Camélia Ben Ramdane
email: benramdane.camelia@gmail.com

Aurélie Julian-Jankowiak

Michel Parlier

Roger Valle

Eric Martin

Pascal Diss

: Camélia

Ben Ramdane

Aurélie Jankowiak

Aurélie

Jankowiak

Pascal

Identification of the mechanical behaviour and damage mechanisms of a Nextel™610/alumina ceramic matrix composite subjected to tensile loading

Keywords: (4): A: Ceramic-matrix composites (CMCs), B: Fibre/matrix bond, B: Mechanical properties, D: Electron microscopy

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Introduction

Two solutions have risen to reduce aircraft greenhouse gases emissions: lightening the aircrafts and increasing the operating temperature of the gas turbine and, as a consequence, of the exhaust gases. Composite materials have been widely used during the last decades to reduce weight. Ceramic matrix composites (CMCs), used for thermostructural applications [START_REF] Levi | Processing and performance of an all-oxide ceramic composite[END_REF], fulfil both requirements: their density is lower than that of conventional nickel-based superalloys and they are capable of maintaining excellent strength and fracture toughness at high temperatures [START_REF] Jackson | Compressive creep behavior of an oxide-oxide ceramic composite with monazite fiber coating at elevated temperatures[END_REF][3][4] [START_REF] Ruggles-Wrenn | Effects of steam environment on hightemperature mechanical behavior of Nextel™720/alumina (N720/A) continuous fiber ceramic composite[END_REF]. Moreover, these materials, designed for applications in oxidizing environments (turbo-engines), have to be oxidation-resistant and thermodynamically stable [START_REF] Jackson | Compressive creep behavior of an oxide-oxide ceramic composite with monazite fiber coating at elevated temperatures[END_REF]. Non-oxide CMCs, generally SiC/SiC composites, show poor oxidation resistance at intermediate temperatures, around 700-800°C [START_REF] Ruggles-Wrenn | Effect of loading rate on the monotonic tensile behavior and tensile strength of an oxide-oxide ceramic composite at 1200°C[END_REF], due to the presence of carbon at the fibre/matrix interface.

Therefore, oxide fibre/oxide matrix CMCs appear to be good candidates for thermostructural applications in this temperature range [START_REF] Parlier | State of the art and perspectives for oxide/oxide composites[END_REF]. Using oxide fibre/non-oxide matrix or non-oxide fibre/oxide matrix composites does not improve the high temperature oxidation resistance [START_REF] Ruggles-Wrenn | Effect of loading rate on the monotonic tensile behavior and tensile strength of an oxide-oxide ceramic composite at 1200°C[END_REF], the mechanical properties of the non-oxide components being reduced by oxidation. Despite their better oxidation resistance compared to SiC/SiC composites, oxide/oxide composites exhibit lower mechanical properties and cannot be used at temperatures higher than 1000-1100°C. Beyond these temperatures, the degradation of the fibres [START_REF] Cantonwine | Strength of thermally exposed alumina fibers. Part II Bundle behavior[END_REF] induces a decrease in the mechanical properties [START_REF] Zok | Mechanical properties of porous-matrix ceramic composites[END_REF].

It is now commonly assumed that the mechanical performance of CMCs relies on the fibre/matrix interface. When matrix cracks initiate and propagate under loading, it seems necessary to isolate the fibres from the matrix so that they remain intact and ensure the mechanical strength of the composite [START_REF] Koch | Ceramic fiber composites: Experimental analysis and modeling of mechanical properties[END_REF]. Therefore, two approaches exist: weak matrix and weak interface composites [START_REF] Koch | Ceramic fiber composites: Experimental analysis and modeling of mechanical properties[END_REF][10] [START_REF] Mattoni | Effects of matrix porosity on the mechanical properties of a porous-matrix, all-oxide ceramic composite[END_REF]. The weak matrix approach is based on a microporous matrix. In these composites, cracks propagate between the micropores of the matrix, which enables energy dissipation, and, when matrix cracks reach the fibres, they are deflected at the fibre/matrix interface. This leads to fibre/matrix debonding and fibre pull-out, thereby providing high composite toughness [START_REF] Zok | Mechanical properties of porous-matrix ceramic composites[END_REF] [START_REF] Wilson | High performance oxide fibers for metal and ceramic composites[END_REF]. The weak matrix concept assumes a strong fibre/matrix bond, but the high porosity of the matrix ensures a weak fibre/matrix interface.

Weak matrix principle is illustrated in Fig. 1.

Regarding the second approach, it is based on a weak fibre/matrix interface promoting fibre/matrix debonding and sliding. This can be achieved by introducing a weak interphase between the fibres and the matrix, which is generally fully dense. Such an interphase consists of a fibre coating, which increases the process cost [START_REF] Wilson | High performance oxide fibers for metal and ceramic composites[END_REF]. Thus, the weak matrix approach appears to be the most attractive. According to previous studies, the matrix porosity must be in the 30-40 % range to obtain a weak matrix composite [START_REF] Levi | Processing and performance of an all-oxide ceramic composite[END_REF][8] [START_REF] Kerans | Structural ceramic composites[END_REF].

Regarding the fibre volume fraction, it is essential, for turbo-engine applications, to manufacture materials with a high fibre volume fraction in order to enhance the mechanical properties of the composite material, especially the Young's modulus. Hence, the aimed fibre volume fraction and total porosity were 50 % [START_REF] Parlier | Potential and perspectives for oxide/oxide composites[END_REF] and 25 %, respectively.

Most of the oxide/oxide composites described in previous studies were reinforced with commercial Nextel™610 or Nextel™720 fibres embedded in an alumina, an alumina-silica or an alumina-mullite matrix. Nextel™610 alumina fibres offer high strength [START_REF] Wilson | High performance oxide fibers for metal and ceramic composites[END_REF] at low and moderately high temperatures (up to approximately 1000°C) whereas Nextel™720 aluminasilica fibres have good creep resistance up to approximately 1200°C [START_REF] Zok | Mechanical properties of porous-matrix ceramic composites[END_REF]. As for the matrix, the main advantage of alumina compared to alumina-silica and alumina-mullite is that its sintering process can start at temperatures lower than 1000-1100°C. Beyond this temperature range, a loss of mechanical properties of Nextel™610 fibres occurs. Thus, combining Nextel™610 alumina fibres with an alumina matrix offers two advantages: (i) the matrix can be partially sintered without damaging the fibres and (ii) there is no reaction between fibres and matrix, hence avoiding the creation of a strong fibre/matrix interface that would prevent crack deflection at the fibre/matrix interface and therefore fibre debonding.

The objective of this study is to investigate the mechanical behaviour of a weak matrix alumina/alumina composite under tensile loading and to identify the damage mechanisms of this composite [START_REF] Ramdane | Etude et modélisation du comportement mécanique de CMC oxyde/oxyde[END_REF]. Previous studies on the mechanical behaviour of several oxide/oxide CMCs are available in the literature, but a precise description of the damage mechanisms of these composites seems to be absent.

Several techniques were used for the determination of damage mechanisms: optical microscopy and scanning electron microscopy examinations, in situ tensile testing in a scanning electron microscope (SEM), infrared thermography, X-ray tomography and ultrasonic scanning. None of the above damage monitoring techniques seems to be reported in the literature for an application to oxide/oxide CMCs.

