Improving the Ply/Interleaf Interface in Carbon Fibre Reinforced Composites with Variable Stiffness

Henry A. Maplesa, Oluwadamilola Smithb, Christoph Burgstallerc, Paul Robinsonb,*, Alexander Bismarcka,d,*

a Polymer and Composite Engineering (PaCE) Group, Institute of Materials Chemistry, Faculty of Chemistry, University of Vienna, Währinger Straße 42, A-1090 Vienna, Austriab The Composites Centre, Department of Aeronautics, Imperial College London, South Kensington Campus, London SW7 2AZ, UKc Transfercenter für Kunststofftechnik (TCKT), Franz-Fritsch-Straße 11, A-4600 Wels, Austriad Polymer and Composite Engineering (PaCE) Group, Department of Chemical Engineering, Imperial College London, South Kensington Campus, London SW7 2AZ, UK

*Corresponding Authors, +43(1)427771301, a.bismarck@imperial.ac.uk; alexander.bismarck@univie.ac.at (A. Bismarck), +44 (0)20 7594 5078, p.robinson@imperial.ac.uk (P. Robinson)Keywords: Laminate (A); carbon fibres (A); structural composites (A); interleaved composites; interface (B).

Abstract

Polystyrene-interleaved carbon fibre reinforced epoxy composites exhibiting controllable

stiffness have been manufactured. These composites undergo reductions in flexural stiffness of

up to 99% when heated above the Tg of the interleaf layers. Potential applications for such

materials include their use in morphing and deployable structures. Flexural tests at room

temperature indicated that improvements in adhesion between the polystyrene and CFRP layers

are required to prevent premature failure of the composites at low shear stresses. Here we

investigate how modification of the interleaf layer improves the interlaminar shear strength of

the laminates without affecting the stiffness loss at elevated temperatures. Two poly(styrene-co-

maleic anhydride) (SMA) films with different maleic anhydride content were prepared and used

as interleaf films. Thick adherend shear tests showed that the adhesion strength more than

doubled, while flexural tests showed that composites containing SMA interleafs had more than

twice the apparent flexural strength of composites containing pure polystyrene layers at 25°C

and yet still undergo significant reductions in stiffness at elevated temperature.

1. Introduction

Much research has focused on the development of variable stiffness composites for use in shape

adaptive structures [1-3]. Potential applications for these composites include their use as skin

materials in morphing and deployable structures [1-4]. For these applications a composite

material is required whose stiffness can be reduced on demand and which can then undergo

large, reversible deformations without failing. There are many designs of controllable stiffness

composites. We showed previously that polyacrylamide (PAAm) coated carbon fibre reinforced

epoxy composites experience reversible reductions in stiffness of up to 88% when heated above

the glass transition temperature (Tg) of PAAm [4, 5]. This loss in stiffness was due to softening

of the PAAm coating, which allowed the fibres to slide within the composite, thus limiting stress

transfer between fibres and matrix. The stiffness was fully recovered when the composites were

cooled to room temperature. Our concept was since adapted by Bachinger et al. [6, 7] to produce

similar controllable stiffness composites. Their materials are designed to be used as vehicle

components that can soften on demand to provide pedestrian protection during road accidents.

Henry and McKnight [8] developed a form of variable stiffness composites consisting of a

polyurethane shape memory polymer (SMP) matrix with discontinuous steel reinforcement. The

SMP was heated until it softened, which led to reductions in storage modulus of up to 99% when

tested using dynamic mechanical thermal analysis (DMTA). After deforming the composites at

elevated temperature, the original shape of the specimens was partly recovered using the SMP as

actuator. Another form of variable stiffness composite consists of braided flexible matrix

composite (FMC) tubes within an elastomer matrix [9]. When the tubes are pressurised with a

fluid they either expand or contract depending on their design, leading to a change in composite

shape and stiffness.

There have been a number of studies focusing on the development of interleaved composites

with controllable stiffness [10-13]. McKnight et al. [14] patented the concept of a structure

containing layers of constant and variable stiffness elements arranged in an alternating sequence.

