD and 00-specimens
41
4.3.2.1 – INTERLAMINAR SHEAR STRENGTH TEST: EFFECT OF TEMPERATURE
As in the case of three point bending, higher temperatures have a bigger effect on our materials than lower
ones. At 80°C polyamides are more ductile and lower shear strengths are needed for the interfaces in the
materials to break. On the contrary at -20°C an higher stiffness of the material is noticed. Graph 12,Graph
13 and Graph 14 summarize the results obtained. In the case of UD samples, at T=80°C the maximum load
decreases to 4.22kN ±126N, corresponding to a shear strength of 43.8MPa ±0.46MPa, while at T=-20°C a
slight increase of the load, and consequently of the shear strength, is measured (6.32kN ±301N and
61.4MPa ±0.85MPa, respectively). Concerning now cross-ply specimens, at T=80°C the load to failure is
3.9kN ±104N for CP0/90 and 4kN ±120N for CP90/0 (τ=34.1MPa±1.03MPa in the first case and
τ=35MPa±0.95MPa in the second one). At -20°C instead the load recorded is 6.51kN ±330N (τ=54.9MPa
±2.08MPa) in the case of CP0/90, while 6.2kN ±208N (τ=53.9MPa ±1.8MPa) in the case of CP90/0.
It is readily noticeable that the two cross-ply configurations show a very similar behavior at every
temperature.
0.0 0.5 1.0 1.5 2.0
0
1000
2000
3000
4000
5000
6000
7000
UD(80°C)
UD(23°C)
Load (
N)
Displacement (mm)
UD(-20°C)
Graph 12 - Effect of temperature on ILSS unidirectional specimens
42
0.0 0.5 1.0 1.5 2.0
0
1000
2000
3000
4000
5000
6000
7000 CP0/90(-20°C)
Load (
N)
Displacement (mm)
CP0/90(23°C)
CP0/90(80°C)
Graph 13 - Effect of temperature on ILSS cross-ply 0/90 specimens
0.0 0.5 1.0 1.5 2.0
0
1000
2000
3000
4000
5000
6000
7000CP90/0(-20°C)
CP90/0(80°C)
Load (
N)
Displacement (mm)
CP90/0(23°C)
Graph 14 - Effect of temperature on ILSS cross-ply 90/0 specimens
43
4.3.3 – CLASSICAL LAMINATION THEORY
In order to apply CLT and determine the stress distribution inside our materials during the tests, the
mechanical properties of both the short-fibers reinforced Radistrong and the continuous fiber reinforced
film Cetex have to be known. In this work, Young moduli and Poisson ratios were estimated by calculating
the flexural modulus of our specimens at different values of E1 and Poisson ratio and varying, for each value
of E1, the modulus in the transversal direction, E2. The calculated flexural moduli were then compared to
those extrapolated from the three point bending test curves. Noteworthy is to specify that in the case of
the core, E1 is the modulus in the flow direction, while in the case of the skins it is calculated in the
direction of the continuous fibers. First of all the moduli of the core are determined. As can be seen in
Graph 15 and Graph 16, for a fixed Poisson ratio ����� = 0.35 the flexural moduli experimentally obtained
are achieved in the ranges 13.5GPa≤������≤15GPa and 5.8GPa≤��
����≤6.5GPa. The results which fits better
showed to be ������=14.8GPa and ��
����=6.2GPa.
Having defined the Young moduli of the Radistrong core, we can apply the same approach to analyze our
sandwiches in order to find the mechanical properties of the continuous fibers reinforced film. Looking at
Graph 17 it can be seen that in the case of UD-sandwiches there are no effects of the skin’s modulus in the
transversal direction with respect to the fibers, as in this test it plays an irrelevant role. Thus from the
analysis of the UD-sandwiches we can easily found the modulus in the direction of the fibers, that results to
be ������=29.5GPa.
