Adrien Wattripont

François Baudin
email: francois.baudin@sorbonne-universite.fr

Marc De Rafelis

Jean-François Deconinck

Specifications for carbonate content quantification in recent marine sediments using Rock-Eval pyrolysis

Keywords:

 noticed that some of the characteristics of recent sediments tended to produce different results from those of more classical Rock-Eval analyses, causing bias in interpretations. In this study, the thermal stability of fossil and recent marine carbonated sediments was analyzed to identify differences in carbonate decomposition and to underline the importance of accounting for them in Rock-Eval analyses. The state of calcite preservation in recent marine sediments and sedimentary rocks at temperatures between 400 °C and 600 °C was characterized using different techniques (calcimetry, XRD, SEM imaging, etc.) for better interpretation of data obtained with Rock-Eval. Our results highlight a clear difference in the range of calcite decomposition temperatures during Rock-Eval analysis: between 550 °C and 775 °C for bulk clayey hemipelagic sediments versus 650 °C to 840 °C after the same sediments were rinsed to get rid of the salt. During heating, water and hydroxide anions are released from clay minerals and react with salt crystals to form acid vapor. This acid vapor reacts with carbonates to produce CO2. The chemical decomposition of carbonate starts at temperatures that are lower than the typical range of decomposition temperatures, leading to overestimation of mineral carbon content (overestimation of the S5 peak) and to underestimation of organic carbon content (underestimation of the S4CO2 peak) with the Rock-Eval method. In the absence of clay minerals, such as in recent marine pure carbonate oozes, there is no evidence for this effect. It is therefore essential to prepare and rinse recent clay-rich carbonated sediment samples before Rock-Eval analysis to avoid misinterpretation.

 [21,22,23,24,25,26,27 and others]. In 2014, a qualitative and quantitative carbonate recognition method based on the results of Rock-Eval pyrolysis was proposed [START_REF] Pillot | Identification and Quantification of Carbonate Species Using Rock-Eval Pyrolysis[END_REF]. According to this method the morphology of the CO2 production curve obtained during thermal destabilization of carbonated minerals also allows the identification of the type of carbonates present in the rock. However, in a previous study [START_REF] Baudin | Guidelines for Rock-Eval analysis of recent marine sediments[END_REF] we demonstrated that for Rock-Eval analyses of some recent marine sediments are biased, which can lead to misinterpretation of the results (organic matter and / or carbonate content). The aim of the present study was thus to identify the cause of the differences in recent marine sediments. The stability of recent carbonate natural samples subjected to heat treatment (such as those used in Rock-Eval pyrolysis) was compared to the same sample set but previously desalted. In parallel, different sample characterization techniques (calcimetry, XRD, scanning electron microscope imaging, etc.) were carried out to insure a better interpretation of the data obtained by the Rock-Eval and to propose specific conditions for analysis and/or the preparation of recent sediments.

Materials and methods 1.1 Material

As the purpose of this study was to understand the thermal behavior of carbonates in recent sediments, we chose samples of different sedimentary rocks with similar carbonate composition (content, mineralogy) to better identify the origin of the phenomenon. We chose consolidated sedimentary rock samples, recent clay-rich sediment samples and pure carbonate ooze samples. The following samples were analyzed: Douvrend Douvrend corresponds to the last fifteen centimeters of the upper Cenomanian "Craie de Rouen" Formation (Marnière de Douvrend section, Renouval Quarry, Dieppe region, France). It is a white to light beige marly chalk mainly made of coccoliths and foraminifera. It contains very little organic matter (TOC<0.3%) and is characterized by approximately 90.42±2% CaCO3 (calcite) content [START_REF] Ferré | Incidence des événements anoxiques océaniques sur les microfaunes Cénomano-Turoniennes du bassin Anglo-Parisien[END_REF]. The sample is in the form of a homogeneous powder and used as ISTeP laboratory internal standard for stable carbon and oxygen isotope measurements and trace element analyses.

MD01-2376 core

The MD01-2376 core was taken in the Timor Sea (12.314°S, 121.192°E) at a depth of 2906 m below the surface during the IMAGES VII MD122 (WEPAMA) cruise in May 2001 and penetrated the sediment to a depth of 27.90 m. Based on visual examination and smear slide observations the sediment is composed of a siliceous (10-30% diatomaceous) and carbonated (20-45% coccoliths and foraminifera) biogenic phases, clay minerals (25-45%) and organic matter (~ 1%). The color of the sediment (sometimes bioturbed) varies from light to olive gray [START_REF] Liu | Comparison of 240ka long organic carbon and carbonate records along a depth transect in the Timor Sea: Primary signals versus preservation changes[END_REF]. Since its recovery and opening, the MD01-2376 core has been stored in the cold and humid core repository of the French National Museum of Natural History in Paris. Around 5 cc of wet sediment were collected at approximately 10 cm intervals along the MD01-2376 core from its top down to 27.90 m (260 samples). Three larger samples (15 cc corresponding to a 2 cm thick interval) were collected at different depths of the core: 464-466 cm, 1914-1916 cm and 2345-2347 cm, in order to dispose of a sufficient quantity of material for additional analyses (see below). In the present paper, samples from this core are hereafter referred to as "bulk clayey hemipelagic ooze" and "rinsed clayey hemipelagic ooze", respectively. SU90-08 core These samples were collected from the top of the core (0.2 first meters) extracted with the Suroît (Ifremer) research vessel in 1990 during the PALEOCINAT drilling campaign at a depth of 3080 m below the surface north of the Azores (43.353N, -30.408W), in the Atlantic Ocean [START_REF] Vidal | Evidence for changes in the North Atlantic Deep Water linked to meltwater surges during the Heinrich events[END_REF][START_REF] Candelier | Temperature dependence of oxygen isotope fractionation in coccolith calcite: A culture and core top calibration of the genus Calcidiscus[END_REF]. The pelagic sediment in this zone is characterized by approximately 90% calcite content, low TOC (<0.2%) and a very small proportion of clay minerals. Samples from this core are hereafter referred to as "nannofossil pelagic ooze".

Methods

The first batch of sediment was dried at 50 °C in a stove and then crushed with an agate mortar. The second batch was rinsed twice in deionized water to remove residual salt from marine pore water. Conductivity was measured to insure the rinses were effective. The samples were then centrifuged and finally dried in a stove at 50 °C for 10 hours. We believe that this rinsing has no influence on the carbonate content because the procedure is fast (2 minutes maximum before centrifugation) and made with a neutral pH water.

Off-line thermal decomposition

Both bulk and rinsed samples were subjected to the following thermal protocol:

-11 mid-ramp isothermal sections in increments of 20 °C (from 400 °C to 600 °C) at 30 min intervals; -Each powdered sample was divided into 11 aliquots each weighing 0.6 g; -The aliquots were placed together in a muffle furnace at ambient temperature and ambient air atmosphere; -The furnace was then heated to 400 °C; -Then, for each isothermal section, the temperature was stabilized for 30 minutes; -At the end of each isothermal section, an aliquot was removed from the oven, cooled and weighed. This protocol was repeated for each sample, and a total of 60 aliquots (including raw samples) were prepared and analyzed.

Calcimetry

The measuring instrument used in this study was a Mélières manocalcimeter. The calcium carbonate dosage is based on acid attack (here with 30% volume HCl) of 100 mg of sample at ambient temperature. The calcimeter was previously calibrated with pure calcite. The volume of CO2 released during the following reaction was measured with a manometer: CaCO3 +2 HCl -> CaCl2 + H2O + CO2 XRF Ca counts High-resolution geochemical elemental composition records were measured on the archive slab core MD01-2376 surface with an Avaatech XRF Core Scanner at the EPOC Laboratory (University of Bordeaux). After cleaning and preparation of the surface, three separate runs (10, 30 and 50 kV, with currents of 0.6, 1.5 and 2.0 mA, respectively) were performed every cm down-core with a down-core slit size of 1cm. Sampling time was set to 10 s at 10 kV, 15 s at 30 kV and 20 s at 50 kV. The slab surface was covered with a 4mm thick SPEXCerti Prep Ultralene1 foil to avoid contamination of the XRF measurement unit and desiccation of the sediment. Raw data spectra were processed by analysis of the X-ray spectra using the iterative least squares software (WIN AXIL) package from Canberra Eurisys. The elements analyzed included a broad range, from aluminum through uranium [START_REF] Richter | The Avaatech XRF Core Scanner: Technical description and applications to NE Atlantic sediments[END_REF][START_REF] Tjallingii | Influence of the water content on X-ray fluorescence core-scanning measurements in soft marine sediments[END_REF]. In this paper, we only focus on calcium (Ca) results as a marker for carbonates assuming the absence of other calcium rich detrital minerals.

X-ray diffraction

In this study, XRD was used to monitor the amount of calcite present in the samples at each temperature step. All the samples were analyzed after thermal treatment (except for raw samples). Diffractograms provided information on the mineral nature of each sample. Once calcite was identified, intensity, width at half height and (1, 0, 4) peak surface were measured with MacDiff software. The diffractometer used in this study was a Bruker D2 Phaser equipped with a copper anticathode (CuKα radiation). Analyses were carried out according to the so-called "powders" method established by Moore & Reynolds [START_REF] Moore | X-Ray Diffraction and the Identification and Analysis of Clay Minerals[END_REF]. Aliquots of the three samples selected from core MD01-2376 (for both bulk and rinsed clayey hemipelagic oozes) were crushed, and part of the powder (grain size less than 50 µm) was added to deionized water to form mineral suspensions, that were then dispersed ultrasonically for 2 minutes. The fraction of the sediments less than 2µm in size was collected by sedimentation, and clay minerals were characterized from the XRD patterns of the oriented preparations acquired in air-dried conditions, after ethylene glycol solvation and heating at 490 °C for 2 hours. XRD analyses were performed using a Bruker D4 Endeavor X-Ray diffractometer (CuKα radiation). The oriented preparation of the fraction less than 2 µm in size was analyzed from 2° to 30° 2θ for the detailed characterization of the d00l diffractions of phyllosilicates.

