D A Papargyris
email: d.papargyris@postgrad.manchester.ac.uk

R J Day

A Nesbitt

D Bakavos

physical properties of a carbon fibre epoxy composite manufactured

Keywords: Microwave curing, E. Resin transfer moulding, A. Polymer matrix composites, B. Interfacial strength, B. Mechanical properties

Comparison of the mechanical and

Introduction

Fibre-reinforced polymer composites have gained substantial interest over the last years, mainly due to their very high strength-to-weight and stiffness-to-weight ratios. These properties led to the use and application of composite materials in a lot of major industries with the aerospace and automotive industry taking the larger part. The high operational costs involved in combination with the intricacy of the manufacturing techniques currently employed have restricted wider industrial use of composites. For that reason, considerable effort has been made in the direction of finding and developing alternative costeffective routes for curing composite materials.

One such alternative route is microwave processing, which offers several advantages over the conventional thermal processing methods, including rapid, selective and volumetric heating, energy savings, reduced processing time and improved processing control [START_REF] Mijovic | Review of cure of polymers and composites by microwave energy[END_REF][START_REF] Zhou | Research on the technology and the mechanical properties of the microwave processing of polymer[END_REF][START_REF] Ku | Productivity improvement through the use of industrial microwave technologies[END_REF][START_REF] Bykov | High-temperature microwave processing of materials[END_REF]. Microwave radiation generates heating within the material rather than relying on heat transfer through conduction and convection. It is, thus, more appropriate to consider microwave heating as conversion of electromagnetic energy to thermal energy rather than heat transfer [START_REF] Thostenson | Microwave processing: fundamentals and applications[END_REF]. A considerable amount of research has been devoted on microwave curing of different polymer [START_REF] Zhou | Research on the technology and the mechanical properties of the microwave processing of polymer[END_REF][START_REF] Wei | Comparison of Microwave and Thermal Cure of Epoxy resins[END_REF][START_REF] Bai | A comparative study of the mechanical behaviour of an epoxy resin cured by microwaves with one cured thermally[END_REF][START_REF] Mijovic | Comparative calorimetric study of epoxy cure by microwave vs thermal energy[END_REF][START_REF] Marand | Comparison of reaction mechanisms of epoxy resins undergoing thermal and microwave cure from in situ measurements of microwave dielectric properties and infrared spectroscopy[END_REF][START_REF] Fu | Comparative Study of Continuous-Power and Pulsed-Power Microwave Curing of Epoxy Resins[END_REF][START_REF] Boey | Microwave curing of an epoxy-amine system: effect of curing agent on the glass-transition temperature[END_REF][START_REF] Rahmat | Comparison of microwave and thermal cure of unsaturated polyester resin[END_REF][START_REF] Ismail | Comparison between the thermal and microwave curing of bismaleimide resin[END_REF] and composite systems [START_REF] Boey | Microwave radiation curing of a thermosetting composite[END_REF][START_REF] Yue | Influence of thermal and microwave processing on the mechanical and interfacial properties of a glass/epoxy composite[END_REF][START_REF] Nightingale | Flexural and interlaminar shear strength properties of carbon fibre/epoxy composites cured thermally and with microwave radiation[END_REF][START_REF] Qiu | Composite Material Processing in a Single Mode Cavity with Variable Frequency Microwaves[END_REF][START_REF] Rao | Studies on Tensile and Interlaminar Shear Strength Properties of Thermally Cured and Microwave Cured Glass-Epoxy Composites[END_REF][START_REF] Varaporn | Comparison of microwave and thermal cure of epoxy-anhydride resins: Mechanical properties and dynamic characteristics[END_REF][START_REF] Varaporn | Comparison between microwave and thermal curing of glass fiberepoxy composites: Effect of microwave-heating cycle on mechanical properties[END_REF][START_REF] Thostenson | Microwave and conventional curing of thick-section thermoset composite laminates: Experiment and simulation[END_REF][START_REF] Boey | Electromagnetic radiation curing of an epoxy/fibre glass reinforced composite[END_REF][START_REF] Mooteri | Studies on mechanical behavior of microwave and thermally cured glass fiber reinforced polymer composites[END_REF]. The implementation of microwave heating into common manufacturing methods for polymer composites, such as resin transfer moulding [START_REF] Johnson | Microwave assisted resin transfer moulding[END_REF], however, has not yet been fully exploited.

The aim of the present work is to incorporate microwave heating into the resin transfer moulding technique and compare the overall physical and mechanical properties of the fabricated composites with those obtained by conventional thermal resin transfer moulding, in an attempt to assess the feasibility and efficiency of microwave radiation as an alternative power source for curing composites as well as determine any repercussions it might have on the attained composite properties and structure.

Experimental

Materials

The resin system used in this study was Araldite LY5052 / Aradur HY5052 (supplied by Huntsman). LY5052 is a low-viscosity epoxidised novolak resin and Aradur 5052 is a hardener comprising a mixture of polyamines. The resin/hardener mix ratio was 100:38 parts by weight. The reinforcement used was high strength T300 carbon fibre fabric (supplied by Hexcel Composites) with balanced 50/50 warp/weft five-harness satin weave architecture. Seven layers of carbon fibre reinforcement were used for each laminate panel.

