J A Kepler
email: jk@me.aau.dk

SIMPLE STIFFNESS TAILORING OF BALSA SANDWICH CORE MATERIAL

Keywords: B: Mechanical properties, C: Elastic properties, C: Sandwich structures, C. Laminate theory, Balsa wood core

A concept for improving the shear stiffness properties of balsa core material for sandwich structures is presented. The concept is based on utilization of the strongly orthotropic properties of the balsa wood, applying an appropriate transverse layup sequence. The effective core material shear modulus is modeled using basic laminate theory. This is subsequently validated through sandwich beam bending and lap shear experiments. Compared to the standard balsa core systems, a substantial increase in the shear stiffness is demonstrated, whereas the transverse stiffness is reduced. The concept is suitable for mass production, using standard plywood fabrication technology.

INTRODUCTION

Balsa wood is frequently used for core material in sandwich structures, due to the favorable stiffness/weight ratio. Descriptions of the general use of balsa wood in sandwich structures may be found in standard handbooks, such as [START_REF] Zenkert | The Handbook of Sandwich Construction[END_REF] by Zenkert. It is usually oriented with the grain-direction normal to the extension plane of a sandwich panel -one reason for this is retardation of decomposition due to moisture in case of a face-sheet puncture. Also, the high axial (along the grain) stiffness of balsa wood improves the resistance of the sandwich panel to local indentations and face-sheet wrinkling and dimpling. In a recent study by Tagarielli et al. [2], the axial compression properties were shown to be highly strain-rate sensitive, providing a significant increase in dynamic strength compared to the quasi-static behavior. Furthermore, the study demonstrated, in agreement with several previous studies, that the axial compressive strength properties are clearly superior to those of PVC foams of comparable densities.

However, the orientation of the axial direction transverse to the sandwich panel plane gives a poor utilization of the orthotropic properties of the wood with respect to the transfer of transverse (shear) force through the core, and makes the core prone to sudden transverse splitting. A concept for improving the stiffness using balsa elements was presented by Bekisli and Grenestedt [START_REF] Bekisli | Experimental Evaluation of Balsa Sandwich Core with Improved Shear Properties[END_REF].This provided an improved shear stiffness in two primary transverse planes, but required elaborate machining of the balsa elements.

Conceptually, the use of balsa wood in such applications is similar to any periodical distribution of a highly orthotropic material, in order to obtain desirable properties. In the extreme case, an orthotropic material may be approximated by a unidirectional strut, where, of course, the sandwich panel core becomes a truss web structure. A paper by Liu et al. [START_REF] Liu | Design Optimization of truss-cored sandwiches with homogenization[END_REF] elaborates on homogenization of properties for these types of core; the relative performance, as compared to other sandwich core concepts, was studied by Wicks and Hutchinson ([5] and [START_REF] Wicks | Performance of sandwich plates with truss cores[END_REF]).

Simplified Concept

The concept presented here provides an improvement of the shear properties, particularly the effective core shear modulus G xy . It will be effective in only one transverse plane, but is straightforward and economically feasible. The principle is illustrated in figure 1. The balsa plates were purchased through a modeling supplies retailer and matched to provide the best possible consistence (6 plates, 121 kg/m 3 < ρ < 138 kg/m 3 ). In most practical cases, the density variation may be considerably larger.

PARAMETERS FOR PREDICTION OF SHEAR STIFFNESS

Through the use of classical (plane) laminate theory (see e.g. Jones [START_REF] Jones | Mechanics of Composite Materials[END_REF]), the effective shear modulus may be calculated for a stacked balsa core at a given stacking sequence.

The individual lamina may be constitutively described for a given angular orientation through transformation of the principal-axis constitutive relation:

[ ] = =
σ Q Q Q Q Q Q (1)
where, due to symmetry, Q 12 = Q 21 . The formulation is reduced from the full 3Dformulation, under the assumption that out-of-plane effects are insignificant. For the present case, it means that the analysis is invalid for the yz-plane core shear stiffness.

The components of the constitutive matrix, in terms of engineering constants are found

υ υ υ υ υ υ - = - = = E E Q Q and 12 66 G Q = (2)
In the arbitrary orientation, the constitutive relation is found through [START_REF] Bekisli | Experimental Evaluation of Balsa Sandwich Core with Improved Shear Properties[END_REF] where again the constitutive matrix is symmetric, and [T] is the transformation matrix.

[ ] [ ] [ ][ ] = = = - - xy y x T xy y x xy y x xy y x T Q T Q Q Q Q Q Q Q Q Q Q γ ε ε γ ε ε γ ε ε τ σ σ
The derivation may be found in [START_REF] Jones | Mechanics of Composite Materials[END_REF], but is summarized here for completeness.

For the present subject, the term of particular interest is 66 Q , which directly links shear strain and stress. In this capacity, it will be a direct measure of the shear stiffness of the core material at arbitrary angles α, and will be referred to as the effective shear modulus G xy .

The 22 Q term, being a measure of the transverse stiffness, is of relevance in areas of high transverse load intensity, where core indentation may be an issue, and in suppression of compressive face-sheet wrinkling. Deviation from the typical fibre angle (α = 90°) will lead to an increasing effective shear modulus, but this will be gained at the expense of transverse compressional stiffness. In this sense, the specific core layup configuration must be a compromise, dictated by the proportions of global shear load, local transverse load intensity and pure bending.

Engineering stiffness constants E 1 , E 2 , ν ν ν ν 12 , G 12
As implied, the elementary properties of the balsa wood in the principal material coordinates (axial, radial and tangential for wooden materials) must be measured or estimated through empirical formulas, for use in the laminate theory expressions.

Properties in the radial and tangential directions are, by necessity, averaged, since the lengthwise sawing orientation of the balsa wood is variable; the approximate average modulus is referred to as the transverse modulus.

In the book by Gibson and Ashby [START_REF] Gibson | Cellular Solids -Structure & Properties[END_REF], readily available formulas for the stiffness constants are presented, and it is relevant to use these. However, this leads to a brief digression, as the formulas suggested in [START_REF] Gibson | Cellular Solids -Structure & Properties[END_REF] are evaluated against the directly and indirectly available experimental data. The approximate chronological progression is recorded in table 1.

Initially, the formulas (given in row a in table 1) were used to produce data for the specific balsa wood (row b in table 1). Clearly, the dominant factors are E 1 and G 12 .

Using classical laminate theory, the dependence of the apparent shear modulus G xy on the angle α is shown in figure 4 and in the final comparison, presented in figure 11. The transverse modulus E 2 seemed unrealistically low, and direct measurements on the 5 mm thick balsa plates was impractical. Therefore, measurements were performed on a larger plank of balsa wood. For this specimen, the density was 78 kg/m 3 , and the primary extensional moduli were measured (row c in table 1). These were compared with corresponding predicted values (row d in table 1), and the formulas were adjusted accordingly (row e in table 1); in particular, the exponent 3 in the original expression for E 2 was adjusted to 2.3, producing an acceptable fit for the 78 kg/m 3 specimen.

A direct measurement of the shear modulus G 12 for the 129 kg/m 3 balsa was obtained by a lap shear experiment (see below), and used in the modeling. An indirect measurement, by 3 point bending of a sandwich beam (see below), gave a value of 147 MPa.

Axial modulus

E 1 [MPa] Transverse modulus E 2 [MPa]
Poisson´s ratio -components of the stiffness matrix for a single "ply" of balsa wood, shown for angles 0° α 90°.

ν 12 [ ] Shear Modulus G 12 [MPa] a): [8]
One notes the expected dependence on the layup angle α..

EXPERIMENTAL INVESTIGATION

For verification of the predicted shear stiffness, simple lap shear experiments and 3point sandwich beam bending experiments were conducted.

Preparation of Test Specimens

In the present investigation, the parameters H c = 20 mm, b = 5 mm were kept constant. The specimens used for the experimental verifications invariably consisted of six layers with a stacking sequence of [± α] 3 . For specimens where α 90°, this introduces a small amount of shear/twist coupling, which is ignored in the following.

For manufacture of a specimen, strips of balsa wood were cut at the appropriate angle from 5 mm thick plates of balsa wood, with a strip width slightly larger than the final core height (22 mm for a final core height H c = 20 mm). The six layers for a given specimen were then bonded together using a PVA wood glue (Akzo Nobel, Cascol 3301). The glue was diluted with water (1:1 by volume), for facilitation of the glue application (using a painting brush). After adhesion, the specimens were left to dry at room temperature for one week, whereupon the specimens were trimmed to the final height.

Face sheets for sandwich beam specimens were cut from mild steel plate:

Thickness t f = 1 mm, width B = 6b = 30 mm, modulus E f = 210 Gpa Immediately prior to bonding of the face sheets to the core, the face sheets were lightly bead-blasted and degreased using acetone. For bonding sandwich face sheets or lap shear bars to the core specimens, an epoxy adhesive (Araldite 2011) was used. The specimens were kept under pressure for 24 hours after adhesive assembly, then left to post-cure for another 24 hours, i.e. 48 hours at room temperature.