Rock-Eval

We used a Turbo Rock-Eval 6 [START_REF] Lafargue | Rock-Eval 6 applications in hydrocarbon exploration, production, and soil contamination studies[END_REF]. Each individual sample was analyzed according to the Rock-Eval "Basic" cycle in which they are subjected to open pyrolysis at 300 °C to 650 °C with a 25 °C/min ramp and then to oxidation at 300 °C to 850 °C with a 20 °C/min ramp [START_REF] Espitalié | La pyrolyse Rock-Eval et ses applications[END_REF][START_REF] Behar | Rock Eval 6 Technology: Performance and developments[END_REF]. Core samples were also analyzed with the method of Baudin et al. [START_REF] Baudin | Guidelines for Rock-Eval analysis of recent marine sediments[END_REF] using a pyrolysis that starts at 180 °C and a ramp of 30 °C/min. Total organic carbon (TOC) was calculated from the hydrocarbons expelled from the sample during pyrolysis as well as from the amounts of CO and CO2 emitted during the pyrolysis and oxidation phases. The amount of mineral carbon (MinC) was calculated with Rock-Eval 6 as the sum of CO2 released during pyrolysis, at more than 400 °C, and the CO2 released by carbonate decomposition during oxidation, between 650 °C and 850 °C (for further explanations, see [START_REF] Behar | Rock Eval 6 Technology: Performance and developments[END_REF]).

Scanning Electron Microscopy

In this study, we observed only core samples using a Zeiss Ultra 55 Field Emission Scanning Electron Microscope. Three samples corresponding to different isothermal sections for both bulk and rinsed clayey hemipelagic oozes were analyzed by SEM to characterize calcite preservation state during heating at 400 °C, 480 °C and 540 °C (isothermal sections were chosen after XRD analysis to insure we focused on the temperature range of interest).

Results and interpretations 2.1 Calcimetry

The results of calcimetry during heating showed a quasi-steady state of the CaCO3 content of the Douvrend chalk sample in the temperature range studied (loss < 5% at the final temperature) whereas the bulk clayey hemipelagic ooze samples (3 samples taken at different depths of the core) showed a marked decrease in %CaCO3 starting at around 440 °C to reach almost 100% loss at around 600 °C (Fig. 1). However, the decrease in the rinsed clayey hemipelagic ooze samples was much less pronounced, started later (>520-540 °C) and reached a maximum of 30% loss at 600 °C. It is thus clear that rinsing the sediment influences the decomposition of CaCO3 content with increasing temperature. Fig. 1 -Variations in CaCO3 content (determined using manocalcimetry) in a Cretaceous chalk indurated sample (Douvrend) and in bulk and rinsed recent clayey hemipelagic oozes from the Timor Sea, by increasing the temperature from 400 °C to 600 °C in 20 °C increments in a muffle furnace. Three samples corresponding to different core depths (triangles, squares and dots) were used for the recent clayey hemipelagic oozes from the Timor Sea.

X-ray diffraction

Considering the observations made with calcimetry, we thought it would be interesting to analyze the behavior of calcite and sodium chloride (NaCl), the only mineral present in the bulk hemipelagic ooze and absent in the rinsed hemipelagic ooze. The major peaks of calcite (104) and halite (200) of the bulk clayey hemipelagic ooze samples tumbled completely and synchronously between 400 °C and 600 °C (Fig. 2) whereas the intensity of the quartz peak (101) remained steady throughout the heating cycle. These diffractograms highlight the obvious relationship between the presence of salt in the samples and the early decomposition of calcite (at low temperatures) even if they do not make it possible to accurately measure the quantities. The diffractograms of the rinsed clayey hemipelagic ooze showed a slight decrease in the intensity of the calcite peak between 400 °C and 500 °C and a more significant decrease between 500 °C and 600 °C but never above 30% loss at these temperatures. Thus, when salt was removed by successive rinses and the same temperature ranges were used, the intensity of the calcite peak only decreased by half when it tumbled down to zero in the presence of salt. Fig. 2 -Comparison of X-ray diffractograms of a bulk and rinsed clayey hemipelagic ooze from the Timor Sea at three stages of a heating experiment. Note the sharp decrease in the calcite peak (104) and in the halite peak (200) between 400 °C and 500 °C and its disappearence at 600 °C in the bulk sample. COLOR FIGURE. The main components of the clay fraction of the hemipelagic ooze were of R0 type illite/smectite mixed-layers hereafter referred to as smectite (about 70%), illite (15%), chlorite (5-10%) and kaolinite (10%) (Fig. 3). Traces of R1 type illite-smectite mixed layers and chlorite-smectite mixed-layer were probably also present. Quartz and traces of feldspars equally occur in the clay fraction. Fig. 3 X-ray diffractograms of the clay fraction of hemipelagic oozes samples in core MD01-2376.

Scanning Electron Microscopy

SEM observations confirmed that the carbonate fraction of clayey hemipelagic oozes was mainly composed of coccoliths and foraminifera, and subordinate micrometric calcite crystals of indeterminate origin. The coccoliths observed belong to the species Gephyrocapsa oceanica, characterized by a crossbar in the center of the coccolith ring. When we examined the conservation status of the nannofossil at the three temperatures previously used for the XRD on the same bulk and rinsed sediment, the differences were drastic. The bulk clayey hemipelagic ooze was subjected to severe degradation of its coccolithic content between 400 °C and 480 °C and even more at 540 °C (temperature at which it becomes almost impossible to identify the organism). Moreover, the coccolith crossbar already decomposed 480 °C. Conversely, the coccoliths in the rinsed clayey hemipelagic ooze were much better conserved with increasing temperature, the structure of the coccolith and even the crossbars were still visible at 540 °C (Fig. 4). This microscopic observation points in the same direction as the XRD results: the presence of seawater and therefore of sodium chloride is correlated with the decomposition of carbonated elements. Fig. 4 -SEM images of a coccolith of the same species (Gephyrocapsa oceanica) at three stages (400 °C, 480 °C and 540 °C) of a heating experiment conducted with bulk and rinsed (desalted) powdered sediment from the Timor Sea.

Rock-Eval

The results of Rock-Eval in this study highlight the difference in the behavior of sediments components subjected to heating cycles. Looking at the CO2 curves corresponding to decomposition of carbonates during the oxidation step (the so-called S5 peak) we observed two clearly different behaviors. Decomposition of the calcite in the fossil sediment samples Douvrend and IFP160000 (used as Rock-Eval standard) started at around 650 °C and reached zero around 830 °C, as also demonstrated by Pillot et al. [START_REF] Pillot | Identification and Quantification of Carbonate Species Using Rock-Eval Pyrolysis[END_REF]. The nannofossil pelagic ooze (SU90-08) had almost exactly the same CO2 signature whether it was rinsed or not. The S5 peak of all the samples showed very similar shapes. On the other hand, our results revealed a disparity between the bulk and the rinsed clayey hemipelagic ooze. While the rinsed sample signal was almost identical to those discussed previously, the bulk signal shifted on the temperature scale. In the bulk clayey hemipelagic ooze, carbonate decomposition started early at around 400 °C increased at around 550 °C and peaked at 680 °C (Fig. 5). Fig. 5 -Infrared CO2 cell-detector response (corresponding to the S4CO2 and S5 peaks of Rock-Eval) for two recent sediment samples, one from the Timor Sea (clayey hemipelagic ooze/MD01-2376, ~30% CaCO3) and the other from the central Atlantic Ocean (nannofossil pelagic ooze/SU90-08, ~90% CaCO3). Each sample was analyzed before (as bulk sediment, solid line) and after rinsing (desalted sediment, dashed line). The S4CO2 and S5 peaks for the chalk sample (Douvrend) and the IFP 160000 are shown for the sake of comparison. COLOR FIGURE .     3. Discussion When a sediment or a sedimentary rock is heated, several modifications take place and are accompanied by mass loss. The loss of mass may be due to the moisture in the sample, the presence of structural water or hydroxide (OH -) anions in minerals such as clays, destruction of both organic and mineral phases, etc.

The organic phase is often the first to be destroyed, followed by the decomposition and/or the neoformation of mineral phases. This study focused on identifying the factors responsible for the early thermal destruction of carbonates.

This study highlight the relationship between the presence of salt (NaCl) in the samples and the early decomposition of calcite. Therefore, the hypothesis we proposed in a previous study [START_REF] Baudin | Guidelines for Rock-Eval analysis of recent marine sediments[END_REF] according to which the thermally labile character of poorly crystallized minerals in recent sediments are responsible for the early decomposition of carbonates is not validated. This study revealed different calcite decomposition behaviors in two different recent sediments: a nannofossil pelagic ooze and a clayey hemipelagic ooze (Fig. 5). Calcite decomposition in the nannofossil pelagic ooze occured in the same temperature ranges as fossil sediments such as Douvrend (chalk) and IFP16000, whether bulk or rinsed. On the other hand, the bulk clayey hemipelagic ooze calcite decomposed at lower temperatures although with a similar curve shape. This observation led us to the crucial identification of the cause of these different behaviors, i.e. the presence or absence of clays and hence of structural water and hydroxyl groups in the sediment.

The other methods of analysis used in this study confirmed the responsibility not only of the presence of salt in the sediments but particularly the water and/or hydroxide anions in the early stage of decomposition of the carbonates. XRD analysis (Fig. 2) showed that in the rinsed clayey hemipelagic ooze, the decomposition of calcite is not early even if a small amount of salt remains in the sample. However a small reduction in the calcite principal peak (104) intensity at 600 °C was observed corresponding the classical thermal decomposition ranges of this mineral [START_REF] Pillot | Identification and Quantification of Carbonate Species Using Rock-Eval Pyrolysis[END_REF]. On the other hand, the bulk clayey hemipelagic ooze calcite started decomposing between 400 °C and 500 °C and was already absent from the sample at 600 °C. These observations were visually confirmed by SEM and calcimetry analyses (respectively Fig. 4 and1). According to the results of calcimetry based on the samples heated in the muffle furnace, the calcite content of the fossil sample (Douvrend) and rinsed clayey hemipelagic ooze started decreasing around 550 °C (in agreement with the onset reported by Pillot et al. [START_REF] Pillot | Identification and Quantification of Carbonate Species Using Rock-Eval Pyrolysis[END_REF]), whereas in the bulk clayey hemipelagic ooze, carbonate content reached almost zero at the same temperature. When the sample was rinsed and dried the water present in it evaporated. It is therefore the presence of water and/or hydroxide anions combined with chloride salt in our samples that triggered the early decomposition of carbonates.

