Julie Lorthios
email: julie.lorthios@edf.fr

Matthieu Mazière

Xavier Lemoine

Philippe Cugy

Jacques Besson

Anne-Françoise Gourgues-Lorenzon
email: anne-francoise.gourgues@mines-paristech.fr

Anne- Françoise Gourgues-Lorenzon

Fracture behaviour of a Fe-22Mn-0.6C-0.2V austenitic TWIP steel

Keywords: Plastic yield anisotropy, ductile fracture, anisotropic fracture criterion, TWIP steels, forming limit diagram

Introduction

Among Advanced High Strength Steels (AHSSs), high carbon, high manganese austenitic twinninginduced plasticity (TWIP) steels offer an excellent compromise between strength and formability (see e.g. [Cornette et al, 2005]) particularly suitable to lightweight automotive applications. Relationships between their microstructure and mechanical properties have been recently reviewed [Bouaziz et al, 2011]. The high work-hardening of these steels relies on complex interactions between solid solution carbon atoms, dislocation glide and intense mechanical twinning, owing to their low stacking fault energy close to room temperature.

The high carbon content of these steels induces thermal activation of dislocation glide (contribution of ~70 MPa to effective stress close to room temperature for a Fe-22Mn-0.6C steel [Bouaziz et al, 2011]) and pronounced dynamic strain ageing (DSA) attributed to complex interactions between carbon atoms and dislocations [Lee et al, 2011]. Serrated yielding and negative strain rate sensitivity have commonly been reported under uniaxial tension under low strain rates [e.g. Bayraktar et al, 2004, Scavino et al, 2010]. The high strain hardening stems from mechanical twins acting as obstacles to dislocation glide [see e.g. Rémy, 1978, Idrissi et al, 2010]).These steels also exhibit a strong anisotropy in plastic yielding, in spite of weak initial texture in the cold rolled + annealed state [START_REF] Gutierrez-Urrutia | The effect of grain size and grain orientation on deformation twinning in a Fe-22wt.% Mn-0.6wt.% C TWIP steel[END_REF], Barbier et al, 2009]. Constitutive modelling of TWIP steels made use of either Hill's anisotropic criterion [e.g. [START_REF] Chung | Formability of TWIP (twinning induced plasticity) automotive sheets[END_REF] or anisotropic criteria developed by Barlat et al. [Chung et al, 2011, Xu et al, 2012] and satisfactorily reproduces the uniaxial tensile and hole expansion behaviour of these materials by proper fitting of model parameters [START_REF] Chung | Formability of TWIP (twinning induced plasticity) automotive sheets[END_REF], Xu et al, 2012].

In spite of their high formability, high manganese TWIP steels are sensitive to abrupt slant fracture under a variety of loading modes. This uncommon behaviour makes diffuse necking criteria [Dorn andThomsen 1947, Swift 1952] as well as localised necking criteria [Hill, 1952, Marciniak and[START_REF] Marciniak | Limits strains in the processes of stretch-forming sheet metal[END_REF]] not suitable to describe the formability of this material [START_REF] Chung | Formability of TWIP (twinning induced plasticity) automotive sheets[END_REF]. Intense deformation bands that develop under uniaxial tension have been reported to assist initiation of abrupt fracture [De Cooman et al, 2008, Scavino et al, 2010, Wang et al, 2011]. Under biaxial stretching [Scavino et al, 2010] as well as stamping of real parts [Allain et al, 2008], no localisation band could be experimentally detected whereas the sensitivity to loading rate is negative and slant fracture still may occur without any noticeable reduction in thickness. Fracture surfaces exhibit two populations of dimples. Coarser dimples, often elongated along the rolling direction (RD), have been reported [Xu et al, 2012, Lorthios et al, 2010]. They have been attributed to either decohesion of MnS stringers from the matrix [Xu et al, 2012] or to the coalescence of groups of smaller voids, whose anisotropic shape was consistent with anisotropy in plastic strain [Lorthios et al, 2010]. The volume fraction of such voids is yet very low (typically below 0.2% even close to the fracture surface) and the vast majority of the fracture surface at room temperature exhibits submicrometer-sized dimples. Their nucleation and coalescence mechanisms, although attributed to the coalescence of nano-voids by [Abbasi et al, 2009], are still unclear [e.g. Bouaziz et al, 2011]. Except for the very low quantity of coarser voids, no sign of damage has been evidenced by metallographic observation of fractured specimens.

Phenomenological fracture criteria have been developed to tentatively predict the forming limit diagram of TWIP steels [e.g. Chung et al, 2011b]. They make use of one, or several of the following parameters: the maximum shear stress, τ max (Tresca-type criterion) [Bouaziz et al, 2011], the stress triaxiality, the equivalent stress and the Lode parameter, which is a function of the second and third invariants of the deviatoric stress tensor that enables e.g. axisymmetric stress states to be conveniently distinguished from plane strain. While the plastic strain anisotropy of high-manganese TWIP steels has been extensively studied [Xu et al, 2012, Barbier et al, 2009], their anisotropy in fracture has only been scarcely addressed [Lorthios et al, 2010, Xu et al, 2012]. In the above mentioned fracture criteria, invariants of stress tensors have generally been derived using a von Mises-based formalism, whatever the formalism of constitutive equations. On the other hand, a fracture criterion expressed with the same formalism as that of constitutive equations has already been used (with a Hill-type yield criterion) by [Beese et al, 2010] to predict slant fracture of an aluminium alloy but, to the authors' knowledge, such an approach has never been reported for TWIP steels.

The objective of this study is to characterise both yield and fracture anisotropy of a Fe-22Mn-0.6C-0.2V TWIP steel, to derive a failure criterion that could be applied for the determination of forming conditions. The fracture criterion makes use of the same formalism as that used in constitutive equations to provide a homogeneous framework for the mechanical analysis of forming tests. In addition, this study aims at getting more insight into the physical fracture mechanisms, including the effect of mechanical twinning and stress state.

Materials and experimental procedures

Material

A Fe-22Mn-0.6C austenitic TWIP steel was provided as a bare 1.5-mm-thick cold-rolled and annealed sheet. Its chemical composition is reported in Table 1. The average grain size is about 3 µm (Fig. 1a), with some segregation bands visible throughout the sheet thickness (Fig. 1b). Due to the low thickness of these bands, no electron probe microanalysis was conducted to determine their chemistry but they are supposed to stem from manganese segregation. They did not interact with the appearance of fracture surfaces, so that they are no longer considered in the following of this paper. The average chemical composition yields a low value of the stacking fault energy (about 20 mJ.m -2 at room temperature) together with thermally-activated dislocation glide, so that this material may deform with intense mechanical twinning without any detectable austenite-to-martensite phase transformation [e.g. Scott et al, 2005, Allain et al, 2004] in agreement with Schumann's diagram [Schumann, 1972]. Vanadium carbides are known to increase the yield strength by interacting with dissociated partial dislocations [Yen et al, 2012]. A resulting contribution to the effective stress of around 140 MPa has been reported for a steel of similar chemical composition and grain size [Scott et al, 2011]. Carbides also decrease the sensitivity of the steel to delayed hydrogen-assisted cracking [Malard et al, 2012]. In the following, RD, TD, ND and DD respectively denote the rolling direction, transverse direction, sheet normal and diagonal direction (i.e. at 45° from RD in the (RD, TD) plane). All investigations were carried out in laboratory air, at room temperature with full-thickness specimens, unless otherwise stated.

Uniaxial tensile tests

Uniaxial tensile (UT) dogbone specimens (60 x 12.5 mm 2 in gauge) were cut along RD, TD and DD and pulled in tension with an elongation rate of 4.10 -4 , 4.10 -3 and 4.10 -2 s -1 in a 250 kN MTS servohydraulic machine equipped with a 50 kN load cell. Two specimens were used per testing condition. Axial elongation was recorded using a 9-mm-gauge extensometer. The reduction in width was monitored using a clip extensometer attached to the central part of the gauge region. The Lankford coefficient, r d (d being the loading direction), was calculated as the ratio between logarithmic plastic strain along the specimen width to that along ND. To this aim, an isochoric assumption was made, consistently with reported very low level of porosity even at fracture [Lorthios et al., 2010].

High strain rate uniaxial tension tests were carried out along RD using an Instron high-speed machine, with a specific dogbone specimen having a 65 x 9 mm 2 or a 65 x 11 mm 2 gauge part. One specimen was used per testing condition and geometry. Both geometries yielded similar results. The load was calculated from elastic deformation of strain gauges fixed to one end of the specimen. A laser extensometer was used for axial strain measurements. For the two selected elongation rates (50 s -1 and 110 s -1 ), self-heating of the specimen did not exceed 100°C during the test, which should not change the deformation mechanism of this material. To investigate the possible influence of mechanical twinning on the fracture mechanism, additional UT tests were carried out at 400°C along RD and TD with a 100 kN MTS servohydraulic machine fitted with a radiation furnace. The temperature was monitored using a K-type thermocouple spotwelded onto the centre of the gauge part. Two other thermocouples were spot-welded at either end of the gauge part. A smaller specimen geometry (30 x 4 mm 2 in gauge) was selected and a muffle was placed around the specimen to reduce the temperature gradient to less than ±10°C. The average axial strain was continuously monitored using an extensometer attached to the gauge part (gauge length 10 mm). As deformation under UT loading may be very heterogeneous in Hadfield steels [START_REF] Dastur | Mechanism of Work Hardening in Hadfield Manganese Steel[END_REF], Owen and Grujicic, 1998, Canadinc et al, 2008, Renard et al, 2010] and TWIP steels [Alain, 2004, Chen et al, 2007, Lebedkina et al, 2009, Zavattieri et al, 2009], the effect of specimen geometry was tested at room temperature, along RD (one specimen), with an elongation rate of 4.10 -2 s -1 by comparison with the above mentioned UT tests on larger specimens. Tensile curves were superimposed with reproducibility better than 3%. At 400°C, the strain rate sensitivity was evaluated along RD with elongation rates of 10 -5 , 4.10 -5 , 8.10 -5 , 4.10 -4 and 4.10 -2 s -1 . Two specimens were used for each condition and yielded very reproducible results, both for tensile curves and fracture modes.

Reverse shear tests

The Bauschinger effect was evaluated by reverse shear tests carried out with a 45 x 30 mm 2 specimen geometry and a loading rate of 500 N.s -1 . The shear direction was parallel to TD and displacement along RD was precluded. Fiducial marks together with a camera were used to derive the engineering shear strain, e 12 . The engineering shear stress, σ 12 , was calculated as the applied load (i.e. along TD)

divided by the specimen width and thickness. Monotonic tests were carried out up to fracture, which occurred close to the grips. As a consequence, these tests could not be used to assess damage and fracture under shear loading. The Bauschinger effect was determined by loading the specimen up to a maximum value, e max , of e 12 (0.025, 0.05, 0.10, 0.15 or 0.20), then unloading and applying immediately the load along the reverse direction until e 12 reached at least -0.1.