Lap shear experiments

The shear stiffness for two configurations (α = 90°, 45°) was investigated using a standard lap shear test. The experimental setup is shown in figure 5. The specimen dimensions were: length L = 240 mm, width B = 30 mm, height H c = 20 mm One weakness of the lap shear tests, as the one illustrated in figure 5, is that a uniform shear stress over the length of the specimen cannot be fulfilled near the free edges of the specimen. In the present case, the balsa core material is highly orthotropic, with the direction of maximum stiffness of the individual core-layer indicated by the angle α, and a specimen length detriment has been introduced, according to figure 6. Fig. 7: Shear stress/strain curves for specimens with α = 90° and 45°.

Calculation of the effective shear modulus is done as follows: [START_REF] Liu | Design Optimization of truss-cored sandwiches with homogenization[END_REF] where: F: shear force (see figure 6) L rep : representative length, as discussed above, B: specimen width γ: shear strain The shear strain γ is, for small deformations, calculated by

γ γ τ ⋅ ⋅ = ⋅ = ⋅ = B L F G B L F G rep xy rep xy

Sandwich beam experiments

As mentioned above, the free edges of the lap shear test specimens introduce a zone of nonuniform shear stress. For compensation, the representative length of the specimen was reduced according to the orthotropic characteristics of the specimen. For a separate verification of the core properties, sandwich beam specimens subjected to symmetrical 3-point bending were studied, using stacked balsa core at three different angles (α = 90°, 67.5°, 45°). The beam geometry is shown in figure 8. As indicated in figure 8, core inserts of steel were used at the load introduction locations in order to obtain negligible core compression. The face-sheets were manufactured from 1 mm thick steel plate.

The test was performed in a 20 kN test machine with the specimen supported on a moving crosshead, as shown in figure 9. In the first three tests (# = 1) in figure 10, the specimens were moderately loaded to provide data for the specimen compliances. In the second three tests (# = 2), the specimens were loaded to failure.

Of immediate interest, the failure strength varies only moderately between the three specimens, with a nominal failure shear stress τ fail between 3.8 and 4.2 MPa. However, all three failed specimens exhibited failure in the interface between the core and the face-sheets, where the interface bond rather than the core strength has been the determining factor. As the present investigation deals primarily with the shear stiffness, this is of secondary importance.

Using first order shear deformation theory (FOSDT), the measured specimen compliance was compared with a calculated compliance, with the shear stiffness G xy of the core material as the single unknown quantity.

The apparent shear stiffnesses were Fig. 11: Comparison of apparent shear modulus G xy (α) found by: laminate theory with material properties estimated according to [START_REF] Gibson | Cellular Solids -Structure & Properties[END_REF], lap shear test, sandwich beam 3-point bending test and, finally, laminate theory using material parameters measured or found through modified versions of formulas suggested in [START_REF] Gibson | Cellular Solids -Structure & Properties[END_REF].

When using base orthotropic properties predicted directly according to [START_REF] Gibson | Cellular Solids -Structure & Properties[END_REF], the level of shear stiffness is substantially and consistently higher than the measured values.

Instead, the values measured directly or extrapolated from measured data may be used: E 1 = 2700 MPa, E 2 = 83 MPa, G 12 = 134 Mpa, ν 12 = 0.35 (0.23-0.49) and one obtains a better agreement.

At α = 90°, where the direct result for the lap shear test was used as input for the laminate theory model, the agreement between theory and experimental data must obviously be good. In the intermediate range, specifically comparing the result from 3point sandwich beam bending at α = 67.5° with the prediction from laminate theory, the agreement is less convincing, but still within reasonable limits.

DISCUSSION AND CONCLUSIONS

The results of the modeling and measurements are not in full agreement -one reason may be found in the Love-Kirchoff assumptions of classical laminate theory which state that deformation-, strain-and stress-effects out of the stacking plane may be ignored.

With the thickness (b = 5 mm) and the relative softness of the balsa core "plies", these assumptions may not permit an adequate description of the deformation state.

Applying First-Order Shear Deformation Theory (FOSDT) for extracting the shear modulus in the 3-point bending tests, the shear stress distribution was assumed uniform across the core. Furthermore, the assumptions of a simple FOSDT are not properly satisfied in the transition region between steel inserts and balsa core; the face-sheet shear stiffness contribution can not be ignored completely. This, as expected, produces a higher apparent shear stiffness as compared to the lap shear test value.

The use of formulas from "Cellular Solids" by Gibson and Ashby [START_REF] Gibson | Cellular Solids -Structure & Properties[END_REF] necessitated some elaboration on the appropriateness of said formulas -in particular, prediction of the transverse modulus deviated significantly from the measurable values. This formula is based on representation of the cell wall deformation by idealized beam-bending (hence the exponent 3), which may not be appropriate. Only a moderate amount of experimental data is available in the low-density range ( [START_REF] Gibson | Cellular Solids -Structure & Properties[END_REF], figure 10.12), and the scatter is considerable.

The method of manufacture involved diluted PVA glue. It may be that the exposure to excess moisture degraded the mechanical properties of the balsa wood (see e.g. [START_REF] Strand | Wet Balsa Core[END_REF]).

However, as the balsa wood is relatively open for diffusion and was stored in a dry room at room temperature for one week after core manufacture, this is deemed unlikely.

The measured apparent shear stiffnesses were 134-147 MPa for α = 90°, 350 MPa for α = 67.5° and 655-670 MPa for α = 45°, indicating an approximately fourfold increase in shear stiffness. This was adequately supported by simple modeling. Manufacture can be automated without disproportionally large waste or time consumption. The compressional stiffness will inevitably be significantly reduced, compared to a 90° grain-to-plane angle.

The fracture morphology has not been the primary issue of the present investigation, but studies of the failed specimens indicate that stacking at alternating angles ± α increases the resilience, as the fibres of crossing layers of balsa core tend to bridge across transverse core cracks.

Fig. 1 :Fig. 2 :Fig. 3 :

 123 Fig. 1: Comparison (balsa wood fiber direction indicated by thin-line hatching)

Fig. 5 :

 5 Fig. 5: Lap shear test setup. The parts are: a -steel shear bars, b -shear specimen

Fig. 6 :

 6 Fig. 6: Determination of length detriment according to the specimen grain direction,

  lap shear plate displacement (as measured using clip gauges, see figures 5, 6) H c : specimen height The corresponding effective shear moduli, derived from the lap shear experiments as the limit secant moduli for γ 0, were: G xy (α = 45°) = 655 Mpa, G xy (α = 90°)= 134 Mpa

Fig. 8 :

 8 Fig. 8: Sandwich beam specimen. L = 300 mm, L 1 = 130 mm, L ins = 20 mm, B = 30

Fig. 9 :

 9 Fig. 9: Closeup picture of 3-point sandwich beam bending test, showing:

Fig. 10 :

 10 Fig. 10: Force / deflection curves for three-point sandwich beam bending tests. Fα,# indicates the #th test of a specimen with core angle α.

G

  xy (α = 45°)= 670 Mpa, G xy (α = 67.5°)= 350 Mpa, G xy (α = 90°)= 147 Mpa COMPARISON OF RESULTS In figure 11, predicted values of the apparent shear modulus G xy are presented alongside the values inferred from the experiments, corresponding to the test specimen range of angles α.

Table 1 :

 1 Derivation of engineering stiffness constants. In all formulas: cell wall density ρ s = 1500 kg/m 3 and axial cell wall modulus E

	N/A	N/A	N/A	134

s = 35 GPa (

[START_REF] Gibson | Cellular Solids -Structure & Properties[END_REF]

)

Using the expressions in table 1 row e, the following engineering constants are predicted for ρ = 129 kg/m 3 : E 1 = 2700 MPa, E 2 = 83 MPa (67 -99 MPa), ν 12 = 0.35.

The shear modulus as measured directly through a lap shear experiment on the relevant material density, yielded G 12 = 134 MPa (table 1 row f). For a detailed account of the experimental procedure, see below.

Table 1

 1

	row b
	Lap shear test
	3 pt bending
	Table 1 rows e, f

2.7 -3.9 0.23 -0.49


J González-Julián

Y Iglesias

A C Caballero

M Belmonte
email: mbelmonte@icv.csic.es

L Garzón

C Ocal

P Miranzo

M I Osendi

Multi-scale electrical response of silicon nitride/multi-walled carbon nanotubes composites

Keywords: A. Carbon nanotubes, A. Ceramic-matrix composites (CMCs), B. Electrical properties, D. Atomic force microscopy (AFM), Silicon nitride

Dense silicon nitride (Si 3 N 4 ) composites with various amounts (0 -8.6 vol%) of multiwalled carbon nanotubes (MWCNTs) are electrically characterised by combining macroscopic dc-ac and nanoscale conductive scanning force microscopy (C-SFM) measurements. In this way, a coherent picture of the dominant charge transport mechanisms in Si 3 N 4 /MWCNTs composites is presented. A raise of more than ten orders of magnitude in the electrical dc conductivity compared to the blank specimen is measured for MWCNTs contents above 0.9 vol%. Semiconductor and metallic-like behaviours are observed depending on both the temperature and the MWCNTs content. Macroscopic measurements are further supported at the nanoscale by means of C-SFM. The metallic-type conduction is associated to charge transporting along the nanotube shells, whereas the semiconductor behaviour is linked to hopping conduction across nanotube-nanotube contacts and across intrinsic defect clusters within the nanotubes.