It is also interesting to note the difference between the temperature of calcite decomposition in the bulk clayey hemipelagic ooze sample between calcimetry (around 450 °C) and the Rock-Eval oxidation cycle (starting around 450 °C but only becoming significant at around 550 °C). This shift in carbonate decomposition temperatures can be simply explained by the pyrolysis phase undergone before oxidation during the Rock-Eval analysis. Considering this cycle brings the sample up to a temperature of 650 °C (Fig. 6), unlike the calcimetry and SEM post thermal protocol analyses, some of the carbonates have already been attacked during pyrolysis. Rock-Eval analyses of rinsed clayey hemipelagic ooze confirmed this observation by showing lower mineral CO2 values during pyrolysis and standard degradation of calcite in chalk-like temperature ranges. Moreover, the difference in pyroMINC values between rinsed and bulk samples decreased in accordance with the depth of the sample (factor of about 2.5 between pyroMINC values for a sample located at 120cm from the top of the core versus 2 for a deeper sample located at 2260cm from the top of the core), which is to be expected since the sediments contain progressively less water and salt with compaction due to the depth. Fig. 6 -Pyrograms of two core samples (at 120 cm and 2260 cm-depth respectively) highlighting the onset of decomposition of carbonates in bulk sediments during Rock-Eval pyrolysis. COLOR FIGURE. The phenomenon of early decomposition of carbonates in some recent sediments can be explained as follows. The clays have a layered structure that can contain water. Clay minerals consist of silicon, aluminum, oxygen and hydroxyl (OH) ions, organized in layers of two types depending on whether the oxygen atoms or hydroxyl groups are associated in tetrahedra or octahedra. The octahedral and tetrahedral layers coalesce along planes, pooling O and OH, so the sheets are separated by interfoliar spaces. Positive charge deficit can occur due to possible partial substitutions of Si 4+ by Al 3+ in the tetrahedra, Al 3+ by Mg 2+ , Fe 2+ or Fe 3+ , in the octahedra. This is then offset by the insertion of cations or water in the interfoliar spaces. Water molecules are in fact no longer present in the structure but have been expelled as a result of thermal decomposition [START_REF] Fayed | A simplified, non-rigorous, tabular classification of clay minerals with some explanatory notes[END_REF]. When a sediment containing clay minerals is heated, dehydration takes place around 100 °C, releasing water from the clay interlayers. Given the fact that the clay fraction of hemipelagic oozes consists mainly of smectites (Fig. 3), dehydration had to provide a significant amount of water vapor. At higher temperatures around 500 °C, depending on the nature of the clay minerals, dehydroxylation occurs [START_REF] Kenzie | Differential Thermal Analysis, A review[END_REF][START_REF] Chantret | Révision critique de l'utilisation des méthodes thermiques en sédimentologie : cas des smectites (montmorillonites)[END_REF][START_REF] Caillère | Minéralogie des argiles, 1. Structure et propriétés physicochimiques[END_REF][START_REF] Smykatz-Kloss | Differential Thermal Analysis[END_REF][START_REF] Ptáček | The non-isothermal kinetic analysis of the thermal decomposition of kaolinite by thermogravimetric analysis[END_REF][START_REF] Sahnoune | Thermal analysis of dehydroxylation of Algerian kaolinite[END_REF] Moreover Du et al. [START_REF] Du | Thermal stability of the eutectic composition in Nacl-CaCl2-MgCl2 ternary system used for thermal energy storage applications[END_REF] measured melting temperature of chloride salt (NaCl) at 420 °C. These temperatures coincide perfectly with the early decomposition temperatures of calcite observed in our experiments.

Dehydroxylation is significant notably in chlorite and kaolinite, two clay minerals that occur in significant quantities in the samples of hemipelagic oozes collected from the Timor Sea and analysed in the present study. Consequently, and when salted pore water is also present in the sediment, released gaseous water (H2O) and/or hydroxide anions react with it to form acid vapor (such as hypochlorous acid (HClO) or hydrochloric acid (HCl)). In turn, this acid vapor, reacts with the carbonates (CaCO3) present in the sediment causing chemical decarbonatation, which manifests itself as the release of CO2. In this case, the phenomenon of carbonate destruction is therefore chemical and non-thermal, giving rise to early release of CO2 and hence to artificial peaks during Rock-Eval analysis. The CO2 peak corresponding to the destruction of the carbonate phase of the sediment (S5) thus shifts on the temperature scale, leading to overestimation of mineral carbon and underestimation of organic carbon (Fig. 7). Fig. 7 -Plausible origin of the artefactual S5 peak that appears during Rock-Eval analysis (oxidation phase) of recent marine sediment containing hydrated clay minerals, salt and carbonates. An acid reaction between NaCl and water provided by the dehydration/dihydroxylation of clay minerals causes the dissolution of calcite, with the resulting generation of CO2 at temperatures ranging between 550 °C and 775 °C. COLOR FIGURE. However, it is important to really understand the shift of this oxidation CO2 peak. It would be easy to think that the entire peak of CO2 shifts which would lead to a reverse observation, i.e. an overestimation of the TOC and underestimation of the MinC since part of the peak associated with the mineral carbon would be included in signal integrated as organic carbon. It is in fact only the mineral part of the oxidation CO2 peak which is shifted, the formation of acid vapors having no influence on the decomposition temperatures of the organic carbon. Thus, when the mineral part of the CO2 peak shifts to lower temperatures, the Rock-Eval cursor defining the boundary between the portions of the curve integrated as organic and the one integrated as mineral seeks a new minimum on the curve (Fig. 8). As this minimum is most often located at much lower temperatures, the software integrates a large part of the organic signal as mineral carbon against a small part of the mineral signal as an organic carbon thus explaining the phenomenon of overestimation/underestimation observed. Fig. 8 -Illustration of the position of the cursor separating the organic carbon from the mineral carbon on the CO2 signal in oxidation for a classical analysis (A) versus an analysis showing the shift of the mineral part of the CO2 (B) following an early decomposition of the carbonates by chemical attack. COLOR FIGURE.

Application

An example of application is presented below to highlight the importance of rinsing recent carbonate sediment samples before analysis. Figure 9 shows four curves representing the percentage of CaCO3 obtained with different methods (calcimetry, Rock-Eval, XRD and XRF) along a core of hemipelagic sediments collected from the Timor Sea. The curve obtained with XRF analyses (light grey line) is a reference curve, since it is not influenced by any other factors than the carbonate content.

Rock-Eval analyses of the bulk sediment (solid line) overestimate carbonate contents while those obtained using calcimetry (dotted line) underestimate them. Indeed, as seen above, the presence of water in the clays and salt in the sediment causes early chemical degradation of the carbonates during the Rock-Eval analysis, inducing a shift of the CO2 peak emitted and hence overestimation of mineral carbon. When calcimetry is used, this phenomenon does not occur because the sample is not heated. However with calcimetry, there is also an error in the estimation of the percentage of CaCO3. Indeed, the presence of salt in the sample during weighing results in overestimation of the weight of the sample and hence to an underestimation of the carbonate content. During Rock-Eval of the rinsed sediments, the analyzed sample no longer contains salt, and its carbonate contents are consequently much closer to those obtained with XRF (black dots). This observation confirms the results of our study and highlights the importance of using properly processed samples before analysis and interpretation of the results. It is also interesting to note the variation in these shifts between the different methods and along the core. We observed indeed that the gap between the CaCO3 percentages produced by Rock-Eval for the bulk sediment and the XRF analyses is much larger in the top part of the core but narrows towards the deeper part. This observation can be explained by simple geological phenomena. Indeed, the upper part of the core (close to the surface) is comprised of highly porous sediments and consequently contain a lot of interstitial marine water, resulting in a large quantity of salt when the sediment is oven dried. Towards the deeper parts, when compaction takes effect, the porosity of the sediment decreases and consequently contains less water and fewer salt crystals after drying, thereby limiting errors caused by their presence. Fig. 9 -Changes (5 point-smoothing) in CaCO3 contents along a 27.9 m-long core of hemipelagic sediments from the Timor Sea. The data were acquired using a calcimetric bomb (dotted line) or were derived from the MinC of Rock-Eval (solid line) of the bulk sediment. Selected samples were rinsed using deionized water and re-analyzed using Rock-Eval (black dots). Variations in the Ca content (light grey line) derived from XRF are shown for the sake of comparison. The latter are in good agreement with CaCO3 determined in rinsed sediments. Beyond overestimation of mineral carbon and underestimation of organic carbon, the analysis of recent bulk sediments may lead to other misinterpretations. In fact, a certain number of parameters such as the oxygen index (OI) used to define the origin of organic matter are calculated from the CO2 emission curves obtained during the Rock-Eval analysis. Thus, if this analysis is distorted by a phenomenon such as the early destruction

MD01-2376

of carbonates through a chemical reaction, the CO2 emission curves shift and the parameters based on them are no longer reliable. This is highlighted in figure 10 below, where the OI reaches 1400mg CO2/g TOC for bulk sediment whereas it does not exceed 600 in the rinsed sediment, a drastic difference for an interpretation based on the pseudo van Krevelen diagram, for example. Fig. 10 -Comparison of organic matter typing using a modified van Krevelen diagram (HI-OI) for bulk and rinsed samples of recent hemipelagic sediment from the Timor Sea. The extremely high OI values for bulk sediment are artefactual (see Fig. 5 for explanation).

Finally, an important point when comparing bulk and rinsed sediments that has not yet been discussed in this paper is the mass effect. Indeed, the weight of a bulk sample is not the same as that of a rinsed sample because the salt disappears along with part (soluble) of the organic carbon. This effect must of course be taken into account but accounts for a minimum part of the observed differences between the bulk and rinsed samples. Figure 8 shows that the OI of the average bulk samples is around 600 mg CO2/g TOC versus 300mg CO2/g TOC for rinsed samples. If the cause of this very large difference was only linked to the loss of mass, this would imply a factor of 2 between TOC values of the bulk and rinsed samples. However it is very unlikely that rinsing causes such a difference in TOC between the bulk and the rinsed sample, let alone that the mass loss due to the removing of salt and the loss of part of the organic carbon is entirely responsible for such a difference. Moreover, core analyses allow us to definitely disregard the impact of this phenomenon in this study. Indeed, the TOC of sample 1-120 (120 cm from the top of the core) was 0.6% for the bulk sample versus 1.15 for the rinsed sample. This result confirms the underestimation of organic carbon discussed previously in this article but also make it possible to hold the mass effect responsible for this difference. However, sample 18-110 (taken deeper in the core at 2260 cm) was much more compacted than sample 1-120 and therefore did not contain enough water to cause the early chemical dissolution reaction of the carbonates. The TOC values for the rinsed and non-rinsed samples were almost identical (~ 1.11%), thereby confirming the negligible effect of mass loss in this situation.