Tensile tests on notched specimens

Double edge, U-notched tensile specimens with a minimum width of 10 mm, a notch depth of 2 mm and a notch radius of 2 mm (DENT2), 1 mm (DENT1) or 0.5 mm (DENT05) were pulled in tension using the same set-up as for room temperature UT tests at low strain rates. For every notch radius, two specimens were tested along each of the RD, TD, and DD directions, with prescribed load line displacement rate of 0.5 mm.min -1 . The notch opening (recorded using clip extensometers) vs. time curves (one for either notch) did not differ by more than 3.4%, so that only values of notch opening averaged over both notches are reported here for a given specimen. The effective stress (i.e. load divided by initial value of the area between notches) was plotted as a function of notch opening. For some additional specimens (one DENT2 and one DENT05 for both RD and TD directions), a digital image correlation system was used to monitor strain fields. It consisted of a 2448 x 2050 pixel² camera fitted with the Aramis TM software [GOM]. A random speckle was deposited on one side by use of acrylic painting. 0.35 x 0.35 mm 2 regions were used for image correlation, with a recording rate of 0.2 up to 1 frame per second. Logarithmic strain fields were derived from displacement field calculations, with a spatial resolution of 0.2 to 0.4 mm depending on specimen geometry.

To investigate a variety of loading modes by still using a UT testing facility, more complex specimens were also selected. An in-plane double edge notched shear (DENS) specimen was designed after that used by [START_REF] Brünig | [END_REF] for aluminium alloys, both with or without spherical-cap-shaped side grooves (specimen geometry illustrated below together with the results).The aim of these tests was to characterise the fracture behaviour at near-zero hydrostatic stress. Specimens were cut along RD and TD and pulled in tension with an electrical-mechanical Zwick testing machine, with prescribed load line displacement of 0.17 mm.s -1 . In contrast to [START_REF] Brünig | [END_REF], specimen ends were clamped in grips to avoid premature failure that could occur at high loads around holes. Some slipping was sometimes observed, so that the strain field monitoring technique was systematically used in the same conditions as previously mentioned.

Double side notched tensile (DSNT) specimens were used to localise strain within a band tilted by 45° with respect to the sheet plane and loading direction, to reproduce the geometry of slant fracture [McClintock & Argon, 1966]. The notches were perpendicular to the loading direction, but not in front of each other. The digital image correlation technique was used to monitor strain field at one edge of the specimen. Two specimens were tested along each of the RD and TD directions, with prescribed load line displacement rate of 0.1 mm.s -1 .

To tentatively interrupt crack propagation and study damage development ahead of the crack tip, single edge notched Kahn-type specimens [START_REF] Astm | B 871-96 standard test method for tear testing of aluminum alloy products[END_REF] were pulled in tension parallel to RD, the crack propagating along TD (RD-TD orientation) or vice versa (TD-RD orientation). Owing to the high forming ability of the steel, extensive buckling occurred and hindered mechanical analysis of these tests. However, crack propagation could be arrested to investigate damage development by metallographic techniques.

Metallographic observation of tested specimens

All fracture surfaces were examined by scanning electron microscopy (SEM), in the secondary electron imaging mode, using either a Leo 1450VP or a Shottky-field emission gun Zeiss DSM 982 Gemini microscope fitted with X-ray energy dispersive spectrometry (EDX). Kahn-type specimens containing arrested cracks were observed in cross-section after conventional polishing followed by a colloidal silica finish. Microtomographic observations close to one of these cracks have already been reported elsewhere [Lorthios et al, 2010]. The high levels of strain hardening, fracture strength and fracture elongation, typical of TWIP steels, are illustrated in Fig. 2 and Table 2. The strain rate sensitivity was negative at lower strain rates with the main contribution coming from work hardening. For higher strain rates, the sensitivity became positive. At room temperature, serrations were observed for strain rates lower than 4.10 -2 s -1 and axial elongation higher than about 20%. For this reason, only values of strain averaged over the gauge part of smooth specimens are reported in the following. Some experimental scatter was found for yield strength (less than 8%) and tensile strength (less than 7%) due to strain instabilities. These were related to propagation of Piobert-Lüders-type bands along the gauge part of the tensile specimen [e.g. Allain et al, 2008]. Such phenomena are referred to as "dynamic strain ageing" (DSA) in the following. They induced strong variation in the values of Lankford coefficients. Consequently, some additional UT tests were carried out using the DIC technique to extract the Lankford coefficients reported in Table 2.

Experimental results and discussion

Uniaxial tensile behaviour

Anisotropy in strength was low (about 30 MPa) at the onset of plastic deformation, consistently with the very weak crystal orientation texture and morphological texture in this material. In a fine-grained Fe-22Mn-0.6C TWIP steel, [Barbier et al, 2009] showed that anisotropy increases with work hardening, due to the gradual development of texture components induced by both plastic slip and intensive twinning. Anisotropy in plastic strain was higher than that in strength.

For low strain rates, fracture occurred abruptly, right after propagation of a Piobert-Lüders-type band, probably somewhat before the Considère strain instability criterion was reached. This induced scatter in fracture elongation. For this reason, all available experimental values have been reported in Table 2.

At 400°C, strain hardening was still high but strength and fracture elongation were much lower than at room temperature. Pronounced necking occurred well before fracture, indicating a more stable ductile fracture process than at room temperature. No serration was detected in tensile curves. The strain rate sensitivity was negative for strain rates lower than 4.10 -4 s -1 . At 400°C, pseudo-DSA is thus still present, in certain conditions, even if no mechanical twinning is expected to occur [e.g. [START_REF] Bouaziz | [END_REF]. Conversely, mechanical twinning may occur together with positive strain rate sensitivity (e.g. for higher strain rates at room temperature, see Fig. 1b). Correlation between DSA and mechanical twinning is thus unclear in this material. There is yet a clear correlation between the occurrence of mechanical twinning and sensitivity to abrupt fracture.

Bauschinger effect

A strong Bauschinger effect was evidenced in shear loading (Fig. 3). It increased with increasing e max (Fig. 3a). The kinematic contribution to hardening was about 50% of the applied stress, so that its evolution with e max was not linear. This result is in agreement with those reported by [Bouaziz et al, 2008] in a vanadium-free Fe-22Mn-0.6C alloy. The contribution of isotropic hardening increased almost linearly with e max , the slope being about 1740 MPa per unit strain (Fig. 3b). ) and isotropic hardening (r 12 ) to the shear stress. σ 12_0 is the initial yield stress in these shearing conditions. Data for the vanadium-free alloy are taken from [Bouaziz et al, 2008].

Tensile behaviour of notched specimens

Double edge notched specimens

As in UT tests, only a slight anisotropy in strength and in fracture elongation was found with DENT specimens (Fig. 4a). The reproducibility of curves was better than 5 MPa (for a given notch opening) except for DENT05 specimens along TD, for which a difference by about 10 MPa was noticed for notch opening lower than 0.5 mm. Notch opening at fracture differed by less than 0.1mm between the two tests per condition. The notch radius did not significantly influence the load vs. notch opening curve (Fig. 4b) but strongly affected notch opening at fracture. No serration was detected in these curves. Constraints induced by the notches may have forced intense deformation bands, if any, not to propagate across the entire specimen. Although fracture occurred in an unstable manner, the DIC system was able to capture one image of a DENT05 specimen loaded at 0.5 mm.min -1 , with one crack initiated from each of the notch roots (Fig. 4c). The logarithmic reduction in thickness at fracture, evaluated in the SEM, was about -0.45±0.05 at each notch root, i.e., at crack initiation, and almost the same (-0.42 ±0.02) along the fracture ligament (representative of crack propagation). The values were identical to those measured in UT for strain rates higher than 4.10 -2 s -1 (-0.45±0.01), whereas for lower strain rates (4.10 -4 and 4.10 -3 s -1 , respectively), the logarithmic reduction in thickness was slightly lower (-0.39±0.03 and -0.42 ±0.02, respectively).

Double edge notched shear specimens

The experimental behaviour of DENS specimens is reported in Fig. 5, together with model predictions that are introduced in Section 4. Load vs. notch opening curves of ungrooved DENS specimens (Fig. 5a) exhibited a slight upward curvature for displacements higher than about 3 mm. This was due to a gradual change in specimen geometry that induced stiffening (Fig. 5b). This was not observed by Brünig et al. [Brünig08], whose aluminium alloy did not exhibit the outstanding strain ability of TWIP steels. Anisotropy in strength was not negligible (about 6%). Significant rotation of the specimen was precluded by the fixture system. However, the bending stiffness of the machine was not high enough to suppress translation of the specimen ends perpendicular to the loading axis (Fig. 5b). At fracture, deformation was not constrained to the central part of the gauge region but extended from one notch root to the other, with a maximum value of about 0.4.

Grooved specimens fractured with much lower notch opening than ungrooved ones (Fig. 5c), and showed an anisotropy in strength of about 9%. Unlike for the other tests, specimens loaded along RD were slightly stronger than those loaded along TD. Fracture initiation occurred close to the centre (Fig. 5d) and strain was kept constrained between the grooves (Fig. 5d), with a maximum shear strain of about 0.59 at fracture. At the centre of the grooves, pure shear strain (at least from in-plane components) dominated throughout the test, so that this specimen was considered to be tested in conditions close to pure shear.

Double side notched tensile specimens

The experimental behaviour of DSNT specimens is reported in Fig. 6, together with model predictions that are introduced in Section 4. Load vs. load line displacement curves (Fig. 6a) showed no change in curvature, a slight anisotropy in strength and, again, abrupt fracture with no detectable softening. The local behaviour (Fig. 6b) was determined by using an effective stress (i.e., load divided by specimen width and by the distance between notch roots along the sheet thickness, i.e. 1 mm) and by averaging axial elongation over the gauge region of DIC measurements (length 3.42 mm, centred at specimen centre). In the model, due to the mesh geometry, axial elongation was averaged over 3.50 mm. The slight difference (about 2%) between gauge length used in DIC measurements and in the model, respectively, is not expected to significantly affect the results. The crack propagated at about 47° from the loading axis (see inlet in Fig. 6b). Determination of fracture strain was out of reach of DIC measurements due to a limit in spatial resolution but it was higher than 0.3. With the spatial resolution of DIC, strain localisation seems to have occurred between the two side notches (Fig. 6c). No variation in width was detected. The loading mode was thus close to generalised plane strain conditions.

