Fan Sun
email: fan.sun@chimieparistech.psl.eu

Laurence Jordan

Alexandre Da Silva

Frantz Martin

Frédéric Prima

Revisiting the effects of low-concentration hydrogen in NiTi self-expandable stents

Keywords: NiTi self-expandable stent, hydrogen redistribution, heterogeneous deformation, stress-induced martensitic transformation, dislocation slip

Superelastic NiTi is widely used in self-expandable transcatheter devices such as cardiovascular stent, neuro thrombectomy retriever and aortic valve scaffold. Hydrogen at a concentration of CH < 100 wppm is almost always contained in these devices due to the chemical surface treatments during production. The survey of hydrogen concentrations in 9 commercial stent samples show H concentrations values ranging from 6 to 55 wppm. Correspondingly, hydrogen (8 wppm and 40 wppm) is charged to NiTi samples at two thermal-treatment states (500°C-annealed and 800°C-annealed) for comparative studies. The effects of hydrogen are revisited to clarify the hydrogen-induced changes of the operating deformation mechanism and the subsequent reduction of ductility in NiTi of both states. In-situ investigations at mesoscopic length scale (50-500 µm) are performed using electron backscatter diffraction (EBSD) mapping during tensile deformation of the 800°C-annealed samples.

The deformation microstructure and the operating deformation mechanism are further studied at microscopic length scale (< 5 µm) by transmission electron microscopy (TEM) for both 500°C and 800°C annealed samples. The results display that the influence of 8 wppm H on final material properties is insignificant in 500°C and 800°C annealed samples. However, for both, a concentration of 40 wppm H noticeably decreases the martensitic transformation temperature and the macroscopic

Introduction

For decades, nickel-titanium (NiTi) has been used in medical devices because of the large superelastic strain (about 7%) due to a reversible thermoelastic martensitic transformation [START_REF] Morgan | Medical shape memory alloy applications -The market and its products[END_REF][START_REF] Zheng | Metallic Biomaterials: new directions and Technologies[END_REF]. This unique property is widely used in biomedical devices such as self-expandable cardiovascular stents. The deployment of a stent immediately after angioplasty has been shown to significantly decrease the risk of restenosis [START_REF] Duerig | An overview of nitinol medical applications[END_REF]. Currently, the majority of NiTi stents are of the self-expandable type. These devices, usually fabricated into a mesh tubular stent by laser process of a NiTi tube, are annealed at about 500°C into a diameter, which is higher than the targeted artery. The austinite finish (Af) temperature of the shaped stent is usually around 22°C, lower than body temperature (37°C) for superelastic self-expansion. The final NiTi stent needs to be crimped (radially deformed by compression) into a small diameter catheter for transportation in the cardiovascular system. Before crimping, the stent is fully austenitic (B2, CsCl-type crystal structure). After crimping, the effect of this deformation (the maximum local strain can reach ε=0.06 -0.1) causes the transformation of the austenite into martensite (B19', monoclinic structure). After being located at the implantation site in the artery, the NiTi device is pushed out of the catheter and immediately and automatically self-expands by the reversion from B19' (stress-induced martensite, SIM) to its parent B2 phase in order to conform to the walls of the artery vessel with a steady radial force [START_REF] Mani | Coronary stents: A materials perspective[END_REF]. Then, the stent can be subjected to cyclic pulsating loads due to heart beating (typically 4.10 7 cycles/year).

To achieve this functionality (with the suitable subsequent radial force), the NiTi stents are required for partially recrystallized microstructure with a characteristic austenite finish (Af) temperature about 22°C inferior to body temperature (37°C) for stable superelastic recovery and a ductility up to about ε=0.1 in tensile/compression strain for being crimped safely into the catheter.

Conventional defects in stents have been intensively studied and strictly inspected in production, e.g. non-metallic inclusions [START_REF] Sun | On the high sensitivity of corrosion resistance of NiTi stents with respect to inclusions: an experimental evidence[END_REF] and surface defects [START_REF]Standard Practice for Surface Preparation and Marking of Metallic Surgical[END_REF], to ensure the long-term structural integrity and superelastic functionality of every single product. However, during the production stages, hydrogen absorption occurs frequently during acid washing and electropolishing processes [START_REF] Sheriff | Hydrogen effects on nitinol fatigue[END_REF]. This fact may be critical since it has been often reported that the contamination of hydrogen can lead to several detrimental effects such as: (1) suppression of martensitic phase transformation [START_REF] Sarraj | Mechanical behavior of NiTi arc wires under pseudoelastic cycling and cathodically hydrogen charging Mechanical behavior of NiTi arc wires under pseudoelastic cycling and cathodically hydrogen charging[END_REF][START_REF] Runciman | Effects of Hydrogen on the Phases and Transition Temperatures of NiTi[END_REF][START_REF] Letaief | Tensile behaviour of superelastic NiTi alloys charged with hydrogen under applied strain[END_REF][START_REF] Robertson | On the Mechanical Properties and Microstructure of Nitinol for Biomedical Stent Applications[END_REF][START_REF] Asaoka | Effect of Hydrogen on Martensite Phase Deformation in Ti-Ni Shape Memory Alloy[END_REF], [START_REF] Zheng | Metallic Biomaterials: new directions and Technologies[END_REF] noticeable increase of the transformation stress threshold [START_REF] Sheriff | Hydrogen effects on nitinol fatigue[END_REF][START_REF] Sarraj | Mechanical behavior of NiTi arc wires under pseudoelastic cycling and cathodically hydrogen charging Mechanical behavior of NiTi arc wires under pseudoelastic cycling and cathodically hydrogen charging[END_REF][START_REF] Letaief | Tensile behaviour of superelastic NiTi alloys charged with hydrogen under applied strain[END_REF][START_REF] Robertson | On the Mechanical Properties and Microstructure of Nitinol for Biomedical Stent Applications[END_REF][START_REF] Rokbani | Hydrogen effects on Ni-Ti fatigue performance by self -heating method[END_REF][START_REF] Gamaoun | Effect of hydrogen on the tensile strength of aged Ni -Ti superelastic alloy[END_REF] and (3) embrittlement [START_REF] Sheriff | Hydrogen effects on nitinol fatigue[END_REF][START_REF] Letaief | Tensile behaviour of superelastic NiTi alloys charged with hydrogen under applied strain[END_REF][START_REF] Rokbani | Hydrogen effects on Ni-Ti fatigue performance by self -heating method[END_REF][START_REF] Yokoyama | Hydrogen embrittlement of Ni-Ti superelastic alloy aged at room temperature after hydrogen charging[END_REF][START_REF] Yokoyama | Hydrogen embrittlement behavior induced by dynamic martensite transformation of Ni-Ti superelastic alloy[END_REF][START_REF] Ogawa | Effect of aging at room temperature on hydrogen embrittlement behavior of Ni-Ti superelastic alloy immersed in acidic fluoride solution[END_REF][START_REF] Asaoka | Hydrogen Embrittlement of Nickel-Titanium Alloy in Biological Environment[END_REF]. Recently, Duerig et al. [START_REF] Duerig | A Practitioner's Perspective of Hydrogen in Ni-Ti Alloys[END_REF] have pointed out several aspects of hydrogen in Ni-Ti and highlighted both localized and highly dynamic effects. It was shown that a content as low as 100 wppm hydrogen at fully homogenized distribution may completely suppress the austenitemartensite transformation. It was also reported that the ductility was drastically affected in a time-dependent manner based on the diffusion of hydrogen: the tensile samples are fully ductile after charging, but become fully brittle hours later.

In the industrial practice, at "high concentrations" (CH > 100 wppm), the hydrogen contaminated stents can possibly be ruled out during quality inspection by Af temperature measurements or mechanical compression tests. In contrast, for "low hydrogen concentrations" (CH < 100 wppm), the hydrogen-containing stents may pass the quality inspection to be finally crimped into the catheter for implantation. According to the conclusion proposed recently by Duerig et al. [START_REF] Duerig | A Practitioner's Perspective of Hydrogen in Ni-Ti Alloys[END_REF], the safety limit of 50 wppm hydrogen in average, required by ASTM F2063 [START_REF]Standard Specification for Wrought Nickel-Titanium Shape Memory Alloys for Medical Devices and Surgical Implants[END_REF] standard for wrought NiTi material, was still considered to be sufficient for the finished NiTi material and devices under condition that the hydrogen was homogeneously distributed. The excessive hydrogen, generally or locally, may lead to the formation of micro-cracks and to struts fracture during crimping. These structural defects can result in implantation failure due to the difficulty of releasing the stent from the catheter and the incomplete self-expansion. High-risk surgery is usually required to remove the failed stent, but sometimes leads to the severe adverse events to the patients whose conditions are not suitable for invasive operation.

The objectives of this study are to study the effects of hydrogen in austenitic NiTi exhibiting superelasticity for different hydrogen concentrations. The concentrations (8 and 40 wppm) are selected according to the actual hydrogen level in the current commercial NiTi stents. In order to study the H-charged samples by in-situ and ex-situ methods, two different starting microstructures (partially and fully recrystallized) are submitted to microstructural and mechanical investigations.

Towards the understanding of mechanisms, the effects of hydrogen concentration and redistribution are revisited in the NiTi materials for the implication on NiTi stent safety with respect to the present acceptance of hydrogen concentration.

Materials and methods

Samples realization and specification

For this study, samples are prepared by using commercial ELI grade (Extra Low Inclusion grade) NiTi plate (0.5 mm thickness) from Johnson Matthey®. The raw NiTi plate is received at as-rolled state. All the NiTi samples are cut from the same NiTi plate and annealed to obtain partially and fully recrystallized states under vacuum (2x10 -5 mbar). The partial recrystallization is conducted at 500°C for 6 minutes to obtain the same starting microstructure and austinite finishing temperature (Af) as commercial stents, hereafter called "500°C-annealed". The full recrystallization is obtained after 10 minutes spent at 800°C (i) to desorb the hydrogen potentially in the raw NiTi plate and (ii) to obtain an equiaxial grain structure with average grain-size of 25 µm for in-situ electron-backscattered diffraction (EBSD) studies. The fully recrystallized state will thereafter referred to as "800°C-annealed".

Both annealing treatments are finished by water quenching.

