Jennifer M Ullbrand

José M Córdoba

Javier Tamayo-Ariztondo

María R Elizalde

Mats Nygren

Jon M Molina-Aldareguia

Magnus Odén

Thermomechanical Properties of Copper-Carbon Nanofibre Composites Prepared by Spark Plasma Sintering and Hot Pressing

Keywords: Carbon nanofibre, copper, metal matrix composite, thermal conductivity, coefficient of thermal expansion

Several types of carbon nanofibres (CNF) were coated with a uniform and dense copper layer by electroless copper deposition. The coated fibres were then sintered by two different methods, spark plasma sintering (SPS) and hot pressing (HP). The Cu coating thickness was varied so that different volume fraction of fibres was achieved in the produced composites. In some cases, the CNF were pre-coated with Cr for the improvement the Cu adhesion on CNF. The results show that the dispersion of the CNF into the Cu matrix is independent of the sintering method used. On the contrary, the dispersion is directly related to the efficiency of the Cu coating, which is tightly connected to the CNF type. Overall, strong variations of the thermal conductivity (TC) of the composites were observed (20-200 W/mK) as a function of CNF type, CNF volume fraction and Cr content, while the coefficient of thermal expansion (CTE) in all cases was found to be considerably lower than Cu (9.9-11.3 ppm/K). The results show a good potential for SPS to be used to process this type of materials, since the SPS samples show better properties than HP samples even though they have a higher porosity, in applications where moderate TC and low CTE are required.

INTRODUCTION.

Composites (metal matrix composite (MMC), ceramic matrix composite (CMC)) are a class of engineering materials that offer opportunities to tailor properties and meet specific requirements [START_REF] Chu | Thermal conductivity of SPS consolidated Cu/diamond composites with Cr-coated diamond particles[END_REF][START_REF] Chu | Fabrication and effective thermal conductivity of multi-walled carbon nanotubes reinforced Cu matrix composites for heat spreader applications[END_REF][START_REF] Cordoba | Properties of Ti(C,N) cermets synthesized by mechanically induced self-sustaining reaction[END_REF][START_REF] Waku | A ductile ceramic eutectic composite with high strength at 1,873 K[END_REF], e.g. in heat spreaders. The ideal material for heat spreader applications combines a high thermal conductivity (TC) with a high heat capacity to absorb thermal spikes, and a low coefficient of thermal expansion (CTE), to remove heat from, for example, electronic circuit boards while avoiding stresses due to thermal cycling. Metals have a high thermal conductivity and hence are commonly used as heat spreaders, but their CTE is high [START_REF] Asthana | Processing effects on the engineering properties of cast metal-matrix composites[END_REF][START_REF] Legzdins | MMCX -An expert system for metal matrix composite selection and design[END_REF].

Materials such as CuMo, CuW, Al/SiC, AlN, FeNi, FeNiCo and diamond are usually used when a heat spreader with high TC and low CTE is needed. However, these materials suffer from various drawbacks, such as high density (MoCu, WCu), relatively high CTE (Al/SiC, AlN), poor thermal performance (FeNi, FeNiCo) and in the case of diamond, a high price [START_REF] Hammel | Carbon nanofibers for composite applications[END_REF][START_REF] Kitzmantel | Wetting and infiltration behaviour of copper and copper alloys on carbon nanofibers[END_REF] and low machinability. Composite materials, especially metal matrix composites, have the potential to fulfill the requirements, because a careful control of the reinforcing filler to matrix ratio allows good adjustment of physical and mechanical properties [START_REF] Chung | Materials for thermal conduction[END_REF]. In this context, carbon nanofibres (CNF) are currently being considered as a potential reinforcement for various composite materials due to their excellent mechanical, thermal, and electrical properties [START_REF] Endo | Vapor-grown carbon fibers (VGCFs) -Basic properties and their battery applications[END_REF][START_REF] Carneiro | Production and assessment of polycarbonate composites reinforced with vapour-grown carbon fibres[END_REF][START_REF] Heremans | Thermal conductivity and thermopower of vapor grown graphite fibres[END_REF][START_REF] Heremans | Electrical conductivity of vapor-grown carbon-fibres[END_REF][START_REF] Cordoba | Morphology influence of the oxidation kinetics of carbon nanofibers[END_REF]. More specifically, CNF reinforced copper composites constitute an ideal candidate to work as an electrical contact material and/or substrate for semiconductor devices [START_REF] Ngo | Thermal interface properties of Cu-filled vertically aligned carbon nanofiber arrays[END_REF] due to their potentially high electrical and thermal conductivity, high wear resistance, low coefficient of thermal expansion, good machinability and low price.

One of the main difficulties to manufacture such a material resides on the poor wetting between copper and carbon, which makes it difficult to fabricate dense composites with homogeneously distributed fibres, through infiltration methods. It is anticipated that the wetting problem could be overcome by coating the carbon nanofibre with copper before the consolidation process. Different coating techniques can be used, such as sputtering [START_REF] Neubauer | The influence of mechanical adhesion of copper coatings on carbon surfaces on the interfacial thermal contact resistance[END_REF], electrodeposition [START_REF] Arai | Carbon nanofiber-copper composite powder prepared by electrodeposition[END_REF][START_REF] Wan | Effects of fiber volume fraction, hot pressing parameters and alloying elements on tensile strength of carbon fiber reinforced copper matrix composite prepared by continuous three-step electrodeposition[END_REF], wetness technique [START_REF] Ma | Characteristics of copper particles supported on various types of graphite nanofibers[END_REF], but the method of choice for metallization and interconnection which has emerged in recent years is electroless deposition due to its low cost, fast deposition rate, good filling capability, good uniformity and low process temperatures [START_REF] Yung | Plating of copper into through-holes and vias[END_REF][START_REF] Lu | Electrochemical and thermal behavior of copper coated type MAG-20 natural graphite[END_REF][START_REF] Cordoba | Growth and characterization of electroless deposited Cu films on carbon nanofibers Surf[END_REF][START_REF] Mallory | Electroless plating. Fundamentals and Applications[END_REF]. The deposition of Cr interlayers onto the CNF before Cu deposition has also been explored with the objective of improving the adhesion of Cu on CNF. Cr is known to react with C and Cu to form the corresponding Cr-C carbide and Cr-Cu alloy. Previous work [START_REF] Mayerhofer | Adhesion promotion of Cu on C by Cr intermediate layers investigated by the SIMS method[END_REF][START_REF] Mortimer | The wetting of carbon and carbides by copper alloys[END_REF][START_REF] Mortimer | The wetting of carbon by copper and copper alloys[END_REF] have shown that the interface properties between C and Cu using Cr is improved by the presence of interlayers. The interlayer should although be thin since impurities in the Cu-matrix decrease the thermal conductivity.

Hence, in this paper we report on the fabrication of carbon nanofibre reinforced copper composites where the fibres are coated by electroless deposition prior to densification by spark plasma sintering or hot pressing. The difference between the used fibre types is how the relative orientation of the graphene sheets with respect to the fibre direction is aligned. The performance of the composites was evaluated in terms of their coefficient of thermal expansion and thermal conductivity.

EXPERIMENTAL.

Two types of carbon nanofibres have been used in the analysis: nanofibres with a herringbone structure (300 nm in diameter and specific surface area of 165 m 2 /g, Future Carbon GmbH) denoted HB and a platelet structure (150 nm in diameter and specific surface area of 72 m 2 /g, Future Carbon GmbH) denoted PL. In the PL CNF, the graphene sheets are perpendicular to the longitudinal direction, while they are inclined in the HB CNF. In some cases, the HB and PL fibres were heat treated at 3023 K for 30 min in helium atmosphere prior to the composite fabrication and those will be denoted HBHT and PLHT, respectively.

With the intention of improving the Cu-C interface, some of the PLHT CNF were coated with Cr prior to the Cu electroless deposition. Cr was deposited by dc magnetron sputtering from a 99.95% pure Cr target resulting in a 24% weight increase of the fibres. These fibres are denoted PLCr.