Conclusion

Rock-Eval analyzes, widely used in the oil industry, were recently found to be appropriate for inferring the CaCO3 content of a sample and identifying the type of carbonates it contains [START_REF] Pillot | Identification and Quantification of Carbonate Species Using Rock-Eval Pyrolysis[END_REF]. However, the relevance of the CO2 production curve shape and temperature ranges obtained during thermal destabilization of carbonates is called into question by the transposition of this method to recent clayey sediments. Indeed, calcite (and probably all carbonates in recent sediments) becomes unstable at temperatures much lower than those of the fossil sedimentary rocks on which the mineral recognition records in the literature are based. Our study allowed us to highlight a shift in peak CO2 temperatures observed with Rock-Eval and corresponding to the destruction of the carbonate phase within a sediment. This phenomenon is due to the combined presence of salt and water Oxygen Index (mg CO 2 /g TOC) Hydrogen Index (mg HC/g TOC) and/or hydroxide anions expelled from clay minerals that leads to an acid reaction and early decomposition of carbonates itself leading to overestimation of mineral carbon and underestimation of organic carbon. Moreover, this CO2 curve can also be misinterpreted if one relies on the CO2 peaks obtained with Rock-Eval to identify the type of carbonate present in the sample according to the method of Pillot et al. [START_REF] Pillot | Identification and Quantification of Carbonate Species Using Rock-Eval Pyrolysis[END_REF]. The same goes for the OI that is overestimated due to underestimation of organic carbon. Care should thus be taken when analyzing recent hemipelagic samples and the samples should be thoroughly rinsed before analysis and interpretation to avoid any misinterpretation of the results.

  Calcimetry bulk sed.Rock-Eval bulk sed. Rock-Eval rinsed sed.


Diego Alfonso Godoy Pérez

Alberto Moreira

Jorge Junior

Gabriel Hitoshi Asato

Jean-Claude Leprêtre

Virginie Roche

Claudemiro Bolfarini

Walter Jose Botta

Jean- Claude Lepretre

Surface anodization of the biphasic Ti13Nb13Zr biocompatible alloy: Influence of phases on the formation of TiO2 nanostructures

Keywords: Ti13Nb-13Zr alloy, Anodization, TiO2 nanotubes, TiO2 nanopores, Surface modification

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

Introduction

Titanium and its alloys have increased their use as biomaterials due to their low elastic modulus compared to Co-Cr alloys and stainless steels. Such a smaller elastic modulus reduces the stress shielding effect, and, consequently, the possibility of bone resorption and implant loosening [START_REF] Sabirov | Bulk Nanostructured Materials with Multifunctional Properties[END_REF]. Also, titanium alloys are increasingly used in medicine due to their excellent biocompatibility and high specific mechanical properties [START_REF] Bortolan | Fatigue strength of Ti-6Al-4V alloy with surface modified by TiO2 nanotubes formation[END_REF].

Ultrafine-grained titanium alloys, produced by severe plastic deformation, may improve wear resistance, mechanical properties and the bioactivity in accelerating apatite precipitation and amplified biological responses when compared to their coarse-grained counterpart material because of Hall-Petch strengthening, higher surface energy, increased number of nano-sized grooves, and a significant increase in the cell adhesion/proliferation [START_REF] Kim | In vitro biocompatibility of equal channel angular processed (ECAP) titanium[END_REF].

The passivation phenomenon contributes to the above characteristics and makes titanium alloys suitable for use in dental and orthopedic implants, without promoting adverse reactions locally or systemically [START_REF] Hernández-López | Correlation of the nanostructure of the anodic layers fabricated on Ti13Nb13Zr with the electrochemical impedance response[END_REF]. Such adverse reactions could be related to the interactions of the surface layer of the TiO2 on the implant and bone tissue [START_REF] Hazan | Bone growth to metal implants is regulated by their surface chemical properties[END_REF]. However, the biological response of an implant depends on the physicochemical properties of the surface [START_REF] Oliveira | Functionalization of titanium dioxide nanotubes with biomolecules for biomedical applications[END_REF]. Surface modification techniques are required for biomedical materials to improve their surface properties and biological activity [START_REF] Bauer | Engineering biocompatible implant surfaces: Part I: Materials and surfaces[END_REF].

The use of TiO2 nanotubes is becoming one of the most exciting techniques as surface modification of metallic biomaterials and several authors have reported the improvement of bioactivity through changes on the surface of titanium alloys by growing TiO2 nanotubes [START_REF] Stróż | Production, structure and biocompatible properties of oxide nanotubes on Ti13Nb13Zr alloy for medical applications[END_REF][START_REF] Córdoba-Torres | Electrochemical impedance analysis of TiO2 nanotube porous layers based on an alternative representation of impedance data[END_REF][START_REF] Bai | The effect of annealing temperatures on surface properties, hydroxyapatite growth and cell behaviors of TiO2 nanotubes[END_REF][START_REF] Das | TiO2 nanotubes on Ti: Influence of nanoscale morphology on bone cell-materials interaction[END_REF][START_REF] Oliveira | Obtaining self-organized nanotubes on biomedical Ti-Mo alloys[END_REF][START_REF] Zhou | Anodic TiO2 nanotube layers: why does self-organized growth occur-a mini review[END_REF][START_REF] Campanelli | Fatigue resistance, electrochemical corrosion and biological response of Ti-15Mo with surface modified by amorphous TiO2 nanotubes layer[END_REF]. Anodization is based on employing a voltage (some tens of volts) connecting a sample and an inert counter electrode. Both sample and electrode are soaked in an electrolyte comprising water and fluoride ions (F -). Inside the solution, there is a competition between the oxidation and oxide dissolution by the action of F -ions, which may lead to the development of a self-organized layer of TiO2 nanotubes. The applied voltage, anodization time, type of the electrolyte, concentration of water and fluoride ions, temperature, or pH can change the tubes morphology (length, diameter, spacing, or wall look) [START_REF] Roy | TiO2 nanotubes: synthesis and applications[END_REF][START_REF] Kulkarni | Titanium nanostructures for biomedical applications[END_REF][START_REF] Macak | TiO2 nanotubes: Self-organized electrochemical formation, properties and applications[END_REF][START_REF] Lee | One-dimensional titanium dioxide nanomaterials: nanotubes[END_REF]. The sequence of nanotubes' formation can be well visualized in, for instance, Fig. 3a of [START_REF] Macak | Mechanistic aspects and growth of large diameter self-organized TiO2 nanotubes[END_REF] and references therein]. In short, in the first anodization stage, there is the formation of a thin and dense TiO2 layer, which possesses some statistically distributed positions where, close to the dense oxide layer, arises an accelerated local dissolution, nucleating pores. Such nucleation appears on nearly the whole surface. On some sites, there are visible destabilized stretches of the first compact oxide. After this, pores have a random look. Following this stage, pore bottoms start to have a marked dissolution, deepening them. At this point, there will be no vestige of the initial compact TiO2 layer on the surface. Continued anodization leads the porous structure to promptly change to a nanotubular structure, fully evolving a self-organized nanotubes' layer.

Anodic ordered nanoporous TiO2 layers were recently reported [START_REF] Kim | Formation of a Non-Thickness-Limited Titanium Dioxide Mesosponge and its Use in Dye-Sensitized Solar Cells[END_REF][START_REF] Lee | Anodic formation of thick anatase TiO2 mesosponge layers for high-efficiency photocatalysis[END_REF][START_REF] Roy | Size-selective separation of macromolecules by nanochannel titania membrane with self-cleaning (declogging) ability[END_REF]. Such nanoporous layers were of some tens of microns and had channels ranging from some nanometers to tens of nanometers in diameter. These nanostructured morphologies have good adhesion to the substrate, but so far they have not been exploited. However, the formation of such nanoporous layers are produced by etching the first formed thick TiO2 oxide layer produced by anodization and does not follow the above sequence, and they do not even have the morphology mentioned above. Nonetheless, one can generally say that arrays of selforganized oxide tubes or pores can be achieved by an anodization process of a proper metal.

Hence, related to the alloy studied in this work, self-organized oxide nanotube/pore layers of acceptable properties have been published on binary alloys, such as TiNb [START_REF] Ghicov | TiO2-Nb2O5-Nanoröhren mit elektrochemisch einstellbaren Morphologien[END_REF][START_REF] Ghicov | Lattice Widening in Niobium-Doped TiO2 Nanotubes: Efficient Ion Intercalation and Swift Electrochromic Contrast[END_REF], TiZr [START_REF] Yasuda | Control of morphology and composition of self-organized zirconium titanate nanotubes formed in (NH4) 2SO4/NH4F electrolytes[END_REF][START_REF] Yasuda | Mechanistic aspects of the self-organization process for oxide nanotube formation on valve metals[END_REF][START_REF] Yasuda | Electrochemical formation of self-organized zirconium titanate nanotube multilayers[END_REF][START_REF] Yasuda | Formation of Self-Organized Zirconium Titanate Nanotube Layers by Alloy Anodization[END_REF], and on ternary [START_REF] Feng | Electrochemical formation of selforganized anodic nanotube coating on Ti-28Zr-8Nb biomedical alloy surface[END_REF][START_REF] Macak | Self-organized nanotubular oxide layers on Ti-6Al-7Nb and Ti-6Al-4V formed by anodization in NH4F solutions[END_REF] and multicomponent complex alloy systems [START_REF] Tsuchiya | Self-Organization of Anodic Nanotubes on Two Size Scales[END_REF][START_REF] Tsuchiya | Nanotube oxide coating on Ti-29Nb-13Ta-4.6Zr alloy prepared by selforganizing anodization[END_REF].