Experimental procedure

Material

The material consisted of a micro-porous alumina matrix reinforced with woven Nextel™610 alumina fibres, without coating. The fibre fabric (8 harness satin weave) was first infiltrated with a water-based alumina matrix slurry, then dried and cut into twelve plies. These prepreg plies were then infiltrated with water, laid up in a mould, hot pressed at 100°C in a vacuum bag and finally heat treated at 1200°C (pressureless sintering) to obtain approximately 2.5 mm thick plates with the required microporosity and fibre volume fractions.

Mechanical testing and damage monitoring

In-plane mechanical properties of the Nextel™610/alumina composite were studied through uniaxial monotonic and cyclic loading in the fibre direction (weft direction) and in ±45° fibre orientation. All tests were conducted at room temperature in laboratory air environment, using a servocontrolled Z150 Zwick/Roell Materials testing machine. The tests were performed in stroke control with a constant displacement rate of 0.5 mm.min -1 .

Straight-sided specimens of 16 mm x 150 mm x composite thickness, with a gauge length of 70 mm were used [START_REF] Ramdane | Etude et modélisation du comportement mécanique de CMC oxyde/oxyde[END_REF].

Strain measurements were accomplished simultaneously by digital image correlation (DIC, VIC-3D™), a uniaxial extensometer of 25 mm gauge length (Schenck) and strain gauges (Vishay Precision Group, CEA-06-250UW-350). After test result analysis, it appeared that the uniaxial extensometer slid on the surface of the specimens. Besides, the strains measured by DIC and by strain gauges were very similar. Thus, only the strains measured by DIC were taken into account.

Young's moduli were determined from stress-strain curves and damage thresholds were defined as the stress and strain corresponding to the elastic limit, for each tested specimen.

In-plane shear moduli (GPa) were calculated through tensile tests performed on specimens with ±45° fibre orientation, using the following equation:

= 2 - × 10 (1)
where is the unidirectional force (N) applied in the loading direction, and are, respectively, the thickness (mm) and the width (mm) of the sample, and are, respectively, the longitudinal and transverse strains. Both longitudinal and transverse strains must remain within the elastic domain, according to the stress-longitudinal strain and stresstransverse strain curves.

For cyclic loading tests, the specimens were sequentially loaded at different stresses, with a return to zero load between two cycles. The sequential stresses were determined from monotonic tensile tests and correspond to different points on the stress-strain curves: a point slightly before the average damage threshold, a point just before the average failure and intermediate points. Interlaminar properties are critical for laminated composites with 2D woven reinforcements.

For
The interlaminar shear strength (ILSS) of the composite was measured by 3-point bending tests on short beams according to BS EN 658:5-2002 standard, in the fibre direction (0°).

The schematic diagram of the 3-point bending test geometry is given in Fig. 2. The dimensions of ILSS test specimens were 25 mm x 10 mm x composite thickness. The cylindrical rollers had a diameter of 4 mm and the outer support span was 15 mm. A displacement rate of 0.5 mm.min -1 was used for these tests. The ILSS (MPa) is calculated using the following equation:

= 3 4 ℎ (2)
where is the shear failure force (N), is the mean test specimen width (mm) and ℎ is the mean test specimen thickness (mm).

In situ tensile tests were carried out in a SEM (Zeiss DSM-960) in order to (i) corroborate the post mortem observations and (ii) ascertain the chronology of the damage mechanisms. In situ testing offers two advantages: real time damage monitoring and observation of the material under load, which prevents crack closure when returning to zero load. This phenomenon cannot be avoided in the case of post mortem observation. But in situ testing also have disadvantages. First of all, the technique requires the use of small straight sided specimens (3 mm x 40 mm x composite thickness) with a 3 mm x 10 mm x composite thickness gauge, which might not be representative of the macroscopic behaviour of the composite, the geometrical unit cell of the composite being 8 mm x 8 mm x composite thickness. Moreover, the observation of damage being performed on the specimen edge, the observed mechanisms are thus occurring in a plane stress state. The observed edge of the specimens was mirror polished prior to testing (without resin impregnation). The samples were observed in real-time during loading tests and images of the total gauge length were acquired after each 100 N loading step (approximately 13 MPa considering the section of the specimens), while the specimens were maintained under constant load.

Microstructural characterization

Porosity of the as-processed composite was measured by mercury intrusion porosimetry and by Archimedes' method, in water. The results obtained by these two techniques were similar.

Five samples from five different plates were analysed.

The microstructure of the material was also studied through 3D computed tomography analysis. A phoenix v|tome|x L300 (General Electric Sensing and Inspection Technologies) In order to identify the damage mechanisms of the studied composites under tensile loading, SEM examination (Zeiss DSM-962 and Zeiss Gemini FEG-SEM) were performed on asprocessed and damaged composites (post mortem observations) [START_REF] Ramdane | Etude et modélisation du comportement mécanique de CMC oxyde/oxyde[END_REF]. The damaged composites consisted of specimens loaded at different stresses, e.g. 22 %, 59 %, 63 % and 85 % of the average ultimate tensile stress.

X
As for cyclic loading tests, these stresses were determined from monotonic tensile tests and correspond to different points on the stress-strain curves: a point slightly after the average damage threshold, a point just before the average failure point and two intermediate points.

The gauge lengths of these specimens were observed to identify the damage mechanisms of the composite. The observation of as-processed material and damaged specimens required a specific preparation. First, they were impregnated by an epoxy resin (in a vacuum chamber for a good impregnation of the microporosity), in order to (i) avoid further damaging of the samples during polishing and (ii) have a flat polished surface, the fibres and the matrix being equally polished. Then, the samples were cut in the middle planes, parallel and perpendicular to the loading direction (parallel and perpendicular to the fibres in the case of as-processed material). Finally, the obtained cross-sections were mirror polished for SEM observation.

In order to detect the possible manufacturing defects and to avoid using damaged plates, non-destructive inspections were performed, in transmission configuration, by infrared pulse thermography and ultrasonic scanning in immersion. The infrared camera used was a CEDIP Jade LWIR (Long Wave InfraRed) covering a 7-9 µm wavelength range, with a 20 mK noise equivalent temperature difference and a frequency acquisition of thermal images of 200 Hz.

As for ultrasonic scanning technique, normalised transmitted energy maps were studied to detect damage. Two Panametrics transducers (General ElectricSensing) were used: M309 type emission transducer and V309 type reception transducer, with a 5 MHz nominal centre frequency. The plates were scanned with a 1 mm step, horizontally and vertically. Both nondestructive evaluation (NDE) methods enabled the detection of large defects, such as delamination. The pre-damaged areas of as-processed plates were removed. Only undamaged areas were used to extract testing samples.

Two non-destructive testing (NDT) methods have been used in this study for the determination of the damage threshold and for damage monitoring: passive infrared thermography and observation of the edge of the specimens by optical microscopy. In both cases, damage was detected at a very late stage, just before failure (for both 0° and ±45° loading directions).