Damage development and fracture mechanisms

Dimpled ductile fracture occurred for all specimens (Fig. 7). As already evidenced in this material [Lorthios et al, 2010, Xu et al, 2012], coarse cavities, strongly elongated along RD whatever the specimen geometry and loading direction, were found in fracture surfaces (Fig. 7a). Microtomography measurements [Lorthios et al, 2010] showed that (i) these cavities resulted from coalescence of smaller voids aligned along RD, (ii) their volume fraction was very low even close to the fracture surface (about 0.06% in average and locally not higher than 0.2% over 1mm from the fracture surface), and (iii) their anisotropic shape was consistent with plastic yield anisotropy as measured in UT with Lankford coefficients. They were mainly located between mid-thickness and quarter-thickness. A possible influence of the banded structure cannot be ruled out even if the distance between such coarse cavities (about 20-50 µm) does not scale with that between segregation microbands (a few µm, see Fig. 1b). No damage could be detected in any metallographic cross-section, even at the tip of arrested cracks (Fig. 7d). Fracture surfaces mainly consisted of submicrometer-sized dimples (0.1-0.5µm in size), involving very limited void growth, as reported in literature [Bayraktar et al, 2004, Abbasi et al, 2009, Scavino et al, 2010, Chung et al, 2011] Consequently, void initiation mechanisms were difficult to observe. In some UT and DENS specimens, tiny spherical particles (about 40 nm in diameter) were found inside dimples (Fig. 7c). Their size was compatible with the size distribution of vanadium carbides (VC) in vanadium bearing Fe-22Mn-0.6C austenitic steels [Dumay, 2008], with particles having an average diameter of 7 nm, 3% of the particles being in the range 20-60 nm. In a vanadium-free material, larger dimples of about 1-2 µm in size were found instead [Allain, 2004].

To determine whether void nucleation in the investigated material could entirely be attributed to VC particles higher than about 20 nm in size (i.e. to about 3% of the VC particles), the number density of these particles was roughly estimated. It was assumed that almost all vanadium atoms were trapped into stoichiometric (NaCl type) precipitates. The number density of precipitates was estimated by considering that all precipitates had the same diameter (7 nm). Then, by only considering 3% of these particles, the distance between coarser precipitates was calculated to about 30 µm. Even if strong approximations were made in this simple model because the actual size and spatial distributions of VC particles were not known, such a distance is by far higher than the size of finer dimples in Fig. 7. This suggests that voids could initiate from either much finer precipitates that could not be imaged in the SEM, or even elsewhere, owing to e.g. complex interactions between carbides and mechanical twins. Even if cross-section observations in the TEM would be possible by taking a thin foil close to the fracture surface (e.g. by focused ion beam milling techniques), further deformation due to void coalescence may have strongly modified the deformation microstructure (dislocation distribution and microtwins) after dimple nucleation.

Fracture surfaces of specimens tested at room temperature were similar whatever the elongation rate (Fig. 7, Fig. 8), hence, whatever the strain rate sensitivity (Fig. 5). This suggests that the void nucleation and coalescence mechanisms are not significantly affected by the strain rate sensitivity, possibly because intense strain localisation induces high strain rates at a very local scale, whatever the loading conditions applied in the present study.

At 400°C, slant fracture was also observed. Primary dimples were more widely open than at 20°C (Fig. 9). The origin of these elongated cavities was clearly evidenced to be coalescence of smaller, aligned cavities as already observed at room temperature. Finer dimples were significantly larger (more than 1 µm in average size) than those found at room temperature (compare Fig. 7 with Fig. 9). This confirms that despite its negative strain rate sensitivity (at least for low elongation rates), the considered material was less sensitive to strain localisation at 400°C than at room temperature.

Mechanical analysis of experiments and derivation of a fracture criterion

To derive a fracture criterion from available experimental data, stress and strain fields had to be determined at fracture initiation as accurately as possible. To this aim, mechanical analysis of all tests was carried out by using a finite element approach using the general purpose, implicit code "Z-set" [START_REF] Besson | [END_REF]. As physical mechanisms of fracture initiation were still unclear, a phenomenological model was selected. Unlike recent micromechanical approaches [Allain et al, 2004, Bouaziz et al, 2008, Gil Sevillano, 2009, Soulami et al, 2011, Favier and Barbier, 2012, Steinmetz et al, 2013], and unlike models based on dislocation and twinning theory (e.g. [Liu et al., 2015]), the constitutive behaviour of this material was also described using a phenomenological approach to provide a unique framework for plastic flow and fracture modelling. The DSA phenomenon was not taken into account in this model. As fracture occurred abruptly at room temperature, without significant softening or extended damage development, only a fracture initiation criterion was considered. No coupling between plastic yield and damage development was thus introduced in the model.

Constitutive equations

From experimental results, anisotropy in strength appears from the beginning of plastic yielding and gradually develops, so that it was chosen to represent it by an anisotropic yield criterion, together with either only isotropic ("R" model) or mixed isotropic and kinematic ("R + X" model) contributions to hardening. Isotropic elastic properties were assumed (Young's modulus of 170 MPa and Poisson ratio of 0.3). Three yield criteria were considered, namely, a von Mises yield criterion, the anisotropic criterion developed by Barlat et al. [Barlat et al, 1991] and the Bron-Besson anisotropic yield criterion [Bron and Besson, 2004]. Full identification of constitutive parameters and numerical simulation of all tests showed very similar results for the two anisotropic criteria, so that only results obtained using the criterion by [Barlat et al., 1991] are reported here. Isotropic hardening was described using a Vocetype formulation; description of kinematic hardening was derived from the Armstrong-Frederick formulation. Note that such a model has already been used to describe the anisotropic plastic flow in various materials under complex loading paths [Laurent et al, 2009;Shinohara et al, 2010;[START_REF] Vladimirov | Anisotropic finite elastoplasticity with nonlinear kinematic and isotropic hardening and application to sheet metal forming[END_REF]]. In the following, the reference frame is that defined by RD, TD and ND. All calculations were carried out using a finite strain framework but for the sake of simplicity, model equations presented below are expressed using small strain formalism. The finite strain formulation adopted in this work was based on the use of a local objective frame as proposed in [Besson et al., 2009].

Observer invariant stress and strain rate measures were defined by transport of the Cauchy stress and strain rate into the corotational frame characterized by rotation Q (function of both space coordinates and time).

In this model, following [Barlat et al, 1991], a linear transformation (fourth-order tensor, S ) was used to define a modified deviator dev T of any second-order tensor T ; for a symmetric tensor T , dev T is expressed as follows using Voigt notations:

T S T dev : = , with                         + - - - + - - - + = 6 5 4 2 1 1 2 1 3 1 3 2 3 3 2 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 0 3 3 3 0 0 0 3 3 3 0 0 0 3 3 3 c c c c c c c c c c c c c c c S (1)
In Eq. (1), c i (i = 1-6) are adjustable material parameters that are linked to the (RD,TD,ND) reference frame. The anisotropic scalar measure, T B , of tensor T is a function of the eigenvalues

dev T 1 ≥ dev T 2 ≥ dev T 3 of dev
T , with a an adjustable material parameter:

a a dev dev a dev dev a dev dev B T T T T T T T / 1 1 3 3 2 2 1 2 1       - + - + - = (2)
The viscoplastic potential is a function of the stress tensor, σ , the kinematic contribution to hardening, X , and the isotropic contribution to hardening, R, as follows:

( ) ( ) R X R X f B - - = σ σ , , (3)
The viscoplastic strain rate tensor, p ε& , is calculated using a normality rule on function f, together with a power-law flow equation for the viscoplastic strain rate, p & :

σ ε ∂ ∂ = f p p & & with ( ) n f p         = 0 0 , 0 max σ ε& & and 0 ε& = 1 s -1 . (4)
In Eq. ( 4), 0 σ and n are constitutive material parameters.

Isotropic hardening is a function of the cumulative viscoplastic strain, p, using a linear term and a nonlinear term (Voce-type):

( ) ( ) bp Q p H R R - - + + = exp 1 0 (5)
For the "R+X" model, the kinematic contribution X to hardening is given by the additive combination of a linear term 1 X and a non-linear term 2 X , as follows:

2 1 X X X + = (6) 1 1 1 3 2 α & & C X = with p ε α & & = 1
for the linear term, and ( 7)

2 2 3 2 α& & C X = with 2 2 α ε α D p p & & & - =
for the non-linear term.

(8)

It was not possible to satisfactorily represent reverse shear test results using constant values of C and D, so that heuristic evolutionary equations were proposed as follows:

0 0 2 3 D D C C = and ( )                 - - - + = 0 0 1 0 exp 1 p p D D D D (9)
In Eq. ( 9), C 0 , D 0 , C 1 , D 1 and p 0 are material parameters.

The total number of parameters for the "R" model is 13: 7 for the viscoplastic potential, 2 for the viscoplastic flow rule, and 4 for isotropic hardening. Five additional parameters are needed to describe kinematic hardening in the "R+X" model.

Identification of constitutive parameters

Identification of constitutive parameters was based on quasistatic uniaxial tension results (tensile curves as well as Lankford coefficients, one test per direction i.e. 3 tests), reverse shear tests (5 tests) and load vs. notch opening curves for the DENT specimens (9 tests: 3 per direction, 3 per notch geometry). UT and reverse shear tests were modelled using a single representative volume element. DENT specimens were meshed in three dimensions by considering one-eighth of each specimen with usual symmetry conditions. Linear eight-node bricks with full integration were used with a mesh size fine enough for the load vs. notch opening curves to be represented without any dependence on the size of finite elements. To account for the very low strain rate sensitivity of the material, low values were fixed for both 0 σ and n, to keep a negligible viscous contribution to stress while facilitating integration of constitutive equations with respect to an elastic-plastic formulation. An implicit integration scheme was used for all calculations. In absence of experimental results involving mainly loading along ND, the values of c 5 and c 6 were arbitrarily set to 1.0. Remaining material parameters were optimised using a Levenberg-Marquardt algorithm. The resulting sets are reported in Table 3.

Table 3: Optimised sets of constitutive parameters for the "R" and "R+X" models.

Modulus Anisotropic yield potential Flow rule Isotropic hardening

Kinematic hardening The value of parameter a was obtained by numerical fitting. It is higher than the value of 8 that is recommended for in-plane loading of face-centred cubic structures by [Hosford, 1979, Hosford, 1985]. This value of 8 was fitted to upper-bound calculations of yield loci, assuming rotationally symmetric textures and conventional slip systems. The general method is described in detail in [Logan and Hosford, 1980]. As the considered TWIP steel was processed by a rolling + annealing schedule, rotational symmetry in crystal texture cannot be assumed in the present case. Moreover, twinning should be taken into account as a significant deformation mechanism in this material.