For hydrogen charging, two concentrations are selected for being representative of the commercial products on the market (low range and high range). To evaluate the H compositional range into these products, a complete thermal desorption spectroscopy (TDS) investigation has been conducted on nine different self-expandable NiTi peripheral stents from eight different trademarks (kept anonymous in the frame of this paper). These stent products have been stored at room temperatures for serval months since the date of production. The observed results display a hydrogen concentration range spreading out from 6 to 55 wppm (Table 1). It can be noticed that 8 of the 9 stents exhibit hydrogen concentrations below the 50 wppm ASTM F2063 standard acceptance; only one is slightly over. The TDS spectra obtained from the stent samples are similar in form. The spectrum of Stent 4 analyzed by Gaussian peak deconvolution is shown in Section 3.4 to show the hydrogen distribution.

Characterization methods

The hydrogen concentrations of the NiTi samples are accurately quantified by using the melting extraction technique on a Bruker™ Galileo G8 apparatus. The stent samples and some of the NiTi samples are analyzed by thermal desorption spectrometry (TDS) under an initial vacuum of 5.10 is used to prepare the TEM thin foils.

Results

DSC measurements and tensile tests

The martensitic transformation process is characterized by DSC measurements on the two 500°C-annealed (partially recrystallized) samples (8 and 40 wppm H) prior to deformation. The corresponding DSC curves are shown in Figure 1a. The phase transformations and transformation temperatures are summarized in Table 2. Both of the 500°C-annealed samples exhibit two-step martensitic transformation (B2-R-B19'). The R-phase, trigonal structure, is presented as a competing martensite phase to B19'. In the 500°C-annealed NiTi from as-rolled state, the high density of rearranged dislocations and precipitates present in the structure (shown in section 3.3) tend to suppress B19' transformation with respect with R-phase transformation [START_REF] Otsuka | Physical metallurgy of Ti-Ni-based shape memory alloys[END_REF]. In presence of hydrogen, the transformation peaks of both transformations reduce in area when the hydrogen concentration raises from 8 wppm to 40 wppm. Moreover, the Ms (R-B19') temperature shows a slight shift towards lower temperatures when hydrogen increases to 40

wppm. Meanwhile, the Af temperature (22°C) does not seem to be affected for the two samples. The phenomena underline a suppression effect of hydrogen on both transformations, but the B19' transformation seems more sensitive to the hydrogen concentration than the R-phase transformation.

Table 2: Results of DSC measurements for each microstructure and H content.

Sample Hydrogen concentration (wppm)

Transformation sequence Ms (B19') Af

500°C-annealed The DSC results are summarized in Table 2. The martensitic transformation is the classic one-step process after full recrystallization between austenite (A) and martensite (M) [START_REF] Asaoka | Hydrogen Embrittlement of Nickel-Titanium Alloy in Biological Environment[END_REF]. The three samples are fully austenitic at RT (22°C), the stress-induced martensitic (SIM) transformation is triggered by tensile stress.

8 ± 2 B2-R-B19' -40°C ≈22°C 40 ± 5 B2-R-B19' -47°C ≈22°C 800°C-annealed Before H charge B2-B19' -15°C ≈4°C 8 ± 4 B2-B19' -15°C ≈4°C
The single-loading curves of NiTi samples are shown in Fig. 2b, exhibiting typical "double-yielding" curves with a plateau in the elastic regime due to stress-induced martensitic transformation [START_REF] Otsuka | Physical metallurgy of Ti-Ni-based shape memory alloys[END_REF][START_REF] Otsuka | Shape Memory Materials[END_REF].

The annealed and 8 wppm H samples exhibit almost identical characteristics regarding the stress-strain curves. The 40 wppm H sample exhibits a higher SIM plateau at about 420 MPa, (+23.5%) than the other two samples and a significant loss of ductility (fracture at ε=0.1).

Consistently with single-loading tests, the cyclic-loading results show similar behaviors for the 800°C-annealed and 8 wppm H samples and drastic ductility reduction in the 40 wppm H sample (Fig. 2c and Table 2). When comparing the behaviors between 500°C-annealed and 800°C-annaled samples, it can be noticed that the superelastic behavior of NiTi can be very dependent on microstructure and temperature [START_REF] Chen | Tensile Deformation of Superelastic NiTi Wires in Wide Temperature and Microstructure Ranges[END_REF].

In-situ EBSD characterizations on the 800°C-annealed samples

In order to further investigate the hydrogen influence on the reversible martensitic transformation, in-situ SEM/EBSD observations under tensile strain are performed to compare the microstructural evolutions during tensile loading. The 800°C-annealed samples allow using EBSD to clarify the fundamentals about the hydrogen influences to the martensitic phase transformation in B2 grains without perturbations of dislocation or nanoprecipitates. However, the EBSD technique is limited on the 500°C-annealed sample due to the sub-micrometric subgrain structure and high density of dislocations and nanoprecipitates (studied by TEM in Section 3.3). The Force-Elongation curves recorded during the in-SEM traction (Fig. 3a and4a) exhibit the same shape and ductility as observed in the tensile tests (Fig. 1 and2). The image of martensite band front (MBF) is captured in Fig. 3b, when it is propagating over the observation surface at ε=0.035. The angle between MBF and loading axis is close to 55°, consistent with previous studies [START_REF] Sedmák | Grain-resolved analysis of localized deformation in nickel-titanium wire under tensile load[END_REF]. Fig. 3c shows a series of micrographs of the surface relief changes at ε=0.06, 0.1 and 0.25 after the MBF has passed over the observation surface.

The fresh fracture surface of the 8 wppm H sample shows dimple-like morphologies, a few micrometers in diameter, flat and uniformly distributed (Fig. 3d). The homogeneous dimple-like features suggest a ductile-type fracture of the material, in agreement with its good ductility. Similar to the "macroscopic" tensile test results (Fig. 2b), the elongation to fracture of the 40 wppm H sample as observed in-situ in the SEM set-up is also drastically reduced to ε=0.1 (Fig. 4a). No clear MB can be seen activated near the observation zone (shown in Fig. 4b) at the very beginning of the deformation (ε=0.015). Three SEI micrographs in Fig. 4c In-situ EBSD mapping is performed in the ε=0-0.1 range to monitor the overall surface fraction of martensite as a function of tensile strain, and to trace the local deformation mechanism on loading.

Before tensile loading, the 8 wppm H sample presents a fully recrystallized microstructure of B2 austenite (Fig. 5a). As shown by IQ+phase EBSD maps, in Fig. 5b, plate-like B19' martensite is stress-induced upon loading in the B2 parent grains along with the MBF propagation. The SIM growth is featured by progressive multiplication and thickening of martensite plates (Fig. 5b). It should be noticed that the propagation of the martensite band seems reluctant when comparing to the classic propagation manner, in which the martensitic transformation stops after the MBF passes over. The phenomenon could probably relate to the reduced gauge size of in-situ tensile samples. Or another hypothesis is that the hydrogen could promote the localization of deformation in MBs but suppress the macroscopic deformation localization. At ε=0.1, the EBSD mapping (Fig. 5c) confirms that most of the B2 grains have transformed to martensite. The surface fraction of the non-transformed B2 phase is about 3.2%. Attention is paid to the formation of {114}B2 type B2 twinning from {201}B19' type deformation twin precursor as accommodation mechanism to the increasing local strain [START_REF] Heller | Beyond the strain recoverability of martensitic transformation in NiTi[END_REF][START_REF] Chen | Recoverability of large strains and deformation twinning in martensite during tensile deformation of NiTi shape memory alloy polycrystals[END_REF][START_REF] Polatidis | Deformation mechanisms in a superelastic NiTi alloy: An in-situ high resolution digital image correlation study[END_REF].

However, the {114}B2 twin is not observed during the in-situ loading process, neither in 8 wppm H nor in 40 wppm H 800°C-annealed samples. In-situ EBSD characterizations are similarly performed on the 40 wppm H samples (see micrographs in Figure 6). The initial microstructure appears to be similar to the 8 wppm H sample (Fig. 6a), i.e. fully austenitic and recrystallized. First SIM plates form around ε=0.015 at the observation zone (Fig. 4b and6b). When comparing to 8 wppm H sample, clear differences in microstructural evolutions start to be obvious when ε reaches 0.06. At mesoscopic scale, the stress-induced transformation seems concentrated in band-like area, forming a MB (shown in Fig. 6c). Inside and surrounding the MB, the martensitic transformation is inhibited in many B2 regions, like "islands" (Fig. 6b). One of the islands is observed about its formation (marked by the yellow dashed line in Fig. 6b), during which martensite plates are reluctant to grow towards the core of the island. The core is a triple-junction of three neighboring grains. In the view of a larger area (Fig. 4c), the total area fraction of the non-transformed B2 "islands" is 14.8% before the sample breaks in the test (ε=0.1). To summarize observation made on the 800°C-annealed samples, the area fractions of non-transformed B2 phase versus tensile strain (Fig. 7a) can be compared between the two H-charged samples. The minimum fraction of non-transformed B2 phase is achieved at εt=0.05-0.06 in the 40 wppm H sample which is earlier than 8 wppm H sample (after εt=0.08) (Fig. 7a). The slope of the stress-strain curve (dσ/dε) is plotted in range εt=0.01-0.1, including the SIM plateau and the superelastic-plastic transition (Fig. 7b). It can be seen that the two slopes are almost identical before

TEM observations on the 800°C and 500°C annealed samples

In order to investigate the operating deformation mechanism and the non-transformed B2 islands, TEM observations are performed on the 800°C-annealed, hydrogen-charged, samples after tensile deformation to ε=0.1 (Fig. 8). [111]B2 is shown in Fig. 8f with the SA diffraction pattern in the inset. Since the two zone-axes are not exactly parallel, the corresponding stereo projection (Fig. 8g) is made to reconstruct the OR. The determined OR is close to (10-1)B19' // (21-1)B2 and (010)B19' // (011)B2, consistent with the classic ones [START_REF] Otsuka | Physical metallurgy of Ti-Ni-based shape memory alloys[END_REF].

In the 40 wppm H 800°C-annealed sample, a non-transformed B2 island is found in the TEM foil (Fig. 9a), marked by yellow dashed line. When compared to the 8 wppm H samples, much higher density of <001> screw dislocations can be observed with considerably reduced inter-spacing between adjacent dislocations. Three different Burgers vectors, noted as b1=a[001], b2=a[100] and b3=a [010] in Fig. 9, are identified. Among them, the b1 dislocations are in much higher quantity than b2 and b3, respectively. Owing to the local high dislocation density, it suggests that the SIM mechanism may not be the dominant deformation mechanism operating in the non-transformed B2 island. It has been reported that the hybrid slip/transformation activity in NiTi undergoing cyclic deformation [START_REF] Dong | On interfacial energy of macroscopic domains in polycrystalline NiTi shape memory alloys[END_REF],

results in high density of dislocations due to repeating transformations between B19' and B2.