Results and discussion

Microstructure

The total open porosity was 24 ± 2 %. The fibre volume fraction, deduced from the total porosity, was 49 %. The main advantage of mercury intrusion porosimetry compared to Archimedes' method is that it provides information about pore distribution (see Fig. 3). The composite exhibited two types of porosity: matrix microporosity and macroporosity. The limit between micro-and macroporosity was set at 0.1 µm. The microporosity, on which the weak matrix principle is based, was 13 ± 1 % of the total porosity of the composite and represented a volume fraction of 32 ± 1 % of the matrix. This fraction stands in the 30-40 % porosity range required for a weak matrix. The microporosity of the material is visible in the SEM micrograph presented in Fig. 4.a. The macroporosity was 12 ± 1 % of the total porosity of the composite. The macroporosity is due to a lack of matrix (see Fig. 4.b). The presence of macropores in oxide/oxide CMCs was reported in other studies [START_REF] Simon | Progress in processing and performance of porous-matrix oxide-oxide composites[END_REF] but not quantified. Fibre volume fraction and porosity data measured on the studied material are summarized in Table 1.

Concerning the macropores, two categories were identified: intraply and interply macropores.

This distinction was assessed by means of X-ray tomography analysis and SEM observations, performed on as-processed materials. Approximately 65 % of the macroporosity, measured by analysis of X-ray tomography data, consisted in interply macropores whose distribution was heterogeneous. On the contrary, intraply macropores appeared to be randomly distributed. An example of X-ray tomography analysis is given in Fig. 5.

Besides macroporosity, as-processed materials exhibited matrix sintering shrinkage microcracks attributed to the sintering process (see Fig. With regard to NDE techniques, the results obtained by infrared thermography and ultrasonic scanning were similar and showed that the plates were heterogeneous, some of them showing delaminated areas. This is in accordance with the previous observation by X-ray tomography pointing out the heterogeneous distribution of interply macroporosity. The example of a plate showing intensive delamination, identified by infrared thermography and ultrasonic scanning, is given in Fig. 6.

Mechanical properties

In-plane mechanical properties of the Nextel™610/alumina composite were studied through uniaxial monotonic and cyclic loading in the fibre direction (warp direction) and in ±45° fibre orientation. Uniaxial loading in the fibre direction provides information on the fibre-dominated properties whereas uniaxial loading in off-axis ±45° direction provides matrix-dominated properties. Nine specimens were cut from six different plates for monotonic tensile tests in the fibre direction and seven specimens were cut from three different plates for ±45° fibre orientation tests. For both directions, three specimens cut from different plates were used to perform cyclic tensile tests (three plates for the fibre direction tests and two plates for ±45° fibre orientation tests). As expected considering the heterogeneous microstructure of the material, curves and mechanical properties of the composite were scattered. Nevertheless, tensile curves exhibited the same shape for each direction of loading. The mechanical properties of the material are summarized in Table 2.

Mechanical properties in the fibre direction

Average ultimate tensile stress and strain of the composite are 260 ± 37 MPa and 0.30 ± 0.09 %, with a Young's modulus of 134 ± 19 GPa. The response of the studied composite under monotonic tensile loading in the fibre direction shows a first linear segment, a second quasi-linear segment and, in some cases, a third non-linear segment (see Fig. 7 and Fig. 8). The first linear segment constitutes the elastic region of the material. The damage threshold of the composite generally corresponds to the elastic limit. In most cases, the failure occurs at the end of the second quasi-linear segment. Such a behaviour, described as nearly linear up to failure, is common on oxide/oxide composites and has been reported by several authors [START_REF] Ruggles-Wrenn | Effects of steam environment on hightemperature mechanical behavior of Nextel™720/alumina (N720/A) continuous fiber ceramic composite[END_REF][8][17] [START_REF] Zawada | Characterization and high-temperature mechanical behavior of an oxide/oxide composite[END_REF]. However, in these studies, the slope of the second segment is slightly lower than the Young's modulus whereas, in the present case, the slopes of the first two segments are clearly different, as shown in Fig. 8.

The curved part of the tensile curves, between the elastic region and the second quasi-linear segment, indicates a brutal damage increase and can be considered as a damage threshold.

The same assumption can be made regarding the non-linear third segment exhibited by some specimens. Young's modulus and maximum stress and strain of the studied composite are higher than those of all-oxide composites, reinforced with Nextel™610 fibres and with similar fibre volume fractions, reported in the literature [START_REF] Simon | Progress in processing and performance of porous-matrix oxide-oxide composites[END_REF][17] [START_REF] Jurf | Advances in oxide-oxide CMC[END_REF]. However, the several alloxide CMCs described do not exhibit the same porosity and fibre volume fraction and were produced by different processes, which makes the comparison difficult.

The damage threshold, which is generally not reported in the literature, is essential in order to compare different materials and to design thermostructural parts. The stress and strain at the damage threshold deduced from the tensile curves are 48 ± 8 MPa and 0.04 ± 0.01 %, respectively (see Table 2).

Cyclic loading tensile tests indicate low residual strains (see Fig. 9.a). The thermal residual stress and strain evaluated by the common intersection point method, established by Steen [START_REF] Steen | Tensile mastercurve of ceramic matrix composites: significance and implications for modelling[END_REF], are 63 MPa and 0.03 %. The hysteresis loops exhibited by the material indicate frictional sliding between fibres and matrix. Such a phenomenon could be expected in the studied material since the weak matrix enables fibre/matrix debonding, hence increasing frictional sliding.

Mechanical properties in ±45° fibre orientation

The response of the composite under monotonic tensile loading in ±45° fibre orientation shows a first linear segment, corresponding to the elastic region, and a curved second segment (see Fig. 7). As reported in previous studies on different materials [START_REF] Zok | Mechanical properties of porous-matrix ceramic composites[END_REF] and the material exhibits hysteresis loops. The area of these loops is slightly larger in this loading direction. This phenomenon can be attributed to a more extensive fibre/matrix debonding occurring for this fibre orientation. Indeed, as this will be developed in further section 3.3, fibre/matrix debonding occurs simultaneously for both fibre directions, warp and weft, when the composite is loaded in ±45° fibre orientation. Thus, there is more frictional sliding and the area of the hysteresis loops is larger in this direction than in the fibre direction.

Out-of-plane mechanical properties

The interlaminar shear strength of the composite, as measured by 3-point bending tests, is 20 ± 2 MPa. This value is higher than those reported by Simon [START_REF] Simon | Progress in processing and performance of porous-matrix oxide-oxide composites[END_REF] for Nextel™610/mullite composites (12.5 -14 MPa) and Nextel™720/mullite composites (10 -11.5 MPa). The material developed at Onera is thus expected to be more resistant to delamination than the composites described in the literature with lower interlaminar shear strengths.

Identification of the damage mechanisms

As aforementioned, the damage mechanisms of the material were studied through post mortem SEM observations and in situ SEM tensile testing; these complementary methods enabled to suggest the following damage scenario of the composite under tensile loading.

At early stage of loading, matrix cracking occurs. Considering the SEM observations on the as-processed material, it can be assumed that, when the damage threshold is reached, preexisting matrix microcracks propagate while new cracks are initiated. Nevertheless, it is difficult to assess which phenomenon occurs first. The two phenomena probably occur simultaneously, depending on the local stress state in the material. We can also assume that the matrix cracking phenomenon starts at stresses lower than the apparent damage threshold (i.e. deduced from the stress-strain curves). In this case, early damage is not sufficient to induce a macroscopic loss of elasticity. Besides, matrix cracking occurs throughout loading, up to the composite failure. This phenomenon is clearly visible in the SEM micrographs (see Fig. 10) acquired during in situ tensile testing. According to Kostopoulos et al. [START_REF] Kostopoulos | On the identification of the failure mechanisms in oxide/oxide composites using acoustic emission[END_REF] who identified the failure mechanisms in oxide/oxide composites using acoustic emission, matrix cracking occurs all along loading with an abrupt increase at a certain critical load.