E (GPa) ν a c 1 c 2 c 3 c 4 c 5 c 6 0 σ (MPa) n R 0 (MPa) H (MPa) Q (MPa) b C 0 (MPa) C 1 D 0 D 1 p 0 "R
The yield locus, expressed by normalising stresses with respect to the yield stress under uniaxial tension along RD, has been plotted in Fig. 10 with those from [START_REF] Chung | Formability of TWIP (twinning induced plasticity) automotive sheets[END_REF], Xu et al., 2012].

For comparison purposes, the yield locus identified by [Busch et al., 2014] on a so-called TWIP/TRIP steel is also added. The higher value of parameter a in the present study leads to a less rounded shape of the yield surface (note that a value of a = 4 yields to an ellipse). The main point, here, is the lower symmetry (with respect to the (σ RD = σ TD ) axis) of the curve from the present model, compared with the other three ones. This is linked to the set of optimised parameters c i describing the anisotropic yield potential.

Figure 10: In-plane yield locus predicted from the present study (1), compared with those from literature data: (2) [START_REF] Chung | Formability of TWIP (twinning induced plasticity) automotive sheets[END_REF]; (3) [Xu et al., 2012]: (4) [Busch et al., 2014]. Stresses along RD and ND have been normalised by the yield stress under uniaxial tension along RD.

The UT behaviour is well described by the model, namely, within a relative error in strength of 5-7% for an elongation of 10%; this difference decreases with increasing elongation down to less than 1% at fracture. For a given elongation, lateral strain was only slightly overestimated (relative difference lower than 4%) and slightly underestimated (relative difference lower than 2%) for tests along RD and TD, respectively, so that strain anisotropy was well reproduced. The relative difference between measured and predicted load vs. notch opening curves of DENT specimens was lower than 6%, and better than 4% for a notch opening lower than 1 mm. The loading part of reverse shear tests was well described using the model, while isotropic hardening was obviously overestimated in the unloading part. Predictions of the "R+X" model were much closer to experimental measurements for this test, yet at the expense of accuracy in strain anisotropy under uniaxial tension (lateral strain overestimated by 7% and underestimated by 6% for tests along RD and TD, respectively). From Figs 5 and 6, load vs. displacement curves, as well as local strains close to the fracture locus for the last recorded experimental data before fracture, are very satisfactorily described with the "R" model. The "R+X" model is thus significantly more accurate than the "R" model only if unloading (or significant change in the loading mode) occurs during the tests.

Strain paths and estimated stress state at fracture initiation

Determination of fracture initiation conditions with the "R" model

Due to marked instability at high strains, UT specimens were not considered in the analysis of fracture initiation conditions. In fact, strain localisation bands induce local stress concentrations that require taking pseudo DSA and strain rate sensitivity into account to be properly described [Lorthios, 2011].

Here, focus was made on a model that is more easily usable in forming applications. Numerical simulation of tests on notched specimens were carried out using refined meshes (typical size of elements close to the notch root: 0.1 mm in the sheet plane and 0.2 mm along ND), in order to obtain accurate values of stress and strain fields close to fracture initiation sites.

Ungrooved DENS specimens were meshed using quadratic brick elements with reduced integration (mesh size 0.55 mm in the sheet plane, 0.5 mm along ND). Grooved DENS specimens were meshed with quadratic tetrahedral elements with reduced integration (element size in the gauge part: 0.16 mm in the sheet plane, 6 elements along ND). Displacement was prescribed along both RD and TD according to experimental measurements at the edges of the DIC window, close to the specimen fixtures (see inlets in Fig. 5). This allowed taking the bending stiffness of the load line into account.

DSNT specimens were modelled using quadratic brick elements with full integration. The size of elements in the notched part was 0.7 mm in the (RD,ND) plane. As the reduction in width was negligible in DSNT tests, the size of the elements was reduced down to 1 mm along TD together with generalised plane strain conditions (i.e., homogeneous value of deformation along TD for any slice of the specimen perpendicular to TD). For all of these tests, load vs. displacement curves and strain fields were in very good agreement with experimental data (Figs 4,5,6).

The value of the stress triaxiality ratio η B (i.e. the ratio of the hydrostatic stress, ξ, to the anisotropic equivalent stress, σ B ), the equivalent strain p (here denoted as ε B to explicitly make reference to the Barlat yield criterion), the ratio of minor to major strain α B , were determined at integration points that were closest to the fracture initiation site, for the same value of displacement as that experimentally observed at fracture. The values are reported in Table 4. For DENT specimens, the values were taken at the notch root, at mid-thickness. For ungrooved DENS specimens, the values of α B taken at the edge and at the centre of the notched region are very different, evidencing a strong gradient in α B . As the fracture initiation point could not be experimentally determined, these specimens were not used for quantitative determination of the fracture criterion. For grooved DENS specimens, the values were taken at mid-thickness, between the groove roots. Except for the DSNT specimen, the value of α B at the fracture initiation point remained almost constant throughout the test. For the DSNT specimen, α B started from near zero up to a cumulative strain of about 0.1 and then turned into -0.18. Such a nearly radial strain path was not the case, however, everywhere in the tested specimens: the high deformability of the TWIP steel led to a noticeable change in specimen geometry before fracture. For instance, in the central region of DENT specimens, a transition from plane strain tension toward uniaxial tension occurred during stress redistribution. Consequently, the stress and strain conditions at fracture were the same whatever the notch radius, yet with avoiding macroscopic strain instability observed in uniaxial tension. The DENT geometry could thus be regarded as useful to investigate fracture initiation under nearly uniaxial tension for such ductile materials exhibiting strain instabilities.

As global load vs. displacement curves were well predicted with the "R" model, and due to nearly constant values of α B at the fracture initiation point, it was decided to only use the "R" model to derive a fracture initiation criterion. In the following, data processed with the "R" model will be denoted as "experimental", in the forming limit diagram. They are reported with symbols in Fig. 11. Table 4: Characteristic values of stresses and strains at fracture initiation estimated with the "R" model, compared to experimental data. α is the ratio of minor to major strain. All specimens are loaded along RD. Superscripts * and a denote tests results from [START_REF] Dell | [END_REF] (also reported in [Bouaziz et al, 2011]) and internal ArcelorMittal tests [Bouaziz et al, 2011]

Influence of anisotropic plastic yield on macroscopic fracture initiation conditions

To quantify the difference between fracture initiation conditions determined using such a complex "R" model and those determined from much simpler, more commonly used modelling, the UT test along RD with an elongation rate of 4.10 -4 s -1 was modelled with a von Mises yield criterion and a Hollomon-type evolutionary equation of the flow stress:

H n p H K ε σ = , ( 10
)
ε p being the von Mises equivalent plastic strain (equal to the uniaxial plastic strain, in uniaxial tension, with the von Mises yield criterion). The fitted values of K H = 2605 MPa and n H = 0.43 allowed accurate prediction (better than 2% difference in strength) of the tensile curve for strains higher than about 0.1 (which is much lower than the fracture strain). This model, denoted as "VM-H" in the following, was used to analyse the same experimental database, using the same numerical procedure as for the "R" model except for the constitutive behaviour.

At fracture, compared to the more accurate "R" model, the "VM-H" model overestimates the value of stress triaxiality by 10%, and the value of cumulative plastic strain by 10 to 20% (in relative error values). It underestimates the value of equivalent stress (compared to that of σ B ) by about 200 MPa.

Thus, taking anisotropic yielding into account is essential to accurately determine fracture initiation conditions for this material. Consequently, only values determined with the "R" criterion are considered in the following.

"Experimental" forming limit diagram

To complete the experimental database of the previous section, the results of additional mechanical tests carried out using the same material were also processed using the "R" model to extract fracture initiation conditions for a wider range of loading paths [Dell andGese, 2007, Bouaziz et al, 2011]. The experimental procedure is reported in [START_REF] Hooputra | [END_REF], Dell et al, 2008] and results are also reported in Table 4. They included Marciniak-type forming tests with two different values of strain ratio α (experimentally measured from optical strain field monitoring) [Bouaziz et al, 2011], and a series of tests [START_REF] Dell | [END_REF] involving plane strain tension with a double side notched tensile specimen, equibiaxial stretching, uniaxial tension of a 20-mm-wide, 110-mm-long rectangular specimen bearing a central hole of 10 mm in diameter, shear loading of a specimen with two side groves, and "tilted plane tension" of a specimen with two side groves (in front of each other), perpendicular to the sheet thickness and making a 45° angle (in the sheet plane) with the RD loading direction. To process these additional data while minimising numerical issues such as contact between tooling and specimens, these tests were simulated using a single representative volume element with prescribed value of α. This value was determined experimentally for the Marciniak-type tests. An approximation with analytical formulae using a "VM-H" assumption was made for the tests from [START_REF] Dell | [END_REF].

Fracture initiation criterion

The fracture initiation criterion was determined from the above mentioned tests that involve nearly constant strain paths and loading along RD. The following results, therefore, should not be directly applied to complex loading paths encountered in forming operations.

As reported by [START_REF] Wierzbicki | [END_REF], Zadpoor et al, 2009] in an aluminium alloy sensitive to shear fracture, a criterion based on the maximum shear stress, τ max , may in certain cases represent the fracture FLD with good accuracy based on a minimum number of calibration tests. In the present case, Fig. 11a and Tables 4 and5 show that a critical value of either σ B (as determined using the "R" model), or of the maximum shear stress, τ max , cannot represent the "experimental" forming limit diagram with an accuracy better than 10% and ±7%, respectively. Such a critical shear stress was mentioned, but not used by [Xu et al, 2012]. In a Fe-22Mn-0.6C-0.2V TWIP steel, a critical shear stress of 1015 MPa, determined from a variety of mechanical tests analysed using a "VM-H" model, allowed a prediction of the FLD with an accuracy better than about 20% [Bouaziz et al, 2011]. In spite of the ductile fracture mode, void growth is very limited even for high values of stress triaxiality ratio so that criteria such as those of [McClintock, 1968, Rice and Tracey, 1969, Oyane et al, 1980] are not relevant for this material. The candidate fracture criterion should be able to address a wide range of values for stress triaxiality ratio η, including those close to zero. To this aim, criteria developed for modelling either rock fracture or fracture of metal alloys very sensitive to shear fracture (even if less ductile than TWIP steels) were considered.