0 5000 10000 15000 20000
10000
12000
14000
16000
18000
20000
Efl
ex
E2
13500MPa 14000MPa 14500MPa 15000MPa 15500MPa
ETPB
=14400MPa
E1
Graph 15 - Variation of flexural modulus for core-L specimens (ν12=0.35)
44
0 5000 10000 15000 20000
0
2000
4000
6000
8000
10000
12000
14000
16000
18000
20000
22000
ETPB
=6700MPa
Efl
ex
E2
Graph 16 - Variation of flexural modulus for core-T specimens (ν12=0.35)
0 5000 10000 15000 20000
14000
16000
18000
20000
22000
28000MPa 29000MPa 30000MPa 31000MPa
E1E
fle
x(M
Pa)
E2(MPa)
ETPB
=20500MPa
Graph 17 - Variation of flexural modulus for UD specimens (ν12=0.3)
Graph 18 andGraph 19 show the results of the calculations of the flexural modulus for CP0/90 and CP90/0
sandwiches. Considering the ranges of accordance between experimental and analytical values of the
transversal modulus (11.7GPa≤������≤18.7GPa for CP0/90, 11.6GPa≤��
����≤15.6GPa for CP90/0), we can
conclude that, concerning the Cetex film, ������=29.5GPa and ��
����=11.7GPa.
45
0 5000 10000 15000 20000
14000
15000
16000
17000
18000
19000
20000
21000
25000MPa 30000MPa
E1
ETPB
=18800MPa
Efl
ex (
MP
a)
E2 (MPa)
Graph 18 - Variation of flexural modulus for CP0/90 specimens (ν12=0.3)
0 5000 10000 15000 20000
12000
13000
14000
15000
16000
17000
18000
19000
20000
25000 MPa 30000 MPa
E1
ETPB
=17400MPa
Efl
ex (
MP
a)
E2 (MPa)
Graph 19 - Variation of flexural modulus for CP90/0 specimens (ν12=0.3)
Once the basic mechanical properties of each material have been found, it is possible to finally apply
classical lamination theory. Graph 20 to Graph 24 show the results: the stresses are calculated using the
maximum load recorded during the three point bending tests. First of all it can be inferred that no
significant differences in the values of stresses exist across the thickness of core-L and core-T specimens,
despite the fact that the latter is able to bear a substantially lower load. This is an evidence that with the
change in the orientation of the short fiber layers, the maximum internal stress in direction 1 and the
maximum shear stress are reached with a much lower applied force. Concentrating now on the sandwiches,
it can be readily noted that the 0°-skin is the most load-carrying layer, as it can be expected. It becomes
now clear the reason why the most critical part in our laminates was the interface between the 0°-layer and
the others, as discontinuities are detected due to the property mismatch between the layers.
46
Moreover, these discontinuities are sharper between 0°- and 90°-skins rather than between the 0°-skin and
the core, explaining the better adhesion obtained in the UD-sandwiches.