Matrix cracks have two origins: (i) low strength of the microporous matrix and (ii) macroporosity. Indeed, macropores act as stress concentrators and, located in matrix-rich regions or in fibre tows (i.e. intraply macropores due to a lack of matrix), are responsible for crack initiation resulting in extensive matrix cracking. Crack initiation at macropores is visible in Fig. 11.a and Fig. 11.b, in specimens loaded to failure in fibre direction and in ±45° fibre orientation.

Matrix cracks then propagate in the matrix and are deflected at the fibre/matrix interface when reaching a fibre. This phenomenon is in accordance with weak matrix principle. Matrix cracks first propagate around the fibre tows (regardless of the loading direction, as observed in Fig. 11.a and Fig. 11.b), then in the fibre tows. When the fibres of a tow are not in contact, cracks are not deviated around the tow but penetrate readily in the tow, as illustrated in Fig.

12.

Tow peripheral cracks, by debonding entire tows from the matrix, facilitate their movement within the matrix, leading to fibre extension. In the case of fibre direction loading, the extension of the longitudinal fibre tows leads to an increase in the undulation amplitude of the transversal tows. The internal stress system resulting from this phenomenon is responsible for interply matrix cracking which, considering fibre/matrix debonding in this weak matrix composite, contributes to delamination. According to Kostopoulos et al. [START_REF] Kostopoulos | On the identification of the failure mechanisms in oxide/oxide composites using acoustic emission[END_REF], the disorientation and disruption of the woven reinforcement occurring after fibre/matrix debonding starts shortly before the composite final failure, when fibre failure becomes the dominant mechanism.

In the case of off-axis loading, the movement of the tows enhances in-plane shear in the different plies, thus increasing delamination. Regardless of the loading direction, the presence of interply macropores, revealed by SEM observations and X-ray tomography, enhances delamination.

After penetrating the fibre tows, cracks propagate within the tows, isolating the fibres from one another, especially in the case of load bearing tows.

Finally, failure occurs when the amount of fibres bearing the load is not sufficient. The damage mechanisms of the composite are illustrated in Fig. 13.

Nextel™610/alumina composite exhibits extensive pull-out and, in many cases, delamination. SEM observations of the fracture surface show that fibre failure events are uncorrelated (see Fig. 14), the distance between the fracture surfaces of different fibres belonging to the same tow attaining a few millimetres. This asserts the fact that the weak matrix allows separating the fibres from each other by fibre/matrix debonding within the tows.

Levi et al. [START_REF] Levi | Processing and performance of an all-oxide ceramic composite[END_REF], Zok et al. [START_REF] Zok | Mechanical properties of porous-matrix ceramic composites[END_REF] and Simon [START_REF] Simon | Progress in processing and performance of porous-matrix oxide-oxide composites[END_REF] reported the same phenomenon for other weak matrix oxide/oxide composites.

In this material, macropores, regardless of their dimension, seem to be the main cause of extensive matrix cracking, leading to a loss of stiffness of the composite. Nevertheless, inter-and intraply macropores do not have the same effect, even though both act as stress concentrators. Intraply macropores are responsible for a loss of stiffness, increasing the strain capacity of the material, whereas interply macropores promote delamination, which makes them prejudicial to the 2D composites.

Besides the previous observations, it is noteworthy that the damage pattern is correlated to the direction of the reinforcement: cracks are parallel and perpendicular to the fibre tows, regardless of the loading direction.

Concluding remarks

The microstructure of Nextel™610/alumina composite was studied through several techniques, namely porosimetry, X-ray tomography and SEM observations, which appear to be complementary.

The as-processed material exhibits a heterogeneous microstructure. The composite is predamaged, showing shrinkage matrix microcracks and macropores. Two types of macropores can be differentiated: inter-and intraply macropores. The quality of as-processed plates and samples was assessed by X-ray tomography, infrared thermography and ultrasonic scanning techniques. These last two techniques provided similar information. However, infrared thermography inspection being faster, it thus appears to be the most suitable technique for initial damage detection.

An extensive study of the damage mechanisms of the composite, through post mortem SEM observations and in situ tensile testing in a SEM, showed that intraply macropores promote crack initiation, enabling energy dissipation and leading to a loss of stiffness, whereas interply macropores, acting as stress concentrators, promote delamination. In-plane and outof-plane mechanical properties of the studied Nextel™610/alumina composite were higher than those reported in the literature. However, considering the difference in fibre volume fraction and porosity of the materials described in different studies, the comparison is rather difficult.

In order to avoid delamination in the Nextel™610/alumina composite, reducing the amount of interply macropores appears mandatory. This can be achieved by improving the manufacturing process in order to increase the cohesion between the plies. Enhancing fibre/matrix bonding could also restrain delamination, but it would also deteriorate the inplane mechanical properties of the composite by limiting crack deflection at fibre matrix interface. Limiting the amount of macropores is currently under study at Onera.

The thermal in situ damage monitoring of the mechanical tests by passive thermography did not enable to visually locate the progressive damage of the tested oxide/oxide composite specimens. The late detection of damage, i.e. just before failure, may be attributed to the high amount of energy released by extensive fibre failure, whereas crack propagation in the matrix, occurring throughout loading, is not energetic enough to be detected. Random and isolated fibre failures occur during loading but the failure of a single filament is not energetic enough to be detected, contrary to the failure of entire tows which can be detected and lead to the specimen failure. The determination of the damage threshold of such materials will require an analysis of the local temperature evolution during mechanical loading.

Adapting other damage monitoring methods to oxide/oxide CMC is also currently under consideration.

grateful to D. Boivin, Y. Renollet and N. Horezan (Onera, DMSM) who co-developed and conducted SEM in situ testing, and to M. Bejet (Onera, DMSC) who designed specific jaws adapted to the in situ tensile testing of parallelepipedic shaped composite specimens.

Figures:

Fig.

Tables:

Table 1. Porosity and fibre volume fraction as determined by mercury intrusion porosimetry (five samples from five different plates). Table 2. Mechanical properties of Nextel™610/alumina composite subjected to tensile loading in the fibre direction and in ±45° fibre orientation.

  -ray tomograph was used for data acquisition, with a 180 kV source, enabling a 10.16 mm x composite thickness.

  4.c). These cracks were predominantly normal to the direction of the woven reinforcement and concentrated in the matrix-rich regions. The presence of sintering shrinkage cracks is common in CMCs and has been reported in several studies[1][3][8][9][10][START_REF] Ruggles-Wrenn | Creep behavior of Nextel™610/monazite/alumina composite at elevated temperatures[END_REF].

1 .Fig. 2 .Fig. 3 .Fig. 4 .

 1234 Fig. 1. Schematic diagram of crack deflection at the fibre/matrix interface in a weak matrix CMC.

Fig. 5 .Fig. 6 .

 56 Fig. 5. X-Ray tomography analysis of Nextel™610/alumina composite. a. Reconstruction of the zone of interest. b. Distinction between macroand micropores. c. Example of the two largest macropores of the zone of interest.