The stress state was expressed in terms of Haigh-Weestergaard (or Lode) coordinates in the space of maximum principal stresses as in [Fricker and Meckbach, 1999, Xue and Wierzbicki, 2008, Bai and Wierzbicki, 2010]. The first invariant B I 1 of the stress tensor calculated with the "R" model at fracture initiation, B σ , and the second and third invariants, B J 2 and B J 3 , of the modified deviator B dev σ , were first calculated. Superscript "B" denotes that unlike in e.g. [Bai and Wierzbicki, 2010], the fracture criterion was set by using the same "R" model (including the Barlat-type yield criterion) as that used to extract fracture initiation conditions from experimental data. Note that such expression "loses" the actual loading axes with respect to the principal directions of the sheet material. In other words, the forming limit diagram (FLD) cannot be uniquely expressed as a function of the invariants of stress tensors, as for an isotropic material. As a consequence, only data of tests involving loading along RD were considered with the exception of shear, for which principal strains are rotated with respect to the (RD, TD) frame. Results of shear tests will thus be only qualitatively interpreted here. The FLD determined in this study will be termed the "RD-FLD" to denote loading along RD. Great care must thus be taken when using this RD-FLD for more complex loading paths (or even loading along other directions) such as in forming conditions. Although being less accurate, a FLD determined using an isotropic (e.g. von Mises) yield criterion overcomes this difficulty.

For the above mentioned invariants, and for a von Mises yield criterion, Haigh-Weestergard coordinates are expressed as follows: For the von Mises yield criterion, the Lode parameter may be defined either using the stress tensor as in Eq. ( 11), or using the strain rate tensor With the yield criterion used in the "R" model, definitions of the Lode parameter based respectively on the stress and on the strain rate tensors are not equivalent. As the considered material is very sensitive to strain localisation, the strain-rate based definition was chosen to represent a strong local variation in strain rate (but not a discontinuity as in [Rudnicki et al, 1975, Needleman andRice, 1978]). θ ) close to 1 (exactly 1 for shear, ≈ 0.99 for plane strain tension). Thus, four tests were used to identify this fracture criterion. The FLD was then estimated using this criterion for a variety of loading conditions as reported in Fig. 11. Using such a function remarkably improves the prediction of fracture with respect to the criteria based on either σ B or τ max . However, Eq. ( 15) underestimates the resistance to fracture in shear and overestimates it for

Conclusions

An extensive experimental database has been built for the mechanical behaviour of a 22Mn-0.6C-0.2V TWIP steel. Several fracture criteria were identified on this basis. Main conclusions are as follows:

1. Anisotropic yield is satisfactorily represented with an anisotropic (Barlat-type) yield criterion for a variety of in-plane loading modes. 2. Taking the strong contribution of kinematic hardening into account is mandatory in the case of significantly non-radial loading path, but not useful in the case of monotonic tensile loading of smooth and notched specimens, for which a simple non-linear isotropic hardening equation provides very good estimates of macroscopic response and local strain fields. 3. Fracture occurs suddenly by nucleation and early coalescence of submicrometer-sized voids, although some damage develops as scarce, coarser voids well before fracture. Vanadium carbides as well as the occurrence of mechanical twinning are probably key parameters governing void nucleation. 4. The forming limit diagram was estimated using several fracture criteria together with the same formalism at that of the constitutive model. A simple function of hydrostatic stress, anisotropic equivalent stress and Lode angle (with a strain rate-based definition) satisfactorily represent fracture initiation conditions. 5. Such criterion relies on a particular loading direction (here, RD) and further work is required to develop a fracture initiation criterion based on the same formalism but more directly usable for the prediction of forming ability in more complex conditions such as deep drawing.

Acknowledgements

Figure 1 :

 1 Figure 1: Microstructure of the Fe-22Mn-0.6C-0.2V TWIP steel. (a) Grain morphology (light optical micrograph, electrolytic etch with a Struers® A2 solution), (b) segregation bands (scanning electron micrograph, secondary electron imaging, colloidal silica finish).

Figure 2 :

 2 Figure 2: Engineering stress-strain curves under uniaxial tension. (a) Effect of the loading direction; strain rate sensitivity (b) at room temperature, (c) at 400°C, and (d) at both (loading along RD).

Figure 3 :

 3 Figure 3: Room temperature shear test results. (a) Stress-strain curves; (b) respective contributions of kinematic (x 12 ) and isotropic hardening (r 12 ) to the shear stress. σ 12_0 is the initial yield stress in these

Figure 4 :

 4 Figure 4: Plastic yielding and fracture of DENT specimens. (a,b) Effects of loading direction and notch radius (curves have been artificially shifted by 1 mm for better readability); (c) image captured with the camera showing cracks initiated from both notch roots (NT0.5 pulled along RD at 0.5 mm.min -1 superimposed to corresponding major (axial) engineering strain field measured by DIC.

Figure 5 :

 5 Figure 5: Results of tests on DENS specimens (continuous lines) compared to predictions of the "R" model (dotted lines). Load and displacement of central point vs. notch opening curves of (a) ungrooved and (c) grooved specimens (same curve labelling for (a) and (c)). Meshing with boundary conditions at specimen ends (displacement along directions 1 and 2) and appearance of specimen at fracture are shown as inlets. In-plane shear strain at fracture in specimens loaded along RD: DIC measurements and model predictions for (b) ungrooved and (d) grooved specimens.

  Disp lacem ent of poin t C // direc tion 1 D is p la c e m e n t o f p o in t C // d ir e c ti o n 2 A p p li e d lo a d experimental displacement // 1 of poin t C // direc tion 1 D is p la c e m e n t o f p o in t C // d ir e c ti o n 2 A p p l i e d lo a d Groove (root at point C) Displacement of central point , C (mm) of poin t C // direc tion 1 D is p la c e m e n t o f p o in t C // d ir e c ti o of poin t C // direc tion 1 D is p la c e m e n t o f p o in t C // d ir e c ti o n 2 A p p l i e d lo a d Groove (root at point C) Displacement of central point , C (mm) of poin t C // direc tion 1 D is p la c e m e n t o f p o in t C // d ir e c ti o of poin t C // direc tion 1 D is p la c e m e n t o f p o in t C // d ir e c ti o n 2 A p p l i e d lo a d Groove (root at point C) Displacement of central point , C (mm) of poin t C // direc tion 1 D is p la c e m e n t o f p o in t C // d ir e c ti o of poin t C // direc tion 1 D is p la c e m e n t o f p o in t C // d ir e c ti o n 2 A p p l i e d lo a d Groove (root at point C) Groove (root at point C) Displacement of central point , C (mm)

Figure 6 :

 6 Figure 6: Results of tests on DSNT specimens. (a) Load vs. load line displacement curves; (b) Experimental and predicted effective stress vs. elongation curves (elongation averaged over 3.42 mm from DIC measurements and over 3.50 mm for model prediction), edge view of undeformed FE mesh and of fractured specimen; (c) major strain at fracture: DIC measurements and prediction of the "R" model (edge view).

Figure 7 :

 7 Figure 7: Fracture surfaces of specimens tested at room temperature with low loading rates. (a) Coarse cavities and finer dimples observed after UT along (a) RD and (b) DD directions. (c) Higher magnification view of smaller dimples with some tiny spherical particles indicated with black circles. (d) Tiny microcracks ahead of the crack tip in a cross-section of a Kahn specimen (test interrupted during stable crack propagation).

Figure 8 :

 8 Figure 8: Fracture surface of a UT specimen pulled along RD at 110 s -1 .

Figure 9 :

 9 Figure 9: Fracture surface after UT testing at 400°C along RD: coarser cavities and finer dimples.

Figure 11 :

 11 Figure 11: Experimental determination (symbols) and prediction (lines) of the FLD of the Fe-22Mn-0.6C-0.2V TWIP steel (main loading along RD). (a) Major vs. minor strains; (b) minor vs. major stresses; (c) equivalent strain vs. stress triaxiality; (d) equivalent stress vs. stress triaxiality. All quantities are evaluated using the "R" model. Same labelling for (a-d).

  "R" model, the definition of ξ, denoted here as B ξ was not changed but the values of ρ and Lode angle θ were modified to be consistent with the formalism of constitutive equations. The value of B ρ was taken proportional to the anisotropic measure of stress, σ B , and equal to the value of uniaxial fracture stress B RD UT // σ in the case of uniaxial tension (again, along RD):

  in the definition of the Lode parameter to ensure that the resulting value ranges between (-1) and (+1). Consequently, the Lode angle B approximation, inspired by very limited void growth in this material, the fracture criterion was tentatively described as a function of B ρ and B ε θ , i.e., the influence of hydrostatic pressure was neglected. The simplest possible function of these two variables that could reasonably account for experimental data was that of a semi-elliptical cylinder in the Haigh-Weestergaard coordinates: B a was fitted to 2200 MPa by averaging values obtained from uniaxial and equibiaxial tension tests with sin (3 B ε θ ) = 0. The value of B b was fitted to 1850 MPa, by averaging values obtained from shear and plane strain tests with sin(3 B

  ε

  Technical help from A. Meddour, Y. Auriac (MINES ParisTech) and G. Petitgand, S. Douchamps and B. Remy (ArcelorMittal) is gratefully acknowledged. Introduction of kinematical hardening parameters varying with cumulative plastic strain was suggested by J.L. Chaboche (MINES ParisTech).

Table 1 :

 1 Chemical composition of the investigated TWIP steel (wt%)

	Fe	Mn	C	V	S
	base	22	0.6	0.2	0.0008

Table 2 :

 2 Room temperature uniaxial tensile properties of the Fe-22Mn-0.6C-0.2V material. Values of strength for higher elongation rates are only indicative due to load oscillations (propagation of elastic waves)

	Loading	Elongation	Yield	Tensile	Fracture	Lankford	Logarithmic
	direction	rate (s -1 )	strength	strength	elongation	coefficient,	reduction of
			(MPa)	(MPa)	(%)	r	thickness at
							fracture
	RD	4.10 -4	442	1137	35; 58	0.87	0.37; 0.41
		4.10 -3	483	1101	42; 61		0.41; 0.43
		4.10 -2	525	1066	45; 46		0.43; 0.45
		50	(oscillations) ~1115	45		0.45
		110	(oscillations) ~1130	47		0.45
	TD	4.10 -4	514	1170	43; 41	1.40 down	(not
		4.10 -3	539	1162	43; 52	to 1.26	available)
		4.10 -2	554	1093	48; 50	(average
						1.30)
	DD	4.10 -4	515	1149	48; 56	0.91	(not
							available)
		4.10 -3	545	1114	45; 52
		4.10 -2	575	1086	59; 50

  , respectively.