It is remarkable that, for the majority of elements and alloys, ordered tubular and nanoporous morphologies have been observed after anodization. Some of them can also shift from one to another (for instance, TiAl, TiNb, and TiTa) [START_REF] Ghicov | Lattice Widening in Niobium-Doped TiO2 Nanotubes: Efficient Ion Intercalation and Swift Electrochromic Contrast[END_REF][START_REF] Berger | Transition from Nanopores to Nanotubes: Self-Ordered Anodic Oxide Structures on Titanium-Aluminides[END_REF][START_REF] Wei | High aspect ratio ordered nanoporous Ta2O5 films by anodization of Ta[END_REF][START_REF] Song | Anodic oxidation behaviour of Al-Ti alloys in acidic media[END_REF]. Nevertheless, due to chemistry differences (solubility of the oxide in fluoride and the corresponding solubility of metal fluorides in distinct solutions), a preferred morphology is observed in specific elements or alloys.

In other words, a specific morphology prevails under the majority of anodization conditions. For example, nanotubular anodic oxides are typically observed on zirconium and hafnium while, nanoporous are usually produced on aluminum, niobium, and tantalum.

Regarding alloys, the morphology is dominated by the matrix. Also, considering the wellknown formation mechanism of TiO2 nanotubes [START_REF] Roy | TiO2 nanotubes: synthesis and applications[END_REF][START_REF] Kulkarni | Titanium nanostructures for biomedical applications[END_REF][START_REF] Macak | TiO2 nanotubes: Self-organized electrochemical formation, properties and applications[END_REF][START_REF] Lee | One-dimensional titanium dioxide nanomaterials: nanotubes[END_REF][START_REF] Macak | Mechanistic aspects and growth of large diameter self-organized TiO2 nanotubes[END_REF][START_REF] Minagar | A review of the application of anodization for the fabrication of nanotubes on metal implant surfaces[END_REF][START_REF] Huo | Fabrication, modification, and biomedical applications of anodized TiO2 nanotube arrays[END_REF][START_REF] Regonini | A review of growth mechanism, structure and crystallinity of anodized TiO2 nanotubes[END_REF][START_REF] Mor | Fabrication of Tapered, Conical-Shaped Titania Nanotubes[END_REF], it is notorious that i) for a specific electrolyte: the higher the voltage, the larger diameter and the longer the nanotubes, ii) regarding organic electrolytes: the less dilute the solution, the smoother the nanotubes wall.

Ethylene glycol electrolyte containing a small water amount will significantly thicken the nanotubes layer.

Following anodization, generally, the oxide composition is consonant with the ratio of alloying elements. Occasionally, anodic oxides can also exhibit small amounts of oxide mixtures. TiZr was the only alloy reported to have a complete conversion to a mixed oxide [START_REF] Yasuda | Formation of Self-Organized Zirconium Titanate Nanotube Layers by Alloy Anodization[END_REF].

Just-anodized nanotubes/nanoporous are amorphous, but they can be crystallized through heat treatment into the anatase phase or a mix of anatase and rutile phases, which depends on the treatment temperature [START_REF] Bai | The effect of annealing temperatures on surface properties, hydroxyapatite growth and cell behaviors of TiO2 nanotubes[END_REF][START_REF] Roy | TiO2 nanotubes: synthesis and applications[END_REF][START_REF] Huo | Fabrication, modification, and biomedical applications of anodized TiO2 nanotube arrays[END_REF]. Both nanotubes size and crystallinity alter surface properties and, inevitably, leading to changes in the bioactivity response.

The work herein aimed to analyze the effect of the processing route comparing structural and microstructural evolutions of the Ti13Nb13Zr alloy. Part of the samples was strained by the SPD method of high-pressure torsion (HPT), which led to an ultrafine grained (UFG) microstructure. The samples' surfaces were modified for micro-nano topographic conversion by anodization, in an attempt to bioactivate them. Anodized surfaces were heattreated and contrasted to non-anodized samples. The resulting surface' morphologies of anodized samples were compared with non treated surfaces of deformed and undeformed conditions. Therefore, the purpose of the investigation was to determine the effect of the microstructure and phases present on the formation of nanostructures on samples' surfaces, changes in bioactivity, and the corrosion behavior before and after the surface modification.

Corrosion properties were assessed and, finally, the resultant in-vitro bioactivities were analyzed considering a possible synergy between processing and surface treatment. Corrosion properties should be improved, and surface bioactivity should eventually lead to osseointegration paths into the bone and more improved in samples processed by SPD for both processing conditions.

Material and methods

The material used in this study was the biomedical β-rich Ti-13Nb-13Zr alloy (ERCATA GmbH, Development Co., Ltd., Germany). As the β transus temperature for this alloy is 735 °C [START_REF] Davidson | New surface-hardened, low-modulus, corrosion-resistant Ti13Nb13Zr alloy for total hip arthroplasty[END_REF], specimens cut in the form of a disk with 10 mm in diameter and 1 mm in thickness were heat-treated from the β field at 750 °C for 1 hour and then fast cooled in water.

Our previous results [START_REF] Pérez | Ultrafine-Grained Ti13Nb13Zr Alloy Produced by Severe Plastic Deformation[END_REF] showed that non-equilibrium metastable martensite (α') phase and a retained β phase appeared after heating and quenching.

A set of specimens were mechanically deformed by High-Pressure Torsion (HPT) using a load of 4.5 GPa and three turns, to refine the microstructure and to observe the effect of the redistribution of α' and β phases on the formation of nanopores/nanotubes. X-ray diffraction (XRD) analysis was conducted using a Siemens D5000 diffractometer with CuKα radiation and a scanning speed of 2°/min from 20° to 90°.

The surfaces of the samples, and cross sections and plan views of anodic coatings were examined by scanning electron microscopy (SEM) using the Zeiss Ultra-55 FEG microscope equipped with an Energy Dispersive X-ray (EDX) microanalysis system with an SDD detector (BRUKER AXS-30mm²).

Orientation phase mapping was performed by Transmission electron microscopy (TEM) using an FEI TECNAI G2, operated at 200 kV, equipped with a NanoMEGAS ASTAR and Digistar P1000 system, which allow Automated Crystal Orientation Mapping (ACOM) in nanoprobe mode enabling orientation/phase maps with a spatial resolution of about 1 nm.

Before anodization experiments, disk samples were successively ground with P320, P600, P1200, and P2400 grade SiC emery paper, and then cleaned with deionized water (Milli-Q®), ethanol and acetone.

Electrochemical anodization setup to produce nanotubes on surfaces of non-deformed and HPT-deformed samples was the same and consisted of a conventional two electrodes system with a platinum grid as the counter electrode, a Ti13Nb13Zr disk as the working electrode, and a power supply (TCR Electronic measurements inc.). The counter electrode and the working electrode were always kept at the same distance (~3 cm).

During the development of this work, we have tried different conditions for the formation of nanotubes as described in our previous work [START_REF] Hilario | Influence of morphology and crystalline structure of TiO2 nanotubes on their electrochemical properties and apatiteforming ability[END_REF]. After evaluation of the reported effects of voltage, time and types of solutions, and solution dilution [START_REF] Roy | TiO2 nanotubes: synthesis and applications[END_REF][START_REF] Kulkarni | Titanium nanostructures for biomedical applications[END_REF][START_REF] Macak | TiO2 nanotubes: Self-organized electrochemical formation, properties and applications[END_REF][START_REF] Lee | One-dimensional titanium dioxide nanomaterials: nanotubes[END_REF][START_REF] Macak | Mechanistic aspects and growth of large diameter self-organized TiO2 nanotubes[END_REF][START_REF] Minagar | A review of the application of anodization for the fabrication of nanotubes on metal implant surfaces[END_REF][START_REF] Huo | Fabrication, modification, and biomedical applications of anodized TiO2 nanotube arrays[END_REF][START_REF] Regonini | A review of growth mechanism, structure and crystallinity of anodized TiO2 nanotubes[END_REF][START_REF] Mor | Fabrication of Tapered, Conical-Shaped Titania Nanotubes[END_REF],

we varied the anodization voltage from 10 V to 60 V, the time from 45 min to 1 hour and the dilution from 7.5 vol% to 25 vol% of water, solutions of glycerol, containing 0.25 M of NH4F and ethylene glycol and containing 0.25 M of NH4F.

The best condition, which will be presented in this paper, was obtained for the anodization carried out applying a potential of 20 V for one hour using glycerol as a working electrolyte, containing 25 vol% of water and 0.25 M of NH4F. All other conditions did not produce any tubular morphology (porous or nanotubes) in the β phase. The electrolyte was kept at 25 °C throughout the anodization process using a microprocessor controlled refrigerated heating bath with an air-cooled refrigerating unit from Huber Kältemaschinenbau AG, Germany, which can control the temperature within the range of -25 to 200 °C.

After anodic oxidation, as the anodized TiO2 nanostructures are amorphous and of very low mechanical strength, the samples were annealed in air at 550 °C for two hours to crystallize TiO2 nanostructures aiming their structural reinforcement.

Electrochemical analyzes were carried out using a conventional three-electrode cell, where the sample was the working electrode, a platinum grid the counter electrode, and a Saturated Calomel Electrode (SCE) the reference electrode. A Simulated Body Fluid (SBF) was used as electrolyte (NaCl 8.035 g L -1 , KCl 0.225 g L -1 , NaHCO3 0.355 g L -1 , K2HPO4•3H2O 0.231 g L -1 , MgCl2•6H2O 0.311 g L -1 , CaCl2 0.292 g L -1 and Na2SO4 0.072 g L -1 ) [START_REF] Kokubo | How useful is SBF in predicting in vivo bone bioactivity?[END_REF]. All the electrochemical analyses were performed in triplicate to confirm the reproducibility of results. Open Circuit Potential (OCP) was always registered throughout 1 h of resting in the electrolyte. The corrosion resistance was estimated through potentiodynamic polarization curves recorded within the range of -0.03 to 1.5 V versus the SCE, using a scan rate of 0.5 mV s -1 .

The in vitro formation of apatite is an indicator of the in vivo bioactivity of a material surface. Thus, to evaluate the bioactivity of different samples used in this study, they were soaked in 40 mL of SBF solution under a static condition for 1, 7 and 14 days. The surface with nanotubes was placed perpendicularly to the bottom of the beaker because the SBF is very unstable and may spontaneously precipitate apatite on the specimen surface. After soaking, samples were gently rinsed with deionized water and naturally dried in a desiccator.