Introduction

Silicon nitride (Si 3 N 4 ) based materials are well known for their extraordinary strength and hardness at high temperatures that make them suitable for applications under wear conditions, like cutting tools or ball bearings [START_REF] Riley | Silicon nitride and related materials[END_REF]. The full potentiality of these materials is only partially achieved by the difficulty inherent to their machining. Recent works [START_REF] Liu | Effect of the electrical discharge machining on strength and reliability of TiN/Si 3 N 4 composites[END_REF][START_REF] Guo | The effect of different sintering additives on the electrical and oxidation properties of Si 3 N 4 -MoSi 2 composites[END_REF] have shown the possibility of the efficient electric discharge machining (EDM) for Si 3 N 4 materials when adding an electric conductive second phase, such as TiN [START_REF] Liu | Effect of the electrical discharge machining on strength and reliability of TiN/Si 3 N 4 composites[END_REF] or MoSi 2 [START_REF] Guo | The effect of different sintering additives on the electrical and oxidation properties of Si 3 N 4 -MoSi 2 composites[END_REF]. The amount of such second phases required to get a reasonable reduction in resistivity was around 30 vol%. Besides, it has been pointed out that problems caused by static electricity in Si 3 N 4 bearings could be avoided by the addition of highly conducting carbon nanotubes (CNTs) [START_REF] Tatami | Electrically conductive CNTdispersed silicon nitride ceramics[END_REF].

Therefore, the change from highly electric insulator to conductor in Si 3 N 4 materials without degrading other properties would have great technological interest.

The step raise in electrical conductivity is the common trend observed in ceramic/CNTs composites once percolation of the CNTs network is achieved [START_REF] Tatami | Electrically conductive CNTdispersed silicon nitride ceramics[END_REF][START_REF] Zhan | Electrical properties of nanoceramics reinforced with ropes of single-walled carbon nanotubes[END_REF][START_REF] Cs | Development of CNT/Si 3 N 4 composites with improved mechanical and electrical properties[END_REF][START_REF] Shi | Effect of multiwall carbon nanotubes on electrical and dielectric properties of yttria-stabilized zirconia ceramic[END_REF][START_REF] Poorteman | Colloidal processing, hot pressing and characterisation of electroconductive MWCNT-alumina composites with compositions near the percolation threshold[END_REF][START_REF] Ahmad | Electrical conductivity and dielectric properties of multiwalled carbon nanotube and alumina composites[END_REF][START_REF] Rul | Percolation of single-walled carbon nanotubes in ceramic matrix nanocomposites[END_REF][START_REF] Flahaut | Carbon nanotubemetal-oxide nanocomposites: Microstructure, electrical conductivity and mechanical properties[END_REF]. Due to their high aspect ratio, quite low percolation thresholds, between 0.64 and 3.0 vol%, have been reported for ceramic/CNTs composites [START_REF] Shi | Effect of multiwall carbon nanotubes on electrical and dielectric properties of yttria-stabilized zirconia ceramic[END_REF][START_REF] Poorteman | Colloidal processing, hot pressing and characterisation of electroconductive MWCNT-alumina composites with compositions near the percolation threshold[END_REF][START_REF] Ahmad | Electrical conductivity and dielectric properties of multiwalled carbon nanotube and alumina composites[END_REF][START_REF] Rul | Percolation of single-walled carbon nanotubes in ceramic matrix nanocomposites[END_REF]. Conductivity values after percolation seem to depend much on the CNT type and purity as well as on the composite processing procedure, which may damage the nanotubes. For single-walled carbon nanotubes (SWCNTs) containing alumina composites, electrical conductivities ( ) of 10 3 S•m -1 have been reported [START_REF] Zhan | Electrical properties of nanoceramics reinforced with ropes of single-walled carbon nanotubes[END_REF], whereas in composites containing multi-walled carbon nanotubes (MWCNTs) much lower values, in the order of 10 S•m -1 , have been attained [START_REF] Tatami | Electrically conductive CNTdispersed silicon nitride ceramics[END_REF][START_REF] Cs | Development of CNT/Si 3 N 4 composites with improved mechanical and electrical properties[END_REF][START_REF] Shi | Effect of multiwall carbon nanotubes on electrical and dielectric properties of yttria-stabilized zirconia ceramic[END_REF][START_REF] Poorteman | Colloidal processing, hot pressing and characterisation of electroconductive MWCNT-alumina composites with compositions near the percolation threshold[END_REF][START_REF] Ahmad | Electrical conductivity and dielectric properties of multiwalled carbon nanotube and alumina composites[END_REF]. For MgAl 2 O 4 /SWCNTs composites processed by in-situ catalytic chemical vapour deposition (CVD) [START_REF] Rul | Percolation of single-walled carbon nanotubes in ceramic matrix nanocomposites[END_REF], was well fitted by the scaling law of the percolation theory as = k(p-p c ) t , giving a percolation threshold, p c , of 0.64 vol% and an exponent, t, of 1.73. In Si 3 N 4 /MWCNTs composites, electrical conductivity values about 10 S•m -1 for nanotube contents between 1 and 5 wt% have been reported [START_REF] Tatami | Electrically conductive CNTdispersed silicon nitride ceramics[END_REF][START_REF] Cs | Development of CNT/Si 3 N 4 composites with improved mechanical and electrical properties[END_REF], but values for p c have not been given yet.

In the present work, we describe the electrical properties of Si 3 N 4 with various contents of MWCNTs in dc and ac conditions for a wide temperature range, highlighting the effect of nanotube content and their connection within the network. The local electrical response of the composites is also analyzed and correlated with their nanostructure by means of conductive scanning force microscopy (C-SFM).

Experimental procedure

Si 3 N 4 materials containing MWCNTs, in concentrations ranging from 0.9 to 8.6 vol%, were prepared as described elsewhere [START_REF] Osendi | Dense and Homogenous Silicon Nitride Composites Containing Carbon Nanotubes[END_REF]. In short, CVD synthesized MWCNTs of 30 nm diameter and 1-5 µm length, according to the supplier (Nanolab Inc., USA), were thoroughly mixed with Si 3 N 4 (SN-E10, UBE Industries, Japan) plus liquid forming sintering additives, 2 wt% of Al 2 O 3 (SM8, Baikowski Chimie, France) and 5 wt% of Y 2 O 3 (Grade C, H.C. Starck GmbH & Co., Germany). Compositions were spark plasma sintered (Dr. Sinter, SPS-510CE, Japan) at 1585 ºC for 5 min in vacuum (6 Pa), applying a pressure of 50 MPa. Specimens were discs of 20 mm diameter and about 3 mm thick. A blank sample without nanotubes was equally processed for comparison. All specimens had densities above 99 % of the theoretical, good nanotube dispersion and no evidence of degradation, as previously reported [START_REF] Osendi | Dense and Homogenous Silicon Nitride Composites Containing Carbon Nanotubes[END_REF].

Microstructures on both polished and plasma etched (CF 4 /5 vol% O 2 at 100 W for 40 s) and fracture specimens containing MWCNTs were observed using a field emission scanning electron microscope (FESEM, Hitachi S-4700, Japan). Besides, samples were also prepared for observation in the transmission electron microscope (TEM, 200kV JEOL JEM 2000 FX) following usual cutting, dimpling and ion thinning procedures.

Micro-Raman spectra of the original MWCNTs and the composites were taken using the 540 nm laser wavelength excitation (In Via, Renishaw equipment, UK).

For the macroscopic electrical characterization, either two or four probe method was used depending on the electrical conductivity level initially measured. Resistivity under dc conditions was calculated from the current density versus electric field curves measured in a Keithley Sourcemeter 2410 model. For the ac characterization an Agilent 4294A Precision Impedance analyzer was employed scanning the frequency range 10-10 7 Hz. Measurements at different temperatures, from 273 to 573 K in 50 K steps and at 2 K•min -1 heating rate, were carried out in a home-made furnace that can be fitted to each probe configuration.

Scanning force microscopy (SFM) measurements were performed under low humidity conditions (2% RH, obtained in a N 2 atmosphere) using a commercial head and software from Nanotec. [START_REF] Horcas | A software for scanning probe microscopy and a tool for nanotechnology[END_REF]. Commercial silicon Cr/Pt coated probes and boron-doped diamond probes with force constants k = 0.2 and 3.0 N•m -1 were used for both morphological and conductive SFM measurements. To check tip-sample conditions, the adhesion force was systematically determined from force versus distance curves prior to and after each conductivity experiment.