Resin transfer moulding (RTM) processing

Conventional thermal RTM processing

Conventional thermal RTM processing was based on the Hypaject Mark II RTM Injection system (Plastech), as illustrated schematically in Fig. 1. The system comprised an inlet valve, a heated homogenizer, a pneumatic valve, a metal mould with 16 electric-cartridge heaters (8 for each mould part) giving a total power output of 4 KW, two K-type thermocouples and two PID temperature controllers.

Prior to the curing process, the metal mould was coated with Frekote 44-NC release agent. The fibres were then placed inside the mould cavity (200 x 300 x 3 mm dimensions) and the mould was inclined at 45 0 from horizontal in a steel stand in order to minimize air entrapment during the injection.

The resin was drawn into the homogenizer through the inlet valve and left for 10-15 minutes under vacuum in order to remove any air bubbles created at the initial resin-hardener mixing stage. The resin was then injected under low pressure (0.5 bar) into the mould through the inlet valve. After completion of the injection stage, the resin-loaded mould was heated for 3 hours at 100°C via the electriccartridge heaters, with the temperature controllers maintaining the desired cure temperature.

Microwave RTM processing

Microwave RTM processing differed from conventional thermal RTM processing in how the resin-loaded mould was heated after the injection stage. Instead of traditional thermal heating, microwave radiation was employed in order to cure the composites. Metal is known to reflect microwave energy, thus a moulding tool suitable for use inside the microwave field was manufactured from Macor. Macor is a machinable glass ceramic with good dielectric properties, exhibiting at room temperature a loss tangent and a dielectric loss of 7.1x10 -3 and 4x10 -2 at 8.5 GHz, and 4.7x10 -3 and 2.8x10 -2 at 1 KHz, respectively (as given by supplier, Precision Ceramics). These properties allowed microwave energy to be absorbed predominantly by the resin and not the mould material, whilst providing the required dimensional stability to withstand the temperature and pressure cycles during processing. The cavity dimensions of the Macor mould was 200 x 200 x 3 mm. Preparation of the Macor mould prior to resin injection involved applying Chemlease C15 sealer agent and PMR Chemlease release agent. The fibres were then placed inside the cavity and the mould was again inclined at 45 0 from horizontal during resin injection, which was carried out using the Hypaject RTM system described in 2.2.1.

The microwave RTM heating equipment comprised a network analyser (NA) (Hewlett Packard HP8720ET) used as the power source, a travelling wave tube (TWT) amplifier, an isolator, a multi-mode microwave applicator (brass cavity), a Macor RTM mould and 4 fluoroptic probes (Luxtron Corp.) connected to a Luxtron 790 fluoroptic thermometer. The microwave RTM system set-up is illustrated in Fig. 2. The microwave signal from the NA was amplified by the TWT amplifier and then fed to the microwave applicator via an isolator in order to protect the TWT amplifier against any reflected power from the applicator. The maximum amplifier output power that could be used in this configuration was 250 W. A computer control system was employed to ensure both a uniform heating and maintain a stable cure temperature. Microwave power, excitation frequency and sample temperature were used as the process control parameters. Temperature stability was achieved by adjusting the microwave power while uniformity of the electric field distribution within the applicator was provided by varying the excitation frequency. Ten different excitation frequencies ranging from 4-8 GHz were used for processing. These frequencies were dictated by the microwave response of the applicator loaded with the Macor moulding tool, resin and carbon fibre reinforcement just before curing and were selected in order to minimise the reflected power.

Microwave power and frequency were adjusted via a General Purpose Interface Bus (GPIB) connection between the computer and the NA. The computer control system was designed in such way that the NA would continuously step through the excitation frequencies, while simultaneously regulating the microwave power in order to maintain the desired cure temperature. The sample temperature was constantly monitored by the computer via a serial connection (RS-232) with the fluoroptic thermometer. The computer control system regulated the microwave power according to the following quadratic algorithm:

2 min - - ∝ T T T T Power cure cure (1)
where, T was the maximum of the measured sample temperatures, T cure was the cure temperature and T min was a pre-set temperature limit. As can be seen from Eq.1, the amplifier output power used was proportional to the square of the difference between the set cure temperature and the actual temperature.

At temperatures below T min full power was applied, whereas when the temperature exceeded T cure the power was turned off.

Mechanical testing 2.3.1. Flexural modulus and flexural strength

The flexural properties of the composite samples were determined according to ASTM Standard D 6272 [START_REF] Astm D | Standard test method for flexural properties of unreinforced and reinforced plastics and electrical insulating materials by four-point bending[END_REF]. The apparatus used was an Instron 4301 equipped with a four-point bend jig and a 5 KN load cell. For each composite panel six samples were tested and an average value was taken.

Interlaminar shear strength

Interlaminar shear testing of the composite samples was conducted on an Instron 4301 equipped with a three-point bend jig and a 5 KN load cell, according to ASTM Standard D 2344/D 2344M [START_REF] Astm D | D 2344M Standard test method for short-beam strength of polymer matrix composite materials and their laminates[END_REF]. For each composite panel eight samples were tested and an average value was taken.