  Fig. 7.Stress-strain curves of Nextel™610/alumina composite under tensile loading in 0° and ±45° fibre orientations. For each loading direction, extreme and average curves are represented.

  Fig. 8.Stress-strain curves of Nextel™610/alumina composite under tensile loading in 0° fibre direction: distinction between the different domains of the curves, namely the elastic domain (first linear part), the second linear part and, in some cases, the third part (generally non linear). The damage threshold is determined as the end of the elastic domain.

  Fig. 9. Stress-strain curves of Nextel™610/alumina composite under cyclic tensile loading in a. 0° fibre direction and b. ±45° fibre orientation.

Fig. 10 .

 10 Fig. 10. In situ SEM observations of Nextel™610/alumina composite under tensile loading: matrix crack opening and propagation throughout loading. At elevated load, matrix cracks span the transverse fibre tows.

  Fig. 11. SEM observations of Nextel™610/alumina composites tensile tested to failure. a. Longitudinal section of a specimen loaded in the fibre direction: matrix cracks perpendicular to the plies deflected at the fibre/matrix interface, crack initiation at macropores. b. Longitudinal section of a specimen loaded in ±45° fibre orientation: matrix cracks perpendicular to the plies deflected at the fibre/matrix interface, crack initiation at macropores.

Fig. 12 .

 12 Fig. 12. Schematic diagram of the deflection of a matrix crack arriving on a fibre tow when a. the fibres are in contact or b. the fibres are separated by the matrix.

Fig. 13 .

 13 Fig. 13. Schematic diagrams of the damage mechanisms of Nextel™610/alumina weak matrix composite subjected to tensile loading. a. As-processed material, showing a macropore and a sintering matrix crack. b. Opening of the pre-existing matrix crack, matrix cracking, crack initiation phenomenon at the macropore. c. Crack deflection at the tow/matrix interface. d. Propagation of the deflected cracks around the fibre tows. e. Crack propagation within the fibre tows.

Fig. 14 .

 14 Fig. 14. Fractography of a specimen loaded in the fibre direction: the single filaments belonging to the same tow are separated, resulting from intensive fibre/matrix debonding within the tow.

Acknowledgements

The collaboration with Herakles is gratefully acknowledged. This work was supported under the PRC Composites, a French research project funded by DGAC, involving Safran Group, Onera and CNRS. The authors would like to thank Safran Composites for conducting X-ray tomography on our specimens. The authors are thankful to Dr. F. Laurin (Onera, DMSC) for developing the method of analysis of X-ray tomography data used in this study. The authors would like to thank M.-H. Ritti (Onera, DMSC) for her help in materials processing, F. Passilly, Dr. J.-M. Roche and Dr. B. Lamboul (Onera DMSC) for NDT inspections and A. Mavel (Onera, DMSC) for the mechanical testing of the specimens. The authors are also


Cătălin M Jureschi
email: catalin.jureschi@eed.usv.ro

Jorge Linares
email: jorge.linares@uvsq.fr

Aurelian Rotaru
email: rotaru@eed.usv.ro

Marie Hélène Ritti
email: marie-helene.ritti@onera.fr.

Michel Parlier
email: michel.parlier@onera.fr

Marinela M Dîrtu
email: marinela.dirtu@uclouvain.be

Mariusz Wolff
email: mariusz.wolff@uclouvain.be

Yann Garcia
email: yann.garcia@uclouvain.be

Pressure sensor via optical detection based on a 1D spin transition coordination polymer

Keywords: spin crossover, pressure sensors, optical detection, smart materials, material characterization

We have investigated the suitability to use the 1D spin crossover coordination polymer [Fe(4-(2'-hydroxyethyl)-1,2,4-triazole) 3 ]I 2 •H 2 O, known to crossover around room temperature, as a sensor via optical detection using various contact pressures up to 250MPa. A dramatic persistent colour change is observed. The experimental data, obtained by calorimetric and Mössbauer measurements, have been used for a theoretical analysis, in the framework of the Ising-like model, of the thermal and pressure induced spin state switching.

I. Introduction

Due to the current need of processing and storage of an increasing amount of information it becomes necessary to use miniaturized devices including multifunctional components. Particularly, molecular materials are considered as a promising class for implementation in molecular nano-devices. To this aim, spin crossover (SCO) materials including 3d 4 -3d 7 transition metals were selected due to their potential miniaturization and suitable applications in memory devices, sensors, switches, and displays [START_REF] Gutlich | Spin crossover -An overall perspective[END_REF][START_REF] Linares | Pressure and Temperature Spin Crossover Sensors with Optical Detection[END_REF][START_REF] Bousseksou | Molecular spin crossover phenomenon: recent achievements and prospects[END_REF][START_REF] Gutlich | Spin state switching in iron coordination compounds[END_REF]. These materials can switch between two distinguished states, namely the low-spin (LS) state and the high-spin (HS) state by applying external perturbations such as temperature, pressure, light or a magnetic field [START_REF] Gutlich | Spin state switching in iron coordination compounds[END_REF]. As a result, different optical, magnetic, vibrational and structural properties are observed, which are more pronounced in the case of Fe(II) complexes. The most widely used perturbation to induce a spin state transition is the temperature, but a pressure stimulus also induces modifications on the spin state. Thus SCO materials are considered as pressure sensors [START_REF] Linares | Pressure and Temperature Spin Crossover Sensors with Optical Detection[END_REF][START_REF] Ksenofontov | Spin crossover behavior under pressure of Fe(PM-L) 2 (NCS) 2 compounds with substituted 2′-pyridylmethylene 4-anilino ligands[END_REF]. A remarkable family concerns 1D Fe(II) coordination polymers with 1,2,4-triazole ligands [START_REF] Roubeau | Triazole-based one-dimensional spin-crossover coordination polymers[END_REF][START_REF] Garcia | Crystal engineering of Fe II spin crossover coordination polymers derived from triazole or tetrazole ligands[END_REF], which show marked colour changes making this family of materials suitable to be used as temperature and pressure sensors with optical detection [START_REF] Linares | Pressure and Temperature Spin Crossover Sensors with Optical Detection[END_REF]. In Figure 1a, the pressure applied to the Fe(II) material is lower than the critical pressure, P c , and the SCO compound appears white (HS state). The white surface can reflect the laser light to the light detector, which can be connected to a computer. In Figure 1b, the pressure applied to the material under study is above the critical pressure, P c , causing the SCO compound to turn purple (LS state). The purple surface absorb the light or reflect it to a lesser degree, depending on the wavelength of the light used. The change in pressure can thus be detected optically at a critical point. The continuous developments of chemical systems in our laboratory allows the selection of the best SCO systems with special requirements concerning the SCO temperature region, which need to be located around the room temperature region as well as the width of their bistability domain which need to be as large as possible. Following these conditions, we decided to focus on the 1D chain [Fe(hyetrz) 3 ]I 2 •H 2 O (hyetrz = 4-(2'-hydroxyethyl)-1,2,4-triazole) as a suitable candidate. This material is known to exhibit a cooperative and thermochromic spin transition from LS (S = 0, violet) to HS (S = 2, white) around the room temperature region [START_REF] Garcia | Synthesis and spin crossover characteristics of polynuclear 4-(2'-hydroxy-ethyl)-1,2,4-triazole Fe(II) molecular materials[END_REF]. In this work we show that [Fe(hyetrz) 3 ]I 2 •H 2 O can be used as a visual detector of strong mechanical contact pressure from 250 to 2500 bars. This result is supported by thermal dependence of the optical density measurements, 57 Fe Mössbauer spectroscopy and differential scanning calorimetry. A correlation of SCO properties with the pressure contact is made thanks to an Ising-like model. These results provide the basis for the construction of a marker pressure device.