The conducting tip was placed in direct contact with the sample, under controlled load, i.e. by using a normal force feedback, and the current was measured between tip and sample. The sample was always grounded and the voltage was applied to the tip (V tip ). Direct electric contact to ground was established through a metallic clamp attached at the sample border. We note that the ability to control the applied load permits avoiding any undesirable tip-induced effect, separating and controlling the mechanical response of the system under study. The conducting response of the sample surface was obtained by following different strategies [START_REF] Munuera | Scanning force microscopy three-dimensional modes applied to conductivity measurements through linear-chain organic SAMs[END_REF]: i) simultaneously acquiring topographic images z (x,y) and current maps I (x,y) over a given region at a given voltage, and ii) acquiring I-V characteristics curves at selected (x,y) locations on the surface. I-V curves were performed at least on 5 different regions for each sample and, within each region; 30-40 I-V curves were recorded.

Results and discussion

According to previous results of present authors [START_REF] Osendi | Dense and Homogenous Silicon Nitride Composites Containing Carbon Nanotubes[END_REF], the processing route used for these Si 3 N 4 /MWCNTs composites produces reasonably good dispersion of the MWCNTs within the Si 3 N 4 matrix, as FESEM observations of fractured surfaces revealed [START_REF] Osendi | Dense and Homogenous Silicon Nitride Composites Containing Carbon Nanotubes[END_REF]. Besides, the spark plasma sintering conditions allow high relative densities, 99.0%, while avoid nanotubes degradation, as micro-Raman spectroscopy data confirmed by showing similar intensity ratios for the characteristics bands of MWCNTs, i.e. D/G and G'/G [START_REF] Osendi | Dense and Homogenous Silicon Nitride Composites Containing Carbon Nanotubes[END_REF], in both the composites and original nanotubes.

TEM observations indicate that MWCNTs are bent and twisted at Si 3 N 4 grain boundaries (Fig. 1). Then, we can expect that the bulk electrical transport behaviour of the Si 3 N 4 /MWCNTs composites could be controlled by the nanotubes network but not by the nonconductive Si 3 N 4 matrix.

Figure 2 shows the dc electrical conductivity as a function of the MWCNTs content.

The value in the plot for the blank Si 3 N 4 corresponds to the detection limit of our experimental set-up, 10 -13 S•m -1 ; therefore, the real conductivity must be even lower. For the 0.9 vol% MWCNTs specimen, a sharp conductivity increase of more than ten orders of magnitude (4x10 -2 S•m -1 ) is registered as compared to the reference Si 3 N 4 specimen, which infers a p c < 0.9 vol%. For the samples with MWCNTs contents of 1.8 and 5.3 vol% , the conductivity steadily increases reaching values of 0.3 and 14 S•m -1 , respectively, and levels off at a content of 8.6 vol% to a saturation value of 17 S•m -1 . Similar rises in electrical conductivity have been reported but for higher MWCNTs contents, 1.8 wt% [4] and 3.0 wt% [START_REF] Cs | Development of CNT/Si 3 N 4 composites with improved mechanical and electrical properties[END_REF] (about 3 and 5 vol%, respectively), for two different Si 3 N 4 /MWCNTs composites sintered by hot pressing at higher temperatures than in the present work. The scattering in the reported percolation threshold and in electrical conductivity data [START_REF] Tatami | Electrically conductive CNTdispersed silicon nitride ceramics[END_REF][START_REF] Zhan | Electrical properties of nanoceramics reinforced with ropes of single-walled carbon nanotubes[END_REF][START_REF] Cs | Development of CNT/Si 3 N 4 composites with improved mechanical and electrical properties[END_REF][START_REF] Shi | Effect of multiwall carbon nanotubes on electrical and dielectric properties of yttria-stabilized zirconia ceramic[END_REF][START_REF] Poorteman | Colloidal processing, hot pressing and characterisation of electroconductive MWCNT-alumina composites with compositions near the percolation threshold[END_REF][START_REF] Ahmad | Electrical conductivity and dielectric properties of multiwalled carbon nanotube and alumina composites[END_REF][START_REF] Rul | Percolation of single-walled carbon nanotubes in ceramic matrix nanocomposites[END_REF][START_REF] Flahaut | Carbon nanotubemetal-oxide nanocomposites: Microstructure, electrical conductivity and mechanical properties[END_REF] are commonly due to differences in CNTs dimensions and type (SWCNT or MWCNT), their dispersion degree in the composite and the sintering temperature, as exposure to high temperatures may damage the nanotubes. For ceramic composites, percolation threshold values mostly lie between 0.64 vol% [START_REF] Poorteman | Colloidal processing, hot pressing and characterisation of electroconductive MWCNT-alumina composites with compositions near the percolation threshold[END_REF][START_REF] Rul | Percolation of single-walled carbon nanotubes in ceramic matrix nanocomposites[END_REF] and 0.80 vol% [START_REF] Ahmad | Electrical conductivity and dielectric properties of multiwalled carbon nanotube and alumina composites[END_REF].

As it is shown in Figure 2 for present composites, data at room temperature fit the scaling law using same parameters as Rul et al., [START_REF] Rul | Percolation of single-walled carbon nanotubes in ceramic matrix nanocomposites[END_REF] i.e., p c of 0.64 vol% and t of 1.73. The value of t lower than the characteristic of a tri-dimensional network (t = 2.0) was associated to a fluctuation induced tunnelling mechanism [START_REF] Ahmad | Electrical conductivity and dielectric properties of multiwalled carbon nanotube and alumina composites[END_REF][START_REF] Rul | Percolation of single-walled carbon nanotubes in ceramic matrix nanocomposites[END_REF][START_REF] Li | Dominant role of tunneling resistance in the electrical conductivity of carbon nanotube-based composites[END_REF], where thermally induced hopping transport between the CNTs separated by an insulating barrier dominates the conductivity behaviour. Monte Carlo simulations [START_REF] Li | Dominant role of tunneling resistance in the electrical conductivity of carbon nanotube-based composites[END_REF] points out that the tunnelling resistance caused by the insulating barrier between nanotubes or their aggregates plays a dominant role in the effective electrical conductivity of nanotube-based nanocomposites. According to those simulations, the tunnelling resistivity augments with increasing dielectric constant of the insulating matrix film between nanotubes, but this does not seem determinant because of the small differences in dielectric constants of these types of ceramics [START_REF] Nalwa | Handbook of low and high dielectric constant materials and their applications[END_REF], explaining the similar percolation thresholds observed for different nanotube-ceramic composites [START_REF] Poorteman | Colloidal processing, hot pressing and characterisation of electroconductive MWCNT-alumina composites with compositions near the percolation threshold[END_REF][START_REF] Ahmad | Electrical conductivity and dielectric properties of multiwalled carbon nanotube and alumina composites[END_REF][START_REF] Rul | Percolation of single-walled carbon nanotubes in ceramic matrix nanocomposites[END_REF].

The temperature dependence of the dc current-voltage curves of the composites with the lowest and highest MWCNTs content (Fig. 3), i.e., specimens with 0.9 and 8.6 vol% of MWCNTs, reveals notable differences in their electrical responses. For the 0.9 vol% composite (Fig. 3a), the I-V curves are non-linear in the whole range of measured temperatures. It should also be noticed that the electrical conductivity increases with temperature, which is consistent with classical negative temperature coefficient (NTC) type semiconductor behaviour [START_REF] Pike | Semiconductor grain-boundary admittance-theory[END_REF][START_REF] Blatter | Carrier transport through grain-boundaries in semiconductors[END_REF]. The origin of the non-ohmic behaviour is the existence of potential barriers that control the charge transport, which, in CNT containing composites, may form at nanotube-nanotube contacts as well as in regions within nanotubes showing high concentration of atomic defects. The composites microstructure can be described by the brick layer model [START_REF] Donald | Impedance Spectroscopy Theory, Experiment and Applications[END_REF], where the conducting MWCNTs network is found at grain boundaries (Fig. 1) surrounding the much less conductive Si 3 N 4 matrix. This agrees with the scenario also assumed for Al 2 O 3 /MWCNTs composites [START_REF] Ahmad | Dramatic effect of multiwalled carbon nanotubes on the electrical properties of alumina based ceramic nanocomposites[END_REF], described as an insulating matrix with large interconnected conducting regions at the grain boundaries in which potential barriers are built-up. On the other hand, the composite with 8.6 vol% MWCNTs (Fig. 3b) exhibits a linear current-voltage response at temperatures up to 423 K changing to a non-linear dependence at higher temperatures. Note also that up to 373 K, a small positive temperature coefficient (PTC) behaviour is measured, in accordance with the relevant metallic-like conduction mechanism. Furthermore, a slight increase in resistivity is detected at 523 K. These results lead us to infer that the electrical transport in these materials is a balance between semiconductor and metallic-like mechanisms, whose weights depend on both the temperature and the MWCNTs content. It should be noticed that both kinds of charge transport mechanisms have been reported for individual carbon nanotubes [START_REF] Ebbesen | Electrical conductivity of individual carbon nanotubes[END_REF] and buckypapers [START_REF] Kulesza | High-temperature electrical transport properties of buckypapers composed of doped single-walled carbon nanotubes[END_REF].