The electroless copper deposition of the samples was carried out following the optimized parameters and reagents described by Cordoba et al. [START_REF] Cordoba | Growth and characterization of electroless deposited Cu films on carbon nanofibers Surf[END_REF]. Sintering was performed using two different methods: spark plasma sintering SPS (Dr Sinter 2050) and hot pressing HP (Sintris 10STV). The SPS samples were manually pre-pressed and sintered in a 12 mm graphite die. The inside of the die and the punches were covered with a graphite layer. In the case of HP, the graphite punches in contact with the sample were covered by a layer of boron nitride to avoid reaction with the sample. The HP machine heats mainly by heating wires, but also apply a low intensity DC current through the die. Both SPS and HP temperatures were measured by pyrometer. Table I summarizes the different samples synthesized and the sintering conditions (maximum temperature, holding time and pressure).

The morphology of the surfaces of the samples, polished down to 1 m diamond paste, were determined by optical microscopy (Olympus BX 60 equipped with a digital camera (Olympus E410)) and scanning electron microscopy (JEOL LEO 1550 Gemini).

X-ray diffractograms (Philips PW-1729) of powders and polished surfaces were recorded with Cu K radiation over 41-47 2 range and the FWHM of the Cu (111) diffraction peak was used for calculating the average diameter of the coherent diffracting domain according to the Scherrer equation [START_REF] Scherrer | Estimation of the size and internal structure of colloidal particles by means of röntgen[END_REF].

The density of the CNF was determined by pycnometry [START_REF] Viana | About pycnometric density measurements[END_REF] in an AccuPyc 1330. The equipment was calibrated with a standard volume. Each run consisted of ten He-purges prior to five readings of the sample volume. The average volume was used to extract the density of the five measurements. The standard deviation of one run was less than 0.0008 cm 3 for sample volumes in the range of 0.1119-0.8223 cm 3 .

The bulk densities of the sintered composites were measured by the water immersion technique (Archimedes' method) using a Sartorius ME2359 balance and taking the temperature effect on the water density into account. The maximum achievable densities were estimated based on the rule of mixtures.

The thermal diffusivity of the composites was measured at room temperature in air using an ordinary light flash and an infrared (IR) camera. A light sensor coupled to an oscilloscope tracked the starting point of the measurement. The thermal diffusivity values were determined based on the time to reach half of the maximum temperature on the rear side of the sample and the sample thickness [START_REF] Nagpure | Thermal diffusivity study of aged Li-ion batteries using flash method[END_REF]. The thermal conductivity was calculated according to the following equation [START_REF] Araki | Transient characteristics of thermal conduction in dispersed composites[END_REF][START_REF] Kerrisk | Thermal Diffusivity of Heterogeneous Materials.2 Limist of Steady-State Approximantion[END_REF]:

TC = ( Cu C Cu ƒ Cu + CNF C CNF ƒ CNF + Cr C Cr ƒ Cr ) ( 1
)
where TC is the thermal conductivity, is the measured thermal diffusivity, CNF and Cu are the densities of the copper and carbon nanofibre, respectively, and ƒ Cu and ƒ CNF are the volume fractions of copper and carbon nanofibre, respectively. The specific heat capacities, C CNF , C Cu , and C Cr used for the fibres, copper and chromium were 930, 385, and 450 J/kgK respectively.

The coefficients of thermal expansion were determined with a dilatometer (Netzsch DIL 402C). Heating cycles were performed in He atmosphere from room temperature to 300ºC with a heating and cooling rate of 10 ºC/min. The coefficient of thermal expansion, CTE, was extracted from the linear slope according to equation 2 [START_REF] Hummel | Electronics Properties of Materials[END_REF][START_REF] Barsoum | Fundamentals of Ceramics[END_REF][START_REF] Nix | The thermal expansion of pure metal: Copper, Gold, Aluminium Nickel, and Iron[END_REF]:

L/L = CTE T (2)
where, L is the initial specimen length, L is its dilatation and T is the temperature increase.

RESULTS AND DISCUSSION

The densities of the as-received and treated CNF fibres determined by pycnometry were: 2.21 g/cm 3 for the CNF-HB, 2.24 g/cm 3 for the CNF-PL, 2.18 g/cm 3 for the CNF-HBHT, and 2.17 g/cm 3 for the CNF-PLHT. The density of the PLCr powder was calculated by the rule of mixtures, using the experimental density value for the CNF-PL and the standard density of Cr (7.14 g/cm 3 ) resulting in a value of 2.68 g/cm 3 .

The morphology of the as received and heat treated CNF with and without Cu coatings have been reported elsewhere [START_REF] Cordoba | Morphology influence of the oxidation kinetics of carbon nanofibers[END_REF][START_REF] Cordoba | Growth and characterization of electroless deposited Cu films on carbon nanofibers Surf[END_REF]. In general, Cu depositions onto heat treated CNF (HBHT, PLHT, and PLCr) results in a complete coverage of the fibres while as received fibres show some naked areas. This is an effect of the reorientation of the outermost graphite sheets during heat treatment and the addition of Cr, which promotes the adhesion of Cu to the CNF.

The measured densities of the metal matrix composites are presented in Table II. As expected, the density decreases linearly with CNF reinforcement since the total density is the volume fraction average of the density of the matrix and reinforcements. The measured densities for the 7% wt , 14% wt and 24% wt reinforcement of CNF were 6-7 g/cm 3 , 4.5-6.0 g/cm 3 and 4.0-5.5 g/cm 3 , respectively. The composite density is strongly affected by the choice of sintering method, with HP displaying significantly higher densities (95-100 %) than SPS 82-97 %).

The size of the Cu crystallites, determined by the Scherrer equation, is shown in Table II. Samples sintered by SPS show smaller crystallite sizes than those sintered by HP. This is a consequence of the higher temperatures used during the hot pressing process. The dispersion of the CNF in the Cu matrix was assessed by optical microscopy. Micrographs in Figure 1(a)

to (e) show examples over sample surfaces, where darker areas represent higher CNF content and brighter areas represent higher content of copper. In all cases, the CNF dispersion observed was homogeneous, with the exception of some agglomerates seen for both of the sintering methods used. The main difference in dispersion is connected to the type of CNF used. For instance, the CNF agglomerates observed were few but large (350 m) for PL CNF, while for HB CNF the agglomerates were observed more frequently but with a smaller size (60 m) see Figure 1 HB CNF (sample HP-HB24) and by SPS using 14% PLHT CNF(sample HP-PLHT14) respectively. In Figure 1(f) a clear preferential orientation of the CNF in planes perpendicular to the pressing direction is observed, which has previously been documented as a consequence of the uniaxial pressing, [START_REF] Barcena | Microstructural study of vapour grown carbon nanofibre/copper composites[END_REF][START_REF] Korb | Properties and microstructure of short carbon fibre reinforced copper matrix composites[END_REF][START_REF] Geffroy | Structural and thermal properties of hot pressed Cu/C matrix composites materials used for the thermal management of high power electronic devices[END_REF]. Figure 1(g) shows the interface between the Cu matrix and a fibre. Despite the poor wetting between C and Cu [START_REF] Lee | Effect of alloying elements on the interfacial bonding strength and electric conductivity of carbon nano-fiber reinforced Cu matrix composites[END_REF], a dense C/Cu interface was achieved. No interdiffusion or segregation over large distances (> 1 m) could be detected by energy dispersive x-ray spectroscopy (EDX). However, reactions at the interface cannot be excluded, especially in the presence of Cr.

Table II summarises the values of the thermal diffusivity measured parallel to the prior pressing direction. A large variation between 3 and 60 mm 2 /s is observed depending on the CNF type, volume fraction and sintering method. Figure 2 shows the thermal conductivity of the samples and similar dependences are seen here. Due to the presence of chromium, the weight fraction of CNF for the PLCr samples is 6%, 10% and 18%, instead of 7%, 14% and 24%, which is the case for the other samples. The thermal conductivity decreases sharply with increasing weight fraction of fibres. Hence the maximum thermal conductivity is observed for PLCr6 with ~200 W/Km, which is considerably lower than pure Cu (385 W/Km) but substantially higher than what has been observed for Cu/CNF composites before [START_REF] Silvain | Novel processing and characterization of Cu/CNF nanocomposite for high thermal conductivity applications[END_REF].