-400 -200 0 200 400
-1
0
1
Th
ickn
ess (
mm
)
xx
(MPa)
core-0°
core-90°
core-90°
core-0°
0 100 200
-1
0
1
xz
(MPa)
Thic
kness (
mm
)
core-0°
core-90°
core-90°
core-0°
Graph 20 - Stress distribution across the thickness of core-L samples
-400 -200 0 200 400
-1
0
1
Th
ickn
ess (
mm
)
xx
(MPa)
core-90°
core-0°
core-0°
core-90°
0 100 200
-1
0
1
Th
ickn
ess (
mm
)
xz
(MPa)
core-90°
core-0°
core-0°
core-90°
Graph 21 - Stress distribution across the thickness of core-T samples
-800 -600 -400 -200 0 200 400 600 800
-1
0
1
Th
ickn
ess (
mm
)
xx
(MPa)
skin-0°
core-0°
core-90°
core-90°
core-0°
skin-0°
0 100 200 300 400
-1
0
1
Th
ickn
ess (
mm
)
xz
(MPa)
skin-0°
core-0°
core-90°
core-90°
core-0°
skin-0°
Graph 22 - Stress distribution across the thickness of UD-samples
47
-600 -400 -200 0 200 400 600
-2
-1
0
1
2
Th
ickn
ess (
mm
)
xx
(MPa)
skin-90°skin-0°
core-0°
core-90°
core-90°
core-0°
skin-90°skin-0°
0 100 200 300
-2
-1
0
1
2
Th
ickn
ess (
mm
)
xz
(MPa)
core-0°
skin-90°skin-0°
core-90°
core-90°
core-0°
skin-90°skin-0°
Graph 23 - Stress distribution across the thickness of CP0/90 samples
-600 -400 -200 0 200 400 600
-2
-1
0
1
2 skin-90°
skin-90°
Th
ickn
ess (
mm
)
xx
(MPa)
skin-0°
core-0°
core-90°
core-90°
core-0°
skin-0°
0 100 200 300
-2
-1
0
1
2
Th
ickn
ess (
mm
)
xz
(MPa)
skin-0°skin-90°
core-0°
core-90°
core-90°
core-0°
skin-0°skin-90°
Graph 24 - Stress distribution across the thickness of CP90/0 samples
48
4.4 – FALLING WEIGHT IMPACT TEST
The impact properties of the Radistrong core and of the cross-ply sandwiches were also investigated. We
have seen from the previous experimental work that Radistrong behaves in a brittle manner, as expected
from a semi-crystalline polymer. This was observed also in this case: in fact, when the impact energy was
high enough, the core specimens were perforated by the impactor without signs of matrix yielding or
plastic deformations (Figure 28). The dominant mechanisms of fracture in short fiber reinforced composites
usually are:
- Fiber-matrix interfacial debonding,
- Fiber fracture,
- Matrix fracture,
- Fiber pull-out.
Each of these failure mechanisms contributes in varying degrees to the fracture toughness of the composite
as well as to the energy absorption under impact loading, but they do not necessarily operate
simultaneously for a given fiber-matrix system. In a composite containing brittle fibers in a brittle matrix, as
our material, debonding and pull-out energies are considerably higher than matrix and fiber fracture
energies. Moreover if the fibers are randomly oriented, there may be additional energies absorbed due to
bending and shear deformations of the fibers (17).
On the other hand, in the case of sandwiches the energy absorption mechanisms of the short fiber
reinforced core are overwhelmed by those of the continuous-fiber reinforced skins. Delamination and fiber-
matrix debonding are the most evident phenomena occurring during our tests (Figure 29). The first involves
the separation of the skins from the core, especially at the edges of the samples, while the second displays
itself on the outer layer on the non-impacted side of the specimens. Depending on the impact energy, fiber
breakage can also be observed. Graph 25 shows the differences in energies absorbed between the core and
the cross-ply samples during impact tests performed at room temperature. It can be readily noted the
strengthening effect of the application of the two orthogonal skins: while the bare core reaches an average
absorbed energy of 11.4J ±2.6J, the mean energy that the cross-ply sandwiches are able to absorb is 34.9J
±0.9J (CP0/90) and 29.4J ±14J (CP90/0). Looking at the data we can infer that the two CP configurations
behave approximately at the same manner.
From the shape of the curves in the plot we can also arrive to an important conclusion: the core specimens
are penetrated by the falling weight while the cross-plies are not. In fact, while the firsts rise until a plateau,
the curves of the sandwiches attain a constant value of energy only after a peak. The latter corresponds to
the energy dissipated by the specimen to make the impactor bounce back.
Figure 28 - Core specimen after falling weight impact test at T=23°C: impacted side (left) and bottom side (right)
49
Figure 29 - Cross-ply specimen after falling weight impact test at T=23°C: impacted side (left) and bottom side (right)
0.00 0.02 0.04
0
5
10
15
20
25
30
35
40
CP0/90
CP90/0
Energ
y (
J)
Time (s)
core
Graph 25 - Impact energy versus time at T=23°C
4.4.1 – FALLING WEIGHT IMPACT TEST: EFFECT OF TEMPERATURE
As for flexural tests, also in this case the behavior of our materials at different temperatures was analyzed.