Fibre and void volume fraction

The percentage fibre volume fraction V r and the percentage void volume fraction V v were determined using matrix digestion with nitric acid in a microwave oven according to ASTM Standard D 3171 [START_REF] Astm D | Standard test method for constituent content of composite materials[END_REF]. Prior to that, the density of the samples was determined according to Airbus UK Test Specification [START_REF]Airbus UK Test Specification[END_REF]. For each panel ten samples were tested and an average value was taken.

Rheology

The viscosity of the neat LY/HY5052 resin system was determined as a function of time using a RMS 800 rheometer. The neat resin was heated between two parallel 50 mm aluminium plates under a 1

Hz oscillating frequency at 2°C/min, 5°C/min and 10°C/min from room temperature until gelation occurred.

Differential scanning calorimetry (DSC)

Differential scanning calorimetric studies were carried out on a TA Instruments MDSC 2920 to ensure that the fabricated composite panels were fully cured. Small size samples (< 10 mg) were heated from 20°C up to 250°C at 10°C/min in an inert nitrogen atmosphere. Samples were taken from different locations on each composite panel.

Dynamic mechanical thermal analysis (DMTA)

Dynamic mechanical thermal analysis was carried out on a TA Instruments DMA Q800 to obtain the glass transition temperatures for both conventional thermal and microwave cured composites.

The composite samples (60 x 8 x 3 mm) were heated from -120°C up to 250°C at 2°C/min while the frequency of the applied oscillating stress was 1 Hz. For each composite panel six samples were tested.

Scanning electron microscopy (SEM)

The composite samples were examined using Philips XL30 FEG SEM microscope. Small sections of the fracture surfaces were cut from test samples subjected to mechanical testing and placed on metal stubs using double-sided carbon tabs. The stubs with the fracture surfaces were then coated with a very thin layer of gold using Edwards S150B sputter coater in order to make them electrically conductive.

Results and discussion

Based on previous work by Yusoff [START_REF] Yusoff | Microwave assisted RTM processing of carbon/epoxy composites[END_REF], the cure cycle employed for conventional thermal processing was 100°C for 3 hours. DSC results confirmed that the composite panels fabricated using this profile were fully cured. Determination of the cure cycle for microwave processing was based on the dielectric properties of the neat resin, as they were measured according to the cavity perturbation technique [START_REF] Nesbitt | Development of a microwave calorimeter for simultaneous thermal analysis, spectroscopy and dielectric measurements[END_REF]. Fig. 3 shows the evolution of the dielectric loss and dielectric constant of the neat resin at 100°C and 2.45 GHz as a function of time.

The measurement of the dielectric properties was carried out at a set frequency of 2.45 GHz and without taking into consideration the effect of the carbon fibre reinforcement. However, these properties should be indicative of the dielectric response of the composite system at the desired curing temperature.

Given that the absorbed microwave power is proportional to the dielectric loss of the material inside the microwave field, it can be assumed that the curing reaction has concluded once the dielectric loss reaches a constant value over time. As depicted in Fig. 3, the dielectric loss stabilised after about 90 minutes, thus the cure cycle employed for microwave processing was 90 minutes at 100°C. DSC analysis of the composite panels produced by microwave RTM under the selected cure cycle showed that the samples were fully cured, indicating sufficient cure cycle time as well as uniform microwave field distribution inside the applicator; otherwise localized hot spots within the field would have resulted in composite panels with inferior mechanical and physical properties, owing to uneven degrees of cure.

This 50% cure cycle time reduction between conventional and microwave RTM processing was achieved due to the different nature of interaction between the resin and microwave radiation, which in turn yields a different heating mechanism compared to conventional thermal heating. Generally, microwave heating is regarded as a volumetric heating technique in which the microwaves interact with the material and heat is generated directly as a result. Hence heating is efficient, because energy is deposited directly into the material and the effects of convective and conduction heat transfer are less significant than with some other heating techniques. In carbon fibre composites, the situation is probably more complicated as the skin depth is very small. Calculations based on the dielectric data measured for Hexply 6376 by Atkinson [START_REF] Atkinson | Investigation of potential methods for improved processing of carbon/epoxy composites[END_REF] show that the skin depth at the frequencies used in the present work is between 0.5 and 2.2 cm. While the resin and fibre system used by Atkinson were different it is likely that the apparent dielectric properties were dominated by the conductivity of the carbon fibre and that the skin depth for the system used in this study is similar. Therefore, the microwaves do not penetrate a long way into the composite at all frequencies. They interact instead with a surface layer. This means that there is still good heat transfer into the surface of the composite. Within the composite it is likely that the microwave electric field generates an electric current in the carbon fibres and thus resistive heating. Resin will be heated through dielectric heating at the surface of the composite. Within the bulk it will be heated due to the resistive heating of the fibres and through thermal conduction. The cure cycle required for microwave heating in this work was half than that needed for conventional thermal heating, suggesting that despite the likely skin depth heat transfer through microwave heating is still faster than through conventional heating.

The key question, therefore, is whether this substantial cutback in the cure cycle time of the microwave cured composites would have any detrimental effect on their overall mechanical and physical properties. Table 1 shows the average values of flexural modulus, flexural strength, interlaminar shear strength, fibre volume fraction, void content and panel thickness obtained for composite panels manufactured by both conventional thermal and microwave RTM processing. The average panel thickness of the microwave cured composite panels was slightly higher compared to those cured thermally due to small dimensional differences of the Macor mould cavity attained at machining. For that reason, given the same number of carbon fibre layers used to manufacture composites with both methods, the microwave cured samples exhibited slightly lower fibre volume fraction, as they contained more resin per unit volume.