	Specimen	Logarithmic strain along ND:	Strain ratio, α B	σ B (MPa)	η B	ρ B (MPa)	Lode angle, θ B	ε B	τ max (MPa)
		experiment					(rad)
		/ model
		prediction
	UT	-0.39	-0.47	2584	0.29	2277	0	0.65
	DENT05	-0.45 / -0.37 -0.48	2470	0.362	2176	0	0.62
					down
					to
					0.306
	DENT1	-0.45 / -0.37 -0.46	2494	0.330	2197	0	0.60
					down
					to
					0.300
	DENT2	-0.45 / -0.37 -0.45	2557	0.300	2253	0	0.57
					down
					to
					0.295
	DENS,	Fracture	-0.47		0
	ungrooved	initiation	(edge),
		point not	-1
		well	(centre
		localised	of gauge
	DENS,	part) -0.59 / -0.58 -1	2357	0	2077	π/6	0.52
	grooved
	DSNT	not	-0.18	2124	0.38	1871	0.41	0.41
		available / -
		0.37
	Center hole	-0.34	-0.50 VM	2443	0.29	2152	0	0.56
	tensile *
	Plane strain *	-0.39	0 VM	2032	0.42	1790	0.48	0.38
	Tilted plane	-0.30	-0.20 VM	1916	0.35	1688	0.32	0.33
	strain * Equibiaxial stretching * Shear * Marciniak 1 a	-0.75 ε 12 = 0.48 -0.41	1 VM -1 VM 0.08	2615 2221 2028	0.60 0 0.56	2304 1957 1787	π /3 π /6	0.66 0.46 0.38
			(exp.				0.61
			value)
	Marciniak 2 a	-0.64	0.68	2446	0.60	2155	0.95	0.57
			(exp.
			value)

Table 5 :

 5 Fracture initiation criteria parameters for loading along RD and accuracy of model predictions.

	Fracture criterion Parameters	Critical values of stress: σ B τ max	Function of σ B and Lode angle (Eq. 15) B a B b	Function of hydrostatic pressure, σ B and Lode angle (Eq. 16) B a' B b' h B ξ 0
		(MPa)	(MPa)	(MPa)	(MPa)			(MPa)
	Adjusted values ∆ calculated using tests involving	2276	1050	2200	1850	1.12	0.93 0.13	15515
	minor strain negative or positive	0.55	0.54		0.41		0.27
	minor strains at fracture initiation
	locus) ∆ calculated using tests involving
	negative minor strains at fracture	0.36	0.35		0.24		0.19
	initiation locus only

plane strain tension, by making a compromise between these two tests, for which

. Such a difference between these two experimental results could arise from (i) the direction of principal strains being different from the RD, TD, ND directions or (ii) a possible effect, yet rather weak, of hydrostatic pressure. The latter effect could be simply taken into account by introducing a weak dependence on ξ. This was done using a linear function, the fracture criterion becoming a cone, with apex at (ξ = ξ 0 , ρ = 0) and aperture angle tan -1 (h B ): 5.

The ratios between minor and major axes of the elliptic section of the two fracture criteria are close to each other (1.23 for Eq. ( 15) and 1.20 for Eq. ( 16)). Predictions of notch opening at fracture are nonconservative with the model of Eq. 15. The model of Eq. ( 16) yields conservative predictions of displacement at fracture for all tests except for ungrooved DENS specimen, for which the loading path is far from being radial in the gauge region. For all other specimens, fracture is predicted to initiate at the same integration point (i.e., same location in the specimen) by the two models. The difference between "experimental" values and model predictions, ∆, was quantified as follows:

In Eq. ( 17), for any test i and for any fracture initiation criterion crit, (namely, critical value of σ B , critical value of τ max , Eq. ( 15) or Eq. ( 16)), exp i σ and crit i σ are respectively the "experimental" value of (Barlat) equivalent stress at fracture and the corresponding model prediction using criterion crit.

Introducing the hydrostatic pressure into the fracture initiation criterion divides the value of ∆ by two (Table 5). Predictions are mainly improved in the left part of the RD-FLD, at the expense of biaxial stretching. Further work is necessary to also improve prediction in the right part of the RD-FLD. This point is not critical, however, for this material and targeted automotive application, in usual forming conditions. The model could further be improved by using a more sophisticated function of hydrostatic pressure. This is, however, beyond the scope of this study due to the low number of "experimental" results involving similar values of Lode angle and significantly different values of hydrostatic pressure. The model of Eq. 16 was thus considered to satisfactorily represent the RD-FLD for the considered material. Note that many tests (17) were used to identify the constitutive parameters (but a lower number could have been used, in particular in reverse shear). On the other hand, only four tests were used to identify the fracture criterion, some of them being the same as those used for identification of constitutive parameters.


Hosna Titah-Benbouzid

Mohamed Benbouzid
email: mohamed.benbouzid@univ-brest.fr

Marine Renewable Energy Converters and Biofouling: A Review on Impacts and Prevention

Keywords: Marine renewable energy converters, biofouling, corrosion, microorganisms, marine environment

In recent years, a number of studies have been performed to assess the damages caused by biofouling, which is simply the attachment of organisms to a surface in contact with water for a period of time. This explanation sounds fairly straightforward, but there are several organisms that cause biofouling, many different types of affected surfaces, and therefore many solutions dealing with this problem. Regarding the marine renewable energy emerging and promising area of research, this paper aims to provide a review of the biofouling issue in the context of Marine Renewable Energy Converters (MRECs). The proposed review will specifically highlight biofouling impacts on MRECs and solutions to prevent fouling. In addition, a discussion will highlight challenges that MRECs market needs to undertake to overcome the biofouling problem.

I. INTRODUCTION

Corrosion and biofouling have been impetus for investigating interactions between microorganisms and solid surfaces. This phenomenon is essential in a marine environment, where zooplankton is relatively diverse and abundant, but finds, although in smaller proportions, in freshwater environments. Marine biofouling can affect boat hulls, harbor constructions and piers as well as underwater engineering installations [START_REF] Rolland | Synthesis and characterization of perfluoropolyether graft terpolymers for biofouling applications[END_REF]. Fishing and fish farming are also affected, with mesh cages and trawls harboring fouling organisms. In Australia, biofouling accounts for about 80% of the pearling industry's costs [2]. Gold-and silicon-based components of micro-electrochemical drug delivery devices are susceptible to biofouling, as are machines in the papermaking and pulping industries and underwater instrumentation [START_REF] Anderson | Biocompatibility and biofouling of MEMS drug delivery devices[END_REF]. Minimally adhesive coatings are currently being investigated for use as fouling release coatings on marine equipment. The attachment of organisms to the ship hull can dramatically increase drag and therefore fuel consumption. A mechanically stable, non-toxic coating is desired to prevent the adhesion of foulants [START_REF] Kolari | Attachment mechanisms and properties of bacterial biofilms on non-living surfaces[END_REF]. Organism groups that contribute to marine biofouling include seaweeds, bivalves, crustaceans and barnacles. Marine biofouling can be divided into two groups: 1) Microfouling organisms and 2) Macrofouling organisms.

A. Microfouling Organisms

Biofouling process is not as simple as it seems because the agencies do not stick on the substrate as a sucker. Indeed, the complex process often begins with the production of a biofilm. It is a film made of bacteria, such as Thiobacilli or other microorganisms, which forms on a material when conditions are right [START_REF] Gehrke | Interactions between microorganisms and physicochemical factors cause MIC of steel pilings in harbours[END_REF]. These organisms are primarily bacterial and microbial in nature and quickly colonize any substrate placed in seawater. They form part of a sticky coating commonly referred to as a biofilm. Biofilms are a considerable nuisance, accumulating in thicknesses sufficient to obscure marine surfaces and adding considerably to the difficulties of operating subsea. They also provide both a food source and a convenient interface to which the larger organisms, the macrofoulers, can adhere.

B. Macrofouling Organisms

The growth of a biofilm can progress to a point where it provides a foundation for the growth of seaweed, barnacles and other organisms [START_REF]Biofouling formation and remedial measures[END_REF]. Macrofoulers cause additional and even more severe problems for subsea operators. This grouping includes many larger animals and plants that may attach as individuals or in large colonies, such as barnacles, mussels, polychaetes, and various species of bryozoans and hydroids: Fouling and abrasive suspended particles growth [7]. When the biofouling growth is in contact with the marine renewable devices, it will affect wave machines more than submerged tidal systems. However, even at 5 m submerged depths, the tidal systems and in particular the rotor hubs (lowvelocity parts) will be covered with fouling [START_REF] Hellio | Advances in Marine Antifouling Coatings and Technologies, Understanding the Biofouling of Offshore and Deep-Sea Structures[END_REF].

In this particular context, this paper aims to provide a state of the art review that highlight the biofouling phenomena; from its understanding to the prevention solutions with a focus on marine renewable energy converters.