However, the current sample is not cyclically deformed. The activation of dislocation slip could probably relate to the presence of hydrogen. The TEM observations are also performed on the 500°C-annealed sample, to study the starting microstructure and operating deformation mechanism, before and after single tensile deformation.

(e) and (h), zone of interest (ZOI) is marked by black dash

square at which the zone axis is used to establish the dislocation contrast.

TDS results and hydrogen redistribution

The TDS thermogram of hydrogen in the stent sample (Stent #4 among the commercial stents mentioned in Table 1) is shown in Fig. 11a. Gaussian peak fitting is used to deconvolute the multiple desorption peaks [START_REF] Lisowski | Thermal desorption of deuterium from modified carbon nanotubes and its correlation to the microstructure[END_REF]. The thermogram can be fitted by three Gaussian peaks centered at 617°C, 657°C and 733°C. The peaks at different desorption temperatures are signatures of the hydrogen interaction sites with different binding energies with the matrix. Because the stent sample has been stored at room temperature (RT) for several months before TDS measurement, the diffusion from surface to the core should have been already completed according to the experimental results obtained by Duerig et al. [START_REF] Duerig | A Practitioner's Perspective of Hydrogen in Ni-Ti Alloys[END_REF]. Therefore, a time-dependent redistribution of hydrogen in the NiTi microstructure may be the cause of three desorption peaks. The hydrogen redistribution phenomenon is also observed in the 40 wppm H 800°C-annealed sample during RT aging. Fig. 11b shows the TDS thermogram and Gaussian peak fitting of the 40 wppm H sample after 24 hours RT aging. The 24 hours aging allows about 0.25 mm diffusion depth of hydrogen from surface towards core [START_REF] Duerig | A Practitioner's Perspective of Hydrogen in Ni-Ti Alloys[END_REF], which corresponds to a "core charging" condition from both surfaces in the 0.5mm-thick plate used in this study. The 24 hours-aged sample displays a symmetrical single peak centered at 29 652°C, that can be fitted by a single Gaussian type peak (Fig. 11b). It indicates that hydrogen atoms in the material, at this stage, reside in sites with similar bonding-energy. These sites have been proposed to be the pseudo octahedral sites of B2 lattice [START_REF] Duerig | A Practitioner's Perspective of Hydrogen in Ni-Ti Alloys[END_REF]. After aging at room temperature for 192 hours, the TDS thermogram presents a peak partially shifted towards higher temperatures. The thermogram can be well fitted by three Gaussian type peaks centered at 640°C, 700°C and 754°C.

Among them, the main peak is the one centered at 640°C, almost at the same temperature as the peak in the 24 hours-aged sample: it represents 76.4% of the total desorbed H2. The other two peaks, at higher temperatures of 700°C and 754°C, represent 15.7% and 7.9% of the total H2 desorption, respectively. This observed evolution of the TDS thermograms with aging suggests an isothermal hydrogen redistribution can occur in fully recrystallized NiTi at room temperature. Accordingly, it would be reasonable to assume that such room temperature redistribution could also happen in the 500°C-annealed stent. However, possible redistribution sites should be very complex due to the microstructure complexity (Fig. 10a) and will be the subject of further works.

Discussion

Hydrogen effects on the Ms temperature and SIM stress

The hydrogen effects on the martensitic transformation temperatures have been widely studied [START_REF] Sheriff | Hydrogen effects on nitinol fatigue[END_REF][START_REF] Tomita | Hydrogen thermal desorption behavior of Ni-Ti superelastic alloy subjected to tensile deformation after hydrogen charging[END_REF][START_REF] Adachi | Effect of Hydrogen on the Shape Memory Effect and Transformation Behavior of Ti-Ni Alloy[END_REF][START_REF] Mao | In situ EBSD investigations of the asymmetric stress-induced martensitic transformation in TiNi shape memory alloys under bending[END_REF] and it was previously reported with consistency that the hydrogen content notably reduces the Ms temperature of NiTi. In the 800°C-annealed samples, the Ms temperature of 40 wppm H sample is 10.6°C lower than for 8 wppm H sample, directly impacting the stability level of austenite in the sample with high hydrogen content. Logically this fact may induce a higher onset stress with respect to the martensitic plateau. Correspondingly, the SIM stress of the 40 wppm H sample is increased of about 25% when compared to the annealed sample (no detectable H) and to the 8 wppm H sample. Very similar results have been observed in literature [START_REF] Sarraj | Mechanical behavior of NiTi arc wires under pseudoelastic cycling and cathodically hydrogen charging Mechanical behavior of NiTi arc wires under pseudoelastic cycling and cathodically hydrogen charging[END_REF], which reports an increase of approximately 60 MPa when comparing a 50 wppm H sample to a sample with 9 wppm H. In the 500°C-annealed samples, a similar increase of the stress plateau and decrease of Ms (R-B19') are observed (Fig. 1). These results corroborate previous findings [START_REF] Asaoka | Effect of Hydrogen on Martensite Phase Deformation in Ti-Ni Shape Memory Alloy[END_REF][START_REF] Duerig | A Practitioner's Perspective of Hydrogen in Ni-Ti Alloys[END_REF] showing that hydrogen addition leads to an increase of the austenite phase chemical stability. As a matter of fact, the hydrogen solute acts as B2 stabilizer, hence delays the martensitic transformation during both thermal cooling process and mechanical deformation.

Hydrogen redistribution

In NiTi, it has been reported that hydrogen atoms are mobile (diffusion coefficient = 3.6 × 10 -15 m 2 /s at 25°C [START_REF] Schmidt | Hydrogen solubility and diffusion in the shape-memory alloy NiTi[END_REF] at concentrations in the 50-300 wppm range [START_REF] Moitra | The location of atomic hydrogen in NiTi alloy: A first principles study[END_REF]. The TDS results on the room temperature-aged 40 wppm H sample (Fig. 11b and11c) show a transition from a single site type distribution to a peak separation after 192 hours aging. The increase of the desorption temperature indicates a reduction of the free energy of hydrogen which are trapped in the lattice defects, providing extra binding energy. In the fully recrystallized sample with single B2 phase, the potential trapping sites could probably be grain boundaries. Based on the in situ EBSD measurements, the GB characteristics can be obtained in the 800°C-annealed sample (40 wppm H). By assuming that GBs can act as localized trapping sites for hydrogen [START_REF] Chen | Observation of hydrogen trapping at dislocations, grain boundaries, and precipitates[END_REF][START_REF] Aucouturier | Grain boundary segregations and hydrogen embrittlement[END_REF], the GB map (Fig. 11b) is colored for two groups of grain boundaries according to their interfacial energies Egb [START_REF] Yazdandoost | Tilt grain boundaries energy and structure in NiTi alloys[END_REF] as: Random GBs with Egb > 600 mJ/m 2 (blue) and Low energy GBs with Egb = 200-500 mJ/m 2 (red). Among the low energy ones, the energies of Σ11(332), Σ13(510) and Σ13(320) coincidence site lattice (CSL) GBs are close to 500 mJ/m 2 . The lowest energy GBs are Σ9(114), Σ3(112) and low angle GBs at 200-300 mJ/m 2 . Meanwhile, the map of non-transformed islands (Fig. 12b) is shown side by side to the GB energy distribution map. The IPF map is presented along the tensile test axis (RD) to illustrate the grain orientations in the islands in addition to Fig. 6c (along ND). It can be noticed that every B2 island is formed around GBs and triple junctions (Fig. 6 Despite this, it is known that the hydrogen trapping at grain boundaries and dislocations eventually leads to the decrease of overall ductility in metallic alloys [START_REF] Moitra | The location of atomic hydrogen in NiTi alloy: A first principles study[END_REF]. When considering the 500°C-annealled NiTi, the interpretation of TDS thermogram (Fig. 11a) by hydrogen trapping would be complicated due to the complexity of the partially recrystallized structure. The structural defects, such as sub-GBs, dislocations, Ti3Ni4 precipitates, can serve as potential trapping sites as reported in the steels (ref 28). Moreover, the situation can be more complicated since the subgrains and dislocations/Ti3Ni4 precipitates are distributed heterogeneously in distinct microstructure zones (Zone A and Zone B shown in Fig. 10). Furthermore, by considering the time-dependent nature of the trapping process, an "incubation" period exists, during which the hydrogen diffuses firstly from surface to the core then redistributes in the microstructure. The phenomena of progressive loss of ductility after hydrogen-charging has been reported [START_REF] Duerig | A Practitioner's Perspective of Hydrogen in Ni-Ti Alloys[END_REF], confirming the incubation period of the H impact to the mechanical properties of NiTi. Such incubation further increases the danger of H uptake/amount in stents, because the quality inspection is usually performed soon after the stent surface treatment.

Change of the local deformation mechanism in 800°C-annealed NiTi

The results of in-situ and ex-situ observations of the operating deformation mechanism indicate the tendency of the local promotion of dislocation slip in the 40 wppm hydrogen-charged samples, in both partially and fully recrystallized microstructures. Taking the fully recrystallized sample for analysis, it is pointed out that the slope of the stress-strain curve beyond the plateau (Fig. 7b) is higher (+3 GPa) and the plastic deformation of martensite starts at higher stress in 40 wppm H sample. This suggests that the plastic deformation of the stress induced martensite phase could be partially suppressed at both mesoscopic and microscopic scales. The contrast of the slope is thought to originate from the differences of the microstructural evolution. That could be due to the change of local deformation mechanism at mesoscopic scale. The B2 matrix almost fully transforms into martensitic phase B19' at the end of superelastic regime in the 8 wppm H sample after the propagation of MBF. In contrast, the SIM in 40 wppm H sample seems concentrated in the MB, the propagation of which seems interrupted until fracture (ε = 0.1). Such mesoscale localization of strain in annealed NiTi has been reported [START_REF] Chen | Recoverability of large strains and deformation twinning in martensite during tensile deformation of NiTi shape memory alloy polycrystals[END_REF] but the presence of hydrogen seems to affect it by the formation of non-transformed B2 islands. These B2 islands form in 40 wppm H sample both inside and outside the mesoscopic band microstructure leading to a mixture of 85% martensite and 15% non-transformed B2 when passing to the plastic regime. From microscopic point of view, the dislocation slip is promoted in the B2 islands. By gathering all the observations, there seems to be a two-fold effect of the absorbed hydrogen on the localization of macroscopic deformation as well as on the localization of deformation on individual grains. Nevertheless, it is difficult to determine the relationship between the B2 islands and the interruption of MB propagation since EBSD results are not sufficient to measure the localization of mesoscale deformation. What can be suggested is that, in the 40 wppm H sample (800°C-annealed), the heterogeneous deformation happens during the superelastic-plastic transition as a result of deformation contrasts at MB and its surrounding, and at slip dominant non-transformed B2 islands and their environs. Such heterogeneous deformation is usually associated with strong macroscopic strengthening effect due to microscopic stress concentration [START_REF] Zhu | Perspective on hetero-deformation induced (HDI) hardening and back stress[END_REF]. The strengthening effect is observed in the 40 wppm H sample (800°C-annealed) that the plateau stress increases more rapidly than that in 8 wppm H sample (800°C-annealed) (Fig. 7b). About the hydrogen's effect to the localization of dislocation slip, a theory called Hydrogen-Enhanced Localized Plasticity (HELP) is worth mentioned to explain enhanced tendency for localization of dislocation slip in the non-transformed B2 islands due to the absorbed hydrogen, similar to many hydrogen-charged alloys near stress concentrations areas, e.g. crack tips [START_REF] Birnbaum | Hydrogen-enhanced localized plasticity-a mechanism for hydrogen-related fracture[END_REF].