In a previous paper we showed that polystyrene (PS) interleaved carbon fibre reinforced epoxy

composites undergo considerable reductions in flexural stiffness when heated above the Tg of the

polystyrene layers using the carbon fibres as Joule heaters [15-17]. At these elevated

temperatures the composites can be deformed significantly and then fully recover upon removing

the load. However, the problem with interleaved composites is that the adhesion between the

interleaf and matrix materials is often (very) weak, resulting in low toughness [18-21]. Some

thermoplastic interleaf materials have low surface energies, resulting in poor wetting by uncured,

liquid thermosetting resins used in prepregs, which can lead to poor adhesion strength at the

interleaf/CFRP interface. There are a number of ways to overcome this problem, including

plasma and chemical treatments [22, 23]. Heness et al. [18] showed that plasma etching and

chemical treatment does improve the adhesion of Nylon interleaf films to an epoxy matrix. Nay

et al. [19] used a plasma induced polymerisation technique to deposit thin polyisoprene and

polybutadiene layers onto polyimide interleaf films. These polymer layers improved the adhesion

between the polyimide films and the carbon fibre reinforced epoxy plies. In our case, the issue of

poor adhesion between polystyrene layers and CFRP plies in interleaved laminates leads to

composites prematurely failing in shear during flexural tests. Here we selected a modified

polystyrene (random poly(styrene-co-maleic anhydride) (SMA)) as interleaf material, which

should provide better adhesion to the carbon fibre-epoxy layers than pure PS. The incorporated

maleic anhydride (MAH) moieties in SMA provide sufficient reactivity and make SMA a

suitable compatibilising or crosslinking agent. Maleated polymers are often used as matrices for

glass or natural fibre reinforced composites [24, 25]. We will investigate how this difference in

adhesion affects the mechanical properties of the composites. We chose to use a different

interleaf material as opposed to post treatments, such as plasma etching, as these would

complicate potential scale-up and add cost. For this investigation we used extruded PS and SMA

films.

2. Flexural strength and stiffness predictions

Simple beam theory was used to predict the apparent flexural modulus and strength of the

interleaved material using equations described previously [15]. Room temperature (RT)

predictions (where, RT < Tg of CFRP and thermoplastic) and high temperature (HT) predictions

(where, Tg of thermoplastic < HT < Tg of CFRP) were calculated.

3. Experimental details

3.1 Materials

Pure polystyrene pellets (PS 124N/L, Styrolution, Frankfurt, Germany) were supplied by TCKT

(Wels, Austria). Poly(styrene-co-maleic anhydride) (SMA) pellets (XIRAN SZ 08250 [50 PS]

and XIRAN SZ 15170 [70 PS]) were supplied by Polyscope Polymers (Geleen, The

Netherlands). Unidirectional (UD) prepregs (Hexply 914C-TS-5-34% and Hexply

M21/194/34%/T800S) were purchased from Hexcel (Duxford, UK). The materials used for

specimen preparation are listed in Table 1.

3.2 Film Production

All films were produced using an extruder (Plastik Maschinenbau PM 30, Kelburg, Germany)

equipped with a tapered screw (30 mm diameter, 102 mm long) and a flat film die. The

temperatures within the machine at zones 1, 2 and 3 were 200ºC, 220ºC and 240ºC, respectively.

The temperature at the adapter (where the die is attached to the extruder) was 250ºC and of the

die 260°C. The screw speed was set at 50 rpm. The melt pressure was 30 bar for pure PS and 45-

50 bar for the SMA grades. The take-off speed was 5.2-5.4 m/min. The chill roll temperature was

70°C for pure PS and 90°C for the SMA grades because of their higher film brittleness. The

estimated throughput, calculated from film volume and take-off speed, was 7.2 kg/h. The

thickness of each interleaf film is given in Table 1.

3.3 Wetting behaviour of interleaf films

The wettability of the polymer films was determined by contact angle measurements using a

Drop Shape Analyser (DSA-25, Krüss, Hamburg, Germany). The advancing θa and receding θr

contact angles were measured by placing droplets of distilled water with a volume ranging from

2 μl to 20 μl onto the surface of the polymer films and withdrawing water from the droplets,

respectively, using an automatic syringe at a rate of 6 μl/min.

3.4 Elemental analysis of interleaf films

Elemental analysis was performed using a CHNS-O Elemental Analyser (EA 3000, Eurovector,

Milan, Italy). Sample weights between 0.75 and 3.0 mg were taken for each individual analysis

run. At least three specimens of each sample were tested to calculate the average elemental

composition.

3.5 DSC analysis of interleaf films

Differential scanning calorimetry (DSC) was performed on a Discovery DSC (TA Instruments,

Eschborn, Germany). Samples were heated from 0°C to 300°C, cooled back to 0°C and then re-

heated to 300°C at a heating and cooling rate of 10°C/min. All tests were conducted in a nitrogen

atmosphere.