In order to compare the flexural mechanical properties of the manufactured composites, the mechanical data for the microwave RTM composites were normalised to an equivalent fibre volume fraction of 33%. The normalised values are tabulated in Table 1. The normalisation was carried out assuming a linear relationship between fibre volume fraction and flexural strength and modulus. The flexural moduli of both thermally and microwave cured composites were found to be comparable. A small increase in the flexural strength of the microwave cured composites was observed. Nevertheless, this increase is within the experimental error involved in measuring; no distinctive improvement in the flexural strength of the microwave cured samples can be concluded, with both types of composites exhibiting equivalent flexural properties. The value, however, of the interlaminar shear strength obtained for the composites fabricated via microwave RTM yields a very interesting result. As can be seen in Table 1, the interlaminar shear strength for the microwave cured samples was found to be 9% higher than those cured thermally, possibly signifying that microwave processing improves the interfacial bonding between epoxy resin and carbon fibre reinforcement. This is in agreement with other work by Fang [START_REF] Fang | Investigation of microwave energy to cure carbon fiber reinforced phenylethynyl-terminated polyimide composites, PETI-5/IM7[END_REF] and Rao [START_REF] Rao | Studies on Tensile and Interlaminar Shear Strength Properties of Thermally Cured and Microwave Cured Glass-Epoxy Composites[END_REF].

The enhancement in the interlaminar shear strength can be ascribed to a lowering in the resin viscosity at the initial curing stage of microwave processing, before the cross-linking density starts to rise rapidly and the resin becomes very viscous; in which case no significant change in the composite physical properties could occur. Previous work suggests that by lowering the resin viscosity, improved resin-fibre adhesion and fibre wet-out can be achieved [START_REF] Huang | Influence of ultrasonic treatment on the characteristics of epoxy resin and the interfacial property of its carbon fiber composites[END_REF][START_REF] Liu | Effect of ultrasound on wettability between aramid fibers and epoxy resin[END_REF][START_REF] Davies | Effect of cure cycle heat transfer rates on the physical and mechanical properties of an epoxy matrix composite[END_REF]. Fig. 4 shows the viscosity change against time at 3 different heating rates (2°C/min, 5°C/min and 10°C/min) for the neat LY/HY5052 resin system. As the heating rate increased the minimum viscosity of the neat resin decreased considerably from 0.12 Pa-s to 0.068 Pa-s and finally to 0.043 Pa-s. Microwave heating is known to be rapid and volumetric, generating heat within the material, mainly due to ionic polarization and dipole orientation [START_REF] Mijovic | Review of cure of polymers and composites by microwave energy[END_REF][START_REF] Decareau | Microwave Processing and Engineering[END_REF], rather than relying on heat transfer through conduction or convection, although the effect in carbon fibre composites is more complicated, as noted earlier. A higher heating rate would result in a decrease in the resin viscosity, which in turn would facilitate resin flow, making it easier for the resin to impregnate the carbon fibre surface during the initial stage of the curing process. A lower viscosity would improve the interfacial bonding and enhance the mechanical interlocking between the resin and the fibres. The increase in the interlaminar shear strength could be also partly attributed to extended plastic deformation and hence increased crack-tip blunting, as a result of the greater amount of resin between the fibre layers in the microwave cured samples compared to those cured thermally, given the lower fibre volume fraction of the former. In addition, the improvement in the resin flow has a direct effect on the void content obtained for the microwave cured composites, as shown from the low void content value of 0.6% (Table 1). The fluidity of the resin is thought to provide better contact between the resin and the fibre surface contour, allowing the resin to effectively enter any formed micro-holes which where previously occupied by trapped air. However, panels produced conventionally also yielded an acceptable low void content of 1.7%.

Typical DMTA curves at 2°C/min scan rate and 1 Hz frequency for the thermally and microwave cured composites are illustrated in Figs. 5 and6. Regarding the average peak values of the higher temperature -transition, no significant differences existed for both types of composites, with the microwave cured samples exhibiting a slightly higher glass transition temperature (T g ) of 130 ± 1 °C, as opposed to 128 ± 2 °C for those cured thermally. However, a 15°C difference in the average peak values of the lower temperature -transition was observed between the two composite types (-72 ± 3 °C and -57 The fracture surfaces of composite samples subjected to four-point bending and interlaminar shear testing were also examined using scanning electron microscopy. Representative topographic features of both types of composites are shown in Figs. 789. Careful examination of the fracture surface morphology of thermally and microwave cured samples revealed some differences between the two types of composites. It can be seen from Figs. 7a and7b, that the fracture surface of the composites produced with the conventional method exhibits poorer interfacial bonding and more extensive fibre pull-out compared to that produced using microwaves. This is more obvious in Figs. 8a and8b, where there is a significant amount of clean fibres present as opposed to fibres thoroughly coated with resin, for the thermally and microwave cured composites respectively. Brittle matrix failure regions were also identified for both types of composites, as shown in Figs. 9a and9b. However, due to the complexity of woven fabric composites compared to unidirectional composites, the interlaminar fracture behaviour under Mode I, Mode II and mixed Mode I and II loading could also be investigated in order to have a more in depth understanding of the main failure modes.