Partially recrystallized NiTi used in stent application and the effect of hydrogen

The microstructure used in NiTi for stent application is partially recrystallized from as-drawn raw material. The structural defects in the raw material, i.e. high density of dislocations, are partially recovered through annihilation and rearrangement to form subgrain structure during the thermomechanical treatments. The treatment is performed at intermediate temperature, much inferior to the recrystallization temperature, for limited duration and under geometrical constrain.

The treatment is called "shape-setting" to fix the stent geometry via the subgrain formation and to adjust the Af temperature. The mechanical properties of NiTi in wide microstructure ranges have been well-studied by Y. Chen et al. [START_REF] Chen | Tensile Deformation of Superelastic NiTi Wires in Wide Temperature and Microstructure Ranges[END_REF], showing that the stability of the cyclic superelastic deformation is of high sensitivity to the microstructure and dislocation activities. Moreover, the deformation is concentrated on the mesoscopic scale since the profile of the stent strut is in the dimensions of 100µm. Therefore, the subtle balance between shape-setting and Af adjustment is of great importance for performance of the stent in the scope of complete self-expansion and suitable radial force. Thus, the presence of hydrogen in such highly complex microstructure (Fig. 10) adds extra variables which drastically increase the difficulty of accurate control of the deformation behaviors of stent.

In this study, due to the difficulty of applying in-situ EBSD characterization on the 500°C-annealed NiTi, much less information of the deformation mechanism can be obtained to directly evidence the mesoscopic mechanism change by hydrogen. The changes in microstructures, i.e. the fracture surface (Fig. 1b and1c) and local dislocation density (Fig. to the same ductility loss (Fig. 1b), similar suppression of martensite transformation (Fig. 1a) and promotion of dislocation activity in specific region of the microstructure (Fig. 10c) as observed in 800°C-annealed NiTi. The fact raises the concern about the structural damage of hydrogen containing NiTi stent undergoing the crimping process.

Conclusion

The present work revisits the effects of low-concentration hydrogen on mechanical properties of NiTi materials, with regards to hydrogen contents found in commercial stents (6 to 55 wppm). Two

From this set of

  values, two different hydrogen concentrations, 8 wppm (representative of the low concentrations range) and 40 wppm (representative of the high concentrations range) were selected and used as target for the present study. The hydrogen-containing samples are prepared by immersion in the modified Kroll's reagent (HF: HNO3 = 1:3) at 60°C for different durations to obtain the target concentrations. In order to quantify the hydrogen concentration, 5 mm length specimens are cut from each sample. The hydrogen charged samples are aged at room temperature (RT) for 192 hours before experiments. For comparison, a 40 wppm H sample is also analyzed by TDS after only 24 hours RT aging to verify the hydrogen redistribution.

Figure 1 .

 1 Figure 1. Thermal and mechanical characterizations of 500°C-annealed NiTi after hydrogen-charging. (a) DSC curves of 8 wppm H sample (black) and 40 wppm H sample (red), a two-step martensitic

Figures

  Figures 1b and 1c display the cyclic tensile curves of the two hydrogenated 500°C-annealed samples.Perfect superelastic recoveries are achieved by both samples up to ε=0.06, which is the range for stent service in the artery. However, a drastic ductility loss for the 40 wppm H sample (fracture at about 0.07) is noticed when comparing to the 8 wppm H one (fracture at over 0.3). When comparing the results to the maximum crimping strain (ε=0.06-0.1, marked in yellow in Fig.1b and 1c), a risk of fracture can be a concern for the 40 wppm H sample. The micrographs of fracture surface are shown as insets near the two tensile curves. Features of flat dimples and facets can be seen on the 8 wppm H sample, whereas morphologies like blocks and pits can be noticed in addition to the flat dimples on the 40 wppm H sample. The blocks and pits are of about 10 µm diameter in sizes. The change of the fracture surface may result from a change in deformation mechanisms.

Figure 2 .

 2 Figure 2. Thermal and mechanical characterizations of 800°C-annealed NiTi after hydrogen-charging. (a) DSC curves of as-800°C-annealed sample (red), 8 wppm H sample (blue) and 40 wppm H sample (green) following the cooling-heating passage at 5°C/min, a single step martensitic transformation (B2-B19') is observed in each sample; (b) Curves of uniaxial single loading-to-fracture tensile test of the as-800°C-annealed sample (red), 8 wppm H sample (blue) and 40 wppm H sample (green) at Af+15°C; (c) Curves of uniaxial cyclic loading-unloading tensile test of the three samples under the same conditions; the unloading cycles are performed at ε=0.02, 0.04, 0.06, 0.08… until fracture. The deformation rate is 10 -3 s -1 in both single loading and cyclic loading-unloading tests. The fracture points on the tensile curves are marked by the black triangles.

  at ε=0.06, 0.08 and 0.1 present the surface relief change during the in-SEM loading at the same observation zone as shown in Fig.4b. Fig.4dpresents the fresh fracture surface exhibiting brittleness features, such as ridges and facets. The size and distribution of the dimples are non-uniform and in agreement with previous studies of the fractographies observed after hydrogen charging[START_REF] Runciman | Effects of Hydrogen on the Phases and Transition Temperatures of NiTi[END_REF][START_REF] Asaoka | Effect of Hydrogen on Martensite Phase Deformation in Ti-Ni Shape Memory Alloy[END_REF][START_REF] Asaoka | Hydrogen Embrittlement of Nickel-Titanium Alloy in Biological Environment[END_REF]. The fracture surface change, between 8 wppm H and 40 wppm H, may again suggest a change of deformation mechanisms in the 800°C-annealed samples, depending on the hydrogen amount.

Figure 3 :

 3 Figure 3: Force-elongation curve and micrographs of 8 wppm H sample (800°C-annealed) during in-situ traction in SEM. (a) Sample schema and in-situ force-elongation curve of the 8 wppm H sample; (b) Secondary electron imaging (SEI) micrograph (200x) of the martensite band front (MBF) crossing the observation area at ε = 0.035. The MBF line is marked by red dash line, the propagation direction of the MBF is marked by red arrow, the angle between MBF line and loading axis is about 55°; (c) A series of SEI micrographs (500x) at the same observation area captured at ε = 0.06, 0.1, 0.25; (d) The SEI micrographs of the fresh fracture surface, on which shallow dimples are observed. The fracture points on the tensile curve are marked by the black triangle. The holding points for EBSD mapping are marked by vertical arrows.

Figure 4 :

 4 Figure 4: Force-elongation curve and micrographs of 40 wppm H sample (800°C-annealed) during in-situ traction in SEM. (a) Sample schema and in-situ force-elongation curve of the 40 wppm H sample; (b) Secondary electron imaging (SEI) micrograph (200x) of the observation area, no clear martensite band front (MBF) can be seen; (c) A series of SEI micrographs (500x) at the same observation area are captured at ε = 0.06, 0.08, 0.1 before fracture; (d) The SEI micrographs of the fresh fracture surface, on which flat dimples in mixture with ridges and facets are observed. The fracture points on the tensile curve are marked by the black triangle. The holding points for EBSD

Figure 5 :

 5 Figure 5: EBSD maps of 8 wppm H sample (800°C-annealed) recorded during in-situ traction process. The EBSD scans are performed by holding the deformation strains without unloading. The deformation microstructures are traced by five EBSD scans at ε = 0, 0.015, 0.025, 0.035 and 0.1, respectively. (a) Inversed pole figure (IPF) map of B2 phase before deformation at the 500x zone shown in figure 3b; (b) IPF+IQ maps of B2 phase and the corresponding phase maps of B19' martensite at ε = 0.015, 0.025, 0.035 at 1000x magnification at the zone marked by white square in (a). The micrographs show the progressive growth of B19' martensites in B2 grains when martensite band front (MBF) propagates across the observation zone (the same 200x zone shown in figure 3b); (c) the IPF+IQ map of B2 phase at ε = 0.1 at the same observation zone (500x). The area fraction of

Figure 6 :

 6 Figure 6: EBSD maps of 40 wppm H sample (800°C-annealed) recorded during in-situ traction process. The EBSD scans are performed by holding the deformation strains without unloading. The deformation microstructures are traced by five EBSD scans at ε = 0, 0.015, 0.06, 0.08 and 0.1 before fracture, respectively. (a) Inversed pole figure (IPF) map of B2 phase before deformation at the 500x zone shown in figure 5b; (b) IPF+IQ maps of B2 phase and the corresponding phase maps of B19' martensite at ε = 0.015, 0.6, 0.8 at 1000x magnification at the zone marked by white square in (a). The micrographs show the progressive formation of a non-transformed B2 island (marked by yellow dash line); (c) the IPF+IQ map of B2 phase at ε = 0.1 at the same observation zone (500x). A MB can

  the minimum fraction of non-transformed B2 phase is reached in 40 wppm H sample. Afterwards, the slope of the 40 wppm H sample increases rapidly to 3 GPa higher than that of 8 wppm H sample. The 3 GPa difference is kept until tensile fracture of the 40 wppm H sample. This difference could be related to the change of deformation mechanism between the two samples. At ε=0.1 (under loading before fracture), it can be observed from EBSD (Fig.3cand Fig.4c) that 40 wppm H sample presented a high fraction (14.8 area%) of B2 islands whereas the microstructure of the 8 wppm H sample is composed by much less non-transformed B2 phase (3.2 area%).