3.6 DMTA of interleaf films

DMTA was conducted using an RSA-G2 (TA Instruments, Eschborn, Germany). Interleaf films

(40 mm × 5 mm) were tested in tension from 30°C to 140°C at a heating rate of 5°C/min and at a

frequency of 1 Hz.

3.7 Composite production

The interleaved laminates consisted of 8 interleaf layers and 9 CFRP 0° layers (each layer

consisting of 1 ply) arranged in alternating sequence. A pure CFRP 0° control composite

laminate consisting of 16 plies was also manufactured. This was used to compare mechanical

properties of the interleaved laminate and also to provide the elastic modulus and flexural

strength of the composite used for the beam theory predictions. The composite plates were 100

mm × 150 mm with the 0° fibre direction parallel to the 100 mm side. The composite panels

were cured in an autoclave (Quicklock Thermoclave, LBBC, Leeds, UK) at 175°C and 7 bar for

1 h.

3.8 Preparation of DMTA, flexural and short beam shear test specimens

The cured panels were cut into 80 mm × 10 mm specimens for flexural testing, 50 mm × 10 mm

specimens for DMTA and 20 mm × 10 mm for short beam shear testing using a dry diamond

bladed cutter (Tilt Arbour Saw, Startrite, Derbyshire, UK). The 0° direction was parallel to the

longer specimen dimension in all cases.

3.9 Preparation of thick adherend shear test specimens

To manufacture specimens for thick adherend shear tests (TAST), two types of CFRP prepreg

were selected: Hexply 914C and Hexply M21. The M21 prepreg was used as there was a limited

amount of Hexply 914C available. The aim of these tests was to determine the shear properties in

the 914C CFRP/interleaf region, therefore the interleaf films were placed in between layers of

914C. The M21 was used as a sandwich material. Figure 1a shows the composite layup for

TAST. The two prepregs were cured separately and then bonded together with a toughened

epoxy adhesive film (SA 80, Gurit, Isle of Wight, UK).

Four laminates of 914C (150 mm × 300 mm) were manufactured, each comprising of four plies.

Of these four laminates, two formed the outer layers of the TAST specimen and the other two

formed the inner layers, between which the PS (8mm × 300mm) was placed. Two laminates of

M21 (150 mm × 150 mm) were manufactured, each comprising of 26 plies. Two laminates of

914C and one laminate of M21 made up one adherend, 6 mm each. The 914C laminates were

cured as described in section 3.7. The M21 laminates were also cured in the autoclave at 180°C

for 2 h at a pressure of 7 bar then left to cool to 60°C before the pressure was removed.

The surfaces of the laminates were then abraded to improve bonding with the adhesive (SA 80).

The M21 was grit blasted and the 914C was abraded lightly using sandpaper as the laminate

thickness was only 500 μm (sand blasting could have penetrated the laminate). After abrasion the

surfaces were cleaned to remove any dust or impurities. The SA 80 adhesive film was cut and

placed between the laminates as shown in Figure 1a. This sandwich structure was then placed in

hot press at 120°C at 1.4 bar for 30 min to bond the laminates together.

The TAST specimens (Figure 1b, c) were cut to 110 mm × 25 mm using a water jet cutter (Mach

2/2031b, Flow Waterjet, Leicestershire, UK). 12 mm diameter loading holes were drilled (Bench

pillar drill, Meddings, Devon, UK) 80 mm apart into the specimens. Care was taken to ensure the

holes were drilled at the centre to prevent misalignment, which could result in rotation and

uneven loading during testing. Slots were then milled 5 mm apart either side of the specimens,

cutting through the interleaf layer. This enabled testing of the CFRP/interleaf interface.

3.10 DMTA of composites

The composite specimens were tested using an RSA-G2 (TA Instruments, Eschborn, Germany)

in 3-point bending from 30°C to 230°C and at a heating rate of 5°C/min. The frequency was set

at 1Hz.

3.11 Flexural testing of composites

The flexural properties of the composites were determined at room (20°C) and elevated

temperature (140°C) using three-point bending tests in accordance with ASTM D7264-07. The

tests were performed using an Instron 5985 (Bucks, UK) equipped with a 10 kN load cell. A 32:1

span-to-thickness ratio and a crosshead speed of 1 mm/min were used. The loading nose and

supports had a diameter of 6 mm. Specimens were loaded at room temperature (RT1) to a

maximum deflection of 1 mm (at this deflection failure did not occur) and then unloaded. The

tests were then repeated in an environmental chamber (SFL, Eurotherm, West Sussex, UK) at

140°C. Selected specimens were tested again at room temperature (RT2) to determine if the

flexural modulus recovered. Eq. 1 was used to calculate the apparent flexural modulus E f of the

specimens:

E f=L3m

4 b h3 (1)

where L is the support span, b the beam width, h the beam thickness and m the gradient of the

linear portion of the load - displacement curve between zero and 1 mm displacement. Five

specimens of each sample were tested to calculate the average flexural modulus of the

composites.