The decrease in thermal conductivity with respect to pure Cu might come from several factors: (i) the crystal quality of the CNF, (ii) the orientation of the graphene planes within the CNF, (iii) the interfacial thermal resistance of the Cu-C interface, (iv) alignment of the CNF in the pressing direction during processing, (v) inhomogeneous fibre dispersion and (vi) the amount of porosity in the final composite.

First of all, considering the expected anisotropy of the CNF and the observed preferred alignment of the CNF, perpendicular to the pressing direction, we expect a higher conductivity in the perpendicular pressing direction than in the parallel pressing direction of the sample. However, the thermal conductivity was only measured in the parallel direction due to constrains in the sample size. Therefore, the measured thermal conductivities should be interpreted with caution. For instance, the strong decrease in thermal conductivity with CNF volume fraction could be attributed to this anisotropy effect, especially when considering that no major changes have been observed in density and/or CNF dispersion with volume fraction.

Even with the limitations addressed above, the relative differences in thermal conductivity observed raise some interesting points. For instance, the HB fibres are more bent and entangled than PL fibres [START_REF] Cordoba | Growth and characterization of electroless deposited Cu films on carbon nanofibers Surf[END_REF] resulting in a decrease in the net heat transport which could explain why the thermal diffusivity values of the PL composites are better. The dramatic effect of the high temperature heat treatment of the CNF is due to the graphitisation [START_REF] Cordoba | Morphology influence of the oxidation kinetics of carbon nanofibers[END_REF],

which results in improved Cu deposition efficiency. Finally, it is also clear from the results that the introduction of a Cr interlayer between the Cu and the CNF has a beneficial effect on the thermal conductivity of the composite, at least in the samples processed by SPS. This is likely due to the improvement of Cu adhesion onto the Cr-C interface, resulting in an improved Cu coating process and a reduction of the thermal resistance of the interface.

However, in the samples processed by HP, the beneficial effect of the Cr addition is not that clear. In this case, the processing temperatures are considerably higher and it is well known that the dissolution of Cr into the Cu matrix would have a detrimental effect on the thermal properties.

Table III summarizes the coefficients of thermal expansion (CTE) of selected samples, measured from room temperature to 250ºC. The CTE of the Cu/CNF composite depends upon the CNF type, the volume fraction of CNF, the CNF orientation, the quality of the dispersion, the porosity, and the interface resistance. In all studied cases, the CTE values are in a narrow range, 9.9-11.3 ppm/K, independently of the fibre content and the sintering method. The addition of Cr on the CNF surface results in a decrease in the CTE (PLCr6: HP, 10.1 ppm/K; PLCr6: SPS, 9.9 ppm/K), which is consistent with the expected improved Cu adhesion.

Overall and compared with other techniques (powder metallurgy [START_REF] Korab | Thermal conductivity of unidirectional copper matrix carbon fibre composites[END_REF][START_REF] Prakasan | Thermal expansion characteristics of cast Cu based metal matrix composites[END_REF], tape casting [START_REF] Geffroy | Preparation by tape casting and hot pressing of copper carbon composites films[END_REF],

copper salt chemical decomposition [START_REF] Silvain | Novel processing and characterization of Cu/CNF nanocomposite for high thermal conductivity applications[END_REF]), electroless copper deposition as a pre step before sintering has been shown to be a superior method for the enhancement of the thermomechanical properties of CNF/Cu composites. Also when comparing with other Cu reinforcement materials [START_REF] Geffroy | Preparation by tape casting and hot pressing of copper carbon composites films[END_REF][START_REF] Kang | Achieving highly dispersed nanofibres at high loading in carbon nanofibre-metal composites[END_REF] such as with SiC [START_REF] Schubert | Interfacial design of Cu-based composites prepared by powder metallurgy for heat spreader applications[END_REF], W [START_REF] German | A model for the thermal properties of liquid phase sintered composites[END_REF], AlN [START_REF] Lee | Thermomechanical properties of AlN-Cu composite materials prepared by solid state processing[END_REF] and with Mo [47], the carbon nanofibres have shown better thermal conductivity and lower CTE.

We note that, although SPS samples in general show lower relative density (higher residual porosity) than the HP samples they display comparable values of thermal conductivity and CTE. The spark plasma sintering and hot pressing parameters (temperature, pressure and holding time) affect the interface between CNF and copper as well as the CNF degradation, therefore further detailed studies of the C-Cu interface is needed in order to understand the influence of the chosen sintering method. An explanation for the higher densities but worse properties for the HP samples might be the high compressing temperatures and pressures that can cause Cr and Cu to melt and flow away from the CNF surface reducing the interface properties and thermal conductivity. The effect of the current in the SPS sintering is relatively unknown and might influence the interfacial properties as well.

However, our results suggest that the SPS technique is a promising sintering method to obtain superior composites for heat spreader applications, since further densification would yield an even higher thermal conductivity. Additionally, the use of highly graphitic fibres (fibres subjected to high temperature pre-treatment) is preferable.

  (a) and (b). The samples with heat treated CNF and PLCr show the least agglomeration (Figure1(c) and (e) compared to untreated fibres Figure1 (a) and (d)). The amount of fibre agglomeration in the composites is directly related to the Cu-coating process, so that the best fibre dispersion is achieved for the most homogeneously Cu coated fibres.

Figure 1 (

 1 Figure 1 (f) and (g) show SEM images of the composite processed by HP using 24%

Figure 1 .

 1 Figure 1. Representative optical images of the samples, (a) HP-PL7, (b) HP-HB7. (c) SPS-

Figure 2 .

 2 Figure 2. Thermal conductivity of the composites reinforced with Platelet (PL) and

Figure 2

 2 Figure 1

Table I .

 I Summary of the sample compositions (HB, CNF-herringbone; PL, CNF-platelet; HBHT, CNF-herringbone heat treated; PLHT, CNF-platelet heat treated and PLCr, CNFplatelet previously coated with Cr by PVD) and used experimental conditions (T, temperature;

	p, pressure; t hold , holding time)
	Sample ID	Composition (%wt.)	T (ºC)	SPS P (MPa)	t hold (min)	T (ºC)	HP P (MPa)	t hold (min)
	HB7	7%CNF HB +93%Cu	600	50	5	850 122	3
	HB14	14%CNF HB +86%Cu	600	50	5	900 100	3
	HB24	24%CNF HB +76%Cu	600	100	5	975 122	4.2
	PL7	7%CNF PL +93%Cu	600	50	5	750	80	3
	PL14	14%CNF PL +86%Cu	600	50	5	800 100	3
	PL24	24%CNF PL +76%Cu	650	100	3	975 138	3
	HBHT7	7%CNF HBHT +93%Cu	600	50	5	750 122	3
	HBHT14	14%CNF HBHT +86%Cu	600	50	5	900 100	3
	HBHT24	24%CNF HBHT +76%Cu	650	100	3	975 122	4.2
	PLHT7	7%CNF PLHT +93%Cu	600	50	5	750	80	3
	PLHT14	14%CNF PLHT +86%Cu	600	50	5	900 100	3
	PLHT24	24%CNF PLHT +76%Cu	650	100	3	-	-	3
	PLCr6	6%CNF PL +92.7%Cu+1.3%Cr 600	50	5	750 100	3
	PLCr10 10%CNF PL +87.8%Cu+2.2%Cr 600	50	5	900 122	3
	PLCr18	18%CNF PL +78%Cu+4%Cr	550	100	1	975 138	3

SPS, Spark Plasma Sintering; HP, Hot Pressing

Table III .

 III Thermal conductivity (TC) [±1%] and normal coefficient of thermal expansion (CTE) [±1%] from room temperature (rt) to 250ºC measured parallel to the pressing direction.Only the composites showing the highest thermal conductivity are reported.