Temperature study

The colour change of the sample from white (HS) to violet (LS) was monitored through variabletemperature optical reflectivity measurements in a dry nitrogen atmosphere at 2 K/min. The thermal dependence of the optical density has been recorded spectroscopically at a quasimonochromatic wavelength of λ = 550(50) nm (Figure 2). The hysteretic SCO behaviour around room temperature is confirmed for [Fe(hyetrz

) 3 ]I 2 •H 2 O at T c ↑ = 292 K and T c ↓ = 275
K by optical reflectivity measurements (Figure 2). The high colour contrast of the sample observed on cooling can be tracked by recording temperature dependence of the reflectance spectra. In the LS state, the reflectance spectrum shows a band centred on 680 nm, in good agreement with the pink colour of sample. By increasing the temperature, a blue shift of the spectrum maximum is observed, followed by a broadening of the reflectance spectrum, in good agreement as well with the white colour of the sample in the HS state. The optical characterization has been completed by differential scanning calorimetric (DSC) measurements that were undertaken to confirm the spin transition temperature range as well as to determine, quantitatively, the thermodynamic parameters to be used in our Ising model described hereafter. DSC profiles were recorded at 10 K/min in the heating and cooling modes around the room temperature region as shown in Figure 3. The DSC curve shows an endothermic peak at T max ↑ = 293 K and an exothermic peak at T max ↓ = 280 K, in fair agreement with the transition temperatures observed from optical measurements. The difference seen in the lower branch of the spin transition curve is associated to the lower scan rate (2K/min) used for the optical measurements. The enthalpy and entropy variations were determined as ∆H = 16.42 kJmol -1 and ∆S = 57.42 J.mol -1 K -1 . The experimentally measured entropy variation accounts for an electronic contribution, Rln5 = 13.4 J.mol -1 K -1 , and a vibrational contribution of 41.77 J.mol -1 K -1 . These values were corrected taking into account the active sites fraction associated to the spin transition which was accurately determined from complimentary Mössbauer spectroscopy measurements detailed below. Indeed on cooling to 78 K, a single quadrupole doublet with isomer shift δ LS = 0.50(1) mm.s -1 was observed indicating 100% LS ions. The presence of a quadrupole splitting ∆E Q LS = 0.24(1) mm.s -1 , indicates a distortion for the octahedral as expected within a 1D chain [START_REF] Dîrtu | Iron(II) spin transition 1,2,4-triazole chain compounds with novel inorganic fluorinated counteranions[END_REF]. At 313 K, a temperature for which the compound is expected to have undergone the spin transition, according to optical reflectivity measurements (Figure 1), the spectrum shows a major quadrupole doublet (87%) attributed to HS Fe II (δ = 0.99(1) mm.s -1 and ∆E Q = 2.49(2) mm.s -1 ) and a minor one (13%) corresponding to LS Fe II (δ = 0.36 mm.s -1 and ∆E Q = 0.17

Pressure study

Pressure experiments were carried out on a home-made micromechanical device. The sample was deposited on a metal plate and covered with a scotch tape (Figure 5). As shown in Figure 5, the material is white at room temperature and ambient pressure which is characteristic of the HS state. The pressure was applied at selected spots on the sample, released and a photograph was taken again. As it can be seen, for a threshold value of the applied pressure of ca. 30-40 MPa, the SCO powder switches from white (HS) to pink (LS) and retain its colour (spin state) when pressure is released.

Phase diagram based on Ising like model

Various physical methods, such as magnetic susceptibility measurements, Mössbauer spectroscopy, optical and vibrational spectroscopy, X-ray diffraction and heat-capacity measurements [START_REF] Gutlich | Spin state switching in iron coordination compounds[END_REF], are used to describe the switching behaviour of SCO materials. Due to the expensive cost of such equipments used in physical characterizations, several theoretical models are used to predict the SCO properties. This includes models such as atom-phonon coupling [START_REF] Nasser | First order high-spin/low-spin phase transition induced by acoustic-phonons[END_REF][START_REF] Nasser | Two-step spin conversion and other effects in the atom-phonon coupling model[END_REF][START_REF] Boukheddaden | Elastic interaction among transition metals in onedimensional spin-crossover solids[END_REF][START_REF] Rotaru | Size and pressure effects in the atomphonon coupling model for spin crossover compounds[END_REF][START_REF] Rotaru | Re-entrance phase and excited metastable electronic spin states in one-dimensional spin crossover compounds explained by atom-phonon coupling model[END_REF][START_REF] Boukheddaden | Static and Dynamical Aspects of a Phonon-Induced Spin-Crossover Transition -Local Equilibrium Approach[END_REF], Ising-like [START_REF] Wajnflasz | Low spin »-« High spin » dans les complexes de Fe 2+[END_REF][START_REF] Slichter | Pressure-Induced Electronic Changes in Compounds of Iron[END_REF][START_REF] Bousseksou | Ising-like model for the 2step spin crossover[END_REF][START_REF] Linares | Analytical solution of 1D Ising-like systems modified by weak long range interaction -Application to spin crossover compounds[END_REF][START_REF] Jureschi | Thermo-and piezochromic properties of [Fe(hyptrz)]A 2 •H 2 O spin crossover 1D coordination polymer: towards spin crossover based temperature and pressure sensors[END_REF] or mechano-elastic [START_REF] Nasser | Two-step spin conversion and other effects in the atom-phonon coupling model[END_REF] which have been used to describe the role played in SCO behaviour by the lattice architecture, the influence of short-and long-range intermolecular interactions when applying external perturbation such as a temperature and/or a pressure variation.

In the Ising like model, to represent the two states of a SCO molecule, Wajnflasz and

Pick [START_REF] Wajnflasz | Low spin »-« High spin » dans les complexes de Fe 2+[END_REF] have introduced a fictitious spin operator σ which can take the value 1 + when the molecule is in the HS state or 1 when the molecule is in the LS state with respective degeneracies HS g and LS g . For a system with non-interacting molecules the Hamiltonian is given by:

1 2 N i i H σ = ∆ = ∑ (1)
where 1 N i= ∑ denotes the sum over all SCO molecules, N is the number of molecules and ∆ represents the energy gap between the HS and LS states.

Over the last few years an important effort has been done in order to explain the influence of both short-and long-range interaction which induce different types of spin transition shapes [START_REF] Linares | Analytical solution of 1D Ising-like systems modified by weak long range interaction -Application to spin crossover compounds[END_REF][START_REF] Jureschi | Thermo-and piezochromic properties of [Fe(hyptrz)]A 2 •H 2 O spin crossover 1D coordination polymer: towards spin crossover based temperature and pressure sensors[END_REF][START_REF] Chiruta | Analysis of long-range interaction effects on phase transitions in two-step spin-crossover chains by using Ising-type systems and Monte Carlo entropic sampling technique[END_REF][START_REF] Chiruta | Analysis of the Hysteretic Behaviour of 3D Spin Crossover Compounds by Using an Ising-Like Model[END_REF][START_REF] Chiruta | Lattice architecture effect on the cooperativity of spin transition coordination polymers[END_REF][START_REF] Klokishner | Effect of hydrostatic pressure on phase transitions in spincrossover 1D systems[END_REF], with or without hysteresis. The wideness of the hysteresis loop is determined by the long-range interaction parameter.