To determine the flexural strength, previously untested specimens were loaded to failure at room

temperature and 140°C. Eq. 2 was used to calculate the apparent flexural strength:

σ= 3 PL2bh2 (2)

where σ is the maximum stress at the outer surface at the mid-span of the specimen (MPa) and P

the applied force at failure (N). Five specimens were loaded to failure to determine the apparent

flexural strength of each sample.

3.12 Short beam shear testing of composites

Short beam shear tests were carried out in accordance with ASTM D2344-00 to determine the

interlaminar shear strength (ILSS) of the composites. The tests were performed on an Instron

5985 with a 10 kN load cell. A crosshead speed of 1 mm/min was used. The diameter of the

loading nose and supports was 6 mm, which meant using a span-to-thickness ratio of 6:1. ILSS

was calculated using Eq. 3. Five specimens of each sample were tested to determine the average.

ILSS= 3 P4bh (3)

3.13 Thick adherend shear testing of the interleaf/ply interface

Tests were performed on an Instron 5985 (Bucks, UK) equipped with a 250 kN load cell. Tests

were performed at RT and 140°C using an environmental chamber. Tensile load was introduced

via pins that were initially 80 mm apart prior to loading. The test was performed at a speed of 0.5

mm/min. Eq. 4 was used to calculate the average shear stress τ (MPa):

τ12=P

l·w (4)

where P is the applied load (N), l the overlap length (mm) and w the width of the interleaf layer

(mm).

4. Results and Discussion

This investigation focuses on improving adhesion at the ply/interleaf interface in composites

with controllable stiffness to prevent premature shear failure of the composites during flexural

testing. Here we outline the findings and show how this improvement in adhesion affects the

mechanical properties of composites with controllable stiffness.

4.1 Surface and elemental analysis of interleaf films

The advancing and receding water contact angles are shown in Table 2. The pure PS films are

much more hydrophobic than the SMA films as indicated by their higher θa and θr. In general, the

more hydrophilic the polymer films are, the better their adhesion to the epoxy matrix of the

CFRP layers should be [26]. 70 PS was slightly more hydrophilic than 50 PS and this can be

explained as 70 PS contains a higher proportion of MAH than 50 PS (Table 1). The MAH

content of the films was determined by measuring their oxygen content using elemental analysis

(Table 2). The oxygen content of the pure PS, 50 PS and 70 PS films was 0.05%, 3.7% and

7.3%, respectively. The recorded MAH content for the 50 PS and 70 PS films was therefore

1.2% and 2.4%, respectively. The difference in oxygen (MAH) content accounts for the different

wetting behaviour of the films. Unlike the pure PS films, the oxygen containing SMA films can

form hydrogen bonds with water. They will also form chemical and H-bonds with the CFRP

plies (epoxy matrix) in a laminate during cuing of the resin, resulting in better adhesion at the

ply/interleaf interface.

4.2 Thermal properties of interleaf films

The glass transition temperatures of the interleaf films were recorded using DSC. Tg of the pure

PS, 50 PS and 70 PS films were 87°C, 111°C and 127°C, respectively. The difference in T g for

the two SMA films was due to variations in molecular weight and MAH content. The films were

also analysed using DMTA. Similar Tg, determined from the peak in the loss modulus, were

obtained (Table 1). The mechanical Tg of pure PS, 50 PS and 70 PS were 87°C, 112°C and

132°C, respectively.

4.3 Viscoelastic properties of composites

The pure CFRP composite has a single Tg at 194°C that corresponds to that of the epoxy resin

(Figure 2). The storage modulus at 30°C was 93 GPa and the initial stiffness loss occurred at

approximately 140°C. At 30°C the moduli of the interleaved laminates were much lower than

that of the pure CFRP composite due to differences in fibre volume fraction, i.e. the interleaf

films replaced carbon fibre plies. The interleaved laminates were only tested up to 140°C to

minimise the stiffness loss contribution from the CFRP layers softening. The interleaved

composite containing pure polystyrene layers had a storage modulus of 61 GPa at 30°C, which

dropped by 98% to 1 GPa when heated through the Tg of PS to 140°C. The composite containing

50 PS interleaves had a storage modulus of 67 GPa at 30°C. When the composites were heated to

140°C the stiffness dropped by 97% to 2 GPa. The 70 PS-interleaved composite underwent a

60% loss in stiffness from 58 GPa at 30°C to 23 GPa at 140°C. The reduction in stiffness for the

70 PS-interleaved composite was less than that of the other composites as the interleaf layer had

not fully softened at 140°C.