The apatite-forming ability was estimated by FEG-SEM observations, EDS, and XRD, and by weighing the sample before and after immersion [START_REF] Kokubo | How useful is SBF in predicting in vivo bone bioactivity?[END_REF].

Raman Spectroscopy was performed on samples' surfaces using the RENISHAW InVia Raman Microscope employing an Argon Ion laser of 20 mW with a wavelength of 514 nm.

Results and discussion

Figure 1 exhibits XRD patterns of Ti13Nb13Zr alloy samples in conditions, undeformed and deformed by HPT. The indexation of Bragg's peaks denotes the presence of the hcp α' (ICSD # 253841) and the bcc β (ISCD # 645545) phases before and after deformation. Figure 1b indicates that the fraction of β phase had a slight augment after processing by HPT. Such an increase was probably due to solute redistribution and mechanical activation, driving to a rapid phase transformation of the coarse non-equilibrium metastable martensite (hcp α') during severe plastic deformation. Also in Figure 1b, one may note the broadening of peaks after HPT processing, pointing to enormous grain refinement.

Figure 2 presents morphological aspects of the alloy microstructure before (Fig. 2a) and after deformation (Fig. 2b). Figure 2a displays an SEM image showing that, before deformation, the microstructure of Ti13Nb13Zr alloy is formed by retained β phase and the coarse non-equilibrium metastable acicular martensite phase (hcp-α'), precipitated after heating and quenching, confirming the XRD analysis in Fig. 1a. Grain size measurements gave an average initial size of ~250 μm for the β phase and an average length of ~20 μm for the martensite laths of α' with widths ranging between 1 and 2 μm.

The microstructural refinement and phases' distribution after HPT was evaluated by ACOM-TEM analysis. Figure 2b presents a virtual TEM bright field image superimposed to a mapping of high-angle grain boundaries (> 15 o ), denoting recrystallized grains. As stated before, when discussing Figure 1b, Figure 2b shows a huge grain refinement and spatial redistribution of α' and β phases after HPT processing. Additionally, one may observe an inhomogeneous distribution of grain sizes. Determination of the grain size distribution resulted in an average size of ~203 nm for martensite laths and ~50 nm for the β phase.

Statically assessment of the amount of each phase revealed that, after deformation, the Ti13Nb13Zr alloy sample was composed of about 77% of hcp α' and 23% of β. 3b). Figures 3b and3c indicate that the self-organized arrays of nanotubes were obtained with the following average dimensions: diameter of ~80 nm, a length of ~2 µm and a thickness of ~18 nm. By comparing Figures 3d and3a, one may observe that HPT processing caused the refinement and spatial redistribution of α' and β phases. The formation of the nanotubes was kept sole within the α' phase (Figure 3e), however, with the difference that nanotubes are not so organized and are deformed in their longitudinal direction (Figure 3f) suggesting that their growth follow the spatial distortion of the α' phase. This result shows in an unambiguously way that nanotubes uniquely form on the α' phase.

Such results, even the unsuccessful ones described in the methodology, are of interest since, as stated in the introduction [START_REF] Ghicov | TiO2-Nb2O5-Nanoröhren mit elektrochemisch einstellbaren Morphologien[END_REF][START_REF] Ghicov | Lattice Widening in Niobium-Doped TiO2 Nanotubes: Efficient Ion Intercalation and Swift Electrochromic Contrast[END_REF][START_REF] Yasuda | Control of morphology and composition of self-organized zirconium titanate nanotubes formed in (NH4) 2SO4/NH4F electrolytes[END_REF][START_REF] Yasuda | Mechanistic aspects of the self-organization process for oxide nanotube formation on valve metals[END_REF][START_REF] Yasuda | Electrochemical formation of self-organized zirconium titanate nanotube multilayers[END_REF][START_REF] Yasuda | Formation of Self-Organized Zirconium Titanate Nanotube Layers by Alloy Anodization[END_REF][START_REF] Feng | Electrochemical formation of selforganized anodic nanotube coating on Ti-28Zr-8Nb biomedical alloy surface[END_REF][START_REF] Macak | Self-organized nanotubular oxide layers on Ti-6Al-7Nb and Ti-6Al-4V formed by anodization in NH4F solutions[END_REF][START_REF] Tsuchiya | Self-Organization of Anodic Nanotubes on Two Size Scales[END_REF][START_REF] Tsuchiya | Nanotube oxide coating on Ti-29Nb-13Ta-4.6Zr alloy prepared by selforganizing anodization[END_REF][START_REF] Berger | Transition from Nanopores to Nanotubes: Self-Ordered Anodic Oxide Structures on Titanium-Aluminides[END_REF][START_REF] Wei | High aspect ratio ordered nanoporous Ta2O5 films by anodization of Ta[END_REF][START_REF] Song | Anodic oxidation behaviour of Al-Ti alloys in acidic media[END_REF], due to the solubility of the metal oxide in fluoride and the solubility of metal fluorides in distinct solutions, a favored morphology is expected to prevail under the majority of anodization conditions for most elements. In the case of this work, it would be expected that titanium and zirconium should produce nanotubular anodic oxides and nanoporous should grow on niobium. In the Ti13Nb13Zr alloy, where Nb and Zr are in solid solutions in both phases, the morphology should be dominated by titanium [START_REF] Ghicov | TiO2-Nb2O5-Nanoröhren mit elektrochemisch einstellbaren Morphologien[END_REF][START_REF] Ghicov | Lattice Widening in Niobium-Doped TiO2 Nanotubes: Efficient Ion Intercalation and Swift Electrochromic Contrast[END_REF][START_REF] Yasuda | Control of morphology and composition of self-organized zirconium titanate nanotubes formed in (NH4) 2SO4/NH4F electrolytes[END_REF][START_REF] Yasuda | Mechanistic aspects of the self-organization process for oxide nanotube formation on valve metals[END_REF][START_REF] Yasuda | Electrochemical formation of self-organized zirconium titanate nanotube multilayers[END_REF][START_REF] Yasuda | Formation of Self-Organized Zirconium Titanate Nanotube Layers by Alloy Anodization[END_REF][START_REF] Feng | Electrochemical formation of selforganized anodic nanotube coating on Ti-28Zr-8Nb biomedical alloy surface[END_REF][START_REF] Macak | Self-organized nanotubular oxide layers on Ti-6Al-7Nb and Ti-6Al-4V formed by anodization in NH4F solutions[END_REF][START_REF] Tsuchiya | Self-Organization of Anodic Nanotubes on Two Size Scales[END_REF][START_REF] Tsuchiya | Nanotube oxide coating on Ti-29Nb-13Ta-4.6Zr alloy prepared by selforganizing anodization[END_REF][START_REF] Berger | Transition from Nanopores to Nanotubes: Self-Ordered Anodic Oxide Structures on Titanium-Aluminides[END_REF][START_REF] Wei | High aspect ratio ordered nanoporous Ta2O5 films by anodization of Ta[END_REF][START_REF] Song | Anodic oxidation behaviour of Al-Ti alloys in acidic media[END_REF], which is the principal element even in the β phase. In fact, as presented in Table 1, which summarizes some results from the literature using similar conditions to this work, nanotubes, with a diameter close to the ones found in this work but with half-length, grew in the Ti6Al4V alloy, thus confirming the prevalence of Ti presence in the formation of nanotubes. From this finding, one may suppose that the amount of alloying elements can also make the difference and not only the major alloying component forming the matrix, as it is the case of Ti in the β phase.

From the above discussion one can speculate that the distinct nanostructures obtained in this work was due to a competition between the dissolution and formation of oxides as a function of the different chemical composition of both phases [15-19, 36-39, 46], or, in other words, a rate-controlling factor is influencing the tubes' growth in the β phase.

Regarding the tubes' growth mechanism explained in the introduction [START_REF] Macak | Mechanistic aspects and growth of large diameter self-organized TiO2 nanotubes[END_REF], one can presume that diffusion can be split into two components: (i) inside tubes and (ii) in the adjacent layer surrounded by the electrolyte. Due to changes in the diffusion layer thickness, the growing shape will be affected thus changing the concentration at the pore entrance. From the present results it is impossible to distinguish what is the critical factor, the internal diffusion of F -or the external one of [TiF6] 2-. However, based on a model described elsewhere [START_REF] Yasuda | Control of morphology and composition of self-organized zirconium titanate nanotubes formed in (NH4) 2SO4/NH4F electrolytes[END_REF] and well-done experiments in [START_REF] Macak | Mechanistic aspects and growth of large diameter self-organized TiO2 nanotubes[END_REF], one can suggest that the diffusion rate is controlled by the [TiF6] 2-. Therefore, throughout the oxide growth, it will take place an increase of fluorides at pore boundaries, sensitizing these regions. Fluoride-rich regions become prone for selective chemical dissolution, thus etching out pore boundaries and leading to isolated tube shapes.

From the above reasoning, one may assume that, because of titanium fluorides' solubility in water, these areas are crucial for nanotubes' generation instead of porous structures. Consequently, under the same good anodization condition found in this work, β and α' phases will inevitably produce different nanostructures due to the high solubility of Ti-F or Ti-O-F compounds in aqueous media unlike Nb-F and Zr-F species that are much less soluble. Thus, an ordered nanoporous oxide will be formed on the Ti-poorer β phase (less Ti fluorides), but nanotubular morphology will be found for the Ti-richer α' phase (more Ti fluorides). We shall come back to this subject when discussing the structure of nanotubes.

Finally, despite also growing nanopores, the dimensions of formed nanotubes in this work presented sizes equivalent to the ones found in the literature for Ti and Ti6Al4V alloy, using similar electrolyte and applied voltage. Actually, our experiments produced nanotubes with an average length larger than for some of those materials (Table 1).

Figure 4a presents XRD patterns of the Ti13Nb13Zr alloy after anodization. In this condition, amorphous peaks of TiO2 are observed (ICSD # 154605), indicating that nanotubes and nanopores are amorphous. When TiO2 nanotubes and nanopores are heat-treated at 550 °C for two hours a mix of anatase and rutile was obtained (Figure 4b). Indexation of Bragg peaks indicates the presence of the phases TiO2-anatase (ICSD # 154601, 2θ = 25.4°, 38.3°, 48.1°, and 55.2° [START_REF] Ji | Fabrication of double-walled TiO2 nanotubes with bamboo morphology via one-step alternating voltage anodization[END_REF]) and TiO2-rutile (ICSD # 9161, 2θ = 27.4° and 36.0° [START_REF] Ibrahim | Anatase/Rutile TiO2 composite prepared via sonochemical process and their photocatalytic activity[END_REF]).