	Sample	TC [W/Km]	CTE [10 -6 K -1 ] rt 250ºC	Relative density [%]	Processing conditions T P (ºC) (MPa)
	PL7 (HP)	144.7	13.1	98	750	80
	PL7 (SPS)	115.2	10.3	94	600	50
	HBHT7 (HP)	171.8	10.4	96	750	122
	HBHT7 (SPS)	157.3	11.5	89	600	50
	PLHT7 (HP)	147.1	12.7	97	750	80
	PLHT7 (SPS)	144.0	10.2	92	600	50
	HBHT14 (HP) 106.2	11.0	100	900	100
	PLHT14 (HP)	102.5	11.9	99	900	100
	PLCr6 (HP)	104.5	10.1	99	750	100
	PLCr6 (SPS)	194.6	9.9	95	600	50

Sample sintered by hot pressing (HP) or spark plasma sintering (SPS)

Acknowledgements: Gerhard Traxler and Thomas Placzek from Dept. of Materials Technology, ARC Seibersdorf Research, and Dr. Volker Brüser and Sina Kutschera from Dept. Plasmaprozesstechnik, Leibniz-Institut für Plasmaforschung und Technologie e.V. The Basque Government (project Etortek-inanogune, project PI-07/17) is also acknowledged for the funds received. J.T-A is grateful to the Basque Government for his PhD Grant. The financial contribution of EU FP6 031712 INTERFACE project is gratefully acknowledged *Corresponding Author. Phone: +46 13281780. E-mail address: josga@ifm.liu.se

CONCLUSIONS

In this paper we report on the fabrication of carbon nanofibre reinforced copper composites where the fibres are coated by electroless deposition prior to densification by spark plasma sintering or hot pressing, using different types of CNF: platelet PL, herringbone HB, before and after a high temperature (HBHT and PLHT) graphitization process, and also CNF precoated with Cr (PLCr).

Compared with conventional powder metallurgy and infiltration techniques, a very good dispersion of the CNF into the Cu matrix can be achieved with this approach, independently of the sintering method used. The results show that the quality of the dispersion depends on the efficiency and homogeneity of the Cu coating, which depends on the CNF type, HB shows the worst dispersion while HBHT, PLHT and PLCr shows superior dispersion.

SPS sintering, which is carried out at considerably lower temperatures (550-650ºC) than hot pressing (750-975ºC), results in slightly lower densities, but the final thermal properties obtained for both routes are comparable. This indicates that the SPS technique stabilizes the adhesion of Cu on the Cr-C interface possibly by the current or by the lower operating temperature. The SPS sintering technique is therefore, according to this study, the more promising sintering method. Overall, strong variations of the thermal conductivity of the composites were observed (20-200 W/mK) as a function of CNF type, CNF volume fraction and Cr content, while the coefficient of thermal expansion (CTE) in all cases was found to be considerably lower than Cu (9.9-11.3 ppm/K).

The highest thermal conductivity (194 W/Km) was obtained for Cr pre-coated CNF and a weight fraction of 6% processed by SPS. All in all, the results show a good potential for SPS to be used to process this type of materials in applications where moderate thermal conductivity and low CTE are required.


Sheikh M Uddin

Tanvir Mahmud

Christoph Wolf

Carsten Glanz

Ivica Kolaric

Christoph Volkmer

Helmut Höller

Ulrich Wienecke

Siegmar Roth

Hans-Jörg Fecht

Minhaz Sheikh

Uddin

Effect of size and shape of metal particles to improve hardness and electrical properties of carbon nanotube reinforced copper and copper alloy composites

Keywords: A. Carbon nanotubes, A. Metal matrix composites, B. Mechanical Properties, B. Electrical Properties, E. Hot-pressing, Copper and Copper alloys

des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

Introduction

Carbon nanotubes (CNTs) have received considerable interest in research due to their unique atomic structure and fascinating properties. Till now, the theoretical and experimental results achieved indicate that CNTs are the most unique and versatile materials so far discovered in the world. CNTs have an extremely high Young's modulus of 0.5-2 TPa [START_REF] Yu | Tensile loading of ropes of single wall carbon nanotubes and their mechanical properties[END_REF][START_REF] Yu | Strength and breaking mechanism of multiwalled carbon nanotubes under tensile load[END_REF][START_REF] Wong | Nanobeam mechanics: elasticity, strength, and toughness of nanorods and nanotubes[END_REF][START_REF] Treacy | Exceptionally high Young's modulus observed for individual carbon nanotubes[END_REF][START_REF] Yakobson | Nanomechanics of carbon tubes: instabilities beyond linear response[END_REF][START_REF] Lu | Elastic properties of carbon nanotubes and nanoropes[END_REF], large ultimate tensile strength of 20-150 GPa [START_REF] Yu | Tensile loading of ropes of single wall carbon nanotubes and their mechanical properties[END_REF][START_REF] Yu | Strength and breaking mechanism of multiwalled carbon nanotubes under tensile load[END_REF][START_REF] Treacy | Exceptionally high Young's modulus observed for individual carbon nanotubes[END_REF][START_REF] Walters | Elastic strain of freely suspended single-wall carbon nanotube ropes[END_REF][START_REF] Wagner | Stress-induced fragmentation of multiwall carbon nanotubes in a polymer matrix[END_REF][START_REF] Belytschko | Atomistic simulation of nanotube fracture[END_REF] and usually high flexibility [START_REF] Wong | Nanobeam mechanics: elasticity, strength, and toughness of nanorods and nanotubes[END_REF][START_REF] Iijima | Structural flexibility of carbon nanotubes[END_REF]. Other than their exceptional mechanical properties, CNTs have an extremely high aspect ratio, excellent chemical stability as well as superior thermal and electrical properties. All these characteristics have made CNTs outstanding reinforcement materials for developing advanced nanocomposites for various applications.

Considerable researches have been conducted in carbon nanotube reinforced polymer matrix composites with a remarkable enhancement in mechanical properties compared to those of monolithic materials. However, limited research has been done in preparation, structural, physical and mechanical properties of metal-CNT nanocomposites. Poor wetting behaviour or weak interfacial bonding to matrix materials, agglomeration among themselves with Van-der-Waals force, inhomogeneous distribution of CNTs in the matrices and degraded thermal stability at high processing temperature are the prime drawbacks to use carbon nanotubes as reinforcements of metal matrix composites. For instance, Kuzumaki et al. reported very little improvement in tensile strength of CNT reinforced Al nanocomposites prepared by conventional powder mixing, hot pressing followed by hot extrusion due to inhomogeneous dispersion of CNTs in the metal matrix [START_REF] Kuzumaki | Processing of carbon nanotube reinforced aluminium composite[END_REF]. Moreover, carbon nanotubes as reinforcements of metal matrix composites have so far been reported poor efficiency [START_REF] Zhan | Single-wall carbon nanotubes as attractive toughening agents in alumina-based nanocomposites[END_REF][START_REF] Xu | Fabrication of aluminum-carbon composites and their electrical properties[END_REF][START_REF] Chen | Carbon-nanotube metal-matrix composite prepared by electroless plating[END_REF] due to agglomeration among themselves with Van-der-Waals force, poor wetting behaviour or weak interfacial bonding to matrix materials [START_REF] Kuzumaki | Mechanical characteristics and preparation of carbon nanotube fiber reinforced Ti composite[END_REF][START_REF] Lau | The revolutionary creation of new advanced materials-carbon nanotube composite[END_REF][START_REF] Lau | Effectiveness of using carbon nanotubes as nanoreinforcements for advanced composite structures[END_REF] and degraded thermal stability at high sintering temperature [START_REF] Zhan | Single-wall carbon nanotubes as attractive toughening agents in alumina-based nanocomposites[END_REF][START_REF] Frankland | The stress-strain behavior of polymer-nanotube composites from molecular dynamics simulation[END_REF][START_REF] Feng | Fabrication and properties of silver-matrix composites reinforced by carbon nanotubes[END_REF][START_REF] Xu | Fabrication of aluminum-carbon nanotube composites and their electrical properties[END_REF].