4.4 Flexural properties of composites

At room temperature the pure PS-interleaved specimens had an apparent flexural strength of 382

± 29 MPa (Table 3) and failed in interlaminar shear. The 50 PS and 70 PS-interleaved

composites had flexural strengths at RT of 1107 ± 38 MPa and 1002 ± 17 MPa, respectively.

This increase in strength was due to improved bonding at the interleaf/CFRP interface. These

composites underwent compression failure on the compressive side of the specimen, which led to

interlaminar shear and also showed breaking of the outer fibres on the tensile side. At RT the

control CFRP composite had a flexural strength of 1814 ± 80 MPa and failed in compression at

the mid-span. Flexural strength vs. flexural modulus at RT for each laminate is shown in Figure

3. At HT, the pure PS-interleaved laminate bent under load without failing, however the load

attained was lower than that of the other composites. A flexural strength for this composite could

not be recorded at HT due to limitations with the testing equipment at high specimen deflections.

The 50 PS-interleaved laminate had a flexural strength of 82 ± 2 MPa at 140°C. The composites

showed formation of micro-cracks between the loading nose and supports. The 70 PS-interleaved

laminate had a flexural strength of 187 ± 12 MPa at HT and failed in compression at the mid-

span. Unlike the pure PS-interleaved composite, these laminates failed at 140°C because of the

improved bonding at the ply/interleaf interface that restricted deformation of the interleaf layers.

The flexural modulus of the 16-ply 0° control sample was 119 ± 6 GPa at room temperature. No

stiffness reduction was observed when the specimens were heated to 140°C. The pure PS-

interleaved composites had an apparent flexural modulus of 56 ± 1 GPa at RT, which dropped by

99% to 0.4 ± 0.1 GPa at 140°C. The apparent flexural modulus of the 50 PS composite at RT

was 66 ± 4 GPa. When heated to 140°C there was a 95% reduction in stiffness to 3 ± 0.1 GPa.

The 70 PS composite underwent an 84% loss in stiffness from 64 ± 3 GPa at RT to 10 ± 2 GPa at

140°C. Again, this reduction in stiffness was lower than that of the other composites as the

interleaf layers had not fully softened due to the higher Tg of the interleaf material (Table 1). The

stiffness of selected (un-failed) composites was re-measured at room temperature (RT2) and in

each case the stiffness returned to its full value, indicating that the stiffness loss process is

completely reversible for all composites.

4.5 Predicted vs experimental flexural values

At room temperature the predicted values for the flexural moduli were higher than the

experimental values (Table 3). This may be because shear distortions of the interleaf layers were

not taken into consideration in the predictions, therefore giving an overestimation of the values

(Figure 4). At high temperature, the predicted values were lower than the experimental values. A

possible explanation is that the interleaf layers retained a small portion of their stiffness when

heated. The same explanation goes for the flexural strength (Figure 5).

Despite the difference in thickness between the 50 PS and 70 PS films (100 μm and 170 μm,

respectively), the predicted and measured strengths and moduli of the 50 PS and 70 PS-

interleaved laminates are similar. This was expected as the modulus is mostly dependent on the

stiffness of the CFRP plies. The laminate strength is dependent on the adhesion between the

interleaf layer and the CFRP plies and not on interleaf thickness.

From the bar chart comparisons at HT (Figure 4 and 5) the 70 PS-interleaved laminate had the

highest experimental flexural modulus and strength values. This can be attributed to the Tg of 70

PS being close to the temperature at which the specimens were tested, causing it to have greater

residual stiffness and strength compared to the other interleaf materials. The same, but to a lesser

extent, applies for the difference in flexural properties between the composites containing 50 PS

and pure PS layers.