Conclusions

A modification of the conventional RTM method incorporating the use of microwave heating was presented. A ceramic moulding tool with dielectric properties suitable for use inside the microwave field was manufactured. The method was based on the Hypaject Mark II RTM Injection system. The microwave heating equipment comprised a network analyser, a travelling wave tube amplifier, an isolator, a multi-mode applicator and optical temperature sensing. Composites were produced by both conventional thermal and microwave RTM processing in order to compare their physical and mechanical properties.

The curing profile employed in conventional RTM processing was 3 hours at 100°C as opposed to 90 minutes at 100°C in microwave RTM processing, resulting to a 50% cure cycle time reduction through the use of microwave heating. DSC analysis of samples taken from different panel locations confirmed that the microwave cured composites were fully cured, validating the uniformity of the microwave field inside the applicator. Mechanical testing showed similar values of flexural moduli obtained for the two types of composites. A small increase in the flexural strength of the microwave cured composites was observed. However, this increase is not profound as it falls within the experimental error.

The interlaminar shear strength of the composites produced by microwave RTM processing was found to be 9 % higher than that of the thermally cured composites. This enhancement is suggested to be mainly attributed to a lowering of resin viscosity at the initial stage of the curing process due to a higher heating rate, which in turn facilitated the resin flow and improved the interfacial bonding between the resin and the fibre surface. Examination of the composite fracture surface morphology, as illustrated by a series of SEM micrographs of specimens failed after mechanical testing, provided additional evidence of improved interfacial bonding in terms of increased fibre wetting, less fibre pull-out, and better coating of the fibres with resin for the microwave cured specimens compared to those cured thermally. Both microwave and thermally cured composites yielded minimal void contents (<2%), as indicated by the results from the hot acid digestion method. The microwave cured composites exhibited lower void content values than those cured conventionally. DMTA analysis of composites produced by both methods revealed similar glass transition temperatures (T g ) for the higher temperature -transition (128 ± 2 °C and 130 ± 1 °C for the conventionally and microwave cured composites respectively). However, a significant shift of about 15°C in the position of the -transition peak was observed between the two types of composites, probably due to an alteration in the cross-linking path followed.

The results presented in this work underline the potential of microwave heating in improving composite processing and manufacture. Despite the half cure cycle time employed, the mechanical and physical properties of the microwave cured composites were found to be similar and, in some cases, superior compared to their thermally cured counterparts. However, further research is necessary in order to understand and explain the differences in the reaction mechanism between thermally and microwave cured composites. Loss Modulus (MPa) Tan Delta

± 3

 3 °C for the thermally and microwave cured composites respectively). The low -transition in epoxy systems has been observed between -80°C to -40°C [37-39]. More specific in epoxy-amine curing, lowtemperature relaxations are mainly associated with the molecular motion of the hydroxyether portion of the molecule [40]. A possible difference in the molecular structure of the polymer chain could, hence, account for the shift in the position of the -transition peak, which in turn might be caused by an alteration in the cross-linking path followed between the microwave and the thermally produced composites.

Fig. 1 .

 1 Fig.1. Schematic of conventional thermal RTM processing (solid arrows indicate resin flow, dotted arrows indicate process control).

Fig. 2 .

 2 Fig.2. Schematic of microwave RTM processing (solid arrows indicate microwave power flow, dotted arrows indicate process control / temperature measurements).

Fig. 3 .

 3 Fig.3. Dielectric constant and dielectric loss of neat LY/HY5052 resin system measured at 100°C and 2.45GHz as a function of time.

Fig. 4 .

 4 Fig.4. Viscosity versus time at 2°C/min, 5°C/min and 10°C/min for neat LY/HY5052 resin system.

Fig. 5 .

 5 Fig.5. Storage modulus, loss modulus and tangent delta for the LY/HY5052/carbon composite system cured by conventional thermal RTM.

Fig. 6 .

 6 Fig.6. Storage modulus, loss modulus and tangent delta for the LY/HY5052/carbon composite system cured by microwave RTM.

Fig. 7 .

 7 Fig.7. SEM micrographs at same magnification showing fibre pull-out after four-point bending test; a) conventionally cured specimen b) microwave cured specimen.

Fig. 8 .Fig. 9 .

 89 Fig.8. SEM micrographs at same magnification after interlaminar shear test showing; a) clean fibres in conventionally cured specimen b) fibres coated with resin in microwave cured specimen

Fig. 1

 1 Fig.1

  Fig.2

  Fig.3

Fig.7 a) b)

Fig.8 a) b)

Fig.9 a) b)

Acknowledgement

The authors would like to thank Hexcel Composites and Huntsman for supplying the carbon fibres and the epoxy resin, respectively.