For any case, the homogeneous dispersion of nanotubes in metal matrices is the prime requirement before fabricating carbon nanotube based metal matrix nanocomposites. Early studies [START_REF] Uddin | Thermal Expansion Coefficient of Nanotube-Metal Composites[END_REF] of the authors showed that, high energy mechanical alloying involved continuous impact, welding, fracturing and rewelding of powders to disperse nanotubes in the metal powders. Different milling and sintering parameters were optimized for CNT-Al and CNT-Cu nanocomposites. In this contribution, we report on the successful preparation of CNT reinforced copper and bronze composites by mechanical mixing of nanotubes and metal powders followed by well-known hot-press sintering method of powder metallurgy. The effect of shape and size of metal particles to improve hardness of CNT-Cu composites and selection of nanotubes (single-walled or multi-walled) to improve electrical properties of CNT-Bronze composites have been investigated.

Experimental

Materials

Multi-walled nanotube (MWNT) with an average diameter of (10-20) nm, a length of (0,5-200) µm from Ahwahnee Technology Inc. and single-walled nanotube (SWNT) with a diameter of <2 nm, a length of (5-20) µm from Shenzhen Dynanonic Co., Ltd.

were used as nanotube starting materials. Both types of nanotube were fabricated by the catalysis of hydrocarbon with chemical vapour deposition (CVD) process. Very pure copper powder (purity >99.95%) of 3 µm particle size with dendritic shape, 10 µm particle size with spherical shape from Sigma-Aldrich and 45 µm particle size with spherical shape from TLS Technik GmbH were chosen as staring copper metal powders. The metal powders as alloying elements (Sn, Zn, Ni) with 45 µm particle size, spherical shape were also used from TLS. The typical composition for fabricating bronze composites was Cu 79%, Sn 10%, Zn 3% and Ni 8%. Dendritic Copper Powders, more commonly referred to electrolytic copper powders (ECP) were produced by the electrolytic deposition method, whereas spherical metal powders are manufactured by atomization (gas/water) process. Table 1 and 2 shows the specification of nanotubes as well as metal particles.

Mixing powders

Different wt% of CNTs were mixed with metal powders inside a glove box having inert (Ar) atmosphere and milled in a planetary ball-mill machine (PM 400) for different hours with the following parameters: number of balls: 50 (stainless steel, hardness 48-52 HRC), diameter of each ball: 10 mm, ball to powder ratio 10:1, milling speed: 200 rpm. Scanning electron microscopic (SEM, JEOL 6500) investigations were carried out to evaluate the dispersion of CNT in metal powders after the milling process. Different milling times and wt% of CNTs were used to ensure better dispersion of nanotubes in different metal powders.

Composite fabrication

The mixed powders were then sintered in a hot-press sintering machine (DSP 510) filled in inert (Ar) atmosphere. The DSP 510 sintering device using direct resistance heating in the mould produces very high heating speed and the whole composite fabrication process needs maximum 15 minutes. The sintering parameters (temperature, pressure, time, heating and cooling rate) were optimized for each type of composite by approaching their theoretical densities. The densities of the composites were measured using Archimedes' principle. Sintering temperature of 750 °C and pressure of 40 MPa were used for Cu-CNT composites, where as 800 °C and 40 MPa were used for bronze-CNT composites fabrication.

Characterization of the composites

Before characterization, composite samples were ground with consecutive finer SiC abrasive papers followed by polishing with a diamond grit suspension up to one micrometer. The samples were then examined with light optical microscopy (LOM, Axio Imager from Carl Zeiss) etched with Ferric Nitrate or Chloride.

Scanning electron microscopy (SEM, JEOL 6500) attached with Energy Dispersive Xray (EDX) was used to evaluate the distribution of nanotubes in the matrix materials as well as to analyze the elements in the composite materials. The hardness of the composites was measured in the 3000 BLD Brinell hardness tester from Wilson-Wolpert using a 10 mm steel ball as an indenter with a 3000 kgf (30 kN) force.

Electrical conductivity of the composites was measured by a device called Sigmatest 2.069, which measures the electrical conductivity of non-ferromagnetic metals by measuring the eddy current based on the complex impedance of the sample. When unknown test pieces are measured, this instrument converts the complex impedance value to an electrical conductivity value.

Results and discussions

Evaluation of mixing

In the early contribution [START_REF] Uddin | Thermal Expansion Coefficient of Nanotube-Metal Composites[END_REF], the authors showed that the planetary ball mill

Evaluation of composites

After evaluating nanotube-metal mixed powders, composites were fabricated by a hotpress sintering method with optimized sintering parameters. Table 3 shows sintering parameters used for the composite fabrication and their corresponding relative densities.

The density of nanotubes was considered to be 2 g/cm³ [23]. Reducing the particle size of metals, the relative density of the composite is increased. 99.33% of the theoretical density was achieved with the composite made of 10 µm, spherical Cu particles. On the other hand, particles having dendritic shape should have less density (the more regular the shape, the denser the powder [START_REF] Pavlovi | The effect of particle structure on apparent density of electrolytic copper powder[END_REF]) compared to that of spherical structure.

Increasing CNT contents decreases the relative density of the composites due to increased porosities and voids inside the tubes and among the agglomeration of CNTs.

Table 3

The early issue [START_REF] Uddin | Thermal Expansion Coefficient of Nanotube-Metal Composites[END_REF]

Characterization of composites

Hardness

Hardness of CNT-based Cu matrix composites was measured in a Brinell scale (29 kN force and with a steel ball). It was plotted against increasing proportions of CNTs. According to D.L. Erich [START_REF] Erich | Metal Matrix Composites: Problems, Applications and Potential in the P/M Industry Int[END_REF], discontinuous MMCs formed by the mixtures of individual metal phases exhibits strengths as much as 50% higher than those predicted in theory from the strength of the individual constituents. In our case, the hardness of pure Cu fabricated by hot-press sintering was also found almost double (58 HB) compared to that of reference value (35 HB). However, the hardness decreases with increasing amount of CNTs in the composite for both CNT-Cu and CNT-Bronze composites.

Figure 3

Hardness of pure Cu and Cu-0.5wt% MWNT composites with different particle size and shape (45 µm, spherical; 10 µm, spherical and 3 µm, dendritic) is plotted in Fig. 4.

Hardness values of pure Cu composites fabricated with 45 µm, 10 µm and 3 µm remain in the similar range of 56-60 HB. However, hardness of Cu-CNT (0.5wt%) increases significantly by reducing the size of Cu particles. Obviously, there should be a certain effect of Hall-Petch relationship for the increase of hardness with particle size reduction. In our case, we did not find it significantly with pure Cu materials. However, reduction of particle size influenced a lot with nanotube incorporation. Distribution of CNTs in the matrix materials with reduced particle size of metals increases, which in turn plays a role to increase the hardness of the composites. Moreover, the dendritic structure of metal particles enhancing the CNT distribution in the matrix is also expected to improve the strength of the composite materials. composites is increased up to 42% by using 3 µm Cu metal particles. However, the best hardness value (83.5 HB) was achieved in the composite fabricated with 3 µm, dendritic Cu particles and 0.1 wt% MWNT (Fig. 5). Therefore, hardness of Cu-CNT (0.1 wt%) composites can be improved up to 47% by using 3 µm Cu particles in the composites.

Further increasing CNT contents rather decreases the hardness of the composite.

Electrical Conductivity

Improvement of electrical conductivity of a highly conductive metal is still a big question of science. A composite based on aligned, ballistic conducting carbon nanotubes embedded in a metal matrix might work as an ultra-low-resistive material [START_REF] Hjortstam | Can we achieve ultra-low resistivity in carbon nanotube-based metal composites?[END_REF]. Moreover, in a sintering body of powder metallurgy, it is difficult to achieve 100% density of the theoretical density. There will be a certain amount of porosity inside the composite. These porosities are in fact the insulation sites which reduce the conductivity of the composites. Fig. 6 shows the electrical conductivity of the pure copper composite as 58 MS/m, which is a little bit lower compared to reference conductivity (59.59 MS/m). Increasing CNT contents in the Cu composites decreases the electrical conductivity. Conductivity of Bronze composites using both MWNTs and SWNTs were investigated.