Considering first of all the core, Graph 26 shows that temperature has a great incidence on the impact
behavior of our short fiber composite. In fact, at T=-20°C Radistrong can absorb lower energy with respect
to room temperature (3.8J ±1.8J) , while at T=80°C it is increased to an average value of 20.2J ±2.7J. This
demonstrates one more time the embrittling effect of low temperatures and the ability of higher ones to
make composites more ductile. The minimum impact energies used and able to penetrate the samples
were 16J at 23°C, 20.6J at 80°C and 6J at -20°C.
Cross-ply sandwiches instead showed a somehow different behavior as can be seen in Graph 27 andGraph
28 Again at low temperatures we have a lower absorption of energy (15.7J ±1.5J for CP0/90 and 16.8J ±2J
for CP90/0) but this time it is accompanied by a sharper peak with respect to that depicted by the curves at
T=23°C, meaning that the samples have an higher rigidity and then more energy is dissipated by the
50
rebound of the impactor. On the contrary, at T=80°C the peak flattens and the energy absorbed increases
to 37.1J ±4J and 34.6J ±0.03J for CP0/90 and CP90/0 respectively.
0.00 0.02 0.04
0
5
10
15
20
core(80°C)
core(-20°C)
core(23°C)
Energ
y (
J)
Time (s)
Graph 26 - Falling Weight Impact Test on core-specimens: effect of temperature
0.00 0.02 0.04
0
5
10
15
20
25
30
35
40
Energ
y (
J)
Time (s)
CP0/90(23°C)
CP0/90(-20°C)
CP0/90(80°C)
Graph 27 - Falling Weight Impact Test on CP0/90-specimens: effect of temperature
51
0.00 0.02 0.04
0
5
10
15
20
25
30
35
40
Energ
y (
J)
Time (s)
CP90/0(-20°C)
CP90/0(23°C)
CP90/0(80°C)
Graph 28 - Falling Weight Impact Test on CP90/0-specimens: effect of temperature
From visual inspection of the fracture surfaces of the CP samples tested some considerations can be drawn
(Figure 30 andFigure 31). First of all those tested at room temperature have an evident hole due to the
puncture of the falling weight, accompanied by the breakage of the continuous fibers of the skins and of
the fiber/matrix interface. At 20°C instead the penetration of the dart is reduced, but the energy of the
impact is used to develop delamination of the skins on the non-impacted side and again fiber/matrix
interface failure. On the contrary, at 80°C samples still demonstrate a lower penetration damage with
respect to 23°C, but in this case the temperatures, being higher than the glass transition one, allows to
dissipate the impact energy on the impacted side giving rise to buckling of the continuous fibers. Even in
this case fiber/matrix debonding is present, demonstrating how critical this part of the composite is.
It can also be inferred that the lower effect that higher temperatures have on the impact behavior is in line
with the higher importance that they have in flexural tests. In fact, flexural tests were performed at a
displacement rate of 1mm/min: they thus were long tests, in which the microstructure had time to show its
increased ductility, to reorganize and hence influence the response of the specimens to the load. On the
contrary impact events are very brief phenomena, as can be seen by the time scale in the graphs above.
This follows that the increased polymer matrix stiffness at low temperature, which does not need time to
develop its effect, is more prone to have an influence on the tests.