M Lillemose

L Gammelgaard

J Richter

E V Thomsen

A Boisen

Epoxy based photoresist/carbon nanoparticle composites

Epoxy based photoresist/carbon nanoparticle composites

Introduction

Micrometer sized cantilevers can be used for label-free biochemical detection [START_REF] Larvik | Cantilever transducers as a platform for chemical and biological sensors[END_REF]. Surface stress changes on one side of the cantilever due to the adsorption of molecules result in a bending of the cantilever. Detection of for example DNA [START_REF] Fritz | Translating biomolecular recognition into nanomechanics[END_REF], proteins [START_REF] Arntz | Label-free protein assay based on a nanomechanical cantilever array[END_REF], pesticides [START_REF] Alvarez | Development of nanomechanical biosensors for detection of the pesticide DDT[END_REF] and TNT explosives [START_REF] Pinnaduwage | Desorption characteristics of uncoated silicon microcantilever surfaces for explosive and common nonexplosive vapors[END_REF] has been demonstrated with cantilever-based sensors.

Cantilever bending is normally detected by an optical leverage technique (as known from atomic force microscopy) or by integrating a piezoresistor in the cantilever. Cantilever based sensors with integrated piezoresistive readout have traditionally been realized in silicon based materials [START_REF] Rasmussen | Optimised cantilever biosensor with piezoresistive read-out[END_REF]7]. Replacing the silicon based cantilever material with a soft polymer can potentially increase the surface stress sensitivity significantly. Moreover polymer processing is faster and cheaper than silicon processing. The challenge is here to find a soft polymer based piezoresistive material.

For a piezoresistive cantilever consisting of one material with Young's modulus, E, and thickness h, with an infinitely thin piezoresistor placed at the top, the surface stress sensitivity is given by [7]:

∆R R = -K 4 Eh σ s , (1)
where K is the gauge factor of the piezoresistor and σ s is the surface stress induced at the top surface of the cantilever. For a fixed geometry and disregarding the noise, the surface stress sensitivity is proportional to the relationship, ∆R/R∝K/E. To maximize the signal a soft material with a high gauge factor is needed. For single crystal silicon, K Si /E Si ≈150/150 GPa=1 GPa -1 .

Our aim is to fabricate an all polymer cantilever using SU-8 (a negative tone epoxy based photoresist) as the cantilever material. We have chosen to realize the cantilevers in SU-8, which is straightforward to structure with standard UV lithography. SU-8 has a Young's modulus of E≈4.4 GPa, so only a gauge factor of K>4.4 is needed for the piezoresistor in order to compete in sensitivity with silicon cantilevers.

We have previously fabricated SU-8 cantilevers with integrated gold strain gauges [START_REF] Tamayo | Polymeric cantilever arrays for biosensing applications[END_REF][START_REF] Johansson | SU-8 cantilever sensor system with integrated readout[END_REF][START_REF] Johansson | Polymeric cantilever-based biosensors with integrated readout[END_REF]. The gauge factor of gold is approximately K∼3.5 and we thus look for new soft materials with a higher gauge factor. We therefore started to investigate SU-8/carbon NP strain gauges. In our initial work we reported on the realization of an all polymer cantilever with integrated SU-8/carbon NP piezoresistive readout [START_REF] Gammelgaard | Microfabricated photoplastic cantilever with integrated photoplastic/carbon based piezoresistive strain sensor[END_REF]. In this work only one type of carbon particle was tested and the focus was on the possibility of fabricating a complete micromechanical device with an SU-8 based strain gauge.

In the present work we concentrate on the characterization of different SU-8/carbon NP strain gauges with respect to gauge factor and stability.

Experimental

Polymer/carbon NP composites

When adding conductive particles to an insulating polymer matrix the composite undergoes a transition from insulator to conductor. By adding more and more conductive filler particles the average interparticle distance is decreased. At a certain critical filler concentration (the percolation threshold)

an infinitely connected network is formed and the conductivity increases with a power law dependence (see ref. [START_REF] Balberg | Percolation and tunneling in composite materials[END_REF] for a detailed review). When a polymer composite is stretched the interparticle distance increases, some of the conductive paths are broken and the resistance goes up. This is the basic idea behind a polymer composite piezoresistor.

Polymer

Fabrication

For measuring the resistance of the polymer composite a characterization chip has been designed and fabricated using standard microfabrication techniques.

The white square in figure 2 it is still possible to UV-structure. As line width wasn't crucial in our design, we used an exposure dose of 1.53 J/cm 2 to be sure to crosslink the SU-8 composite.

Setup

The fabricated chips are placed in a four-point bending fixture specially designed for the chips. By applying a force to the two upper blades, the inserted chip is pressed down on two bottom blades and subjected to a pure bending moment, which results in a uniform stress stress along the length axis of the chip between the two upper blades (figure 3 -top) The force is applied by loading a weight on top of the clamped bricks (figure 3 -bottom).

The chip is electrically connected to a Keithley 2700 Multimeter through a Flat Flexible Cable. Using the set-up a well defined stress is applied to the fabricated SU-8/carbon NP composites and simultaneously we measure the two-point resistance. For a detailed description of the experimental setup we refer to ref. [START_REF] Richter | Piezoresistance of silicon and strained Si 0.9 Ge 0.1[END_REF].

Four-point measurements are the standard method for measuring resistivity; however, measuring the resistivity is not the main goal of this work. We are interested in measuring the piezoresistive effect and if we assume that the contact resistance is constant during strain, two-point measurements can be used. Furthermore, with our chip design we can measure the resistance for nine different lengths of the resistor, hence if we plot R vs. length we can find the contact resistance as the crossing of the R-axis.