Figure 7 :

 7 Figure 7: Plotting of the non-transformed B2 area fractions and the slope of the stress-strain curve as a function of tensile deformation strain of the 800°C-annealed samples. (a) Plots of the area fractions of non-transformed B2 phase of 8 wppm H sample (blue open square) and 40 wppm H sample (green open square) obtained from in-situ EBSD measurements. Non-linear fitting curves of every samples are shown in dash lines, respectively using the same color; (b) The slope of the stress-strain curve obtained from single tensile loading curves are plotted in dot-lines for 8 wppm H sample (blue) and 40 wppm H sample (green).

Figure 8 :

 8 Figure 8: Micrographs and stereo projection analysis of the TEM observation on the deformation microstructure of the 8 wppm H sample (800°C-annealed). The TEM foil is prepared from the sample after single tensile loading to ε=0.1. (a) Bright-filed (BF) image of the deformation microstructure. (b) and c) BF image of the dislocation contrasts in B2 phase. The axis of the Burgers vectors, noted b1, b2

Figure 9 :

 9 Figure 9: Micrographs of the TEM observation on the deformation microstructure of 40 wppm H sample (800°C-annealed). The TEM foil is prepared from the sample after single tensile loading to ε=0.1. (a) Bright-filed (BF) image of the deformation microstructure (out of zone axis), where a non-transformed B2 island can be seen (marked by yellow dash line). Dislocation characterization is performed at the squared area in the non-transformed B2 island; (b), c) and d) dislocation contrast imaging of b1, b2 and b3.

Fig. 10a presents

  Fig.10apresents the starting microstructure of 500°C-annealed NiTi. The partial recrystallization

Figure 10 :

 10 Figure 10: Micrographs of the TEM observation on the microstructure of 500°C-annealed samples before and after deformation to ε=0.06. (a) The BF imaging of the sample at as-500°C-annealed state. Two zones of different microstructures are observed. Zone A: the zone of subgrains of grain size around 100 nm; Zone B: the zone of dislocations and precipitates (marked by orange dash lines, the zone size is about 5 µm in diameter) without subgrain structure. The diffraction pattern (SA at red dash square) is shown in the inset indicating the presence of two Ti3Ni4 variants in Zone B. The DF image (in the inset) of Ti3Ni4 precipitates; (b) Zone A is magnified and tilted to dislocation contrast conditions g=0-11B2 and g=-110B2, no visible dislocations are observed; (c) Zone B is magnified and tilted to the same contrast conditions, dislocations (b1=a[001], b2= a[100] and b3=a[010]) can be seen; (d) BF image of the deformation microstructure of the 8 wppm H sample (500°C-annealed) after unloading from ε=0.1; (e) Dislocation contrast images at g=0-11 and g=-110 in Zone A; (f) Dislocation contrast images at g=0-11B2 and g=-110B2 in Zone B; (g) BF image of the deformation microstructure of the 40 wppm H sample (500°C-annealed) after unloading from ε=0.1; (h) Dislocation contrast images at g=0-11 and g=-110 in Zone A; (i) Dislocation contrast images at g=0-11B2 and g=-110B2 in Zone B. In (b), (e) and (h), zone of interest (ZOI) is marked by black dash

Figure 11 .

 11 Figure 11. Thermal desorption thermograms of the 500°C-annealed stent sample (Stent #4) and 40 wppm H hydrogenated 800°C-annealed NiTi samples. (a) 500°C-annealed stent sample with 33 wppm H after long-time aging at RT (> 1 month); (b) 40 wppm H 800°C-annealed sample after 24 hours aging at RT; (c) 40 wppm H 800°C-annealed sample after 192 hours aging at RT. The experimental data (black line) can be deconvoluted into Gaussian peaks (in blue). In red is the sum of all Gaussian peaks, indicative of correct fit.

  and 12), more specifically where low energy GBs are concentrated (correlation made in the black squares of Fig 12a and 12b). Nevertheless, the image correlation is insufficient to be conclusive to link the B2 islands, GBs and hydrogen redistribution. Therefore, the question keeps open if the formation of B2 island results from hydrogen redistribution, thus local enrichment, at certain GB or at other structural defects.

Figure 12 :

 12 Figure 12: The analysis of the grain orientation and grain boundary (GB) network of the 40 H wppm sample ε=0.1 (before fracture). (a) The distribution of Random GB (in blue, interface energy > 600 mJ/m2) and Special GB (in red, interface energy = 200-500 mJ/m 2 ); (b) The EBSD IPF map of the non-transformed B2 islands along tensile test axis (RD) of the same map of Fig. 6c. Same squares are marked in each image to help compare the positions of the islands and the nearby GBs. It can be seen

  ), observed under ex-situ conditions (after fracture or after unloading the strained sample) might imply a local change of deformation mechanisms between 8 wppm and 40 wppm hydrogen. Further investigations with suitable in-situ techniques (e.x. Digital Image Correlation DIC) are required to clarify the mesoscopic strain-stress concentration and potentially the change of deformation mechanism of H-charged NiTi with partial recrystallization structure. Nonetheless, the impact of 40 wppm H to the 500°C-annealed NiTi leads

hydrogen concentrations, 8

 8 wppm and 40 wppm, are therefore studied in both 500°C-annealed NiTi and 800°C-annealed NiTi to illustrate the hydrogen influences to the mechanical property and deformation mechanism of NiTi. The results present that, after aging 192 hours at room temperature, the presence of 40 wppm hydrogen reduces drastically the ductility of NiTi with an observable increasing of onset stress of SIM in both 500°C-annealed NiTi and 800°C-annealed NiTi. In the 8 wppm H sample, the mechanical property does not exhibit significant changes. In-situ EBSD coupled with tensile loading and ex-situ TEM observations detect the local changes of operating deformation mechanisms. In the 800°C-annleaed samples, 40 wppm H leads to interruption of MB propagation and the formation of non-transformed B2 islands (14.8 area%) during the superelastic deformation.In the 500°C-annealed sample, 40 wppm H leads to extensive dislocation activity at the microstructural zones containing dislocations and Ti3Ni4 precipitates. TDS confirms that a hydrogen redistribution has happened in the 40 wppm hydrogen during the homogenization process. The mesoscale MB propagation, the formation of the non-transformed B2 islands and the microscale change of deformation mechanism are discussed in relation to the redistribution of the hydrogen. For stent safety, extensive attentions should be payed to the low-concentration hydrogen absorption during stent manufacturing in order to minimize the risk of hydrogen related issues (high radial stress, strut fracture during crimping, incomplete self-expansion, etc.).

Table 1 :

 1 The hydrogen concentrations of current commercial stents in Europe.

	Product	Stent 1	Stent 2	Stent 3	Stent 4	Stent 5	Stent 6	Stent 7	Stent 8	Stent 9
	H (wppm)	15± 1.5	20± 2.0	14± 1.0	33± 3.0	35± 3.5	19± 2.0	7± 0.5	55± 5.5	6± 0.5

Acknowledgement

The authors acknowledge the research collaboration between the team of Métallurgie Structurale IRCP and the team of Laboratoire d'Etude de la Corrosion Aqueuse, CEA Saclay, supported by the Métallurgie Structurale IRCP. The authors acknowledge the great help from Dr. Domitille GIAUME in the collaboration and the discussions with Dr. Philippe VERMAUT about the crystallography of martensitic phase transformation in the shape memory alloys and the techniques for TEM observations.


S Konate

C I Lo

E Kuete

M Sarr

N Amsrtong

A Levasseur

A Caputo

M A Thera

D Raoult

M Million

Tiago P Ribeiro

Fernando J Monteiro

Marta S Laranjeira

Dr Marta De

Sousa Laranjeira

a new bacterium isolated from stool of healthy children living in Mali

Keywords: Biocompatibility, Drug loading, Magnetic hydroxyapatite, MRI contrast agents, PEG, Surface modification

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

Introduction

Magnetic resonance imaging (MRI) is a very powerful tool for tumour detection, presenting a good accuracy to safety ratio. Compared with other techniques, it has higher accuracy than ultrasounds, while being safer than computer tomography since it uses electromagnetic radiation instead of ionizing radiation.

Interestingly enough, for increased accuracy, these imaging techniques take advantage of certain compounds that increase contrast, known as contrast agents. In MRI, the most common contrast agent is gadoteric acid, which is a complex between a chelating agent and gadolinium(III), being heavily used for shortening the T1 relaxation time [START_REF] Wahsner | Chemistry of MRI contrast agents: Current challenges and new frontiers[END_REF]. However, despite being commonly and extensively used, gadolinium is not naturally found in the human body, leading to high levels of toxicity, especially in patients with poor kidney function that can develop a fatal condition known as nephrogenic system fibrosis [START_REF] Semelka | Influence of excess ligand on Nephrogenic Systemic Fibrosis associated with nonionic, linear gadolinium-based contrast agents[END_REF]. Knowing this, researchers shifted towards alternatives, such as magnetic iron-based nanoparticles (NPs). These NPs are good substitutes, since iron is commonly found in the human body and the small size of the particles allowed an easy administration [START_REF] Wei | Exceedingly small iron oxide nanoparticles as positive MRI contrast agents[END_REF]. One example of these are ferrite NPs, which are superparamagnetic oxides and can reduce the T2 relaxation time, with products like Feridex® and Combidex® being already developed [START_REF] Miller | Prediction of anti-cancer nanotherapy efficacy by imaging[END_REF]. To show this, Weng et al. made a comparison between iron oxide nanoclusters and gadolinium based contrast agents, reporting an increased biosafety profile of the iron oxide materials [START_REF] Weng | Toxicological Risk Assessments of Iron Oxide Nanoclusterand Gadolinium-Based T1MRI Contrast Agents in Renal Failure Rats[END_REF]. Yet, despite iron being safer than gadolinium, the high concentration of iron and the low biodegradability of these magnetic NPs make them also connected with bioaccumulation and toxicity [START_REF] Ittrich | Superparamagnetic iron oxide nanoparticles in biomedicine: Applications and developments in diagnostics and therapy, RoFo Fortschritte Auf Dem Gebiet Der Rontgenstrahlen Und Der Bildgeb[END_REF]. To counteract these problems, magnetic hydroxyapatite (HAp) NPs are proposed as a solution due to the high biocompatibility and biodegradability of HAp [START_REF] Barros | Alginatenanohydroxyapatite hydrogel system: Optimizing the formulation for enhanced bone regeneration[END_REF][START_REF] Turon | Biodegradable and Biocompatible Systems Based on Hydroxyapatite Nanoparticles[END_REF]. Several studies have already proved that iron introduction into the structure of hydroxyapatite provides magnetic properties to its usually non-magnetic behaviour [START_REF] Ribeiro | Duality of iron (III) doped nano hydroxyapatite in triple negative breast cancer monitoring and as a drug-free therapeutic agent[END_REF][START_REF] Tampieri | Intrinsic magnetism and hyperthermia in bioactive Fe-doped hydroxyapatite[END_REF][START_REF] Mercado | Paramagnetic iron-doped hydroxyapatite nanoparticles with improved metal sorption properties. A bioorganic substrates-mediated synthesis[END_REF][START_REF] Gloria | Magnetic poly( -caprolactone)/iron-doped hydroxyapatite nanocomposite substrates for advanced bone tissue engineering[END_REF]. Moreover, our previous studies showed that iron doped hydroxyapatite (FeHAp) NPs can be used as contrast agents in MRI [START_REF] Ribeiro | Duality of iron (III) doped nano hydroxyapatite in triple negative breast cancer monitoring and as a drug-free therapeutic agent[END_REF][START_REF] Laranjeira | Different hydroxyapatite magnetic nanoparticles for medical imaging: Its effects on hemostatic, hemolytic activity and cellular cytotoxicity[END_REF]. Therefore, combining the magnetic behaviour, the low iron concentration and the degradability of these NPs, they should be adequate to be used as safer MRI contrast agents.