The change of the electrical response observed for the highest MWCNTs content points out to the possible origin of the potential barriers. Increasing the amount of MWCNTs will decrease the intertube spacing, improving the connectivity of the conducting nanotube network, reducing the thickness of the insulating film and favouring the tunnelling penetration of electrons. Consequently, at low MWCNTs content the contact resistance associated to the insulating film plays a dominant role in the effective electrical conductivity, whereas intrinsic defect clusters within the nanotubes becomes competitive in composites with high MWCNTs contents. Then, the changes observed in the temperature dependence of the electrical conductivity, are determined by the intrinsic electronic properties of MCWNTs.

For gaining some insight of the different conduction mechanisms in these composites, ac impedance spectroscopy measurements versus temperature were done in the 0.9 and 8.6 vol% MWCNTs composites (Fig. 4). According to the proposed brick layer model, where the conductive phase is the grain boundary one, the impedance semicircles bring information only of the MWCNT grain boundary phase. The insulating character of the Si 3 N 4 matrix is corroborated by measurements in the Si 3 N 4 blank specimen, as no impedance signal could be detected up to 1173 K (detection limit of the system is 10 7 ). As seen in Fig. 4a, for the specimen with the lowest nanotube content only one R-C circuit is needed to fit the arc, the dielectric relaxation time being in the order of 10 -6 s and slightly decreasing with temperature.

Conversely, the ac electrical response up to 423 K of the specimen with 8.6 vol% of MWCNTs (Fig. 4b) can be fitted by two series of R-C parallel circuits, whereas above that temperature a single R-C parallel circuit describes the response of this material. The relaxation times for the two R-C parallel circuits are 1.3 x 10 -6 s and 1.4 x 10 -7 s. The latter decreases with temperature down to 9.8 x 10 -8 s and above 423 K the arc vanishes. Zhao et al. [START_REF] Zhao | Frequency-dependent electrical transport in carbon nanotubes[END_REF] reported relaxation times of 8 x 10 -8 s for nanotube ropes of SWCNTs and MWCNTs. Therefore, the smallest relaxation time (~10 -7 s) might be related to the formation of a nanotube bundle network. At 523 K a significant increase in resistivity is measured.

To get experimental evidences of the charge transport mechanisms in Si 3 N 4 /MWCNTs composites, the local electrical response in relation to the microstructure of the composites was investigated by C-SFM. Low magnification simultaneously acquired topographic images and current maps (at 1 V) are shown in Figure 5 for samples with 0.9 vol% (Figs. 5a ,b), just above the percolation limit, and 5.3 vol% of MWCNTs (Figs. 5 c,d), that's it, in the plateau conductivity regime.

We first note that the surface morphology is extremely flat, as corresponds to a thoroughly polished specimen, except at the grain boundaries emerging at the surface, which are viewed as topographic holes. The flat regions do not show any detectable current, as expected from the insulating character of the Si 3 N 4 grains. However, for both specimens, electrical conduction is clearly observed at the hole grain boundary regions, exactly where the nanotubes are located as observed in the TEM micrograph (Fig. 1). In addition, note that, some nanotubes outcrop at the exposed grain boundaries (Fig. 6). The outcropping is better detected on the fracture surfaces (Fig. 6b), as the plasma etching process, necessary for revealing the Si 3 N 4 microstructure at the FESEM, damages the MWCNTs (Fig. 6a).

Provided that the electrical contact (tip counter electrode) is established by direct contact of a metallic clamp at the sample border, the fact of measuring current with the C-SFM indicates that the electrical current flows between the tip-sample contact and the grounded clamp all along the nanotubes. As a result, the MWCNTs must be interconnected forming a network that extends at least over several millimetres (tip to clamp distance). Such a percolated network of nanotubes is confirmed then as the main path for electrical conduction. Further valuable information can be still obtained from the C-SFM data. On one hand, as it can be seen comparing the current maps of the composites with low and high MWCNTs contents (Fig. 5), lower current is detected for the 0.9 vol% case. Moreover, at this low magnification, the conduction detected at the outcropping nanotubes looks like a spikelike signal. The reason has been revealed by magnified images at specific boundaries (Fig. 7).

Interestingly, the current is detected at the perimeter of small (~50-100 nm) circular contours (Fig. 7a) that are compatible with the radial dimensions of the MWCNTs. Though influenced by the unknown tip-nanotube contact resistance, I-V curves at these locations have a linear behaviour typical for a metallic-like conduction (Fig. 7b). To make clearer the correspondence between the nano-structured relief and the local electric conduction, the merged topography plus C-SFM image is presented in Fig. 7c. The result is plain; noticeably, the current is measured only at the external multi-wall shell of the nanotubes. As far as we know, this is the first experimental evidence that current is confined to the outermost layers as models predict [START_REF] Frank | Carbon nanotube quantum resistors[END_REF][START_REF] Tsukagoshi | Multiple-layer conduction and scattering property in multi-walled carbon nanotubes[END_REF]. In fact, a bias dependence of the conductance in MWCNTs has been ascertained by Tsukagoshi et al. [START_REF] Tsukagoshi | Multiple-layer conduction and scattering property in multi-walled carbon nanotubes[END_REF], being the main conduction layer at low-bias voltage (0-0.1 V) the outermost layer; whereas a few layers would contribute to the current conduction for higher-bias voltage (about 2 V).

At the holes, bent nanotubes bundles also emerge to the surface as small sheets of the MWCNTs network. They are observed lying on the surface nearby the grain boundaries as patches with two different electrical conductivities and I-V responses. This is illustrated in the current map of Figure 8 taken at V tip = 1 V. As can be seen in the current profile I (x) (Fig. 8b), zones with different brightness present differences in current of more than one order of magnitude, pointing out different conduction mechanisms, which are corroborated by the I-V characteristic curves. The regions showing higher current levels (Fig. 8d) present I-V ohmic linear behaviour whereas the regions of low current values present non-ohmic I-V responses (Fig. 8c). The conducting circular contours described above (Fig. 7) were seen in the high conductivity ohmic regions placed at the holes, but never detected in the low conductivity regions (Fig. 8a). At these less conducting locations, instead of contacting the nanotubes ends, the tip of our C-SFM does make contact perpendicularly to the lying MWCNTs sheets forming part of the percolated network. Different I-V responses in C-SFM have been reported for individual graphene sheets as well [START_REF] Gomez-Navarro | Electronic transport properties of individual chemically reduced graphene oxide sheets[END_REF], where non-ohmic behaviour was attributed to the presence of graphitic domains separated by defect clusters, giving raise to hopping conduction as the dominant charge transport mechanism.

Therefore, C-SFM results strongly upholds the hypothesis previously presented that electrical current is supported by MWCNTs bundles located at grain boundaries, which present two types of conduction mechanisms: metallic-like and semiconductor. The metallictype one is associated to charge transport along the nanotube shells, whereas the semiconductor behaviour could be linked to hopping conduction across different potential barriers associated to the nanotube-nanotube contacts [START_REF] Li | Dominant role of tunneling resistance in the electrical conductivity of carbon nanotube-based composites[END_REF] and to intrinsic defect clusters within the nanotubes separating metallic domains. These barriers give two different relaxation phenomena in ac measurements. Intrinsic defects within nanotubes are formed in areas where twisting and bending occur, inducing strong modifications of their electronic structure and provoking metallic to semiconducting transitions, as shown recently by Giusca et al. [START_REF] Giusca | Evidence for metal-semiconductor transitions in twisted and collapsed double-walled carbon nanotubes by scanning tunneling microscopy[END_REF]. The semiconducting mechanisms might become dominant if the conducting regions are not large enough. For the specimen with 0.9 vol% of MWCNTs, hopping conduction associated to contacts between nanotubes would be the dominant charge transport mechanism at all temperatures. As nanotube content increases, defect clusters separating metallic domains along them, with smaller relaxation distances [START_REF] Giusca | Evidence for metal-semiconductor transitions in twisted and collapsed double-walled carbon nanotubes by scanning tunneling microscopy[END_REF], would become significant. By increasing the temperature, the release of the nanotubes twist and collapse deformations would explain the transition from two to one R-C circuits observed at 473 K. Startlingly, this temperature coincides with the annealing temperature for releasing of the Wigner energy (energy stored as crystalline dislocations induced by neutron irradiation in graphite). Moreover, thermal relaxation of intersticial-vacancy pair defects has also been observed at that temperature in double-walled carbon nanotubes [START_REF] Urita | In situ observation of thermal relaxation of interstitial-vacancy pair defects in a graphite gap[END_REF].

Even though the study of mechanical properties of Si 3 N 4 /MWCNTs is not the scope of this paper, it can be said that for composites containing 1.8 vol% of MWCNTs, the fracture toughness (K IC ) is slightly higher than for blank Si 3 N 4 specimen of similar / phase ratio, 4.8 vs. 4.5 MPa•m 1/2 [START_REF] Osendi | Dense and Homogenous Silicon Nitride Composites Containing Carbon Nanotubes[END_REF]. Furthermore, ongoing studies show that these composite materials have a substantial improved response in terms of wear resistance, which is of great interest considering the real applications of Si 3 N 4 materials.