52
Figure 30 - Cross-ply specimen after falling weight impact test at T=-20°C: impacted side (left) and bottom side (right)
Figure 31 - Cross-ply specimen after falling weight impact test at T=80°C: impacted side (left) and bottom side (right)
53
5. – CONCLUDING REMARKS
It has been demonstrated that the application of continuous glass fiber reinforced polyamide 6 films over
an injection molded short glass fiber reinforced polyamide 66/MXD6 blends effectively results in a
strengthened material. In fact, both flexural stiffness and impact resistance are higher in the case of
sandwiches. Concerning mechanical properties under bending, we obtained the best results applying only
one layer with continuous fibers parallel to the injection flow of the short fiber reinforced core. Cross-ply
sandwiches instead showed to have lower adhesion to the core and this behavior is attributed to the
presence of the layer with continuous fibers placed at 90° with respect to the injection flow. This can be
inferred as the testing of samples with two 0°-continuous fibers layers ended up in higher interlaminar
shear strength than both UD- and CP-samples, meaning that the processing condition were able to make
two plies adhere to the substrate. However, this higher adhesion is also a drawback as all samples with 0°-
continuous fibers only failed in a catastrophic way, contrarily to cross-ply sandwiches, and this surely is a
problem for a materials thought to satisfy structural applications. Moreover, the impact resistance of UD
sandwiches, both with one or two 0°-plies, should be lower than cross-ply ones as impact loadings induce a
biaxially distributed stress state in the material.
In order to overcome these adhesion problems, further testing can be made considering a different
stacking sequence of the continuous glass fiber reinforced tapes. Placing these plies at different angles
smaller than 90° may avoid the occurrence of an abrupt transition in the distribution of stresses across the
thickness of the samples, as happens between 0°- and 90°-layers. In this case it is suggested not to use
Classical Lamination Theory but more powerful tools like First-Order Shear Deformation Laminated Plate
Theory. As can be seen from the results reported, our materials undergo high deflections during three point
bending tests, thus a model able to include the rotations of the deformed samples may give a better
approximation of the stress distribution.
Changing the approach to the solution, a way to improve the adhesion and the mechanical properties of
these sandwiches could be the use of another production process such as the Spriform process (33). During
this process, continuous fiber reinforced thermoplastics prepregs are firstly thermoformed through heating
in a infrared oven and placing in a mold to give the desired shape. The same mold is then utilized by a
standard injection machine to infuse short fiber reinforced thermoplastics. This production method should
be able to achieve a stronger adhesion between the core and the continuous fiber layers and it is capable
of large mass production.
54
6. – REFERENCES
1. The effect of temperature changes on to quasi-static tensile and flexural performance of glass fiber
reinforced PA66 composites. Butterworth, Ian, et al. 2013. Proceedings of the 11th International
Conference on Manufacturing Research (ICMR2013).
2. Mouti, Zakaria. Localised Low Velocity Impact Performance of Short Glass Fibre Reinforced Polyamide 66
Oil Pans. 2012.
3. Hassan, A., et al. Tensile, Impact and Fiber Length Properties of Injection-molded Short and Long Glass
Fiber-reinforced Polyamide 6,6 Composites. Journal of reinforced Plastics and Composites. 2004, Vol. 23, 9.
4. Advani, Suresh G. and Sozer, E. Murat. Process Modeling in Composite Manufacturing. s.l. : Marcel
Dekker, Inc., 2003.
5. Carlsson, Leif A. Thermoplastic Composite Materials. s.l. : Elsevier Science Publisher, 1991.
6. Mascia, L. Thermoplastics: Materials Engineering. s.l. : Elsevier Science Publisher, 1989.
7. Jeyakumar, A. Solid-state modification of polyamide-6,6. 2012.
8. Marchildon, Keith. Polyamides – Still Strong After Seventy Years. Macromolecular Reaction Engineering.
2011, 5, pp. 22-54.
9. Chanda, Manas and Roy, Salil K. Plastics Technology Handbook. s.l. : Taylor&Francis Group, 2006.
10. Margolis, James M. Engineering Thermoplastics. Properties and Applications. s.l. : M. Dekker, Inc., 1985.
11. Dasgupta, Siddharth, Hammond, Willis B. and Goddard III, William A. Crystal Structures and Properties
of Nylon Polymers from Theory. Journal of American Chemical Society. 1996, 118, pp. 12291-12301.