As reported in the literature [START_REF] Varghese | Crystallization and High-Temperature Structural Stability of Titanium Oxide Nanotube Arrays[END_REF], nanotubes' morphology does not change up to 580 o C, when annealing in the air, with no noticeable changes in the pore diameter and wall thickness after annealing during three hours. Also reported in the literature [START_REF] Regonini | A review of growth mechanism, structure and crystallinity of anodized TiO2 nanotubes[END_REF][START_REF] Munirathinam | Titania nanotubes from weak organic acid electrolyte: Fabrication, characterization and oxide film properties[END_REF][START_REF] Tighineanu | Conductivity of anodic TiO2 nanotubes: Influence of annealing conditions[END_REF][START_REF] Roguska | Improvement of the bio-functional properties of TiO2 nanotubes[END_REF] is that the presence of mixed structures of anatase and rutile phases leads nanotubes to be in the anatase structure, while the rutile structure forms a thick layer at the bottom of such nanotubes.

Thickening of the rutile layer occurs during annealing at higher temperatures [START_REF] Roguska | Improvement of the bio-functional properties of TiO2 nanotubes[END_REF]. In such a case, the nanotube structure tends to contract and disintegrate [START_REF] Bai | The effect of annealing temperatures on surface properties, hydroxyapatite growth and cell behaviors of TiO2 nanotubes[END_REF][START_REF] Varghese | Crystallization and High-Temperature Structural Stability of Titanium Oxide Nanotube Arrays[END_REF][START_REF] Mohan | Electrochemical behavior and effect of heat treatment on morphology, crystalline structure of self-organized TiO2 nanotube arrays on Ti-6Al-7Nb for biomedical applications[END_REF].

It is noteworthy that the mechanical properties of the distinct phases of TiO2 nanostructures are dissimilar. In other words [START_REF] Bai | The effect of annealing temperatures on surface properties, hydroxyapatite growth and cell behaviors of TiO2 nanotubes[END_REF], when the amorphous TiO2 transforms into anatase, there are no considerable changes in hardness and Young's modulus. Conversely, the rutile phase leads to an increase in both properties. Indeed, such a kind of mixed structure appears to be very interesting considering hard tissue substitution. Bearing in mind that [START_REF] Bai | The effect of annealing temperatures on surface properties, hydroxyapatite growth and cell behaviors of TiO2 nanotubes[END_REF] highest hardness increases the wear resistance and smallest Young's modulus minimize stress-shielding, it becomes that the most advantageous structure would be the mix of anatase (outermost surface) and rutile (interface), which also looks to be an attractive implant surface, since the mixed structure may favor the mechanical interaction between the implant and bone.

Moreover, one may anticipate that the close-packed structure of rutile phase, when compared to anatase, improves the corrosion resistance of the implant in a biological environment, hence preventing Ti-ion release [START_REF] Munirathinam | Titania nanotubes from weak organic acid electrolyte: Fabrication, characterization and oxide film properties[END_REF][START_REF] Roguska | Improvement of the bio-functional properties of TiO2 nanotubes[END_REF][START_REF] Zhu | Anodic oxide films containing Ca and P of titanium biomaterial[END_REF].

At last, bone growth is favored by high surface bioactivity, thus being able to avoid premature implant substitution. Hereof, once anatase, which is on the outermost surface, and hydroxyapatite have comparable lattices [START_REF] Oh | Titanium oxide nanotubes with controlled morphology for enhanced bone growth[END_REF][START_REF] Uchida | Structural Dependence of Apatite Formation on Titania Gels in a Simulated Body Fluid[END_REF], the bulk anatase is assumed to be more useful for nucleation and growth of hydroxyapatite than the compact bulk rutile, which will have a crucial role in the corrosion resistance by efficiently isolating the substrate from the body fluid. Furthermore, many works confirm that the mixed structure may be more able to promote apatite growth than the single anatase structure [START_REF] Bai | The effect of annealing temperatures on surface properties, hydroxyapatite growth and cell behaviors of TiO2 nanotubes[END_REF][START_REF] Tsuchiya | Hydroxyapatite growth on anodic TiO2 nanotubes[END_REF]. Such fact is associated with the complementary effect of the crystalline structure and surface morphology on the capability of apatite formation.

Another important fact that can be observed in Figure 4 is the absence of any other oxide other than TiO2, despite the presence of high amounts of Nb and Zr. Such absence of other oxides was also observed in several works in the literature, for instance for the Ti20Nb [START_REF] Jin | Fabrication and characterization of anodic oxide nanotubes on TiNb alloys[END_REF][START_REF] Jang | Electrochemical characteristics of nanotubes formed on Ti-Nb alloys[END_REF], Ti13Nb13Zr [START_REF] Stróż | Production, structure and biocompatible properties of oxide nanotubes on Ti13Nb13Zr alloy for medical applications[END_REF], and Ti15Mo [START_REF] Oliveira | Obtaining self-organized nanotubes on biomedical Ti-Mo alloys[END_REF] alloys. Minor or no explanation at all is given for this fact by the authors of these works. Jin et al. [START_REF] Jin | Fabrication and characterization of anodic oxide nanotubes on TiNb alloys[END_REF], stated that the nonexistence of Nb reflections in the as-anodized and annealed nanotubes is assigned to the unnoticeable quantity of Nb in the nanotubes after both anodization and annealing due to the very low diffusion coefficient of Nb compared to that of Ti. Indeed, by observing Table 2, which presents results of EDX analysis, even not considering that part of Nb and Zr is coming from the substrate due beam penetration, one can confirm that amounts of both elements is very low and the presence of either Nb2O5 or ZrO2, if any, is below the detection limit of the XRD equipment.

The above observation confirms that, as stated before, [TiF6] 2-controls the diffusion rate for nanotubes formation, due to the higher solubility of Ti-F or Ti-O-F compounds in aqueous media than Nb-F and Zr-F species. Therefore, the low solubility of Nb and Zr, and their low diffusion coefficient justify the negligible amounts of both elements in the composition of nanotubes or nanopores, and, consequently, small amounts of their oxides.

Figure 5 presents results of potentiodynamic polarizations that display different corrosion potential values for the various conditions, being higher for anodized and thermally treated samples (TiO2 containing samples), meaning that anodization enhances the stability of the material significantly. Also, one observes a large difference in corrosion current densities, which indicates a lower corrosion resistance for polished samples compared to the anodized and crystallized ones. All samples exhibit very stable passivation characteristics up to ~0.9 V vs. SCE, which evidences the absence of pitting corrosion initiation.

In fact, on average, the corrosion current density (icorr) for titanium substrates is ~120 nA cm -2 , while for samples coated with TiO2 nanostructures such current is decreased to about 3 nA cm -2 . From these results, it is noteworthy that icorr for titanium-based surfaces is exceptionally smaller than for other metallic materials used in longstanding implants, for example ranging from ~8 to ~3000 times lower than the icorr for stainless steel 316L in physiological media (~ 1-10 µA cm -2 ) that is extensively employed in surgical implants [START_REF] Chen | Metallic implant biomaterials[END_REF][START_REF] Sinha | Potentiodynamic polarization study of type316L and 316LVM stainless steels for surgical implants in simulated body fluid[END_REF].

In short, it is visible that ultrafine grained samples have better corrosion behavior than coarse-grained ones, considering or not the presence of the nanostructured coating. However, it is still more noticeable that TiO2 nanostructures enhanced the corrosion resistance when compared to the different titanium substrates. Such an improvement is due to the presence of anatase nanostructures and the compact bulk layer of rutile that creates an interfacial oxide layer barrier. Comparing electrical conductivity and structural arrangement one may be advised that the rutile phase is respectively more conductive [START_REF] Tighineanu | Conductivity of anodic TiO2 nanotubes: Influence of annealing conditions[END_REF] and more compact than the anatase [START_REF] Roguska | Improvement of the bio-functional properties of TiO2 nanotubes[END_REF], thus, as also reported in the literature [START_REF] Munirathinam | Role of crystallinity on the nanomechanical and electrochemical properties of TiO2 nanotubes[END_REF], leading to an improvement of corrosion resistance of the interfacial layer. Therefore, regarding the synergy between corrosion and mechanical properties, it appears that, due to the mixed structure, the nanostructured TiO2 coating can satisfy both properties simultaneously. In other words, high corrosion potential (Ecorr), low corrosion current density (icorr), meaning stability in a physiological environment and, as stated in the literature [START_REF] Bai | The effect of annealing temperatures on surface properties, hydroxyapatite growth and cell behaviors of TiO2 nanotubes[END_REF], attractive values of hardness and Young's modulus.

One can confirm the above statements by examining Table 3, which brings results for cp-Ti and Ti-alloys from the literature. The literature is vast in the nanotubes' subject.

However, the considerable majority of articles are related to cp-Ti and classical Ti6Al4V.

Nevertheless, regarding potentiodynamic polarization, since few results from other alloys were published, we decided to use only the more recent information. 3 for different substrates gave the value of about -380 mV. Considering such a value, one may observe that, except for one value for cp-Ti, corrosion potentials for substrates are electrochemically and statistically similar. However, one must bear in mind that Ecorr is related to surface characteristics. In other words, nobler and more stable surfaces in the medium will lead to higher corrosion potentials. Therefore, differences may be associated with the soaking time, different times before soaking, and distinct surface conditioning previous to polarization. Nevertheless, the above observation is not completely clear for anodized samples. As one may observe in Table 3, in anodized surfaces there would be reciprocated impact of the morphology of TiO2 nanostructures (length, diameter, distribution, and so on) and, of course, their crystallinity influencing surface properties. It is also remarkable that the majority of samples coated with a layer of TiO2 nanostructures usually have higher corrosion potentials, compared to the substrate. This reasoning was not observed for Ti-Nb alloys, where Ecorr was much worsened after anodization. However, it is clear that surfaces of the mixed morphology of nanotubes and nanopores for the Ti13Nb13Zr are much nobler than the majority of the alloys.