Conclusions

The multi-scale electrical behaviour response assembled by combining macroscopic electrical characterization and nanoscale C-SFM measurements describes a coherent picture of the dominant charge transport mechanisms in Si 3 N 4 /MWCNTs composites. A balance between semiconductor and metallic-like mechanisms, whose weights depend on both the temperature and the MWCNTs content, is observed. The metallic-type conduction is associated to charge transport along the nanotube shells whereas the semiconductor behaviour is linked to hopping conduction across nanotube-nanotube contacts and across intrinsic defect clusters within the nanotubes separating metallic domains. When the MWCNTs network is close to the percolation (0.9 vol% composite), inter-nanotubes hopping or tunnelling mechanisms become the rate limiting steps for charge transport. Once this concentration is surpassed the contribution of defect clusters within nanotubes in the highly conductive paths competes with the above mechanism. C-SFM measurements highlight that the high current and linear behaviour measured in outcropping nanotube clusters is due to conduction in the outermost shells, while lower current and non-linear behaviour is assessed when the contact is perpendicular to lying MWCNTs sheets of the percolated network.

Captions

Figure 1 .

 1 Figure 1. TEM micrograph for the composite containing 5.3 vol% of MWCNTs. The

Figure 2 .

 2 Figure 2. Measured dc conductivity as a function of MWCNTs content. Electrical

Figure 3 .

 3 Figure 3. Measured dc current-voltage data as a function of the temperature for the

Figure 4 .

 4 Figure 4. Complex impedance measured at different temperatures for the specimens with (a)

Figure 5 .

 5 Figure 5. Topographic (a, c) and C-SFM images at V tip = 1 V (b, d) for the composites

Figure 6 .

 6 Figure 6. FESEM micrographs showing: a) polished and plasma etched surface corresponding

Figure 7 .

 7 Figure 7. C-SFM image at V tip = 1 V (a), I-V curve at the outcropping MWCNTs (b) and

Figure 8 .

 8 Figure 8. (a) Merged topography plus C-SFM image for the composite with 5.3 vol% of

Acknowledgements

The financial support of Spanish Ministry of Science and Innovation through projects MAT2006-7118, MAT2007-62732 and NANOSELECT CSD2007-00041 is recognized. J. Gonzalez-Julian acknowledges the financial support of the JAE (CSIC) fellowship Program and Y. Iglesias acknowledges the financial support of the FPI Program.

 [START_REF] Ahmad | Dramatic effect of multiwalled carbon nanotubes on the electrical properties of alumina based ceramic nanocomposites[END_REF]


A Borrell
email: a.borrell@cinn.es

V G Rocha

R Torrecillas

A Fernández

Surface coating on carbon nanofibers with alumina precursor by different synthesis routes

Keywords: E. Spark plasma sintering, A. Carbon nanofibers, B. Mechanical properties, A. Coating, A. Nanoalumina

Alumina-reinforced carbon nanofiber nanocomposites were prepared using different routes; powders mixture, colloidal route and sol-gel process followed by spark plasma sintering (SPS). CNFs/xAl 2 O 3 (x=10-50 vol.%) were prepared through nanopowders mixing in a high-energy attrition milling. The main limitations in the preparation of this kind of nanocomposites are related to the difficulty in obtaining materials with a homogeneous distribution of both phases and the different chemical nature of CNFs and Al 2 O 3, which causes poor interaction between them. A surface coating of CNFs by wet chemical routes with an alumina precursor is proposed as a very effective way to improve the interaction between CNFs and Al 2 O 3 . An improvement of 50% in fracture strength was found for similar nanocomposite compositions when the surface coating was used. The improved mechanical properties of these nanocomposites are caused by stronger interaction between the CNFs and Al 2 O 3 .

INTRODUCTION

The development of new alumina-reinforced carbon nanofiber composite materials with excellent mechanical and electrical properties is useful for a wide range of industrial applications. Thanks to their exceptional mechanical, electrical and thermal properties, carbon nanofibers (CNFs) encourage the development of advanced materials through their incorporation in matrices such as polymers [START_REF] Ma | Structure, properties, and processing of polyester/nano carbon composite fibers[END_REF][START_REF] Higgins | Polycarbonate carbon nanofiber composites[END_REF], metals [START_REF] Kim | Fabrication of carbon nanofiber/Cu composite powder by electroless plating and microstructural evolution during thermal exposure[END_REF][START_REF] Hammel | Carbon nanofibers for composite applications[END_REF] and ceramics [START_REF] Matsui | Stress graphitization of C/C composite reinforced by carbon nanofibers[END_REF][START_REF] Seghi | High density carbon fiber/boron nitride matrix composites: Fabrication of composites with exceptional wear resistance[END_REF].

Recently, there have been major developments in polymer matrix composites reinforced by CNFs, where properties achieved have been clearly higher than those corresponding to pure polymeric material. Nevertheless, there have been few advances for CNFs composites in ceramic and metal matrix, because getting good dispersion of CNFs in these systems is technically challenging [START_REF] Hirota | Fabrication of carbon nanofibers (CNF)-dispersed Al 2 O 3 composites by pulsed electric-current pressure sintering and their mechanical and electrical properties[END_REF][START_REF] Maensiri | Carbon nanofiberreinforced alumina nanocomposites: Fabrication and mechanical properties[END_REF][START_REF] Duszová | Zirconia/carbon nanofiber composite[END_REF][START_REF] Borrell | High density carbon materials obtained at relatively low temperature by spark plasma sintering of carbon nanofibers[END_REF]. Alumina is a ceramic commonly used in the development of advanced ceramics [START_REF] Cranmer | Overview of technical, engineering and advanced ceramics. Engineered Materials Handbook, Ceramic and Glasses[END_REF], especially in functional applications because of its relatively high hardness, good oxidation resistance and chemical inertness. There are studies [START_REF] An | Tribological properties of hot-pressed alumina-CNT composites[END_REF][START_REF] Siegel | Mechanical behavior of polymer and ceramic matrix nanocomposites[END_REF][START_REF] Ch | Carbon nanotubes-Fe-Alumina nanocomposites. Part II: microstructure and mechanical properties of the hot-Pressed composites[END_REF] showing that the addition of CNFs to alumina matrix does not produce a reinforcement in the materials obtained. This has been attributed to an inhomogeneous dispersion of CNFs in the composite materials due to the formation of CNFs aggregates as a consequence of strong Van der Waals interactions [START_REF] Jiang | Production of aqueous colloidal dispersions of carbon nanotubes[END_REF][START_REF] Sun | Development of a dispersion process for carbon nanotubes in ceramic matrix by heterocoagulation[END_REF], which sometimes cause anisotropy in the final properties of the materials. In the majority of the previous studies, alumina powders are mechanically mixed with CNFs in order to prepare CNFs/alumina composites that have, in some cases, properties lower than expected. Generally, the addition of carbon nanofibers as reinforcement leads to slight improvements in the fracture strength although the hardness and/or toughness are degraded [START_REF] Hirota | Fabrication of carbon nanofibers (CNF)-dispersed Al 2 O 3 composites by pulsed electric-current pressure sintering and their mechanical and electrical properties[END_REF]. These results are related to the formation of CNFs agglomerates and the weak CNFs/Al 2 O 3 interface. A promising solution for avoiding these problems could be the direct synthesis of alumina nanoparticles on the CNFs surface. There are many processes for synthesizing ceramic nanopowders such as mechanical synthesis, vapor phase reaction, precipitation and combustion or sol-gel methods [START_REF] Hida | Pulse electric current sintering of transition alumina/zirconia composite powders prepared by a novel sol-gel route[END_REF]. Although there are different limitations in all the aforementioned methods [START_REF] Shojaie-Bahaabad | Economical synthesis of nano alumina powder using an aqueous sol-gel method[END_REF][START_REF] Panchula | Mechanical synthesis of nanocrystalline -Al 2 O 3 seeds for enhanced transformation kinetics[END_REF], wet chemical methods are especially useful for this purpose. In many cases conventional processing methods for traditional powder mixture are relieved by chemical synthesis methods that allow the reactions, nucleation and growth of nanocrystals at molecular level to be controlled.