12. Ridgway, James S. Structure-Property Relationships of Ring-Containing Nylon 66 Copolyamides. Journal
of Polymer Science: part A-1. 1970, Vol. 8, pp. 3089-3111.
13. Yoda, Naoya and Matsubara, Ikuo. Studies on Structure and Properties of Aromatic Polyamides. II.
Structural Elucidation of Poly (m-xylylene Adipamide). Journal of Polymer Science: Part A. 1964, Vol. 2, pp.
253-265.
14. Datasheet TDS RADISTRONG® A-X15120 BK.
15. Datasheet Cetex® TC910.
16. Pivi, Francesca. Manufacturing and mechanical characterization of a thermoplastic composite material
for automotive applications. 2016.
17. Kelly, Anthony and Zweben, Carl. Comprehensive Composite Materials. s.l. : Elsevier Ltd., 2001.
18. Hodgkinson, J.M. Mechanical Testing of Advanced Fibre Composites. s.l. : Woodhead Publishing Ltd.,
2000.
19. ISO14125. Fibre-reinforced plastic composites — Determination of flexural properties. 1998.
55
20. ISO14130. Fibre-reinforced plastic composites - Determination of apparent interlaminar shear strength
by short-beam method. 2003.
21. Lee, S.M., Cheon, J.S. and Im, Y.T. Experimental and numerical study of the impact behavior of SMC
plates. [book auth.] I.H. Marshall. Composite Structures. s.l. : Elsevier Ltd., 1999.
22. Ives, G.C., Mead, J.A. and Riley, M.M. Handbook of Plastic Test Methods. s.l. : The Plastic Institute,
1971.
23. ISO6603-1. Plastics — Determination of puncture impact behaviour of rigid plastics. Part 1: Non-
instrumented impact testing. 2000.
24. ISO6603-2. Plastics — Determination of puncture impact behaviour of rigid plastics. Part 2:
Instrumented impact testing.
25. Reddy, J.N. Mechanics of Laminated Composite Plates and Shells. Theory and Analysis. s.l. : CRC Press,
2004.
26. Yasuniwa, Munehisa, et al. Thermal Analysis of the Double-Melting Behavior of Poly(L-lactic acid).
Journal of Polymer Science: Part B: Polymer Physics. 2004, Vol. 42, pp. 25-32.
27. Hassan, Aziz, Rahman, Normasmira A. and Yahya, Rosiyah. Moisture Absorption Effect on Thermal,
Dynamic Mechanical and Mechanical Properties of Injection-Molded Short Glass-Fiber/Polyamide 6,6
Composites. Fibers and Polymers. 2012, Vol. 13, 7, pp. 899-906.
28. Starkweather, Howard W., et al. The Heat of Fusion of 66 Nylon. Journal of Polymer Science: Polymer
Physics Edition. 1984, Vol. 22, pp. 1615-1621.
29. Fereydoon, Maryam, Tabatabaei, Seyed H. and Ajji, Abdellah. Rheological, Crystal Structure, Barrier,
and Mechanical Properties of PA6 and MXD6 Nanocomposite Films. Polymer Engineering and Science. 2013.
30. Zhang, Gongzheng, et al. Phase Transition Behavior of Nylon-66, Nylon-48, and Blends. Polymer Journal.
2003, Vol. 35, 2, pp. 173—177 .
31. Ramesh, C., Kellert, A. and Eltink, S. J. E. A. Studies on the crystallization and melting of nylon-6,6: 1.
The dependence of the Brill transition on the crystallization temperature. Polymer 35 . 1994, Vol. 35, 12.
32. Murthy, Sanjeeva N. Hydrogen Bonding, Mobility, and Structural Transitions in Aliphatic Polyamides.
Journal of Polymer Science: Part B: Polymer Physics. 2006, Vol. 44, pp. 1763–1782.
33. http://www.hbw-gubesch.de.