Nevertheless, these nanosystems are usually intravenously administrated. In the blood system, they must avoid deposition and aggregation, while being hemo-and cytocompatible for an appropriate mode of action. In a previous study of ours, FeHAp NPs did not induce hemolysis, however they were highly thrombogenic, which is a critical factor in intravenous administrated systems [START_REF] Laranjeira | Different hydroxyapatite magnetic nanoparticles for medical imaging: Its effects on hemostatic, hemolytic activity and cellular cytotoxicity[END_REF]. Thrombus formations lead to clogged blood vessels and consequently, surface modification of NPs with biocompatible polymers has been indicated for improved colloidal stability, hemocompatibility and lower toxicity. As an example, a study made by Pinto et al. showed that the adsorption of polymers such as Polyethylene Glycol (PEG) and Polyvinyl Alcohol (PVA) onto the surface of graphene platelets highly increased their cytocompatibility and hemocompatibility [START_REF] Pinto | Polymer surface adsorption as a strategy to improve the biocompatibility of graphene nanoplatelets[END_REF]. In fact, PEG is one of the most studied biocompatible polymers and due to its hydrophilic characteristic, it enables the modification of NPs in order to prevent agglomeration, avoid protein adsorption and improve biocompatibility. With this thought, Lankveld et al.

compared bare gold NPs with PEGylated ones, concluding that the PEG modification induced higher circulation times and made the NPs less prone to be uptake by the reticuloendothelial system, increasing the NPs half-life inside the organism [START_REF] Lankveld | Blood clearance and tissue distribution of PEGylated and non-PEGylated gold nanorods after intravenous administration in rats[END_REF].

PEG can be covalently attached to some NPs or can also be non-covalently added. This last alternative creates a matrix around the NPs which can act as a drug carrier as proven by León et al. [START_REF] León | FTIR and raman characterization of TiO2 nanoparticles coated with polyethylene glycol as carrier for 2methoxyestradiol[END_REF]. Such evidence is of major importance, since one of the main causes for chemotherapy low success rates is its systemic action, versus which drug carriers purpose an enormous advantage. Drug carriers limit the region of action, by transporting their contents to a specific region such as the tumour microenvironment. Such feature enables not only a more efficient profile for the treatment but also the administration of lower doses, ultimately leading to lower economic costs and systemic toxicity.

With this being explained, this work proposes the modification of FeHAp NPs surface with PEG using PEG2000 and PEG8000, in order to avoid the thrombogenecity found in a previous work [START_REF] Laranjeira | Different hydroxyapatite magnetic nanoparticles for medical imaging: Its effects on hemostatic, hemolytic activity and cellular cytotoxicity[END_REF]. The composites were also put into contact with fibroblasts for cytocompatibility assessment, and a chemotherapeutical drug was introduced in order to create a drug carrier.

Materials and Methods

Nanoparticles synthesis and PEGylation

FeHAp NPs detailed production is described elsewhere [START_REF] Ribeiro | Duality of iron (III) doped nano hydroxyapatite in triple negative breast cancer monitoring and as a drug-free therapeutic agent[END_REF]. Briefly, solutions of iron (III) nitrate (98+%, Acros Organics, USA) and orthophosphoric acid (85 wt% solution in water, Acros Organics, USA) were added dropwise into a previously heated to 100 °C, calcium hydroxide (98% extra pure, Acros Organics, USA) solution. The pH was continuously measured and kept at 10 by adding ammonium hydroxide (25% solution in water, Merck, Germany) as needed. After 2h of reaction, the particles were washed by centrifugation with ultrapure water, dried at 60 °C and powdered. Finally, the NPs were sintered for 1h at 1000 °C with a heating rate of 5 °C min -1 and stored in a desiccator for further use. PEGylation of the NPs was performed in a non-covalent manner. With this purpose, 1% PEG2000 (Sigma-Aldrich, USA) and PEG8000 (Sigma-Aldrich, USA) solutions (w/v) were prepared separately, to which FeHAp NPs were added in a 1:1 ratio. The mixture was stirred for 7h at room temperature, further washed by centrifugation at 10 °C, dried under vacuum at room temperature and finally stored at 4 °C.

The PEG coated NPs were denominated PEG2 and PEG8 for PEG2000 and PEG8000, respectively.

Physicochemical characterization of FeHAp and PEGylated NPs

Morphological, atomic composition, crystallographic studies and iron content of FeHAp NPs were done by transmission electron microscopy (TEM), energy-dispersive X-ray spectroscopy (EDX), X-ray diffraction (XRD) and Inductively Coupled Plasma-Atomic Emission Spectrometry (ICP-AES), respectively. The results were already published and can be observed in the supplementary data on Fig. S1, Fig. S2 and Table S1 [START_REF] Ribeiro | Duality of iron (III) doped nano hydroxyapatite in triple negative breast cancer monitoring and as a drug-free therapeutic agent[END_REF].

In order to confirm the formation of a PEG coating, Fourier Transform Infrared (FTIR) spectroscopy analysis was carried out. Infrared spectra were recorded in a FTIR spectrometer (Perkin-Elmer, USA), with a 4 cm -1 resolution and after an averaging of 32 scans per sample. Before analysis, pellets of KBr were prepared by mixing 1 mg of the NPs with 200 mg of potassium bromide and compressed in a hydraulic press (Grasedy Specac, USA). Pellets were produced in a low humidity environment for improved sample quality.

Thermogravimetric analysis

The thermal behaviour of the FeHAp and PEGylated powders was studied by thermal gravimetric analysis (TGA), using a NETZSCH STA 449F3 (NETZSCH, Germany). The thermal cycle, under nitrogen atmosphere, started with a one-minute stage at 25ºC followed by a 10 °C min -1 slope up to 1000 °C.

Size distribution and Zeta Potential analysis

Particle size, size distribution and zeta potential measurements were determined through dynamic light scattering (DLS) in a Zetasizer Nano ZS (MALVERN, UK) with a scattering angle of 173° and a He-Ne 630 nm wavelength laser, using a Fraunhofer's model. For this, the dried particles were dispersed in ultrapure water at a concentration of 1 mg mL -1 . Afterwards, the suspension was diluted 10 times in ultrapure water before analysis. Three measurements were done per sample at 25 °C.

Magnetic resonance imaging

MRI experiments were performed with a 3T clinical scanner (Sigma HDXT 3T, GE Health Care, USA). T2 contrast images were recorded by Sag T2 Fast Spin Eco technique with TR=2500, TE=77,159, FOV=240 mm, eco train of 120, slices of 1 mm and a band width of 244 Hz. Samples were prepared by dispersing the NPs in ultrapure water at a concentration of 500 µg mL -1 . A water filled well was used as the control group for qualitative analysis.

Biological evaluation

Haemolysis assay

Red blood cells were isolated from buffy coats (samples were obtained from Immuno-hemotherapy Service, Hospital S. João, Porto, Portugal). Buffy coats were centrifuged at 4°C at 1500 rpm. The red bottom layer was isolated and washed three times with PBS. The purified red blood cells were diluted to a cell density of 2.0 x 10 8 cells mL -1 . 100 µL of the red blood cell suspension were added to a 96-well plate, to which 100 µL of the NPs suspension were added. Incubation was carried for 3h at 37 °C at 80 rpm on an orbital shaker. Afterwards, the suspensions were centrifuged at 1500 rpm for 15 minutes to collect the supernatant. Finally, 80 µL were transferred to a 96-well black plate and absorbance was read at 540 nm in a SynergyMX fluorimeter (BioTek, USA). The haemolysis percentage was calculated as follows:

Haemolysis (%) = (Hb value of sample/Hb of 100% haemolysis with 1% Triton X-100) x 100. The tested NPs concentrations were 125, 250 and 500 µg mL -1 . Controls with PBS (negative control) and 1% Triton X-100 (Sigma Aldrich, USA) in PBS (positive control for 100% haemolysis) were performed. All assays were performed in triplicate.

Thrombogenicity

Thrombogenicity of the NPs was evaluated using the whole blood kinetic clotting time method. FeHAp and PEGylated NPs, as well as positive (glass) and negative (polypropylene) controls, were used in a concentration of 4000 µg mL -1 . Clotting reaction was activated with the addition of 100 µL of CaCl2 (0.1 M) to each 1 mL of whole blood. Afterwards, 100 µL of the activated blood was put in contact with controls and NPs for 5, 15, 25, 35 and 45 min. At the end of each time point, 3 mL of distilled water was added and incubated for 5 min. The supernatant absorbance was measured at 540 nm, in a SynergyMX fluorimeter (Biotek, USA). Red blood cells that were not trapped in a thrombus were lysed with addition of distilled water (which causes complete haemolysis by hypoosmotic shock), thereby releasing haemoglobin into water for subsequent measurement. The clot size is inversely proportional to the absorbance value.

Cytocompatibility

Human gingival fibroblasts (HGF, ATCC, USA) were cultured in Alpha Modified Eagle Medium (α-MEM, Gibco, Germany) supplemented with 10% (v/v) fetal bovine serum (FBS, Gibco, Germany) and 1% (v/v) penicillin/streptomycin (Biowest, France) and incubated at 37 °C and at 5% CO2. Trypsinization and seeding was performed every three days.