In order to obtain CNFs/alumina dense materials, spark plasma sintering (SPS) was used as the sintering technique. This is also known as field assisted sintering technique (FAST) and it is a new sintering technique that can consolidate powder compacts by applying an on-off dc electric pulse under uniaxial pressure [START_REF] Joanna | Sintering activation by external electrical field[END_REF]. This technique can work at heating rates of hundreds degrees per minute, reaching high temperatures in a short time, and producing dense materials [START_REF] Wang | Study of temperature field in spark plasma sintering[END_REF]. These features allow the achievement of microstructures unattainable by other sintering techniques. Furthermore, spark plasma sintering allows densification of materials that are difficult to sinter. In a previous work it was shown how it is possible to obtain high density carbon materials by SPS sintering of carbon nanofibers [START_REF] Borrell | High density carbon materials obtained at relatively low temperature by spark plasma sintering of carbon nanofibers[END_REF]. Then, this technique allows obtaining CNTs and CNFs monoliths without binder [START_REF] Uo | Multiwalled carbon nanotube monoliths prepared by spark plasma sintering (SPS) and their mechanical properties[END_REF][START_REF] Li | Microstructure and mechanical properties of hot-pressed carbon nanotubes compacted by spark plasma sintering[END_REF]. It means that high density carbon materials could be obtained in very short processing cycles and their mechanical properties could be improved by adding ceramic nanoparticles as second phase. The improved graphitic materials would be very interesting for applications that require a combination of thermal and electrical conductivities with good mechanical strength.

Therefore, the aim of this work was to develop alumina-reinforced carbon nanofiber nanocomposites. In order to improve the affinity between the components, the surface coating of CNFs with an alumina nanoparticles precursor was studied by the sol-gel and colloidal processes. The "surface coating" composites have been evaluated in comparison with materials prepared by conventional powders mixing. The materials obtained were compared in terms of their microstructural evolution, mechanical and electrical properties.

EXPERIMENTAL PROCEDURE

Preparation of composite powders

Mixed powders

The materials used in this study were commercial carbon nanofibers (CNFs) having an average outer diameter of 20-80 nm and lengths >30 m, supplied by Group Antolin Engineering, (Spain). These CNFs were generated via vapor phase growth (VGCNFs) [START_REF] Ci | Carbon nanofibers and single-walled carbon nanotubes prepared by the floating catalyst method[END_REF] using a floating catalyst of nickel in solution (6-8%). α-Al 2 O 3 nanopowder (Taimei TM-DAR Chemicals Co. Ltd, Japan) with an average particle size of 153 nm and a purity of 99.99% was used as a second phase. The powder mixtures, CNFs/Al 2 O 3 = 90/10 -80/20 -70/30 -60/40 -50/50 vol.%, were dispersed in ethanol (Panreac Quimica) with a high-energy attrition milling (Union Process, EE.UU) using alumina media of 2 mm diameter at 400 rpm and milling times of 1 hour. After milling, the resulting slurry was dried at 60 ºC and the dried powder was sieved under 60 m.

Surface coating of CNFs

The surface coating of CNFs was carried out employing two different processing routes, colloidal and sol-gel. Aluminum chloride hexahydrate, AlCl 3 •6H 2 O (>99% purity, Fluka), was used as the alumina precursor in both cases. The alumina precursor was added to a CNFs water dispersion eventually yielding composites with a 10 vol.% Al 2 O 3 . In the colloidal synthesis route or heterogeneous precipitation method [START_REF] Schehl | Alumina nanocomposites from powder alkoxide mixtures[END_REF],

CNFs were dispersed in distilled water using magnetic stirring. The appropriate amount of aluminum chloride, also dissolved in distilled water, was added dropwise to the water dispersion of CNFs. The slurry was first dried maintaining magnetic stirring at 60-70 ºC and subsequently at 120 ºC in order to eliminate wastewater. In the case of the sol-gel method, the aluminum chloride was dissolved in distilled water under stirring and, in order to control the chemical reactivity in the solution, the pH was adjusted with ammonium hydroxide solution (Fluka, 28% in water), keeping a constant pH value (~9) throughout the process in order to control the sol formation. Subsequently, CNFs were added to the aluminum chloride solution and mixed thoroughly by stirring. Finally, the solution was heated gently to form a gel. After aging for 24 h, a gelatinous product was obtained. The resultant amorphous gel was dried at 60-70 ºC and subsequently at 120 ºC for 24 h in order to eliminate wastewater. The CNFs/Al 2 O 3 powders were ground by mortar and pestle, and sieved below 60 µm to achieve a uniform distribution of particle sizes. Thermogravimetric analysis of the modified powders (TGA Star System, Mettler Toledo) were carried out up to 1000 ºC, using a 5 ºC min -1 heating rate under an argon atmosphere in order to check the temperature to remove solvent and chloride wastes.

SPS and characterization of sintered bodies

The powder samples were placed into a graphite die with an inner diameter of 20 mm and cold uniaxially pressed at 30 MPa. Then, they were introduced in a spark plasma sintering apparatus HP D 25/1 (FCT Systeme GmbH, Rauenstein, Germany) under low vacuum (10 -1 mbar) and sintered at 1500 ºC for 1 min under an applied pressure of 80 MPa and a heating rate of 100 ºC min -1 . Bulk density of the sintered bodies was measured by the geometric method from weight and geometric volume of the material.

Relative densities were calculated as the relation between geometrical and theoretical densities. The fracture strength was measured using biaxial testing employing the equations of Kirstein and Woolley [START_REF] Kirstein | Symmetrical bending of thin circular elastic plates on equally spaced point supports[END_REF], Vitman and Pukh [START_REF] Vitman | A method for determining the strength of sheet glass[END_REF], and the standard specification ASTM F394-78 [28]. All tests were performed at room temperature using the universal machine Instron (Model 8562) with a cross-head displacement speed of 0.002 mm s -1 . The samples for hardness analysis were previously polished (Struers, model RotoPol-31) with diamond to 1 m roughness. The hardness of the materials was determined using the indentation technique (Buehler, model Micromet 5103) with a conventional diamond pyramid indenter. The diagonals of each indentation were measured using an optical microscope. Measuring conditions for the Vickers hardness, H v, were an applying load of 2 N for 10 s and the standard specification ASTM E92-72.

The value of H v is the relationship between applied load P and the surface area of the diagonals of indentation [START_REF] Antis | A Critical Evaluation of Indentation Techniques for Measuring Fracture Toughness: I, Direct Crack Measurements[END_REF]. X-ray diffraction (XRD) analysis was performed on a Siemens X-ray diffractometer (Model D 5000), using "Bragg-Brentano" geometry with Cu-k α radiation (λ=0.15418 nm, 40 kV, 30 mA). The measurements were performed in the range of 15-80º with a step size and time of reading of 0.02º and 0.3 s, respectively.

Fracture surfaces of sintered samples were characterized by scanning electron microscopy (SEM, Zeiss DSM 950). The electrical resistivity of monolithic and composite materials was determined according to ASTM C611. The specimens were placed between two sheets of copper connected to a power supply, which allowed working at different current intensities. The measures were carried out by fixing the intensity of the current at 0.5 A using a multimeter of fixed pegs (9.55 mm separation), determining the voltage drop.

RESULTS AND DISCUSSION

Composites obtained from mixing powders

Table 1 summarizes the main properties such as density, hardness and fracture strength of the different CNFs/Al 2 O 3 composites prepared by powder mixing and sintered by SPS at 1500 ºC. Relative densities slightly increase with alumina addition from 10-40 vol.% Al 2 O 3 being always around 90%. However, when 50 vol.% Al 2 O 3 is added the relative density reached ~96%. This indicates that alumina powder addition helps the sinterability of CNFs and this effect is not proportional to ceramic phase content. The density improvement of the nanocomposites can only be observed for alumina contents of 40-50 vol.%. This fact may be due to the activation of sintering mechanisms of the ceramic phase, a high alumina content being necessary.

In good correlation with the densities values it can be seen that the fracture strength only increases when alumina content in the composite is 50 vol.%. Fracture strength of CNFs+50 vol.% Al 2 O 3 composite is 134 MPa, and this value is twice as much as corresponding to 40 vol.% Al 2 O 3 composite. So, when alumina content is equal or greater than 50 vol.% Al 2 O 3 , it can be considered that the alumina behaves as the matrix. Therefore, for a relatively low volume fraction of alumina, the fracture strength of the nanocomposite is similar to that of a carbon material, approximately 60 MPa [START_REF] Borrell | High density carbon materials obtained at relatively low temperature by spark plasma sintering of carbon nanofibers[END_REF], and a similar behavior to ceramic material is observed only when the alumina content is high. However, the hardness of the composites increases gradually with the ceramic phase addition. The reason for the different trend in both mechanical properties with alumina content is related to the influence of second phase nanoparticles on matrix behavior. In the case of hardness, the presence of alumina nanoparticles, characterized by its high hardness, leads to local reinforcement and this effect depends on nanoparticles content. However, the material fracture strength is proportional to its fracture toughness and therefore, to its opposition to crack propagation. Only when the matrix behavior changes as in the case of high alumina content, is a major improvement in material fracture strength observed.