The Hamiltonian for interaction molecules, as proposed in [START_REF] Rotaru | Re-entrance phase and excited metastable electronic spin states in one-dimensional spin crossover compounds explained by atom-phonon coupling model[END_REF] can be written as follows: , , ln 2

B i i j i j i i i j i k T g H J G σ σ σ σ σ ∆ - = - - ∑ ∑ ∑ (2)
where, , i j ∑ is the sum over nearest neighbour spins, B k is the Boltzmann constant, / HS LS g g g = is the degeneracy ratio, J and G represent the short and long rage interaction parameters and σ is the average value of the fictitious magnetization.

In the mean field approach the average of the fictitious magnetization is given by:

2 ln tanh 2 B B k T g k T σ σ  Γ + -∆  =     (3)

where

Jq G Γ = + is the reduced interaction parameter and ( , )

T p ∆ = ∆
is the reduced gap energy. The HS molar fraction is given by: 1 2

HS n σ + = (4)
The pressure dependence of the energy gap is:

0 ( , ) T p p V δ ∆ = ∆ + , ( 5
) where 0 ( , 0) T p ∆ = ∆ = , V
δ is the volume variation and p is the external applied pressure.

From the DSC measurements, and taking into account 87% of spin switching determined by Mössbauer measurements, the entropy value calculated ∆S = 57.42 J.mol -1 K -1 give a degeneracies ratio state of the two spin state, g, such as lng = 6.906 and the enthalpy variation give a gap energy ∆ 0 /k B =1978.6K. For the volume variation (δV) during the spin transition (LS→HS, HS→LS), we used a typical value of volume change δV=100A 3 .

The thermal and pressure dependence of the HS molar fraction, n HS , at ambient pressure and temperature computed for [Fe(hyetrz) 3 ]I 2 •H 2 O, i.e. for a complex displaying an hysteretic spin transition around room temperature is displayed in Figure 6

Figure 6

Thermal dependence at 1 bar (left) and pressure dependence at 300 K (right) of the HS molar fraction, n HS , derived from the Ising like model. The parameters values are ∆ 0 /k B =1978.6K, δV=100A 3 , lng = 6.906 and Γ = 360 K A (p-T) phase diagram was calculated to understand the influence of temperature and pressure on the SCO behaviour. The calculated phase diagram in pressure-temperature coordinates using the same parameters as in Figure 6, is displayed in Figure 7. As can be seen in Figure 7, the transition from the HS to the LS state can be achieved either by decreasing the temperature or by increasing pressure. When the SCO compound located at point A, i.e. at given temperature T 1 and pressure P 1 , is heated until it reaches point B at temperature T 2 ≥ T up at a constant pressure P 1 , a switch back to the HS state is observed. From B, if the temperature is decreased until it gets back to point A (T 1 ,P 1 ) then the compound remains in the HS state. By applying a pressure, the compound switches from HS to LS states at point C (T 1 ,P 2 > P down ). If now the pressure is lowered back to P 1 (point A), the compound will remain in the LS state. In other words, inside the hysteresis width, the compound will keep the state of origin. When coming from the LS state it will remain in the LS state and when coming from the HS state it will retain the HS state as shown in Scheme 1.

Conclusion

In this work we proved the feasibility of pressure detection using a molecular SCO based sensor/marker operating at ambient temperature. For the titled compound we obtained a threshold value of the contact pressure of about 30 MPa to induce irreversibly the colour change of the molecular material, due to the spin state switching form HS to LS state. These results open important perspectives for molecular materials in pressure sensing applications, making possible their insertion in piezo-and thermo-chromic paints that will allow the visual detection of mechanical collisions, a very important issue in the aeronautic and automotive industries.

Figure 1

 1 SCO material used as pressure sensor: (a) the SCO material reflects the light to the detector, indicating that the pressure is below the critical point; (b) the SCO material does not reflect the light to the detector, indicating that the pressure is above the critical point.

Figure 2 .

 2 Figure 2. Thermal evolution of reflectance spectra recorded on [Fe(hyetrz) 3 ]I 2 •H 2 O in the solid state (left) and of the optical density at λ = 550 nm (right).

Figure 3 .

 3 Figure 3. DSC curves for [Fe(hyetrz) 3 ]I 2 •H 2 O over the temperature range 265-303 K.

  [START_REF] Ksenofontov | Spin crossover behavior under pressure of Fe(PM-L) 2 (NCS) 2 compounds with substituted 2′-pyridylmethylene 4-anilino ligands[END_REF] mm.s -1 ).Thus [Fe(hyetrz) 3 ]I 2 •H 2 O undergoes an incomplete ST on warming with 13% of non switching sites, which need to be taken into account for the enthalpy determination associated to the spin state change.

Figure 4 .

 4 Figure 4. 57 Fe Mössbauer spectra for [Fe(hyetrz) 3 ]I 2 •H 2 O at 78 K (left) and 313 K (right).

Figure 5 (

 5 Figure 5 (Left): sample holder showing the SCO compound at room temperature on its sample holder covered with a scotch tape. (Right): Enlarged view of the sample evidencing colour

  . The transitions temperatures at atmospheric pressure are T c ↑ = 292.5 K and up and T c ↓ = 278.6 K, which are in excellent agreement with the ones detected by optical reflectivity (Figure1). The transitions temperatures at 300 K are P c ↑ = 19.2 MPa and up and P c ↓ = 9.3 MPa.

Figure 7

 7 Figure 7Pressure-temperature phase diagram calculated for a SCO system switching around the room temperature region using the following parameters values: ∆ 0 /k B =1978.6K, δV=100A3 , lng = 6.906 and Γ = 360 K.

Scheme 1

 1 Representation of the switching mechanism induced by temperature and/or by pressure.

		/ t p ↑ ↓	LS	/ t p ↑ ↓
	HS	/ t p ↓ ↑	HS	/ t p ↓ ↑	LS

Acknowledgments

We acknowledge financial support from the Fonds National de la Recherche Scientifique (FNRS) (FRFC 2.4537.12), Romanian National Authority for Scientific Research, CNCS -UEFISCDI, (project number PN-II-RU-TE-2011-3-0307), Romanian Academy -WBI, WBI-Roumanie, WBI-Tournesol and the Walloon region (MELMINK). We also acknowledge the COST actions CM1305 and MP1202. M. M. Dîrtu is chargé de recherches of FNRS.