The conduction occurs essentially through the outermost nanotubes. In case of MWNTs, the interactions with the internal coaxial nanotubes may lead to variations in the electronic properties. However, metallic SWNTs can sustain huge current densities (max. 10 9 A/cm 2 ) without being damaged, i.e. about three orders of magnitude higher than copper [START_REF] White | Armchair carbon nanotubes as long ballistic conductors[END_REF]. Figure 7 shows that the electrical conductivity of CNT-Bronze composites is increased up to 10% by adding 0.1 wt % MWNTs and up to 20% by adding same amount of SWNTs. Higher CNT contents do not increase electrical conductivity more, rather decrease.

Conclusion

Carbon nanotube based copper and bronze composites were successfully fabricated by well-established hot-press sintering method of powder metallurgy. The hardness and electrical properties of the composites were evaluated and found up to 47% improved hardness of Cu by adding 0.1 wt% of MWNT in the composites, up to 20% improved electrical conductivity of Bronze by adding 0.1 wt% of SWNT in the composites. Thus, nanotubes can improve the hardness of highly conductive-low strength copper metals, whereas electrical conductivity of low conductive-highly strength copper alloys can be improved by using nanotubes as reinforcements in composite materials.

  employed a strong centrifugal force to develop high energy milling action inside the vial, which in fact involved continuous impact, welding, fracturing and rewelding of CNT-metal powders. Milling for shorter time (30 minutes) did not show good embedment of CNTs in the matrix materials and prolonged milling promoted cold-welding of metal powders producing very large metal particles. By reducing the particle size of metals, the possibility to disperse CNTs in metal powders increases. Milling of 0.5 wt% of CNTs with 3 µm, dendritic copper particles for 1 hour showed very good distribution of CNTs in the metal particles (Fig. 1.a). The dendritic shape of the metal particles with its higher specific surface area [22] in this regard is expected to enhance nanotube distribution within the dendritic arms of the metal powders. Dispersion of CNTs is also promoted by milling with different alloying elements to produce CNT-Bronze composites (Fig. 1.b). CNTs are expected to diffuse in the different metal powders during mechanical alloying.

Figure 1

 1 Figure 1

  already showed that CNTs were very uniformly distributed throughout the whole matrix and well-embedded in the matrix materials. The distribution of nanotubes was analysed from the fracture surface of the composites by scanning electron microscopy. The microstructure of the composites was also evaluated by Light Optical Microscopy (LOM). The micrograph (Fig. 2.a) showed two types of phases, one of which is completely CNT free, where CNTs are not distributed in the matrix and the other is CNT/Cu phase, where CNTs are distributed inside Cu matrix. The distribution of CNTs in the matrix depends on the mixing of CNT-metal powders [25]. If CNTs are well distributed within metal powders during the ball-milling process, more homogeneous distribution of nanotubes is obtained. However, during the solidification process, the primary phase of pure Cu is formed shifting nanotubes towards the grain boundaries, while the other phase contains CNTs in Cu matrix [26]. A clear distinction between the two phases is shown in Fig. 2.a. The primary Cu phase appears as bright, whereas other phase contains CNTs as black dots inside Cu matrix. However, some black dots may also be appeared for copper (II) Oxide. Energy Dispersive X-ray (EDX) analysis on the surface of the same composite also ensured the presence of Cu, O and C in the composite. Since copper does not form carbides, the carbon peak in EDX is an evidence of carbon nanotubes in the sample.

Figure 2

 2 Figure 2

Figure 3

 3 Figure 3 shows the Brinell hardness values of CNT-Cu and CNT-Bronze composites

Figure 4 High

 4 Figure 4

Figure 5

 5 Figure 5

Figure 6

 6 Figure 6

Figure 7 As

 7 Figure 7

Figure 1 (hoursFig. 1 Figure 2

 112 Figure 1

Fig. 2 Figure 5 HardnessFig. 5

 255 Fig. 2 Optical micrograph (a) and EDX analysis (b) of Cu-MWNT (0.5 wt%) composite

Figure 6 ElectricalFig. 6

 66 Figure 6

Figure 7 ElectricalFig. 7

 77 Figure 7

Table 1

 1

Table 2

 2

Table 1 : Specification of carbon nanotubes

 1

	Type	Average	Range of length	Method of	Producer/
		diamete (nm)	(µm)	Synthesis	Supplier
	MWNT	10-20	0.5-200	CVD	Ahwahnee
	SWNT	<2 nm	5-20	CVD	Shenzhen

Table 2 : Specification of metal particles Metal powders Particle size

 2

			Particle	Method of synthesis	Producer/
		(µm)	shape		Supplier
	Cu	3	dendritic	Electrolytic deposition	Sigma-Aldrich
	Cu	10	spherical	Gas/water atomization	Sigma-Aldrich
	Cu	45	spherical	Gas/water atomization	TLS
	Sn, Zn, Ni	45	spherical	Gas/water atomization	TLS

Acknowledgments

The authors are thankful to Fraunhofer Society and Bögra Technologie GmbH for the project support.


J Macutkevic

D Seliuta

G Valusis

J Banys

S Hens

V Borjanovic

V Kuznetsov

O Shenderova
email: oshenderova@itc-inc.org

Effect of thermal treatment conditions on the properties of onion-like-carbon based polymer composite

Keywords: A. Nano composites, B. Transport properties, E. Annealing, E. Heat treatment, Onionlike Carbon

Dielectric properties of onion like carbon (OLC) and polyurethane composite prepared using different procedures were investigated in the frequency range up to 1 MHz. We show that broadband dielectric spectroscopy is powerful tool to determine technological fingerprints in the studied materials. It is demonstrated that cured samples annealed at temperature close to the melting point (450 K) exhibit substantially higher dielectric permittivity and electrical conductivity in comparison with untreated samples. With the increase of temperature of an untreated sample, its dielectric permittivity, electric conductivity and critical frequency increase, while Maxwell-Wagner mean relaxation time and aggregate sizes of OLC decreases. Annealing of the composite at temperature close to the melting temperature produce sample with more homogeneous distribution of OLC. The temperature dependence of conductivity of the homogeneous sample is mainly caused by a weak positive temperature coefficient effect.

Introduction

Electrically percolative polymer-based composites have been attracting much attention because of their potential applications such as electroactive materials, sensitive materials, and electromagnetic coatings. Owing to the advanced electrical, thermal, and mechanical properties, various carbon nanoparticles like carbon nanotubes or carbon black have been widely studied and often used as nanofillers in the past few years. The main advantage of carbon nanotubes versus other carbon nanofillers is their extremely low percolation threshold [START_REF] Sandler | Ulta-low electrical percolation threshold in carbon-nanotube-epoxy composites[END_REF]. Nevertheless, the percolation threshold in carbon nanotubes based composites dependents on many factors such as polymer type and dispersion method [START_REF] Bauhofer | A review and analysis of electrical percolation in carbon nanotube polymer composites[END_REF]. The onion-like carbons (OLCs), consisting of stable defected multishell fullerenes, exhibit high conductivity similar to carbon nanotubes [START_REF] Kuznetsov | Electrical resistivity of grafitized ultra-disperse diamond and onion-like carbon[END_REF]. The OLC have a rare structure allowing for hierarchical assembly of conductive structural units and high concentration of electrically active defects, being able to absorb a wideband of electromagnetic radiation [START_REF] Langlet | Dipole polarizability of onion-like carbons and electromagnetic of their composites[END_REF].