Results

Resistivity

Figure 4 shows a plot of the resistivity, ρ, of the polymer composites as a function of the carbon NP content, when no stress is applied. As the carbon NP amount is increased, the resistivity drops several orders of magnitude and a transition from insulator to conductor is observed. The highest doped composites have resistivities of ρ ∼1.0×10 -2 Ωm, which is in the semiconductor region.

To determine the percolation threshold one would in principle need to scan the whole NP concentration range from insulator to conductor. For XE2 we estimated the percolation threshold to P c =1-2%. This value is quite low; however, XE2 is a 'high' structure carbon black with a OAN value (oil absorption value [START_REF] Astm) D | describe the structure of carbon NPs[END_REF]) value of 380 ml/100 g. High structure carbon blacks form elongated aggregates and are known to give lower percolation thresholds [START_REF] Rubin | Electronic structure: Wide-band, narrow-band, and strongly correlated systems -Critical behavior of the electrical transport properties in a tunneling-percolation system[END_REF][START_REF] Balberg | A comprehensive picture of the electrical phenomena in carbon blackpolymer composites[END_REF].

For Conductex 975U and Conductex 7051U the resistivities for the lower NP concentrations were too high to be measured with the present setup, hence we cannot positively say that the percolation threshold was reached. However; figure 4 suggests that Conductex 975U and Conductex 7051U should have higher percolation thresholds than XE2, because a higher concentration of NPs is needed to reach a resistivity equivalent to XE2. This is also what one could expect since Conductex 975U and Conductex 7051U have 'lower' structure than XE2, with OAN values of 169 ml/100 g and 121 ml/100 g, respectively [START_REF] Astm) D | describe the structure of carbon NPs[END_REF]. Lower structure carbon blacks do not form highly aggregated structures, hence a larger amount of NPs needs to be added to make a conductive composite.

Piezoresistivity

To investigate the basic mechanism of the polymer composite piezoresistor a measurement series was made, where tensile and compressive stress was applied alternately. As can be seen from figure 5 the baseline, R 0 , drifts, but the height of the peaks is stable, which means that the piezoresistive effect is reversible. Such a reversible piezoresistive effect has previously been reported by Knite et al.

for macro-size structures [START_REF] Knite | Reversible tenso-resistance and piezo-resistance effects in conductive polymer-carbon nanocomposites[END_REF]. We thus verify that a similar behavior is found for microstructures.

Gauge factors

We have measured gauge factors for the three different types of carbon NPs at different concentrations. The gauge factors were measured in two different

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ways:

(1) A tensile stress of 160 MPa was applied in a cyclic manner, hence 2 minutes stress off, 2 minutes stress on, 2 minutes stress off and so on. Three complete cycles were made for a total of 14 minutes. Figure 6 (left) shows the measured resistance, R, as a function of time for a 6.5% XE2 composite. As ∆R/R=Kε, where K is the gauge factor and ε is the strain, we can calculate the gauge factor from the values of the graph: ∆R is the height of the peaks, R 0 is the baseline and the strain is related to the stress through Hooke's law, σ = ε E, where E is Young's modulus for silicon. [18]. For the three peaks we find a gauge factor of K=11. The spike observed at the start of each stress cycle is due to the fact, that a weight load is manually placed on top of the four point bending probe. Even though the weight is loaded very gently, the stress goes momentarily up; however, the resistance rapidly stabilizes.

(2) A tensile stress is stepped from zero to 200 MPa in five equidistant steps over a period of 10 minutes. Again we have ∆R/R=Kε, so if we plot ∆R/R as a function of the strain, ε, we can find the gauge factor, K, as the slope of the graph. Figure 6 (right) shows a plot for a 6.5% XE2 composite. We do not observe a linear behavior in the whole strain range. A similar behavior has previously been measured and theoretically explained by Zhang et al. [START_REF] Zhang | Piezoresistance of conductor filled insulator composites[END_REF].

To estimate the gauge factor, we decided to make a linear fit (the dotted line)

in the interval ε ≥ 0.37•10 -3 , since the strain applied in method (1) lies in this interval. We find a gauge factor of K=11. There is good agreement between the gauge factors measured with method (1) and ( 2).

In figure 7 we have plotted the measured gauge factors as a function of the NP content for the three different carbon NPs used. As we go towards lower concentrations the gauge factors increase; however, the spread in the measured values increase for lower concentrations (indicated by the error bars). As an example, for the 10% 975U composite we measured gauge factors in the range K=22-54 on a single chip using different electrode configurations.

For higher NP concentrations we observe lower, but more reliable gauge factors. To illustrate this trend we have plotted the resistance, R, as a function of time for a 3.8% and an 8.6% XE2 composite in figure 8. The measurements shown are similar to the one in figure 6 (left). From the figure we find gauge factors of K=14-19 for the 3.8% composite and K=5-6 for the 8.6% composite; however, the measurements for the 3.8% composite have fluctuations of ≃1%, whereas it is less than 0.1% for the 8.6% composite. This trend can be illustrated if we plot the standard deviation of the resistivity, σ(ρ), as a function of the resistivity, ρ, for different XE2/SU-8 composites. The graph in figure 9 yields a straight line in a log-log plot, hence as the resistivity increases, the fluctuations increase with a power law dependence.