The metabolic activity of HGF, in contact with the PEGylated and bare FeHAp NPs, was assessed by the resazurin assay. Briefly, metabolically active cells can reduce the blue compound resazurin into resorufin, turning it into a pink-purple appearance that can be detected by spectrophotometry [START_REF] Coelho | Heparinized nanohydroxyapatite/collagen granules for controlled release of vancomycin[END_REF]. For this assay, the cells were trypsinized and seeded onto 96-well plates at a density of 5000 cells per well and incubated for 24h. In parallel, the NPs were sterilized in ethanol, washed with sterile PBS and then re-dispersed in complete supplemented α-MEM. After 24h of incubation, regular medium was exchanged for the nanoparticle-rich medium, at concentrations of 125, 250 and 500 μg mL -1 , and the cells were again incubated at 37 °C and at 5% CO2. After 24h and 48h of incubation, each well was loaded with resazurin (10% in α-MEM) and re-incubated for 4h. The medium was then transferred to a black 96-well plate and fluorescence was measured at excitation and emission wavelengths of 566 and 605nm, respectively, in a SynergyMX fluorimeter (BioTek, USA).

Drug loading

In this step all solutions were prepared in PBS pH=7. [START_REF] Miller | Prediction of anti-cancer nanotherapy efficacy by imaging[END_REF]. A 1 mg mL -1 NPs suspension was added to 2 mL of a 1 mg mL -1 doxorubicin (98.0-102.0%, Sigma-Aldrich, Germany) solution and stirred at room temperature in dark conditions for 24h. The suspension was then centrifuged (4000 rpm, 5 min) and the supernatant and precipitate were separated. With this, the amount of drug loaded was determined as follows: Drug loaded (%) = ((total amount of drug -drug in the supernatant) / total amount of drug) x 100.

Statistical analysis

Statistical analysis and graphs were carried out with the GraphPad Prism 5 software (GraphPad, USA).

Data are represented as mean ± standard deviation (SD). Statistical analysis was performed using the analysis of variance One-way ANOVA, except for the thrombogenicity studies were a Two-way ANOVA was employed, followed by Tukey's multiple comparison test, with n=3. Statistical significance was considered when p<0.05 and represented with * for p<0.05, ** for p<0.01 and *** for p<0.001.

Results

Physicochemical characterization

The FTIR spectra, presented on Fig. 1A, shows that all samples comprised the characteristic peaks of HAp. Peaks corresponding to OH -groups can be found at ~3572 cm -1 and at 631 cm -1 . A broad band and a noticeable peak associated with adsorbed water are found at ~3436 cm -1 and ~1630 cm -1 . Additionally, residual carbonates incorporated from atmospheric CO2 can be found at ~877 cm -1 . Lastly, all four modes of PO4 3-can be detected at ~960 cm -1 , ~475 cm -1 ; ~1044 cm -1 , ~1090 cm -1 ; ~570 cm -1 and ~605 cm -1 .

However, regarding the PEGylated NPs, PEG specific peaks can be detected at 845 cm -1 , 1345 cm -1 and 2888 cm -1 , which correspond to C-C, -CH and C-O stretching, respectively. On Fig. 1B is presented the thermogravimetric profile of the NPs. The presence of the polymer was confirmed due to the formation of a slope at ~300 °C that does not occur on the bare NPs, being attributed to polymer degradation. On the other hand, pristine FeHAp NPs presented little to no mass variation (~5% weight loss). Both the modified NPs lost, respectively for PEG2 and PEG8, 8% and 10% until 1000 °C. Taking into the account the mass loss of pristine FeHAp Nps up to that temperature, it can be concluded that the total amount of polymer at the NPs surface was 3% and 5%, respectively for PEG2 and PEG8. Fig. 1C shows the particle size distribution of the NPs in water. For the bare FeHAp NPs three distinct populations were detected, one of them with sizes above 4000 nm which corresponds to particle aggregation. Dispersion after coating with PEG, drastically influenced the dispersion profile, creating one single population for each type of composite. Moreover, particles hydrodynamic size tended to increase when comparing PEG2 (approx. 435 nm) to PEG8 (approx. 875 nm).

Regarding zeta potential analysis (Fig. 1D), the samples presented surface charges of -18.2mV ± 0.22, -16.37mV ± 0.49, -14.63mV ± 0.45 for FeHAp, PEG2 and PEG8, respectively. The variation in surface charge also indicates the presence of adsorbed PEG.

MRI contrast potential

In Fig. 2 a T2-weighted MR image of the bare and PEGylated NPs dispersed in water may be observed. A brightness reduction was observed in comparison with water, clearly indicating the potential of these materials to serve as MRI contrast agents. Interestingly, PEGylation did not negatively impact the contrast potential of the bare FeHAp NPs, however the rapid deposition of the bare NPs induced a small decrease in the contrast potential when compared with the PEGylated ones.

Hemocompatibility analysis

As seen in Fig. 3A, all samples presented high hemocompatibility since the haemolysis percentage (< 2%) of the test groups was very comparable with the negative control (p>0.05). Additionally, no significant (p>0.05) differences were observable between PEGylated and non-PEGylated NPs.

Thrombogenecity evaluation

Fig. 3B presents the kinetics of the blood clotting process. Thrombus formation was evaluated for the PEGylated and bare NPs. In this assay, the size of the blood clot was inversely proportional to the absorbance level. Red blood cells trapped inside the thrombus are not lysed with the addition of distilled water (positive control), while the opposite occurs when no thrombus is formed (negative control). From the test, it was possible to confirm that FeHAp Nps were highly thrombogenic, close to the positive control (p>0.05). However, the PEGylated samples corrected the problem, leading to less thrombus formation, being closer to the negative control (p>0.05). Additionally, significant differences (Fig. 3C) were detected between the bare and the PEGylated samples, clearly indicating the influence of PEG in avoiding thrombus formations. Still, no significant differences were obtained when comparing PEG2 and PEG8 samples, suggesting that the chain length is not a factor during thrombus formation.

Cytocompatibility evaluation

The cytocompatibility of the produced NPs was confirmed with the resazurin assay. From Fig. 4, it was possible to conclude that PEGylation increased the cytocompatibility of the NPs. In fact, after 24h the bare NPs were biocompatible in comparison with the control group (p>0.05), however this profile changed after 48h (p<0.001). It was also at this time point that PEGylation proved to be efficient in reducing the toxicity, since the metabolic activity stayed high and very comparable to the control group (p>0.05). Still, no significant differences (p>0.05) were detected between the different PEGylations with PEG2 and PEG8.

Drug loading

Fig. 5 shows the amount of loaded drug in the different NPs. The bare FeHAp NPs were able to carry a small amount of drug (1.96% ± 0.76) however, the PEGylated NPs could carry a significant (p<0.001) amount (12.93% ± 5.14) of drug with PEG8 being more efficient (30.56% ± 4.37) than PEG2 (p<0.01).

Discussion

To the best of our knowledge, the combination of FeHAp with PEG was never explored to date for in vitro biological effects evaluation. The presence of the polymer at the surface of the FeHAp NPs was confirmed through FTIR and quantified through TGA. FTIR is a technique that characterizes the functional groups at the surface of materials and in the specific case of hydroxyapatite derived NPs, the functional groups that can be detected are phosphate and hydroxyl groups, however with the presence of PEG carbon associated groups, it was possible to confirm the presence of PEG at the surface of FeHAp NPs [START_REF] Coelho | Antibacterial bone substitute of hydroxyapatite and magnesium oxide to prevent dental and orthopaedic infections[END_REF]. Moreover, TGA is a useful technique to determine the presence of residual components in NPs.

This method is based on the observation of a mass loss in response to temperature increase. Since the FeHAp NPs were sintered at 1000 °C no significant mass loss was detected, however on the PEGylated NPs the abrupt mass loss at 300 °C was attributed to PEG degradation, confirming the presence of a PEG layer on the NPs [START_REF] Parab | Biocompatible transferrin-conjugated sodium hexametaphosphate-stabilized gold nanoparticles: Synthesis, characterization, cytotoxicity and cellular uptake[END_REF]. Comparing both PEG2 and PEG8 samples, the latest presented a higher mass loss, which was expected due to the higher molecular weight of PEG 8000. In addition, the DLS analysis indirectly confirmed the presence of the polymer due to zeta potential and size distribution variation.

Regarding the zeta potential analysis, FeHAp NPs usually present a negative surface charge at neutral pH, due to the abundance of hydroxyl groups, which are also present in the terminal ends of PEG [START_REF] Iafisco | Magnetic bioactive and biodegradable hollow fe-doped hydroxyapatite coated poly(l -lactic) acid micro-nanospheres[END_REF].

Despite the negative zeta potential of the PEGylated samples, the value tended to be less negative with PEG chain size increase. We hypothesize that due to the non-covalent bonding of PEG, it masked the negative charge of the bare NPs. In fact, Jiang et al. evaluated the zeta potential of chromium-doped zinc gallate NPs before and after grafting PEG to the surface, concluding that the PEG modified NPs presented a less negative zeta potential than the bare ones [START_REF] Jiang | Hemocompatibility investigation and improvement of near-infrared persistent luminescent nanoparticle ZnGa2O4:Cr3+ by surface PEGylation[END_REF]. Moreover, as mentioned in section 1, NPs must avoid aggregation and as confirmed by the size distribution, bare FeHAp NPs tend to form aggregates (>4000 nm). Visually, after dispersion in water through sonication, the bare NPs formed visible aggregates and tended to rapidly deposit at the bottom of the flasks. On the other hand, the PEGylated samples formed a more stable suspension, being justified by the size distribution, which revealed a single population of NPs. Such evidence, confirmed the potential of PEG in reducing the interactions between NPs, and therefore avoid their aggregation and further deposition.

Despite being used for increased biocompatibility of magnetic NPs, polymer coating can alter their magnetic response. FeHAp NPs are known for presenting a superparamagnetic behaviour which is a required characteristic for MRI contrast [START_REF] Ribeiro | Duality of iron (III) doped nano hydroxyapatite in triple negative breast cancer monitoring and as a drug-free therapeutic agent[END_REF]. From the obtained results their contrast potential is confirmed due to the decreasing brightness. Additionally, the PEG coating provided a more uniform contrast signal due to the increased stability of the NPs suspension, while the FeHAp NPs tended to rapidly form aggregates and deposit, leading to an uneven signal.