In order to study this effect, Figure 1 shows SEM fracture surfaces micrographs of CNFs/Al 2 O 3 composites sintered by SPS at 1500 ºC with 10 (a), 40 (b) and 50 (c) vol.% Al 2 O 3 . Due to the nature, content, size and shape of the CNFs present in these composites, the components identification by SEM is very difficult, and the ceramic phase is not well identified until the alumina content in the composite is around 50 vol.% Al 2 O 3 (Fig. 1c). However, in micrographs corresponding to composites with lower ceramic content, alumina nanoparticles were identified by a circle (Fig 1a). It is important to note the small size of those nanoparticles. Taking into account the starting grain size of the alumina particles (~150 nm) it can be seen that after sintering at 1500 ºC the final size is very similar. This shows that the grain growth of the ceramic component in the nanocomposites is suppressed by CNFs presence. As was previously explained, the fracture mode of the material with a 50 vol.% of Al 2 O 3 is predominantly intergranular, and is reminiscent of the behavior of a ceramic material (Fig. 1c).

Therefore, this micrograph suggests the matrix change, when the alumina volume content achieves a 50 vol.%.

The electrical resistivity of all CNFs/Al 2 O 3 composite materials, measured at room temperature was ~10 -2 cm. This value corresponded to the carbon nanofibers raw material. This result shows that whereas the conductive component content, in this case the carbon nanofibers, is over the critical percolation volume fraction [START_REF] Shackelford | Material science and engineering handbook[END_REF], the electrical conductivity of the nanocomposites remains constant. So, very useful properties have arisen in the CNFs+50vol.% Al 2 O 3 composite. It has the same electrical conductivity but with a considerably higher fracture strength (134 vs. 60 MPa) than those obtained in 100% CNFs material [START_REF] Borrell | High density carbon materials obtained at relatively low temperature by spark plasma sintering of carbon nanofibers[END_REF]. This new composite will be especially useful because it could replace graphite materials [START_REF] Hoshii | Preparation of graphitic materials by spark plasma sintering method[END_REF] in applications in which their uses are now forbidden due to their low mechanical properties. The homogeneous distribution of ceramic and carbon components in the nanocomposite allows the obtaining of a material that combines good electrical and mechanical properties.

Composites obtained from colloidal and sol-gel methods

Another problem in obtaining ceramic-carbon composites is related to their different surface properties, reflected in a weak interface [START_REF] Chan | Fabrication of carbon nanotube reinforced alumina matrix nanocomposite by sol-gel process[END_REF]. So, in order to explore the possibility of improving this feature, the preparation of these nanocomposites by two wet chemical synthesis routes was studied.

Figure 2 shows the thermogravimetric analysis corresponding to the powders synthesized by colloidal and sol-gel methods: the curve of sol-gel powders shows two weight loss steps. In the first step (160-360 ºC) there is a weight loss of ~32 wt.% with T max at 285 ºC. This is due to the residual solvent adsorbed in the powders caused by a three-dimensional network formed in the sol-gel process that retains more water. In the second step (360-680 ºC) there is a weight loss of 8 wt.%. This loss tendency can be associated with the chloride elimination. Finally, at temperatures higher than 720 ºC the weight loss stabilizes.

Table 2 shows the two weight loss steps of the colloidal composite powder. T max takes place at 225 ºC and the first step weight loss ~12 wt.%, is considerably lower than the sol-gel process due to the easier elimination of water in this material during the drying steps. The second weight loss step takes place between 360-680 ºC and the weight loss is very similar to the sol-gel powder ~10 wt.% that could be expected considering that it is related to the chloride removal and the aluminum chloride amount used in both cases was the same. Taking into account these results, the powders were thermally treated at 700 ºC for 2 h in an argon atmosphere, in order to remove precursor wastes without affecting the carbon nanofibers.

Figure 3 shows the XRD-patterns of sintered bodies. Both composite powders led toalumina and graphite after sintering up to 1500 ºC by SPS. The lower relative intensity and wider peaks of alumina in the composite obtained from sol-gel powder, compared with material prepared by the colloidal route is indicative of a smaller crystal size, as they have the same content in ceramic phase.

Table 3 summarizes the properties such as the density, hardness and fracture strength of the different CNFs+10vol.% Al 2 O 3 composites sintered by SPS at 1500 ºC. The surface coating by colloidal or sol-gel method should improve the mechanical properties as it could be predicted from their relative densities. In both cases the relative density of composites (~93%) are slightly higher than that achieved in the material with similar composition obtained by powder mixing (~90%). The sol-gel process allows the formation of a three dimensional network which perfectly interconnects the particles in the whole volume. Consequently, carbon nanofibers which have been previously well dispersed in the solvent will be homogeneously coated by the gel layer. Heat treatment will lead to the formation of alumina nanoparticles that are very well dispersed in the composite. Although the colloidal route process seems to be very similar to the sol-gel process the differences are very significant. In the colloidal route, carbon nanofibers are preferential nucleation sites for the alumina precursor and this mechanism is known as heterogeneous precipitation. However, another process is in competition with heterogeneous precipitation, the nucleation and growth of aluminum hydroxide particles directly from the solution, this being homogeneous precipitation.

The alumina nanoparticles formed by heterogeneous precipitation would be similar to those obtained by the sol-gel method. However, particles formed by homogeneous precipitation are not linked to the carbon nanofibers surface and so i) they can grow, achieving greater size and leading to potential defects and ii) they are not contributing to ceramic-carbon interface reinforcement. In fact, the fracture strength of the sintered body of the sol-gel composite powder is almost 50% higher than that obtained by the mixed powders composite and considerably higher than colloidal processing route composite.

As could be expected, the electrical resistivity of both composites, obtained by colloidal and sol-gel methods is 10 -2 cm, this value being similar to the material obtained by mixing powders. Additionally, the electrical resistivity was measured in both directions, parallel and perpendicular to the direction of the applied pressure during sintering, and in both directions the same results were achieved. This is indicative of a homogenous dispersion of CNFs and alumina.

CONCLUSIONS

The addition of Al 2 O 3 powders as a second phase to the CNFs matrix improves the densification of materials based on CNFs. High alumina contents are necessary to improve mechanical properties of CNFs/alumina nanocomposites by powder mixture, while the surface coating of carbon nanofibers using a precursor of alumina leads to an improved composite even with only a 10 vol.% Al 2 O 3 . All composites show an electrical conductivity similar to the raw carbon nanofibers, therefore alumina could be used as reinforcement phase of carbon composites. The surface coating of carbon nanofibers using a precursor of alumina is a very useful method for obtaining CNFs/Al 2 O 3 composites, due to an improvement of interface between the CNFs and alumina. A major improvement in the mechanical properties of CNFs/Al 2 O 3 nanocomposites was achieved by sol-gel CNFs surface coating with an alumina precursor. This nanocomposite shows almost a 50% higher fracture strength compared to composites prepared by conventional mixing powders.

Figure 1 .

 1 Figure 1. SEM micrographs of fracture surfaces of the CNFs/Al 2 O 3 sintered materials

Figure 2 .

 2 Figure 2. TG analysis of the powders synthesized by colloidal and sol-gel methods.

Figure 3 .

 3 Figure 3. XRD patterns of the powders prepared by the colloidal and sol-gel methods

a

  Temperature of initial weight loss (ºC) b Temperature of final weight loss (ºC) c Difference between Tf and T 2 (ºC) d Difference between T 2 and Ti (ºC) e Temperature of maximum rate of weight loss (ºC) f Weight loss in T 1 g Weight loss in T 2

Table 1 .

 1 Characteristics of CNFs/Al 2 O 3 composites materials obtained from mixing powders and fabricated by SPS at 1500 ºC and 80 MPa for 1 min in a vacuum.

	Samples	Relative density (%)	Fracture strenght (MPa)	Vickers Hardness (Kg/mm 2 )
	CNFs+10vol.% Al 2 O 3	89.4	60.8 ± 2.3	62.5 ± 1.7
	CNFs+20vol.% Al 2 O 3	89.8	59.4 ± 1.9	63.9 ± 1.2
	CNFs+30vol.% Al 2 O 3	90.2	60.9 ± 2.2	95.2 ± 2.2
	CNFs+40vol.% Al 2 O 3	93.5	63.4 ± 2.6	134.1 ± 2.8
	CNFs+50vol.% Al 2 O 3	95.9	134.3 ± 3.2	211.7 ± 2.6

Table 2 .

 2 Thermogravimetric data of the coated CNFs/Al 2 O 3 powders.

	Samples	Relative density (%)	Fracture strength (MPa)	Vickers hardness (Kg/mm 2 )
	CNFs+10vol.% Al 2 O 3 (Mixing powders)	89.4	60.8 ± 2.3	62.5 ± 1.7
	CNFs+10vol.% Al 2 O 3 (Colloidal)	92.3	65.9 ± 2.8	57.5 ± 1.4
	CNFs+10vol.% Al 2 O 3 (Sol-Gel)	93.5	89.5 ± 3.2	60.9 ± 0.9

Table 3 .

 3 Properties of CNFs+10vol.% Al 2 O 3 composites obtained from mixing powders, colloidal and sol-gel methods and fabricated by SPS at 1500 ºC and 80 MPa for 1 min in vacuum.

Acknowledgments

This work has been carried out with financial support of National Plan Projects MAT2006-01783 and MAT2007-30989-E and the Regional Project FICYT PC07-021.

A. Borrell, acknowledges the Spanish Ministry of Science and Innovation for Ph.D. grant.