The noise of the equipment was estimated by measuring the resistance once every minute for 3 hours for (a) an ordinary resistor and (b) a 6.5% XE2 composite resistor, both with a resistance of 650 kΩ. We found that the fluctuations in the composite was around a factor of 100 times larger than in the ordinary resistor. Thus it can be concluded that the observed fluctuations are not instrumental.

Conclusion

Composites of SU-8 and three different types of carbon NPs have been fabricated. The resistivity has been measured and we have observed an insulator to conductor transition, when increasing the NP content of the composites.

The percolation threshold seems to be correlated with the OAN number, such that high OAN number gives a low percolation threshold. Fig. 6. LEFT: The resistance, R, as a function of time for a 6.5% XE2 composite, where a tensile stress of 160 MPa corresponding to a strain of ε = 1.1 • 10 -3 , is applied in a cyclic manner. From the baseline, R 0 , and the height of the peaks, ∆R, we can calculate the gauge factor, K. For the three peaks in the figure we find a gauge factor of K=11. RIGHT: The relative change in resistance, ∆R/R, as a function of the strain, ε. The overall behavior is not linear. We have decided to make a linear fit (the dotted line) in the interval ε ≥ 0.37 • 10 -3 and from the slope we find K=11 for the gauge factor. XE2 composite with a resistivity of ρ=1860 Ωm and 1.5 Ωm, respectively. The resistance was measured using the same method as in figure 6 -left. For the 8.6% XE2 composite peak 1-3 give gauge factors of 5.4, 6.1 and 6.5, respectively. For the 3.8% XE2 composite we get 14.6, 16.7 and 18.9, respectively. However, the resistance of the 3.8% composite fluctuates up to ∼1% (1 MΩ).

  figure2). On each chip we can thus do mesurements with nine different electrode configurations.The chips are fabricated through the following steps: (a) a 150 nm electri-

Figure 5

 5 shows the resistance, R, of an SU-8/carbon NP composite as a function of time. From the figure it is clear that the behavior is as expected: (a) Onset of tensile (compressive) stress, (b) conductive paths are broken (established) and R increases (decreases), (c) tensile (compressive) stress is released and (d) conductive paths are re-established (broken) and R returns to its initial value.

Fig. 1 .

 1 Fig. 1. An SEM image of a 10% CD975U/SU-8 polymer composite. The aggregated NP structures are seen as the bright points.

Fig. 2 .Fig. 3 .

 23 Fig. 2. A picture of the fabricated chip. The white square in the middle of the chip is the structured SU-8/carbon black NP composite. The gold electrodes are also visible.

Fig. 4 .

 4 Fig.4. The resistivity, ρ, as a function of the carbon NP particle content for three different types of particles. A clear insulator to conductor phase transition is observed. The resistance was measured with a Keithley 2700 Multimeter in two-point resistance mode passing a 0.7 µA current through the resistor and measuring the voltage drop across. The graph suggests that a lower OAN value gives lower percolation threshold.

Fig. 5 .

 5 Fig. 5. An 8.65% XE2/SU-8 composite, where tensile and compressive stress is applied alternately. (a) Tensile stress applied, (b) conductive paths are broken and R increases, (c) tensile stress released and (d) conductive paths are re-established and R returns to initial value. The peaks correspond to time intervals of nonzero tensile stress and the valleys to periods of nonzero compressive stress.

Fig. 7 .

 7 Fig.7. The measured gauge factors as a function of the NP content for the three different carbon NPs used. The gauge factors increase for lower concentrations; however it can be seen that the spread in the measured values increase (indicated by the error bars). For 7051U only two data points were obtained as we could not measure gauge factors for concentrations lower than 20%. The data presented represents measurements for several chips and at different electrode spacings.

Fig. 8 .

 8 Fig.8. The resistance, R, as a function of time for 3 tensile stress cycles of 180 MPa. The measured values are for a 3.8% (right R-axis) and 8.6% (left R-axis) XE2 composite with a resistivity of ρ=1860 Ωm and 1.5 Ωm, respectively. The resistance was measured using the same method as in figure6-left. For the 8.6% XE2 composite peak 1-3 give gauge factors of 5.4, 6.1 and 6.5, respectively. For the 3.8% XE2 composite we get 14.6, 16.7 and 18.9, respectively. However, the resistance of the 3.8% composite fluctuates up to ∼1% (1 MΩ).

Fig. 9 .

 9 Fig. 9. The standard deviation of the resistivity, σ(ρ) as a function of the resistivity for the XE2 composites. The graph shows a straight line in a log-log plot indicating a power law dependence.

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µm Figure(s)

We have shown that the composites are piezoresistive and measured stable and reversible gauge factors of K=5-9, but gauge factors of up to K=40-50 have also been measured. A gauge factor of K∼5 is sufficient for an all polymer based cantilever to compete in sensitivity with a silicon based cantilever.

We have shown that for a specific NP, composites with a lower NP content gives larger, but less stable gauge factors compared to composites with a higher NP content.

Future work involves integration of the polymer piezoresistor in an al polymer based cantilever sensor system. The noise level of the composites needs to be investigated thoroughly.

Acknowledgements

This work is funded by the European Union as a part of the NOVOPOLY project (NMP-3.4.2.3-1). We thank Degussa Corporation and Columbian Chemicals Company for supplying carbon nanoparticles.