Experimental

[Fe(hyetrz) 3 ]I 2 •H 2 O was prepared as a solid white powder as described in ref [START_REF] Roubeau | Triazole-based one-dimensional spin-crossover coordination polymers[END_REF]. The reflectance was determined by means of an optical microscope (Olympus BX51) equipped with a photodiode and a 543 nm interferential filter, while the optical spectra have been recorded by using an UV-Vis Ocean Optics spectrometer mounted on the optical microscope. The sample temperature was controlled using a Linkam THMS600 liquid nitrogen cryostage. During the experiment the temperature was changed at a rate of 2K/min. Differential scanning calorimetry measurements were carried out in a He (g) atmosphere using a Perkin-Elmer DSC Pyris 1 instrument equipped with a cryostat and operating down to 98 K. The purge gas was N 2(g) . Temperatures and enthalpies were calibrated over the temperature range 98-300 K using the solid/solid and liquid/solid transitions of pure cyclopentane (P99%, Acros) [START_REF] Rotaru | Calorimetric measurements of diluted spin crossover complexes [Fe x M 1-x (btr) 2 (NCS) 2 ]•H 2 O with MII = Zn and Ni[END_REF]. The calibration sample was introduced in an aluminum pan and hermetically sealed using an encapsulating press. The calibration was made at a scan rate of 10 K/min. The characteristic temperatures, which were assigned to the crystal/crystal transitions of cyclopentane, were obtained by the extrapolation of the onset peak temperatures. An empty aluminium pan, identical to the one used for the sample, was used as a reference to obtain a reliable baseline. The system produces or uptakes energy in order to keep the temperature of the compound identical to the reference. This energy difference between two resistances (in mW) is transformed by PYRISTM DSC Software 7.0 in specific heat C p (J /mol/K). The DSC measurement of the sample was carried out at a scan rate of 10 K/min, in warming and cooling modes. 27.7 mg were encapsulated at room temperature in aluminum pans and hermetically sealed. The sample was maintained at room temperature for 5 min in order to allow the system to equilibrate, and was further cooled down from 298 to 98 K. The sample was maintained at 98 K for 5-10 min to reach equilibrium, followed by a similar scanning mode as on cooling between 98 K and room temperature. 57 Fe Mössbauer spectra were recorded in transmission geometry with a Wissel Mössbauer spectrometer equipped with a 57 Co(Rh) radioactive source (Cyclotron Ltd, Russia) operating at room temperature, and a Reuter-Stokes proportional counter. The samples were sealed in iron free aluminum foil and mounted on a nitrogen Oxford Instruments bath cryostat. The spectra were fitted to the sum of Lorentzian functions by a least-squares refinement using Recoil 1.05 Mössbauer Analysis Software [START_REF] Lagarec | Mössbauer Spectral Analysis Software for Windows 1.0[END_REF]. Pressure tests were performed on a micromechanical device developed at Onera. The cylindrical indenter with a contact surface of 4 mm 2 is fixed in the upper part and the sample holder is placed on the movable part. This latter receives the load sensor and two deformable blades are used to guide the column designed to transmit the load between the sample holder and the pressure sensor. A stepper motor is used to move the platform via a screw-nut system. An Entran load sensor (FC231100000250L) with a capacity of 1100 N is used. During contact between indenter and sample, the moving part stops automatically when a force of 1 N is measured. For indention testing, the moving speed was 5 µ/s and the displacement is stopped when the required load is reached.


Assessment of a sequential extraction protocol by examining solution chemistry and mineralogical evolution

Nicolas Maubec (1), Hélène Pauwels (1), Hervé Noël (2), and Xavier Bourrat (1)

(1) BRGM, 3 Avenue Claude Guillemin, 45060 ORLEANS Cédex 2, FRANCE, (2) Géo-Hyd, 101 rue Jacques Charles, 45160 OLIVET, FRANCE Knowledge of the behavior of heavy metals, such as copper and zinc in sediments, is a key factor to improve the management of rivers. The mobility of these metals, which may be harmful to the environment, depends directly on their concentration and speciation , which in turn depend on physico-chemical parameters such as mineralogy of the sediment fraction, pH, redox potential, salinity etc ... [START_REF] Anderson | Column leaching and sorption experiments to assess the mobility of potentially toxic elements in industrially contaminated land[END_REF][START_REF] Sterckeman | Factors affecting trace element concentrations in soils developed on recent marine deposits from northern France[END_REF][START_REF] Van Oort | Impacts of long-term waste-water irrigation on the development of sandy Luvisols: consequences for metal pollutant distributions[END_REF]. Several methods based on chemical extractions are currently applied to assess the behavior of heavy metals in soils and sediments. Among them, sequential extraction procedure is widely used in soil and sediment science and provides details about the origin, biological and physicochemical availability, mobilization and transports of trace metals elements. It is based on the use of a series of extracting reagents to extract selectively heavy metals according to their association within the solid phase [START_REF] Cornu | Extractions séquentielles et spéciation des éléments trace métalliques dans les sols naturels[END_REF] including the following different fraction : exchangeable, bound to carbonates, associated to oxides (reducible fraction), linked to organic matter and sulfides (oxidizable fraction) as well as silicate minerals so called residual fraction [START_REF] Hickey | Chemical partitioning of cadmium, copper, nickel and zinc in soils and sediments containing high-levels of heavy metals[END_REF][START_REF] Tessier | Sequential extraction procedure for the speciation of particulate trace-metals[END_REF]. Consequently sequential extraction method is expected to simulate a lot of potential natural and anthropogenic modifications of environmental conditions [START_REF] Arey | Immobilization of uranium in contaminated sediments by hydroxyapatite addition[END_REF][START_REF] Brannon | Fixation, transformation, and mobilization of arsenic in sediments[END_REF][START_REF] Hickey | Chemical partitioning of cadmium, copper, nickel and zinc in soils and sediments containing high-levels of heavy metals[END_REF][START_REF] Force | Redistribution of trace elements from contaminated sediments of Lake Coeur d'Alene during oxygenation[END_REF][START_REF] Tessier | Sequential extraction procedure for the speciation of particulate trace-metals[END_REF]. For three decades, a large number of protocols has been proposed, characterized by specific reagents and experimental conditions (concentrations, number of steps, extraction orders and solid/solution ratio) [START_REF] Das | Metal speciation in solid matrices[END_REF][START_REF] Gomez Ariza | Selectivity assessment of a sequential extraction procedure for metal mobility characterization using model phases[END_REF][START_REF] Quevauviller | Evaluation of a sequential extraction procedure for the determination of extractable trace-metal contents in sediments[END_REF][START_REF] Rauret | Extraction procedures for the determination of heavy metals in contaminated soil and sediment[END_REF][START_REF] Tack | Chemical speciation and fraéctination in soil and sediment heavy-metal analysis -a review[END_REF], but it appeared that several of them suffer from a lack of selectivity of applied reagents: besides target ones, some of them are able to leach several solid phases.

In this context, the aim of the present study is to investigate the effectiveness and the selectivity of different reagents for metal extraction from target geochemical fraction. It is based on solid analyses with the use of X-ray diffraction and a scanning electron microscopy (SEM) coupled to a microRaman spectrometer in conjunction with chemical analyses of extracting solutions at each step. This methodology provides the opportunity to assess more accurately the effect of each reagent. The study focuses on extraction of Cu and Zn from sediment samples collected at two sites from river banks and characterized by presence of Quartz, Feldspar K, Micas, Kaolinite but with differences regarding accessory phases (pyrite, organic matter, iron oxyhydroxide, calcite). The interaction of the samples with eight different reagents was assessed and compared (Ca(NO 3 )2 and CaCl2 for the exchangeable fraction; buffered solutions of sodium acetate/acetic acid at pH = 5.5 and pH = 5 for the acido-soluble fraction; hydroxylamine hydrochloride and a solution of ammonium oxalate/oxalic acid for reducible fraction; hydrogen peroxide and sodium hypochlorite for the oxidizable fraction. In-depth characterization of solid residue at each step allowed proposing the best protocol for both metals.