Haemolysis in the presence of bare and PEGylated FeHAp NPs was very similar to the negative control, proving to be non-haemolytic. The potential of NPs to harm erythrocytes derives from their physicochemical characteristics (e.g. chemical composition, size, shape, surface charge). NPs with sharper edges are more prone to damage erythrocytes, however the produced NPs have a slightly round shape and the low haemolytic effects can derive from that [START_REF] Ribeiro | Duality of iron (III) doped nano hydroxyapatite in triple negative breast cancer monitoring and as a drug-free therapeutic agent[END_REF][START_REF] Pinto | Smaller particle size and higher oxidation improves biocompatibility of graphenebased materials[END_REF]. Moreover, PEGylated NPs presented very similar results (p>0.05) to bare NPs, which was expected since PEG is known for increasing the hemocompatibility of NPs, as shown by Ruan et al., and therefore should not have any negative contribution [START_REF] Ruan | PEGylated fluorescent carbon nanoparticles for noninvasive heart imaging[END_REF]. However, due to the lack of hemocompatibility studies on FeHAp NPs it is very difficult to make comparisons, yet Yildirim et al. studied the hemocompatibility of mesoporous silica NPs and obtained very similar results regarding the haemolytical activity. In that study, it was reported that there is no correlation between surface charge and haemolysis, being the exposed functional hydroxyl groups responsible for interacting with red blood cells membrane and induce damage [START_REF] Yildirim | Impact of mesoporous silica nanoparticle surface functionality on hemolytic activity, thrombogenicity and non-specific protein adsorption[END_REF]. In the case of FeHAP NPs, there are also hydroxyl groups at the surface, meaning that FeHAp NPs should prove haemolytical, However, these amount of exposed hydroxyl groups may not be sufficient to induce cell membrane damage. In fact, Lin et al. reported that mesoporous silica NPs were significantly less haemolytical than non-porous silica NPs, justifying this occurrence with the amount of exposed hydroxyl groups [START_REF] Lin | CH2-Hemolysis and size different MSN[END_REF].

These blood interaction results mimic those obtained in a previous work, where the main problem was thrombus formation [START_REF] Laranjeira | Different hydroxyapatite magnetic nanoparticles for medical imaging: Its effects on hemostatic, hemolytic activity and cellular cytotoxicity[END_REF]. Fortunately, from the obtained results, PEGylation solved this problem.

Thrombus formation is the last product from the activation of the coagulation cascade [START_REF] Lippi | Arterial thrombus formation in cardiovascular disease[END_REF]. The high thrombogenicity of FeHAp NPs is directly linked to the trivalent characteristic of Fe 3+ ions. These ions easily react with hydroxyl groups of water and produce powerful hydroxyl radicals. Furthermore, these radicals are the most biologically damaging ones and when in contact with blood increase the formation of blood clots, deriving from the conversion of soluble fibrinogen into fibrin-like polymers [START_REF] Lipinski | Interaction of Fibrin with Red Blood Cells: The Role of Iron[END_REF]. It is possible that Fe 3+ ions are released from the NPs and that those are sufficient to start the clotting process. However, PEG serves as a protective layer around the NPs, ultimately leading to less interaction with biological fluids and therefore reducing the release of iron ions, which could explain the nonthrombogenic behaviour of the PEGylated NPs in comparison with their bare counterparts [START_REF] Cauda | Bio-degradation study of colloidal mesoporous silica nanoparticles: Effect of surface functionalization with organo-silanes and poly(ethylene glycol)[END_REF]. It is also important to notice that in this study a high concentration of NPs (4 mg mL -1 ) was tested, being higher than the usual dose of intravenously administrated contrast agents (50 mg Kg -1 ) which roughly correspond to 0.5 mg mL -1 , clearly indicating their capability for being intravenously administrated [START_REF] Yildirim | Impact of mesoporous silica nanoparticle surface functionality on hemolytic activity, thrombogenicity and non-specific protein adsorption[END_REF].

Additionally, due to Fe 3+ and Ca 2+ ions release, FeHAp NPs proved to be toxic to cells after 48h. It is known that Ca 2+ ions are released from hydroxyapatite NPs during the degradation process that occurs both intracellularly and extracellularly [START_REF] Meena | Effects of hydroxyapatite nanoparticles on proliferation and apoptosis of human breast cancer cells (MCF-7)[END_REF]. In the case of FeHAp, both Fe 3+ and Ca 2+ ions are expected to be released, leading to ion imbalances than can disrupt cellular mechanisms [START_REF] Ribeiro | Duality of iron (III) doped nano hydroxyapatite in triple negative breast cancer monitoring and as a drug-free therapeutic agent[END_REF]. Ca 2+ excess is usually associated with mitochondrial damage and activation of an apoptotic response in cancer cells, however on healthy cells such overload is quickly counterbalanced in order to avoid toxicity [START_REF] Meena | Effects of hydroxyapatite nanoparticles on proliferation and apoptosis of human breast cancer cells (MCF-7)[END_REF][START_REF] Han | Different inhibitory effect and mechanism of hydroxyapatite nanoparticles on normal cells and cancer cells in vitro and in vivo[END_REF]. On the other hand, Fe 3+ is related with reactive oxygen species (ROS) production usually via the Fenton reaction, which leads to the formation of hydroxyl radicals [START_REF] Shen | Fenton-Reaction-Acceleratable Magnetic Nanoparticles for Ferroptosis Therapy of Orthotopic Brain Tumors[END_REF]. From the obtained results, since bare FeHAp NPs did not induce toxic effects after 24h, it was hypothesized that it is derived from the low degree of degradation and therefore low ions release. However, after 48h a significant decrease in the metabolic activity was detected, showing the toxic profile of the bare FeHAp NPs, which is probably derived from the release of Fe 3+ ions. However, the PEGylated NPs did not impact the cell metabolic activity at 48h of exposure, which could be explained by lower cellular internalization and delayed degradation, since PEG produces a protective layer around the NPs, limiting their direct contact with biological fluids [START_REF] Cauda | Bio-degradation study of colloidal mesoporous silica nanoparticles: Effect of surface functionalization with organo-silanes and poly(ethylene glycol)[END_REF].

Finally, regarding the drug loading experiments, PEGylation induced higher drug uptake. Hydroxyapatite NPs have already been explored as drug carriers, especially due to their high biocompatibility [START_REF] Thomas | Synergistic anticancer efficacy of Bendamustine Hydrochloride loaded bioactive Hydroxyapatite nanoparticles: In-vitro, ex-vivo and in-vivo evaluation[END_REF][START_REF] Zhang | Preparation and characterization of hydroxyapatite nanoparticles carrying insulin and gallic acid for insulin oral delivery, Nanomedicine Nanotechnology[END_REF]. In this case, it was observed a slight level of drug adsorption on the bare FeHAp NPs was observed, probably due to interactions between the negative surface charge of the NPs and the positively-charged amine moiety of doxorubicin. More interestingly, due to the polymer matrix on the PEGylated samples, the loaded drug percentage increased significantly, which may be explained through charge interactions and hydrogen bonding. Additionally, when comparing both PEGylated samples, PEG8 showed increased drug loading, which may be explained by the longer chains of PEG, that contributed to the matrix size increase.

Conclusions

This work describes an easy solution for increasing the biocompatibility of FeHAp NPs, while creating a drug carrier for cytostatic drugs with potential application in some cancers. The PEG modified surfaces successfully decreased the NPs agglomeration, avoided their high thrombogenic profile up to 4 mg mL -1 , increased their cytocompatibility up to 48h with concentrations as high as 500 µg mL -1 and even increased the amount of loaded drug. It is important to emphasize that between PEG2000 and PEG8000 no significant differences were detected on the biologic assays, meaning that the choice between both polymers only varies depending on the amount of drug that is intended to be loaded. Therefore, PEGylation is in fact an easy method to provide all the necessary conditions for FeHAp NPs to be used as a blood administered MRI contrast agent that could also carry a drug for a complementary therapeutic effect. Additionally, due to the simplicity of the proposed PEGylation method, it can be expanded to several other types of NPs, increasing their biological applicability.

Fig. 1 .

 1 Fig. 1. Physicochemical characterization of the NPs. A) FTIR spectra of bare and PEGylated FeHAp NPs, a-FeHAp, b-PEG2, c-PEG8; B) Thermogravimetric analysis at and heating rate of 10 °C min -1 , a-FeHAp, b-PEG2, c-PEG8; C) Particle size distribution in water; D) Zeta potential analysis in water with the data being presented as mean ± standard deviation (n=3).

Fig. 2 .

 2 Fig. 2. T2-weighted MR images and the colour coded versions of bare and PEGylated FeHAp NPs at 500 µg mL -1 . The image of water was taken together for the purpose of comparison. In the colour coded image, high and low water intensities were indicated in red and blue, respectively.

Fig. 3 .

 3 Fig. 3. Hemocompatibility assays. A) Hemolysis of red blood cells after 3h of incubation at 37 °C with PBS as the negative control, Triton X-100 as the positive control and several NP concentrations (125, 250 and 500 µg mL -1 ). B) Measurement of whole blood kinetic clotting time for bare and PEGylated FeHAp NPs at a concentration of 4 mg mL -1 . The results are presented as mean ± standard deviation (n=3). C) Statistical analysis matrix of the thrombogenicity assay. * corresponds p<0.05, ** to p<0.01 and *** to p<0.001 while NS means non-significant.

Fig. 4 .

 4 Fig. 4. Metabolic activity of HGF cells after 24h and 48h of incubation with bare and PEGylated FeHAp NPs at concentrations of 125, 250 and 500 µg mL -1 . *** correspond to p<0.001 in comparison with the control group. Data are presented as mean ± standard deviation (n=3).

Fig. 5 .

 5 Fig. 5. Loaded drug into bare and PEGylated FeHAp NPs. ** corresponds to p<0.01 and *** to p<0.001. Data are presented as mean ± standard deviation (n=3).

Acknowledgements

The authors acknowledge project Biotherapies (NORTE-01-0145-FEDER-000012), supported by Norte Portugal Regional Operational Program (NORTE2020), under the PORTUGAL 2020 Partnership Agreement, through the European Regional Development Fund (ERDF). This work was also financed by FEDER -Fundo Europeu de Desenvolvimento Regional funds through the COMPETE 2020 -Operational Programme for Competitiveness and Internationalization (POCI), PORTUGAL 2020, in the framework of the project "Institute for Research and Innovation in Health Sciences"(POCI-01-0145-FEDER-00727). The authors also acknowledge the support of the i3S Scientific Platform Biointerfaces and Nanotechnology (UID/BIM/04293/2019). Additionally, the authors acknowledge the collaboration of João A. M. Santos and Pedro Conde from the Medical Physics, Radiobiology and Radiation Protection Group, IPO Porto Research Center (CI-IPOP), Portuguese Oncology Institute of Porto (IPO Porto), Porto, Portugal, during MRI image acquisition.