Calculation of the average value regarding Ecorr presented in Table

Regarding corrosion current (icorr) for substrates, except for unexpected results for the Ti15Mo and Ti6Al7Nb alloys, all other alloys presented icorr in the same order of magnitude, with similar values, being electrochemically equivalents. Although it would be expected a decrease in icorr after anodization, this was not the case for one cp-Ti, Ti-Nb alloys (where there was a considerable increase in icorr), and Ti15Mo alloy. By observing the micrographs in the related articles, it was possible to note that all surfaces of these alloys were very irregular and, in some cases, with voids, which can lead to crevice corrosion. In this way, one could infer that such irregularities could be one of the reasons for such an increase. Nonetheless, the authors of these articles did not give any explanation regarding this fact. Excluding these exceptions, it is noteworthy that nanotubes on the surface of one of the Ti6Al4V alloy, and nanotubes + nanopores on the surface of the Ti13Nb13Zr alloy, led to very small and equivalent corrosion currents when compared to other alloys.

A qualitative analysis of apatite formation performed by FEG-SEM observations and EDS analysis can be seen in Figures S1,S2, and S3 in the supplementary material file. As one may observe, no apatite layer was found on polished surfaces (Fig. S1), indicating that, as already stated in the literature [START_REF] Bauer | Engineering biocompatible implant surfaces: Part I: Materials and surfaces[END_REF], they are not bioactive. Regarding anodized surfaces, Figure S2 shows that there is apatite precipitation on anodized surfaces of the coarse-grained material between the first and seventh day. Meanwhile, there was apatite precipitation on the anodized surface of ultrafine grained material from the first day of immersion in SBF solution. This behavior indicates that the nanostructured oxide layer produced on fine-grained surfaces is more bioactive than on surfaces of the coarse-grained material. The EDX analysis of Figure S3 and the Raman spectroscopy in Figure 6a confirm the presence of apatite from the first day of immersion.

Comparing the micrographs in Figure S2, one may observe globular precipitation of apatite on the first day for the deformed sample, which grows fast up to the seventh day and thickens in the next seven days. However, in the same Figure, it is possible to see that the precipitation of apatite on the surface of the undeformed sample is very thin after seven days, but also thickens with continued immersion in SBF solution.

Table 4 presents results of periods of precipitation of apatite on surfaces coated with nanotubes from the literature comparing with the results of this work for deformed and undeformed samples of the (α'+β)-Ti-13Nb-13Zr alloy coated with nanopores and nanotubes.

As one may observe, the materials coated with nanopores and nanotubes of the present work acquired higher bioactivity than most of the materials coated with nanotubes found in the literature and have similar bioactivity with some of them. As all of the materials possessed mixed anatase and rutile structures, the most probable reasons for this behavior may be related to different diameters of nanotubes (nanopores) and the degree of ordering of TiO2 nanostructures present in each condition used for growing nanotubes. Both reasons are directly linked to the specific surface area that may favor apatite precipitation [START_REF] Li | Synthesis and bioactivity of highly ordered TiO2 nanotube arrays[END_REF].

A quantitative evaluation of apatite formation was carried out by weighing samples before and after immersion in SBF. Values in Figure 6b represent the average of three measurements while error bars consider the standard deviation and measurement uncertainties of the weighing balance (0.2 mg). Figure 6b shows apatite mass gain for annealed TiO2 nanostructures on (α'+β) Ti samples compared to results for a cp α-Ti previously published by our group [START_REF] Hilario | Influence of morphology and crystalline structure of TiO2 nanotubes on their electrochemical properties and apatiteforming ability[END_REF] (alloy number 3 in Table 4), after 1, 7 and 14 days of immersion in SBF. As one may observe, no apatite was deposited on the surface of coarse-grained (α'+β) Ti-alloy during the first day of soaking, while on other samples, the mass gain of apatite increased with the immersion time from the first 24 hours. However, despite continuously growing, one can note that during the second week of immersion the growth rate of apatite was reduced compared to the first week. From the first day immersed in SBF, cp α-Ti samples showed higher bioactivity than samples of (α'+β) Ti-alloys, demonstrating the effect of the different kind of surfaces (only nanotubes and the mix of nanotubes and nanopores). In this case, it is possible to compare bioactivity related to surfaces because the parameters, procedures, and conditions used to produce the nanostructures on the surface of different alloys were precisely the same. Nevertheless, comparing coarse-grained samples, α-Ti is much more bioactive than the coarse-grained (α'+β) Ti-alloy. Comparing (α'+β) Ti samples, it is also noticeable that the smaller the grain size, the more important the mass gain, which also demonstrate the importance of the kind of the surface for bioactivity, and, in this case, the redistribution and sizes of α' and β phases on the sample surface caused by HPT.

Conclusions

The results of the anodization of the Ti13Nb13Zr alloy showed for the first time that there is the formation of two different nanostructures, where nanotubes grow in the α'-phase and nanopores in the β-phase, irrespective to the spatial distribution of both phases.

Even with the different nanostructures formed in α' and β phases after anodization, potentiodynamic polarizations revealed an increase of corrosion resistance for the anodized Ti13Nb13Zr samples when compared to polished ones, without pitting corrosion initiation.

A comparison between anodized (α'+β)-Ti coarse-grained alloy and (α'+β)-Ti ultra-fine grained one indicates that nanostructuring the (α'+β)-Ti alloy by HPT helped to improve the bioactivity when compared with the coarse-grained sample. 4), coarse-grained (α'+β)-Ti-13Nb-13Zr alloy and ultra-fine grained (α'+β)-Ti13Nb13Zr alloy processed by HPT.

Tables' captions

Figure 3

 3 Figure 3 presents in SEM images the morphologies of the top and lateral views of nanostructures after anodization. Such images show self-organized TiO2 nanotubes formed in the α' phase and nanoporous of TiO2 in the β phase (see Fig.3b). Figures3b and 3c indicate

undeformed and deformed samples Table 3 .Table 4 .Figure 1 .

 341 Figures' captions

Figure 2 .

 2 Figure 2. (a) SEM image showing the microstructure of the non-deformed Ti13Nb13Zr alloy. (b) ACOM-TEM analysis, showing a virtual TEM bright field image for the Ti13Nb13Zr alloy after deformation, superimposed to a mapping of high-angle grain boundaries.

Figure 3 .

 3 Figure 3. SEM micrographs of anodized Ti13Nb13Zr alloy. (a, b, c) Initial condition. (a) Top view image, (b) Top view image (higher magnification), (c) Lateral view of the mechanically fractured region. (d, e, f) HPT deformed samples. (d) Top view, (e) Top view image of nanopores, (f) Lateral view of the mechanically fractured region.

Figure 4 .

 4 Figure 4. XRD pattern for the Ti13Nb13Zr alloy in the initial and deformed conditions with (a) as-formed nanotubes and nanopores and (b) annealed TiO2 nanostructures at 550 °C in air.

Figure 5 .

 5 Figure 5. Potentiodynamic polarization curves, recorded in SBF at 37°C.

Figure 6 .

 6 Figure 6. (a) Raman spectroscopy of the surface of the ultrafine-grained samples, showing the presence of hydroxyapatite (Hap) on the surface of nanotubes and nanopores of the alloy (α'+β)-Ti13Nb13Zr. (b) Apatite mass gain of anodized samples when immersed in SBF, as a function of immersion time, comparing coarse-grained cp α-Ti [42] (alloy number 3 in Table

Table 1 .

 1 Summary of some results from the literature using similar conditions to this work

Table 2 .

 2 Results of EDX analysis of anodized and heat treated nanostructures coating

Table 1 .

 1 Summary of some results from the literature using similar conditions to this work

	Nanotubes'dimensions
	(nm)

Table 2 .

 2 Results of EDX analysis of anodized and heat treated nanostructures coating undeformed and deformed samples

		Condition
	Element	Undeformed (at. %)	Deformed (at. %)
	Ti	26.30	27.93
	Zr	2.55	2.65
	Nb	3.14	2.78
	O	68.01	66.64

Table 3 .

 3 Corrosion parameters of annealed nanotubes compared to substrates (NP = nanopores, NT = nanotubes, CG = Coarse grained, UFG = ultrafine grained, GP= Ground and polished)

	Material	Morphology	Condition	Icorr (nA/cm2) Ecorr (mVxSCE)	[ref]
	Ti13Nb13Zr13 CG	GP	Substrate	161	-451
	Ti13Nb13Zr13 UFG			81	-541	This work
	Ti13Nb13Zr13 CG	NP+NT	Annealed at 550 o C	2.5	-3
	Ti13Nb13Zr13 UFG			3.4	-71
	cpTi cpTi	GP NT	Substrate Annealed at 500 o C	100 180	-145 80	[63]
	cpTi cpTi	GP NT	Substrate Annealed at 450 o C	300 200	-523 -389	[64]
	Ti10Nb			200	-367
	Ti20Nb	GP	Substrate	200	-367
	Ti30Nb			350	-337
	Ti40Nb Ti10Nb Ti20Nb	NT	Annealed at 600 o C	100 28,000 15,000	-493 -1323 -1000	[59]
	Ti30Nb			30,000	-1266
	Ti40Nb			30,000	-811
	Ti6Al7Nb Ti6Al7Nb	GP NT	Substrate Annealed at 600 o C	600,000 8000	-218 -317	[65]
	Ti6Al4V Ti6Al4V	GP NT	Substrate Annealed at 500 o C	4334 549	-143 -89	[66]
	Ti6Al4V Ti6Al4V	GP NT	Substrate Annealed at 650 o C	150 2	-250 -125	[67]
	Ti15Mo Ti15Mo	GP NT	Substrate Annealing ?	1100 1500	-397 -206	[14 ]

Acknowledgments

The authors are grateful to the French Committee for the Evaluation of Academic and Scientific Cooperation with Brazil (COFECUB) and the Brazilian CAPES (Process number #88887.137755/2017-00) for funding the bilateral cooperation, and also FAPESP under the grant number for the partial funding. FAPESP Thematic # 2013/05987-8.The authors thank the Laboratory of Structural Characterization (LCE/DEMa/UFSCar) for using their facilities.

AMJJ would like to acknowledge the CNPq (Brazil) under grant #301429/2017-0.

Alloy

Period of precipitation [reference] 13 Ti-6Al-4V Between the 7th and 14th day [START_REF] Sarraf | In vitro bioactivity and corrosion resistance enhancement of Ti-6Al-4V by highly ordered TiO2 nanotube arrays[END_REF]