4.6 Interlaminar shear properties of composites

The ILSS values at room temperature obtained for the pure PS, 50 PS and 70 PS-interleaved

laminates were 16 ± 7 MPa, 57 ± 5 MPa and 54 ± 4 MPa, respectively. The ILSS for the pure

CFRP composite was 94 ± 4 MPa. At RT the composites containing pure PS debonded at the

PS/CFRP interface while the 50 PS and 70 PS-interleaved laminates failed in a similar manner as

the pure CFRP specimen, i.e. they fractured on the compression side leading to interlaminar

shear. Images of SBS specimens after failure at RT show complete separation of the PS interleaf

from one side of the CFRP layer (Figure 6a) while for the composite containing 50 PS (Figure

6b), fracture within the interlayer layer in some regions was observed. For the 70 PS-interleaved

laminate (Figure 6c), microcracks in the interleaf layer occurred with a thin layer of the 70 PS

interleaf film remaining on the surface of both CFRP plies. For a good bond, fracture in the

interleaf layer is desired, such that the interleaf films remains bonded to both neighbouring

CFRP plies. At HT, failure of the composites was not recorded, however pure PS was seen

flowing out of the specimens during the test. This was attributed to excessive compressive force

and the fact that the test temperature was a lot higher than the Tg of PS, thereby causing it to

soften more than the SMA films.

4.7 Shear properties of composites from thick adherend shear tests

At room temperature the shear strength of the pure PS-interleaved composite was 12 ± 4 MPa

(Table 4). The laminates containing 50 PS and 70 PS interleaves had much higher shear strengths

of 28 ± 12 MPa and 28 ± 2 MPa, respectively. The ILSS values from short beam shear tests were

higher than the shear strength values from TAST. Short beam shear test can give artificially high

values of ILSS due to the compressive forces that exist in this test [27, 28]. At 140°C the shear

strength of the pure PS-interleaved composite dropped to 0.04 ± 0.03 MPa due to softening of

the PS layers. The 50 PS and 70 PS-interleaved laminates had shear strengths at 140°C of 1.5 ±

0.5 MPa and 2.1 ± 0.2 MPa, respectively. The SMA interleaved composites had significantly

higher shear strengths at both room and elevated temperature due to improved bonding at the

CFRP/interleaf interface.

5. Conclusion

The interleaf/CFRP interface in carbon fibre reinforced epoxy composites with controllable

stiffness was investigated. Premature failure occurred when polystyrene interleaved composites

were tested at room temperature in three point bending. This was due to poor bonding at the

polystyrene/CFRP interface. This issue was resolved through the incorporation of more

hydrophilic poly(styrene-co-maleic anhydride) interleaf films. Laminates were tested at room

and elevated temperature (140°C) using flexural, short beam shear and thick adherend shear tests

to verify the improvement in the interleaf/CFRP bond. Laminates containing SMA interleaf

layers had significantly higher adhesion, flexural and shear strengths than composites containing

pure PS layers. Images of the fractured specimens from short beam shear tests show a change in

failure mode confirming an improvement in bonding.

Acknowledgements

We thank the University of Vienna for funding this research, Klaus Göschl (UniVie, 2014) for

the contact angle measurements and Johannes Theiner (UniVie, 2015) for performing elemental

analysis.

References

1. Kuder IK, Arrieta AF, Raither WE, Ermanni P. Variable stiffness materials and structural

concepts for morphing applications. Progress in Aerospace Sciences. 2013;63:33-55.

2. Thill C, Etches J, Bond I, Potter K, Weaver P. Morphing skins. The Aeronautics Journal.

2008;112:117-39.

3. Sofla AYN, Meguid SA, Tan KT, Yeo WK. Shape morphing of aircraft wing: Status and

challenges. Materials and Design. 2010;31:1284-92.

4. Tridech C. Smart fibre coatings: stiffness control in composite structures. Ph.D. Thesis. Imperial

College London; 2011.

5. Tridech C, Maples HA, Robinson P, Bismarck A. High performance composites with active

stiffness control. ACS Applied Materials and Interfaces. 2013;5:9111-19.

6. Bachinger A, Marklund E, Rössler J, Hellström P, Asp LE. Stiffness-modifiable composite for

pedestrian protection, ECCM16, Seville, Spain; June 2014.

7. Bachinger A, Marklund E, André A, Hellström P, Rössler J, Asp LE. Materials with variable

stiffness, ICCM20, Copenhagen, Denmark; July 2015.

8. McKnight G, Henry C. Variable stiffness materials for reconfigurable surface applications,

Smart Structures and Materials 2005: Active Materials: Behavior and Mechanics, San Diego,

US; March 2005.