Nevertheless, the first investigations of OLC based composites have shown low dielectric permittivity and electrical conductivity similar to the pure polymer matrix parameters [START_REF] Macutkevic | Dielectric properties of onion-like carbon based polymer films: experiment and modeling[END_REF][START_REF] Macutkevic | Dielectric properties of a novel high absorbing onion-like-carbon based polymer composite[END_REF][START_REF] Macutkevic | Terahertz probing of onion-like carbon-PMMA composite films[END_REF]. It was observed, however, that the dielectric permittivity and electrical conductivity drastically increase in onion-like-based composite approaching the percolation point [START_REF] Macutkevic | High dielectric permittivity of percolative composites based on onion-like carbon[END_REF]. It is interesting to note that in the OLC polyurethane (PU) composite, the percolation point of a 10% volume fraction is estimated [START_REF] Macutkevic | High dielectric permittivity of percolative composites based on onion-like carbon[END_REF], and this is lower than a prediction from an effective medium approximation for the inclusions with a spherical shape (the percolation threshold for spherical inclusions is close to 1/3 volume fraction [START_REF] Rychetsky | Dielectric properties of microcomposite ferroelectrics[END_REF]). Carbon black topology is similar to OLC. For carbon black, the percolation threshold depends on the preparation technology and can be similar to carbon nanotubes, i.e., it can be also very low [START_REF] Adrianse | High-dilution carbon black/polymer composites: hierarchical percolating network derived from Hz to THz AC conductivity[END_REF]. Therefore, for OLC applications it is very important to find optimal technological ways to further reduce the percolation threshold [START_REF] Macutkevic | High dielectric permittivity of percolative composites based on onion-like carbon[END_REF] and to study the effect of technological procedures on the electrical/dielectrical properties of the structures.

In this paper, we therefore, investigate an influence of preparation conditions on OLC/PU dielectric properties in order to find a most favorable technology for producing coatings with a low percolation threshold. In more detail, we explore the role of heating of the samples during mixing of OLC and PU solution as well as annealing effect of the cured films.

Sample preparation technology and experimental part

The OLC material was produced by annealing of detonation nanodiamond powder in vacuum at 1800 K, as it described by Kuznetsov et al. [START_REF] Kuznetsov | Electrical resistivity of grafitized ultra-disperse diamond and onion-like carbon[END_REF]. Nanodiamonds were purchased from New Technologies, Chelyabinsk. Average primary particle size of DND was 4 nm and the average volumetric aggregate size was 120 nm. To achieve fully transformed DND into OLC, DNDs were heated in vacuum 10 -4 torr at 1800 K for 2 h. Size of primary OLC particles is around 6---7 nm, while average aggregate size is 130 nm for OLC dispersed in N-methylpyrrolidone solvent as measured by dynamic light scattering technique.

Samples of OLC were mixed with commercial formulation of oil-based polyurethane "Minwax clear satin" containing 60% of volatile compounds. All samples contained 15wt.% of OLC. Six samples were prepared under different thermal treatment conditions (more detailed information can be found in Table 1). First, after manual mixing with the OLC powder, the polymer suspension was magnetically stirred at 400 rpm during 24 hrs at 50 °C. Then, before casting a suspension, samples 2 and 3 were additionally heated at 70 o C and 90 o C, correspondingly, on a hot plate with a purpose to improve uniformity in distribution of OLC in a polymer solution with a lower viscosity. Then suspensions were casted on a Teflon substrate and dried at ambient conditions at 45 °C overnight. Half of the dried samples ("b" series) were additionally annealed at 150 o C during 1 hr under pressure in an oven. A motivation for this treatment originated from our observation that after testing of dielectric properties of the samples at high temperatures (for example, 450 K), dielectric constants and conductivities of the samples, when cooled to the room temperature, were increased. Thus, we introduced an annealing stage during sample preparation. Thickness of OLC/PU films varied between 250 and 700 µm. m. Distribution of OLC in PU matrix was characterized by scanning electron microscope SEM Zeiss Supra 25. To characterize the samples we have employed broadband dielectric spectroscopy. This technique was already proved to be powerful tool to characterize these materials [START_REF] Macutkevic | Dielectric properties of onion-like carbon based polymer films: experiment and modeling[END_REF][START_REF] Macutkevic | Dielectric properties of a novel high absorbing onion-like-carbon based polymer composite[END_REF][START_REF] Macutkevic | High dielectric permittivity of percolative composites based on onion-like carbon[END_REF]. The complex dielectric permittivity ε * = ε'-iε" was measured by a capacitance bridge HP4284A in the frequency range 20 Hz-1 MHz. Measurements have been performed on cooling and heating at a rate of about 0.5 K/min. Each measurement started from heating, and after reaching 450 K, the sample was kept at this temperature for ~0.5hr, and then the samples were cooled down to room temperature.

Results and discussion

Figure1 illustrates SEM images of the top view of the samples 3a and 3b at the same magnification. As can be seen from the images, OLCs are uniformly distributed within the matrix and OLC aggregate sizes are smaller in the annealed sample 3b.

The frequency dependence of dielectric permittivity ε' and electrical conductivity ' at room temperature for composites prepared under different conditions with OLC concentration of 15% is shown in Fig. 2. It can be seen that both dielectric permittivity ε' and electrical conductivity ' is strongly influenced by preparation conditions. In particular, it is more pronounced at low frequencies. All annealed samples (1 b -3 b) exhibit higher electrical conductivity ' and dielectric permittivity ε' in comparison with not treated samples in the frequency range up to 1MHz. The highest electrical conductivity ' and dielectric permittivity ε' are observed in 3b sample. The increasing of dielectric permittivity ε' at low frequencies (below 300 Hz) in the annealed sample is obviously caused by Maxwell-Wagner relaxation [START_REF] Liu | Dielectric properties and Maxwell-Wagner relaxation of compounds ACu 3 Ti 4 O 12[END_REF]. In order to explain an effect of annealing on the dielectric properties, measurements of dielectric properties were performed under heating at different frequencies. These results for the most representative samples 3a (not annealed) and 3b (annealed) are shown in Fig. 3. As one can see, the dielectric permittivity ε' and electrical conductivity ' increase on heating in both annealed and untreated samples, however, in the untreated sample the increase is more pronounced. It is worth noting that more pronounced changes of dielectric properties take place at low frequencies. The electrical conductivity also exhibits a kink in sample 3b at nearly 380 K. It is probably associated with the fact that polymer molecules can move more freely at this temperature, because conductivity and dielectric permittivity of pure polyurethane also exhibit kinks close to this temperature. Usually, such temperature is associated with alfa (glass) transition in polymers [START_REF] Alvarez | Glass transition temperature and fragility in a side-chain liquid crystalline polymer: a study by TSDC and DSC[END_REF]. To get deeper insight into processes, frequency dependence of dielectric permittivity ε' and electrical conductivity ' on the heating of composites 3a and 3b are plotted and presented in Fig. 4a and Fig. 4b, respectively. It can be seen that the dielectric permittivity ε' measured at room temperature strongly decreases with the frequency, meanwhile electrical conductivity strongly increases with frequency in the untreated sample. Thus, it is not possible to estimate DC conductivity in our frequency range. On further heating of untreated sample frequency-independent (DC) conductivity becomes observable only at low frequencies.

To analyze the data, we have used an approach in which the frequency dependence of can be fitted according to the Almond-West type power law with a single exponent:

= DC +A s , ( 1
)
where DC stands for the DC conductivity and A s is the AC contribution. It is evident from Fig. 4a and 4b that the DC conductivity DC coincides with conductivity at 129 Hz in annealed samples (3b) and in untreated sample (3a) at higher temperatures. In the untreated sample the low frequency dielectric dispersion (below 1 kHz), mentioned above as Maxwell-Wagner relaxation, is strongly temperature-dependent. The Maxwell-Wagner relaxation can be described by Cole-Cole formula [START_REF] Cole | Dispersion and Absorption in Dielectrics I. Alternating Current Characteristics[END_REF]:

α ωτ ε ε ε - ∞ + ∆ + = 1 * ) ( 1 CC i , ( 2
)
where 0 1 are the parameters of width of Cole-Cole distributions functions, CC is the mean and most probable Cole-Cole relaxation time, is the dielectric strength, is the high frequency dielectric permittivity. However, in our case the Cole-Cole fit for low frequency dielectric dispersion cannot be performed, because the mean relaxation CC is very long and 1/ CC is substantially below of our low frequency limit. Nevertheless, it is easy to conclude that the mean relaxation CC decreases on cooling in both annealed and untreated sample. In the Maxwell-Wagner relaxation case, CC =RC, where R and C is the resistance and the capacitance of grains, respectively. The capacitance of OLC aggregate, a sphere with radius r, can be described as C=4 0 r. Therefore, the decrease of the mean Cole-Cole relaxation time CC is related to the decrease of the mean size (radius) of the OLC aggregate.