II. BIOFOULING DEVELOPMENT MODES

Attachment and growth of living organisms on man-made immersed surfaces is a complex colonization process. In marine environment, any MREC could be colonized by bodies such as bacteria, diatoms, protozoans, algae and invertebrates [START_REF] Mineur | Changing coasts: Marine aliens and artificial structures[END_REF][START_REF] Polagye | Screening for biofouling and corrosion of tidal energy device materials[END_REF]. Figure 1 summarizes the four sequences of a typical fouling installation process: Biochemical conditioning, bacterial colonization, installation of single-celled species, and installations of multicellular species [START_REF] Wahl | Marine epibiosis. I. Fouling and antifouling: some basic aspects[END_REF]. Biochemical conditioning describes the adsorption of dissolved chemical compounds (mostly macromolecules) to any surface in the first moments after contact with natural seawater. This instant (hereafter simply called "immersion") may be the extrusion of a growing sea grass blade from its sheath, the appearance of a new crustacean carapace after moulting, the emergence of a fresh rock surface after breakage or the experimental immersion of a glass slide, etc. The concentration process of organic molecules at interfaces (solid/liquid, liquid/gas) is purely physical and "spontaneous" [START_REF] Baier | Early Events of Micro-Fouling of all Heat Transfer Equipment[END_REF]. Bacterial colonization of the area consists of one reversible approach phase (adsorption) and an irreversible fixing phase (adhesion). At the first step in the sequence of colonization (biochemical conditioning) bacterial adsorption is mainly governed by physical forces: Brownian motion, electrostatic interaction, gravity, Vander-Waal forces [START_REF] Dexter | Influence of substrate wettability on the formation of bacterial slime films on solid surfaces immersed in natural sea water[END_REF][START_REF] Dexter | Influence of substratum critical surface tension on bacterial adhesion -In situ studies[END_REF][START_REF] Fletcher | Influence of substratum characteristics on the attachment of a marine pseudomonad to solid surfaces[END_REF][START_REF] Walt | The effect of gravity on initial microbial adhesion[END_REF]. When the two bacteria cells from surface, which constitute the adsorbed macromolecular film, are mainly negatively charged [START_REF] Marshall | Mechanisms of adhesion of marine bacteria to surfaces[END_REF][START_REF] Corpe | Secretion of adhesive polymers and attachment of marine bacteria to surfaces[END_REF][START_REF] Wicken | Bacterial Cell Walls and Surfaces[END_REF], the opposing repulsion and attraction electric forces tend to immobilize the cell at a distance of 15 to 20nm [START_REF] Fletcher | Microbial attachment to nonbiological surfaces[END_REF]. This electrostatic barrier may be filled by the production of polysaccharide fibrils (mainly of glucose and fructose). The establishment of covalent bonds between the bacterial glycocalix and macromolecular film pushes the adsorption phase to blend in an adhesion phase. Similar mechanisms have been described for yeasts [START_REF] Douglas | Adhesion of pathogenic Candida species to host surfaces[END_REF], unicellular algae [START_REF] Tosteson | Aggregation-adhesion enhancing macromolecules and the specificity of marine microbial surface interactions[END_REF], spore [START_REF] Fletcher | Microbial attachment to nonbiological surfaces[END_REF], and polychaetes larvae (Fig. 2) [START_REF] Kirchman | Lectins may mediate in the settlement and metamorphosis of Janua (Dexiospira) brasiliensis Grube (Polychaeta: Spirorbidae)[END_REF]. The colonization by unicellular eukaryonts step includes the arrival of yeasts, protozoa and diatoms, with a clear quantitative dominance of the last one [START_REF] Marshall | Mechanisms of the initial events in the sorption of marine bacteria to surfaces[END_REF][START_REF] Corpe | Attachment of Manne Bacteria to Solid Surfaces[END_REF][START_REF] Cuba | The initial development of a marine fouling assemblage on a natural substrate in a subtropical estuary[END_REF]. Benthic diatoms attach by the secretion of mucus [START_REF] Cooksey | The Attachment of Microfouling Diatoms[END_REF][START_REF] Ferreira | The colonization process of algal epiphytes on Ruppia maritima L[END_REF] and densely cover the areas of broad substrate and contribute significantly to the chemical/biological evolution of the substrate. Subsequently, the colonizers protozoa feed of microorganisms (bacteria, yeasts, diatoms and other protozoa) [START_REF] Cuomo | Fungal flora of Posidonia oceanica and its ecological significance[END_REF]. One to several weeks after immersion the substratum bears a highly differentiated and 3D structured microbiotic community. Several days to weeks after biochemical conditioning the last and longest phase of colonization, which is, colonization by multicellular eukaryonts, begins with the settlement of meroplanktonic larvae and algal spores. This overlaps with the continued recruitment and evolution of the microepibiotic community [START_REF] Wahl | Marine epibiosis. I. Fouling and antifouling: some basic aspects[END_REF]. After these 4 successive stages, the community soiling evolves continuously by mechanisms such as disturbance, facilitation, inhibition, tolerance, etc. [START_REF] Connell | Mechanisms of succession in natural communities and their role in community stability and organization[END_REF][START_REF] Railkin | Marine Biofouling: Colonization Processes and Defenses[END_REF].

In summary, most of the mechanisms may be -and frequently are -combined to form a multifactorial antifouling adaptation, which effectively covers the range of potential colonists (Fig. 1).

III. MARINE RENEWABLE ENERGY CONVERTERS PERFORMANCE ALTERATION

Biofouling proliferation concerns most artificial constructions in marine environment consisting in non-natural materials, such as treated wood, metal, glass, rubber, rigid plastic, concrete, or fiberglass. It is defined as unwanted growth on materials; in contrast, development and succession of befouling in some artificial reefs is wanted for marine fauna and flora.

In our context, it has been shown an increase in species abundance with increasing structural volume and complexity of artificial reefs [START_REF] Potts | Structural influences of artificial and natural habitats on fish aggregations in Onslow Bay, North Carolina[END_REF][START_REF] Spieler | Artificial substrate and coral reef restoration: What do we need to know to know what we need?[END_REF]. For example, fouling on artificial structures can cause large economic costs by impairing equipment performance or life span [START_REF] Langhamer | Artificial reef effect and fouling impacts on offshore wave power foundations and buoys -A pilot study Estuarine[END_REF].

Biofouling can easily cause obstructions in MRECs and/or increase the weight and drag, thus significantly affecting the device performance (Fig. 3) [START_REF] Jouanne | Wave energy industry research and ocean testing solutions to advance the wave energy industry[END_REF].

Wave power buoys floating on the surface may be heavily overgrown by epibiotic assemblages and this may, literally, become a technical burden. The dynamics and thus the buoys ability to extract energy from an ocean wave is determined by the size and shape of the buoy, the mass of the moving parts, together with the power take-off system [START_REF] Eriksson | Hydrodynamic modelling of a direct drive wave energy converter[END_REF]. Biofouling will change the mass of the buoy and the flow of water around the buoy. It is therefore important to investigate how it will affect the energy absorption of a wave energy converter, and ideally, wave power buoys should be formed in a way that fouling has a negligible impact on performance [START_REF] Langhamer | Artificial reef effect and fouling impacts on offshore wave power foundations and buoys -A pilot study Estuarine[END_REF]. Fouling may negatively affect buoy buoyancy and dynamics but also include the upper parts of the rope. The extra weight in temperate waters, mainly caused by mussels, e.g. Blue mussels (Bivalvia) and barnacles (Cirripedia) may be more than 10 kg/m 2 thus adding extra weight (Fig. 4), whereas algae more likely only will affect buoy dynamics and not its weight [START_REF] Leijon | An electrical approach to wave energy conversion[END_REF][START_REF]Regional Coastal Monitoring Programmes[END_REF][START_REF] Langhamer | Colonization of Blue Mussels (Mytilus edulis) on Offshore Wave Power Installations[END_REF]. Two potential performance issues for marine current turbines are the roughening of the turbine blades due to impact, cavitation or scour due to particulates, and the fouling of the turbine blades by marine growth [START_REF] Fraenkel | Power from marine currents[END_REF][START_REF] Ng | 2002-2012: 10 years of research progress in horizontal-axis marine current turbines[END_REF]. In this context, there is a clear need for high reliability given the difficult maintenance access issues in the underwater environment.

There have been few studies to investigate roughness effects or fouling on marine current turbines. In [START_REF] Orme | Investigation of the effect of biofouling on the efficiency of marine current turbines[END_REF], potential effects of barnacles have been investigated. The lift and drag coefficients for an aerofoil covered with idealized barnacles of different sizes and distribution densities were determined using a wind tunnel. The lift to drag ratio decreased with both increasing barnacle size and distribution density. It was concluded that the presence of barnacles would have a detrimental effect on turbine efficiency. In [START_REF] Batten | The prediction of the hydro-dynamic performance of marine current turbines[END_REF], the potential effects of an increase in blade roughness or blade fouling using a numerical model have been investigated. It has been assumed that the presence of roughness or fouling would increase the drag coefficient by up to 50% [START_REF] Walker | Experimental and numerical studies of blade roughness and fouling on marine current turbine performance[END_REF]. Biofouling has been shown to reduce turbine blades efficiency and therefore decreasing the overall power generation [START_REF] Orme | Investigation of the effect of biofouling on the efficiency of marine current turbines[END_REF]. Similarly, biofouling may decrease the power conversion rate for wave energy converters through added inertia [START_REF] Eriksson | Hydrodynamic modelling of a direct drive wave energy converter[END_REF]. Furthermore, biofouling may damage protective coating and interfere with sensitive areas necessary for monitoring and maintenance [45]. Fig. 4. Heavily fouled directional waverider buoy [START_REF]Regional Coastal Monitoring Programmes[END_REF].

IV. BIOFOULING PREVENTION

In the case of marine renewable energy equipment, the adverse effects caused by the biological settlement (biofouling) are well known: High frictional resistance, due to generated roughness, which leads to an increase of weight and subsequent potential speed reduction and loss of maneuverability. To compensate for this, higher consumption is needed, an increase of the frequency of dry-docking operations, i.e. time is lost and resources are wasted when remedial measures are applied. A large amount of toxic wastes is also generated during this process [START_REF] Abbott | Cost-benefit analysis of the use of TBT: The case for a treatment approach[END_REF][START_REF] Rouhi | The squeeze of tributyltin[END_REF], deterioration of the coating so that corrosion, discoloration, and alteration of the electrical conductivity of the material are favored [START_REF] Cooney | Copper oxide impregnated textiles with potent biocidal activities[END_REF], introduction of species into environments where they were not naturally present (invasive or non-native species) [START_REF] Brancato | Impacts of invasive species introduced through the shipping industry[END_REF][START_REF] Reise | Introduced marine species of the North Sea coast[END_REF].

Since antiquity and until the beginning of the era of the great marine sail, large ships were severely disabled by Lepadomorpha and other marine organisms. To combat them, Phoenicians, Egyptians and Romans used several techniques: dubbing of the hull by a second shell in wood, copper or lead [51]. In addition to tars, first known recipes paintings protective and biocides were invented by the shipyards of a former navy sailing: Lime-based paints have initially been utilized [52]. At Christophe Colomb time, waxy coatings were used (mixture of pitch, tallow or other grease with the beeswax) [52]. Arsenic was then widely used in inorganic form [54]. Mercury and organomercury (long known as agricultural pesticides) have also been used by sailing ships [54].

From the beginning of the 18 th century, it has been started using copper sheets nailed on the submerged parts of the shell of certain ships, which had the same effect more sustainably but more expensive. Cinnabar red and white lead (very toxic) could be added between the copper and wood, to effectively combat shipworms [52]. Tributyltin has been very used between 1960 and 1990 (almost all boats wore it during 1970 years [55]). Too toxic, it leaves place to copper salts (primarily Cu20). However the Cu20 is easily bioaccumulated [START_REF] Olsen | Inorganic precursor peroxides for antifouling coating[END_REF]. Since 1990, diuron and irgarol are among the first alternatives to organotin compounds, but other alternatives were tested, including for example peroxides, which reacts with seawater to create hydrogen peroxide and very soluble metal ions in the water. Peroxides of strontium, of calcium, of magnesium, of zinc were therefore tested but they were found much less ecologically toxic, but often less effective, more expensive or less durable [START_REF] Olsen | Inorganic precursor peroxides for antifouling coating[END_REF][START_REF] Manzo | Toxic effects of irgarol and diuron on sea urchin Paracentrotus lividus early development, fertilization, and offspring quality[END_REF]. At this time, studies conducted in the USA on different types of vessels show that there are alternatives economically and environmentally acceptable, including for submarines or very large ships such as aircraft carriers, for example with the organometallic copolymers [58].