9. Shan Y, Philen M, Lotfi A, Li S, Bakis CE, Rahn CD, Wang KW. Variable stiffness structures

utilizing fluidic flexible matrix composites. Journal of Intelligent Material Systems and

Structures. 2009;20:443-56.

10. Raither W, Bergamini A, Gandhi F, Ermanni, P. Adaptive bending-twist coupling in laminated

composite plates by controllable shear stress transfer. Composites Part A: Applied Science and

Manufacturing. 2012;43:1709-16.

11. Gandhi F, Kang SG. Beams with controllable flexural stiffness. Smart Materials and Structures.

2007;16:1179-84.

12. Murray G, Gandhi F. Multi-layered controllable stiffness beams for morphing: energy, actuation

force, and material strain considerations. Smart Materials and Structures. 2010;19:045002.

13. Gandhi F, Murray G, Kang SG. Flexural stiffness control of multilayered beams. AIAA Journal.

2009;47:757-66.

14. Henry C, McKnight G. Deformable variable-stiffness cellular structures, Patent US 7678440 B1;

2010.

15. Maples HA, Wakefield S, Robinson P, Bismarck A. High performance carbon fibre reinforced

epoxy composites with controllable stiffness. Composites Science and Technology.

2014;105:134-43.

16. Robinson P, Bismarck A, Zhang B, Maples HA. Exploring the potential of controllable stiffness

hybrid composites, ECCM16, Seville, Spain; June 2014.

17. Maples HA. Composites with controllable stiffness. Ph.D. Thesis. Imperial College London;

2014.

18. Heness G, Stevens M, Dossett B. The effect of surface treatment on delamination for a nylon

interleaving material. Material Forum. 2007;31:90-95.

19. Nay JC, Pitt WG, Armstrong-Carroll E. Improving adhesion in interleaf composites using

plasma processing. Journal of Applied Polymer Science. 1995;56:461-69.

20. Riefler RS, Powers JW. Interleaf layers in fiber reinforced resin laminate composite, Patent US

4954382 A, September, 1990.

21. Hojo M, Matsuda S, Ochiai S, Murakami A, Akimoto H. The role of interleaf/base lamina

interphase in toughening mechanisms of interleaf toughened CFRP, ICCM12, Paris, France; July

1999.

22. Wolf R, Sparavigna AC. Role of Plasma Surface Treatments on Wetting and Adhesion.

Engineering. 2010;2:397-402.

23. Friedman M, Walsh G. High performance films: Review of new materials and trends. Polymer

Engineering and Science. 2002;42:1756–88.

24. Keener TJ, Stuart RK, Brown TK. Maleated coupling agents for natural fibre composites.

Composites Part A: Applied Science and Manufacturing. 2004;35:357–62.

25. Rijsdijk HA, Contant M, Peijs AAJM. Continuous-glass-fibre-reinforced polypropylene

composites: I. Influence of maleic-anhydride-modified polypropylene on mechanical properties,

Composites Science and Technology. 1993;48:161-72.

26. Cantrell JH. Hydrogen bonds, interfacial stiffness moduli, and the interlaminar shear strength of

carbon fiber-epoxy matrix composites. AIP Advances. 2015;5:037125.

27. Turner S. General principles and perspectives. In: Hodgkinson JM, editor. Mechanical testing of

advanced fibre composites. Cambridge: Woodhouse Publishing; 2000. p. 27.

28. Feraboli PJ, Kedward KT. In search of the true interlaminar shear strength. In: Nineteenth

technical conference of the joint American society for composites/American society for testing

and materials committee D30, Atlanta, USA; October 2004.

Table 1: Materials used for specimen preparation.

Material Tg/°C† Thickness/μm MAH content/%*

Hexply 914C-TS-5-34% (UD CFRP) 194 ± 3 125 -Hexply M21/194/34%/T800S (UD CFRP) 195‡ 184 -PS 124N/L (Pure PS) 87 ± 2 100 -XIRAN SZ 08250 (50 PS) 112 ± 3 100 1.2XIRAN SZ 15170 (70 PS) 132 ± 6 170 2.4

†Tg determined using DMTA‡Tg from material data sheet*Maleic Anhydride (MAH) content according to elemental analysis

Table 2: Contact angles and elemental analysis of interleaf films.