From the Almond-West type fit it is possible to calculate critical frequency f c , at which ( ) deviates from the dc plateau. Experimental value for f c , have been defined as the frequency at which the value of conductivity is 10% higher than the DC conductivity value. The results are plotted in Fig. 5. As can be seen, the critical frequency value f c of untreated samples (2a and 3a) at room temperature is very low, and falls below low frequency limit (20 Hz) of the equipment used. On heating above 350 K, the critical frequency for these samples strongly increases, but still remains lower in comparison with annealed samples.

To estimate an activation energy the temperature dependence of at 129 Hz (DC conductivity) for different OLC samples are presented in Fig. 5. The activation energy E of the conductivity and the conductivity prefactor 0 was evaluated using the Arrhenius law: (

The obtained parameters are listed in Table 2. It is evident that the activation energy E of the DC conductivity is higher in untreated samples and pure PU than that in annealed samples. The physics behind that can be associated with an electrical transport via electron diffusion through conducting OLC aggregates and (or) tunneling through a nonconductive (at low temperature) polyurethane matrix. The DC conductivity is a result mainly of electron tunneling through nonconductive polyurethane matrix, because no continued network of OLC clusters is observed in the composite (Fig. 1). Knowing that the activation energy E is a potential barrier for the carriers tunneling, it is reasonable that it is lower in more homogeneously annealed samples (Table 2). The average time fl , which a charge carrier needs to travel through the OLC aggregate, is directly proportional to the aggregate size d. The biggest OLC size is observed in an untreated sample (Fig. 1), therefore, it is reasonable to assume that:

fl c f τ 1 ∝ . ( 4
)
For frequencies f<f c charge carriers can travel over different OLC aggregates within one electromagnetic wave period. In contrast, for frequencies above f c the charge carriers can move mainly within a length of one OLC aggregate. Therefore, the critical frequency f c and its temperature dependence reflect the mean OLC aggregate size. Hence, the observed increase in critical frequency f c above 350 K manifests that OLC aggregate size is decreased. It can be thus concluded that annealing is a very important treatment for obtaining a homogeneous composite.

It deserves to emphasize that in contrast to heating measurements the temperature dependence of electrical conductivity and dielectric permittivity of both annealed and untreated samples on cooling is very weak (Fig. 6). Pronounced discontinuity between electrical conductivity and dielectric permittivity data on heating and on cooling indicate that the main composite homogenization process happens during isothermal annealing at 450 K. This temperature is close to the melting point of pure polyurethane (melting temperature of pure polyurethane was established in additional experiment and found to be 489 K). A small increase of electrical conductivity on cooling in an untreated sample (above temperature 400 K) and in an annealed sample obviously is caused by a positive temperature coefficient effect [START_REF] Strümpler | Polymer composite thermistor for temperature and current sensors[END_REF]. The effect results from the break of the conductive links due to the thermal expansion of the polymer matrix. Some small decrease of dielectric permittivity ε' and electrical conductivity on cooling is observed in an untreated sample at low temperatures (below 400 K). This phenomenon is related to a negative temperature coefficient effect [START_REF] Strümpler | Polymer composite thermistor for temperature and current sensors[END_REF], which is caused by residual redistribution OLC aggregate in polymer matrix.

Finally, we have considered the frequency dependence of electrical conductivity on cooling. To analyze the data Eq. ( 1) was used. It was found that the critical frequency only slightly increases on cooling as it is seen from Fig. 7. The low-frequency dielectric dispersion (Maxwell-Wagner relaxation) is also almost stable on cooling. Therefore, one can note that the diffusion of OLC aggregates through polyurethane matrix mainly vanishes in cooling measurement and, as a result, a homogeneous composite is obtained.

Conclusion

The influence of a preparation procedure on dielectric properties of OLC-polyurethane composites was investigated. Two procedures were implemented: (i) mechanical stirring of OLC and PU in a solvent at increased temperature and (ii) annealing of the cured samples. It was concluded that heating of the suspension during mixing of OLC and PU has not so significant effect on the enhancement of the dielectric properties of the composites, as compared to an annealing of the cured sample. Samples annealed close to the melting point (450 K) exhibit substantially higher dielectric permittivity and electrical conductivity in comparison with an untreated sample. Dielectric permittivity, electric conductivity and critical frequency increased in the untreated sample during heating while Maxwell-Wagner mean relaxation time and average OLC aggregate size decreased. All parameters exhibited an offset close to the alpha (glass) transition temperature (380 K). Annealing of composites close to the melting temperature results in more homogeneous distribution of OLC. The temperature dependence of conductivity of a homogeneous sample is mainly caused by a weak positive temperature coefficient effect. Figure 2. Frequency dependence of the real parts of complex dielectric permittivity ' and complex electrical conductivity prepared under different conditions at room temperature. Figure 3. Temperature dependence of the real parts of complex dielectric permittivity ' and complex electrical conductivity of 3a (a) and 3b (b) sample on heating. Figure 4. Frequency dependence of the real parts of complex dielectric permittivity ' and complex electrical conductivity of 3a (a) and 3b (b) sample on heating. Figure 5. Temperature dependence of conductivity at 129 Hz and critical frequency f c on heating. Figure 6. Temperature dependence of the real parts of complex dielectric permittivity ' and complex electrical conductivity of 3a (a) and 3b (b) sample on cooling. Figure 7. Temperature dependence of DC conductivity and critical frequency f c on cooling.

List of Captions for figures and tables

Figure 1 .

 1 Figure 1. SEM images of OLC/PU composites prepared under different conditions: untreated sample 3a (a) and sample 3b annealed at 150 o C (b).Figure2. Frequency dependence of the real parts of complex dielectric permittivity ' and complex electrical conductivity prepared under different conditions at room temperature. Figure3. Temperature dependence of the real parts of complex dielectric permittivity ' and complex electrical conductivity of 3a (a) and 3b (b) sample on heating. Figure4. Frequency dependence of the real parts of complex dielectric permittivity ' and complex electrical conductivity of 3a (a) and 3b (b) sample on heating. Figure5. Temperature dependence of conductivity at 129 Hz and critical frequency f c on heating. Figure6. Temperature dependence of the real parts of complex dielectric permittivity ' and complex electrical conductivity of 3a (a) and 3b (b) sample on cooling. Figure7. Temperature dependence of DC conductivity and critical frequency f c on cooling.

Table 1 .

 1 List of investigated OLC/PU composites.

Table 2 .

 2 Parameters of Arrhenius law obtained from DC conductivity.

	sample	OLC-PU mixing conditions	treatment of a cured sample
	1a	1 day at 50 0 C	untreated
	1b	1 day at 50 0 C	annealing at 150 0 C
	2a	1 day at 50 0 C + 70 0 C 2hr	untreated
	2b	1 day at 50 0 C + 70 0 C 2hr	annealing at 150 0 C
	3a	1 day at 50 0 C + 90 0 C 2hr	untreated
	3b	1 day at 50 0 C + 90 0 C 2hr	annealing at 150 0 C

Acknowledgements

This work was partially supported by NATO Collaborative Linkage Grant (Reference No.CBP.EAP.CLG983910). The work was also partially supported by the NATO Science for Peace program (grant SfP 981051). ITC acknowledges the partial support through Air Force Office of Scientific Research under grant N66001-04-1-8933.