Among all the different solutions proposed throughout navigation history, tributyltin self-polishing copolymer paints (TBT-SPC paints) have been the most successful in combating biofouling on ships. The widespread use of these paints, estimated to cover 70% of the present world fleet in [START_REF] Gerigk | The present status of TBT copolymer antifouling paints versus TBT-free technology[END_REF][START_REF] Champ | The Status of the treaty to ban TBT in marine antifouling paints and alternatives[END_REF], has led to important economic benefits [START_REF] Abbott | Cost-benefit analysis of the use of TBT: The case for a treatment approach[END_REF], [START_REF] Champ | A review of organotin regulatory strategies, pending actions, related costs and benefits[END_REF].

Unfortunately, TBT-SPC paints affect the environment. In this context, numerous anti-biofouling measures such as mechanical, chemical, and biological methods are in practice but their effects on the biofouling are not remarkable. In addition, the commercially available antifouling paints such as tributyltin (TBT) and copper sulphate are highly toxic to the non-target aquatic organisms [START_REF] Anderson | TBT-free anti-fouling coatings in 2003: For better or for worse?[END_REF]. The TBT contributes to the development of antimicrobial tolerance and impose pseudo hermaphroditism in marine invertebrates [START_REF] Smith | Male characteristics on female mud snails caused by antifouling bottom paints[END_REF].

In 1990, the International Maritime Organization (IMO) adopted a resolution recommending governments to adopt measures to eliminate anti-fouling paints containing TBT. In October 2001, IMO adopted a new international convention on the control of harmful anti-fouling systems on ships, which will prohibit the use of harmful organotins in anti-fouling paints used on ships and will establish a mechanism to prevent the potential future use of other harmful substances in antifouling systems [64]. After that, it was established that a good biocide for use in an anti-fouling system has the following characteristics: Broad spectrum activity, low mammalian toxicity, low water solubility, no bioaccumulation in the food chain, not persistent in the environment, compatible with paint raw materials, favorable price/performance [START_REF] Tong | Tributyltin distribution in the coastal environment of peninsular Malaysia[END_REF].

Recently used anti-fouling techniques include methods based on the combined activity of copper ions and a booster biocide (e.g. diuron). However, these technologies have an effectiveness duration estimated between 2.5 and 5 years, which require shroud and cover at regular intervals [64]. In this context, electrochemical foul prevention was undertaken. Indeed, it is perceived to have lower environmental impacts when compared to biocide release coatings and have a muchincreased operational life spans resulting from no consumption of coating based chemicals. Experimental validation of such a biofouling prevention approach has been undertaken at the Hatfield Marine Science Center and demonstrated minimal electrochemical parameter degradation. In addition, no algae growth was observed over the course of the testing time. This performance was identical to that of fresh, un-aged biocidal coatings (Fig. 5) [START_REF] Jouanne | Wave energy industry research and ocean testing solutions to advance the wave energy industry[END_REF].

All the above-cited anti-fouling new technologies are however still in test and there is not yet enough feedback to ascertain their environmental impact.

The main anti-fouling approaches are summarized in Table I.

V. CASE STUDY FROM OREGON STATE UNIVERSITY

In the context of MRECs, a critical issue is to evaluate biofouling impact on their different components/materials. Indeed, marine renewable energy devices are composed of several materials and each of them will have a different impact (Fig. 4).

For that purpose, the Northwest National Marine Renewable Energy Center of the Oregon State University has conducted a study from April 2009 to February 2010. In this context, several materials of a typical tidal energy installation have been immersed. In fact, material coupons, which might be used in the rotor, drive train, or foundation were deployed in the seabed. Electrochemical coating with no power applied; Three fresh; Un-aged conventional biocidal coatings [START_REF] Jouanne | Wave energy industry research and ocean testing solutions to advance the wave energy industry[END_REF].

The test matrices consist of uncoated and coated coupons, nominally 6.35cm wide, 5.72 long, and 0.32cm thick. For expediency, coupons are attached to a fiberglass plate using a marine grade adhesive. Each plate is secured to the leg of an instrumentation tripod, as shown in Figure 6. The primary purpose of this tripod is to characterize the physical and biological environments at tidal energy sites [START_REF] Polagye | Screening for biofouling and corrosion of tidal energy device materials[END_REF].

A. Composite 1) Glass fiber composite. This composite material could be used for a MREC rotor, hub, or duct/shroud. In general, glass fiber composite performed well, with limited surface fouling after 10 months of deployment. Biocide release not constant, poor selfsmoothing, little activity during idle times, higher costs due to necessity of sealer coat on recoats [START_REF] Yebra | Antifouling technologypast, present and future steps towards efficient and environmentally friendly antifouling coatings[END_REF]. Slow drying time [START_REF] Terlizzi | Environmental impact of antifouling technologies: State of the art and perspectives[END_REF].

In-water cleaning difficult as brushes may damage silicone, foul release coatings are prone to abrasion damage [START_REF] Lewis | Proceedings: National Shipping Industry Conference[END_REF] A barnacle adhered to the edge of one material coupon after 5 months (Fig. 7). In this case, it has been shown a general trend that biological fouling is more common on edges and in crevices than on smooth surfaces.

2) Carbon fiber composite. This composite material could also be used for a MREC rotor, hub, or duct/shroud. As with glass fiber composite, the carbon fiber composite developed minimal surface fouling after up to 10 months of deployment.

B. Aluminum

This material could be used for a MREC rotor. In this case, more than 90% of the exposed surfaces on all aluminum coupons were oxidized during each 3-4 month deployment, as shown in Fig. 8.

After 10 months deployment in marine environment, the surface is almost entirely oxidized and embrittled to the point that one corner of the coupon broke away during routine handling.

C. Stainless Steel

This material could be used for a MREC hub or shroud. In this case, the tested sample developed superficial corrosion along the contact surface between the stainless steel and the fiberglass panel, and a barnacle attached to one edge (Fig. 9). Other stainless steel hardware used to secure instrumentation to the measurement tripod developed more severe corrosion during some deployments. After three months of immersion, the bracket has been oxidized to a failure point (Fig. 10).

D. Steel

1) Common steel. This material is commonly used for a MREC support structure. In addition to bare steel, protection by a zinc anode screwed into the center of the coupon is tested. The surfaces of both coupons were almost entirely oxidized upon retrieval. Anodic protection reduced but did not eliminate oxidation (Fig. 11).

2) Structural steel. As with common steel, in addition to a bare surface, protection by a zinc anode screwed into the center of the coupon was also tested. The protection significantly reduced but did not entirely eliminate oxidation (Fig. 12).

E. Bearing Materials

Three types of potential bearing materials were tested in terms of operational biological fouling. High-density polyethylene experienced minimal surface and edge fouling. No fouling was visible on either fiber reinforced phenolic resin or low friction liner on stainless backing [START_REF] Polagye | Screening for biofouling and corrosion of tidal energy device materials[END_REF].

VI. CONCLUSION AND PERSPECTIVES

This paper has proposed a state of the art review that highlight the biofouling phenomena; from its understanding to the prevention solutions. A focus on marine renewable energy converters has been attempted even if the marine renewable energy conversion market is a slow-developing one with no effective tidal or wave installed farms [START_REF] Titah-Benbouzid | Ocean wave energy extraction: Up-to-date technologies review and evaluation[END_REF][START_REF] Benelghali | Marine tidal current electric power generation technology: State of the art and current status[END_REF]. In this context, it has been shown that there are several techniques for biofouling prevention/control that are well accepted as they are most least environmentally.

Research and promising assessments are still ongoing. In particular, tests are still ongoing to determine an effective and durable antifouling respecting the environment. Nevertheless, they remain significant aspects that should be taken into account in these tests. In particular, the numerous interactions that happen during the antifouling use. The environment, the engineered coating, and the substrate mainly cause these. Figure 13 summarizes the different interactions that must be taken into account.

Fig. 1 .

 1 Fig. 1. Highly schematized colonizing sequence leading to the establishment of a fouling community [9].

B:

  Brownian motion; C: bacterial cell; DL: electrostatic double layer; DS: downsweep (microturbulence); E: electrostatic repulsion; F: bacterial fibrils anchored to adsorbed macromolecules; G: gravity (for a horizontal surface); M: bacterial motility; OF: organic film of adsorbed macromolecules; S: currents and macroturbulence; SUB: substratum; V: viscous boundary layer; W: Van-der-Waal's forces; WA: water body.

Fig. 2 .

 2 Fig. 2. Bacterial adsorption: dominating forces as a function of distance from the previously adsorbed macromolecular layer [11].

Fig. 3 .

 3 Fig. 3. Colonization by mussels on protective structures of the Horn Rev offshore wind farm in Denmark.

Fig. 5 .

 5 Fig. 5. From top to bottom: Aged electrochemical coating with applied power;Electrochemical coating with no power applied; Three fresh; Un-aged conventional biocidal coatings[START_REF] Jouanne | Wave energy industry research and ocean testing solutions to advance the wave energy industry[END_REF].

Fig. 6 .

 6 Fig.6. Material samples attached to instrumentation tripod[START_REF] Polagye | Screening for biofouling and corrosion of tidal energy device materials[END_REF].

Fig. 7 .

 7 Fig. 7. Barnacle attached to edge of glass fiber composite sample [10].

Fig. 8 .

 8 Fig. 8. Aluminum coupon corrosion [10].

Fig. 9 .

 9 Fig. 9. Stainless steel edge corrosion and biofouling [10].

Fig. 10 .

 10 Fig. 10. Stainless steel bracket corrosion [10].

Fig. 11 .

 11 Fig. 11. Common steel coupon corrosion [10].

Fig. 12 .

 12 Fig. 12. Structural steel coupon corrosion [10].

TABLE I .

 I COMPARISON OF KEY ANTI-FOULING SYSTEMS[START_REF] Chambers | Modern approaches to marine antifouling coatings[END_REF].

	Antifouling System	Leaching Rate	Life	Erosion Rate	Cost (US $/m 2 )	Problems
			Time		[67]
	(TBT) self-polishing	Chemical reaction	4 to 5 years	< 3µm/month [68].	$680,884	Banned 2008 [69]
	copolymer paints	through hydrolysis.	[54]
		Reaction zone of
		ablation 5µm deep
	(Tin-free) self-	Chemical reaction	5 years	N/A	$1,382,670	Lifetime shorter then
	polishing copolymers	through hydrolysis of				TBT-based paint
		copper, zinc, and silyl				systems; therefore
		acrylate				increasing the ship
						maintenance overall
						cost
	(Tin-free) conventional	10 µg/cm 2 d [70]	12 to 18	N/A	N/A	Hard non-polishing
	paint		months			performance leads to
						coating build-up.
						Performance only
						suitable for low fouling
						environments [67].
	Control depletion	Physical dissolution,	3 years	Matrix erodes due	$1,357,786
	polymers: copper paint	soluble matrix.		to dissolution of
				coating binder.
	Foul release	Low energy surface,	2 to 5 years	N/A	N/A
		leached silicone oils
		possible use [72]