Contact angle Elemental analysis

Interleaf film θa†/° θr

†/° C (wt.%)

H (wt.%)

N (wt.%)

S (wt.%)

O (wt.%)

PS 124N/L (Pure PS) 93 ± 1 71 ± 1 91.80 8.13 <0.05 <0.02 0.05XIRAN SZ 08250 (50 PS) 79 ± 1 43 ± 1 88.47 7.41 <0.05 <0.02 3.70XIRAN SZ 15170 (70 PS) 77 ± 1 42 ± 1 85.13 7.01 <0.05 <0.02 7.33

†Advancing (θa) and receding (θr) contact angles

Table 3: Flexural properties and ILSS strengths of interleaved laminates.

Composite Apparent flexural strength/ MPa Apparent flexural modulus/ GPa ILSS†/ MPaRT1† 140°C % loss RT1 140°C % loss

16-ply 0° control Measured 1814 ± 80 714 ± 173 61 119 ± 6 121 ± 3 - 94 ± 4

Pure PS Measured 382 ± 29 - - 56 ± 1 0.4 ± 0.1 99 16 ± 7Predicted 1099.5 21.1 98.1 72.2 0.3 99.6

50 PS Measured 1107 ± 38 82 ± 2 93 66 ± 4 3 ± 0.1 95 57 ± 5Predicted 1177.8 24.7 97.9 77.2 0.3 99.6

70 PS Measured 1002 ± 17 187 ± 12 81 64 ± 3 10 ± 2 84 54 ± 4Predicted 1144.2 23.0 97.9 75.2 0.3 99.6†Room temperature (RT)‡Interlaminar shear strength (ILSS), tests performed at room temperature

Table 4: Shear strength measured using thick adherend shear tests.

Composite Shear strength/MPaRT† 140°C

Pure PS 12 ± 4 0.04 ± 0.0350 PS 28 ± 12 1.5 ± 0.570 PS 28 ± 2 2.1 ± 0.2

†Room temperature (RT)

(a) (b) (c)Figure 1: Composite layup for TAST (a) and geometry of the TAST composite specimens [top

view (b), side view (c)].

SA 80

Interleaf layer

M21 (26 plies)

Key

914c (4 plies)

42.5

51

6.55

25

110

12

Loadingholes

12

(a) (b)

(c) (d)Figure 2: DMTA curves showing the storage modulus (E′), loss modulus (E″) and tan δ as a function of temperature for (a) pure CFRP (b) pure PS (c) 50 PS and (d) 70 PS-interleaved

laminates.

Figure 3: Apparent flexural modulus vs. apparent flexural strength of the interleaved composites at room temperature.

50 100 150 2000

20

40

60

80

100

120

140E/

GPa

Temperature/ °C

0.00

0.02

0.04

0.06tan

E''

E'

Tan

40 60 80 100 120 1400

10

20

30

40

50

60

70

E/ G

Pa

Temperature/ °C

0

1

2

Tan

E'

tan

E''

40 60 80 100 120 140 160 1800

10

20

30

40

50

60

70

E/ G

Pa

Temperature/ °C

E'

tan

E''0

1

2

Tan

40 60 80 100 120 140 160 1800

10

20

30

40

50

60

E/ G

Pa

Temperature/ °C

E'

E''

tan

0.0

0.5

1.0

Tan

0 500 1000 1500 200040

60

80

100

120 16-ply 0° control Pure PS 50 PS 70 PS

App

aren

t Fle

xura

l Mod

ulus

/ GPa

Apparent Flexural Strength/ MPa

Figure 4: Comparison of measured and predicted flexural modulus values for interleaved

composites.

Figure 5: Comparison of measured and predicted flexural strength values for interleaved

composites.

(a) (b) (c)

Figure 6: Images of pure PS (a), 50 PS (b) and 70 PS-interleaved laminates (c) tested to failure at

room temperature in short beam shear.

RT 1 measured RT 1 predicted0

10

20

30

40

50

60

70

80

90

100Fl

exur

al M

odul

us/ G

Pa Pure PS 50 PS 70 PS

140°C measured 140°C predicted 0

2

4

6

8

10

12

Flex

ural

Mod

ulus

/ GPa

Measured Predicted0

20

40

60

80

100

% st

iffne

ss lo

ss fr

om R

T1 to

140

°C

RT 1 measured RT 1 predicted0

200

400

600

800

1000

1200

1400

Flex

ural

Stre

ngth

/ MPa

Pure PS 50 PS 70 PS

140°C measured 140°C predicted 0

20

40

60

80

100

120

140

160

180

200

Flex

ural

Stre

ngth

/ MPa

Measured Predicted0

20

40

60

80

100

% s

treng

th lo

ss fr

om R

T1 to

140

°C